NASA/CP--2001-210948

Proceedings of the Sixth Applied Diamond

Conference/Second Frontier Carbon

Technology Joint Conference

(ADC/FCT 2001)

July 2001
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NASA/CP--2001-210948

Proceedings of the Sixth Applied Diamond

Conference/Second Frontier Carbon

Technology Joint Conference

(ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa,

Y. Koga, K. Kobashi, and G.A.J. Amaratunga, editors

Proceedings of a conference held at Auburn University

cosponsored by Auburn University,

Alabama Microelectronics Science and Technology Center,

Space Power Institute,

Alabama Space Grant Consortium,

NASA Glenn Research Center,

Japan New Diamond Forum, and

Japan Fine Ceramics Center

Auburn, Alabama

August 6-10, 2001

National Aeronautics and

Space Administration

Glenn Research Center

July 2001
Contents were reproduced from the best available copy

as provided by the authors.

Trade names or manufacturers' names are used in this report for

identification only. This usage does not constitute an official

endorsement, either expressed or implied, by the National

Aeronautics and Space Administration.

NASA Center for Aerospace Information

7121 Standard Drive

Hanover, MD 21076

Available from

National Technical Information Service

5285 Port Royal Road

Springfield, VA 22100

ISBN 0-9710327-0-X Available electronically at bttpb_!gltrs,_c,nas_.govd'GLTRS
FOREWORD

The Applied Diamond Conference (ADC) was established in 1991 to promote practical applications

and to enable sciences and technologies of diamond-related wide bandgap semiconductors, and superhard

materials. The first ADC was held at Auburn University and has since been held biennially in Japan,

Europe, and the United States. The exploration and rapid advances in sciences and applications of

diamond-related materials expanded into significant and novel carbon materials and nanostructures such as

fullerenes, nanotubes, C-N, B-C-N compounds, and so on, leading to the establishment of the International

Conference on Frontier Carbon Technology (FCT) that was first held jointly with the Fifth Applied

Diamond Conference (ADC/FCT '99) in Tsukuba, Japan, in 1999.

These are the Proceedings of the Sixth Applied Diamond Conference/Second Frontier Carbon

Technology Joint Conference hosted by Auburn University from August 6 to 10, 2001. The diamond CVD

process was first reported by Dr. Spitsyn in 1981 and Prof. S. Iijima reported his discovery of carbon

nanotubes in 1991. In the past years, both diamond-related materials and novel carbon materials have

attracted considerable interest by the scientific, technological, and industrial community. Many practical

and commercial products of diamond materials are reported in these proceedings. A broad variety of

applications of carbon nanotubes and novel carbons have also been explored and demonstrated. Having

more than 200 invited and contributing papers by authors from over 20 countries for presentations at

ADC/FCT 2001 clearly demonstrates that these materials, due to the combination of their superior

properties, are both scientifically amazing and economically significant.

Invited speakers, contributing authors, and participants comprising leading scientists, R&D managers,

engineers, researchers, investors, and funding officers will gather at the Auburn University Hotel and

Dixon Conference Center to exchange their visions and discuss manufacturing, application, technological,

and enabling fundamental issues on diamond materials, nitrides, carbides, carbon nanotubes, and related

nanomaterials, structures, devices, and systems. Due to the multidisciplinary nature of this conference, five

pre-conference short courses are offered. It is our wish that by means of this conference, accelerated

commercialization as well as novel and practical applications of this wonderful class of materials will be

achieved.

We thank all of the sponsors, invited speakers, instructors, contributors, attendees, committee

members, and session chairs who have made this conference a success.

Y. Tzeng

Conference Chairman

M. Murakawa

Organizing Committee, Co-chairman

K. Miyoshi

Program Committee, Chairman

M. Yoshikawa

Organizing Committee, Co-chairman

Y. Koga

Program Committee, Co-chairman

K. Kobashi

Program Committee, Co-chairman

G.A.J. Amaratunga

Program Committee, Chairman

NASA/CP--2001-210948 iii August 6, 2001

PROCEEDINGS OF THE SIXTH APPLIED DIAMOND

CONFERENCE/SECOND FRONTIER CARBON TECHNOLOGY

JOINT CONFERENCE (ADC/FCT 2001)

ORGANIZING COMMITTEE

Y. Tzeng, Chairman

Auburn University, United States

M. Murakawa, Co-chairman

Nippon Institute of Technology, Japan

J.L. Davidson

Vanderbilt University, United States

S.-T. Lee

City University of Hong Kong, China

R.S. Sussmann

King's College London, United Kingdom

L. Schaefer

Fraunhofer IST, Germany

K. Kobashi

Kobe Steel and Frontier Carbon Technology

Project/Japan Fine Ceramics Center, Japan

P. Pehrsson

Naval Research Laboratory, United States

J. Sung

Kinik, Taiwan

M. Yoshikawa, Co-chairman

Japan Fine Ceramics Center, Japan

K. Miyoshi

NASA Glenn Research Center, United States

Y. Koga

National Institute of AIST, Japan

R.J. Nemanich

North Carolina State University, United States

J.C. Angus

Case Western Reserve University, United States

J.E. Butler

Naval Research Laboratory, United States

O. Zhou

University of North Carolina, United States

A. Grill

IBM, United States

M.P. D'Evelyn

GE Corporate R&D, United States

PROGRAM COMMITTEE

K. Miyoshi, Chairman

NASA Glenn Research Center, United States

Y. Koga, Co-chairman

National Institute of Advanced Industrial

Science and Technology, Japan

G.A.J. Amaratunga, Co-chairman

University of Cambridge, United Kingdom

O. Zhou

University of North Carolina, United States

K. Kobashi, Co-chairman

Kobe Steel and Frontier Carbon Technology

Project/Japan Fine Ceramics Center, Japan

P. K. Bachmann

Philips Research Laboratories, Germany

Y. Vohra

University of Alabama in Birmingham, USA

K. Prassides

University of Sussex, United Kingdom

NASA/CP 2001-210948 v August 6, 2001
A.Hatta

KochiUniversityofTechnology, Japan

T.Ito

OsakaUniversity, Japan

M.Kitabatake

Panasonic, Japan

A.Grill

IBM,UnitedStates

T.Imai

SumitomoElectric, Japan

R.Hauge

RiceUniversity, UnitedStates

M.Uemura

Toyohashi UniversityofTechnology, Japan

A.Hirata

TokyoInstituteofTechnology, Japan

N.Otake

TokyoInstituteofTechnology, Japan

S.Yugo

Electro-Communications University, Japan

M.Capelli

StanfordUniversity, UnitedStates

S.Matsumoto

National InstituteforResearch inInorganic

Materials, Japan

Z.Sitar

NorthCarolinaStateUniversity

K.Komaki

JapanNewDiamondForum, Japan

K.Dohda

GifuUniversity, Japan

INTERNATIONALDVISORYCOMMITTEE

S.Iijima

NEC& MeijyoUniversity, Japan

F.Lu

BeijingUniversityofScienceandTechnology,

China

A.Sawabe

AoyamaGakuinUniversity, Japan

R.Kalish

Technion, Israel

H.Ohkushi

AIST,Japan

P.Koidl

FIAF,Germany

J.F.Prins

UniversityofWitwatersrand, SouthAfrica

K.H.Chen

AcademiaSinica, Taiwan

C.-P.Klages

FhG-IST, Germany

M.Schreck

UniversityofAugsburg, Germany

J.M.Kim

SamsungSDI,SouthKorea

J.S.Foord

UniversityofOxford,UK

T.Imai

SumitomoElectricIndustries, LTD.,Japan

B.Spitsyn

InstituteofPhysical Chemistry, Russia

J.Robertson

Cambridge University, UnitedKingdom

M.Kamo

National InstituteofMaterialsScience, Japan

NASA/CP2001-210948 vi August 6,2001
T.Sakai

ToshibaCorp.,Japan

S.Prawer

UniversityofMelbourne, Australia

J.Angus

CaseWesternReserveUniversity, USA

LOCAL ARRANGEMENTS COMMITTEE

G.W. Blanks

Auburn University, United States

E.H. Ridgway

Auburn University, United States

C. Cutshaw

Auburn University, United States

J.H. Brady

Auburn University, United States

C. Liu

Auburn University, United States

Y.C. Chen

Auburn University, United States

Candi Smith

Auburn University, United States

C. Ellis

Auburn University, United States

PUBLICATION COMMITTEE

L.C. Feher

InDyne, Inc., at NASA Glenn Research

Center, United States

K. Soboslay

InDyne, Inc., at NASA Glenn Research

Center, United States

K. Joyce

InDyne, Inc., at NASA Glenn Research

Center, United States

NASA/CP 2001-210948 vii August 6, 2001

TABLE OF CONTENTS

1. REVIEW OF RESEARCH PROGRAMS

NSF FUNDED RESEARCH PROGRAMS ON HARD COATINGS AND CARBON NANO TECHNOLOGY

Join Larsen-Basse, National Science Foundation, U.S.A .............................................. 3

ONR RESEARCH PROGRAMS ON CARBON NANOTUBE

John Pamk, Office of Naval Research, U.S.A ....................................................... 4

CURRENT STATUS OF R&D IN FRONTIER CARBON TECHNOLOGY (FCT) PROJECT--OVERVIEW

M. Yoshikawa, Frontier Carbon TechnoIogy Project/Japan Fine Ceralnics Center, Japan ..................... 5

ACTIVITIES OF NOVEL MATERIALS SYNTHESIS TECHNOLOGY GROUP IN FRONTIER CARBON

TECHNOLOGY PROJECT

Yoshinori Koga, National h_slitute ofAdvat)ced It_dusi;rml Science and Technology, Japan ................... 9

RESEARCH AND DEVELOPMENT OF HARD CARBON FILMS AND THEIR TRIBOLOGICAL

PROPERTIES IN THE FRONTIER CARBON TECHNOLOGY PROJECT

Akihiro Tanaka, Natienal Institute of Advanced Industrial Science and Technology, Japan .................. 11

CURRENT STATUS OF R&D IN FRONTIER CARBON TECHNOLOGY (FCT) PROJECT

ELECTRICALLY FUNCTIONAL MATERIALS

K. Kobashi, Frot_tier CaFoon Technology Pmject/lapat_ Fit)e Ceramic_ Center, Japan ........................ l 6

R&D OF DIAMOND FILMS IN KOBE STEEL

K. Kobashi and _5 Yokota, Frontier Carbon Technology, Preject/Japan Fine Ceramics Center, Japan; and

N. Kawakami, T. Tachibana, K. Hayashi, and K. Innue, Kobe Steei, Ltd., Japan .......................... 24

R&D OF CVD DIAMOND IN SUMITOMO ELECTRIC INDUSTRIES

T. imai, Sumitomo Electric industries, LTD, Japan .................................................. 31

THE DAWN OF DIAMOND AGE

_ames Cbien-Min Sung, KINIK Company and National Taipei University oUl-echaology, ikiwan, R.OC ...... 32

APPLICATIONS OF CARBON NANOTUBES TO ENERGY STORAGE

Se'_mg M; l_,ee_ Kay Hyeok An, Kwan Ku _eon, \_/i)t_ Seok Kim, _eong-Mi Moon_ Ymmg So(; Park, an'l

52nmg Hee Lee, Sungkyankv_an L' nivers ily, South Korea; and Chil-Hom- Doh, .)ang-H,a'an Lin-, ar, d

Seong-In Meon, Korea Electrotechnni,'_gy Research Instimle, South Korea .............................. 33

2. MARKETS AND TECHNOLOGIES OF DIAMOND MATERIALS

CVD DIAMOND: THE INDUSTRIAL LANDSCAPE

R.S, Sussmann, King's College Londom U.K ......................................................... 43

DIAMOND AND DIAMOND-LIKE CARBON FILMS AND COATED PRODUCTS_ TECHNOLOGY

AND MARKET UPDATE

T. Abraham, Business Communications Co., U.S.A ................................................. 44

THE PRODUCTION AND FEATURES OF INDUSTRIAL DIAMOND MADE IN CHINA

Zhibin Zhan, Queenbee Diamond industrial Co. LTD, China; andAiguo Yao and Bairu Xia, China

University of Geosciences, China ................................................................... 45

NASA/CP--2001-210948 ix August 6, 2001
3. DIAMOND APPLICATIONS

3.1 Diamond Electronics and Sensors

CHARACTERIZATION OF A1N AND LiNbO 3 THIN FILMS DEPOSITED ON DIAMOND SUBSTRATES

M. Ishihara, T. Naka_a_;a, Y. Koga, and F. Kok_i, Nal:or, al h-stimte ,,_fAdvansed lnddstria! Science and

Technology, Japan ........................................................................... 51

HIGHLY ORIENTED DIAMOND THIN FILMS FOR SILICON-ON-INSULATOR TECHNOLOGY

S.D. Wol_,er, G.N. "_klshin, F. Okttzumi, R. Schlesser, and Z. Sitar, Noah Carolina State Universi_,y,

U.S.A.; LT. Prater, Army R_-search Office, U.S.A.; and B.R. Stoner, MCNC, U.S.A ....................... 55

CVD DIAMOND FOR HIGH TEMPERATURE ELECTRONICS

Susan L Heidger, Navxd J. Bm'aly, and Jesepb A. Weimer, U.S. Air Force, Wright-Patters,'m Air

Force Base, I).S.A ........................................................................... 61

POLYCRYSTALLINE DIAMOND FILMS FOR HIGH-TEMPERATURE ELECTRONICS AND MEMS

R. Ramamurti, V. Shanov, R.N. Singh, M. Samiee and P. Kosel, Ueiversity of Cincinnati, U.S.A ............. 62

DIAMOND MEMBRANE ELECTRON MULTIPLIER

Nick N. Dzbanovsky5 Pavel Mmakov, at_d Nikolay V. S_etln, Moscow State Utqversity, ICalssia; and

EogenJ A. Poltorztsk5; Gent_a'ly S RJchkov, Eageny A. l!'h:hev, awl Sergey A. Gavrilov, Zelenograd's

Physical Pr_blems !nstilute, Russ:a ............................................................. 68

ELECTRICAL PROPERTIES OF BORON-DOPED DIAMOND FILMS GROWN BY BIAS METHOD

S. Yugo, T. Tashiro, and D. Saito, University of Electric-Commueications, Japan ......................... 7i

NOVEL ELECTRONIC DEVICES USING DIAMOND THIN FILMS

A. Vaseashta, Marshall University, U.S.A ......................................................... 72

ELECTRON AND HOLE CONTRIBUTION TO THE RESPONSE OF CVD DIAMOND NUCLEAR

DETECTORS

Marco Marinelli, E. Milm'.i, A. Pao!etti, G. Pucella, A. Tucciarone, A. Stelitano, G. Verona-Rinati,

University of Rome "Tor Vergata," Italy; and S. Albergo, V. Bellini, A. Musumarra, R. Potenza,

C. Randieri, M.L. Sperduto, and C.M. Sutera e A. Tricomi Universit_ di Catania, Italy ..................... 78

TEMPERATURE DEPENDANCE OF TRAPPING-DETRAPPING EFFECTS IN CVD DIAMOND

PARTICLE DETECTORS

Mm'_:_; Mm'inelli, E. MJlani, A Pa_;ie,lti, A. Tucciarone, and G Ve,r,'ma-Rina_i_ Univers_ta di Roma, Italy;

and M. Apgelone and M. Pil!or,, Associazione EURAI'OM-ENEA s_lla Fusione, Italy ..................... 80

3.2 Diamond Micro-Electro-Mechanical Systems (MEMS)

ON THE FABRICATION AND BEHAVIOR OF DIAMOND MICROELECTROMECHANICAL SENSORS

(DMEMS)

K. Holmes, J.L. Davidsop, W.P. Kang, and M. Howell, Vanderbilt University School of Engineering, U,S,A .... 82

DIAMOND MEMS_ROM PROOF-OF-CONCEPT TO APPLICATION

E. Kal-a, M. Adamscl-ik, and R Schmid, l)niversily of Ulm, Gem-any; and S. Ertl mad A. Fl6ter,

GesellschaR fth" Dimnm,tt;redukte , Gerrnar, y ...................................................... 90

MICROMACHINING OF CVD DIAMOND FOR MEMS APPLICATIONS

Takayu!d Shibata, Hokkaido University, Japan ..................................................... 98

NASA/CP--2001-210948 x August 6, 2001
3.3 Electron Emitters

FABRICATION AND BEHAVIOR OF DIAMOND FIELD EMITTER TRIODE UTILIZING SILICON-ONINSULATOR TECHNOLOGY AND CVD DIAMOND

W. P Kaog, A. Wis:ts(;ra-al, m'd J.!,. Davidsoo, VaoderbJlt (Jmversity, U.S.A ; and D. V Kems,

Olin Co!lege of Engineering, U.S.A ............................................................ 104

ELECTRONIC PROPERTIES OF PHOSPHORUS-DOPED TRIODE-TYPE DIAMOND FIELD

EMISSION ARRAYS

Chia-Fu Chen, Chia-Lun Tsai, and Chien-Liang Lin, National Chiao Tung Universir¢, Taiwan .............. 109

MICROFABRICATION OF VARIOUS ELECTRON EMITTERS ON SINGLE CRYSTAL DIAMOND

Y. Nisb%a-¢ashl, Y. Ando, and K. Kcbasbi, Osaka University, Japan; K. Meguro and T. Imai_ Sumitomo

Electrk: In&lslries_ Lid., Japing; ar_'l T. Hirao add K, Oora, Osaka 1,1nivers;ty, Japan ........................ 110

FIELD EMISSION FROM GRAPHITE-DIAMOND COMPOSITES

L. LeT, J.B. Cui, M. Stammler, and J. Ristein, University of Erlangen, Gennal,.y . ........................ 114

FIELD EMISSION FROM DIAMOND FILMS CONTAINING ION-IMPLANTED GRAPHITIC REGIONS

N. Koenlgsfeld, R. Kalish, 1st, Hofthaan, and R. _lavlediani, Technion, _;srael .......................... 120

FIELD EMISSION CHARACTERISTICS OF ION-IMPLANTED CVD DIAMOND

i'. Ono, i. Sakai, N. Sakmr, a, L. Zhar,g, and H. Yosbida, Frontier Carbop iechnology Project/Japm_ Fine

Ceramics Center and Toshiba Corporation, Japan; and M. Hasegawa and H. Okushi, Electrotechnical

Laboratory, Japan .......................................................................... 121

THERMIONIC FEEM, PEEM AND I/V MEASUREMENTS OF HYDROGEN TERMINATED N-DOPED

CVD DIAMOND SURFACES

F.A.M. K6ck, J.M_ Garguiio, B. Brown, and R.J. Nemanich, Noah Carolina State University, U.S.A ......... 126

N-DOPED CVD DIAMOND FILMS AS A LOW TEMPERATURE THERMIONIC-FIELD ELECTRON

SOURCE

F.A.M. K6ck, J.M. Gacguil,% B. Brown, R.J. Nemanich, North Carolina State i,Jniversity, U.S.A ............. !32

FIELD EMISSION, PEEM AND FEEM MEASUREMENTS OF EMITTING SITES OF MPCVD GROWN

NANOCRYSTALLINE DIAMOND FILMS

J.M. Garguilo, F.A.M. K/Sck, Billyde B_o_xm, and RoJ. Nemanlch, Noah Carolina State University, U.S.A ..... 133

INVESTIGATION ON FIELD EMISSION OF Si TIPS WITH NANOCRYSTALLINE DIAMOND COATINGS

W.L. \Vang, K.J. L{ao, C.G. Ha, m_d Z. Li, Ch_;ngqit3g _,;niversily, RR. China ........................... 139

PARAMETER ESTIMATION FOR ELECTRON FIELD EMISSION

D.G. Walker, T.S. Fisher, J.L. Davidson, and W.R Kang, Vanderbilt University, U.S.A ..................... 144

3.4 Diamond Electrodes

CONDUCTIVE DIAMOND: A UNIQUE ELECTRODE MATERIAL FOR ANALYTICAL APPLICATIONS

Akira Fujish:ma, The Unive_s:ty of Tokyo, Jal;aO ................................................... 150

ADVANCED ELECTROCHEMISTRY WITH DIACHEM ® ELECTRODES

L. ScMfer, M. Fryda, D. Herrmann, and I. Tr6ster, Fraunhofer Institute for Surface Engineering and Thin

Films, IS% Gem3any; and W. Hfinni andA. Perret, Centre Suisse d'Electronique et de Microtechnique SA

CSEM, Switzerland ......................................................................... i58

NASA/CP--2001-210948 xi August 6, 2001
ELECTROCHEMICAL PROPERTIES OFSULFUR-TREATED DIAMOND

SallyC. Eaton, Alfred B. Anderson, and John C. Angus, Case Western University, U.S.A.; and Yulia E.

Evstef_eva and "_ktri V. Pleskov, Frumkin institute of Electrochemistry, Russia .......................... 164

STM AND ELECTROCHEMICAL INVESTIGATION OF HOMOEPITAXIAL BORON-DOPED CVD

DIAMOND FILMS

John B. Cooper, Jason A. Moulton, Sacharia Albh', and Bip_g Xiao, Old Don:inim'_ University, U.S.A ......... 170

HTHP DIAMOND ELECTRODES: ELECTROCHEMICAL IMPEDANCE CHARACTERISTICS OF

INDIVIDUAL CRYSTAL FACES

"_kl.V. Pleskov, M.D, Krotova, _ Ya. Mish,3k, and Yu, E. Evstefeeva, Russian Academy of Sciences, Russia;

V.A_ Laptev, All-Pcassian Research _nstitu_,e for Synthesis of Mineral Rmv Materials, Russia; and _4_. N.

Palyan_;v and Yu. M Borzdov, Russiat_ Academy of Sciences, Russia .................................... 17_

3.5 Diamond Optics

FAR INFRARED CVD DIAMOND OPTICS: PROPERTIES AND APPLICATIONS

T.R Molia_ and K.L. Lewis, Defence Research and Evaluation Agency, U.K ............................ i77

THIN FILM DIAMOND DETECTORS FOR INDUSTRIAL APPLICATIONS IN NEXT GENERATION

157 nm LITHOGRAPHY AND HIGH POWER EXCIMER LASER BEAM DIAGNOSTICS

Michael D Whitt_eld and Nadeem Rizvi, Exi_ech Ltd., U.K,; Rober_ D, McKe_g, Cent_onk:, U.K; aad

Snlm'l P. Lansley_ Oii'eier Gmldin_ and Rish_cd B. Jackmm_, U ni_,ersily College Lond_;n, U.K ............... 178

SINGLE-PIXEL DIAMOND-BASED VUV PHOTODETECTORS FOR SPACE APPLICATIONS

A. De Sio, E. Pace, and A. Pini, Universit_ di Firenze, Italy; and S. Scuderi, Ossera, atorio Astrofisico di

Catania, italy .............................................................................. 179

DIAMOND THIN-FILM FABRY-PEROT OPTICAL RESONATORS

R.A, Booth and D.K. Reinhard, Michigan Stale University, U.S.A ...................................... 189

SURFACE FLASHOVER CHARACTERISTICS OF POLYCHRISTALLINE DIAMOND AND DLC

DIELECTRIC THIN FILMS IN VACUUM: A REVIEW

Hulya Kirkdci, Auburn University, U.S.A ........................................................ !86

OPTICAL PROPERTIES AND RESIDUAL STRESS OF DIAMOND FILMS

YN. Sun, J.H_ Zhang, J.Q, Li, WIT. Guo, and X. Gao, Lanzhou Ins_,imte of Physics, China ................. 197

3.6 Diamond Machine Tools

STATE-OF-THE-ART AND FUTURE TRENDS FOR CVD-DIAMOND COATED TOOLS FOR CUTTING

OPERATIONS

O. Lemrner, D. Breidt, M. Frank, and T. Leyendecker, CemeCon GmbH, Gem__any ....................... 198

ULTRASONIC END MILLING OF HARD FERROUS MATERIALS USING DIAMOND TOOL

Masao Murakawa, Hiroyuki Noguchi, and Masahiko Jin, Nippon instir, ute of Tecbnology, Jaaan ............ 204

DIAMOND CUTTING TOOLS FOR DRY MACHINING APPLICATION

Prabtm Ammugam, Ajay P. Malshe, and Steve Balze_, [Jr_i,_e_s'_ty ,'ff Arkaasas, (J.S.A ..................... 210

CMP PAD DRESSER: A DIAMOND GRID SOLUTION

James Sung, KINIK Company, Taiwai'l, R.O.C .................................................... 211

NASA/CP--2001-210948 xii August 6, 2001
THERECIPROCATIVE SAWINGOFGRANITE

JamesC.Sting,KLNIKCompanyandNational Taipei UniversityofTechnology, Taiwan, R.O.C............ 213

THECASTINGOFCVDDIAMOND

JmDesSungandSbiao-KaoChang, KYNIKCompanyandNational Tail-el UmversitvofTechnology,

i'aiwan,R.O.C............................................................................. 214

ANANALYSISOFTHECATALYSISOFFe,Ni orCoONTHEWEAROFDIAMONDS

M.Uemura, Toyohashi UniversityofTechnology, Japan............................................ 215

4. DIAMOND SYNTHESIS, MODELING, AND CHARACTERIZATION

4.1 High-Pressure High-Temperature Diamond Processes

HIGH-TEMPERATURE ANNEALING OF COLOUR CENTRES IN DIAMOND

Alan T. Collins, King's College London, U.K ..................................................... 223

DEFECT AGGREGATION AND DISSOCIATION IN DIAMONDS DURING ANNEALING kINDER HPHT

CONDITIONS

Filip De W_erdt, Hoge Raad voor Diamant, Belgmm ....................................................

FULLERENES AS A CO-CATALYST FOR HPHT SYNTHESIS OF DIAMOND

A. _(a. Yul' and S.V. Kidalov, Ioffe Physico-Technical Institute, Russia; V.M. Da,Adenko m'.d V.A. Yashin

Iiyich Abrasive Company, Russia; and S.S. Ordanyan, St. Petersburg Technological Ulfiversig,, Russia ....... 237

FORMATION OF SYNTHETIC SINGLE AND POLYCRYSTALLINE DIAMONDS

Vladimir Prokofievieh Poliakov, Federal Universi W of Pa,rana, Brasil ................................. 239

THE MULTI-ANVIL SPHERICAL TYPE DEVICE OF HIGH PRESSURE WITH CUBIC COMPRESSION

CHAMBER

Guerold Bobr,'_vnitdfii, Labo_amr 2 of Advm_ced Materials of Stale Universi D oflhe North F hmrhlense,

Brazil; and Eduard Persikov, Laboratory of High Pressm:e of Institute ef Experimental Mineralogy of AS

of Russia, Russia ........................................................................... 245

THE OPTIMIZED CELL DESIGN FOR HIGH PRESSURE SYNTHESIS OF DIAMOND

James C. SLing, K1NIK Company and National Taipei University of Technology, Japan ................... 253

LIQUID PHASE SYNTHESIS OF DIAMOND IN HYDROGEN ATMOSPHERE

Hsiao- Kuo Ctnmg, KINIK Con-par, y, Taiwan_ R.O.C.; Miog-Sbow Wi;ag, I'ang-Hwa University_ Taiwan,

R.O.C.; Yong Liu, Chinese Academia Sinica, Taiwan, R.O.C.; and James C. Sung, IcANIK Company and

National Taipei University of Technology, Taiwan, R.O.C ........................................... 254

4.2 Ultradispersed Diamonds

THE INVESTIGATION OF FORMATION AND AGGREGATION OF UDD (ULTRADISPERSED)

DIAMONDS ON SPECTRA OF THEIR IR (INFRARED) ABSORPTION

A. Kocetz, E. Mironev, VI ,Malyi, Krasnoyarsl_ .'State _ieehmcal University, Russia; and E. Petrov,

Federal Scientific Resem:d- Cer,{er "Almy," R_ssia ................................................ 255

CHEMICAL STATES CARBON ATOMS ON UPPER LAYERS OF ULTRA DISPERSION DIAMOND

(UDD) AND NATURAL DIAMOND (ND) AFTER H-TREATMENT

A.E Dementjev, K.I. Maslakov, and O.O. Zabusov, RRC "Kurchatov institute," Russia ................... 257

NASA/CP--2001-210948 xiii August 6, 2001
4.3 Diamond Nucleation

DIAMOND (HETEROEPITAXIAL) NUCLEATION BY ION BEAM IMPACT: A REALITY

S T. Lee, I. BalM, and C.S. Lee, Cfly Universily of Hong KJ;ng, Cbma ................................... 262

HRTEM STUDY OF THE EARLY STAGES OF DIAMOND NUCLEATION ON Si(111) DIMPLES

J.C. Amault, S. Peceraro, J. Werckmann, and F. Le Nennand, Institut de Physique et Chimie de Strasbourg,

IPCMS-GSI, France; and N. Motta and R. Polini, Universita Tot Vergata, Italy .......................... 263

ULTRADISPERSE DIAMOND STRUCTURE, PHASE TRANSFORMATION AND APPLICATION FOR CVD

DIAMOND FILMS GROWTH

A,E. Alexenskii, MN. Baidakova, N,A. Feoktistov, V.G. Golubev, S. Grudinkin, A_T, DkteikJn, V.G.

Melehin, V Y_. Osipov, VI. Siklitski, and A. "_. V_d _, ioft_z Physmo-Technieal institute, Russia; A,M

Panich and I. Shames, Ben.G_rkm Universlty o:_ the Negev, israel; W. KempmsM, institute of Molecular

Phvsics, Po!ish Acadmr.v of Sciences, Poland; aad A. go'iiger, E. Osawa, and M. Ozawa, Toyohashi

UniversiD of Technology, Japan ............................................................... 264

THE PROPERTIES OF NUCLEATION AND GROWTH FOR CVD DIAMOND ON MIRROR-POLISHED

Si PRE-SEEDED WITH NANOSIZED DIAMOND POWDER

Zengs_m Jim Changzhi Ga, Yue Stm, and Ye Sun, Jilin University, China ............................... 266

EFFECT OF CYCLIC TIME INTERVAL OF CH 4 FLOW RATE ON THE DIAMOND NUCLEATION

DENSITY ON GLASS SUBSTRATE

S.-H. Khn and E.-J. Bae, Silla University, Korea; T.-G. Kim, Mirym_g University, Korea; 3'.-H. Khn,

POSTECH, Korea; and T. Hosolni, T. Maki, and T. Kobayashi, Osaka University, Japar .................... 267

DIAMOND NUCLEATED BYARC DISCHARGE AND GROWN BY MWPCVD

Changzhi Gu and Zengstm Jin, Jilin University; China ............................................. 273

4.4 Homoepitaxial Growth of Diamond

HOMOEPITAXIAL DIAMOND GROWTH AT HIGH AND LOW TEMPERATURES: DESIGNER

DIAMONDS TO LARGE CRYSTALS

Yogesh K. Vohra, University of Alabama at Bimaingham, U.S.A ...................................... 274

BORON-DOPED HOMOEPITAXIAL DIAMOND (100) FILM INVESTIGATED BY SCANNING

TUNNELING MICROSCOPY

Bing Xiao, Weihai Fu_ and Sacharia Albin, Jason Moulton, and Jobn Cooper, Old Dominion

Unive: sity, U.S.A ............................................................................... 275

CATHODOLUMINESCENCE TOPOGRAPHY OF HOMOEPITAXIAL DIAMOND MADE BY

MICROWAVE PLASMA CVD

Itisao Kanda, Satoshi Koizumi, Kenji Watanabe, National Institute for Materials Science, Japan; and

Tokuyuki Teraji, Osaka University, Japan ........................................................ 279

4.5 Growth and Manufacturing of Diamond

ORIGIN, EVOLUTION AND SOME PROBLEMS OF DIAMOND CVD

B.X/_ Spitsyn and A.E, Alexepko, Insl'Xnte of Phvsical Chemistry RAS, Russia ........................... 284

COMPUTER SIMULATION OF LARGE AREA DIAMOND DEPOSITION BY DC ARC JET WITH ARC

ROOTS ROTATING AND OPERATING AT GAS RECYCLING MODE

F.X. Lu, T.B. Huang, W.Z. Tang, J.H. Song, and Y.M. Tong, University of Science and Technoiogy

Beijing, ER. China ......................................................................... 285

NASA/CP--2001-210948 xiv August 6, 2001
GAS-PHASECONCENTRATIONS ANDTEMPERATURE MEASUREMENTS OFC2INADIAMOND

DEPOSITINGDC-ARCJETCVDSYSTEM

J.A_Smith_KN.Rosser, andP.W.Ma_; University of Bristol, U.K.; H. Yagi, University of Ehime,

Japan; andA.J. OrroEwmg at_d M.N.R. Ashtbld, Uoiversily of Br_st_;i, U.K ............................... 291

NUMERICAL SIMULATION AND REALISATION OF NOVEL MICROWAVE PLASMA REACTORS

FOR DIAMOND CVD

C. Wild, E. Pleuier, W. Miiller-Sebert, and P. Koidl, Fraunhofer-instimt ,_ir Angewandte Festk6rperphysik,

Gem_any ................................................................................. 292

ELABORATION AND MECHANICAL PROPERTIES OF DIAMOND COATINGS OBTAINED BY

FLAME PROCESS

D. Pa'_dmier and M. Schmitt, LPMM-ERMES, France; and M. Mem_ou_, LEPMI-ENSEEG k_;rnalne

Llniversitaire, France ............................................................................ 293

OPTICAL QUALITY DIAMOND WAFER SYNTHESIS BY MULTI-CATHODE DC PACVD

Kwang Yong Eun, Jae-Kap Lee, and Young-Joon Baik, Korea Institute of Science and Technology, Korea .... 304

METASTABLE GROWTH OF DIAMOND IN MOLTEN ALLOY

Yen-Te Chert, Kai-Hon Hsu, Hsiao-Kuo Chung, and James C. Sung, t';XNIK Company and National

'ihipei Universily of Tecbnol_;gy, 'Fa_wan, R.O.C ..................................................... 305

FABRICATION OF DIAMOND MICROPATTERNS BY ECR OXYGEN PLASMA WITH A METAL

NAPTHENATE MASK

Shuji Kiyohara and Katsumi Mori, Science University of Toks,o in Yamaguchi, Japan .................... 317

MORPHOLOGICAL CONTROL OF DIAMOND FILMS IN A 60 kW MICROWAVE PLASMA

CVD REACTOR

Y. Yokota, Y. Ando, T. Tachibana, A. Watanabe, YI Nishibayashi, K. Kobashi, T. Hirao, and K. Oura,

Osaka Universits; Japon ......................................................................... 323

A HIGHLY EFFECTIVE SETUP FOR DIAMOND COATING DEPOSITION

V.K. Pash]'_ev, O.A. Opalev, 1.!. Yyrovets, ViE. Stcel'nitskij, and V.A. Belous, National Science Centre,

Ukraine; Z.i. Koiupaeva, Kharkov State Polytechnical University, Ukraine; and Yu. F. Shmal'ko,

Institute o f Mechanical Engineering Problems, Ukraine ............................................ 327

HIGH GROWTH RATE AND HIGH QUALITY CVD DIAMOND GROWTH

H.-G. Jentsch_ S.M. Ros_wal, R.F. Sk_ger, Friedrich-Ale_ander Urdvers_t/it Erlat_gen-Nfm_berg, Gern_any;

and H. Eibisch, Neae Mate6ahea Fti_th GmhH, Ge_m_r_y .............................................. 333

THE RAPID GROWTH OF THIN TRANSPARENT FILMS OF DIAMOND

Itsiao-Kuo Chung and James C. Sung, K125IiK Cornpany, Taiwan ..................................... 334

DIAMOND PARTICLES SYNTHESIZED ON GLASSY AND OXIDE SUBSTRATES

Yoshi!d Takagi and Fumitomo Onishi, Teikyo University of Science and Technology, Japan ................ 335

HIGH QUALITY DIAMOND GROWTH ON IRON AND STEEL SUBSTRATES

E_I N_km_ll_ra_ Mar_at'_ Ohyama_ an'l KeNi K. Hirakori, Tokyo Denki L'mversilv, Japan; Not'_kl

Mu_sol_an'a and "Yoshio Mach'_, T,'_kyo Denk'_ Uni'_e_si_, _apan; and Gerpot l_rie,'lbacher, se!_enna

U]fiversit_/of Yechnologs', Austria .............................................................. 340

THE UNIFORMITY OF 8 INCH CVD DIAMOND FILMS

Sheng-Yi Lee, Shiao-Kuo Chang, Jow-Lay Hung, and James C. Sung, gdNIK Company, Taiwan, R.O.C ...... 343

NASA/CP--2001-210948 xv August 6, 2001
SYNTHESISOFDIAMONDUNDERHIGHPRESSURE USING HOT-FILAMENT CVD METHOD

Eri Nakamm'a, Toshihiro Kobayashi, and KeNi K. Hirakuri, Tokyo Denki University, Japan; Nobuki

MuT,sukura and Yoshio Machl, Tokyo Denkl University, Japan; and Gemot Friedbacher, Vienna

Upiversity of Tectmology, Austria .................................................................... 344

DIAMOND SYNTHESIS BY MICROWAVE AND DC HYBRID PLASMA CVD

Keiji F,4ji , Taku Surnmitom,% and Akimits-d Hatta, Kochi University of Technology, japan ................. 348

CHEMICAL PROCESSES IN ECR-MICROWAVE PLASMAS CAPABLE OF DIAMOND FILM

DEPOSITION

l_Jk Blumen_,hal, Auburn UniversiD, USA ....................................................... 349

EXPERIMENTAL CONFIRMATION OF POSSIBILITY OF SYNTHESIS OF DIAMOND SINGLE

CRYSTALS SIMILAR TO NATURAL

T.A. Nachalna2, V.G. Malogolovets, G,A. Pod2Nm'ey, S.A. Ivahlqlenko, mzd O.A. Zanevskiy, Institute for

Superhard Materia!s of NAS ef Ukraine, Ukraine; m'_d I.A. Novoselova, Institute of General & InorgapAc

Chemistry of NAS of Ukraine, Ukxaine ......................................................... 355

THE FTIR STUDIES OF (100) ORIENTED DIAMOND FILMS GROWN ON Si SUBSTRATE USING

TEMPERATURE GRADIENT ACROSS THE SUBSTRATE

E, Titos, iDdian Institute of'_chnology; iDdi_; A.K. Sikder, (Jmversity of South Florida, U.S.A.; and

U. Palmika_, M.K Singh) an'l D.S. Misra_ indian Iastku_e of'ieehnolo_', India ........................... 360

DIAMOND MICRO CRYSTALS GROWN BY CVD METHOD

Itsiao-Kuo Chung and James C. Sung, KESIiK Company, Taiwan, R.O.C ............................... 364

DIAMOND SYNTHESIS IN PERSPECTIVE

James Sung, China Grinding Wheel Corao_azion and National Talpei University of Technology,

Taiwan, R_O.C ............................................................................. 365

THE EASTERN WIND OF DIAMOND SYNTHESIS

.)ames C. S_mg, KfNIK Company, Taiwm-, R,O.C.; and Jol-a C. Chen, Fine Abrasives Company of

Taiw_n, Taiwan, R.O.C ...................................................................... 366

MODERN THERMODYNAMICS FOR ACTIVATED LOW PRESSURE CVD DIAMOND SYNTHESIS

Ji-Tao Wang, Fudan University, China .......................................................... 367

4.6 Characterization of Diamonds

PREPARATION OF DIAMOND SURFACES

James E. Bm!er, Naval Research Laboratory, U.S.A ................................................ 368

INFLUENCE OF ELECTRON BEAM IRRADIATION ON WETTABILITY OF DIAMOND SURFACE

Kazuya O_'i and Akira Tonegawa, Tokai University, Japan; and Nobuhiro Iwataka and Yoshitake

Nishi, Tokai UniversiD; Japan ................................................................. 369

TIME-RESOLVED ELECTRICAL RESISTANCE MEASUREMENTS OF GRAPHITIZED LAYERS IN

DIAMOND DURING PULSED LASER IRRADIATION

S.M. Piroen_;% VNI K_monenko, E,V. Zavedeev, and V.I. Konov, Geveral Physics !nstim_e, Russ'_a;

M. Kehr!i, W. Lfithy, V. Romano, and H.P. Weber, Uni,_ersity of Bern, Switzerlandl and R.A. Khmelnitskiy,

RN. Lebedev Physicai Institute, Russia .......................................................... 370

NASA/CP--2001-210948 xvi August 6, 2001
MECHANISMOFELECTRICALCONDUCTIVITYAPPEARED1NSURFACEREGIONOFUNDOPED

DIAMONDFILM

Takeshi Hosomi, KiyoshigeTsuji,andTakeshi Kobayashi, OsakaUniversityandCP, EST,JapanScience

andTecbru_Jogy Co_;ration, Japan................................................................ 372

APPLICATIONOFRAMANSPECTROSCOPY TODETERMINESTRESSINPOLYCRYSTALLINE

DIAMONDTOOLSASAFUNCTIONOFTOOLGEOMETRYANDTEMPERATURE

J.D.Cemins, R.M.Erasmusal,.d\( Mofokeng, UniversityoftheWitwatersrand, SouthAfrica;and

Z.Martin,UniversityofOxford,U.K........................................................... 376

EVALUATIONOFCRYSTALDEFECTSIN(100)AND(111)SINGLE-CRYSTAL DIAMONDSBY

CATHODOLUMINESCENCE SPECTROSCOPY

S_dao Takeuch_ amt Masao Murakawa_ Nippon lns6tute c CFeehnology, }apart ........................... 382

VACANCY IN BULK DIAMOND AND DIAMOND FILM: A MOLECULAR DYNAMICS STUDY

X.J. Hu, Y.B. Dai, H.S. Shen, Z.M. Zhm'_g, _(:Z. Wan, and X.C. He, Shanghai Jim; Tong University, China .... 386

5. APPLICATIONS OF DLC, NANO-DIAMOND, AND NOVEL CARBONS

5.1 Electrical and Electronics

THIN FILM DIAMOND-LIKE CARBON DIELECTRICS

Richard L.C. \Via, H. Kosai, and J. "0¢rb_nek, K Systems Cold, oration, U.S.A.; M. Freeman, DARPA

TACTICAL Teehnolo_./, U.S.A.; S. Fries-Ca_ and J. Weimer, Wright-Patterson AFt3, U.S.A.; and

EB. Kosel, University of Cincinnati, U. S.A ...................................................... 395

FROM DLC TO DLC-SiO 2 HYBRID LOW-k DIELECTRICS FOR ULSI INTERCONNECTS

Alfred Grill, iBM, U.S.A ..................................................................... 397

5.2 Electron Emitters

ENHANCED FIELD EMISSION FROM NANOSTRUCTURED CARBON FILMS

Igor Paviovsky, Ri. Fink, L.F. Thuesen, ap.d Zvi Yaniv, Field Emission Picture Elemep_t Technoiogy, Inc.,

U.S.A.: and Robert Espinosa, Microwave Power Technology, U.S.A ................................... 399

CONDUCTING ION TRACKS IN DIAMOND-LIKE CARBON FILMS

J. Krauser, A. WeidiDger, M Waibl_nger, V. Hoffmatm, Hahn_Meimer-Ir_stiialt, Ge_narw; C. TrautmaDn

Gese!lschafi t_r Scbwerieoenfo_sc_amg mbH_ Germany; B. Scholt_:cb, l_rmmh,'ffer-!nstilut IWS, GermaDy;

and H. Hofsclss and C. R_;nnmg, Universit'iit G6ttingep, Germany ..................................... 400

FIELD ELECTRON EMISSION FROM NANOSTRUCTURE DIAMOND ON POROUS SILICON

W.L. Wang, K.J. Liao. C.G. Hu, and J. Xu, Chongqing University, RR. China ........................... 406

ELECTRON FIELD EMISSION PROPERTIES OF NANO-DIAMONDS SYNTHESIZED BY CHEMICAL

VAPOR DEPOSITION PROCESS

I-Nan 15n, Gia-MiDg Lin, and Tong Hsu, Naliot)N Tsing_Hua Universit); Taiwan ......................... 410

5.3 Protective DLC Films

SUB 10 nm THICK AMORPHOUS CARBON OVERCOAT FOR THIN FILM MAGNETIC MEDIA

Takavu!d Yamamoto and HiroyuYJ Hyodo, Fijitsu Laboratories Ltd., Japan; and Takashi Toyoguchi,

Yamagata F Nitsu Ltd., Japan ................................................................. 412

NASA/CP--2001-210948 xvii August 6, 2001
IMPROVEMENTS OFTRIBOLOGICAL PROPERTIES OFMETALEVAPORATEDTAPESFORDIGITAL

RECORDINGBYDLCPROTECTIVECOATING

HiroyuMOsaki,SowCorporation, Japan....................................................... 413

PROPERTIES OFFLEXIBLEDLCFILMDEPOSITEDBYAMPLITUDE-MODULATED RFP-CVD

i'akabiroNakahigasbl, Ni'pon1TI_,h-c.,.lapar,................................................... 423

HIGH PERFORMANCE DLC COATINGS FOR SEVERE WEAR APPLICATIONS

Suri A. Saszri, Niiesh Gunda, and Richard Cooke, Surmet Corporation, U.S.A.; and Samosh Jha and

Vimal Desai, University of Cemral Florida, U.S.A ................................................. 424

LOW-TEMPERATURE TECHNOLOGY OF OBTAINING THE PROTECTING LAYER FROM

DIAMOND-LIKE FILM

K. Kassabian, Zh, Panosya_l, A. Stepat_ymL G Torosyao, and Y Yengibaryat_, State Engineering

University of Armenia, A_:_enia ............................................................... 431

THE APPLICATIONS OF AMORPHOUS DIAMOND DEPOSITED BY CATHODIC ARC

James C. Sung, KD_IK Company and Nationai Taipei University of Technology, Taiwan, R.O.C ............ 436

5.4 Wear and Tribology

SUPERLOW FRICTION CARBON FILMS FROM HYDROGEN-RICH HYDROCARBON PLASMAS

Ali Erdemir, Argor, ne Nati,;nal Labo:amry, {J.S.A ................................................. 437

FRICTION AND WEAR OF NANO-SIZE DIAMOND FILMS WITH AMORPHOUS PHASE

Koichiro Wazumi, Japan Fine Ceramics Center, Frontier Carbon Technoiogy Project/Japan Fine Ceramics

Cemer, Japan; and Yoshinori Koga and Akihiro Tanaka, National Institute of Advanced Industrial Science

and Technology, Japan ....................................................................... 447

LOW TEMPERATURE GROWTH OF NANOSTRUCTURED DIAMOND FILMS ON METALS

Paal A. Baker, Sbane A, Cadedge, an'l Y0gesh K Vahra_ Univers:k:_ of A!abama at B innmgham_ U. S.A ....... 452

THE TRIBOLOGY OF AMORPHOUS DIAMOND

Shenq Y. Luo, Huafan University, Taiwan, R.O.C.; Jui-Kang Kuo, _rian Yeh, and James C. Sung, KINIK

Company, Taiwan, R.O.C.; Chuang-Wen Dai, Tamkang Universiw, Taiwan, R.O.C.; and Tsung J. Tsai.

Huafan University .......................................................................... 453

PEEM AND NEXAFS ANALYSIS OF WEAR DEBRIS OF Si INCORPORATED DIAMOND LIKE

CARBON FILMS IN VARIOUS ENVIRONMENTS

Seflm Park, Kwang-Vbng Elm, ami Kwang-Ryeo! Lee, Korea _nstitute of Sciepce m_d Technoh;gy, Korea;

and An_'eas Scholl, FrithjofNelting, m'.d H. Padmore, University ef _erkeley, t,.S.A ..................... 454

6. DEPOSITION AND CHARACTERIZATION OF DLC AND NOVEL CARBONS

6.1 Nitrogen-, Silicon-, and Metal-Containing DLC Films

MOLECULAR DYNAMICS SIMULATIONS OF THE STRUCTURE OF AMORPHOUS CARBON

CONTAINING NITROGEN

D.R. McKenzie, University ef Sydney, Australia; and A. Merchant and D.G. McCulloch, R.M.I.T.

University Melbourne, Austraiia ............................................................... 457

NASA/CP--2001-210948 xviii August 6, 2001
EFFECTOFNITROGEN1NAMORPHOUSCARBONFILMSDEPOSITEDBYELECTRONCYCLOTRON

RESONANCEPLASMASPUTTERINGMETHOD

T_OhanaandT.Nakamura, Nar.ional instltu_eofAdvancedIndustrial ScienceandTechnology, Japan;

A.TanakaandY:K_;ga, National InstimleofAdvancedIndustrial Scienceand'_ehnologyandFrontier

CarbonTechnology Pn;jecl/)apan FisCeramicsCenter, Japat_; m_dA.Golo,FroDtier CarbonTechnology

Pn_eci,/Japan FineCeram,csCentec, Japat_....................................................... 458

AN INVESTIGATION OF Si-DLC FILM STRUCTURE AND CHEMISTRY RELATIONSHIPS TO

DEPOSITION PROCESS CONDITIONS AND MECHANICAL PROPERTIES

Ryan D. Ew, ns, Jeffrey T. Glass, and Phiiip W. Morrison, Jr., Case Western Reserve University, U.S.A.;

and Gary" L. Doll, The Timken Company, US.A ................................................... 460

METALORGANIC CHEMICAL VAPOR DEPOSITION OF RUTHENIUM-DOPED DIAMONDLIKE

CARBON FILMS

M.K. Sunkara and M. Ueno, Ur, ive_sity of l_,m_isville, _,'.S,A.; and G. Lian m_d E.C. Dickey, Un,versity

of Kentucky, _,,.S.A ......................................................................... 466

NICKEL-CONTAINING CARBON FILMS DEPOSITED USING THE SCREEN GRID TECHNIQUE IN

AN ELECTRON CYCLOTRON RESONANCE CHEMICAL VAPOR DEPOSITION SYSTEM

Q.F. Huang, S.F. ¥oon, Rusli, Q Zhang, and J. Ahn, Nanyang Technological University, Singapore; and

E,J. Teo, T Os_powicz, and F. WatL National 1,}niversity of Singapore, Singapore .......................... 473

6.2 DLC Deposition Processes

GROWTH OF DLC AND BCN FILMS BY MAGNETRON SPUTTERING USING DIAMOND AND

C-BN TARGET

N. Ohtake, H. Momose, and X. D. Yang, Tokyo Institute of Technology, Japan .......................... 479

FORMATION OF a-C THIN FILMS BY PLASMA-BASED ION IMPLANTATION

Toshiya Walanabe_ Frontier Carbon Teehnolo©" project/Japan Fine Ceramics Center, Japan; and Kam_b_ro

Yammn,'_lv,, Yoshinori Kega, m_d Aldbm) Tatmka, National Institute _;{'Advanced Industrial Scieace and

Tectn-ology, Japan .......................................................................... 485

HIGHLY ADHESIVE DIAMOND-LIKE CARBON FILM BY ION MIXING METHODS

Kenji Fuchigarni, Kats_mi Suzuki, Masaaki Horiuchi, Kazuo Uematsu, and Tadashi Sasa, IshikawajimaHarima Heavy IndusTries Co., L_d., Japan; and Koichiro Wam,,,mi, Frontier Carbon Technology Project/'

Japan Fine Ceramics Center, Japan ............................................................. 486

PREPARATION OF DIAMONDLIKE CARBON FILMS BY HIGH-INTENSITY PULSED-ION-BEAM

DEPOSITION

X.X. Mei and T.C. Ma, Dalian University of Technology, Cl-ipa ...................................... 491

DEPOSITION OF DIAMOND-LIKE CARBON FILM IN A CLOSED-SPACE CH 4 rf PLASMA

N. Mutsu!cara and Y. Handa, Tokyo Den_i University, Jaaan ......................................... 492

A METHOD FOR DEPOSITION OF AMORPHOUS CARBON

Jawad Haidar, CSIRO 'Felecornmlmieations and In_h_stnal Physics, Australia ........................... 498

THE INFLUENCE OF SUBSTRATE BIAS AND ARGON FLOW ON THE PROPERTIES OF AMORPHOUS

DIAMOND DEPOSITED BY CATHODIC ARC

Wen Ting Yeh, KINIK Company and National Taipei University of Technology, Taiwan, R.O.C.; and

Sea-Fue Wang, National Taipei University of Technology, Taiwan, R.O.C.; and James Sung, KINiK

Company and National Taipei Universizy of Technology, R.O.C ...................................... 504

NASA/CP--2001-210948 xix August 6, 2001
STRUCTUREANDMECHANICALPROPERTIES OFCARBONFILMSPREPAREDBYMASSSEPARATEDIONBEAMDEPOSITION

KazuhiroYamamoto, Yoshlnori Koga,andShuzoF_.ljiwara, National instituteofAdvancedindustrial

ScienceandTechnology, Japan; and_bshiyaWai:anabe, KoicbiroWazumi, andFunfioKol,ai,FronOer

CarbonTechnology Pn;jecl/]apan FisCeramicsCenter, Japan........................................ 505

ONTHEEFFICIENCYOFPLASMAFILTERSFORTHECATHODICARCSYNTHESISOFDLCFILMS

I.I.Aksenov, D.Yu.Zaleskij, andV.E.Strel'nitskij, National ScienceCenter"KharkovInstituteofPhysics

andTechnolog3_" Ukraine.................................................................... 509

AMORPHOUSDIAMONDEPOSITEDBYCATHODICARCWITHCONTROLLEDASPERITY

Ming-Chi KanandJaw-LayHaung, National Cheng-Kung Universky, Taiwan, RO.C.;andWenTing

YahandJamesS'_mg, KINIKCompanyat_dNational TaipeiUniversityofTechnology; Taiwan, R.O.C........ 516

ELASTICONSTANTSONNANOMETRICDIAMOND-LIKECARBONFILMSBY SURFACE

BRILLOU1N SCATTERING

M.G. Beghi, C.E. Bottani, and A. LiBassi, Politecnico di Milano, Italy; B.K. Tanner, University of

Durham, U.K.; andA.C. Fen-ari and J. Robertson, Cambridge University, U.K ........................... 521

STRUCTURE ANALYSIS OF TETRAHEDRAL AMORPHOUS CARBON FILMS USING SYNCHROTRON

RADIATION LIGHT SOURCE

Chin!- Seung Lee, Kwang Yimg EmL and Kwang..Ryeol Lee, Korea Institute e{' Science and Tecbm;iogy,

Korea; KL Hyur, Y.'nm, Vbnsei University, Korea; and E. Gullicksm- apd H. Padroo_e, Universi_' _;f

Berke!e_/, U.S.A ............................................................................ 522

6.3 Novel Carbon Materials

NANO-STRUCTURAL PROPERTIES OF CARBON MATERIALS OBTAINED FROM ORGANIC

COMPOUNDS

Mutsumasa Kyotani and Cbiharu Yamagucb_, Froni:ier Carbon Technology Laboralory/Japan F_ne Ceramics

Center and Osaka Gas Co., Japan; Aklko Gala, Frontier Cad)on Technology Laboratory/Japan Free

Ceramics Centre; Japan; Katsdh'.ro Sasaki, Fr{,ptier Caxb{m Tecbn{}i{,gy Lab{}rak,ry/Japap Fine Ceramics

Center and Osaka Gas Co., Japan; Hisaji Matsui, Osaka Gas Co., Japm'4 Kazuo Akagi, UniversiW of

Tsukuba, Japan; and Yoshinori Koga, National Institute of Materials and Chemical Institute, Japan .......... 523

CARBONIZATION OF POLYACETYLENE BY PYROLYSIS

Addko Goto, Mutsamasa Kyotani, Kazuo Tsugawa, and ¥oshinori Koga, National institute ()fAdvanced

In_bxstrial Science and Technology, Japan; and Kaz'ao Akagi and G_Jangzhe Piao, Universfiy of

Tsuku ha, _apan .................................................................................. 528

A FLORID CARBON STRUCTURE

Y.K. Yap, M. Yamaoka, M. Yoshimnra, Y. Mori, and T. Sasaki, Osaka University, Japan ................... 533

SOLID LUBRICANT PROPERTIES OF CARBON ONIONS PREPARED FROM DIAMOND CLUSTERS

AND PARTICLES

A. Hirata, M. _;ga_ashi, and T. Kaito, Tokyo Institute of Technology, Japan .............................. 538

AB-INITIO DESCRIPTION AND PREDICTION OF PROPERTIES OF CARBON-BASED AND

OTHER NON-METALLIC MATERIALS

D. Bagayoke, G.L. Zhao, and S. Hasan, Southern University/and A&M College, U.S.A ................... 544

NASA/CP--2001-210948 xx August 6, 2001
INVESTIGATIONFDEFECTSTRUCTURE1NMULTILAYEREDCAGECARBONUSINGX-RAY

EMISSIONSPECTROSCOPY

L,G.BulushevaandA.V.Okorzub, InstituteofinorganicChemistrySBRAS,Russia; _,_L. Kuznetsov

andYu.VButenko, BoreskovInslm_teofCalalys_s SBRA,S,R_ssia; A. Fonseca_ Facuiles Universitaires

Notre-Dm:le de la Paix, Belgim:l; and M.I. Heggie, Universily of Sussex, U.K ............................. 550

HRTEM STUDY OF CARBON AGGREGATES UNDER IRRADIATION

J.C. Amault, M. Romeo, G. Ehret, and F. Le Nornaand, Instimt de Physique et Chimie de Strasbom'g,

IPCM5-G SI, France; and F. Banhart and P. Kohier-Redlich, MPI Stuttgart, Gen'nany . .................... 551

SYSTEMATIC STUDY OF CONDITIONS FOR DEPOSITION OF VARIOUS CARBON STRUCTURES

PREPARED BY ION BEAMS

N.G. Shang, F Y Meng_ Q l_,i, X.F. Duan, Y. Lifshitz, C S Lee, S.T. Lee, and I. Bellc_ City University of

Hong Kong, China ............................................................................... 552

ENERGY-CONTROLLED DISCONTINUOUS STRUCTURES IN CARBON MATERIALS

D. Shvets, Institute for Sorption and Problems of Endoecology of NAS Ukraine, Ukraine ................. 553

POSSIBILITIES OF IDENTIFICATION OF ADVENTITIOUS CARBON ON A CARBON SURFACE BY

XPS AND XAES

A.P Dementjev, K.I. Maslak_;_; A. V Naumkm, and O.O. Zabusov, IRTM, RRC Kurcbamv Institute, Russia ,.. 554

7. NITRIDES, CARBIDES, AND B-C-N MATERIALS

7.1 Cubic Boron Nitride

SYNTHESIS OF cBN FILMS BY PLASMA JET CVD IN AN Ar-N2-BF3-H 2 GAS SYSTEM

S. Mar_sumoto, W,J. Zhang, and J. Yu, National Instimrx_ for Research in _;norganic Materials, Japan .......... 56 i

STRUCTURE VARIATION OF CUBIC BORON NITRIDE FILMS IN A TWO-STEP PROCESS BY

CHANGING THE BIAS VOLTAGE AND GAS COMPOSITION

W.J. Zhang, City University of Hong Keng, China; S. Matsumoto, Nationa! Institute for Materials

Science, Japan; and Quan. Li, S.T. Lee, and I. Bello, City University ef 1tong Kong, China ................ 565

ION BEAM DEPOSITION AND SURFACE PROCESSING OF CUBIC BORON NITRIDE

C. Ronning, H, Felde_ann, and H Hofs'ass, University of G6ttingem Germany ......................... 571

ELECTRON FIELD EMISSION OF CUBIC BORON NITRIDE THIN FILMS

P.A, Epcarnacx0t_, Y, Sheog, A. Danie!, an'l R, Cial_ke, The {Jnxversi*_ of Michigan, U.S.A ................. 577

RECTIFYING CHARACTERISTICS OF W/p-type c-BN FILMS

K.J. Liao, W.L. Wang, C.G. Hu, and J. Xu, Chongqing University, RR. China ........................... 579

SOME PROPERTIES AND CUTTING PERFORMANCE OF POLYCRYSTALLINE CUBIC BORON

NITRIDE WITHOUT ADDITIVES

Tadakazu Ohashi_ Hktetoshi Nakajima, _tkfichi Hamada, Katswoshi Omino, at_d Kazuo Yammru;to,

Mitsut-ish_ Materials Coq'oration CeDtrN Research institl,lt% Japan ..................................... 583

CUBIC BORON NITRIDE COMPOSITE COATING FOR ADVANCED MACHINING APPLICATION

Ajay R Malshe, S.N. Yedave, and W.D. Brewl,, University of Arkansas, U.S.A.; and William Russell,

Valenite Inc., U.S.A ......................................................................... 589

NASA/CP--2001-210948 xxi August 6, 2001
SUPERHARDCOMPOSITES BASEDONCUBICBORONITRIDE

AnatolyBochko, insututeofMetal'_.ur_" andMaterialsofNational AcademyofScienceofRussia,

Russia;ndGueroldBobrovnitchii, LaborazoryfAdvancedMazerialsofStateUniversityofthe

NorthFbJminense, Brazil........................................................................ 590

7.2 Carbon Nitride

REACTIVE MAGNETRON SPUTTERING OF HARD C-N and Si-C-N FILMS AND THEIR PROPERTIES

J. Vicek, M. Kotmunda, J. Cizek, and V. Hajek, University of West Bohemia, Czech Republic .............. 595

THE SPECULATION OF SUPERDIAMONDS

James Sung, YdNIK Company and National Taipei University of Technology, Taiwan RO.C ............... 596

SYNTHESIS AND DIELECTRIC PROPERTIES OF CARBON NITRIDE THIN FILMS

Wiilian- C. Lanter, hroovalxve Scier, tific Solutions C,'_]_., U.S.&; l-3avid C Iograro, Ohio Upive:csity,

U.S.A.; Chides A. DeJoseph, U.S. Air Force, Wright-Pattersol'_ Air Force Base, U.S.A.; al'_d Richard L.C.

Wu, K Systems Corp., U.S.A .................................................................. 597

STRUCTURAL PROPERTIES OF CARBON NITRIDE FILMS DEPOSITED BY REACTIVE PULSED

LASER DEPOSITION TECHNIQUE

A.R. Pbani and LE. KrzaDowsld, Universii:y of New Hampsbire, U,S,A .................................. 598

MICRO-WEAR PROPERTIES OF ION-PLATED CARBON NITRIDE THIN FILMS

Shuichi Watanabe, Shojiro Miyake, and Masao Mm'akawa, Nippon h'_stitate of Techno!ogy, Japan .......... 599

DEMONSTRATION OF GROWTH OF ULTRA-HARD CARBON NITRIDES FILMS BY HOT-FILAMENT

CHEMICAL VAPOR DEPOSITION

Saibal Mitra, Justin Ha_man, and Tobias Hanrath, University of Tulsa, U S_A ........................... 604

PROCESS PARAMETERS FOR SYNTHESIS OF CARBON NITRIDE THIN FILMS USING ELECTRON

CYCLOTRON RESONANCE PLASMA SPUTTER METHOD

Eifi Kan-:ijo, Yoshifuroi Ao'_, Talsuya Hh:al-ara, ar,'l 'rakashi 'iochm-oto, Rs'uk,'_kd Universits' _ Japan .......... 609

OPTICAL EMISSION STUDIED TO THE CARBON NITRIDE FILM DEPOSITION BY ECR PLASMA

SOURCE ENHANCED PLANAR MAGNETRON SPUTTERING

Jun Xu, Tencai Ma, Jialiang Zhang, Xianxiu Mei, and Xinlu Deng, Dalian University _f Technology,

RR China ................................................................................ 615

AN XPS STUDY OF PULSED LASER DEPOSITED CN x FILMS

F. Le Nommnd and 3. HorameL lnstitut tie PhysNue et Cbin-ie de Slrasbolr% IPCMS-GSI, Frm-ce;

Y. Szorenyi, CNRS-PHASE, France, and Research Group on Laser Physics, Hungary; and C. Fuchs and

E. Fogarassy, CNRS-PIIASE, Frap_ce ........................................................... 616

C3N 4 CONTROVERSIALS AND THE CHINESE CONNECTION

James C. S_mg, KIN1K Company and National Taipei University of Technology, Talwan, R.OC ............ 617

7.3 Silicon Carbide and Boron Carbide

IMPROVED INVERSION CHANNEL MOBILITY FOR 4H-SiC MOSFETS FOLLOWING HIGH

TEMPERATURE ANNEALS IN NITRIC OXIDE

G.Y. Chung, C.C. Tin, and J.R. Willian:s, Auburn University, U.S.A.; K. McDop_ald, R.K. Chanap_a,

R.A. Waller, S.T. Pantelides, L.C. Feldrnan, Vanderbilt University, U.S.A.; and O.W. Holland,

M.K. Das, and J.W. Palmour, Cree Research Inc., U.S.A ............................................ 618

NASA/CP--2001-210948 xxii August 6, 2001
STUDYOFGAPSTATEDISTRIBUTIONS 1NHYDROGENATED AMORPHOUSSILICONCARBIDE

GROWNUSINGECR-CVD

K.Chew, Rusli,S_F_Yoon,J.Ahn,Q.Zhang, andVLigatche_, NanyangTechnological University,

RepublicofSingapore; andE.L'Fec_ T. Oslpowie% and F. \Vatl, National University _;f Singapore,

Republic of Singapore ........................................................................... 621

THE CARBON IMPLANTATION IN SILICON BY CATHODIC ARC

Wen Ting Yeh, KE'_IK Company and National Taipei University of Technology, Taiwan, R.O.C.;

Ming-Chi Kan. National Cheng-Kung University, Taiwan, R.O.C.; Sea-Fue Wang, National Tai-oei

University of Technology, Taiwan, R.O.C.: and James Sung, KINIK Company and National Taipei

University of Technology; Taiwan, R.O.C ........................................................ 627

ELECTROMAGNETICALLYACCELERATED PLASMA SPRAYING FOR BORON CARBIDE COATING

Kilarou_a, Frot_tier Carbot_ 2_chnolog?¢ P_oject/Japat_ Floe Ceramics Cenle_ and National IostJtu_e o_'

Advar_ced In_h_strial Scierce and Technology, Japan; and S. U suba, Y. Kakutlate, H. Yokci, K. Yaroarooto,

and S. Fuj iwam, National Institute of Advanced Industrial Scier.ce m'_d TechnologiC, Japar. ................. 632

DEPOSITION OF BORON CARBIDE THIN FILMS AND THEIR HARDNESS

F,.anio Kokai and Manabu Taniwa!d, Frontier Carbon Technolo©" Project/Japan Fine Ceramics Center

and National institute of Advanced industrial Science and Technology, Japan; and Masatou ishihara and

YosbiDor_ Koga, National institute of Advanced Industrial Sc_et_ce and Technology, }apart .................. 636

THE SPATIAL-RELATIVE JOINT DENSITY OF STATE FOR AMORPHOUS CARBIDES

Z.Q. Ma, Xinjiang L'niversity, RR. China ........................................................ 638

7.4 B-C-N Films

ELECTRICAL AND OPTICAL PROPERTIES OF TERNARY BCN THIN FILMS

Guanghua Chen, ShenNun Zhang, and Xuemei Song, Beijing Polytechnic University, China; and

Jinshon Yue_ Lanzhou Universily, RR China ......................................................... 642

MONATOMIC STRUCTURES OF B, C, N, and O: FIRST-PRINCIPLE STUDY OF RELATIVE

STABILITIES AND BULK MODULI

James C.M. Sung, NatienaI Taipei University of Technology, Taiwan, R.O.C.; and B.R. Wu, S.L. Lee,

and M.F. Tai, National Chung Cheng University, Taiwan R.O.C ...................................... 64,3

8. CARBON NANOTUBE PRODUCTION, PROPERTIES, AND APPLICATIONS

8.1 Carbon Nanotube Production and Applications

RECENT PROGRESS IN DEVELOPMENT OF LARGE SCALE SYNTHESIS OF MULTI-WALLED

CARBON NANOTUBES OF THE FRONTIER CARBON PROJECTS

Motoo Yumura, Satoshi Ohshima, HirokA Ago, and Kunio Uchida, National Institute of Materials and

Chemical Research, Japan; and Hitoshi lnoue and Toshiki Komatsu, Jaaan Fine Ceramics Center, Japan ...... 65

CATALYTICALLY GROWN CARBON NANOTUBES AND THEIR APPLICATIONS

Ctmnrning Niu, Bob Ho_:h_ add Howard Tenr_em;, Hyperion Catalysis international, U. S.A .................. 652

NUMERICAL ANALYSIS ON THE DISPERSION PROCESS OF CARBON CLUSTERS IN GAS

EVAPORATION SYNTHESIS USING D.C. ARC DISCHARGE

S. Usuba. Y. Kakudate, H. Yokoi, H. Suz"a!d, and "3. Fujiwara, National Institute of Advanced Industrial

Science and Technology, Japan ................................................................ 657

NASA/CP--2001-210948 xxiii August 6, 2001
8.2 Carbon Nanotube Magnetics

AN ELECTRICALLY CONDUCTING BORON-DOPED CARBON NANOTUBE POLYSTYRENE

COMPOSITE

P.CP. Vvatls, W.K Hsu, H.Vvl Kroto, and D.R.M Waltz;n, Umversilv of Sussex, U.K.; and GZ. Chen

and I-3..LF_ay, University of Cambridge, U.K ..................................................... 661

MAGNETIC COUPLING OF CARBON NANOTUBES

Trinh Vo, Tai-Chou Lee, Behnaz Payandeh, and Mark Robert, Rice University, U.S.A ..................... 667

8.3 Carbon Nanotube Storage Systems and Capacitors

LITHIUM STORAGE IN NANOMATERIALS

O. Zhc-u, _Jmversitv ofNocth CaroUna, L1.S.A ....................................................... 668

ELECTROCHEMICAL HYDROGEN STORAGE IN SINGLEWALLED CARBON NANOTUBES

Won Seek Kim, Kay Hyeok An, Yeung See Park, Kwan Ku Jeon, Hee Jin Jeong, Jeong-Mi Moon, and

Young Hee Lee, Sungksrankwan University, South Korea ........................................... 669

HIGH CAPACITANCE SUPERCAPACITORS USING SINGLEWALLED CARBON NANOTUBE AND

POLYPYRROLE HYBRID ELECTRODES

Kay Hyeok An, Jeong-Mi Moon, Won Seek Kim, Young Min Shin_ add Yi-ut_g Hee Lee, Stmgkyunkwan

Ur, i,_e_sity, Smnh Korea; and Yom_g See Pa_k, Kwan Ktl Jeon, Hee Jip J_;ng, Seoung yel Jeong, S_;ng

Chu Lira, and Dong Jae Bae, Jeonbuk National University, South Korea ............................... 675

8.4 Processing of Carbon Nanotube

HIGH-PRESSURE POLYMERIZATION OF SINGLE WALL CARBON NANOTUBES

M. Popov, M. Kyotam, and v Koga, F_ontler Carbon Technology Research/Japan Fine Ceramics

Center, Japan; and R.I Nemanieh, North Carolina State Umversilv, U.S. A ............................... 681

COVALENT SIDEWALL FUNCTIONALIZATION OF SINGLE WALL CARBON NANOTUBES

I.W. Chiang, R.K. Saini, E.T. Mickelson, \_:E. Billups, R.H. Hauge, and J.L. Margrave, Rice

University, U.S.A ........................................................................... 687

LOW-TEMPERATURE FLUORINATION OF CARBON NANOTUBES

A.V. Okotrub, N.F. Yudanov, L.G. Bulusheva, A.I. Romapenko, ipstinlte of inorganic Chemistry

SB RAS, Russia; and A.L. CN_wlin, Boreskov institum of Cai:alysis SB RAS, Russia ....................... 694

SYSTEMATIC STUDY FOR PURIFICATION OF SINGLEWALLED CARBON NANOTUBES

Jeong-Mi Moon, Kay Hyeok An, Chulsu 7o, _(bung Hee Lee, Sungkymlkwan University, Korea;

Young See Park, Hee Jin Jeong, Dong Jae Bae, and Seong Chu Lirn, Jeonbuk National University,

Korea; ai'ld Gyeong-Su Park, Sainsurlg Advanced Institute of Technology, Korea ........................ 695

8.5 Sensors and Gas Separation

CARBON NANOTUBE BIOSENSORS FOR SPACE MOLECULE DETECTION AND CLINICAL

MOLECULAR DIAGNOSTICS

._ie P2an, NASA Ames Research Centeg L',S,A .................................................... 700

SEPARATION OF CO 2 FROM FLUE GASES BY CARBON-MULTIWALL CARBON NANOTUBE

MEMBRANES

Rodney Andrews, Merit Jagtoyen, and Eric Grulke, University of Kentucky, U.S.A.; and Ki-Ho Lee,

Zugang Mao, and Susan B. Sinnott, University of Florida, U.S.A ..................................... 701

NASA/CP--2001-210948 xxiv August 6, 2001
8.6 Electron Emitters

NEW EMITTER TECHNIQUES FOR FIELD EMITTER DISPLAYS

N.S. Lee, et ai., Sarnslmg Advat_ced Instimle of Tecbnology, Korea ..................................... 707

FUNDAMENTAL PROPERTIES OF NANOTUBE FIELD EMITTERS FOR LARGE AREA ELECTRON

SOURCES

Kennefl_ A. Dean, Babu R. Chalamala, Bernard F. Coll, A. Alec Talin, Johann Trujillo, Yi Wei, and

James E. Jaskie. Motorola Inc., U.S.A ........................................................... 708

MULTIWALLED CARBON NANOTUBES GROWTH IN NANO-TEMPLATE AND THEIR

APPLICATION TO A FIELD EMISSION DEVICE

Soo-Hwan ,_eong_ Hee Young Hwat_g, and K_m_Hong Lee, Pohang University of Seience and

Technology; and Won-Ki Cho and Kwarlg..Yi;ung Kim, LG Electronk:s inc., Korea .................... 716

FIELD EMISSION FROM CARBON NANOTUBES GROWN ON PLASMA TREATED NICKEL/

SILICON SUBSTRATE

W.K. Wong, Frederick C.K. Au, I. Bello, C.S. Lee, and S.T. Lee, City University of Hong Kong, U.S.A ....... 717

FIELD EMISSION PROPERTIES OF ULTRA-FINE MULTI-WALLED CARBON NANOTUBES

GROWN BY MICROWAVE PLASMA CVD

Weihai Fu, Biog Xiae, at_d Sacharia Alt-in_ Ol'l Dorninion University, U.S.A ............................ 7 l 8

CHARACTERISTICS OF CARBON NANOTUBES SYNTHESIZED BY LASER ABLATION PROCESS

Ilsiu-Fung Cheng, National Taiwan Nonnal University, Taiwan; Teng-Fang Kno, National Tsing-Hua

University, Taiwan; Tzai-Hsin Lai, Chung-Yuan Christian University, Taiwan; Chun-Homg Tsai, National

Tsing Ilua University, Taiwan; and I-Nan Lin, National Tsing-Hua University, Taiwan .................... 724

FIELD EMISSION FROM CARBON NANOTUBES GROWN BY THERMAL CVD ON NiCr COATED Si

Bing Xiao, Weibai Fu, and Sacharia Albm, Old Dominion i,q_iversiW, U.S.A ............................ 726

ALIGNED AND OPENED CARBON NANOTUBE ARRAYS AND THEIR FIELD EMISSION PROPERTIES

Zheng Wei Pro'. m'.d Zhong L. Wang, Georgia Institute of Technology, U.S.A.; S.T. Lee, Cit_/University

of Hong Kol,.g, P.R. China; and S.S. Xie, Chinese Academy of Sciences, P.R. China ...................... 730

MICROWAVE PLASMA ASSISTED BRAZING OF CARBON NANOTUBES AND DEPOSITION OF

CARBON FILMS ON IRON ELECTRODES FOR APPLICATIONS AS ELECTRON FIELD EMITTERS

Yophua Tzeng, Chao Liu, and Calvin Cutsbaw, Aubt_ Universfiy, U.S. A ................................ 736

HYSTERESIS OF ELECTRON FIELD EMISSION FROM SINGLE-WALLED CARBON NANOTUBES

BRAZED ON IRON SUBSTRATES

Chao Liu, Calvin Cutshaw, and Yonlraa Tzeng, Auburn University, U.S.A .............................. 742

SECONDARY ELECTRON EMISSION OF MAGNESIUM OXIDE ON MULTIWALLED CARBON

NANOTUBES

Won Seok Kim, Young Min Shin, Hee Jin Jeong, and Yo'_mg Hee Lee, Stmgkyunkwan University,

South Korea; Whikun Yi, SeGJ _l, Taewon Jeong_ Jeonghee Lee, Sunghwan Jin, Jlmgna Heo, and

M. KJrn, Samsung Advansed Institute of Technoh-gy, Korea .......................................... 747

WORK FUNCTION AT THE TIPS OF MULTIWALLED CARBON NANOTUBES

I- 1 I- 1 1 R.P. Gao, Z.W. Pap, and Zhong L. Wang, ,aeorgia InsJtute of Techno_og3, U.S.A ........................ 752

NASA/CP--2001-210948 xxv August 6, 2001
8.7 Electrical and Electronics of Carbon Nanotube

CHARGE TRANSFER, STRUCTURE, AND ELECTRONIC PROPERTIES OF K-DOPED SINGLE-WALL

CARBON NANOTUBES

Cbulsu _o aad "_2nmg Hoe Lee, Sung!<yut_l,:v_an Universily, Korea ....................................... 753

ELECTRICAL RESISTIVITY OF ANISOTROPICALLY ALIGNED SINGLEWALLED NANOTUBE FILM

Dong Jae Bae, Kae Hyeok AI,, Jeong Mi Moon, Seol,_g Chu Lira, and Young Hoe Lee, Sungkyunk_a,an

University, Repubiic of Korea; and Jeonbuk National University, Republic of Korea ...................... 758

OXIDATION OF CARBON NANOTUBES AND CHANGE OF THEIR ELECTRONIC AND GEOMETRIC

STRUCTURES

Seong C_m Lira, Hoe Jin Jlmg, Dong Jae Bae, and Eun Kyung Suh, Jeonbuk National Umversily, South

Korea; Yi-ut_g Mm Shin, K_y HyeokAt_, m_d Young Hoe Lee_ Slmgkylmkwat_ University, South K,'_rea;

.)aewa Churl, Louis:ana State Ur, i,_e_sity, U.S.A .................................................... 759

ELECTRONIC STRUCTURE MODIFICATION OF MULTIWALLED CARBON NANOTUBES BY ION

BOMBARDMENT AND THERMAL TREATMENT

Jaewu Choi, Louisiana State University, U.S.A.; Young Chul Choi, Samsung SDI, Korea; and

Young Hee Lee, Sungky',mkwan University, Korea ................................................ 764

8.8 Carbon Nanotube Composites

WORKING TOWARD NANOTUBE COMPOSITES

Sivaram Arepalli, Pavel Nikoiaev, and Olga Gorelik, GB TecK/Lockheed Martin, U.S.A.; Victor G.

Itadjiev, University of Houston, U.S.A.; and Carl D. Scott and Bradley S. Flies, NASA Johnson Space

Cen_.er, U.S.A .............................................................................. 769

COMPOSITE MATERIALS FROM MODIFIED CARBON NANOTUBES

l_;dney Andrews and Terry Rantell, University of Kentucky_ tl. SA.; Robert C. Haddon, Universfiy of

Cahtbraia, Riverside, U.S.A ; and Elizabeth C. Dickey, R_,t-eO; G Bergosh, Hui Hu, Chad Lar_d_s, m_d

Mark S. Meier, Ur, i,eersit y of Kentucky, U.S.A .................................................... 7'71

8.9 Mechanical Properties and Applications

MECHANICS OF CARBON NANOTUBES

Rod Ruoff_ Noehwestem University, U.S.A ...................................................... 772

FRACTURE STRENGTH OF CARBON NANOTUBES IN HYDROGEN ENVIRONMENT

Long Gaang Zhou, Chinese Academy of Sc:ences, China; and San-Qiang Shi, Hong Kong Polytechr, ic

Uni,eersity, China ........................................................................... 773

FRACTURE STRESS ENHANCEMENT BY EB TREATMENT OF CARBON FIBER

Atsushi Kimura, Akihiro Mizutani, Takashi Toriyama, Kazuya Oguri, Akira Tonegawa, and Yoshitake

Ni_hi, Tokai University; Japan ................................................................. 779

9. GROWTHAND MODELING OF CARBON NANOTUBES AND NANOSTRUCTURES

9.1 Carbon Nanotubes and Related Materials and Structures

COMPUTATIONAL STUDIES ON FORMATION AND PROPERTIES OF CARBON NANOTUBES

Tahir _agin, Jianwei Che, Weiqiao Deng, Xin Xu, and William A. Goddard IH, California Institute of

Technology; U.SA .......................................................................... 787

NASA/CP--2001-210948 xxvi August 6, 2001
NANOMACHINEDCARBON: NANOTUBES, FULLERENES, PEAPODS

David Tomfinek, Michigan State University, U.S.A ................................................ 793

9.2 Carbon Nanotube Synthesis and Patterning

GROWTH OF CARBON NANOTUBES ON PATTERNED CATALYSTS: TOWARDS A CARBON

NANOTECHNOLOGY

G.A.J. An:aratunga, et al., Cambridge University, U.K.; and J.J. Papworth and C.J Kiely, Universi,q/of

Liverpool, U.K ............................................................................. 794

SURFACE DECOMPOSITION MECHANISMS ON SiC (0001)

M. Kusunoki, T. Suzuld, To Hirayama, and N, Shlbar_a, Frontier Carbon Technology Project/Japan Fine

Ceramics Center, Japan ........................................................................... 795

CARBON NANOTUBES AND CARBON NANOFIBERS SYNTHESIZED BY SUBLIMATING

DECOMPOSITION OF SILICON CARBIDE WITH CATALYSTS

Kaz-ao Tsugawa and Akiko Gore, Frontier Carbon Technology Project/Japan Fine Ceramics Center,

Japan; and Yoshinori Koga, National Institute of Advanced Industrial Science aed Technology, Japae ........ 800

SYNTHESIS OF CARBON NANOTUBES BY MICROWAVE HEATING

Eot) Hwa Hot, g, Klm_Hong Lee, Sang Ho Oh, Chat_-Gyung Park, and Cbang-Mo Ryu, Pohang University

of Scien_:e and Enginemmg, Kore, a; and Jong Hur, klan and _ae Eun Ryu, Iljkl NaDolech Co., Ltd., Korea ..... 805

MULTI-WALLED CARBON NANOTUBES FROM ETHYLENE DIFFUSION FLAMES

Liming Yuan and Kozo Saito, University of Kentucky, U.S.A.; and Wenchong Hu and Zhi Chen,

University at California at San Diego, U.S.A ..................................................... 810

OPTICAL PLASMA CONTROL DURING ARC CARBON NANOTUBE GROWTH

L Hinkov and S. Farhat, L_MHP Universi_2 Paris, France; M. Lamy De La Chapelle and S.S. Fan,

Tsinghlm University, China; H X. Han and G.H. Li_ Chinese Academy of Sciences, China; and C D. Scott,

NASA l_,yndon B Vohnseo Space Center, U.S.A ....................................................... 816

BOUNDARY-LAYER CONTROL DURING CARBON NANOTUBE GROWTH BY ARC METHOD

Samir Farhat, Ul,.iversity of Paris, France ........................................................ 822

GROWTH OF ALIGNED CARBON NANOTUBES ON SILICON BY RF-PLASMAASSISTED

PULSED-LASER DEPOSITION

YK. Yap, M. Yamaoka, M. Yosbimura, Y. Mori_ T. Sasak_, T. Hanada, H. Fumta_ T. H_rao, S. Honda,

and K. Oura, Osaka Ur, ive_sity, ,lapar, ............................................................... 823

DIRECT SYNTHESIS OF CARBON NANOTUBES ON PLASTIC SUBSTRATES

Etm Hwa Hong, Kun-Hong Lee, Sang Ho Oh, Chan-Gyung Park, and Chang-Mo R?_a, Pohang University

of Science and Engineering, South Korea; and Jong Hun Han and Jae Eun Ryu, Iljin Nanotech Co., Ltd.,

South Korea ............................................................................... 828

A NOVEL APPROACH TO GROWTH OF LARGE SCALE ALIGNED CARBON NANOTUBES

OR CARBON NANOSHEETS ON Si SUBSTRATE BY ECR CHEMICAL VAPOR DEPOSITION

Chao Hsm_ Lin, Hoi Lin Cbang_ Mmg Her Tsai, m_d Cheng Tzu KI10, National Chiao Ttmg

University, 'l-aiwar, .......................................................................... 829

RADIAL AND VERTICAL VAPOR GROWN CARBON FIBERS (VGCF)

Maheshwar Sharon and Debabrata Pra&han, Indiae Institute of Technology, India ........................ 830

NASA/CP--2001-210948 xxvii August 6, 2001

1

REVIEW OF RESEARCH PROGRAMS

Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

NSF Funded Research Programs on Hard Coatings and Carbon Nanotechnology

Jorn Larsen-Basse

Surface Engineering and Material Design Program

National Science Foundation, 4201 Wilson Blvd., Arlington, VA 22230, USA.

ABSTRACT

The National Science Foundation supports basic research in materials science and engineering through a

number of different divisions and programs and in a number of different funding modes. As part of that

activity, research in surface engineering has been growing quite rapidly during the past ten years, much of

it focused on hard coatings for tribological applications which will be a primary focus of this presentation.

NSF funds primarily university researchers. Most projects are quite small and are carried out by a

professor with one or two graduate students, but small groups and even quite large centers are also

supported, with many of the centers focused on encouraging industry-university collaboration in research.

Projects are selected for support based on research proposals, most of them unsolicited. They are evaluated

by other members of the research community. This system is intended to bring forth the best quality basic

research, with no specific "mission" application, as determined by the research community and not by

bureaucrats in Washington.

In this paper I will summarize a sampling of past and current research projects on diamond and DLC

coatings, on hard coatings in general, and on carbon nanotubes and carbon nanotechnology. The topics

cover the range from materials science and processing, to mechanics of thin film and non-destructive

testing. Some of the examples are:

in-situ sensors for CVD process control

use of buckyballs as boundary lubricant films

use of buckyballs in "shot peening" of growing optical films

deposition and properties of carbon nitride films

ion beam implantation for various property modifications

wear of C* films on hard disks

nanolayered films for improved hardness

polishing of CVD diamond films

measuring properties of coating/substrate interface with shear waves

nanoindentation and nanoscratching to determine mechanical properties of thin films

mechanics of structured and layered thin films

nanomechanics

preparation of carbonaceous and diamond films from SiC

techniques for determining internal stresses in thin films

In the past few years, a major national initiative has been launched on the general topic area of

nanotechnology. NSF is a major participant in this activity and a large number of proposals have recently

been evaluated. A number of those deal with carbon nanotubes in a variety of intended applications, from

fibers in composite materials to heat transfer films, and with a number of ways of making, testing, and

manipulating the fibers. Some of the recently initiated research projects in this area will also be discussed.

The nanotechnology initiative has generated much public interest and it appears that it will receive

significant funding for several years to come. This should provide excellent opportunities for researchers

in the area of carbon and other nanotubes as well as other forms of carbon nanotechnology.

Keywords: NSF, hard coatings, nanotubes, nanotechnology initiative, mechanical properties,

nanomechanics

Jorn Larsen-Basse

National Science Foundation

4201 Wilson Boulevard, Rm. 545

Arlington, VA 22230, USA

JlaI serd_a_ast" gev

Fax: +001-703-292-9053, Phone: +001-703-292-7016

NASA/CP--2001-210948 3 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ONR RESEARCH PROGRAMS ON CARBON NANOTUBE

John Pazik

Office of Naval Research, USA

The Solid State and Materials Chemistry program at the Office of Naval Research (ONR) is currently

investing in a number of research activities related to carbon nanotubes. The emphasis is on the science and

technology (S&T) of single walled nanotubes (SWNT) although development of multi walled nanotube S&T

(MWNT) is considered when unique opportunities exist.

The program seeks to address S&T issues in nanotubes from both supply and demand. From the demand

perspective, the program is interested in new concepts and ideas where nanotubes (particularly SWNT) exhibit a

unique and demonstrable capability/property. For example, the program has supported activities to demonstrate the

strength ofnanotubes, examine field emission properties, and demonstrate unique device operation. From a supply

perspective the program is interested in activities that enable the synthesis and use of these materials.

This presentation will provide an overview of current activities and recent accomplishments of ONR

investigators as well as a perspective on the future impact of carbon nanotubes in Navy/Marine Corp systems.

NASA/CP--2001-210948 4 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CURRENT STATUS OF R&D IN FRONTIER CARBON TECHNOLOGY (FCT)

PROJECT - OVERVIEW

M. Yoshikawa

FCT Project/Japan Fine Ceramics Center

Puros-Nishi-Shinbashi, 3F, 2-4-3 Nishi-Shinbashi, Minato-ku, Tokyo 105-0003, Japan

ABSTRACT

This paper will address an overview of current R&D activities in Frontier Carbon Technology (FCT) Project

operated by Japan Fine Ceramics Center (JFCC) under the auspices of New Energy and Industrial Technology

Development Organization (NEDO) and Ministry of Economy, Trade and Industry (METI). The FCT Project

started in 1998 as a 5-year project, but it will be renewed a year earlier, following the entire reorganization of the

Japanese government.

Keywords: FCT Project, diamond, carbon nanotube, diamond CVD reactor

INTRODUCTION

Our interest in R&D is a new class of carbon materials (that we call Frontier Carbons) represented by such

materials as diamond, DLC, carbon nanotube, and fullerene that have new and unique characteristics in mechanical,

electronic, and chemical properties that have not been known for conventional carbon materials. It is thus expected

st

that Frontier Carbons can be new basic materials after metals and ceramics for an industrial use in the 21 century.

The purpose of the FCT Project is to develop basic technologies, common to different industrial sectors, of the

new carbon materials. The Project basically consists of two categories: (1) New materials synthesis and (2) Materials

processing. The latter is further divided into (a) Electrically fimctional materials processing and (b) Mechanically

functional materials processing.

R&D SUBJECTS

The Project is carried out in two Central Research Laboratories: one in the National Institute of Advanced

Industrial Science and Technology (AIST) at Tsukuba, and the other in Osaka University. In addition, there are five

Satellite Research Laboratories in private companies to utilize their facility. The member organizations include 15

private companies, two research organizations as well as 5 national research laboratories, 6 universities, and three

overseas research groups.

Followings are major research subjects in the three groups:

(1) New Materials Syntheses

- Synthesis of new materials by Electron Beam-excited CVD, multiple ion beams, detonation, and laser

ablation

- Large-scale production of carbon nanotube

- Design of reaction paths and materials syntheses

- Synthesis of anti-corrosive films in high temperature environment

(2a) Electrically Functional Materials Processing

- Control of film morphology using 60 kW microwave plasma CVD reactor

- Field emission mechanism from diamond powder

- Fabrication of sharp emitter tips using single crystal diamond

- Synthesis ofhomoepitaxial diamond slab

- Synthesis ofheteroepitaxial film

(2b) Mechanically Functional Materials Processing

- Deposition of graded-composition films

- Film coatings on complex-shape substrates

- Large-area film coatings

NASA/CP--2001-210948 5 August 6, 2001
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Center for Advanced Research Projects

Osaka University

Fig. 1 R&D subjects in the FCT Project

NASA/CP--2001-210948 6 August 6, 2001
LARGE CVD REACTORS

Since one may able to obtain a catalogue that describes the project detail (Please contact: fct.jfcc.kamejima @

nifty.ne.jp), and the latest results will be later presented by the Group Leaders in this conference, we would like

describe two major microwave plasma CVD reactors for diamond deposition under development in the FCT Project.

Figure 2 is a 60 kW 915 MHz CVD reactor made by ASTeX/Seki Technotron. This system is installed in the

Central Research Lab at Osaka University. The maximum substrate size is 6 inches, and the standard operational

conditions are as follows: 60 kW, 100 Torr, 5%CH4/H2, and the substrate temperature 850 °C. The substrate holder

is water-cooled, and the substrate temperature is more precisely controlled by changing the plenum pressure

(between the substrate holder and the substrate). The plasma shape is disk-like and quite uniformly covers the

substrate, though there still exists a temperature gradient across the 6-inch substrate. The system runs quite

smoothly without problems. The chamber diameter is about 1 m.

._--- Source gas

Exaust I._ _ _1 Cao!ingwa!er

Prenunl pressure

co ntro l

Figure 2. 60 kW 915MHz microwave plasma CVD reactor (1)

The second reactor is custom-made by Sumitomo Electric, where the microwave scale is similar to the above.

This reactor was designed to make a thick homoepitaxial diamond layer with a minimal defect density. Thus, the

gas flow and the electromagnetic distribution are carefully evaluated in the design, and it has more mechanisms than

those of Fig. 1 for fine-tuning the growth conditions. This system is installed in Sumitomo Electric.

Microwave

(60 kW, glfi MHz)

Ouartzwinduw _J_- -- Plasma Water-cooled chamber

Viewing port Y _(gold coated)

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Figure 3. 60 kW 915MHz microwave plasma CVD reactor (2)

NASA/CP--2001-210948 7 August 6, 2001
CONCLUSION

Since the functions of Frontier Carbons are diverse, many practical applications can be considered. The

synthesis and processing technologies for Frontier Carbons will give great benefits for new industries. For example,

in the energy field, Frontier Carbons will be used for high-performance lightweight batteries and hydrogen storage

materials. In the electronics, information and communication fields, Frontier Carbons will be used for electron

emitter materials and surface acoustic wave filters. In the machinery, aeronautics, and space fields, they will be

applied for maintenance-free systems and cutting tools due to their abrasion resistant properties. In the optical field,

Frontier Carbons will be used for optical window materials due to their high transparency over a wide wavelength

range. In the chemical and environmental fields, they will be used for anticorrosive coatings. Lastly, in the medical

field, Frontier Carbons will be utilized for artificial medical materials due to their affinity with biological substances.

Furthermore, judging from the electronic, mechanical, optical, and other properties, diamond is expected to be

widely used in many aspects of medical applications.

Finally, it is expected that this presentation will encourage other countries to start new and large-scale projects

for Frontier Carbons and other new materials to further stimulate international collaborations.

This work was supported by FCT Project, which was consigned to JFCC by NEDO.

NASA/CP--2001-210948 8 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ACTIVITIES OF NOVEL MATERIALS SYNTHESIS TECHNOLOGY GROUP IN FRONTIER

CARBON TECHNOLOGY PROJECT

Yoshinori Koga

Research Center for Advanced Carbon Materials, National Institute of Advanced Industrial Science and Technology

AIST Tsukuba Central 5, 1-1-1 Higashi Tsul_ba, Ibaraki-ken, 305-8565, Japan

Keywords: Carbon nanotube, nanofibers, polymerization of nanotube.

In the group of the novel materials synthesis technology, R & D is carried out to synthesize novel carbon

materials such as nanotube, pressure-induced polymer of single carbon nanotube, heterodiamond, carbon nitride etc.

using thermal decomposition, MWCVD and high pressure techniques.

In this paper we report (1) the highly oriented carbon nanotubes by sublimation decomposition from single

crystal of SiC, (2) the highly oriented carbon nanofibers and hollow nanofibers by MWCVD from methane and

argon mixture gas and, (3) the polymerization of single carbon nanotubes by high pressure method. __

(1) Highly Oriented Carbon Nanotubes by Sublimating Decomposition of SiC film: (Dr. M.Kusunoki et al.)

Carbon nanotube (CNT) has novel electric, optical, chemical and mechanical properties according to its unique

configuration. CNTs have been produced by carbon-arc discharge technique, catalytic pyrolysis of hydrocarbons and

condensed phase electrolysis. In many cases CNTs are usually oriented randomly with a wide distribution of

diameter and length. Kusunoki(JFCC) et al. have developed a very simple method of producing highly oriented

CNTs. The alignment method is based on the self-organization by sublimation decomposition of single crystal SiC

in a vacuum at the temperature between 1500 and 1700 °C. Fig. 1 shows the TEM micrograph of CNT film on the

surface of an _-SIC(0001) wafer heated at 1700 °C for 30 min. The formation mechanism of CNT is proposed

according to the results of the observations of high-resolution transmission electron microscopy. In the initial

process of the formation of CNT the graphite nanocaps of 2-5 nm in size were observed. The surface of the SiC is

oxidized by residual oxygen in the chamber and SiO gas and solid carbon are generated. With the evaporation of

SiO gas the graphite sheets are successively deformed to form nanocaps of graphite. And then cylindrical graphite

sheets with the diameter of the nanocaps grow perpendicularly on the SiC. The CNTs grow, eroding the SiC single

crystal.

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FIGURE 1.HIGHL Y ORIENTED CNTS PRODUCED AT 1700 °C FOR 30 MIN.

(2) Highly oriented carbon nanofibers and hollow nanofibers by MWCVD: (Mr. F.Hoshi et al.)

In Figs. 2(a) and 2(b) highly oriented carbon nanofibers and hollow carbon nanofibers grown by MW ECRCVD method using methane and argon mixture gas at a temperature of 550 °C are shown. The carbon nanofibers

and the hollow carbon nanofibers were grown perpendicularly on Si substrates and on Si substrates coated with Ni

catalyst, respectively. From TEM analysis the diameter and length of the nanofibers are about 60 nm and 15/4m ,

1

respectively. Raman spectra of these highly oriented carbon nanofibers showed new bands of 1340 and 1612 cm- of

the first-order Raman scattering and 2660, 2940, and 3220 cm -1 of the second-order Raman scattering. By the

measurement ofXPS CIs band energies of 284.6 for the carbon nanofiber and 284.7 eV for the hollow carbon

nanofiber indicate mainly sp 2 carbon component. Field emission characteristics of the highly oriented carbon

NASA/CP--2001-210948 9 August 6, 2001
nanofibersandhollowcarbonnanofiberswereinvestigatedandthecurrentdensitieswere7.25mA/cm2and

0.69mA/cm2at12.5V/pm,respectively.

" -i.

FIGURE 2. HIGHLY ORIENTED CARBON FIBERS AND CARBON HOLLOW FIBERS PRODUCED AT

550 °C USING ECR-MWCVD METHOD.

(3) Polymerization of single carbon nanotubes by high pressure method: (Dr. M.Popov et al.)

The P-SWNTs were synthesized by applying a shear deformation under load in a diamond anvil cell; the

procedure of stress tensor variation. Single wall carbon nanotubes(SWNT) were pressurized up to 55 GPa.

Experimental evidence of pressure-induced polymerization of SWNT under the pressure were obtained. A new

superhard composed ofpolymerized single wall carbon nanotubes (P-SWNT) has been synthesized which exhibits a

bulk modulus exceeding or comparable with that of diamond.

5 10 15 20 25 30 35 40 45

Displacement into surface (rim)

FIGURE 3. FORCE-DEPTH INDENTATION CUR VES OF P-SWNT (MARKED B Y SOLID CIRCLES),

CUBIC BN FACES (100) (MARKED BY HOLLOW CIRCLES) AND (111) (MARKED BY CROSSES) AND

FUSED SILICA (MARKED B Y HOLLOW TRIANGLE) ARE PLOTTED.

The process of polymerization is accompanied by irreversible changes in the Raman specra. Bulk modulus of 462 to

546 Gpa was found out for P-SWNT from the Raman mode of 1590cm -1. Hardness measured by nanoindentation

method was in between 62 to 150 Gpa from the comparative study of P-SWNT and the single crystals of diamond

and cubic boron nitride.

NASA/CP--2001-210948 10 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

RESEARCH AND DEVELOPMENT OF HARD CARBON FILMS AND THEIR TRIBOLOGICAL

PROPERTIES IN THE FRONTIER CARBON TECHNOLOGY PROJECT

Akihiro TANAKA

Research Center for Advmaced Caxhon Materials, National Institute of Advanced Industrial Science and Technology

AIST Tsukuba Central 5, Tsukuba 305-8565, Japma

e-mail: a-tanaka@aist.go.jp, Fax: +81-298-61-4636, Phone: +81-298-61-7069

ABSTRACT

The "Frontier Carbon Technology (FCT)" project, focused on R&D in processing technologies for materials with

excellent mechanical properties such as low friction, wear resistance, corrosion resistance, and so on, was conducted.

Processing technology for materials with excellent tribological properties was developed by both antenna-induced

microwave plasma CVD and ion-plating. Nano-size diamond film with a low friction coefficient of around 0.1 and

low wear of around 1 "10 -7 mm3/Nm was deposited. The adhesion of DLC film to the substrate was improved using a

system consisting of ECR plasma-enhanced CVD combined with high-energy ion implantation. The ion implantation

of C ions clearly improved the adhesion of DLC film to a metal substrate. Processing technology for carbon films

deposited uniformly on complexly shaped substrates was developed using plasma-based ion implantation. DLC films

deposited on surfaces with different microwave incident angles had fairly uniform thickness. The effects of factors

such as environments, mating materials, and normal loads on friction and wear properties of DLC films deposited by

ion-plating and RF plasma CVD were determined using ball-on-plate friction testers.

To develop carbon films usable in a high-temperature corrosive environment, combined films composed of B4C

and c-BN were studied using electromagnetically accelerated plasma spraying and ion beam sputtering. Diamond

films deposited uniformly on large-area substrate were developed by microwave plasma CVD.

Keywords: tribology, friction, wear, DLC film, microwave plasma CVD, ion plating, plasma-based ion implantation,

high-energy ion implantation

INTRODUCTION

In Japan, a national project on highly functional carbon and related materials technology, "Frontier Carbon

Technology (FCT)" was conducted from 1998 following a five-years plan. The FCT project consisted of 2 types of

R&D, i.e., R&D on new materials synthesis technologies and material processing technologies. R&D on material

processing technologies included 2 processing technologies for electrically and mechanically functional materials. In

R&D on mechanically function materials, processing technologies for materials with excellent tribological properties

or corrosion resistance at high temperatures were developed. Film deposition technologies on complexly shaped

substrates or large-area substrates were developed. This paper details the progress of R&D in mechanically

functional materials processing technologies, focusing on those with excellent tribological properties.

PROCESSING TECHNOLOGY FOR CARBON FILMS WITH LOW FRICTION

AND HIGH WEAR RESISTANCE 1'2)

Nano-size diamond films were prepared by antenna-induced microwave plasma CVD enabling deposition of a

uniformly thick film. The substrate was a Si wafer and the material gas a mixture of methane and hydrogen. The

surface morphology of film samples deposited under a hydrogen gas flow of 100 sccm and a methane gas flow of 10

sccm (C1) and 20 sccm (C2) is shown in Fig. 1. The C1 surface shows 50-200 nm size grains, and the C2 surface

very fine grains of 20-50nm. X-ray diffraction measurement showed diamond peak from (111), (220), and (311)

planes in both C1 and C2. The average surface roughness of C1 and C2 measured by AFM was 24 nm and 7 nm,

whose average hardness measured by nano-indentation was 83 and 94 GPa.

Friction and wear properties of nano-size diamond films were examined by a ball-on-disk friction tester using

SiC and steel balls (SUJ2). Fig. 2 shows the friction behavior of C1 and C2 sliding against the SiC ball. The C2

sample with very fine grains had a fairly low friction coefficient of around 0.16, although the friction coefficient of

C1 with coarse grains was considerably higher, perhaps due to the abrasive action of coarse grains. When a SUJ2 ball

was used as the mating ball, the friction coefficient was slightly higher than that of the SiC ball. Specific wear rate of

NASA/CP--2001-210948 11 August 6, 2001
C2 was a low 8*10-Smm3/Nm, but that of C1 could not be measured due to the existence of many adherent materials.

A nano-size diamond film was deposited with a methane content of 23at%. XRD analysis of this film showed a fairly

broad diffraction peak suggesting the existence of an amorphous phase. The average crystal size was 32nm and

average microhardness was 55GPa. Friction and wear tests sliding against steel balls showed a low friction

coefficient of 0.11, a low film wear of 1 * 10 .7 mm3/Nm, and a low ball wear of 5" 10-7mm3/Nm.

Active dangling bonds generally occur on the surface of diamond and may increase adhesion between sliding

materials. To reduce the friction of nano-size diamond against a SUJ2 ball, nitrogen-doped nano-size diamond films

were deposited. The addition of nitrogen to nano-size diamond films decreased microhardness and slightly increased

surface roughness, although XPS analysis showed the nitrogen content in film to be very low value at 0.65 at% even

at a maximum nitrogen flow of 40 sccm. The friction coefficient of nitrogen-doped films is shown in Fig. 3. The

friction of films against SUJ2 balls decreased with increasing nitrogen content. XPS analysis showed the ratio of Fe

residue to carbon atoms on the friction track decreased, with increasing nitrogen content. These results suggest that

nitrogen in nano-size diamond film restricts the transfer of Fe and reduces friction.

In the addition of R&D to carbon films by antenna-induced microwave plasma CVD, amorphous C and CN films

were developed using ion-plating. Recent results showed that radiation of Ar ions during the deposition of a-CN

films improves friction and wear properties, as will be detailed elsewhere in the near future.

• , • . .._o •

Fig. 1 Surface morphology of nano-size diamond films

[ ..... : .... : ....

9.5 _r Camp:e,C:

Fig. 2 Friction coefficient of C1

0.3

_0.25

0.2

._0.15

0.1

0.05

0

Si G

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N2 Flow- rate(seem)

4O

Fig. 3 Friction coefficient of nitrogen-doped nano-size diamond films

NASA/CP--2001-210948 12 August 6, 2001
PROCESSING TECHNOLOGY FOR DLC FILMS WITH EXCELLENT ADHESION 3'4)

The strength of DLC film adhesion to substrates was improved using an apparatus consisting of ECR plasmaenhanced CVD combined with high-energy ion implantation (Fig. 4). The advantage of this apparatus is that it

conducts CVD coating and ion implantation simultaneously. DLC film was deposited on a Ni-base alloy of Inconel

718 using ethylene gas. Carbon ions with a high energy of 50 keV were implanted in DLC film. AES analysis

showed that DLC with ion implantation had a mixed layer with a carbon concentration of around 40 at.%.

Friction and wear properties were evaluated using a ball-on-disk friction tester. DLC film with ion implantation

maintained a low friction coefficient of less than 0.2 until about 45,000 revolutions (Fig.5), while a film without ion

implantation increased to the friction coefficient of more than 0.2 only after about 5,000 revolutions. The wear of

DLC film with ion implantation was low at around l*10-7mm3/Nm even at 30,000 revolutions, while that of film

without ion implantation was around 3*10-7mm3/Nm at 5,000 revolutions. This low wear of film with ion

implantation was maintained until the friction revolution at which the film was fully worn. These results suggest ion

implantation is very effective to the improvement of the strength of DLC film adhesion.

.......,I, III

High voltage

power supply

Ion soume

Fig. 4 Apparatus

0.4

0.2

DLC (without ion mixing)

_¢jDLC (wi_

, , , , , , , , , , , , , , ,

5000 10000 15000 20000 25000 30000 35000 40000 45000

revolutions of friction

Fig. 5 Friction coefficient of DLC films with and without ion implantation

PROCESSING TECHNOLOGY FOR CARBON FILMS DEPOSITED

5 6) ON COMPLEXLY SHAPED SUBSTRATES '

Carbon films deposited uniformly on complexly shaped substrates were developed using plasma-based ion

implantation (PBII) developed in the FCT project (Fig. 6). ECR plasma with a mirror field was used to generate a

high-density plasma. A substrate was connected to a high-voltage power supply with a pulse generator. Negative DC

bias and high-voltage pulses were supplied simultaneously to the substrate. The substrate was a Si wafer placed on a

hexagonal prism holder of stainless steel. Methane was the raw material.

The distribution of DLC film thickness on each face of the hexagonal prism against the microwave incident angle

to the surface is shown in Fig. 7. Film thickness is fairly uniform, despite the thickness scattering exceeding 20%.

NASA/CP--2001-210948 13 August 6, 2001
Ramanspectrumanalysisshowedthatwhenthemicrowaveincidentangledecreased, theID/IGratioincreased

slightly,perhapsduetotheslightdifferenceintheratiooffilmthicknessdepositedduringnon-pulsebiasandthe

amount ofirradiatedionsbychangingthemicrowaveincident angle.

FrictionpropertiesofDLCfilmsdepositedbychangingthemicrowaveincidentanglewereevaluatedusinga

ball-on-diskfrictiontester.Fig.8showsthedependence ofthefrictioncoefficient onthemicrowaveincident angle.

Thefrictioncoefficient ofDLCfilmdecreased slightlydecreasing microwaveincident angle,similartotheID/IGratio

inRamanspectra, althoughthecorrelationbetweenthefrictioncoefficient andtheID/I_ratioisnotyetknown.

!

Electromagnet coil

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600

500

400

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300 ..................

200 .....................................................................................

100 ...................................

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-150 -100 -50 50 100 150 200

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-150 -100 -50 0 50 100 150 200

Microwave incident angle (deg)

Fig. 8 Distribution of friction coefficient of DLC films versus microwave incident angle

NASA/CP--2001-210948 14 August 6, 2001
EFFECTS OF DIFFERENT FACTORS ON DLC FILM TRIBOLOGICAL PROPERTIES 7's)

To determine the design concepts for obtaining tribologically excellent carbon film, we used ball-on-plate friction

testers to clarify the effects of factors such as environment, mating material, and normal load on the friction and wear

properties of DLC films deposited by ion-plating and RF plasma CVD.

Figure 9 shows the friction coefficient of DLC films deposited by ion plating and RF plasma CVD and Si or F

including DLC films under different humidity. These results suggest that the dependence of the friction coefficient on

humidity differs considerably with the deposition method and conditions or film composition, although it is well

known that the friction coefficient is strongly affected by humidity. When mating ball material hardness was

decreased, the friction coefficient of DLC films increased, perhaps because the contact area between the ball and

DLC film is wider when ball is less.

.0,.20

c

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0.05

0.00

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Fig. 9 Friction coefficient of different DLC films under different humidity conditions

PROCESSING TECHNOLOGYFOR ANTICORROSIVE FILMS UNDER HIGH TEMPERATURE

CONDITIONS AND DIAMOND FILMS DEPOSITED ON LARGE-AREA SUBSTRATES 9)

To develop carbon films usable in a high-temperature corrosive environment, we studied combined films

composed of B4C and c-BN. B4C films were deposited using electromagnetically accelerated plasma spraying. The

advantage of this spraying is that spraying powders are very high-velocity and attain very high temperature, enabling

the spraying of material with a high melting point such as B4C. We developed a spraying apparatus and determined

the effects of spraying conditions on film spraying quality. We also studied c-BH films for protecting B4C by ion

beam sputtering.

Diamond films deposited uniformly on large-area substrates were developed by microwave plasma CVD. We

developed an antenna-induced microwave plasma CVD to deposit diamond film uniformly on large-area substrate

such as glass plate.

ACKNOWLEDGEMENTS

This work was been supported by the New Energy and Industrial Science and Technology Development

Organization of Japan as a part of the FCT project of the Ministry of Economy, Trade and Industry of Japan.

REFERENCES

1) K. Wazumi, A. Tanaka, and Y. Koga, Proc. ITC Nagasaki 2000, (2000).

2) K. Wazumi, Y. Koga, and A. Tanaka, Proc. ADC/FCT 2001, (2001).

3) K. Fuchigami, M. Horiuchi, K. Takashima, and K. Uematsu, Proc. ITC Nagasaki 2000, (2000)

4) K. Fuchigami, M. Horiuchi, K. Takashima, and K. Uematsu, Proc. ADC/FCT 2001, (2001).

5) T. Watanabe, K. Yamamoto, Y. Koga and A. Tanaka, Proc. ITC Nagasaka 2000, (2000).

6) T. Watanabe, K. Yamamoto, Y. Koga, and A. Tanaka, Proc. ADC/FCT 2001, (2001).

7) Y. Ozmen, A. Tanaka and T. Kumagai, Proc. ADC/FCT'99, (1999) 289.

8) B. Na and A. Tanaka, Proc. Ibaraki Dist. Conf. JSME, (2000) 45.

9) J. Kitamura, S. Usuba, Y. Kakudate, H. Yokoi, K. Yamamoto, and S. Fujiwara, Proc. ADC/FCT 2001, (2001).

NASA/CP--2001-210948 15 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CURRENT STATUS OF R&D IN FRONTIER CARBON TECHNOLOGY (FCT)

PROJECT - ELECTRICALLY FUNCTIONAL MATERIALS

K. Kobashi

FCT Project/JFCC

Center for Advanced Research Projects, 6F, Osaka University, 2-1 Yamada-oka, Suita, Osaka 565-087 l, Japan

ABSTRACT

The R&D subjects in the Electrically Functional (EF) Group of the FCT Project at the Central Research

Laboratory (CRL)in Osaka University are focused on diamond. They include (1) morphology control of diamond

films grown by a 60 kW 915 MHz microwave plasma CVD reactor, (2) fabrication of sharp emitter tips on single

crystal diamond and (3) characteristics control of field emission from diamond particles. Both Sumitomo Electric

(SEI) and Kobe Steel (KSL) have Satellite Research Labs in the companies and are associated with the CRL: (4)

SEI develops a new CVD reactor to grow large single crystal diamond with a low density of defects by CVD, while

(5) KSL undertakes heteroepitaxial growth technology development of diamond films on Pt.

Since the start of the Project, there are remarkable achievements in each subject, and this paper will present the

latest results obtained in the Group.

INTRODUCTION

The diamond research in the EF Group at Osaka University CRL consists of 5 major subjects: (1) diamond

growth by a 60 kW 915 MHz microwave plasma CVD reactor, (2) fabrication of sharp emitter tips, (3) field

emission from diamond particles, (4) development of a new CVD reactor for homoepitaxial growth, and (5)

heteroepitaxial growth on Pt. It must be emphasized that these technologies are developed not for scientific purposes

but to actually use for real products and production in the near future. To this end, our efforts were focused more on

finding new regimes in the process parameters and the identification of optimized regions than understanding the

mechanisms. This is because there are a number of possible hi-tech applications of diamond and diamond films, as

shown in Fig. 1, and it is most urgent at the present stage to undertake feasibility studies for each item.

DIAMOND FILM GROWTH ON LARGE AREA USING A 60-kW REACTOR

The 60-kW, 915-MHz microwave plasma CVD reactor was installed by ASTeX/Seki Technotron, as shown in

Fig. 2 (ref. 1). Usually, the reactor was operated using CH4/H2 under the conditions of {60 kW, 100 Torr, 800 - 900

°C}. The recommended substrate is a 12-mm thick Si plate, but we also use standard Si wafers. The plasma

generation is very smooth and stable, and so far there was no serious troubles in the reactor. Our purpose of this

project is to study various film morphologies by changing the CH4 concentration c and the substrate temperature Ts

that is a fimction of the gas pressure and the microwave power. It might appear that given c and Ts, the domains for

different film morphologies would be similar to those of small (i.e., 1.5 kW) reactors.

It was however found that in the c-Tsplane, the domains of<100>- and <11 l>-oriented growth in the c-Ts plane

are totally reversed between the 60 kW and small reactors (ref. 2). For small reactors, <11 l>-oriented domain is

known to exist at higher temperature than the <100>-domain, while it is opposite in Fig. 3. It is not certain at the

present stage how the crossover of the domain takes place as the gas pressure is increased from about 30 Torr for

small reactors to 100 Torr for the 60 kW reactor.

It also is of interest that the film morphologies are unique. For the <100>-oriented film (top of Fig. 3), the

comers surrounded by (111) faces are sticking out of the film surface. On the other hand, for the <11 l>-oriented film

(bottom of Fig. 3), the majority of the diamond faces is (111) that seem to be mutually twinned.

The a parameters were evaluated from the change in the isolated crystal shape with time, and the approximate

curves for a = 1.5 and 3 are shown in Fig. 3. Compared with the results ofref. 1, it is seen that the curves are

shifted toward the high temperature side. Further, there was indication that when a > 3, <11 l>-oriented diamond

films were grown, which is in contradiction to the conventional result that the orientation is < 100>.

One of our future tasks is to elucidate those interesting phenomena.

NASA/CP--2001-210948 16 August 6, 2001
R&D and Future Applications of

Diamond and Diamond films

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PLASMA ETCHING OF DIAMOND

Fabrication of a sharp tip and their array using diamond would be useful for field emission tips, various sensors

and devices, actuators, and optical components. In particular, a use of single crystal diamond is advantageous over a

use of polycrystalline diamond films, because the tips are single crystal so that a uniform function is expected by

each tip. In the present study, we used a standard reactive ion etching system. For the etching gas, a mixture of CF4

and 02 was used, and the A1 masks were photolithographically fabricated on single crystal (100) diamond surface.

One of the etched samples is shown in Fig. 4. It is obviously seen that the length of the diamond column is

9.5 gm, its diameter is 1.2 gm, i.e., the aspect ratio is 8. The etching rate was as high as approximately 9.5 gm/h.

Emitter fabrication technologies combined with diamond CVD also are ongoing (ref. 3).

NASA/CP--2001-210948 17 August 6, 2001
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Figure 3. <111> and <100>-oriented domains determined by X-ray diffraction and the film morphologies

observed by SEM. The rhombus points indicate the experimental points.

NASA/CP--2001-210948 18 August 6, 2001
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Figure 4. RIE etching of single crystal diamond

FIELD EMISSION

Field emission from diamond is currently a hot topic for applications of flat panel displays and vacuum electron

emitter due to the negative electron affinity of properly surface treated diamond. In the FCT Project, we use diamond

particles sprayed on Mo or other substrates, followed by diamond CVD for 10 - 15 min. We believe that the use of

diamond powder is industrially more feasible because of cost and processing (ref. 4).

The electron emission from diamond

powders can be visible by imaging on a 200 gm

size plate. Figure 5 is the result under a bias of

1.9 kV/mm. Figure 5(a) is an image under a HgXe light illumination, while Fig. 5(b) is that of a

standard field emission. One may be able to notice

an intense emission spot (hot spot) in the upper

left of Fig. 5(b). Such hot spots have been

observed by Prof. R. Nemanich at NCSU by

PEEM using similar samples. From Fig. 5, it

was confirmed that it is diamond particles that are

emission sources.

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Figure 5. Field emission with Hg-Xe illumination (a) and without light(b)

Figure 6 is the results of electron energy distribution due to field emission. This work was done in

collaboration with Dr. S. Kono of Tohoku University (ref. 5). The experimental setup is shown in the inset and, the

spectra are plotted as a function of extraction voltage V_x. The vertical line indicates the Fermi energy of Si as a

reference. It is seen that the spectral peak shifts to the lower energy side as V_x was increased. Although the

mechanism of the peak shift has not been fully understood yet, we assume that this is related with the band structure

at the interface between diamond and the substrate.

We have scan the surfaces of diamond particles by scanning tunneling microscopy (STM) and measured the IV characteristics at each point. It was then found that even on a san-face of 1 gm diamond particle, there were

basically two sites that exhibit different I-V characteristics, as shown in Fig. 7.

NASA/CP--2001-210948 19 August 6, 2001
Theupper-left graphin Fig.7 is thecaseof narrow

bandgapofapproximately 1eV,whiletheupper-right graph

is thecasewherethebandgapisapproximately 5eV.The

lower-left figureis aSTMtopographicimageof 3.8gm

square, showingtheshapeofadiamondparticle.Thelowerrightfigurecorresponds tothelower-right figurexcept that

thebandgaps(>5eV,1 5eV,0 1eVandghostimages)

aremappedwith25nmsquarepixcels.It wasfoundthatthe

siteswithabandgapofabout5eVarelocatedalongthetop

edgeofthediamondparticle.Webelievethathistechnique

will beusefulto identifytheemissionsitesandtheir

electronicstates.

Now,it isassumedthathedataofFig.6suggest that

a strongbandbendingexistsattheinterfacebetweenthe

substrateandthediamondthatareseparated byunidentified

insulator. Usingthisasaworkinghypothesis, weundertook

thefollowingexperiments: wedepositedthin, undoped

diamondlayersonW substrates, andthenseededmicronsizediamondparticles. It wasfoundthattheemissionwas

greaterwhentheundopeddiamondlayerthicknesswas0.5

gmthantheyare0.1gmornone.Thisresultshowsthatthe

insulatinginterlayer betweentheW baseandthediamond

powderplaysanimportant roleinthefieldemission.

Figure 6.

emission. The inset is the experimental setup.

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NASA/CP--2001-210948 20 August 6, 2001
DEVELOPMENT OF SINGLE CRYSTAL GROWTH

The purpose of this project is to develop (1) the homoepitaxial technology of large single crystals with a

minimal defect density, and a specialized CVD reactor for this purpose. Figure 8(a) is the reactor designed by

Sumitomo Electric. The microwave is guided from the generator on the left, and introduced from the top of the CVD

Figure 8(a) CVD reactor for homoepitaxial growth Figure 8(b)Plasma seen from the viewing port

chamber. The schematic structure of the system is shown in the paper by Prof. Yoshikawa. Figure 8(b) is the plasma

seen via the viewing port. It may be seen that the plasma is located in the center of the chamber above the diamond

substrate. The design of the chamber was carried out in two ways: (i) gas flow analysis and (ii) electromagnetic

analysis. Aplasma emission analysis indicated that under the conditions of {1% CH4/H2, 10 kW, 100 Torr}, the

plasma spectrum was uniform up to 9 mm from the center (18mm in diameter)

The plasma uniformity was experimentally studied by evaluating a parameters across the diamond crystal

substrate. This was undertaken by first fabricating diamond columns of 6 gm in diameter and 5 6 gm in height by

photolithography and plasma etching, and growing diamond by CVD on the substrate. As a result, the columns

change to cubo-octahedrons without the bottom half, and a parameters can be evaluated from the shape. Again, this

experiment showed that the plasma was uniform up to 8 mm from the center.

As a preliminary test, type Ib single crystal diamonds of 3 mm square are placed in an array of 4×4, and the

CVD was done using the reactor shown above. The results are presented in Figs. 9(a) and (b). Figure 9(a) is the

boundary before CVD, where an approximately 10 gm gap existed between adjacent diamond plates. By contrast,

after CVD, such a gap entirely disappeared to be a coalesced single plate. Even though grain boundaries still

remains, a large diamond of 12 mm square will find industrial uses.

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NASA/CP--2001-210948 21 August6,2001
HETEROEPITAXIAL GROWTH

For the growth of spontaneously coalesced diamond (SCD) on Pt(111), we have used almost just (zero off

axis) substrates, but investigated the off-axis substrate effects on the coalescence in the following: It is seen that

while there are randomly-oriented diamond grains in Fig. 10(a), the surface became progressively flat as the film

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!_" :i:i:i:i:i:i:!:i:i:i:i:i:!:i:i:i:i:i:!:i:i:i:i:i@_:i:}: :)i:

:.: ::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

..:_.::::.. ::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

....::;4"::::::!:: i:i:i:i:::::::::::::::::::::::::::::::::::::::::::::::::::::

Film thickness

iiiiii!ii___i_!i_!ii!iiiiii_ii_iii_iii!iiiiiiii_iii!ii_ii_iii_!_iii!}!!iii_iiiii_:i!_!ii!ii!ii_!ii!i!_i!!_____!__.__iiiiii!__iiii!i!!iii,_ } (a) 1.5_tm, (b)3_tm, (c)

6gm, (d) 9gm, (e) 12gm

Figure 10. Film surface morphology for different film thickness on 2 ° off-axis Pt(lll)

thickness was increased. There are however two major issues to be solved: the first is to increase the coalesced area

over the 10 mm substrate. We have achieved a uniform network of partial coalescence over the entire substrate, but

there still remain triangular pits and uncoalesced regions locally. The second is the existence of anti-phase

boundaries. In fact, the off-axis experiments were intended to investigate the possibility of removing the anti-phase

boundaries, and there was indication that diamond (111) faces are oriented in one direction as the film thickness is

further increased.

Subtle changes of the process parameters resulted in better coalescence of the diamond films. Figure 11 shows

the surface morphology after 100 h growth. Unlike the results of Fig. 10, there was no pits and the surface was

covered with bunched steps. 25 30 % of the substrate surface was covered by the morphology. On the other hand,

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

i._:.?:.:.:._._.:.:.:._.:._._::.;.:.:+:.;._._._.?..;.:o::.:.:._.:._,

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,.,...%.o-, .,.,%..,,-., .%°.. %%. %.o. %. ,....,..o*.O.-

............................ ,o.,.

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y..:.:.:+;.:.;.;.:.1-:.;.:o:-:.::.;.:.:.:.:...:.'.:.:.:.:.:.:.2.:.:.;.:

%o.,,.Oo,,,o.,,.,.,,.°°,.o, .,.. ° ,.oO.,,..,...,,,,o,,.,,-,,,,o,°,

i!!i!ii!ii!ili!!!!i!il;i

Figure 11. Surface after 100 h growth

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Figure 12. Surface grown hy adding B_H6

NASA/CP--2001-210948 22 August 6, 2001
WefoundthatanadditionofB2H6inthesourcegasaccelerates thecoalescence betweendiamond(111)faces,as

seenin Fig.12.Judgingfromtheresultsdescribedsofar,it is consideredthatfurtheroptimizationfordiamond

growthwillmakeafullycoalesced diamondfilmwithoutgrainboundaries.

CONCLUSIONS

It seemsthatfordiamondapplications, (i)development ofproductionscalereactorswithpreciseprocess

controland(ii)largesinglecrystalsynthesis byhomo-andheteroepitaxyremostimportant atthepresent sage

alongwithfeasibilitystudiesofdeviceandotherapplications. It is ourexpectationthattheR&Dactivities

ongoingintheEFGroupoftheFCTProjectwillgiveusefulcontributionfordiamondapplications.

REFERENCES

1. Tachibana, T., Ando, Y., Watanabe, A., Nishibayashi, Y., Kobashi, K., Hirao, T., and Oura, K.,

Diamond and Related Materials, to be published.

2. Wild, Ch., Kohl,R., Herres, N., Mtiller-Sebert, W., and Koidl, P., Diamond and Related Mater. 3 (1994)

373.

3. Ando, Y., Nishibayashi, Y., Kobashi, K., Meguro, K., Imai, T., Hirao, T., and Oura, K., to be presented

at Diamond 2001.

4. Watanabe, A., Deguchi, M., and Kitabatake, M., to be published in Diamond and Related Materials.

5. Kono, S., Goto, T., Sato, K., Abukawa, T., Kitabatake, M., Watanabe, A., and Deguchi, M., to be

published on Surf. Sci.

NASA/CP--2001-210948 23 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

R&D OF DIAMOND FILMS IN KOBE STEEL

K. Kobashi and Y. Yokota

Frontier Carbon Technology Project/JFCC

Center for Advanced Research Projects (6F), Osaka University, 2-1 Yamada-oka, Suita, Osaka 565-0871, Japan

N. Kawakami, T. Tachibana, K. Hayashi, and K. Inoue

Kobe Steel, Ltd, Electronics Research Laboratory, 1-5-5 Takatsuka-dai, Nishi-ku, Kobe 651-2271, Japan

ABSTRACT

Kobe Steel, Ltd. (KSL) started R&D of diamond films in 1985, aiming at electronic applications. Its R&D

history is approximately divided into flaree stages: the first stage was from 1985 up to 1995, when basic technologies

for CVD and microfabrication were developed. The second stage was after 1995 up to now, when the developed

technologies are employed to make individual test devices such as gas sensors, ultraviolet (UV) sensors, and field

effect transistors (FETs) for feasibility studies. The third stage began in 1998 with the Frontier Carbon Technology

(FCT) Project of Japan Fine Ceramics Center (JFCC) and runs in parallel to the second stage. In this stage, both

heteroepitaxial growth and large area deposition of diamond films are the main interests of KSL that has been

scientifically studying the morphology and growth of diamond films since 1985.

Keywords: electronic devices, transistor, UV emission, UV sensor, heteroepitaxial growth

INTRODUCTION

In this paper, we will present the current R&D status of diamond film technologies and applications in KSL. In

particular, the latest results on three topics, gas sensors, UV sensors, and field effect transistors (FETs), will be

described in detail. Current diamond film activities in KSL are open on Internet, and thus the present paper will be

described along the contents of the homepage:

(htto://www.kobdco.co.jl'/showr_,_cm/:q'O802e/'nl'O8025e.hlm-#Ir_dex).

HIGHLY ORIENTED DIAMOND FILMS

The first topic, shown in Fig. 1, is the highly oriented diamond (HOD) film (ref. 1), where (100) faces of

diamond are aligned in the film. The surface of a 30 pm thick film consisted of 10-20 pm square crystal grains.

FWHM values of {004} and {022} diffraction peaks of the 30 pm thick film were approximately 8 and 12 degrees

(ref. 2), respectively, as determined by the polar X-ray diffraction. Consequently, the grain boundary density in

HOD was significantly smaller than that of regular polycrystalline films, and the electrical properties of HOD,

represented by hole mobility, was high. The average surface roughness (ea) of the films was in the range of 100 -

200 nm. It also is an advantage that the film surface is more flat than regular polycrystalline films, which facilitates

electrode formation and other microfabrication on HOD. More precise control of the alignment, further coalescence

of adjacent (100) faces still need to be studied.

SPONTANEOUSLY COALESCED DIAMOND FILMS ON Pt(111)

Spontaneously coalesced diamond films can be grown by microwave plasma chemical vapor deposition on

(111) domains of Pt foils (ref. 3), bulk single-crystal Pt(lll), and single-crystal Pt (111) films deposited on

strontium titanate (111), as seen in Fig. 2. Such films are referred to as "spontaneously coalesced diamond (SCD)

films. In all cases, Pt surfaces have been ultrasonically roughened to enhance diamond nucleation prior to CVD.

Nevertheless, (111) faces of diamond are azimuthally aligned, and a significant coalescence is developed between

adjacent (111) faces. A TEM study by Osaka University revealed that diamond crystals were in an epitaxial

relationship with the basal Pt lattice, and the diamond surface contains screw dislocations of only 10 s/cm 2.

For industrial applications of SCD films, we developed a new substrate where a platinum (111) films of 4 pm

thickness were sputter-deposited on sapphire (0001). Furthermore, a substrate structure ofPt(111)/Ir(111)/Pt(111)

NASA/CP--2001-210948 24 August 6, 2001
iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii! ................................................................................................... __-_'_-_'_-_'_-_'-'_i i

,:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.: ...... __---_ :,

iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii Wehavesuccessfully synthesized (001)- iii

:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i oriented diamond films on Si (001). Highly II_"¢#_-..:I___;....'_NftttJ_ :i:

iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii oriented diamond films have low inter-and __.._ iii

:::::::::::::::::::::::::::::::::::::::::::::::::: intJca-grain defect densities, and hence the _::::_::_ i::

iiiiiiiiiiii_i_{_Niiiiiiiiiiiiiiii Hall mobilities are as high as 280 cm2/V.s. __ :ii

i:i:i:i:i:i:_i:i:i:i:i:i: By comparison, conventional polycrystalline _.._..._i i:i

iiiiiiiii_._iiiiiiiii diamond films have a high density of grain _:.. _ :_ iii

:i:i:i:i:"K";"_;'_'_u_i:i:i:i:i boundaries and thus their Hall mobilities are _i_'_i :i:

:.:.:.:.:_:.:.:.: - "._ 2.......... ___i :': ,'.'.'.'._X_I_N,I._I.I_X.'.'.'.: omv around _u cm / v .s. m_mv onemea '.'.:.:.i+'.'-'. :.:..-:.:o'.'.'.'.'.."- -'.'.'.'.:.._Xo:.:.:.:.. ",

:::::::::::it::::::::::::: dmmond films can be used for sensors and :.:::':'::.:::.::::::::_::_::_:.:::::::_.:::::":.:'.;-:.:"::l :::

iiiiiii_iiiiiiiii i actuators, i

',iiiiiiiiiiiiiiiiiiii Niiiiiiiiiiiiiii .................... ',

::iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii Figure,..ighlyorienteddiamond ii::

i':iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii ...................................................................................................... ................................................:-........

:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:, --------------------------z . ,

iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii The growth of highly-coalesced

i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i: diamond film on Pt(111) was _. _l_l_.:-'k_.'_'_

:::::::::::::::::::::::::::::::::::::::::::::::::: discovered by Prof. Y. Shintani at the ___ _q_x_,._,,_.,_l.N.

p%°o°o°o°o°o°o°o°o°o°o°o°o°o°o°o°o°o°o°o°o°o°o%1 _ _. •

:::::::::::::::::::::::::::::::::::::::::::::::::: University of Tokushima and we are __.__

........... :.:.:.:.:.:.:.:.:.:.:.:.:.:. . . , ...,_ .,,,

::::::::::::::::::::::::::::::::::::::::::::::::: collaboratmg m research. The growth ___

i:i:i:i:i:i:_:_i:i:i:i:i:i: technology of (l l l)-oriented diamond __'_

::::::::::::::::::::::::: films without grain boundaries on a ___

iiiiiii__iiiiii: Pt(l l l) substrate of l 0 mm in diameter __

::iiiiii_iiiiiii is a R&D target in the FCT Project. We __

:i:i:i:iI_]_:i:i:i: foresee that highly-functional electronic _i.':._l_ ._

i:i:i:i_:i:i:i:i: devices will be fabricated using _'_.el.":'_'-I__NI_I_

:i:i:i:i:i:i::.. ..::i:i:i:i:i:i:i heteroepitaxial diamond films. __-_._._

,................................................

iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii Figure 2. Spontaneously coalesced diamond film

,:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:

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We are developing UV-light emitting 6 5 4 Photon3 Energy (eV t '.' :::

devices using semiconducting diamond '_--_: II i

films. Since diamond is an indirect " I_

transition-type semiconductor, it has not

,_ ,,

been considered as a light emitting _ ,- _., (b) 16 _z _I

device material. We found, however, _ ', '

that heavily boron-doped diamond films _ '__

emit a 248-nm (5.0 eV) ultra-violet light c

co L_

, ]

at room temperature. The energy of -_ " ,'

light is close to the bandgap of diamond _ ,' ' ,'

,,

(5.5 eV). Once the mechanism of UV ,- (a) 1,8_mI

light emission is made clear, we will be _d'_'_im _e __.__.,

able to fabricate solid-state UV LEDs. _-- T _ _ ..........

_oo _oo _oo _oo _oo 7oo _ooiii

Wa,_elength (nrn) ,.,"

Figure 3. Cathodoluminescence and UV LED :::

/ I

T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T_T**_

NASA/CP--2001-210948 25 August 6, 2001
/sapphire(0001) was newly designed (ref. 4). It was then found that the FWHM of X-ray diffraction rocking curve of

diamond (111) was reduced to 1.1 °. In this case, the coalescence of diamond (111) faces was extensive even when

the fill thickness was only a couple of microns. Furthermore, a fill delamination from the substrate was avoided.

CATHODOLUMINESCENCE AND UV-LED

Since SCD films on Pt(111) have (111) diamond faces, it is possible to heavily dope (3x102° cm -3) diamond

films with B. Figure 3 shows two cathodoluminescence (CL) spectrum of a film deposited only to 1.8 gm thickness

(a) and that to 16 gm (b)(ref. 5). Both measured at room temperature. It is of interest that sample (a) has a band edge

emission at 248 nm (5.0 eV) as well as band A at 400 - 600 nm. By contrast, in sample (b), only a 248 nm band is

seen, while band A is absent. These results are in contrast to previous data that the band edge emission was virtually

extinct at room temperature. By contrast, the temperature effect on the present 248 nm band intensity was found to

be very weak, and hence it is likely that the SCD film can be used for UV-LED. Indeed, we found the 248 nm

emission from Metal-intrinsic diamond-Semiconducting diamond (MiS) diode structures, though the intensity was

still too small for practical applications. At the present stage, it is necessary to know the electron transition processes

for the 248 nm emission occur at room temperature.

FIELD EMISSION FROM DIAMOND FILMS

In KSL, field emission from diamond films has been studied over years under the guidance of Prof. J.

Ishikawa, Kyoto University (ref. 6). The results are published in ref. 6. Shown in Fig. 4 is an array of diamond tips

made by the mould method. The field emission measurements of this and other specimens as well as the S-K

analyses have been undertaken by Prof. Ishikawa.

GAS SENSORS

Conventionally, electrolyte-type systems are used at manufacturing sites for sensing semiconductor doping

gases such as PH3 and ASH3. It is however expected that both the size and the maintenance cost would be greatly

reduced, if they could be replaced with solid-state sensors. In association with New Cosmos Electric Co., Ltd., KSL

was able to fabricate polycrystalline diamond film gas sensors that have very high sensitivities (the change in

electrical resistance was more than 20 %, and >100 % in certain cases). The selectivity against alcohol and hydrogen

was reasonably high. This is in contrast to a previous work where the sensitivity is only a few %. The reason for the

high sensitivity in our case is presumably due to the proprietary pretreatments to activate the fill surface.

NASA/CP--2001-210948 26 August6,2001
Au-paste \

Pt electmde_ _---"-- Pt wire

. u 222m07

1 1 i i i i i ...._. Pt heater

Figure 5. Schematic structure of gas sensor

Figure 6 shows the response (change in the sensor resistance) upon

introduction of 0.1 - 0.5 ppm PH4. It is seen that the resistance

increased quickly by the introduction of 0.1 ppm PH4, and is almost

proportional to the PH4 concentration. After the gas shut-off, the

decay time was about 10 min. The mechanism for the sensitivity is

still unknown, and a further study will be needed to achieve the best

sensitivity, and selectivity. (ref. 7)

[ '-"

63

C

CO i

¢

/

P

/

..... .,_-J ,

•_ _o

=

=

i

I

1

i

"[]me

Response of diamond gas sensor to PH3

Figure 6. Response of gas sensor

CHEMICAL ELECTRODE

Either B-doped semiconducting or heavily B-doped "metallic" diamond films are recently found to have a

number of characteristics that have not been seen in metallic or graphitic electrodes. Diamond film electrodes are

featured by a wide potential window in aqueous and non-aqueous solutions, a low backN:ound current, and

robustness against acidic and alkaline solutions. Numerous studies have been done in this field, and the reader can

refer to ref. 8.

Shown In Fig. 7 is a result of a collaborativ e work with Prof. G. Swain, Michigan State University, where Pt

particles of 10 - 500 nm are fixed on a B-doped diamond film. It was then found that Pt particles are stable at least

up to 2000 cycles in 0.1 M HC104 (ref. 9).

:::::::::::::::::::::::::::::::::::::::::::::::::::::::: ............................ Heavily B-doped diamond film has a low resistivity and exhibits very high

::::::::::::::::::::::::::::::::::::::::::::::::::::::: performance as a chemical electrode. It is (1) dimensionally stable because

:::::::::::::::::::::::::::::::::::::::::::::::::::::::: diamond is resistant to chemical agents, (2) high overpotential, and (3) low

:::::::::::::::::::::::::::::::::::::::::::::::::::::::: noise. These features are most suitable for bio- and chemical sensor

i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i: applications. We foresee that diamond sensors will be widely used for

:::::::::::::::::::::::::::::::::::::::::::::::::::::::: biosensors for medical inspections and treatments.

i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i: _,,z_[

............................. _ ,,,_" f'_'x \ !, ::';:-:':'::"::':'::'::': _ ,:.-':..,:.:,:.:,:;:

.............................. , .1 '_

:::::::::::::::::::::::::::::::::::::::::: _,,,,,,. ,,.:,,.

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:::::::::::::::::::::::::::::::::::::::::::::::::::::::: : l'_ "7..X; ,,

:::::::::::::::::::::::::::::::::::::::::::::::::::::::: :L_'_,!_'_,_" _ ---: ,-? _ "'-------:w-_,_

::::::::::::::::::::::::::::::::::::::::::::::::::::::: ;_ :>-',S_': : : : , B-doped

::::::::::::::::::::::::::::::::::::::::::::::::::::::: " , : , , : : l diamond

i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i: "_ '_' "_'" '_"_' _'" _'!_"

iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii _.;_,_.:,.,;_,, ,_,,,:i, ,:..,,,_,.,.._r_,

::::::::::::::::::::::::::::::::::::::::::::::::::::::: Figure 7. CV characteristics of Pt particle seeded B-doped diamond film

,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.,.%,,

UV SENSORS

Diamond film is known to be a most appropriate material for UV detection with solar-blind characteristics. In

association with Prof. A. Gicquel, LIMHP-CNRS and her colleagues, we fabricated UV sensors that have a pair of

interdigited electrodes on highly oriented, undoped diamond films. Its temporal photoresponse was measured by

NASA/CP--2001-210948 27 August 6, 2001
illuminating the UV sensor with an ArF excimer

laser (X = 193 nm, pulse width = 5 ns) or dye

laser (X = 313 nm, pulse width = 7 ns).

.;.;_.7<.7...7,7.:.707.:.:.:.7.7.;q.7.:,7.;,:o;,:.' Interd glted Pt or AIelectrodes ,,.._" _,'_4_. _

::::_>.._:...................,....,...

 iiiil

_iiii] ::::__:,x,x_..'.x:_,s

(al Cross-section (b) Interdigited electrodes

Figure 8. Schematic structure of UV sensor

g

8

8

,=

2,5-

2 _

,#

1,5" 4 %,

,iF ,_ Bias xo_ge = 50 V

1._ "x 0,5"

.......... "B'_ .....................................................

"nine(s)

Figure 9(a). Photoresponse against 193 nm (6.42 eV) laser pulse

......... t]_14_ ........................................................

rl,0 4. ,0_.

n[]_ .t

3'

rln3

,7,

,# .0 Bias xoltage = 513 V

4_

8 0.[11 . %

::......... ',13-,[11........................................................ ::

"rime(s)

Figure 9(b). Photoresponse against 313 nm (3.96 eV) laser pulse

Figures 9(a) and (b) show the photoresponses of a diamond UV sensor against 193 and 313 nm lasers,

respectively, where the gap between the interdigited electrodes was 10 gm. In both cases, the laser power was 25 gJ.

In Fig. 9(a), the dotted curve is the measured data for 193 nm laser, while a thick curve near the time axis is the

measured data for the 313 nm laser. The same data is replotted in Fig. 9(b). Hence, it is obvious that the signal

intensity ratio for 193 nm/313 nm is approximately 100, indicating that the diamond film UV sensor is solar blind.

NASA/CP--2001-210948 28 August 6, 2001
DIODEANDTRANSISTORS

A diode consisting of a metal electrode, undoped (intrinsic) diamond, semiconducting diamond, and silicon or

back contact is named MiS diode, and has been a subject of extensive study previously. The same structure has been

recently used to make a UV LED. The reader can refer to ref. 11 on this matter.

In the present section, a new type of field effect transistors (FETs), called p +-i-p + FETs, are presented (ref. 12).

They were fabricated using electron beam lithography combined with B-ion implantation to make gate and sourcedrain structures, respectively, as shown in Figs. 10(a) and 10(b). The gate length is only 0.4 gm.

The uniqueness of the present FET lies in the fact that holes go through undoped (i) diamond layer under the

gate by the space charge limited current (SCLC) mechanism. By using this device structure and the conduction

mechanism, basic problems associated with a standard MES-type FET have been solved. We thus believe that the

present FETs with the SCLC mechanism will be one of the standard structures of diamond FETs.

Channel length = 0.4/__rn

j Gate electrode

_ Gate S & Dinsulator electrodes

_ i.Diamond p÷-Diamond

Diamond substrate

Figure 10(a). Device structure Figure 10(b). SEM photograph of the FET

The I-V characteristics of a fabricated FET are shown in Figs 1l(a) and (b). It is clearly seen that the device

exhibits FET characteristics and the transconductance gm was 0.3 mS/mm. Due to the still high resistance of the

source and drain electrodes, however, the rise in the drain current Id vx drain voltage Vd was not sharp. Figure 1 l(b)

is a log-log plot for the same I-V characteristics of Fig. 1 l(a). It is obviously seen that the log I-log V slope is about

unity for drain voltage Vd < 1 V, while in the device operation voltage of Vd > 1 V, the slope is approximately two.

This is clear indication and the first FET demonstration that holes travel across the i-diamond layer by the SCLC

mechanism.

8

E

E

E

® 4

i-

,m

13

.qj_ ,-" 9.3 :r,_._)'._

-20

-10 -20 -30

Drain voltage (V)

Figure 11 (a). I-V characteristics

E

E

Ill

¢J

tl

10-2

10"4

10'6

1o'S

0.0001 0.001 0.1 1 10

Drain voltage (V)

Figure ll(b). Log 1-Log V plot

.... -12

Slope =2

100

NASA/CP--2001-210948 29 August 6, 2001
CONCLUSION

In the present paper, we described various applications of diamond fills under development in KSL. It is very

encouraging that the number of applications for diamond films has increased significantly, and thus it would be fair

to expect that real products will steadily appear on the market in the coming years. In order for this actually happen,

it seems that there are two most important technologies to be established: heteroepitaxial growth of diamond on at

least one to two inch wafers, and production scale CVD reactors with precise controllability and energetic

efficiency. We hope that R&D will put more emphasis along this line.

ACKNOWLEDGEMENT

Finally, the authors wish to sincerely thank collaborations with Prof. J. Glass of Case Western Reserve

University for HOD development, Prof. Y. Shintani of Tokushima University for heteroepitaxial growth of diamond

on Pt, Prof. Y. Takai of Osaka University for TEM measurements, Prof. J. Ishikawa and Prof. Y. Goto of Kyoto

University for field emission research, Prof. A. Gicquel of LIMHP-CNRS and her colleagues for UV sensor

research, Prof. G. Swain of Michigan State University for electrochemical studies, Prof. R. Nemanich of North

Carolina State University for PEEM and other experiments, and New Cosmos Electric Co., Ltd. for gas sensor

development.

This work was supported by FCT Project, which was consigned to JFCC by NEDO.

REFERENCES

1. Stoner, B. R., Ma, G.-H. M., Wolter, S. D., and Glass, J. T., Phys. Rev. B 45 (1992) 11067.

2. Tachibana, T., Hayashi, K., and Kobashi, K., Appl. Phys. Lett., 68 (1996) 1491.

3. Shintani, Y., J. Mater. Res. 11 (1996) 2955; Tachibana, T., Yokota, Y., Miyata, K., Onishi, T., Kobashi, K.,

Tarutani, M., Takai, Y., Shimizu, R., and Shintani, Y., Phys. Rev. B56 (1997) 15967.

4. Tachibana, T., Yokota, Y., Kobashi, T., and Yoshimoto, M., J. Cryst. Growth, 205 (1999) 163.

5. Yokota, Y., Tachibana, T., Mikyata, K, Kobashi, K., Hatta, A., Ito, T., Hiraki, A., and Shintani, Y.,

Appl. Phys. Lett., 73 (1998) 1493.

6. Goto, Y., Kondo, T., Nagao, M., Tsuji, H., Ishikawa, J., Hayashi, K., and Kobashi, K., J. Vac. Sci. Tecllol.

B18, (2000) 1018.

7. Takada, T., Fukunaga, T., Hayashi, K., Yokota, Y., Tachibana, T., Miyata, K., and Kobashi, K., Sensors

and Actuators, 82 (2000) 97; Hayashi, K., Yokota, Y., Tachibana, T., Miyata, K., Kobashi, K., Fukunaga,

T., and Takada, T., Jpn. J. Appl. Phys. 39 (2000) L22.

8. New Diamond and Frontier Carbon Technology, Vol. 9, Nos. 3 and 4 (1999).

9. Wang, J., Swain, G. M., Tachibana, T., and Kobashi, K., Electrochem. Solid-State Lett. 3 (2000) 286.

10. Hayashi, K., Yokota, Y., Tachibana, T., Kobashi, K., Achard, J., Gicquel, A., Olivelo, C., Castex, M.-C.,

and Treschalov, A., Diamond and Related Materials, to be published.

11. Miyata, K., Kobashi, K., and D. L. Dreifus, D. L., Diamond and Related Materials, 2 (1993) 1107.

12. Kawakami, N., unpublished.

NASA/CP--2001-210948 30 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

R&D OF CVD DIAMOND IN SUMITOMO ELECTRIC INDUSTRIES

T. Imai

Itami R&D Laboratories, Sumitomo Electric Industries, LTD.

Koya-Kita, Itami, Hyogo, 664-0016, Japan

ABSTRACT

Sumitomo Electric Industries (SEI) has developed several kinds of diamond products by using CVD methods.

The R&D mainly have been performed in Itami R&D Laboratories since 1981. SEI's CVD diamond products

includes speaker diaphragms, tape-automated bonding tools (TAB), 3-inches mirror-polished diamond coated

wafers and surface acoustic wave (SAW) devices are released from 1985. And the results of research consists the

investigation of the atomic structure on epitaxial-grown diamond, the well-controlled single domain epitaxial

growth, CVD-grown high quality free-standing diamond with 8mm diameter, the forming of sharp diamond emitter

tips with 200nm diameter at the tops, and the diamond coated and polished bearings. The latter there themes were

started as sub-themes of Frontier Carbon Projects since 1998.

The status ofSEI's R&D are as follows,

(1) The 4-inches-diameter diamond-coated wafer is developed for next generation diamond SAW devices. The

diamond layer of the wafer is about 10 microns in thickness. Their surface are mirror-polished and the roughness

is under 5nm in average.

(2) The Lithium Niobate (LN) piezoelectric film are developed for high performance and wide-band diamond SAW

devices. The LN film which highly oriented to c-axis is deposited by KrF excimer laser ablation method. The

structure of laminated LN/diamond are expected to obtain higher electro-mechanical coupling coefficients than

ZnO/diamond structure as well as high acoustic velocity.

(3) The plasma assisted CVD reactor optimized for the homoepitaxial growth of single crystalline diamond is

designed and constructed. The uniformity of plasma parameter and homoepitaxial growth conditions are enough

to grow large size single crystalline diamond as 20mm diameter. The large free-standing single crystalline

diamond with 8mm diameter is already obtained.

(4) A diamond coated bearing is manufactured, and its inner-diameter is 10mm and length is also 10mm. About 10

microns CVD diamond are coated on Si3N4 substrate and polished under roughness of 1 micron.

Some of these works were supported by FCT Project, which was consigned to JFCC by NEDO.

Keywords: homoepitaxial growth, single crystal, piezoelectric, surface acoustic wave, bearing

Takahiro IMAI

Itami R&D Laboratories, Sumitomo Electric Industries, LTD.

Koya-Kita, Itami, Hyogo 664-0016, Japan

Lro_a_i_-_t__'_h___o__s_eLc_,o_j_.

Fax: +81-727-71-0407, Phone: +81-727-71-0621

NASA/CP--2001-210948 31 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

The Dawn of Diamond Age

James Chien-Min Sung

Kinik Company, Taipei, Taiwan, ROC

Taipei University of Technology, Taipei, Taiwan, ROC

Abstract

Diamond is the miracle substance that excels all other materials in many extreme properties. Currently industrial

diamonds are limited by size and geometry, so they are serving primarily in mechanical industry as superabrasives (e.g.,

for grinding wheels). But with the availability of large sized diamond films and complicated DLC coatings many

functional applications of diamond are either commercially available or to be debut. Thus, diamonds are the key

component for tribological surfaces (e.g., bearing linings), electro-optical devices (e.g., field emitters), semiconductors

(e.g., electrodes), thermal management (e.g., heat spreaders), optical gates (e.g., X-ray windows), acoustic vibrators

(e.g., SAW filters), dialectical medium (e.g., capacitor interlayer), medical implants (e.g., hip joims), nuclear detectors

(panicle sensors), military weaponries (e.g., missile radomes), chemical barrier (e.g., acid container) and consumer

products (e.g., razor blades), to name just a few. When many of these exotic diamond applications become routine,

human's materials civilization will reach the climax---diamond age. As there is no other material that can be more

versatile than diamond, diamond will eventually become forever.

Key Words: Diamond, Diamond Applications, Diamond Age

NASA/CP--2001-210948 32 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

APPLICATIONS OF CARBON NANOTUBES TO ENERGY STORAGE

Seung Mi Lee, Kay HyeokAn, Kwan Ku Jeon, Won Seok Kim,

Jeong-Mi Moon, Young Soo Park, and Young Hee Lee*

Department of physics, Institute of Basic Science, Sungkyunkwan University, Suwon, Kyunggi 440-746,

South Korea

Chil-Hoon Doh, Jung-Hwan Lim, and Seong-In Moon

Battery Research Group, Korea Electrotechnology Research Institute, Changwon 641-120, South Korea

ABSTRACT

We have investigated key factors for the applications of singlewalled carbon nanotubes (SWNTs) to various

energy storage systems, such as hydrogen storage, Li-ion secondary, and supercapacitor. In electrochemical hydrogen

storage, the maximum discharging storage capacity of 450 mAh/g (1.63 H-wt%) was obtained at a cutoff voltage of

0.8 V by controlling the amount of transition metals during synthesis and post-armealing process. For Li-ion battery

using SWNT electrodes, SWNT electrodes showed about 560 mAh/g (Lil.56C6) of the maximum reversible specific

capacity, significantly higher than the ideal value of LiC6 (372 mAh/g) for graphite. In case of supercapacitor using

SWNT electrodes, the SWNT electrodes annealed at 1000 °C for 30 min have a maximum specific capacitance of

180 F/g. Singlewalled carbon nanotube (SWNT) and Polypyrrole (Ppy) hybrid electrode was fabricated to more

improve the specific capacitance of the supercapacitor. We obtained a maximum specific capacitance of 265 F/g from

the SWNT-Ppy hybrid electrode containing 15 wt% of conducting agent.

Keywords: Carbon nanotube, hydrogen storage, supercapacitor, Li-ion battery,

INTRODUCTION

Recently, there have been considerable attempts to use carbon nanotubes (CNTs) for active materials of energy

storage systems, such as hydrogen storage by hydrostatic pressure and electrochemical approaches (ref. 1 to 2), Liion secondary battery (ref. 3 to 5, 11), and supercapacitors (ref. 6 to 8). The CNTs are attractive materials for energy

storage systems due to their superb characteristics of chemical and mechanical stability, low mass density, low

resistivity, and large surface area. Our research laboratory has been pursuing application of CNTs to energy storage

system, such as hydrogen storage by hydrostatic pressure and electrochemical approach, Li-ion secondary battery,

and supercapacitors. Here, we will introduce our effort to develop application of CNTs to energy storage system.

Hydrogen has cognized as an excellent energy source, because hydrogen is very abundant element on the earth

surface and the use of hydrogen as a fuel and energy carrier produces neither air pollution and nor green house gas

emission. More recently, the CNTs have been reported to be very promising candidates for hydrogen uptake. The

CNTs take advantage of the gas-on-solids adsorption of hydrogen as well as electrochemical approaches. We will

demonstrate experimentally that the storage capacity is strongly dependent on the sample-preparation conditions in

electrochemical approaches.

A lithium ion cell is the outstanding battery system because of good performances such as high operating

voltage, high energy density, long cycle life, no memory effect, and environmental compatibility. It has been well

known that the performance of Li-ion battery strongly depend on active materials used for cathode. Therefore, much

more effort has been paid to the development of active materials for cathode with high performance. Since discovery

of CNTs as a novel form of carbon, the progress in synthesis of CNT has ignited significant worldwide investigation

on new application of CNTs for cathode material of Li-ion battery. CNTs have been suggested as candidate anode

materials with high energy density for Li-ion battery. We will report the relation of the specific capacity of Li-ion

battery using singlewalled carbon nanotube (SWNT) electrode and the mechanism of lithium intercalation taking

into account the structure and the micro texture of SWNTs.

Supercapacitors (also called electrochemical capacitors, electric double layer capacitor, and ultJcacapacitors) have

many advantages compared to the secondary battery, for instance, long cycle life (> 100,000 cycles), simple

principle and mode of construction, short charging time, safety, and high power density. Among them, the power

density of supercapacitors is the most remarkable property. The power density of supercapacitors is larger than that

of the secondary battery by about ten times, although the energy density of supercapacitors is smaller than that of the

secondary battery, which is a drawback to be applied for practical devices. Now the supercapacitor's capability with

high power density extends its application to various other novel devices of load leveling and pulse power devices,

NASA/CP--2001-210948 33 August 6, 2001
andelectricvehicles. Inorderto applythesupercapacitors to variouspractical devicesevento electricvehicles,

development ofsupercapacitors withhighpowerdensityandhighenergydensityisnecessary. Wewill demonstrate

thekeyfactorsdeterminingtlaeperformance ofsupercapacitors withhighenergydensityandpowerdensityusing

SWNTelectrodeandSWNT-polymer compositeelectrode.

ELECTROCHEMICAL HYDROGEN STORAGE IN SWNTs

We have investigated key factors for electrochemical hydrogen storage in SWNTs to determine storage

capability by controlling the amount of transition metals during synthesis and post-armealing process. The influence

of morphology and crystallinity of SWNTs on the storage capacity is also studied. We limit the weight percent of

the transition metals to 5 wt% (lower limit) and 20 wt% (upper limit) during the synthesis of SWNTs by DC arcdischarge in order to investigate the effect of transition metals on the hydrogen storage capacity. Although the yield

might be lowered with large amount of transition metals, uniform distribution of the transition metals in the sample

may enhance the conductivity of the electrode and eventually increase tlae storage capacity.

The SWNT electrodes are tested electrochemically under static current of 50-100 mA/g and the lower cutoff

voltage of 0.4 V for a reference test with above three samples (Fig. 1). It shows tlaat the pellet with large amount of

transition metals, sample (c), can reserve most hydrogen up to around 140 mAh/g at cut-off voltage of 0.4 V, as

shown in Fig. 1(c), where the mass introduced here implies the total weight of electrodes from SWNTs, amorphous

carbons, and transition metals. Theoretical estimation on SWNTs shows that the metallic SWNTs become

insulator with hydrogen adsorption (ref. 9 to 10). This suggests that the additional conducting materials compensate

the electron transport in the electrode and eventually increase the storage capacity. The maximum capacity down at

zero voltage reaches 450 mAh/g. Optimization of the maximum storage capacity with the amount of transition

metals is further necessary in the future.

1.g

1.4

_1.2

_" 0.8_

0.6.

0.4-

20 40 0 80 100 120 140

Specific capacity (mAh/g)

Fig. 1. The charge-discharge tests of the as-prepared sample of (a) SWCNTs (web) with 5 wt% transition

metals, (b) SWCNTs (cathode) with 5 wt% transition metals, and (c) SWCNTs (cathode) with 20 wt%

transition metals

In order to remove the poly (vinyl alcohol) (PVA) used as a binder from the electrode, we heated the sample by

two steps. We first preheated the sample in air at 300 °C for five hours. This oxidizes the PVA. The degree of

oxidation is also important, as shown in Fig. 2(a). The plateau region in the voltage extends with increasing

annealing hours. Thus a complete oxidation is important in order to remove the PVA more effectively in the second

step annealing. This preheat treatment further prevents tlae sample from being cracked by the second step annealing.

In the second-step annealing, most of PVAs are burned and some are transformed into carbon particles. This high

temperature annealing also improves the crystallinity of the pristine SWNTs. We annealed the sample at different

temperature in Ar atmosphere for 5 hours. The elevation of annealing temperature in Ar extends the plateau region,

as shown in Fig. 2(b). At temperature greater than 900 °C in Ar atmosphere, tlae electrode is cracked, resulting in

the decrease of the storage capacity of hydrogen and tlae plateau region. We emphasize that the preheat treatment is

an essential step in order to remove the PVA binder.

NASA/CP--2001-210948 34 August 6, 2001
Specific capacity [mAh/g]

1.6

64

&

_1.2

1.0

0.8

0 10 20 30 40 50 60 70 80

Specific capacity [mAh/g]

Fig. 2. The specific charge-discharge curves in terms of (a) the annealing time in air with the sample

containing 20 wt% transition metals and (b) the annealing temperature in Ar with the sample containing

5 wt%.

LI-ION BATTERY USING SWNT ELECTRODES

We have surveyed the relation of the specific capacity of Li-ion battery using SWNT electrode and the

mechanism of lithium intercalation taking into account the structure of SWNTs. Two types of SWNT electrodes are

prepared in order to investigate Li-intercalation site in SWNTs. The one is fabricated with no treatment of as-grown

SWNT powder, the other is prepared with the SWNT powder ball-milled for two hours to solve and brake the

bundles of as-grown SWNTs.

_)] 14oo

CS5153Hl=CNT,PVDF/PP//Li, 1.3mA-lhPJ3-OV,charge&discharge, 15.3446mg,

@ mixing without zirconia ball

1200 ................ ._..._..._..._..._..._..._..A...A.. A _--_..._ ...A..._ ...&..., ...4 ................

L

1000

-_ s00

600 .......... _ ......................................................................................................................

_," °, _._P ....* t _, = °,_ e, t 3 t__ p. _

o_ 400 ...................................................................................................................................

200 ...................................................................................................................... _t.........

0

1200

o

lOcycl e Number15 20 25

1000 100

_o _ _ 8oo

•_ o._600

60 _

° i 400

40

C50102GI=CNT, PVDF/PP//Li, 0.8mA-lhPJ3-0V, charge&discharge, 6.3786 mg,

mixing with zirconia ball 100

._.._._..,..=.=.*._._._._. *.._._._.._' =.._.._._-._' =_:.Z'.._...._.ft A.._

J

8o

6o

..,___.r,.,.,_..,,_:__... ___3 __.i__;L;_, __ ................... ,o _"=:

20

5 10 15 20 25 30 35

Cycle Number

lm Jm

Fig. 3. The specific capacity and charge-discharge efficiency as function of cycle numbers;

(a) as-grown SWNT electrode with no treatment, (b) SWNT electrode with ball milling for two hours,

Qc: specific capacity of charge, Qd: specific capacity of discharge, Ah Eft.: charge-discharge

efficiency.

Figure 3 shows specific capacity, discharge efficiency, and recyclability of SWNT electrodes. The no treat_nent

of as-grown SWNT electrode has 1327 mAh/g of very large charging specific capacity at fist cycle. However the

specific capacity of discharge at fist cycle is 560 mAh/g, indicating 767 mAh/g of large irreversible specific capacity.

NASA/CP--2001-210948 35 August 6, 2001
It is suggested that the undesirably large irreversible specific capacity may be due to Li intercalating into the inner

core of SWNT, indicating tlae intercalated Li in the inner core of SWNT have difficult in coming out from the inner

core of SWNT. Gao et al. have reported that the irreversible specific capacity is attribute to the large surface area of

SWNT electrode, estimated to be -250 m2/g by a N2 BET measurement (ref. 11). After fist cycle, it is found that

the reversible specific capacity for no treatment of as-grown SWNT electrode is uniformly maintained, as shown in

Fig 3(a). We emphasize that the discharge efficiency and recyclability axe also very excellent, which is very

important to apply practical device. The discharge efficiency and recyclability of the ball-milled SWNT electrode

also shows very excellent behaviors. However, the ball-milled SWNT electrode shows smaller both of 1019 mAh/g

for charging specific capacity and 385 mAh/g for discharging specific capacity at fist cycle than the no treatment of

as-grown SWNT electrode, implying the Li-reversible intercalation sites are decreased by ball milling for two hours.

The large drop of reversible specific capacity of SWNT electrode by ball milling may indicate that Li intercalate tlae

interstitial bundles of SWNTs, which attribute to reversible specific capacity. Therefore, after solving, and braking

bundles of SWNTs by ball milling, reversible specific capacity of ball-milled SWNT electrode is decreases, due to

the reduction of Li-reversible intercalation sites.

SUPERCAPACITORS USING SWNT ELECTRODES

We have investigated tlae key factors determining the performance of supercapacitors using SWNT electrodes. In

addition, SWNT and Polypyrrole (Ppy) hybrid electrodes are fabricated to improve the specific capacitance of tlae

supercapacitor. The combination of two types, the electric double-layer and the redox system, have been developed

for enhancing tlae performance of supercapacitors.

Figure 4 shows the specific capacitances of the heat-txeated as-grown SWNT electrodes at various temperatures

as a function of the charging time. Capacitances increase abruptly and reach about 80 % of tlae maximum capacitance

during the initial 10 min, regardless of the heat-treatment temperatures. The capacitances increase gradually fuNaer

and saturate to tlae maximum values at long charging time. Persistent increase of the capacitance over a long time is

generally observed from the porous electrodes and is attributed to tlae existence of various forms of pore and pore

diameters in the electrode. The saturated capacitance increases with increasing heat-txeatment temperatures mad

saturates to 180 F/g at 1000 ° .... We note that the capacitance is directly proportional to the specific surface area

with increasing temperatures. The average pore diameter approaches to 30 A at 1000 °... and moreover the

populations of such small diameters are highly weighted at 1000 °....

_-" 180.

120'

e_ 100.

_ 600°C

--.-- 800°C

_- 9oo_

10001C

;o _oo _;o 200 2;0 300 3;0

Charging time (rain.)

400

Fig. 4. The specific capacitances of the heat-treated electrodes at various temperatures as a function of the

charging time at a charging voltage of 0.9 V.

NASA/CP--2001-210948 36 August 6, 2001
•._.:,_J_0_o ,. ,,:.: -: ..:.-..q_ .:,

:._:.:_-_:.:-X,'.:-:.:@:..I._, :¢._.¢,'_:.:-:,:,1.x'-._:

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_,:.::::.:-_:::::.:.:.'ZO-_:',_:,:':,;;:X_.Z,:,i :::,:.X_

,,,,,..:_:_._.._:.-_.-,,....,........:.._.:.:::.x.:

.: ,'_.'... _,:.. ,.*.'.,.,@ ...-...>_,

.::::::::::::;-.__::;';::: g::: .;:::;:":::::;:;::2;:.:: _...."" :._'_.""." :i:! i_:_!i!:.*_:::_; •... .......... ......... .. .........:,,_ '.'- _: .-... <, ,-..,.,

• "'-'" .'..::_'>_.,i:'.'.'..'_.':.:!.'. ." " -" .'_". "".'-.'." ." ".'_'.'-'.". •:':': "":':'. .':'.':':*" ,':':': " :" ;'_. '.':':v;':*:" "._':':':"':':'_.""

:::.:. ============================================== .':: ;_,:::_:_..:."... :::'.::::_:_.:_::::;.;

•" * * %'°'' %" .%%**'.'.*** %" .%*-* " ." *%_o" °'._i'-' *.* _.,_." " °%'.%'. *

:.... ......::..._............_

•:.:.:.:.. ... ....:.:.:.:.;; ::!:i:i:i::!:":':':::_:_:i ......:.: _i; ...+ "" .:..'_:_: :',_:,: .'._:;::'_.::' .:.:.. _::;:_ .::::'.'.'."¢"":'::;:" •

:::::::::::::::::::::::::::::.:::!_:,_.::: _:..:...;::..::::::: ..

Fig. 5. The FE-SEM images of as-grown SWNT, pure Ppy, and SWNT-Ppy powder.

Figure 5 shows the field-emission scanning-electxon-microscope (FE-SEM, Hitachi $4700) images of as-grown

SWNTs, pure Ppy and SWNT-Ppy powder formed by in-situ chemical polymerization. The as-grown SWNTs are

randomly entangled and cross-linked, and some carbon particles are also observed. The purity of as-grown SWNTs

was roughly estimated to be about 25-30 %. From Figure 5(b), one can see that the pure Ppy synthesized without

SWNT shows a typical granular morphology. The granule size of pure Ppy is about 2-3 gm. Figure 5(c) reveals flae

individual nanotube bundles to be uniformly coated with Ppy, indicating that in-situ chemical polymerization of

pyrrole is effectively achieved.

250.

200.

150.

100.

id

pure Ppy

._ 50. as-grown SWNT

o '8b :

0 20 40 6() 100

Discharge current density (mA/g)

Fig. 6. The specific capacitances of the as-grown SWNTs, pure Ppy, and SWNT-Ppy hybrid electrodes as a

function of discharge current density at a charging voltage of 0.9 V for 10 min.

,"- 270"

_:_

260-

250"

240:

230:

220'

(_ 210-

200"

1_0i

_. 180"

g'_ 170'

0 5 10 15 20

Addition amount of Sup-p (wt%)

Fig. 7. The effects of conducting agent on the specific capacitance of the SWNT-Ppy hybrid electrode

as a function of discharging current density at a charging voltage of 0.9 V for 10 min.

NASA/CP--2001-210948 37 August 6, 2001
Figure 6 shows the specific capacitances of the as-grown SWNTs, pure Ppy, and SWNT-Ppy hybrid electrodes

as a function of discharge current density, where the data were taken from the samples charged at 0.9 V for 10 min.

In comparison to the pure Ppy and as-grown SWNT electrodes, the SWNT-Ppy hybrid electrodes shows very high

specific capacitance by 5 - 10 times, although the specific capacitance of SWNT-Ppy hybrid electrode decreased

gradually with increasing discharge current density. This very large specific capacitance seems to be caused by

uniformly coated Ppy on tlae SWNTs, signifying that active sites on Ppy chains are increased. We emphasize tlaat

the capacitance by EDL is also enhanced by tlae enlarged surface area of Ppy.

The effects of conducting agent on the specific capacitance of the SWNT-Ppy hybrid electrode axe shown in

Fig. 7. The maximum specific capacitance appears at 15 wt% of conducting agent in electrode, where the data were

taken from tlae samples charged at 0.9 V for 10 min. The effect of tlae conducting agent is clearly seen particularly at

high discharging current density. The specific capacitance significantly increases with increasing the amount of tlae

conducting agent up to 15 wt% and drops at 20 wt%. We obtain a maximum specific capacitance of 265 F/g from

the SWNT-Ppy hybrid electrode containing 15wt% of the conducting agent. The conducting agent itself does not

directly contribute to the specific capacitance but by reducing the IR drop. However, an excessive addition of tlae

conducting agent diminishes the effective portion of the active electrode material, resulting in the reduction of tlae

specific capacitance.

SUMMARY

We have intxoduced the applications of singlewalled carbon nanotubes (SWNTs) to various energy storage

systems, such as hydrogen storage, Li-ion secondary, and supercapacitor. In electrochemical hydrogen storage, we

have obtained the maximum discharging storage capacity of 450 mAh/g (1.63 H-wt%) by controlling the amount of

transition metals during synthesis and post-armealing process. For Li-ion battery using SWNT electrodes, we

emphasize that the discharge efficiency and recyclability are very excellent, which is very important to apply practical

device. It is suggested that Li-reversible intercalation sites axe the interstitial bundles of SWNTs. The SWNT

electrode shows about 560 mAh/g (Lil.56C6) of the maximum reversible specific capacity, significantly higher tlaan

the ideal value of 372 mAh/g (LiC6) for graphite. In case of supercapacitor using SWNT electrodes, the SWNT

electrodes annealed at 1000 °C for 30 min have a maximum specific capacitance of 180 F/g. Singlewalled carbon

nanotube (SWNT) and Polypyrrole (Ppy) hybrid electrode was fabricated to more improve the specific capacitance of

the supercapacitor. We obtained a maximum specific capacitance of 265 F/g from the SWNT-Ppy hybrid electrode

containing 15 wt% of conducting agent.

ACKNOWLEDGEMENTS

This project was supported by the Korean Ministry of Science and Technology (MOST) through National

Research Laboratory (NRL) program and in part by Brain Korea 21 (BK21) program.

REFERENCES

* The corresponding author: Young Hee Lee, Department of physics, Sungkyunkwan University,

Suwon 440-746, Kyunggi, South Korea, e-mail: leeyoung@yurim.skku.ac.kr, fax: +82-31-290-5954,

phone: +82-31-290-7053

1. Rajalakshmi, N., et al.: Electrochemical Investigation of Single-Walled Carbon Nanotubes for Hydrogen

Storage, Electrochimica Acta, vol. 45, no. 27, September 2000, 4511-4515

2. Lee, S.M., et al.: Hydrogen Adsorption and Storage in Carbon Nanotubes, Synthetic Metals, vol. 113,

no. 3, June 2000, 209-216.

3. Maurin, G., et al.: Electrochemical Lithium Intercalation into Multiwall Carbon Nanotubes: A MicroRaman Study, Solid State Ionics, vol. 136-137, November 2000,1295-1299.

4. Gao. B., et al.: Electrochemical Intercalation of Single-Walled Carbon Nanotubes with Lithium, Chemical

Physics Letters, vol. 307, July 1999, 153-157.

5. Frackowiak, E. et al.: Electrochemical Storage of Lithium Multiwalled Carbon Nanotubes, Carbon, vol.

37, no.l, 1999, 61-69.

6. Niu, E.K. et al.: High Power Electrochemical Capacitors Based on Carbon Nanotube Electrodes, Applied

Physics Letters, vol. 70, no. 11, March 1997, 1480-1482.

NASA/CP--2001-210948 38 August 6, 2001
7. Diederich, L.,etal.:Supercapacitors BasedonNanostructured CarbonElectrodes GrownbyCluster-Beam

Deposition, AppliedPhysicsLetters,vol.75,no.17,October1999,2662-2664.

8. An,K.H.,etal.:Supercapacitors UsingSingle-Walled CarbonNanotubeElectrodes, AdvancedMaterials,

vol.13,no.7,April2001,497-500.

9. S.M.Lee,etal.:HydrogenStorageinCarbonNanotubes, acceptedtoSyntheticMetals,July2000.

10.S.M.Lee,etal.:AMechanismofHydrogenStorageinSingle-Walled CarbonNanotubes, submittedtoJ.

Kor.Phys.Soc.,Jan2001.

11.Gao.B., et al.:EnhancedSaturationLithiumCompositionin Ball-MilledSingle-WalledCarbon

Nanotubes, Chemical PhysicsLetters,vol.327,September 1999,69-75.

NASA/CP--2001-210948 39 August6,2001

2

MARKETS AND TECHNOLOGIES OF DIAMOND

MATERIALS

Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ABSTRACT

Invited Talk

CVD Diamond: The Industrial Landscape

R S Sussmann, King's College London, The Strand, London WC2R 2LS, UK

It is now generally accepted that CVD diamond has fundamentally changed the way

diamond can be used in industry and research. This new synthesis technique has opened

applications in a wide range of fields including optics, electronics, microwave

transmission, cutting tools, wear parts and others. Depending on synthesis conditions,

CVD diamond can be made as a polycrystalline ceramic or as homoepitaxial single

crystals. Synthesis conditions also determine the type of CVD diamond produced and it

has been demonstrated that CVD diamond &many different grades can be manufactured

ranging from dark material which can be used for mechanical or thermal applications to

extremely pure material with excellent optical transmission or electronic properties. This

paper will review the present range of applications pursued by different manufacturers

and will attempt to survey the status of supply of raw material and products. This is

important information required for potential end-users of CVD diamond.

NASA/CP--2001-210948 43 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Diamond and Diamond-Like Carbon Films and Coated Products - A Technology

and Market Update

T. Abraham

Vice President of Research

Business Communications Co

25 Van Zant St.

Norwalk, CT 06855, USA

Tel: 203/853-4266, extn 313

Fax: 203/853-0348

E-mail: tabraham@bccresearch, corn

After a gestation period of over a decade, the technology for producing diamond coatings

and thin films at low temperatures and pressures and at relatively low cost has emerged

as a viable technology. This has allowed the number of potential applications to increase.

These range from coatings on shaving blades, cutting tools, .wear resistant components,

surgical knives, IR windows and optical devices, sunglasses and ophthalmic lenses,

prosthetic devices, microelectronic substrates and devices, Sensors/SAWs/MEMS, data

storage, recording and print heads, flat panel displays/cold cathode, sports equipment and

protective coatings. Research and development continue to improve the current

technologies as well as create entirely new technologies with high deposition rates.

Researchers are also carrying out R&D at a fast pace to apply these coatings in several

new applications. The presentation discusses technological and industrial driving forces

that have made diamond, diamond-like carbon and CBN films and coatings a future

growth industry. Emerging technologies and applications, current and future market

opportunities, and global competition are discussed. U.S. and world markets and

projectedtill the year 2005.

NASA/CP--2001-210948 44 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

THE PRODUCTION AND FEATURES OF INDUSTRIAL DIAMOND MADE

IN CHINA

Zhibin Zhan

Queenbee Diamond Industrial Co. LTD(Beijing 100029 China)

Aiguo Yao Bairu Xia

China University of Geosciences ( Wuhan 430074 China)

ABSTRACT

The paper briefly introduced the production condition of diamond in China. The features of

diamonds produced with different device and by different techniques were discussed in detail. The top

class of diamond by international standard was made with imported two-piston press. The diamond

made with Chinese six-piston press is good material for making saws like tools. In some applications it

is selected as an alterative for above diamond due to its higher quality and lower price. The fine

diamond manufactured with six-piston device made in China and powder catalysis is one of the best

diamonds for special application, for instance, making a band saw to cut single crystal silicon. The

main product of diamond (Annual output is about 800 million carats.) was produced with Chinese

six-piston device and its size ranges from 40/45 to 80/100, which is suitable for the manufacture of

cutting or grinding tools applied in processing less hard rocks. The diamond directly synthesized with

six-piston device and re-processing diamond are less important products in China, which can be used

as abrasives and its price is very low. CBN, PCD, PDC and CVD are also manufactured in some scale.

Annual output of diamond in China is great enough but its quality need improving.

Keywords: production, features, diamond made in China

INTRODUCTION

The synthetic diamond was first made in China in 1963. Since then the industry of synthetic

diamond and diamond tools developed rapidly. Now the production capability of diamond is about 1.5

billion carats a year and PCD, PDC and CBN also produced in a large scale.

The predominant equipments for producing diamond are hinged six-piston device. There are

about 4000 of 6X8MN, and about 1000 of 6X12MN and 6X16MN at present. The main features of the

equipments are as following:

Table 1. The features of the device mostly used in China

Types Chamber Production Production Production of Annual

space capacity greater than SMD greater production

cm 3 (carat) 60/70(carat) than 60/70(carat) (carat)

6X8 MN 9.5 8.5 6.0 2.0 200,000

6X12 MN 17.2 15.0 11.2 3.3 1000,000

6X16 MN 24.1 20.0 16.0 5.5 1200,000

Note: SMD is a label for higher quality diamond by Chinese national code.

Estimating parameters for diamond currently used in China are: mesh size-- the same as

American code; strength_ompressive strength for a single grain (kg/grain); and stacking density.

The average price of synthetic diamond is less than 0.2USD/carat. Production of CBN is about

50 million carats and average price less than 0.30USD/carat. The total value of production is less than

16 million US dollars.

NASA/CP--2001-210948 45 August 6, 2001
Thequalitiesof diamondsvarylargelydueto usingdifferentequipmentsanddifferent

techniques forsynthesis.

THEFEATURESOFINDUSTRIALDIAMOND

1. Diamond made with two-piston press.

There are about 20 of two-piston device in China. Most of them are imported from west

countries and some of them ranging from 1800T to 5000T are made in China. With 5000T type

equipment made by West techniques, top class of diamond by international standard can be produced.

The diamond shows golden color, transparent crystal, smooth crystal face, and distinct crystal edges

°°°°..%....... o%,.//./.......°,°°%.oO°.°Oo°o°_%%..

(seen Fig. 1). There is very little of impurity inclusions, almost !!!iiiiiiiiiiiiii!iiiii!!iiiiii!iiiii!iiiiiii!!!!iiiiii!i!!i!!!!iiiii!i?i

no bubbles and cracks. The shapes of most diamonds are iiiii::_::_i_!_ii:_:i:_ii::i_::::i]iiii!!_iii_:i!!i::::iii::_:_:_:::ii::ii::i::iii_!ii!ii::

•.%./.°o..°.o......°%°. °.. °°....°°° •.///. ° ....°°°°o°.,°°.

hex-octahedrons with hexagonal crystal faces (111 faces) and i:i:i:i:i:7:!'!'!:i:!:i:!:i:i:i:i:i:i:i:i:!:i:i:i:i:i:i:i:i:i:i:i:ii_i!:!

quadrilateral crystal faces (100 faces). Crystal planes grew

uniformly in three dimensions. The impact strength (TI) of the

diamond for 35/40 mesh size is as high as 90. Thermal impact

strength (TTI) reaches 90. The compressive strength is higher

than 40kg/grain by Chinese standard.

While with the two-piston device made in China, the

quality of diamond is much worse than above diamond. The

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

Figure 1. The best diamond

made in China

main differences exhibit in high impurity, low TI and TTI values. For diamond of 45/50, TI is about 74

and TTI only 65. Annual production for this kind diamond is about 50 million carats and the size

ranges from 40/45 to 45/50.

2.Diamond made with six-piston device by powder catalysis technique

More than 90% of the diamond made in this way has perfect crystal shapes, and most of them

are hex-octahedrons with hexagonal crystal faces and quadrilateral

crystal faces. The shape of the diamond looks good but there is a

little diffusion of inclusions and bubbles. Its TI value reaches

74-81 and TTI is about 63-68. The annual output is about 80

million carats and the grain sizes are 45/50, 50/60 and 60/70. This

kind diamond mainly used for the manufacture of saw like tools. It

perhaps is a good alternative for the high cost products mentioned

in above paragraph in international market on large extent because

of its high quality and lower price (Its price is as low as about 35%

of the diamond above mentioned.)

i:i:i:i:i:i:i:i:i:!:_g_:_:!:i:i:i:i:i_:.".!:!._*i.:i:i:iR

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

::ii ::ili!iiiii:#:ii iii iiiiii!iii::;iiiii i i!i ! 

Figure 2. Diamond made

with six-piston device

3. Fine diamond made with six-piston device by powder catalysis synthesizing technique

The diamond made in this way is a yellow green transparent crystal. 95% of it shows perfect

.°-,-,-°-,-,-,-°-°-,-.-°-,-,-,-¢,-.-. -.-,-°-o-°-.-°:,:°-.-°-o:%_.

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

hex-octahedral, and most hexagonal crystal faces and quadrilateral :::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

crystal faces are of excellent symmetry. The diamond has the little :iiiiii::i:iii!i:ililj:iii:ii!i!ii:ili:iii:iiiiiiii::i!i !,_ii:i:i:;i;

of impurity and magnetic performance, so that it is one of the best ".:.:-:.:..'.;.:.:.:.:..'.:.'.:.'.:.'.'-'.:.:.:_.4-:.;.:.:.:-:.:.:. !

::.:..:.:.'.'.:.:.:::::::::::::::::::::::::::::::::::::::::::::

fine diamonds for special applications The diamond is suitable for

making tools for cutting or polishing precious crystals (for instance,

cutting single crystal silicon.). We have electroplated the band saw

with the diamond for cutting single crystal silicon, which set up a Figure 3.

long life record of continuously used for more than one month. The produced

grain size of the diamond ranges from 100/200 to 270/325. With method.

Fine diamond

by conventional

NASA/CP--2001-210948 46 August 6, 2001
limitationofapplications, theoutputofit confinesinasmallamountstill.

4. Diamond made with six-piston device by slices of catalysis

The diamond is a main product in China. Its shape is also

hex-octahedron but with poor symmetry and there is black

impurity inclusions bubbles and cracks in it in some extent. The

better one may have tiny radial line of impurity. TI of it is only 60;

TTI is about 45; and compressive strength is less than 16kg for

40/50. Its size ranges from 40/45 to 80/100, among them 50/60 and

60/70 are better. The diamond can be used to manufacture the tools

for cutting or grinding less hard marble. Annual output of it is

about 800 million carats and the price is relatively low.

5. Fine diamond directly synthesized with six-piston device

i:i:?" ::::::::::::::::::::::::::::::::::::::::::::::::

.o.o. °°°°:-;._.i.'....-......:.;.:.........'.:.;.;"."

g ii !i!!i:::i!iNNNiii !::iiii!

,.o ::.:.:.:.:.:.::l_.:::.:.:.::::. 1.1.:...;:

Figure 4. Main product of

diamond in China

The diamond is a yellow green crystal ranging from 100/120 to 325/400, most of them falling

into the scope from 170/200 to 200/230. 90% of the diamond is

nearly equal volume crystals, but most of crystals growth

uncompleted. The edges of crystals are not distinct; Crystal face is

not smooth and of less transparency; more defects inside and high

brittleness. The diamond of this kind is suitable for the

manufacture of diamond grinding tools, resinoid wheels and

metallic wheels. The price of the diamond is low and output is

about 20 million carats.

- -:;::

Figure 5. Directly synthesized

fine diamond.

6. Ro- rocossin .iamon. "" "i Fine diamond of high purity can be obtained by breaking the i:i::.:.: ." :"

coarse grains ranging from 35/40 to 50/60, in which a lot of impurity _'_ """:° .e

inclusions exist, exposing the inclusions and then eliminating them _:_:_i_:!_:_:_:_:____!_'i!i_!:i:!:i_'_!_

by chemical method. The fine diamond is often employed as !!i:!::i!!!iii::ii!iiii_i::__'_.

diamond abrasives and the annual output is about 50 million carats.

7. CBN production Figure 6. Re-processing

CBN is produced commonly with six-piston device made in diamond.

China. CBN generally demonstrates black and amber color. The

CBN with good transparency, amber color and less impurity is made up

in China but the crystal shape much worse than diamond. CBN is

divided into two types: I and II. The shape and compressive strength

of type II is better. The annual output of CBN is about 50 million carats

and its size ranges from 50/60 to 325/400. The price of black CBN is

less than amber one.

8.

" ... "'i:jii:!:i:i:i!iii

Figure 7. CBN product.

PCD, PDC, CVD

There are a complete series of PCD, PDC, CVD in Chinese market currently.

FUTURE PROSPECTS

A great progress has been made for the industry of super-hard material not only in the output

but also in technology Some manufacturers has grasped of advanced technique to produce the diamond

of high quality The device of 6X16 MN and 6X20MN increase by 10% annually, and the type of

6X30 MN is designing. In the near future, the device less than 6X8 MN will be eliminated and focus on

promoting the device greater than 9000T. Advanced products and synthesizing techniques will develop

NASA/CP--2001-210948 47 August 6, 2001
firstinnextwoorthreeyearsandremovethedominancefromoutputtoqualityandprice.

References

1. Hn Jiangen and Wang Gnangzn, "Joint in WTO and developing strategy of synthetic diamond

industry", Diamond Industry, 2001,1 (In Chinese)

2. A. Bakon and A. Szymanski, P.Daniel (translation Editor) "Practical uses of diamond", ELLIS

HORWOOD and PWN-POLISH SCIENTIFIC PUBLISHERS, 1993

3. Sun Ynchao, Lin Yibo and Wang Qinsheng, Diamond tools and fundamentals of metallurgy,

Chinese Construction Material Industry Publishing House, 1999

NASA/CP--2001-210948 48 August 6, 2001
3

DIAMOND APPLICATIONS

Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

CHARACTERIZATION OF A1N AND LiNbO3 THIN FILMS DEPOSITED ON

DIAMOND SUBSTRATES

M. Ishihara, T. Nakamura and Y. Koga

Research Center for Advanced Carbon Materials,

National Institute of Advanced Industrial Science and Technology,

AIST Central 5, 1-1-1 Higashi, Tsukuba, Ibaraki 305-8565, Japan

m.ishihara@aist.go.j p

F. Kokai

Joint Research Consortium of Frontier Carbon Technology, JFCC,

c/o AIST Central 5, 1-1-1 Higashi, Tsukuba, Ibaraki 305-8565, Japan

ABSTRACT

Diamond has become interest for the high velocity materials because the elastic constant is highest and the

surface acoustic wave (SAW) velocity is more than 10,000 m/s. Although diamond has not a piezoelectricity, its high

acoustic propagation makes it a desirable substrate for SAW devices when coupled with piezoelectric thin films such

as aluminum nitride (A1N) and lithium niobate (LiNbO3). Highly oriented A1N thin films and LiNbO3 thin fills were

prepared by a reactive sputtering method on polycrystalline diamond substrates at low substrate temperature. The

average of surface roughness (Ra) of the A1N thin fills was below 2 nm by locating the diamond substrates at the

position of 100 mm from the aluminum target. The full width at half maximum of the rocking curve for AIN(002)

peak of X-ray diffraction was about 0.2 degree. SAW characteristics with an interdigital transducer

(IDT)/A1N/diamond structure were investigated. The average surface roughness of LiNbO3 thin films was 5 nm.

When the highly oriented LiNbO3 fills are deposited on diamond substrate, the IDT/LiNbOg/diamond layered

structure is capable of a wide-bandwidth application in SAW devices at high frequencies.

Keywords: Nitride, Oxide, SAW device, Preferential orientation, Surface roughness

INTRODUCTION

The development of high-frequency devices has become technologically important because of the increasing

volume of information and communication media, such as mobile telephone and satellite broadcasting. Many

component devices such as bandpass filters resonators must also conform to this high-frequency band. Dielectric

ceramics and surface acoustic wave (SAW) devices have conventionally been used for high-frequency devices (ref.

1). However, the dielectric filters and duplexers are too large for portable communication devices, SAW devices

have become popular with the advantage of small size and free filter characteristics.

Various approaches for developing high-frequency SAW devices have been attempted. The center frequency of

the SAW devices is determined by the equation Fo Vp/2_, where Fo, Vp and _ are the center frequency, the phase

velocity of the SAW materials and the wavelength, respectively (ref. 2). The wavelength of,_ depends on the size of

lines and spaces of the interdigital transducers (IDTs). New materials with a high phase velocity or a sub-half micron

rule process are required for the development of high-frequency devices. However, reduction of the IDT size raises

problems such as reliability, power durability and fabrication margins in the manufacturing process (ref. 3). Recently,

the use of diamond for high velocity materials has attracted interest, because the elastic constant is highest and the

SAW velocity is more than 10,000 m/s. Although diamond has no piezoelectricity, its high acoustic propagation

makes it a suitable substrate for SAW devices when coupled with piezoelectric thin films such as zinc oxide (ZnO)

(refs. 3 and 4).

Aluminum nitride (A1N) has a hexagonal wulzite structure and the lattice constants are a 3.112 A and c

4.982 A (ref. 5). A1N has been used as an electric insulator with heat radiation property. Since the phase velocity of

A1N is highest (6700 m/s) among those of piezoelectric materials such as ZnO (5700 m/s) that is easy to obtain the caxis orientation, the hig_frequency SAW devices can be expected from the layered structure of A1N thin films on

the diamond substrates (ref. 6). Additionally, the velocity dispersion can be expected small because the difference of

NASA/CP--2001-210948 51 August 6, 2001
phasevelocitybetweenA1NanddiamondissmallerthanthatofbetweenZnOanddiamond. A1Nthinfilmshave

attractedincreasinginterest forapplicationtoSAWdevicesoperatingattheGHzbandinwirelesscommunication

technology(ref.7).It iswellknownthatsingle-crystal A1Nthinfilmscanbeepitaxiallygrownona sapphire

substratebychemical vapordepositionatthesubstratetemperature ofmorethan1273K(refs.8and9).Sincethe

smoothsurfaceofA1Nthinfilmscannot beobtainedbytheCVDmethodbecauseofgraingrowth,it is impossible

tofabricateamicropattern IDTonas-grownA1Nthinfilms(ref.8).

Lithiumniobate(LiNbO3)isaveryimportant ferroelectric material withaCurietemperature of 1483K and

hasarhombohedral structurewithplanarsheetsofoxygenatomsinadistortedhexagonal closepackedarrangement

(ref.10).Ithasuniquepyroelectric, piezoelectric andnonlinear opticalproperties, whichmakeitsidealmaterialsfor

fabricationofsurfaceacousticwave(SAW)andopticaldevices. ThesinglecrystalLiNbO3iswidelyusedinSAW

devicesanditsSAWphasevelocityisabove4000m/s.Particularly, theelectromechanical couplingcoefficients (k2)

ofthesinglecrystalLiNbO3is larger(5.5%)thanthatofquartz,A1NandLiTaO3(0.14,0.7and0.64%,

respectively) (ref.11).WhentheLiNbO3thinfilmsaredepositedonthediamondsubstrate,heLiNbO3/diamond

layeredstructureisapplicabletoawide-bandwidth applicationinSAWdevicesathighfrequency band.

Inthispaper,A1NandLiNbO3thinfilmsweredepositedondiamondsubstrates atlowsubstratetemperature by

reactivemagnetronsputtering. Theinfluenceofthesputteringparameters onthesurfacemorphology andtheacoustic

propagation wasinvestigated.

EXPERIMENTAL

A1N and LiNbO3 thin films with the thickness of 1 [tm were prepared by a DC and an RF magnetron sputtering

system, respectively. A water-cooled A1 target (99.999% purity) and LiNbO3 target (stoichiometric concentration)

with a diameter of 50.8 mm was used as a planer magnetron cathode. The discharge mode of the DC sputtering

system was power-constant mode and the experimental power was 300 W. The RF power supply (Advanced

Energy: RFX-500) was used with a power-regulated mode of 100 and 50 W. The substrates used in this experiment

were diamond-coated silicon (Sumitomo Electric Industries Co., Ltd.) substrates without pre-cleaning such as argon

ion etching. The thickness of the diamond films was about 12 [tm and the surface had been polished mechanically.

The substrates were heated at the substrate temperature of 623 K. The distance L between target and substrates was

50 and 100 mm. After the sputtering chamber was evacuated below 5×10 -7 Torr, sputtering gases were introduced

into the vacuum chamber with regulation by mass flow controllers. At the deposition of A1N thin films, total gas

pressure was controlled to 3×10 -3 Torr by operating a gate valve, and the flow rate of Ar and N2 gases (99.9995%

purity) was 2 and 8 sccm, respectively. At the deposition of LiNbO3 thin films, total gas pressure was controlled to

1×10 -3 Torr by operating a gate valve, and the flow rate of Ar and 02 gases (99.9995% and 99.99% purity,

respectively) was 8 and 2 sccm, respectively.

The crystalline quality and orientations of the films were determined by X-ray diffraction analysis (XRD: CuKa,

40 kV, 30 mA) in a 20-0scan mode and a 0 scan mode. The surface morphology was observed by atomic force

microscope (AFM).

RESULTS AND DISCUSSIONS

A1N thin films

Figure 1 shows typical AFM images for A1N thin films with the thickness of 1 [tm deposited on a diamond

substrate at the distance L of 100 mm. The uniform and smooth surface was obtained on the diamond substrate. The

crystal size and surface roughness Ra were 30 nm and less than 2 nm, respectively. There has been remarkable

progress of surface flatness; diamond substrates with a mechanically polished surface roughness below 3 nm are

commercially available. Since the density of nucleation for the A1N crystal was uniform on the flat substrate, the

A1N thin films grew smooth, too. The average surface roughness was sufficiently level to deposit an aluminum

electrode for the IDT onto the as-grown A1N thin films. The A1N thin films deposited on the diamond substrate

exhibited high adhesion by the peeling test.

Figure 2 shows an XRD pattern of the A1N thin films grown on the diamond substrate at the distance L of 100

nm and the substrate temperature of 623 K. The diffraction patterns show expected diamond(ll 1) plane from the

diamond substrate together with A1N(002) reflection in 20/0 scan mode. No other A1N reflection can be observed in

the measurement conditions. The XRD pattern shows that the A1N thin films are highly oriented along the c-axis

NASA/CP--2001-210948 52 August 6, 2001
normal tothesubstrate. Thefullwidthathalfmaximum(FWHM)oftherockingcurvefortheA1N(002) peakwas

about0.2°andthatfortheSi(400)peakfromthediamondsubstratewasabout0.1° incoscanmode.Theseresults

suggest thatheA1Nthinfilmsdepositedonthediamondsubstratewerehighlyorientedalongthec-axis normal to

the substrate. It is a first report that the c-axis oriented A1N films were deposited on polycrystalline diamond

substrates.

Shikata et aI. (ref. 6) calculated the phase velocity Vp and the coupling coefficient k 2 of the IDT/A1N/diamond

structure at various kh (k = 2rc/)_, )_ is SAW wavelength, h is thickness of A1N thin films). At the optimal value of

kh = 1.57, Vp and k 2 were 11,200 m/s and 0.7 % for the 1st mode wave (Sezawa wave), respectively. Standard IDTs

were fabricated on the surface of the A1N films and a frequency response for the GHz-band filter with the

IDT/A1N/diamond structure was measured on wafer by a network analyzer (ref. 11). The linewidth of the IDT was 1

gm (=)_/4, )_; SAW wavelength), and the number of IDT finger and the aperture length were 100 pairs and 400 gm,

respectively. The experimental Vp obtained from the center frequency and the SAW wavelength was 10,120 m/s

(2.48 GHz), which was almost in agreement with the theoretical Vp. The experimental results of Vp and F0 were

better than that of the A1N/sapphire structure (Vp = 6780 m/s, F0 = 1.211 GHz) (ref. 6). The diamond SAW

substrate was effective in the phase velocity of the A1N/diamond SAW filter.

LiNbO3 thin films

During the deposition process at the RF power of 100 W, many cracks were caused often in the LiNbO3 target

(thickness of 4 mm). It seems that the cracking of the target was caused by a thermal shock from the RF plasma. It

is necessary to lower RF power to prevent the target from the cracking. When the RF power was decreased to 50 W,

the LiNbO3 target was not cracked during the deposition.

The deposition rate increased by locating the substrates close to the target from 100 mm to 50 mm, and the

color of the deposited films changed colorless and transparent. It was found that the low RF power and the distance

L were also effective to prepare the colorless LiNbO3 thin films though the deposition rate slightly decreased in

comparison with that of the deposition at 100 W. A typical AFM image for the LiNbO3 thin films was shown in

Fig.3. The LiNbO3 thin films with a smooth surface and transparency were obtained on the diamond substrate. The

crystal size and surface roughness Ra were 100 nm and 5 nm, respectively. There has been remarkable progress of

surface flatness; diamond substrates with a mechanically polished surface roughness below 3 nm are commercially

available. Since the density of nucleation for the LiNbO3 crystal was uniform on the flat substrate, the LiNbO3 thin

films grew smooth, too. The average surface roughness was sufficiently level to fabricate aluminum electrodes for the

SAW devices onto the as-grown LiNbO3 thin films.

100 nm

Figure 1. Typical AFM image for AIN thin films with the thickness of 1 gm

deposited on a diamond substrate at the distance L of 100 mm.

NASA/CP--2001-210948 53 August6,2001
-3

v

.m

if?

eeo

o

v

z

"_ -0.5 0.0 0.5

C co(deg.)

0

E

D

|

......... J_ .... 1 ...........

20 30 40 50 60

2e (deg.) (Cu KoO

Figure 2. X-ray diffraction patterns of AIN films

deposited on a diamond substrate at

the distance L of 100 mm. The

corresponding rocking curve for the

AIN(002) peak is shown in the inset.

Figure 3. Typical AFM image for LiNbO3

films deposited on a diamond

substrate without substrate heating.

CONCLUSION

In conclusion, A1N and LiNbO3 thin films with smooth surfaces were prepared on polycrystalline diamondcoated silicon substrates at low substrate temperature of 623 K. We have succeeded in fabricating an

IDT/A1N/diamond structure. Highly c-axis oriented and stoichiometric AIN thin films were obtained on a diamond

substrate using reactive DC magnetron sputtering method. The average surface roughness Ra was below 2 nm by

locating the diamond substrates at the position of 100 mm from the A1 target. The experimental value of SAW

velocity was 10,120 m/s. IDT/A1N/diamond SAW devices operating at around 2.5 GHz were prepared and tested.

The LiNbO3 thin films were obtained on a diamond substrate using reactive RF magnetron sputtering method. The

surface roughness Ra of the LiNbO3 thin films was 5 nm.

REFERENCES

1. H. Okano, et al.: Jpn. J. Appl. Phys., vol. 31, 1992, 3446.

2. Y. Shibata, et al: Jpn. J. Appl. Phys., vol. 34, 1995, L320.

3. K. Higaki, et al.: IEEE Trans. Ultrason., Feerroelect. Freq. Contr., vol. 42, 1995, 362.

4. H. Nakahata, et al.: Jpn. J. Appl. Phys., vol. 34, 1998, 2918.

5. S. Strite and H. Morkoc: J. Vac. Sci. Technol. B, vol. 10, 1992, 1237.

6. S. Shikata, et al.: New Diamond and Frontier Carbon Tech., vol. 9, 1999, 75.

7. Y. Kobayashi, et al.: Jpn. J. Appl. Phys., vol. 34, 1995, 2668.

8. K. Tsubouchi and N. Mikoshiba: IEEE Trans. Sonics Ultrason., vol. SU-32, 1985, 634.

9. G. Radhakrishnan: J. Appl. Phys., vol. 78, 1995, 6000.

10. R. S. Weis and T. K. Gaylord: Appl. Phys. A 37, 1985, 191.

11. M. Ishihara, et al.: Jpn. J. Appl. Phys. (submitted).

NASA/CP--2001-210948 54 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

HIGHLY ORIENTED DIAMOND THIN FILMS FOR SILICON-ON-INSULATOR

TECHNOLOGY

S.D. Wolter, G.N. Yushin, F. Okuzumi, R. Schlesser, and Z. Sitar

Department of Materials Science and Engineering, North Carolina State University, Raleigh, NC 27695, USA

J.T. Prater

Army Research Office, Research Triangle Park, NC 27709-2211, USA

B.R. Stoner

MCNC, Materials and Electronic Technologies Division, Research Triangle Park, NC 27709, USA

ABSTRACT

A study of pulsed bias-enhax_ced nucleation of highly oriented diamond on (100) silicon is reported. In

contrast to the conventional direct current substrate biasing, a pre-caxburization was not needed in the processing of

the aligned diamond. Pulse biasing at frequencies from 1 Hz to 2 kHz did not show any frequency dependence and

produced a constant percentage of oriented diamond nuclei of nearly 50%. In contrast, a variation of the duty cycle

of a 60 Hz square waveform from 3% to 75% showed a hyperbolic decrease in the percentage of highly oriented

diamond nuclei, with a maximum of-45% occuring at the shorter duty cycles.

A preliminary study addressing the issue of wafer bonding of diamond to silicon was also performed.

Highly polished, polycrystalline diamond (RMS roughness -15.0 nm) was bonded to single-side polished, (100)

silicon in a dedicated ultrahigh vacuum bonding apparatus. Direct bonding under an applied uniaxial pressure of

-32MPa was observed at temperatures above 950°C, as evidenced by infrared and acoustic imaging techniques.

Partial bonding of the diamond to silicon was observed at 950°C and 1050°C, whereas, complete bonding was

evidenced at 1150°C and 1200°C.

Keywords: Pulsed bias-enhanced nucleation, highly oriented diamond, silicon-on-diamond, fusion bonding

INTRODUCTION

Silicon-on-insulator (SOI) technology has been the subject of interest for its favorable influence on metaloxide-semiconductor (MOS) devices. The conventional SOI technology uses a buried silicon dioxide layer which

effectively reduces parasitic capacitances leading to nearly ideal MOS device operation. One detriment of this

approach is its poor thermal conductivity, leading to less efficient heat dissipation from the active silicon layer [1-4]

and deteriorated device attributes [5].

Diamond's outstanding thermal and dielectric properties seem ideally suited for use as the buried insulator

in the SOI scheme. Bulk, crystalline diamond possesses the unique combination of high room temperature thermal

conductivity and electrical insulating properties. Although the application of chemical vapor deposited diamond

may find utility in such thermal management applications, the thermal transport properties of polycrystalline

diamond are limited by thermal scattering at the grain boundaries [6-8] and at the diamond/substrate interface [9-

11 ]. In this regard, it is imperative that steps be taken to limit the presence of defects within the diamond films.

Highly oriented diamond (HOD) [12,13], with properties close to those of natural diamond, offers a great

potential as the passive heat conducting layer in the SOI structure. The oriented nature of the diamond crystallites is

anticipated to significantly enhance the thermal transport characteristics of the polycrystalline films. This paper

discusses heteroepitaxial diamond growth using a pulsed bias-enhanced nucleation treatment, followed by

presentation of preliminary results on fusion bonding of (100) silicon to highly polished, polycrystalline diamond.

This latter experimentation is prerequisite to wafer bonding of the highly oriented diamond to silicon.

EXPERIMENTAL APPROACH

Pulsed bias-enhanced nucleation

Single-side polished, (100) silicon substrates were used in the nucleation and growth experiments. The

substrates, 2.5 cm x 2.5 cm in dimension, were cleaned in acetone and methanol and thoroughly rinsed in de-ionized

NASA/CP--2001-210948 55 August 6, 2001
Table1.Diamonducleationandgrowthconditions.

Parameters Pulsed bias-enhanced nucleation Diamond deposition

CH4/H2 2% 1%

Pressure 15 Torr 25 Torr

Microwave power 900-1100 Watts 700 Watts

Substrate temperature - 850°C 730°C

'ON' state bias voltage -250 volts 0

'ON' state bias current -0.350 A 0

'OFF' state bias voltage 0 volts 0

'OFF' state bias current 0 A 0

water, placed on a molybdenum substrate holder and immediately transferred to a microwave plasma chemical

vapor deposition system. A 2.45 GHz ASTeX TM microwave power supply with a 1.5 kW capability was used to

generate a flee standing plasma within the growth chamber. The substrate holder containing the samples was

isolated from ground, enabling the substrates to be electrically biased.

The pulsed bias-enhanced nucleation (BEN) experiments involved the use of a square waveform bias

potential. The voltage, provided by a frequency generator, was amplified to a desired value using a voltage

amplifier. An 'ON' and 'OFF' state bias potential of-250 V and 0 V were used, respectively. Both the bias voltage

and effective bias current were monitored using an oscilloscope; the effective bias current by measuring the voltage

across a 10 fl resistor. The process parameters used during the BEN and diamond deposition processes are

summarized in Table 1.

A variation of the bias voltage duty cycle for a 60 Hz square waveform was studied first. Duty cycles

ranging from 3% to 75% were investigated, as shown schematically in Figure 1. A subsequent study addressing the

influence of bias frequency (1 Hz and 2 kHz) was also undertaken employing a square waveform with a duty cycle

of 17%. In both studies, a pre-caxburization step was not used as a means of facilitating oriented diamond

formation; only the pulsed substrate biasing was implemented. The percentage of aligned crystallites and the

duration required to form a diamond film were the parameters monitored in these investigations. The latter data was

qualitatively ascertained by visually inspecting the substrate during the biasing experimentation. Previous

experiments determined the onset of diamond film formation to be evidenced by a dark film formed on the silicon

substrate. Scanning electron microscopy (JEOL 6400f) was used to image the diamond crystallites following the

nucleation and growth experiments.

Silicon-on-diamond fabrication

The bonding experiments were conducted in a dedicated ultrahigh vacuum apparatus. The applied load and

fusion temperature were the important control parameters in this investigation. Highly polished, randomly oriented

diamond films (RMS roughness of-15.0 nm) obtained from Kobe Steel USA Inc. and single-side polished, (100)

silicon wafers (RMS roughness of <1.0 rim) were used in these experiments. The polycrystalline diamond was

nominally 25 gm thick grown on -1 mm thick silicon.

The silicon wafers (1 cm x 3 cm) and the CVD grown diamond (3 mm x 4 mm) were cleaned in acetone

and methanol, and rinsed in de-ionized water to remove surface contaminax_ts. The samples were then affixed to

separate molybdenum sample holders and placed in a cleaned container. The whole procedure was done in a Class

100 pm-ticnlate hood. The container was subsequently transferred to a glovebox connected to a load-lock which was

interfaced to the bonding apparatus. The glovebox was continuously purged with hepa-filtered air enabling entrance

of the samples into the UHV system while limiting the amount of particulate contamination.

The bonding experiments were designed to evaluate the influence of fusion temperature on the quality of

the silicon-to-diamond bonding. Temperatures of 850°C, 950°C, 1050°C, 1150°C, and 1200°C were examined

under an applied pressure of-32MPa and duration of 15 hrs. The samples were subjected to the applied load only

after the desired temperature was attained, which generally took 2 hrs. Following the high temperature bonding, the

bonded specimen was slowly cooled to room temperature over a period of 5 hrs in the absence of any load. The

specimen was subsequently removed from the system and examined using infrared (IR) and acoustic imaging.

NASA/CP_001-210948 56 August 6, 2001
RESULTS AND DISCUSSION

Pulse biasing of highly oriented diamond on silicon

The results of the pulsed bias-enhanced nucleation investigation revealed that the cyclic action of the

substrate bias is essential for obtaining oriented diamond in the absence of a pre-carburization surface treatment.

Past work by this laboratory [12] has found it important to use a carburization pretreatment prior to the conventional

direct current BEN in order to produce close to 50% oriented nuclei.

An indication that the duty cycle is a critical aspect of the nucleation enhancement process as well as the

aligned diamond formation was revealed in the bias time to diamond film formation results. This time was found to

be inversely proportional to the duty cycle; a similar functional dependence was observed for the fraction of oriented

diamond nuclei. Substrate biasing using a 17% duty cycle and a bias time of -60 rain provided nearly 50% oriented

diamond, as seen in Figure 2a. At a duty cycle of a 50%, a bias duration of only -20 rain was required for film

formation, however, only about 20% of the diamond was oriented, as seen in Figure 2b. Although the hyperbolic

trend in our data suggested that the fraction of oriented nuclei can be increased beyond 50% for low duty cycles, the

authors observed a saturation just below this value, as seen in Figure 3. This suggests that the diamond/substrate

lattice mismatch may be limiting the nucleation of epitaxial diamond and can not be improved by optimized

processing.

Wavelimn

peliod Duty cycle _ ,

il 0 v, 'OFF' state 3q_,

.................. -250 v, 'ON' state

6% 1

17% U

25% U--

33% ] [--

50% [__ --

Figure 1. Plots of the various duty cycles examined in this investigation to determine their influence on the

formation of highly oriented diamond. (Note: The waveform period is 0.016 sec)

Figure 2. Scanning electron micrographs of diamond produced using a square waveform substrate bias involving a

(a) 17% duty cycle and (b) 50% duty cycle followed by 20 hrs of diamond deposition.

NASA/CP--2001-210948 57 August 6, 2001
6O ° t

50 t

,_ 40

3O

20 q

_ 10

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9

Duty cyCle

Figure 3. Plot of highly oriented diamond percentage versus duty cycle. A hyperbolic functionality is evident for

most of the data.

?

"¢:3

g

>-,

6O

50 T [ T

40

30

20

10

0- ' ....... I ........ r ' '

lxl0 q lxl0 ° lxl01

..... I ....... I ...... !I .....

lxl02 lxl03 lxl04 lxl05

Bias licequency (Hz)

Figure 4. Plot of the fraction of highly oriented diamond versus bias frequency. The data reveals relatively little

influence on the oriented crystallite formation with bias frequency.

A variation in the bias frequency using a square waveform potential with a duty cycle of 17% was also

performed to evaluate the influence of frequency on the HOD formation. As seen in Figure 4, there is no

appreciable change in the fraction of oriented diamond in the examined frequency range. However, it is interesting

to note that the film formation bias time decreased linearly from -60 min to -40 min when increasing the bias

frequency from 1 Hz to 2 kHz.

Fusion bonding of diamond to silicon

Wafer bonding is a promising approach to obtaining monocrystalline silicon on high quality diamond and

was targeted in this study. The task of obtaining smooth diamond samples is rather difficult based on the inherent

hardness of diamond. For this study the roughness of the polished diamond films displayed extreme roughness

values as large as 50.0 nm, as determined by atomic force microscopy, although typical RMS surface roughness

values closer to 15.0 nm were observed. Surface roughness generally aggravates the bonding process, as has been

experienced in silicon-silicon fusion experiments [14]. Given these relatively large surface roughness values, the

preliminary successes in bonding observed in this study were considered quite promising.

NASA/CP--2001-210948 58 August 6, 2001
Table2.General observations forthesilicon-on-diamond bondingexperimentation.

BondingtemperatureSpecimens bonded Bondedarea

850°C No

950°C Yes partial

1050°C Yes partial

1150°C Yes complete

1200°C Yes complete

Figure5showsaphotograph ofthemolybdenumholdersafterawaferbondingexperiment at1200°C,

showingthefusediamond-coated siliconsubstrate bondedtothelargersiliconwafer.Ingeneral, allthespecimens

subjected to1150°Corhigherbondingtemperatures atfusionpressures of-32MPaweresuccessfully bondedover

thewholearea,asdetermined byIRimaging.Partial bondingwasobservedattemperatures below1150°C.The

general observations fromtheseexperiments areummarized inTable2.

TheacousticimageinFigure6showsthediamond-silicon coupleproducedat1050°C. Itrevealsthathe

waferswereonlypartiallybonded.Theauthorsaxecurrentlyevaluatingwhether amodification intheappliedload

atthelowerfusiontemperatures mayenhancetheareaofbonding. Theacousticimage(andschematic) inFigure6

showcrackingwithinthediamondlayer,asevidenced bythedarklinearfeaturesintheimage.Ithasbeenobserved

thatheapplicationfanon-uniformloadledtoincompletebondingbetweenthediamondandsiliconaswellas

crackingintheregionwheretheloadwasconcentrated.

i!i!i_!!_i_i!i!iii!iiiiii!iiiiiii!iiiii!iiii!!i!!i!ii!!!!!iiiii!iiiiiiiiii!!

iiiiiiiii!iii !:!iiii!iiii!! iiii ii!iii!iiiiii . iiiiii

iiiiiiii!iii?!ii_Z X_2!!!_i_iiiii_i_i.".&.'i_iiiiiiii!iii

iiiiiiiiiiiiiiiiiii iii iiiiii

Figure 5. Diamond and silicon samples after fusion bonding.

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::: __YER ::::Unb;nlded ::::::::

::::::::::::::::::::::::::::::::::::::::: Ci_g_I:::: :::::::::

Ii:i:il

:::: ::::::::::::::::::::::::::::::::::::: :: :::::

_i:i:i:i:i:i:i:i_sl'.ta._e/Ni$$_gi:i:i:i:i:i:i:i:l

Figure 6. Acoustic image (a) of fusion bonded diamond to silicon and (b) the schematic representing the various

aspects of the bond region. (Note: The graphic reveals the diamond was grown selectively into pads of 3 mm x 4

mm in dimension; Each specimen consisted of 1 1/2 diamond pads)

NASA/CP_001-210948 59 August 6, 2001
Summary

Pulsedbias-enhanced nucleationfhighlyorienteddiamondon(100)siliconhasbeenevaluatedinthe

context ofdutycycleandfrequency oftheappliedbias.Thepulsingactionofthesquarewaveformwasimportant

fortheprocessofformingtheepitaxial diamondintheabsenceofapre-carburization surfacetreatment. Theduty

cycleshowedastrongeffectofthebiastimeontheonsetofdiamondfilmformationandfractionoforientednuclei

thatdeveloped. Conversely, thebiasfrequency hadverylittleinfluenceontheepitaxial diamondformation.

Thewaferbondingofdiamond-to-silicon revealedfusionbondingattemperatures above950°Catan

appliedpressureof-32MPa.Uniformbondingovertheentirediamondfilmsurfacewasobservedforfusion

temperatures above1150°C.

ACKNOWLEDGEMENTS

We would like to acknowledge the funding for this research which was provided for by the Army Research

Office, grant number DAAG5598D0003.

REFERENCES

1. M. Berger and Z. Chai, IEEE Trans. Electron. Dev. 38, 871 (1991).

2. L.T. Su, K. Goodson, D. Antoniadis, M. Flik, and J. Chung, Tech. Digest IEDM, 357 (1992).

3. W. Redmax_-White, M.S. Lee, B.M. Tenbroek, MJ. Uren, and RJ.T. Bunyan, Electron. Letters 29, 1180 (1993).

4. D. Yachou and J. Gautier, ESSDEC Conf., 787 (1994).

5. J.-P. Colinge, in Silicon-on-insulator Tecnology: Materials to VLSI, 2 na Edition, Kluwer Academic Publishers,

Boston, 223 (1997).

6. O.W. Kfiding, E. Matthias, R. Zachai, H.-J. Ffi[3er, P. Mfinzinger, Dia. Related Mater. 2, 1185 (1993).

7. W.F. Banholzer and T.R. Anthony, Thin Solid Films 212, 1 (1992).

8. C. Gu, Z. Jin, X. Lu, G. Zou, J. Zhang, and R. Fang, Thin Solid Films 311, 124 (1997).

9. S. Jin, L.H. Chen, J.E. Graebner, M. McCormack, and M.E. Reiss, Appl. Phys. Lett. 63, 622 (1993).

10. J.E. Graebner, S. Jin, G.W. Kammlott, Y.-H. Wong, J.A. Herb, and C.F. Gaxdinier, Diamond. Relat. Mater. 2,

1059 (1993).

11. M. Berger and Z. Chai, IEEE Trans. Electron. Dev. 38, 871 (1991).

12. S.D. Wolter, B.R. Stoner, J.T. Glass, PJ. Ellis, D.S. Bnhaenko, C.E. Jenkins, and P. Southworth, Appl. Phys.

Lett. 62, 1215 (1993).

13. X. Jiang, C.-P. Klages, R. Zachai, M. Hartweg, and H.-J. Mfisser, Appl. Phys. Lett. 62, 3438 (1993).

14. Q.-Y. Tong and U. G6sele, Semiconductor Wafer Bonding: Science and Technology, The Electrochemical

Society, Inc., Pennington, New Jersey, 1999.

NASA/CP--2001-210948 60 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

CVD DIAMOND FOR HIGH TEMPERATURE ELECTRONICS

Susan L. Heidger, Navid J. Baraty and Joseph A. Weimer

U.S. Air Force, AFRL/PRPE, Wright-Patterson AFB, OH 45433

Tel: (937) 255-6016, Fax: (937) 255-3211

Email: Susan.Heidger @ wpafb.af.mil

ABSTRACT

Polycrystalline Diamond films produced by Microwave Plasma Enhanced Chemical Vapor Deposition

(MWCVD) were investigated for high power and high temperature electronic applications. The diamond films were

deposited on tungsten substrates and on silicon substrates coated with metal electrodes. Due to the difference in the

thermal expansion coefficients of diamond and tungsten, cohesive freestanding diamond films ranging from 10btm to

150btm thick were obtained as the diamond films parted from tungsten substrates as the samples cooled after

deposition. Test devices were fabricated on both freestanding diamond and diamond on silicon substrates.

Dielectric characterization was performed as a function of frequency and temperature. Voltage breakdown, I-V

characteristics, insulation resistance, resistivity and thermal conductivity measurements were also performed. The

film quality was assessed using SEM, XRD, Raman spectroscopy and electrical characterization. High quality CVD

diamond films had dielectric constants ranging from 5 to 7 and loss tangents less than 0.001 over the frequency

range of 100 Hz to lxl06 Hz at room temperature, resistivities on the order of 1014 ohm-cm and breakdown voltages

approaching that of natural diamond. Applications of CVD diamond in high temperature and high power passive

devices and switches will be discussed.

Keywords: CVD diamond, high temperature passive, dielectrics

NASA/CP--2001-210948 61 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

POLYCRYSTALL1NE DIAMOND FILMS FOR HIGH-TEMPERATURE

ELECTRONICS AND MEMS

R. Ramamurti, V. Shanov, and R. N. Singh

Department of Materials Science and Engineering, University of Cincinnati, P.O. Box 210012,

Cincinnati, OH 45221-0012

M. Samiee and P. Kosel

Department of Electrical and Computer Engineering and Computer Science, University of Cincinnati,

P.O. Box 210030, Cincinnati, OH 45221-0030

ABSTRACT

High-temperature electronics and MEMS (Micro-Electxo-Mechanical Systems) based on polycrystalline

diamond (PCD) appear to be promising because of its wide bandgap, high thermal conductivity, and large carrier

mobilities. The primary objective of this research is to develop techniques for the synthesis of both doped and

undoped high quality PCD films with good flatness and surface roughness suitable for the fabrication of high

temperature electronic and MEM devices. The PCD films are deposited by the microwave plasma enhanced chemical

vapor deposition (MPCVD) method for the deposition of PCD films on Si (100) substrates using methane in a

hydrogen plasma environment. The influence of various process parameters on the crystal quality, and properties of

the PCD films are studied. The films are characterized by Raman Spectroscopy, SEM, AFM, and X-Ray Diffraction

techniques. Initial evaluation of the electrical properties is done involving measurement of current-voltage and

capacitance-voltage characteristics of test structures fabricated on free-standing 10 gm thick PCD films at both room

....... 11 12

and elevated temperatures. Typmal film reslshvltms axe m the range of 10 - 10 _.cm while the dielectric

constants are near 5.6, which is representative of natural diamond. These PCD films are being used for the

fabrication of single and multilayer PCD capacitors for high temperature characterization and operation.

Keywords: diamond, microwave plasma CVD, electrical properties, capacitors, electronics

_TRODUCTION

High value capacitors used in power electronics are based on ceramic dielectrics, but at high temperatures they

have limited lifetime and often fail in operation. Recent advances in the growth of polycrystalline diamond have led

to the production of thick coarse-grained films on silicon, which have very high resistivity and breakdown strength

approaching that of single crystal diamond (refs. 1 to 5). Several properties of diamond such as: its large bandgap

energy, exceptional mechanical hardness and chemical inertness, high melting point, large thermal conductivity,

very high resistivity, and dielectric strength, make it promising over other semiconductors (ref. 6). These properties

make PCD attractive for use as a dielectric in capacitors for high power electronic circuits and also for MEMS.

Many promising electronic device applications of thin diamond films are expected in the near future (refs. 7 to 12).

The objective of this research is to produce high quality polycrystalline diamond films with good flatness

and low surface roughness, and to investigate ways of the fabricating electronic devices. The devices addressed here

are primarily for high temperature operation while the MEM devices axe for sensor and actuator applications, which

may not necessarily be for high temperature operation. Diamond is almost an ideal candidate material as an insulator

in capacitors. Preliminary research is done to investigate methods for fabricating capacitor test structures on freestanding PCD films, and related characterization techniques for determining of PCD film properties.

EXPERIMENTAL

Experimental Set-Up

An Electron Cyclotron Resonance (ECR)-PMCVD facility has been created and used for PCD thin film

deposition (refs. 13 to 14). This facility is based on an ASTeX magnetized microwave plasma source (ref. 15) as

shown in Fig. 1. It consists of a 1.5 kW magnetron power supply at 2.45 GHz, microwave components,

symmetric plasma coupler with quarts window, window and exit electromagnets for ECR mode, downstream

chamber, and movable r.f. heated platform. In addition, the facility is supported by a turbomolecular pump, gas

flow and pressure control units, cold traps, and a water cooling unit. This system can produce high quality

polycrystalline diamond films in the upper deposition chamber at high pressures such as at 40 torr. Our ECR

NASA/CP--2001-210948 62 August 6, 2001
sourceiscoupledtoaseparatedownstreamdepositionchamber, whichallowsusto generateplasmaatevenlower

pressureranges, downtothe2xl0.5torrlevel.

3 Stub Tuner

Symmetric plasma

coupler

Circulator/

Dummy Load

Window Magnet

ECR Source

Precursor downstream

injection

Substrate graphite plate.

heatable to 900 C

low fteq biased (11313 - 4.513 HZ)

Pumping and

instrumenlalion port

Moto[ dr_ve for heiqht

adjustment

Figure 1. Schematic of the ECR-Microwave Plasma ECR-PECVD

Materials, Deposition Conditions, and Characterization of Diamond Films

Silicon wafers, p-type (100) diced into squares (25 x 25 x 0.5 mm) were used as substrates for deposition of

PCD. Standard wafer cleaning prior to deposition was applied in a controlled environment, which included

sequential ultrasonic agitation in ethanol, acetone, ethanol, de-ionized water, and drying in Ar jet (refs. 16 to 17).

The surface activation procedure prior deposition included treatment of the Si substJcate for 2 hours in an ultrasonic

bath of ethanol and diamond powder (20-40 pm slurry). The process was followed by rinsing in acetone and in

ethanol and drying in a N2 before placing into the reactor. Prior to deposition, a 30 minute etching was done in the

hydrogen plasma to clean the Si substrate off the native oxide layer.

The diamond films were deposited under different microwave-plasma deposition conditions in order to observe

changes in film characteristics because of the different process parameters such as temperature, total pressure, and

flow rates of H2 and CH4. High quality PCD was grown using conventional conditions (ref. 18), which were

modified to fit the nature and geometry of the substrates and the MPCVD system. An optimized set of growth

parameters included; substJcate temperature of 750°C, pressure of 4 kPa (30 torr), hydrogen gas flow of 500 sccm,

methane gas flow of 5 sccm, microwave power of 900 W, and deposition time of 5 hours. The growth rate was

about 0.4 pm/h and a 10 pm thick PCD films were easily obtained in 24 hours using the above listed deposition

parameters. Our approach allows better control of the deposition parameters during PCD growth, which yields

smooth and uniform diamond films (ref. 19).

A scanning electron microscope (Cambridge Stereoscan 90) was used to examine the surface morphology of the

samples. Micro-Raman (Instrument SA, T 64000, Jobin Yvon triple monochromator system, equipped with an

optical multi-charmel detector-CCD array) was employed for identification of the CVD diamond films. The output

power of the Ar + laser (514.5 nm wavelength) was 10 mW, and it was focused up to 2 gm beam size. The surface

roughness was studied by using Atomic Force Microscope, model Autoprobe CP Unit, made by

Thermomicroscope.

Some of the diamond films were studied by x-ray diffraction using an X'pert- Philips model PW3040. The Xray source was a Cu target (Cu K_) and the voltage on the x-ray tube was 25 KV. The range of angle (20) was set

between 40 ° and 145 ° to cover all the important diamond peaks at a scanning rate of 2.5 s per step size of 0.02 °.

NASA/CP--2001-210948 63 August 6, 2001
RESULTS AND DISCUSSION

For the construction of single and multilayer capacitors, the PCD films need to be only thin, in the range of 4-

• 2 • 2 •

6 gm, but they must be umform over areas up to 25.8 cm (4 m ). The growth surfaces of PCD films generally tend

to be rough with typical surface roughness being about 10% of the PCD thickness. In this research, growth

techniques have been developed for the production of uniform PCD films of 10 gm thickness with low surface

roughness.

Some of the deposited films were selected for examination by SEM. Two types of samples were prepared;

fractured cross-sectional view and top view which are shown in Fig.2 (a) and (b). The PCD, depending on the

deposition conditions, had very fine grains and hence low roughness. The morphology at higher magnification was

found to be generally faceted, and consisted of a combination of {111 } and {110} facets (Fig.2 (b). PretJceat_nent

with the 20-40 gm slurry seems to be the best as for as film morphology and flatness are concerned. The crosssectional view of the films shown in Fig.2 (a) indicates that the ultrasonic treatment produces films of uniform

thickness• The films are dense, well adhered, and continuous•

::::::::::::::::::::::::::::::::::::::::::: :>!'!'!?!?:':':':'!':"

(a)

_; ".". .... . _. "-.:.:.._'..

•;." .;.:.:.:.1.;.:.:.'<

_.. ".:

(b)

•...4::'>.::_._

(c)

Figure 2. SEM micrograph of diamond film deposited on Si (100); ultra sonic activation of the substrate

in diamond slurry (20gm grids), 750°C, 30 torr, 1.0 % CH4, 500 sccm H2, 900 W, 24 h growth time; (a)

cross sectional view; (b) plan view; (c) three-dimensional AFM image of the as-grown film

To obtain information on planography and surface roughness, AFM in a contact mode is employed• Figure 2

(c) shows a three dimensional AFM image of PCD film. The obtained root mean square (RMS)2 surface roughness is

about 42.9 nm and the average roughness is about 34.4 nm for a scanned area of about 935 gm. This surface

roughness is closer to nanocrystalline diamond films and about one order of magnitude smaller than the reported

polycrystalline diamond roughness (ref. 20). These smooth diamond films are very useful for further processing of

electronic and MEMS devices•

The CVD diamond films were characterized by Raman spectroscopy• Typical Micro-Raman spectrum of

1

polycrystalline diamond obtained in this research is shown in Fig. 3 (a). The sharp peak at 1332 cm- is the first

.... 3

order phonon peak of dmmond, and is assigned as the dmmond (sp bonding) component of the film. The high

intensity of this peak with respect to the background fluorescence indicates good crystal quality• The broad peak

2 ....

centered around 1550 cm -1 is attributed to the sp bonding of non-dmmond forms of carbon (graphite and possibly

amorphous carbon)• The sensitivity of the Raman technique to sp _ (non-diamond) bond_ed phase of carbon is 50

times greater than that of sp 3 bonded phase (ref. 21). Thus the peak around 1550 cm represents much smaller

amount of the non-diamond phase in our films (ref. 21 to 22).

NASA/CP--2001-210948 64 August 6, 2001
45O

400

u

350

g

0

o 300

25o

J_

j_

jJ

1

I

200 ....

1100 1200 1300 1400 1500

Wavenumber cm-1

1600

N.......... :N"

"iii:iii'ii'!ii::ii'i'ii:ii:::':'!'"2"i'i)ii!5i'i'i

*1.5>2.1.1.>>>i-_._.i.1.1.)i...i.i*i.1.1.i.1*i*_.:*i.M*1°1°i.1*1*:._°_

(a) (b)

Figure 3. (a) Micro-Raman spectrum of PCD film on Si (100); ultra sonic activation of the substrate in

diamond slurry (20gm grids), 850°C, 60 torr, 0.5 % CH4, 300 sccm H2, 900 W, 5 h growth time; (b) PCD

DOT capacitor test structures fabricated on free-standing 10 gm thick PCD film

Fabrication of Free Standing Diamond Films and Test Capacitors Based on PCD

The electrical properties of PCD films were measured by creating test capacitors by (a) high-rate etching of the

crystalline host silicon in sulfur hexafluoride (HF6) plasma, and (b) developing a non-lithographic technique for

patterning the silicon to produce free-standing PCD films on which circular (abbreviated: DOT) test capacitors were

then fabricated. The free-standing PCD areas were produced by etching away the host silicon over window areas of

5x5 mm 2. A simple method adopted using a shadow mask, which consists of long metal stripes separated by spaces

of equal width. Aluminum which was found to be resistant to SF6 plasma is deposited through the shadow mask

over the silicon surface. This produced an array of square windows through which the silicon can be etched away in

the SF6 plasma. After etching, the PCD structure had the appearance of arrays of translucent PCD squares supported

on a rigid silicon grid, which provided physical strength to the structures. Such free-standing diamond films can be

used for pressure sensors based on optical diaphragm displacement. The high strength and high thermal stability of

the PCD allows the sensors to operate at elevated temperatures. Research on such sensors is currently in progress.

The ground plane of the capacitor array was fabricated by sputtering a 0.4 gm thick continuos Ti film on the

rough PCD surface, and sequentially thermally evaporating 0.1 gm thick gold layer on top of Ti. The DOTs on the

opposite (etched Si) side of the capacitors were formed by the same procedure using a shadow mask with circular

windows. Details this fabrication and testing of diamond-based capacitors are given elsewhere (ref. 23). The final

DOT capacitor test structure is shown in Fig. 3 (b).

Electrical Characterization of PCD Test Capacitors

The DOT capacitors built on thin PCD films were very fragile and easily fractured during testing with

conventional sharp-tipped tungsten probes. To overcome this problem, gold coated cats-wisker probes were

constructed and mounted on a hot chuck. PCD samples designated MSE1 have been subjected to the following set

of measurements of the PCD test capacitors: (a) the measurement of individual test capacitors at room temperature,

and (b) the measurement of current-voltage (IV) characteristics at room and elevated temperatures. The capacitances

were measured at 1 MHz at room temperature, and the IV measurements were done at 25°C, 50°C, and 100°C over

the voltage range of 0-200V. The results for MSE1 are summarized in Fig. 4 (a) and (b).

NASA/CP--2001-210948 65 August 6, 2001
g

1.8E-07

1.6E-07

1.4E-07

1.2E-07

1.0E-07

8.0E-08

6.0E-08

4.0E-08

2.0E-08

0.0E+00

_25C ]

--_--50C /

100 C]

C = 9.4 pF /

Sample MSE1 !

Window #5 /

/

/

/

/

/

/ ./

/ /

/I 1,1'

1/

I .

_=a- - ._--_-" ,_/

0 50 100 150 200

Voltage (V)

(a)

1.OE+13,

1.0E+12

1.0E+ll

1.0E+10,

1.0E+9,

1.0E+8

AverageCapacitance=83pc

SampleIVBE1

• "_ckness=5um

• 1Rckness=9um

u u u u u

40 80 120 160 200

Voltage (V)

(b)

Figure 4. (a) Current-voltage characteristic of PCD sample MSE1 (window#5, located in the middle of

the capacitor array); (b) Differential resistivity versus voltage for sample MSE1 at 25°C

The PCD samples yielded very narrow spreads in the capacitor values indicating good uniformity of the film

thickness over the 6.45 cm 2 (1 inch square) size. The results from other windows were quite similar across the

individual test capacitors/windows of each sample. In general, the current increased with increasing applied voltageFig 4 (a). Initial PCD films show a characteristic 150V knee in the current-voltage characteristics. This may be a

feature of the polycrystalline nature of the fills and needs to be further investigated. Plots of differential resistivity

as a function of voltage are shown in Fig. 4 (b). The zero bias resistivities axe then obtained by extrapolation. A

summary of the electrical data is given in Table 1.

Table 5. Electrical characteristics of PCD film MSE1 grown on Si (100); ultra sonic activation of the

substrate in diamond slurry (20gm grids), 750°C, 30 torr, 1.0 % CH4, 500sccm H2, 900 W, 24 h growth time

Sample ID

MSE1

PCD thickness

(gm)

10.2

Average

Capacitance

(p_)

8.3

Dielectric constant

5.41

Zero bias

resistivity (tl-cm)

>1013

The values for the dielectric constant are close to that for natural diamond. However, some error is expected because

the exact thickness of the PCD was difficult to obtain. Similar data regarding the resistivity and the dielectric

constant of diamond fills are reported by other groups (refs. 1 to 5).

SUMMARY AND CONCLUSIONS

An ECR-MPCVD facility for advanced thin fill processing was created. Polycrystalline diamond films on

(100) silicon substrates were deposited from methane in a hydrogen microwave plasma environment. The nature of

the diamond film was characterized by SEM, micro-Raman spectroscopy, AFM, and X-ray diffraction. Excellent

diamond films with very good flatness and law surface roughness were fabricated for application in high temperature

electronics and MEMS. A non-photolithographic technique for the fast fabrication of arrays of parallel-plate (DOT)

........ 12

capacitors on thin free-standing PCD films was developed. High PCD film reslsltlVltles exceeding 10 tl.cm with

dielectric constant of about 5.6 were obtained. These films appeared to be adequate for capacitor and MEMS

applications. Techniques for the fabrication of multilayer capacitors and optical fiber-based pressure sensors for high

temperature operations are under investigation.

NASA/CP--2001-210948 66 August 6, 2001
ACKNOWLEDGMENTS

TheauthorswouldliketothankMr.BobJonesforhelpwiththe AFM. This material is based upon work

supported by the National Science Foundation under Grant No. ECS-0070004. Any opinions, findings, and

conclusions or recommendations expressed in this material are those of the authors and do not necessarily reflect the

views of the National Science Foundation.

REFERENCES

1. Jany C., et al.: Post-Growth Treatments and Contact Formation on CVD Diamond Films for Electronic

Applications. Diamond and Related Materials, vol. 7, 1998, p. 95.

2. Jany C., et al.: Influence of the Growth Parameters on the Electrical Properties of Thin Polycrystalline CVD

Diamond Films. Diamond and Related Materials, vol. 9, 2000, p. 1086.

3. Jany C., et al.: Influence of the Crystalline Structure on the Electrical Properties of CVD Diamond Films.

Diamond and Related Materials, vol. 5, 1996, p. 741.

4. Tzeng Y., et al.: High Electrical Resistivity Diamond Films Deposited from an Oxyacetylene Flame. Appl.

Phys. Lett., vol. 57, no. 8, Aug 1990, p. 789

5. Davidson, J.L., Brown, W. D.; and Dismukes, J. P.: CVD Diamond, A Unique Dielectric, Semiconductor and

Thermal Conductor for Electronic Applications. Fourth International Symposium on Diamond Materials, K.V.

Ravi, J. P. Dismukes, J.L. Davidson, K.E. Spear, R.H. Hauge, and B. V. Spitsyn, editors, May 1995, p. 537.

6. Lin H., and Dandy D.: Diamond Chemical Vapor Deposition, Nucleation and Early Growth. Noyes

Publications, New Jersey, 1995, p. 2.

7. Pan, L. S., et al.: Diamond, Electronic Properties and Applications. Kluwer Academic Publishers, Boston,

MA, 1995.

8. Yoder M. N.: Applications of Semiconducting Diamond and Related Materials. 72/SPIE: Diamond-Fihn

Semiconductors, vol. 2151, 1994, p. 72.

9. Edwards L.M., and Davidson J.L.: Fabrication Process Development and Characterization of Polycrystalline

Diamond Film Resistors. Diamond and Related Materials, vol. 2, 1993, p. 808.

10. Kosel P., et al. : Device-Quality Polycrystalline Diamond by Microwave Plasma-Enhanced CVD. SPIE Proc.

Diamond-Film Semiconductors, vol. 2151, 1994, p. 121.

11. Kosel P., et al.: UV/VIS Photodetectors in Wide-Bandgap Semiconductors: GaAs and Polycrystalline

Diamond. SPIE Proc. Photodetectors: Materials and Devices, vol. 2685, 1996, p. 140.

12. Kosel P., et al.: Device Fabrication in Polycrystalline Diamond. Proc. of '98 CARTS Conf., 1998, p. 77.

13. Shanov V., Tabakoff W., and Singh R. N.: CVD Diamond Coatings for Erosion Protection at Elevated

Temperatures. presented at the ASM International 2000, Materials Solutions Conference, Surface Engineering

Symposium, October 9-12, 2000, St. Louis, Missouri.

14. Shanov V., Singh R. N., and Tabakoff W.: CVD Coatings for Erosion Protection at Elevated Temperatures.

Invited presentation at the International Conference on Metallurgical Coatings and Thin Films, April 10-14,

2000, San Diego, California.

15. ASTeX User Manual, : HPM/M Magnetized HPMS Plasma Source, Version 1.6, Woburn, MA, 1992, p. 5.

16. Behnahi M., et al.: Influence of Mechanical and Chemical Silicon Surface Preparation on Diamond Nucleation

and Growth in CH4AI2 System Discharge, Surface and Coatings Technology vol. 106, 1998, p. 53.

17. Sun C., et al.: Nucleation of Diamond Films by ECR-Enhanced Microwave Plasma Chemical Vapor

Deposition. Diamond and Related Materials, vol. 8, 1999, p. 1410.

18. Wu R., et al.: Synthesis and Characterization of Fine Grain Diamond Films. J. Appl. Phys., vol. 72, no. 1,

1992, p 110.

19. Jayaseelan V. S., et al.: Processing of Polycrystalline Diamond Films by Microwave Plasma CVD. Accepted

paper for presentation at the American Ceramic Society Annual Meeting, April 22-25,2001, Indianapolis,

Indiana.

20. Zhou D., et al.: Control of Diamond Film Microstructure by Ar Additions to CH4/H2 Microwave Plasmas. J.

Appl. Phys., vol. 84, no. 4, Aug. 1998, p. 1981.

21. Bou P. and Vandenbulcke L. Raman.: Investigation on Diamond Films and Crystals Deposited by PlasmaAssisted CVD. J. Electrochemical Soc., vol. 138, no. 10, 1991, p. 2991.

22. Knight D. and White W. Characterization of Diamond Films by Raman Spectroscopy. J. Mater. Res., vol. 4,

no. 2, 1998, p. 385.

23. Kosel P.B., et al.: High temperature polycrystalline diamond capacitors. CARTS 95:15th Capacitor And

Resistor Technology Symposium, 13-16 March 1995, p. 255.

NASA/CP--2001-210948 67 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

DIAMOND MEMBRANE ELECTRON MULTIPLIER

Nick N. Dzbanovsky, Pavel Minakov, Nikolay V. Suetin,

Institute of Nuclear Physics, Moscow State University, Vorobjevy Gory, 119899 Moscow, Russia;

ndzb(&dmw.phys.msu.su

Eugeny A. Poltoratsky, Gennady S. Richkov, Eugeny A. IL'ichev, Sergey A. Gavrilov,

Zelenograd's Physical Problems Institute, Moscow, Russia.

1. Introduction.

The main barrier for electron emission in metal/diamond/vacuum system is at the back metal contact.

Emission is limited by the electron tunneling from the back contact into the diamond conduction band,

rather than from the conduction band into the vacuum. For overcoming this problem we have suggested to

inject electron in diamond membrane at the back side from external electron source. Diamond consists of

carbon atoms, which have low atomic number, and therefore it has high transparency for low energy

electrons. Moreover, diamond is one of most chemical and mechanical stable material and has the highest

secondary electron emission coefficient (more than 100). So, diamond membrane may be used for electron

multiplication and direct studying of diamond emission properties. In this report, we inform about

measurement of secondary electron emission yield from cesiated diamond membrane.

2. Diamond growth and characterization

Diamond films were grown by microwave plasma CVD on p- type Si {100} substrates. The

growth parameters were : MW power 800 1000 W, Wsubstrat e _ 850°C, Total pressure - 80 Torr. The

working gas mixture was composed by hydrogen with ethanol (or methanol) vapor in proportion 94% to

6% accordingly. Boric acid dissolved to saturation in methanol and trimethylboron (B (CU3) 3), dissolved in

ethanol were used as boron sources. For preparing the diamond film membranes, the holes with size of 2

6 mm in diameter were etched in the back side of Si substrates.

The secondary-electron yield was studied on experimental setup shown in Fig. 1.

...

.' 2

4*"

9

3

*- 4 ' _* ' " 5

.'i" ¢r

0_. Ra

Fig.1 Experimental setup for studying the secondary electron emission from diamond membrane (1-

electron gun, 2-electron beam, 3-Si substrate, 4-diamond film, 5-dielectric holder, 6-detector, 7-amplifier,

8-Cs source, 9-voltage supplier).

NASA/CP--2001-210948 68 August 6, 2001
3.Resultsand Discussion.

It was found that the secondary-electron yield is very sensitive to the incident electron energy and the

surface termination with hydrogen and Cs. Influence of external electric field on secondary emission

properties was measured. It was found that the secondary-electron yield essentially increase when

external electric field was applied. Fig. 2 show influence of external field Ut on secondary-electron yield K

for different energy of primary electrons (V-15 kV,, - 20 kV, • - 25 kV and V- 30 kV). From these results

we can conclude that enhancement of secondary electron yield by external electric field is higher for high

energy of primary electrons.

k.

-03 -45 -33 -15

3

K

0 15

Fig.2. Dependence of secondary electron yield on external field Ut.

The influence of cesiation time (thickness of Cs layer) on secondary-electron yield K is presented on Fig. 3.

Fig.3. Dependence of secondary electron yield on the time of Cs deposition

NASA/CP--2001-210948 69 August 6, 2001
Fromthisresultwecanmakethenextconclusions:

- mostpartofCVDpolycrystalline diamondfilmhavepositiveelectronaffinity,

-secondary electronyieldisverysensitivetothethicknessofCslayer,

-Cesiateddiamondmembrane maybeusedaseffectiveelectronmultiplier.

ACKNOWLEDGMENTS Thisworkwaspartiallysupported bygrantsfromtheINTAS(Project 98

1754)andISTCGrantNo.1187

NASA/CP--2001-210948 70 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

Electrical properties of boron-doped diamond films grown by bias method.

S.Yugo,T.Tashiro and D.Saito

University of Electric-Communications

1-5-1,Chofu gaoka, Chfu- shi,Tokyo, Jap an

(e-mail;yugo@diamond.crc.uec.ac.jp)

(Keywords; doped diamond films, bias method, electrical properties)

The new method using the bias was proposed as a method for doping the impurity to the

diamond. Boron dope diamond film was produced in order to evaluate the effectiveness

of this method, and the electrical characteristic was examined. This new doping method

is based on the principle which applies the negative bias in the diamond CVD equipment

in installed electrode for the impurity. The electrode for the impurity with the negative

bias supplies the impurity in the plasma by ion-sputtering. It is possible to promote

diamond growth with the positive bias by the electron shower. As the result, various

impurity addition are enabled, while the diamond thin film is kept high-quality. The

boron-dope diamond film was produced by this method, and the electrical properties were

examined. First, by usual bias processing method, the diamond film for the insulation

was produced on n-Si, and the boron-dope diamond film was produced on the substrate.

In the impurity source, Ti-B composite with the electrical conductivity was used, and the

bias voltage changed from 0 to 300V. As the result, the good boron-doped diamond film

was obtained. The electrical resistivity were about 10-2 f_ -cm, and the Hall mobility

were about 200-1000 cm2/v.s., p-type were recobnized from Hall effect and Zeebeck

effect. The activation energy of the boron was about 0.3eV. It is shown that the

mobility of doped diamond films by this method are very high, and that the method of this

impurity addition is effective for controlling electrical characteristic by the bias voltage.

(Abstractl,ADC2001)

NASA/CP--2001-210948 71 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

NOVEL ELECTRONIC DEVICES USING DIAMOND THIN FILMS

A. Vaseashta

Department of Physics & Physical Sciences

Marshall University

Huntington, WV 25755

ABSTRACT

The excellent electrical, optical, and mechanical property of diamond thin films is the driving force for its

current research in producing thin films of diamond. Based upon the unique properties of diamond, enormous stateof-the-art novel electron devices and sensors devices are possible. Several device configurations are provided of

hybrid microelectronics, opto-electronic, and other novel devices realizable with the diamond material. Thermal

management of electronic devices is extremely important in the design of sensors and high-speed high-density

electronic circuits with power devices fabricated monolithically. A preliminary theoretical computation for

comparison of the power dissipation vs. gate propagation delay for diamond, silicon, and gallium arsenide based

devices is presented. Thermal interaction of semiconductor devices on the substrates as it relates to their close

proximity and spacing is addressed. It is suggested that based upon diamond thin films, as a microelectronic lateral

extension to complexity, RAM on processors, system on chip (SOC), network on chip, and superchips are possible.

Keywords: Diamond, Electronic Devices, Thin Films, Microelectronics, and Hybrid.

INTRODUCTION

The speed, functions, and reliability of the modem day integrated circuits continue to increase with decrease in

chip size. The physical size and the speed consideration are severely limited by the properties of the substrate

materials. GaAs material has higher electron velocity and is used for high-speed devices. However, it suffers from

several disadvantages due to its poor thermal conductivity. Standard semiconductor processing techniques are not

readily applicable for GaAs. Hence, in search of an alternate material, thin films of diamond are currently under

extensive investigation. Due to its physical characteristics diamond has tremendous applications (ref. 1 to 3),

ranging from freestanding monolithic substrate material to high-speed high-density multichip microelectronic

circuits, opto-electronics, packaging modules, optical, biomedical, and tribological applications. The aim of this

paper is to demonstrate the suitability of diamond thin films as a novel electronic material from application

standpoint.

AN OVERVIEW OF DIAMOND CHARACTERISTICS

Thin films of diamond are SP 3 bonded carbon atoms arranged in a cubic or hexagonal crystal structure (ref. 4).

Both the cubic and the hexagonal crystal structures are suitable for electronic and optical applications. Current

research does not indicate any preference towards either structure for enhancement of properties. Diamond is the

hardest material known to mankind. It occurs in nature in Kimberlite rock formation and often displays a

characteristic color of the dominating impurity. Several desirable electrical and optical characteristics result from its

high band gap (5.45 eV). Due to wide energy gap and high electron mobility, semiconducting diamond films can be

used for optical faster device operation such as in communication and computation. Insulating diamond has a high

intrinsic resistivity (_ 1 x 1017 _2-cm.) which allows closely packed devices to have a high degree of isolation, thus

reducing the effects of parasitic capacitance. A dielectric strength on the order of (= 1 x 108 V/cm.) enables the

fabrication of high voltage and high power devices. Other important properties of diamond include high thermal

conductivity of (10-20 W/cm.-K) and controlled high temperature coefficient of resistivity. Hence, diamond

substrates find tremendous applications in thermistors. Diamond has an effective transmission range starting just

above the 230 nm. wavelength range and extending all the way up to 40 _tm. range, the broadest spectral

transmission of all solids. Semiconductor diamond hence is the best candidate for tandem junction photovoltaic

devices, ultra fast room-temperature optical (UV-VIS-IR) radiation detectors, opto-electronic & photonic switches,

and sensors. Due to its wide wavelength range and optical transmissive property diamond in intrinsic form (natural

or synthetic) is considered as a high efficiency window material and is an ideal candidate for fiber-optic couplers.

Employing diamond and other wide bandgap materials, monolithic integration of Opto-electronic Integrated Circuits

NASA/CP--2001-210948 72 August 6, 2001
(OEIC)andApplicationSpecificIC's(ASIC)intosinglechipnowseemslikelytorevolutionize thSystemonChip

(SOC)technologies, whichislikelytoreducethepackagingcostwithincreaseinproduct reliability, ascomparedto

itsdiscretealternatives. Thehighestpeedsoundvelocityandlowspecificdensitypropertyofdiamondhas

potentialsforapplications inupersonic soundtransmission mediaforcouplers. Theimpedance matching

characteristics ofdiamondlayersiidealforacousticsupersonic soundgenerators andtransducers. TableI.shows

thephysical, electrical, optical,andmechanical propertiesofdiamondandotherwidebandgapmaterials. Thetable

showscomparison ofvaluesoffigureofmeritamongSi,GaAs,SiC,& Diamondbaseddevices, aswell.

NUCLEATION AND SYNTHESIS OF DIAMOND FILMS

The mechanism of diamond nucleation and growth is not yet clear. However, diamond deposition techniques

can be broadly categorized into (a): processes operating under non-thermal equilibrium and (b): processes operating

under local thermal equilibrium. An extensive literature exists providing explanation of the methods of diamond

deposition. Most commonly employed deposition techniques such as microwave plasma assisted chemical vapor

deposition, electron cyclotron resonance (ECR), and hot filament CVD (HFCVD) processes operate under nonthermal equilibrium conditions. The deposition rate of such processes range typically from (0.1-5.0 gm/hr). In the

deposition techniques employing LTE e.g. thermal plasma CVD and oxy-acetylene flame the ions and electrons are

at the same temperature. The presence of abundant atomic hydrogen helps etch the graphitic bonds which leads to

enhanced diamond growth and renders the LTE CVD process an attractive technique for diamond deposition. In an

R.F. thermal plasma system, a high power RF field is applied to ionize a confined volume of gas, while in oxyacetylene torch flame temperature creates the ionized plasma. The carrier and the process gases are injected axially

and at an angle by a unique gas dispensing mechanism, so as to produce a swirl or vortex of the reactive species

around substrate. The gas velocities are adjusted such that the reactive species interact with the substrate in less than

2-3 msec (ref. 5). Employing such modification of the existing systems (ref. 6 to 9) growth rates on the order of 200-

300 gm/hr, can be realized. Table 2 lists a summary of diamond deposition growth techniques and typical growth

rates.

The study of radicals contributing to the growth of diamond films is critical to realize the synthesis of good

quality diamond. The radical densities are studied under different growth conditions to clarify the mechanism of

diamond film Growth (ref. 10), allowing controlled deposition from the nucleation stage onwards. A modified

residual gas analyzer (RGA) equipment as radical identification technique (RIT) is employed to detect the

contributing radicals, in an integrated deposition and in-situ characterization system (ref. 11). The unique RIT

design of the system provides detection of radicals and an in-situ Raman evaluation for crystallinity. One of the key

steps in CVD processes is to relate the processing parameters with the characteristics of deposited films. The

experimental permutations necessary to achieve the optimum results are typically prohibitively large. One of the

commonly employed techniques is Taguchi's method. Mathematical simulation of the processes using concept of

designed experiments is typically used to optimize the process yield.

Despite of the large impetus on the growth of diamond films several fundamental issues need to be addressed:

1): Influence of Structure: It is known that the structure and morphology of diamond films depend critically on

deposition conditions. It is also known that the local order or micro-crystallinity of diamond thin films can be

changed by many orders of magnitude. What is the influence of the change in structure, morphology, and microcrystallinity on electronic transport properties? How are the valence band tails, conduction band tails, mid-gap

defects, capture cross-section, electron and hole drift mobilities, and lifetimes affected by the structure? 2): Effect of

Hvdr_en: Some results indicate that atomic H is responsible for preferentially etching the graphitic component and

hence annihilates energetically possible configurations other than diamond. Other evidence indicates that certain

radical bearing carbon and hydrogen are responsible for the growth of diamond. Since the temperature of the

formation is extremely high, hydrogen desorption takes place leaving behind diamond films with long range order.

The exact mechanism for the role of hydrogen in surface reconstruction is not clear at the present time. Is it bonded

both in clusters and randomly in "standard" diamond network? Information about how and where H is bonded and

how its bonding environment changes if H is substituted with F needs further studies. Research on the influence of

H on the electronic properties and effect of substitution by F needs to be systematically studied. 3). Influence of

impurities and doping efficiency: Diamond can be easily doped p-type with boron while doping n-type is subject of

active research (ref. 12 to 15). Nitrogen is a well-known deep donor in diamond, while the only confirmed shallow

n-type dopant is Li, which is believed to occupy an interstitial position and diffuses rapidly at elevated temperatures.

Is doping diamond n-type related simply to a movement of Fermi level, or is it more basic relating to bond

NASA/CP--2001-210948 73 August 6, 2001
Table 1: Comparison of Silicon, GaAs, and Diamond Properties + and Relative Figures of Merit.

Parameter

Band gap 1 (eV)

Hole mobility

(cm2/V-sec)

Electron

mobility

(cm2/V-sec)

Acceptor level

(eV)

Donor level

(eV)

Resistivity

(f2-cm)

Breakdown 2

field (V/cm)

Saturation 3

velocity (cm/s)

(e-electron, h-hole)

Dielectric

Constant

Thermal Conductivity

(Watt/cm-Kelvin) 4

Si GaAs SiC

1.11 1.43 2.9

1350 300

480 8000

Diamond

5.4-5.5

1200

2000

0.045(B) 0.035(Si) 0.40(B)

0.045 (P) 0.0058(Si) 1.7(N), 0.1 (Li)

105_10 -4 108_ 1016 (IIa)_10 -2

3XlO s 4XlO s 2.5X10 6 10 7

1XIO 7 1XIO 7 2.0XlO 7 (lh-2.7e)XlO 7

11.7 15.8 9.7 5.7

1.45 0.46 4.9 20.0

Comparison of Values of Figure of Merit Among Si, GaAs, SiC, GaN, & Diamond

Material Johnson Key Bali_a

Si 1 1 1

GaAs 6.9 0.45 9.5

GaN 281 1.76

SiC 695 5.12

Diamond 8206 31 357

1. High-temperature operation, blue light emission, UV detection, low leakage current devices.

2. High-power, high-density integration.

3. High-frequency operation in high electric fields.

4. High-density integration.

tDiamond Properties

Crystal structure Cubic Fd3m

Lattice Parameter 0.356725 nm

Atom Density 1.77 x 10 27 cm -3

Hardness 5700 - 10,400 Kg/mm 2

Young's modulus 10.5 x 1011 N/m 2

Anisotropy 1.21

Coefficient of thermal expansion 0.8 x 10 .6

Specific heat 6.195 J/°K.mol

Debye Temperature 1890

Refractive index 2.42

Hole effective mass 0.75

Electron effective mass 0.57

NASA/CP--2001-210948 74 August6,2001
TABLE 2 Summary of Diamond Deposition Growth Techniques

Technique Temo (°C) Pressure (torr) Flow (slm) Growth Rate (/_m/hr)

Microwave 800-1000 5-100 0.1-1 0.1-5

ECR 400-600 0.1-50 0.1-1 0.08-0.1

Hot filament 700-1000 10-100 0.1-1 1-5

DC discharge >600 5-760 0.1-1 20-250

DC arc jet 500-760 >10 100-1000

Combustion flame 800-1000 760 1-2 150

RF torch 700-1200 500-760 -10 -50-200

RF plasma 800-1000 1-30 0.1-1 0.5-3

rearrangements? An investigation to identify doping mechanism by shallow donor (< 0.5 eV) needs to be conducted.

FUTURE SCOPE OF DIAMOND BASED DEVICES

Future scope of novel electron devices based on the unique properties of diamond is currently widely

investigated. Predictions based upon their physical properties indicate that devices prepared from these materials

should perform better than currently available devices in terms of power, frequency and short wavelength

applications. Fig. 1 shows a material response chart of several commonly used semiconductors and diamond. It is

evident that optical, electrical, & mechanical characteristics of diamond are far superior to other semiconductors. A

preliminary theoretical computation employing semiconductor transport equations and specific device transfer

characteristics were conducted. A comparison of the power dissipation vs. gate propagation delay for diamond,

silicon, and gallium arsenide based devices is shown in Figure 2(a). Computations of simulated analytical

expressions indicate that diamond devices can withstand higher temperatures of operation without significant

deterioration in electrical performance as shown in Figure 2(b), as reported in literature (ref. 16 to 17) showing

ideality factor for diodes approximately equal to 1.01 even at 1000°C operation temperature. The thermal

management of electronic devices is extremely important in the design of high-speed high-density electronic circuits

especially with power devices fabricated monolithically. Thermal interaction of semiconductor devices on the

substrates as it relates to their close proximity and spacing is particularly important in microminituarized devices.

Organic materials and ceramic substrates, e.g. SiC, A1N, & BN are used in the present high density chip mounting

technology. Diamond substrates are best suited for such applications due to a thermal conductivity of 10-20 W/cm.-

K. Field effect transistors (FET) (ref. 18 to 19) prepared on diamond substrates show improved drain-to-source

current and a uniform transconductance across the substrate. Employing transport parameters of thin diamond films,

thermal models are solved using a modified version of Thermal Analysis of Multi-layer Systems (TAMS) program.

The method uses Fourier Series to solve the full 3-dimensional heat conduction equation in multi-layer structures.

Materials Response function

Transmission

Regions

E/ecMcal

Re_

AConductivh_

MechanicaJ

Hardness

Si

GaAs

Diamond

(in morons) [Ol_cal}

o.1 1.o 5.0 1o 50 10o

Dia. mondpure) (doped)

Si

GaAs

Conductivity (S/cm_ ) [Elec_ical]

le-18 le-10 le_ 1 le5 Dismonlo_

SiC

Gau_

Indentation Hardness /Mechan/ca/]

le2 le3 le4

Typical Range

Fig. 1 shows a material response chart of several commonly used semiconductors and diamond

NASA/CP--2001-210948 75 August 6, 2001
,o,TN_,,_,-_ _ ................. ,_

I- "_ q:!:!:!: :!;_

100/_ E c-_T_\ • ._

"_ \: _EVICE- :,_,

Pe_'er

P

0

W

E

t?

FW

G3As \ , D,_wend Ope,s_

&

o L, ............... 4 .... _- A -; '.,,_,,_

o _oc 2oo 3co ,_oo soc, _u_ 7o0

"£EMPER,_ TURE

Figure 2(a): Gate-propagation delay vs. diamond and related wide bandgap materials based devices,

(b_: Operation characteristics of Si. GaAs. & Diamond devices.

Proof of concept field effect transistors, metal semiconductor FET (MESFET), metal undoped diamondsemiconductor FET (MiSFET), and (MOSFET) have been reported in literature. Boron doped homo-epitaxial

diamond film grown on {100} type IIa natural diamond substrates show improved source to drain current and

transconductance at elevated temperatures.

Hybrid Microelectronics

At room temperature diamond conducts heat better than any other material. Furthermore, diamond is a good

insulating material and its thermal expansion coefficient is very close to that of the silicon and gallium arsenide.

The low dielectric constant of diamond is also extremely useful for semiconductor device applications because of

the reduced cross-talk. Therefore diamond has unrivaled properties for MEMs, heat sink, surface mounted devices,

and applications with large component density. The synthetic diamond heat sinks are made of pure high-pressure

synthesized crystals which contain controlled amount of nitrogen. The surfaces of diamond heat sinks are metalized

for brazing semiconductor chips and for mounting on copper or other metal bases. Such synthetic heat sinks are now

widely used for mounts of laser diodes or high power millimeter wave devices. Insulating or semiconducting wafers

of freestanding diamond can be used for monolithic integration of the discrete or integrated devices, millimeter wave

travelling wave amplifier, and backward wave oscillator. Hence, as a microelectronic lateral extension to

complexity, RAM on processors, system on chip, network on chip, and superchips are feasible.

Opto-electronics

Due to its chemical inertness, high mechanical strength, and high thermal conductivity, diamond can be used for

several unique applications. The ability to grow optical quality diamond on substrates has led investigations towards

the integrated opto-electronic devices, e.g. diamond window for high energy free electron laser (FEL). The use of

diamond in picosecond high-voltage electro-optic switches has recently been demonstrated (Venables, 1990).

Presence of optical centers also makes diamond a suitable candidate for light emission devices. The other related

applications include laser and X-ray windows, diamond windows for synchrotron radiation, monitoring windows for

chemical processes employing toxic and corrosive ambient. The application of diamond for photovoltaic devices and

photonic switches is a nascent area of research. Due to its large optical bandgap and spectral responsivity, it is

recommended for the first time that tandem junction photovoltaic devices with diamond as constituent layer will

produce high open circuit voltage, high responsivity, and broad spectral response. Such devices find applications

ranging from hand held calculators to residential lighting.

Other Related Applications

The velocity of sound propagating in a materials typically employed as an index for the comparison of

diaphragm materials, i.e. c - _]E/p where c is the velocity of the sound, E is the elastic modulus, and p is the density.

Hence a material with lower density and higher Young's modulus provides higher sound velocity. Diamond ranks

higher than any other substance. Hence diamond material could significantly improve the quality of sound and its

applications range from record cartridge, stylus, cantilever, and speaker diaphragm.

NASA/CP--2001-210948 76 August6,2001
Thermal applications includeheatsinkformutitudeofapplication. Afewrelevant examples include, the

heatsinkforImpactionizationavalanche transittime(IMPATT), heatsinksformicrowavediodes, multi-chip

modules(MCM),etc.Diamondallowsimprovedpackingdensityofchipsandmodulesovercurrent aluminaand

siliconbasedsubstrates. DiamondMCMshavebeenemployedformounting2microprocessor chips,staticrandom

accessmemory, andthedataandaddressbuffersonthesamemodule. ThediamondMCMoperates atlower

temperature by10-15C.Basedonthepreliminary calculations andobservationsfloweroperatingtemperature of

diamondbasedMCMs,it issuggestedhatasalateralextensiontomicroelectronics, RAMonprocessors, systemon

chips(SOC)andnetworkonchip,andsuperchips arepossiblewiththediamondbasedsubstrates. Topicssuchas

stress-strain accommodation inmultichipconfiguration andsystemtransitiontechnology implementation employing

intelligent processconcepts andcontrol mechanismtrategies areparticularlyimportant forlargescaleproductionf

freestanding diamondwafers.

CONCLUSION

Innovations in VAJLSI chip architecture, by integrating diamond films is a viable direction to produce high

performance electronic devices. Despite of several advances in the diamond film deposition, a number of problems

need to be resolved for its successful application in electronic devices. The primary technical challenges are:

increasing the deposition rate; increasing the growth rate; identifying a large area heteroepitxial or homo-epitaxial

substrate for single crystal film deposition; and reducing the defects. Although several challenges still exit, the

unique properties of diamond combined with the ability to control the film through CVD makes it a promising

technology for electrical, optical, and thermal application.

REFERENCES

1.

2.

3.

4.

5.

6.

7.

6.

7.

8.

9.

10.

11.

12.

13.

14.

15.

16.

17.

18.

19.

20.

Collins, A.T., Ceram. Int. 22, no. 4, 1996, 321.

Prins, J.F., Materials Science Reports, 7, 1992, 271.

Yoder, M.N., Materials Science Monographs, 1991,287.

Gilkes, K.W.R., et al, J. Applied Physics, 87, 2000, 7283.

Debroy, T., et al., J. of Applied Physics, 68, no.5, 1990, 2424.

Snail, K., and Hanssen, L., J. of Crystal Growth, 112, 1991, 651.

Snail, K., and Craigie, C.J., Applied Physics Letters, 58, no. 17_ 1991, 1875.

Matsui. Y., et al., Jap. J. of Applied Physics, 28, no. 9, 1989, 1718.

Cappelli, M.A., and Paul, P.H., ONR Contract # N00014-89-J-1197 (task # 431A022), Technical Report,

1989.

Cappelli, M.A., and Paul, P.H., J. of Applied Physics, 67, no. 5, 1990, 2596.

Yarbrough, W.A., Stewart, M.A., and Cooper, J.A., (private communication).

Ikeda, M., et all, J. Applied Physics, 82 no. 8, 1997, 4055.

Walter, M.D., and Menningen, K.L., 82, no. 4, 1997, 1900.

Eccles, A.J., et all, Thin Sold Films, 343, 1999, 627.

Okano, K., Diamond Related Mater., 3, 1994, 35.

Rozploch, F., et all, Electron Technol. (Warsaw), 29, 1996, 213.

Aleksov, A., et all, Solid State Electronics, 44, 2000, 369.

Vescan, A., et all, IEEE Electron Devices, 18, 1997, 556.

Vescan, A., et all, 55 th Annual Dev. Res. Conf., 1997, 40.

Holmes, J.S., et all, IEDM, 1994, 423.

Gluche, P., et all, 55 th Annual Dev. Res. Conf., 1997, 42.

Gluche, P., et all, IEEE Electron Dev. Letters, 18, 1997, 547.

NASA/CP--2001-210948 77 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ELECTRON AND HOLE CONTRIBUTION TO THE RESPONSE

OF CVD DIAMOND NUCLEAR DETECTORS

Marco Marinelli, E. Milani, A. Paoletti, G. Pucella, A. Tucciarone, A. Stelitano, G. Verona-Rinati

Dip. Scienze e Tecnologie Fisiche ed Energetiche University of Rome "Tor Vergata" (Italy)

Istituto Nazionale di Fisica Nucleare, Sez. Catania (Italy)

S. Albergo, V. Bellini, A. Musumarra, R. Potenza, C. Randieri, M.L. Sperduto, C.M. Sutera e A. Tricomi

Dipartimento di Fisica, Universitfi di Catania (Italy)

Istituto Nazionale di Fisica Nucleate, Sez. Catania (Italy)

EXTENDED ABSTRACT

High quality synthetic diamonds have been grown by Microwave Plasma Enhanced Chemical Vapor Deposition

on silicon substrates using a 1% CH4 in H2 gas mixture at a temperature of 750°C. The growth time has been

chosen of about 90 hours in order to obtain a film thickness of about 60 gm.

Diamond based particle prototype were then built from the above samples. Circular Au contacts about 7 mm 2

size and 100nm thick were deposited on diamond surface by thermal evaporation. A pre-irradiation with [3 particle

(pumping) has been performed before the measurements in order to enhance the detector performance.

The detectors were tested under the flux of carbon ions at the 15 MV Tandem accelerator of Southern National

Laboratory (LNS) - Italy. In fig 1 the sketch of the set up of the experiment is reported.

ete

Faraday _, _ target /

Diambnd Oe,eOors

pumps

L.N.S.

?_o [_andemaccelerator

Fig 1. Sketch of the experimental setup

Diamond detectors are mounted at 20 ° and 30 ° respectively from the beam, after a gold target 160 nm thick

used as scattering center. Diamond detectors outputs have been connected, trough a charge pre-amplifier and a

shaping amplifier with a shaping time of 2 Its shaping time, to multichannel analyzer. A silicon detector is also

mounted at 20 ° and the beam intensity is monitored during the acquisition trough the measure of the Faraday cup

current.

The energy of the incident particles has been varied from 20 to 90 MeV in order to change the penetration depth

of the radiation from about 8 [tm microns to more then the whole sample thickness.

The charge collection efficiency has been measured in both positive (positive bias on the irradiated side) and

negative polarization in order to selectively evaluate the electron and hole contributions to the detection mechanism.

In fig 2 the charge collection efficiency of the diamond detector under different ion energies is shown. A 120 V

bias voltage is applied.

NASA/CP--2001-210948 78 August 6, 2001
Atlowenergiesacompletelydifferent behavior canbeobservedforthetwobiaspolarizationconditions. In

particular forpositivebiasamuchigherefficiencyisobservedsuggesting thathemajorcontributiontothecharge

collectionprocessisduetoholeratherthenelectrontransport.

45 ,

4-0

35

30

o>,25

¢-

.o_ 20

0

'4-- 15 iii

10

5

0

n • nip , n , n , • , n , n • i , n ,

V

Ib Vbias = +120

• V bias = -120

9 •

o

n • n • n • n • • • n • n • n • n •

10 20 30 40 50 60 70 80 90 100

12C Energy (MeV)

Fig 2. Charge collection efficiency as a function of the incident particles

When the detectors are negatively biased (holes moving towards the irradiated side electrode) the charge

collection due to the holes is limited by the presence of the contact (blocking contacts). When that the penetration of

the impinging particles tends to zero the charge collected is then due to the electron contribution only. As a

consequence, by using the Hecht theory [1], from fig 2 a very poor average drift length of electrons of the order of 3

[tm can be extrapolated. By increasing the ion energy, and thus the penetration depth, a gradual increasing of the

efficiency is observed due to the increasing of the hole contribution to the signal.

Acompletely different behavior is shown for positive bias polarity. In this case for low penetration depth a

much higher efficiency is shown indicating a much higher drift length of holes. In this case, when the ions energy is

increased a lowering of the efficiency is observed due to the generated charge approaching to the other electrode

(opposite to the irradiation side) and to the highly defective layer of diamond close to the substrate side. At about 70

MeV, corresponding to the energy needed to cross all the sample, a slight increase is observed because of the

flattening of the ionization distribution all over the sample For the same reason, at 90 MeV the device became

symmetric with the bias polarization.

The above results clearly demonstrate that hole conduction plays a dominant role in the detection process. In

addition, these behaviors are consistent with a recently proposed model [2,3] in which intra-grain defects and grain

boundaries are considered as the limiting factors in the detection mechanism.

[1] K. Hecht, Z. Phys 77 (1932) 235

[2] Marco Marinelli, E. Milani, A. Paoletti, A. Tucciarone, G. Verona Rinati, M. Angelone, M. Pillon, Appl.

Phys. Lett. 75 (1999) 3216

[3] Marco Marinelli, E. Milani, A. Paoletti, A. Tucciarone, G. Verona Rinati, M. Angelone, M. Pillon, J. Appl.

Phys.. 89 (2001) 1430

Keywords: Synthetic diamond, CVD, Applications, Electrical properties

Corresponding author: Dr. Gianluca Verona-Rinati

Dip. S. T. F. E. - Universitfi di Roma "Tor Vergata"

Via di Tor Vergata 110, 1-00133 Rome ITALY

Tel: +39 06 7259 7232 Fax: +39 06 7259 7145

e-mail: gianluca.veronarinati@uniroma2.it

NASA/CP--2001-210948 79 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

TEMPERATURE DEPENDANCE OF TRAPPING-DETRAPPING EFFECTS IN CVD

DIAMOND PARTICLE DETECTORS

Marco Marinelli, E. Milani, A. Paoletti, A. Tucciarone, G. Verona Rinati

1NFM - Dipartimento di Scienze e Tecnologie Fisiche ed Energetiche,

Universit_ di Roma "Tor Vergata",Via di Tor Vergata, I 00133 Roma, Italy

M. Angelone, M. Pillon

Associazione EURATOM-ENEA sulla Fusione, Via E. Fermi 27, 1-00044 Frascati (Roma), Italy

EXTENDED ABSTRACT

The time evolution of the response of diamond particle detectors is related to the characteristics of defects

present in the detector, which rule carries dynamics because of trapping-detrapping effects. It can therefore give

valuable information on the number and depth of the defects which limit the carriers mean flee path, and the

performance of the detectors. In this paper the study of pulse-shape evolution is carried out using as a probe 5.5 MeV

a-particles impinging on detector grade CVD diamond films deposited on silicon substrates in a microwave tubular

reactor, suitably modified to improve film quality (Refs. 1 and 2). The gas mixture used is 1% cn4 in H2, resulting

in a growth rate of about 0.7 [tm/h at the deposition temperature of 750 °C. The films showed excellent crystal

quality, with very narrow diamond Raman peaks (FWHM about 2.4 cm -1) and extremely low photoluminescence

background (Ref. 2). Particle detectors were then realized by thermally evaporating a circular Au contact about

5 mm 2 in size and 100 nm thick as the top electrode, while Ag paste was used as the silicon backing contact. The

time evolution of the collected charge is measured through a digital oscilloscope, both when the detectors are in the

as-grown state and after pre-irradiation with 9°Sr 13-particles for a total dose of 3 krad, necessary to drive them in a

fully pumped state. In the following, we will report data measured on a 100 [tm thick. The average efficiency of this

film in the pumped state is about 40% under an electric field of 10 kV/cm, corresponding to an average CCD of

about 85 gm.

It is found that the pulse amplitude and shape depend on the field polarity and is dramatically affected by

pumping. In the as-grown state similar amplitudes are observed for positive and negative polarity, and the saturation

value is immediately reached, the collected charge being constant after the 10 ns rise time of the digital oscilloscope.

This implies that once trapped by a defect, electrons and holes are not thermally detrapped, so that the defects

limiting the detector's response in the as-grown state are identified as deep ones. The density of these trapping

centers is such that the CCD is about 12 [tm (Ref. 3). No significant difference could be observed between positive

and negative polarity, even if electrons and holes should have different mean flee paths, because CCD (and therefore

both electron and hole mean flee paths) is lower than the penetration depth G=15 [tm of 5.5 MeV a-particles in

diamond, so that the growth surface boundary plays a very limited role while, being the film thickness much higher

than the CCD the substrate interface boundary plays no role at all. In principle only one kind of traps for electrons

and holes is necessary to explain the behavior in the as-grown state.

In the pumped state no significant change occurs in the case of negative polarity with respect to the unpumped

state. For positive field polarity, however, the amplitude of the pulses is greatly enhanced, reflecting the large

increase in efficiency due to pumping, and breaking the symmetry between positive and negative polarity. It seems

that pumping leads to a much higher mean free path for holes than in the as grown state, while the electron mean flee

path is not strongly affected. Thus, in the case of positive polarity holes are now substantially free to move along the

100 [tm film thickness towards the substrate interface, while for negative polarity they are collected at the upper

electrode after having traveled at most a distance G= 15 [tm. Also, a significant slow component develops, meaning

that another kind of defect becomes important for positive polarity after pumping, namely a relatively shallow trap

for holes, allowing detrapping.

The simplest picture compatible with all experimental facts is as follows. In the pumped state the mean flee

path of holes is greatly enhanced, because pumping saturates most of the deep defects responsible for hole trapping

without detrapping. Thus, shallower defects existing in a lower concentration than saturated, deeper ones, now

become important. Since they allow detrapping, they do not limit the overall amplitude of the pulse, but slow down

the process of charge collection. For negative polarity, the limited changes in the pulse amplitude and shape with

NASA/CP--2001-210948 80 August 6, 2001
respect to theas-grownstateshowthatnosignificant saturationofelectrontrapsoccurs.Holesarein principle

substantially freetomove,butsince5.5MeVa-particlesionizeonlywithintheirpenetration depthG= 15 gm from

the detectors' surface (much less than the hole mean flee path), holes can move only a few microns before being

collected at the upper electrode. They therefore cannot give a great contribution to charge collection, nor can be

trapped and detrapped since their mean free path is now much higher than the real path to the electrode. Therefore the

slow component substantially does not show up, and at the same time the fast one does not increase significantly

with respect to the as-grown case. In conclusion, qualitative analysis of pulse-shapes show that both deep and

shallow defects exist for holes, while only deep ones are present for electrons.

45

40

35

30

,--., 25

20

lo

5

o

-200 0 200 400 BOO 800

t(ns)

Fig.l: Comparison of measured pulses with the simulation for positive polarity for both the as-grown

(lower curve) and pumped states (upper curve).

On these grounds we developed a computer simulation to describe trapping-detrapping processes in the

framework of the trapping-detrapping model originally applied to Si-based detectors (Refs. 4 and 5). The simulated

pulse-shapes agree very well with experiment (Fig. 1) with very reasonable values of the physical parameters

involved, making this technique helpful for studying and identifying defects which are responsible for limitation of

the efficiency of CVD diamond particle detectors.

To better clarify this phenomenon, the changes in the pulse shapes are analyzed when the working

temperature of the detector is varied from 40 °C to 120 °C. A systematic speed-up of the response is found with

increasing temperature, confirming that the slow component is due to thermally activated detrapping from relatively

shallow defects. Since the detrapping time ZDis connected to the activation energy of the defects through the formula

1/z p= s exp(-ED/kT)

where s is the attempt frequency, plotting ln(ZD) VS. T allows to determine the activation energy of shallow

(detrapping) defects. In our case this activation energy is about 0.3 eV.

Keywords: defects, detectors, pumping, CVD

REFERENCES

1) M. Marinelli et al., Appl. Phys. Lett. 75, 3216 (1999)

2) M. Marinelli et al., Proc. of the ADC/FCT Conference, Tsukuba, Japan, August 31-September 3, (1999) 154

3) M. Marinelli et al., J. Appl. Phys. 89, 1430 (2001)

4) K. Zanio, W. Akutagawa and J.W. Mayer, Appl. Phys. Lett. 11, 5 (1967)

5) M. Martini and T.A. McMath, NIM 79, 259(1970)

Corresponding author:

Enrico Milani

Universitfi di Roma "Tor Vergata", Fac. Ingegneria, Dip. STFE

Via di Tor Vergata, 1-00133 Roma, Italy

Fax: +39-06-72597145; Tel: +39-06-72597228; e-Mail: milani@ing.uniroma2.it

NASA/CP--2001-210948 81 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ON THE FABRICATION AND BEHAVIOR OF

DIAMOND MICROELECTROMECHANICAL SENSORS (DMEMS)

K. Holmes, J.L. Davidson, W.P. Kang, M. Howell

Dept. of Electrical and Computer Engineering

Vanderbilt University School of Engineering, Box 99-B, Nashville, TN 37235

ABSTRACT

CVD (chemically vapor deposited) diamond films can be processed similar to "conventional"

semiconductor device fabrication and as such can be used to achieve microelectromechanical structures

(MEMS) also similar to, for example, silicon technology. Very small cantilever beams, membranes,

stripes, tips, etc. can be constructed in doped and undoped diamond films and offer an array of choices in

diamond with its known superior properties such as elastic modulus, high temperature semiconduction,

high thermal conductivity, very low coefficient of expansion and numerous other diamond parameters.

This paper will review the construction and behavior of the second generation DMEMS devices comprised

as an accelerometer with a diamond diaphragm for use in very high G applications and a diamond pressure

sensor for very high temperature and frequency response.

Keywords: diamond, MEMS, sensors

INTRODUCTION

DMEMS derive response from the piezoresistive (PZR) property of doped diamond [ref. 1 to 3].

There are two main components to a DMEMS: a structural member (membrane) and a set ofpiezoresistors.

For these devices, both components are fabricated of diamond. The design principle achieves a membrane

of intrinsic (undoped) diamond which will be strained, for example by pressure or acceleration and thereby

strain the PZRs patterned on the membrane. The operation of these type DMEMS is based on the principle

that there is a thin diamond diaphragm (or membrane) on which doped diamond resistors are made. When

the diaphragm is flexed (pressure or acceleration), the resistors undergo a change in resistance in response

to the strain resulting from the applied stress. The change in resistance from an electric bridge

configuration provides a corresponding voltage change which is modeled and calibrated to measure

pressure or acceleration.

DESIGN

The intended applications for the DPS (diamond pressure sensor) and the DAS (diamond

acceleration sensor) dictate different design considerations as summarized in Table 1.0. Additionally, state

of the art manufacturing techniques impose conditions on the size of sensor components to achieve certain

functions. For instance, high sensitivity for the DPS at the design pressure is at odds with the response

frequency restraint, placing a burden on priorities for the sensor designer. The principle difference between

the pressure sensor and the accelerometer is the intended natural frequency. The DPS is designed for ultrahigh frequency response and high temperature environments, that is, the pressure sensor application

requires a high natural frequency which strongly influences its membrane dimension requirements.

Additionally the operating temperature requirement for the pressure sensor is substantially higher than the

accelerometer, requiring the starting resistance to be higher. The DAS is designed to provide high

frequency information, have a wide dynamic range, high resolution, and demonstrate increased

survivability. The lower natural frequency for the accelerometer allows larger dimensions.

The design of the DPS/DAS requires the definition of three primary response variables:

frequency, deformation, and stress. Frequency is independent of deformation and stress, while deformation

and stress are related. Pressure and acceleration produce deformation which results in stress. The ultimate

capability, and therefore maximum sensitivity of the sensors (DPS and DAS), is dependent on the failure

stress levels for the diamond diaphragm. The primary variables involved in assessment of response for

design considerations are provided in Table 2.0.

NASA/CP--2001-210948 82 August 6, 2001
DPS DAS

Input Signal Magnitude < 0.2 psi < 100kG's

Input Signal Max. 70 kHz 15 kHz

Frequency

Operating Temperature < 680°C 27°C

Sensor Response 210kHz 45kHz

Frequency

Sensor Geometry

- Thickness

(Manufacturing Limit)

- Diaphragm Diameter

(Functional Limit)

- Sensor Diameter

(Functional Limit)

>l.0_t

<1.5 mm

<2.5 mm

>l.0_t

N.A.

N.A.

Table 1.0 Design Information

Response Quantity Primary Variables

Frequency

Deformation

Mass Density of Diaphragm

Proof Mass

Modulus of Elasticity

Thickness of Diaphragm

Diameter of Diaphragm

Force (Pressure/Acceleration)

Proof Mass

Modulus of Elasticity

Thickness of Diaphragm

Diameter of Diaphragm

Stress Deformation

Failure Stress

Table 2.0 DMEMS Response Variable Considerations

NASA/CP--2001-210948 83 August 6, 2001
Initial failure stress data was obtained from pressure burst tests on eight diaphragm test samples

which had resonant frequencies in the 1-5 kHz range. These samples exhibited a significant amount of

geometric non-linear response in analytical evaluations. Using linear analysis techniques, the diamond

failure stress was predicted from the rupture pressure to be in the 1-5 million psi range. However, using

non-linear methods on the same data, the average failure stress calculated for eight samples was 1.23x105

psi. Five additional samples, with significant empirical non-linear pressure response behavior due to slight

modifications in the fabrication process, produced an average failure stress of 1.41x105 psi. Other

published data [ref. 4], for twelve samples with diamond diaphragms possessing linear behavior under burst

pressures, indicated a failure stress of 1.45x105 psi. Natural diamond was reported to have a measured

tensile strength of 0.5x106 psi and a theoretical tensile strength of 16x106 psi [ref. 5]. Documented tests [5]

indicate tensile strength values as high as 3xl 06 psi, but these may be erroneous due to non-linear behavior

that was not included. Additional studies to provide failure stress data for parametric evaluations of

fabrication and material process variables to characterize and optimize diamond failure stress

configurations are the subject of a future paper. Presentations of the calculated failure stress for several

geometries can be seen in Figure 1. The stresses are shown as a function of (thickness/radius) 2, the linear

stress variation parameter for plates in bending.

T able 1.0 indicate s intended re sp onse frequency o f 210kHz for the DP S and 45 kHz for the DAS.

The range of thickness/radius combinations as determined from linear procedures that can provide these

frequencies is indicated in Figure 2. The presented frequencies were determined from linear procedures.

Diaphragms responding in a non-linear fashion would have an instantaneous frequency higher than the

linear frequencies presented. Required sensor size and fabrication capabilities restrict the acceptability

limits of both sensor diaphragm thickness and radius.

Proof masses at the center of a diaphragm could be added to enhance the output signal of

accelerometer components but are not considered here. The centerline deformation of the diaphragm in a

DPS or DAS is a function of the applied force which may be generated from either pressure or acceleration.

Figure 3 plots the relative equivalency of a pressure and an acceleration for a reference diaphragm. The

transient deformation at 100,000 g's is the same as that for a transient pressure of approximately 11 psi for

this case. A parametric series of calculations provided an accelerometer design which required no proof

mass to meet the stress and frequency goals for the DAS. The selected design has a thickness of 23.0

microns and a radius of 2.0mm.

The distribution of strains in the radial and tangential direction for a DPS/DAS configuration is

presented in Figure 4. The normalized strains in this figure are for linear bending only. Nonlinear

membrane actions tend to increase the strains at the center of the plate in proportion to the strains at the

outer boundary. When high frequency responses are required, the nonlinear response associated with

membrane behavior all but disappears so that the strain distributions of Figure 4 are very representative of

the anticipated strains for the specimen.

Using a design stress of 100,000 psi and a modulus of elasticity [ref. 6]of 165 x 106 psi, one

arrives at a design strain of aMax 100,000/165 X 106 0.64 X 10 -3 in/in. Therefore, for linear bending, the

strain at the boundary of the specimen is 0.64 millistrain. Since piezoresistive components will be used to

measure strains in the diaphragm specimens, the required length of the resistance elements plays a

significant role in determining the optimum location for measuring strains. Examining Figure 4 indicates

that if the required length of the elements is more than 20% of the radius, gages at the center of the

specimens may be more effective. If less than 20% of the radius is sufficient, a radial strain measuring

element near the boundary of the membrane is more sensitive.

LAYOUT: Applying the ground rules derived from this modeling, the layout and dimensions of

the diaphragm and resistors were determined. Both the DPS and DAS show the greatest change in strain

within 20% of the edge of the membrane for resistor elements orientated along the radii. Single sensor

description: insertion criteria for the pressure sensor require a die size of less than 2.5 mm in diameter. To

achieve appropriate sensor response frequencies the membrane diameter is 1.5mm. The dimensions for the

resistive elements are aimed at achieving a 1000 ohm bridge at 680°C. On and off membrane resistors are

matched in area with equivalent sheet resistance. Resistors with 3, 6, 12, and 18 squares were designed for

multiple sensor configurations. The individual sensor die uses a 3 square resistor configuration. Based

upon membrane edge proximity requirements, the maximum length of the active resistor is 150 microns.

The minimum width for each resistor is 50 microns. The 3 square configuration consists of two resistors

each 150 microns by 50 microns. The 50 microns critical dimension (width) was used for both the DPS

NASA/CP--2001-210948 84 August 6, 2001
andtheDASto standardize thedesign.Theaccelerometer didnotrequirethesizerestrictionsofthe

pressuresnsor andthuswillhavelargermembrane dimensions.

Figure5isthelayoutofatypicalDMEMSdevice. Numerousconfigurations werecreated. Note

thecomparison inFigure6tothefirstgenerationDMEMSdeviceandthesizereduction

RESULTS

FabricationfcompleteDPSandDASdevicesproceeded throughthewaferlevel.Thecritical

processes ofmembrane formation, PZRdelineation, metallizationdepositionadinterconnect paterning

wereexecuted. Themetallization utilizedwassequentially layeredtitanium/gold, ametalinterconnect

systemalsousedbyothers[ref.7]fordiamondcontact andhightemperature mtallization.

Figure7isanopticalmicroscope pictureof10undiceddevices, onaportionofoneofthewafers.

Thegrayhighlightedareaisenlarged, Figure8,showingtheresistorsandtheTi/Auinterconnect deail.

Thecurrent voltageresponseofexampleintegratedPZRresistorsisshowniFigure9.

ThebasicfabricationschemaoftheDMEMSandthemonolithicdiamondresistorPZRelement

hasbeendescribedpreviously[ref.8]. Thediamondresistorsondiamondsubstrates, achievedfrom

delineatingap-diamondlayeronanundopeddiamondlayersubstrateknowntohaveresistivity> 1013

ohm-cm(25°C),werelectricallycharacterized. Thismonolithic[resistor andinsulating/isolating substrate

arethesamematerialndlayeredsequentially] diamondcomponent isusefulforsensorandruggedresistor

applications. Usinga directetchprocessdescribedelsewhere, patterneddopediamondresistorswere

fabricatedonundoped, insulativediamond.Interconnect metallization, waslayeredTi/Au,sequentially

sputtered. Favorablylowcontact resistance tothep-diamondisachievedwith this configuration. Sensors

could be "diced" from the wafers, packaged and wire bonded, see Figure 10.

Essential to obtaining this resistor construct is achieving complete delineation and electrical

isolation of the body of the p-diamond resistor. Figure 11 shows the current-voltage behavior of a typical

monolithic resistor and the "open" between resistors on a common i-diamond substrate. This isolation has

o

been observed previously to be retained to temperatures over 300 C [ref. 9]. Figure 12 plots the currentvoltage behavior to high loads of a monolithic p-D on i-D resistor. The resistance at nominal load is 60

Kohm, resistivity 6 ohm-cm. Above 300V the current becomes non-linear, similar to the thermal runaway

behavior of the resistors on ceramic. This resistor failed at 340V carrying 12.8 mA, a power density

(surface area) of - 3 MW/cm 2.

"In situ" pressure application (room temperature) of individual "diced" devices under active

electrical probe indicated the PZR bridge was responding to the pressure differential, providing a dV(dP)

output. Typical results are shown in Figure 13 indicating excellent linearity over a wide range of strain.

Detailed gauge factor evaluation and temperature characteristics will be the subject of a subsequent paper.

SUMMARY

Processes and procedures to take CVD diamond processing into a sensor device design and

fabrication project for practical applications and conventional form fit and function configuration have been

described. It is clearly illustrated that existing technology presently available to the silicon industry can be

directly applied to diamond films with predictable and deployable results for advanced rugged diamond

devices with unprecedented performance envelopes.

ACKNOWLEDGEMENTS

Vanderbilt gratefully acknowledge the fiscal support of Physitron, Inc., NASA and the United

States Air Force and for defining the application and needs for advanced devices to pursue visionary

programs with system requirements that push the envelopes for broad ranging benefits from defense to

transportation. The support of certain fabrication steps by the Alabama Microelectronlcs Center at Auburn

University, Auburn, AL under the direction of C. Ellis is greatly appreciated.

REFERENCES

1. J. L. Davidson, D.R. Wur, W. P. Kang, and D. L. Kinser, Sensors, Vol.12, No.11, 32, (1995).

2. D. Wur, J. L. Davidson, W. P. Kang, and D.V. Kems, Proc. of lnt. App. Diamond Conf.'95, NIST, Maryland, USA, August (1995).

3. M. Deguchi, N. Hase, M. Kitabatake, H. Kotera, S. Shima, M. Kitagawa, Diamond and Related Materials, Vol. 6, 367, (1997).

4. M. P. D'Evelyn, K. Zgonc, Diamond and Related Materials, Vol. 6, 812, (1997).

5. D.M. Jassowski, ReportAL-TR-89-044, from Aerojet Techsystems to Air Force Astronautics Lab, November, (1989).

6. J. L. Davidson, R. Ramesham, C. Ellis, J. Electrochem. Soc., Vo1.137, No. 10, (1990).

7. M. Wemer, et al, Diamond and Related Materials', vol.5, p 723, 1996.

8. D. R. Wur. W., J. L. Davidson., W. P. Kang., D. L. Kinser., "Polycrystalline Diamond Pressure Sensor," IEEE._

Microelectromechanical Systems, Vol. 4, No. 1, March, 1995

9. J. L. Davidson, D. R. Wur, W. P. Kang, D. L. Kinser, and D. V. Kerns, Polycrystalline Diamond Pressure Microsensor", Diamond

and Related Materials, 5, pp. 86-92, 1996.

NASA/CP--2001-210948 85 August 6, 2001
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FIGURES

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Figure 6. DMEMS new and original

device comparison (Original _ lcm membrane)

NASA/CP--2001-210948 86 August 6, 2001
Figure 7. Optical microscope picture (- 10X) of 10 undiced devices, from a portion of wafer

Figure 8. SEM enlargement of

area identified in Figure 7

showing the resistors and the

Ti/Au interconnect detail

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NASA/CP--2001-210948 87 August 6, 2001
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Voltage (V)

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common i-diamond substrate

NASA/CP--2001-210948 88 August 6, 2001
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400

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Figure 13. PZR bridge responding to pressure differential (25 C)

NASA/CP--2001-210948 89 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

DIAMOND MEMS - FROM PROOF-OF-CONCEPT TO APPLICATION

E. Kohn, M. Adamschik, P. Schmid

Dept. of Electron Devices and Circuits, University of Ulm,

89069 Ulm, Germany

Tel: +49-731-5026151

e-mail: dept-eb s@e-technik.uni-ulm.de

S. Ertl, A. Fliiter

Gesellschaft ffir Diamantprodukte (GFD),

Wilhelm Runge-Str. 11, 89081 Ulm, Germany

ABSTRACT

Although CVD diamond possesses many characteristics making it attractive for microsystems, it is still a rather

exceptional MEMS material and has not yet entered into main stream applications. Three examples are given, which

may illustrate how the advantages of this new material can be exploited more aggressively for applications in

microdevices.

INTRODUCTION

Diamond possesses many characteristics making it attractive for microsystems. Especially intriguing is the

combination of its exceptional mechanical, chemical and thermal properties with its wide bandgap semiconductor

characteristics. Although electronic active doping is still restricted to boron, producing p-type conduction only, it

allows to tailor the conductivity from insulating to metal-like. Therefore, diamond can be implemented as multifunctional material including actuator and sensor functions. However, these properties apply to monocrystalline

material, which is not available on the waferscale. Therefore deposition of CVD diamond films on foreign substrates

has been investigated intensively during the last two decades, with deposition on silicon resulting in "electronic

grade" films. Three routes have been followed: the deposition of polycrystalline films with random orientation (ref.

1), the deposition of highly oriented films on (100) oriented Si-substrates (ref. 2), and the deposition of films

composed of nano-sized crystals using a high secondary nucleation rate during growth (ref. 3). Techniques used

were MPCVD (refs. 2,4), HFCVD (ref. 3) and the combination of both (ref. 5). Whereas nano-diamond films are

highly conformal with a surface roughness in the nm scale, HOD films are textured with planar oriented islands,

which have partially grown together. Thus, the residual surface roughness allows high resolution lithography to be

employed on as-grown films (ref. 6). To obtain planar surfaces on randomly oriented films these films are frequently

polished.

The various routes in the materials development indicate already the attempt for an optimization for a specific

field of application:

- Randomly oriented films of high phase purity and quality have been used for heat sinks (ref. 7), optical and

microwave windows (ref. 8), and lenses (ref. 9). Other examples pointing already in the direction of microsystems

are DUV, X-ray and particle detectors (refs. 10,11), pressure sensors (ref. 12), gas sensors (ref. 13), flow sensors

(ref. 14) and SAW devices (ref. 15).

- Nanocrystalline films have been used for conformal coatings using molding techniques like in the case of STM tips

(ref. 16), field emitter (ref. 17), capillary wall systems (ref. 18) and electrochemical electrodes (ref. 19).

- Highly oriented films have first been introduced as base technology for Diamond MEMS at the European 1998

Diamond conference (ref. 2). The emphasis was on the demonstration of near ideal materials properties by these

films, the implementation into the device design rules and the development of generic fabrication steps. Initial test

structures were diamond membranes, patterned cantilevers and resistors with high temperature stable contacts.

TECHNOLOGICAL BASIS

The basis of the HOD process is a three step growth process with a nucleation step including bias enhanced

nucleation (refs. 20 to 22), a first outgrowth step to create a (100)-oriented top of the individual crystallites using _-

parameter engineering (ref. 23) and a second outgrowth step, resulting in the overgrowth of misoriented crystals and

NASA/CP--2001-210948 90 August 6, 2001
thecoalescence oftheindividualplatelets. Thegrowthprocedures havebeenoptimizedcontinuously andhave

resultedinthefollowingdataforfilmsofthicknesses rangingfromapprox.10to30gm:Fracturestrengthofupto

4.7GPa,Young'sModulusofapprox.850GPa,surfaceroughness smallerthan200nmrmsandanopticalbandgap

of5.4eV.Stresshasbeencontrolledandbalancedbyspecificdesignoftheprocessflow(ref.24)andcanactually

beusedinthedesignofcurvedcantilever structures.

As-grownfilmsarehighlyresistive, butcanbedopedwithNorB.Boronacceptor dopingisusedforcontacts,

piezoresistor andthermistor activechannels; nitrogen, a 1.7eVdeepelectricallynotactivateddonor,isusedto

compensate residual acceptor doping, tomovetheFermi-level towardsmidgapandto isolateconductivesurface

channelsfrombufferlayerleakage. Surfaceterminationwithoxygenorhydrogenisusedforelectrochemical

electrodes, whichmaybecomeanimportant aspect infuturebio-chemical microsystems (ref.25).

Forpatterningselectiveepitaxyordryetchingprocedures havebeendevelopedandareusedinthevarious

technologies (refs.2,26).Diamondmembranes onSiareeasilyrealizedusingstandardSi-etchingprocesses, which

areselectivetotheSi-diamondinterfaceandarethebasisfordiamondbulkmachinedMEMSdeviceslikepressure

sensors.

Extendingthetechnology tosurfacemachineddevicesseveraldditional stepsareneeded, suchasasacrificial

layertechnologyforfreestandingstructures likebridges. Commonlyusedassacrificial layer(servingatthesame

timeasnucleationlayerforatopdiamondoverlayer) iswetetchedSiO2(ref.26).

TECHNOLOGY-APPLICATION MATRIX

The matrix of verified materials data and developed fabrication steps is the design basis to project the

performance of diamond microdevices. Using HOD films the new technology competes with other

microtechnologies on Si using metal alloys, SiC or dielectric surface films. The implementation of diamond is

advantageous in two areas. First, the outstanding mechanical and thermal materials properties allow design for

heavy duty and extreme conditions. Secondly, the multi-functional property allows an essential reduction in

technological complexity. Especially the second feature allows to design microstructures with low complexity

avoiding heterogeneous layer stacks. This enhances reliability under high mechanical and thermal static and

dynamic stresses.

Aiming at the development of a MEMS component, besides these technological considerations, it is important

to identify the perspective for the future product. Where can diamond be technologically advantageous and enabling

and where can the new product be placed ? Four scenarios are common as sketched in fig.1.

Existing Technology

Existing Market

New Technology

Existing Market

Existing Technology

New Market

New Technology

New Market

Figure 1. Technology-Market Matrix

In the case considered here a new technology has been developed and the two lower options apply:

(1), the new technology allows to improve performance in an existing market, and

(2), there is a new technology for a new market, which both still need to be developed.

Clearly the first configuration offers the possibility of fast acceptance by a broad community, but may result

only in a short term advantage, since the competing technologies also advance. In the second case the enabling

nature of the new technology is decisive and the new idea may be able to dominate with its own rules a just

emerging systems environment.

Reviewing the progress in diamond technology, it is obvious that it is moving from a materials technology

towards a device technology due to progress in synthesis and deposition of thin films on foreign substrates.

Demonstrators published therefore try to show the advantages of this new material applied to classical structures.

Therefore they fall into the first category and have to compete with traditional Si based MEMS technologies,

however lacking the perspective of integration with Si based electronics. Despite its semiconducting properties

NASA/CP--2001-210948 91 August6,2001
diamondisnotanelectronicsmaterial yetandapplicationsarestilllimitedtodiscretecomponents withlow

electronicfunctionality. Applications outsidethereachofSiarethereforeessential, butnoteasilyfound.

In thefollowingthreeexamplesarediscussed, whichhaveemergedfromfirstgenerationexperimental

structures inourlaboratory.

DIAMOND SURGICAL MICROTOOLS

The development of CVD diamond surgical microtools falls into the first category of the development of a new

product. Natural diamond blades for ophaltmology and neural surgery are well established products in a prime

market place. Additionally the advantages of a diamond cutting edge over the metal counterpart, namely sharpness

and durability, are well known to the surgeon. On the other hand, in the case of the blade cut out of a natural stone,

the shapes are restricted to certain facets, the size is limited by the handling during manual production and the cost is

determined by the individual cutting and polishing process. In case these restrictions can be removed by a CVD

diamond blade fabricated by a microtechnology as described above, but with equal or better performance as

compared to the traditional product, it may be readily accepted. Fig.2 a,b shows by comparison a traditional natural

diamond knife, where the round cutting edge is approximated by 10 facetted cuts and a CVD diamond blade with a

freely designed contour. Furthermore, the force free plasma process used for polishing allows to realize microblades

not possible with splinters as shown in fig.3, the world smallest diamond blade as listed in the 2001 Guinness Book

of Records (ref. 27). Clearly these examples show that a new technology should be able to penetrate an existing

market. However, the traditional product is a macroscopic device using natural stones. The step towards a

microdevice based on a microsystems technology is large. The capabilities of this new technology and the resulting

product advantages are not transparent to the user. Thus, the gap between R&D, which is costly, and the product,

which is accepted by the market, is also large. But the new technology may allow to enable a new class of devices.

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Figure 2. Diamond surgical blades. Left: hand polished blade of natural stone;

right: blade made of CVD diamond patterned by plasma polishing.

Figure 3. Smallest diamond scalpel made of CVD diamond.

NASA/CP--2001-210948 92 August 6, 2001
Besidesminiaturization, increasedcomplexityandfunctionalitysthesecondgoalforfutureexpansionfrom

theexistingtechnology. InthecaseoftheCVDdiamondbladethisistheintegrationof activeelementsand

integratedsensors. Usingthemicroheater technologybasedonborondopedresistorstructures, suchabladewith

integratedheaterelement andthermistor temperature controlelement isshowninfig.4.Duetothehighthermal

conductivityandlowheatcapacity, thetemperature distributionacrossthesurfaceofthebladeishighlyuniform

(seefig.5)andexactlyadjustablewithfastdynamicresponse. Nowtheknifecanalsobeusedinthesurgeryofsoft

tissuewithhighbloodvesselcontent,closingthebloodvesselsalreadyduringcuttingorto suppressudden

bleedingveryuniformly.

,_."

Figure4.DiamondCVD blade with integrated heater and temperature sensor.

The center structure is the temperature sensor, the surrounding geometry is the heater.

Figure 5. IR temperature profile over blade surface as shown in fig. 4.

The center blade temperature is 105°C. The blade is mounted into a holder attached on the left side.

This example shows, that embarking from a very conservative concept based on the mechanical materials

properties of CVD diamond, to replace an existing product in an established market, the new microtechnology adds

a new dimension in functionality. It is now enabling the development of a new field of microsurgery tools, which

may already be considered the first step towards higher integrated microsystems, which could not be realized before.

DIAMOND MICROSWITCH

This device may also serve as an example for the first category, where a new technology promises to extend an

established field. In the case of the microrelay, this represents a classical structure. Switches are an important part of

any electric and electronic system. Traditionally they have been electro-mechanical, but have been widely

substituted by solid state diodes and transistor control devices. However at microwave frequencies high power is

still generated by microwave tubes and needs to be controlled by mechanical high speed / high power switches,

which are mainly electrostatically driven. In a 50 _ waveguide environment high voltages and high RF currents

need to be switched with transients in the ps-regime. These extreme requirements can only be met with a complex

combination of high performance materials, centering around ceramics and refractory metals. This makes

miniaturization difficult. In contrast, diamond can combine both properties, being highly insulating and metal like

NASA/CP--2001-210948 93 August 6, 2001
dependingondoping.Fromthefirstgenerationofexperimental diamondrelaystructures, thefollowingadvantages

canbelisted:

allessential partscanberealizedbyCVDdiamondalone

thehighthermal conductivity ofthecantilever allowstotoleratehighthermal losses

thebeamwillnotlooselasticityunderhightemperature / high stress operation

during high current surge, the contact surface will not corrode and will not form an insulating (oxide) compound

the use of diamond results in high switching speed and a high resonance frequency due to its high Young's

Modulus and low mass density.

Based on these properties a microwave power relay structure in co-planar arrangement was designed and

fabricated as shown in fig.6. In this design the powered electrode is the center line with the two adjacent lateral

groundplanes to the right and left. Most microwave relay concepts are based on electrostatically switched capacitors

and are thus not DC-coupled and have high attenuation at MHz-frequencies. This new design based on CVDdiamond will allow DC coupling. In this co-planar arrangement (see fig.6) the center beam needs to be insulated

from the right and left side ground planes with no losses at microwave frequencies. Here, undoped diamond is ideal

as can also be seen from the use as microwave window (ref. 8). The two electrostatic switching capacitors are

located above the right and left ground planes and when activated, do not interfere with the microwave signal on the

centerline. In addition, the "open" capacitance needs to be small for high frequency isolation; therefore the

cantilever is bent upwards using a built-in stress profile.

,,....,....,,.... •.............*....._...,..,.-. ,,_....,..,...,.............:1"_._....,.............,,'..

%'.:e:o:o:°%%:o:o :.- %_ -°:.%%_°°o%%°.'," ° • aw °°°_l_,°," "m%a_

•• o, oO_Oo-°o.-.'. _,,._,_ ...,. _,,......... , ,_" ,,,,o ,..o_-, _.,

:':':'::N':'..N t ::_:::.:._.:.

:::::::::::::::::::::::::::::::::::::::::::::::: :i:_.:.Ni:i:i_..x.."

, ._.,,,. -,.1.,

N

•e,_.e,,,'.-._..,,.. ,',,...... g.'.'.',:_ ,....'...._'X.'.'.'.""-'-"-',

_=_.,,_:g_,_;-::..W_:.._.,.:...:..-:.:.:.:.._ X.:.:.:.:....o -'-g;+_.:.:.:-_ ",:::.

Figure 6. Diamond microwave switch in co-planar arrangement.

The RF groundplanes represent also the switching capacitor driving plates.

Fig.7 shows the simulation of the temperature distribution on the switch for a total power loss of 50W, which

represents a contact current density of 2"106 A/cm 2 across the contact. In this simulation is assumed that the device

will operate in vacuum at a maximum temperature of 600°C.

1mm

I I Temperature (°C)

,_ 20

,:.:.:.:.:.. 84

L.:.;.;.;.;:

150

_; 213

278

342

407

471

1_ 535

600

Figure 7. Simulated temperature distribution of microwave power switch (current density 2"10 6 A/cm _)

NASA/CP--2001-210948 94 August 6, 2001
DIAMOND INK JET

This application concerns a new emerging field in biochemical engineering, where new products are

developed, reaching from highly professional to highly commercial. Microsystems, commonly labeled as "biochip"

or "lab on a chip" play an important part. Here, many established and new technologies compete based on polymers,

glass and Si. Common are arrays of reaction compartments for separation, synthesis, amplification, analysis etc. and

pumps and capillary systems to distribute the biochemical liquids. Most technologies do not combine the chemical

reaction substrate with mechanical, thermal, optical or electronic functions on a single chip yet, but are based on

hybrid integration concepts. For an integrated system the liquid biochemical environment is a new medium and all

components need to be compatible with this medium. In this respect diamond is an ideal material, being chemically

stable and inert with a large window for water hydrolysis (ref. 28). It is therefore able to be used in passive and

active functions, while being in direct contact with the liquids.

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!iiiiiiiiiiiiiiiiiiiiiiiiiii 

Figure 8. Diamond inkjet array with 10 ejector elements and connector leads, and detailed top view of an

individual double inkjet element through the diamond nozzle plate.

Liquids are supplied to the inkjet cavities from the right and left through short capillaries. The gray circles

represent the ejection holes in the nozzle plate.

The first structure designed for use in this environment was the thermal inkjet using a diamond heater (ref.

29). Here again diamond is incorporated as multi-functional material resulting in reliable operation even under

heavy overload (ref. 30). Based on the experience with water, aggressive acids, basic solutions and solvents, a

chemical microreactor array with 2 inkjets in each reaction chamber was designed for the custom designed synthesis

of oligonucleotides, (which are a fraction of a DNA chain) (see fig.8). The diamond inkjets serve as pl-dispenser of

an acidic and basic solution needed in the chemically controlled synthesis reaction in a small reaction compartment

(ref. 31). Fig.9 shows such an inkjet array mounted into a dispenser unit.

4:: :::.. ..: "'.

.........

$i:i:i:i:i:

::::::::::::::::::::::::::::::::::::::::::::::::::::::::: iiiii 

N::':':'::::::::.:.:.:.:.z&:'_':':$_,,>.__:.:.:.:.:.: ::::f.:::::::::::

•:':':'.'.'.'.'.'.'.'.'_f:_:,%N" ".'.'.'.'.',.'.'.'.'.'.'." "

........:.::iiiiiiiiiiiiiii !!!iiiiiii":

"iiiiiiii:: i iiiiiiiiii i iiii i::ii::iiiiiiiiii"

Figure 9. Mounted module of inkjet array on a 2" diameter Si-wafer. The two liquids are supplied through

the tube connectors on the right and left side. The printed circuit board wiring connects each inkjet element

electrically. The inkjet array is located on the rectangular diamond chip in the center.

NASA/CP--2001-210948 95 August 6, 2001
Thisapplicationisagainanapplication, whichmayalsobeobtainedbyaconventional Si-inkjettechnology,

butneedsspecial adaptationinrespect totheliquidsused.Thediamondinkjetdispenser canoperateintheentire

waterbasedpH-rangefrompH-1topH-14withoutcorrosionandhasbeenoperatedinavarietyofsolventsaswell.

Therefore, adaptationwasstraight fotward. Inthefuture,systemsofmuchighercomplexityareneededcontaining

furtherchemical processingsteps,processcontrolandreadout. In addition, besidesthesynthesisofmoleculesa

describedabove,thedetectionandisolationofbio-molecules isofequalimportance. Againherediamondwillbe

advantageous whenusedastrulyinertelectrodematerial.

CONCLUSIONAND OUTLOOK

CVD diamond is moving from a materials technology towards a device and component technology. It is a wide

bandgap semiconductor material combining many aspects of ceramics (hardness, thermal stability) with those of

electronically doped semiconductors. Demonstrators published therefore try to show the advantages of this new

material applied to classical device structures. At present wafer scale diamond films for microsystems devices are

still polycrystalline and different routes are followed, namely deposition as nano-diamond or HOD films. Attempts

to realize large area monocrystalline substrates are underway in several laboratories (refs. 32,33) are however not

yet technically relevant. Each of the approaches will target different areas of application. In this contribution the

focus is on devices developed using HOD films on Si. Three examples are discussed, which may show the spectrum

ranging from basic development to product development.

Especially the third example illustrates, that new fields of application can be found and targeted.

Concerning systems integration, the combination with electronic circuits is essential and thus integration with

Si-electronics on the Si-substrate. This has not been demonstrated yet. However, recent developments in

optoelectronic integration and vertical wafer scale integration indicate that also hybrid wafer integration may

become a interesting solution, especially since it allows to optimize each wafer level in respect to its device needs.

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8. C.S.J. Pickles, T.D. Madgwick, R.S. Sussman and C.J.H. Wort, "Optical performance of chemically vapourdeposited diamond at infrared wavelengths", Diamond and Rel. Mat. 9 (2000) 916.

9. E. W6rner, C. Wild, W. Miiller-Sebert and P. Koidl, "CVD-Diamond Optical Lenses", 11 th European Conf. on

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Abstr. paper 15.2.21

10. P. Gluche, O. Kohn, M. Binder, M. Adamschik, W. Ebert, E. Rohrer, C. E. Nebel and E. Kohn, "High Speed

Diamond DUV-Detector", Proc. IEEE Conf. on Advanced Concepts in High Speed Semiconductor Devices

and Circuits, Ithaca, August 06-08, 1997, NY, USA, p. 314

11. A. Mainwood, "CVD diamond particle detectors", Diamond and Rel. Mat. 7 (1998)

12. M. Werner et al.; "Review on Diamond Based Piezoresistive Sensors"; ISIE'98, Proceedings 147-152

13. W.P. Kang, Y. Gurbuz, J.L. Davidson and D.V. Kern, "A new hydrogen sensor using a polycrystalline

diamond-based Schottky diode", J. Electrochem. Soc. 141 (1994) 2231.

NASA/CP--2001-210948 96 August 6, 2001
14. C.D.Ellis,D.A.Jaworske, R.RameshamandT.Roppel,"Polycrystalline diamondfilmflowsensor", IEEE

Solid-State Sensor andActuatorWorkshop1990,Techn. Digest130.

15. KobeSteelUSA,Inc.,ResearchTrianglePark,NC,USA,"Diamondforelectronics: futureprospectsof

diamondSAWdevices", MRSBull.,Vol.23,No.9,1998,p.49

16. T.Shibata, Y.Kitamoto, K.Unno,E.Makino,"Micromachining ofDiamondFilmforMEMSApplications",

Journal ofMicromechanical Systems, Vol.9,No.1,2000,p.47

17. V.D.Frolov,A.V.Karabutov, S.M.Pimenov, V.I.Konov,"ElectronicProperties oftheEmissionSitesofLowFieldEmittingDiamondFilms",DiamondandRel.Mat.,Vol.9,2000,p.1196-1201

18. P.Rangsten, H.Bj6rkmann, K.Hjort,"MicrofluidicComponents iDiamond", Transducers'99, Proceedings,

p.190-193

19. T.N.Rao,A.Fujishima, "Recent AdvancesinElectrochemistry ofDiamond", DiamondandRel.Mat.,Vol.9,

2000,p.384-389

20. S.Yugo,T.Kanai,T.Kimura,T.Muto,"Generationofdiamondnucleibyelectricfieldinplasmachemical

vapordeposition", Appl.Phys.Lett.58,10(1991),1036-1038

21. S.D.Wolter,T.H.Borst,A.Vescan, E.Kohn,,,Thenucleationofhighlyorienteddiamondonsiliconviaan

alternatingcurrent substratebias",Appl.Phys.Lett.68(1996),3558-3560

22. A. F16ter,H. Grittier,G. Schulz,R. Baumann,C. Lutz-Elsner,R. Zachai,P. Gluche,

M.Adamschik, A.Aleksov,O.Kohn,W.Ebert,E.Kohn,C.E.Nebel,E.Rohrer,"Propertiesofthinhighly

orienteddiamondfilmsandtheiruseinelectronicapplications", to bepublishedinDiamondFilmsand

Technology

23. C.Wild,P.Koidl,W.Miiller-Sebert, H.Walcher, R.Kohl,N.Herres,R.Locher,R.Samlenski, R. Brenn,

"Chemical vapourdepositionadcharacterization ofsmooth{100}-faceted diamondfilms",DiamondandRel.

Mat.2,2-4(1993),158-168

24. E.Kohn,M.Adamschik, P.Schmid,S.Ertl,A.F16ter,"DiamondElectro-Mechanical MicroDevices

Technology andPerformance", Int.Conf.onNewDiamondScienceandTechnology 7(2000),tobepublished

inDiamondandRel.Mat.

25. A.Denisenko, A.Aleksov, E.Kohn,"phSensor BasedonSurfaceBoronDopedDiamond", Diamond2000,to

bepublishedinDiamondandRel.Mat.

26. M.Adamschik etal.;"Electrostatic DamondMicroSwitch",Transducers'99, Proceedings, p.1284-1287

27. GuinessWorldRecords2001,GuinessVerlagGmbH,Hamburg(Germany)

28. J.C.Angus,H.B.Martin,U. Landau, Y. E.Evstefeeva, B.Miller,N.Vinokur,"ConductingDiamond

Electrodes: Applications inElectrochemistry", NewDiamondandFrontierCarbonTechnology, Vol.9,No.3

(1999),175-187

29. P.Gluche,R.Leuner, C.Rembe, S.ausderWiesche, E.P.Hofer,E.Kohn,"NovelThermal Microactuator

BasedonCVD-Diamond Films",Int.ElectronDevicesMeeting, SanFrancisco, Dec.6-9,1998,Technical

Digest, p.483

30. E.P.Hofer,C.Maier,C.Rembe, S.ausderWiesche, E.Kohn,M.Adamschik, "RealisticPerformance Testsof

a DiamondPrintheadby HighSpeedVisualization", IS&TrsNIP15:Int. Conf.onDigitalPrinting

Technologies, October 17-22,Orlando, Florida,Conf.Proc.,1999,p.66

31. M.Adamschik, P.Schmid,C.Maier,M.Hinz,S.Seliger,E.P.Hofer,E.Kohn,"Microdosage Controlled

MicroReactor BasedonCVD-Diamond", 4thInt.Conf.onMicroreaction Technology, IMRET4,March5-9,

2000,Atlanta, GA,Topical Conf.Proc., p.114

32. R.Linares, P.Doering; "Properties oflargesinglecrystal diamond"; DiamondandRel.Mat.8(1999),909-915

33. F.H6rmann, H.Roll,M.Schreck, B.Stritzker,"Epitaxial Ir layersonSrTiO3assubstratesfordiamond

nucleation: depositionof thefilmsandmodificationintheCVDenvironment", DiamondandRel.Mat.9

(2000), 256-261

NASA/CP--2001-210948 97 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

MICROMACHINING OF CVD DIAMOND FOR MEMS APPLICATIONS

Takayuki Shibata*

Electronics and Information Engineering,

Graduate School of Engineering, Hokkaido University

Kita 13 Nishi 8, Kita-ku, Sapporo 060-8628, Japan

ABSTRACT

The excellent mechanical, electrical, thermal, and chemical properties of diamond make it a candidate material

for microsensors and microstructures to be used in microelectromechanical systems (MEMS). There is, therefore, a

growing interest in utilizing chemical-vapor-deposited (CVD) diamond thin film as a material in MEMS, and a

variety of micromachining techniques for diamond thin film have been developed. This paper reviews these

micromachining techniques for the fabrication of diamond MEMS devices. These techniques include the patterning

of diamond thin film by selective deposition or dry etching, surface micromachining technique for the fabrication of

movable microstructures, a silicon mold technique for the fabrication of three-dimensional microstructures, a

bonding technique for assembly, and other techniques required, such as polishing and planarizing, heteroepitaxial or

textured growth, and low temperature growth, in order to fully exploit diamond MEMS. In addition, the application

of micromachining techniques to the construction of some chosen diamond MEMS devices is also illustrated; these

include a diamond microgripper driven by an electrostatic comb actuator, a semiconductive diamond tip for a

scanning tunneling microscope (STM) as well as a machining tool, and a diamond probe for an atomic force

microscope (AFM).

Keywords: CVD diamond, micromachining, MEMS, microstructures, microsensors

INTRODUCTION

There is currently a great deal of interest in the development of microelectromechanical systems (MEMS) where

miniaturized mechanical components are integrated with integrated circuits (IC) on the same silicon (Si) chip. Many

kinds of mechanical microsensors (such as pressure sensors and acceleration sensors) (ref. 1) and microfluidic

devices (such as micropumps and microvalves) (ref. 2) have been developed. These mainly use Si or its compounds,

such as SiO2 and Si3N4. This is because the micromachinlng technologies for the fabrication of such MEMS devices

are based on IC technology. However, for certain applications, the performance of such fabricated microstructures,

including microsensors and microactuators, is limited by the materials used. Therefore, new materials are required if

the range of applications for MEMS is to be extended.

The excellent electrical, thermal, and chemical properties of diamond make it an ideal material for microsensors,

particularly for applications under conditions characterized by high temperatures and in chemically and radioactively

harsh environments. Thermistors (refs. 3 5), pressure sensors (refs. 6 8), acceleration sensors (refs. 9, 10), flow

sensors (ref. 9), chemical gas sensors (ref. 11), and radiation detectors (ref. 12) are all examples of such diamond

microsensors. In addition, it is believed that the mechanical properties of diamond, such as its singular hardness, high

wear resistance and low friction coefficient, offer great potential for its use in movable microstructures such as

rotating micromotors or microgears, and microstructures that have to come into contact with other materials. There

is, therefore, a growing interest in utilizing CVD diamond thin film as a material in MEMS. A variety of

micromachinlng techniques for diamond thin film have been developed in order to realize the diamond

microstructures required for such applications (ref. 13).

This paper reviews a number of micromachinlng techniques including the patterning of diamond thin film,

sacrificial layer etching for the fabrication of movable microstructures to be released from a substrate, a molding

technique for the fabrication of three-dimensional microstructures, and a bonding technique for assembly. In addition,

other techniques necessary to fully realize diamond MEMS, such as polishing and planarizing, heteroepitaxial or

textured growth, and low temperature growth are described.

* Contact address. Tel: +81 11 706 6441; Fax: +81 11 707 6581;

E-mail address: shibata@i_ano.eng.hok_,dai.ac.jp

NASA/CP--2001-210948 98 August 6, 2001
PATTERNING OF DIAMOND THIN FILM

The patterning of diamond thin film is of great importance for structuring diamond MEMS devices. The

patterning techniques can be classified into selective deposition and etching.

Selective Deposition

Selective deposition has been made by using Ar ÷ ion beam etching (ref. 14) or reactive ion etching (RIE) (ref.

15) to decrease nucleation density on extensively damaged substrates. It has also been achieved using an amorphous

Si mask (ref. 15) or SiO2 mask (ref. 16)as an alternative nucleation-suppressing film. Diamond thermistors (refs. 3,

4), pressure sensors (ref. 8), acceleration sensors (ref. 9), and flow sensors (ref. 9), have all been fabricated by

selective deposition. Figure 1 shows a fine diamond pattern with 12 _tm lines and 3 _tm spaces formed on a Si

substrate to a thickness of about 2 _tm using the selective deposition technique (ref. 13). This result indicates that it is

possible to achieve free patterning of diamond film with a feature size of a few micrometers.

Bias-enhanced nucleation (BEN) technique was also found to be effective in increasing diamond nucleation

density on an undamaged substrate, resulting in a diamond nucleation density as high as 10 l° cm 2 on a Si substrate

(ref. 17). A diamond microgear with a diameter of 400 _tm and a thickness of 5 _tm (ref. 18) and a diamond

micromotor with a rotor diameter of 150 _tm (ref. 19) were fabricated by combining BEN with SiO2 mask patterning.

BEN has also been shown to be effective on SiO2 film (ref. 20) and cubic silicon carbide (g-SiC) film (ref. 21), on

which the nucleation densities were about 108 cm 2 and 10 l° 1011 cm 2, respectively. The selective deposition of

diamond film has also been successfully carried out on a Si substrate using silicon-rich nitride (SIN) film (ref. 22).

Under optimized BEN conditions, the diamond nucleation density on a Si substrate has been eliminated and a

diamond film selectively deposited on SiN film.

Other techniques for the patterning of diamond film have been developed using a photoresist mixed with fine

diamond powders 0.1 0.2 _tm in diameter (refs. 23, 24). In this approach, nucleation was selectivity improved by a

Cr mask coated onto a Si substrate. Using this technique, the fine patterning of diamond thermistors (ref. 5) and

diamond microgears (ref. 24) was achieved. Selective deposition has also been made by selectively seeding substrate

with ultrafine diamond powders 5 nm in average size (ref. 25). Both these techniques have the advantage of

preventing damage to the substrate surface.

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:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

:?:i:i:i:!:?:i:i:::i: ;i:!:i:!:!_:_/_.:,_!:!:):i::?i:!:i:!:;i:i:

:.::::::<.::+:.:.:.::::::::::::::::::::::::::::.:. :::

Figure 1. Diamond fine pattern. Figure 2. Diamond microgripper.

Etching

Several techniques have been reported for diamond etching including ion-beam-assisted etching (IBAE) (ref.

26), RIE (refs. 6, 7, 27 30), electron cyclotron resonance (ECR) etching (ref. 31), and microwave plasma etching

(refs. 32, 33). The patterning of diamond film has been achieved by RIE in oxygen (O2) plasma using SiO2 and Si3N 4

film as a mask (ref. 27). An etch rate of diamond film in the range of 10 to 60 nm/min was obtained. Pressure sensors

were fabricated by RIE, using A1 film (ref. 6) or SiO2 film (ref. 7) as the masking material. In general, during the RIE

in O2 plasma, columnar structures were created on the etched surface, while the addition of SF6 to O2 depressed the

formation of the columnar structures (ref. 28). RIE might be suitable for patterning large area and thin diamond films

in diamond MEMS applications. The advantage of the ECR approach is that much lower ion energies can be used

while retaining practical etch rate.

Laser ablation can also be used in the selective etching of diamond film by using an ArF excimer laser (refs. 34,

35). As a simpler technique, diamond film can be patterned through thermal etching in O2 ambience using a Si3N 4

mask. By this method, an etch rate of 4.7 nm/min was obtained at 700°C (ref. 23).

NASA/CP--2001-210948 99 August 6, 2001
SURFACEMICROMACHINING

MostMEMSdevicesrequirefreestanding microstructures sothat mechanical movement can be achieved. The

use of a sacrificial layer is a key technique in releasing movable micromechanical components from a substrate (ref.

36). This technique, where microstructures are built up, layer by layer, on the surface of a substrate, is the so called

surface micromachining.

Figure 2 shows SEM images of a fabricated diamond microgripper driven by an electrostatic force (ref. 37). It

was fabricated by patterning diamond film to a thickness of about 2 btm onto a sacrificial SiO2 layer by selective

deposition and releasing the movable parts by etching of sacrificial layer. The fabricated microgripper had an

electrostatic comb actuator whose teeth width and gap were 10 btm and 5 btm, respectively. It was successfully

released from the substrate by etching of sacrificial layer, although it was slightly bent due to internal stress induced

in the thin film during deposition.

In fabricating such movable structures, Si3N4, polysilicon, and metals with a high melting point such as Ta, W

and Mo would also be candidate materials for the sacrificial layer. However, for a film to be used as a sacrificial

layer, it is necessary for it to be easily patterned and dissolved with high selectivity against any material to be used as

a substrate or other films to be deposited. Using the sacrificial layer etching technique, microswitch devices (ref. 38)

and microstructures such as beams (ref. 39) and air bridges (ref. 10) have also been successfully fabricated.

A unique technique has been proposed for the fabrication of single-crystal diamond microcomponents (ref. 35).

In this process, a sacrificial layer is formed in a single-crystal diamond substrate by implanting 4 5 MeV C + or O +

ions at 196°C and direct patterning of the epitaxial diamond film is then made by ArF excimer laser ablation.

Diamond microgears 400 btm in diameter and 10 btm in thickness have been successfully fabricated by this process.

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!!!!!!i!iiiiiiiiiiiii!

iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii iiiiiiiii ii iiiiiiiiiiiii i i i! ! ii

Figure 3. Diamond tip. Figure 4. STM image of narrow

grooves on Si surface.

FABRICATION OF THREE-DIMENSIONAL MICROSTRUCTURES

Three-dimensional diamond microstructures were achieved by using a Si mold. Figure 3 shows an SEM

photograph of a pyramidal diamond microstructure fabricated using the molding technique (ref. 40). The tip radius

of the pyramid was less than 50 nm. Using this pyramidal diamond microstructure, we carried out the machining of

polished Si to form nanoscale grooves 100 nm deep and up to 1.7 btm in length by moving the pyramid horizontally,

and then scanning tunneling microscopy (STM) measurements of the machined structures in situ (ref. 40). The

fabricated pyramid was a boron-doped p-type semiconductive diamond with a resistivity of about 1x 103 f2cm. As

shown in Figure 4, an STM measurement using the same pyramid showed that a narrow groove 200 nm wide and 40

nm deep had been obtained using this scratching process. At the end point of the formed groove, we observed a

protuberance caused by plastic flow. This result indicated that the semiconductive diamond pyramid was effective

both as a machining tool and as an STM tip.

The applicability of such pyramidal diamond microstructures to form a field emitter array has also been studied

(ref. 41). A similar technique has been applied to the fabrication of diamond microstructures with V-shaped grooves

to be used as a heat sink for a diode laser (ref. 42). More complex diamond microstructures based on this Si molding

technique such as capillaries and fluidic cooling devices, have also been demonstrated (refs. 43, 44).

NASA/CP--2001-210948 100 August 6, 2001
BONDINGTECHNIQUE

Asmentionedpreviously, themoldingtechniquecanbeusedtoformthree-dimensional microstructures. Inorder

tofabricateMEMSdeviceswithmorecomplexfeatures, however, thedevelopment ofaccuratebondingtechniques

isnecessary. Bondingtechniques willalsobeimportant forthepackagingofdiamondMEMSdevices.

Wehavepreviouslyreportedthatdiamondfilmcouldbeanodicallybondedtoglass(Pyrex7740)usinganA1

filmasanintermediate layerandaTi filmasanadhesivelayer,eventhoughthesurfaceroughness ofthediamond

filmwasashighas200300nm(peaktovalley)(ref.45).Figure5showsSEMphotographs ofdiamondprobesfor

anAFMfabricatedbytheselectivedepositiontechniqueforproducinga diamondcantilever patternandtheSi

moldingtechniqueforproducingasharpdiamondtip(ref.45).Aglassbackingplatewassuccessfully bondedtothe

diamondbasetoaprecisionofbetterthan10btm by the anodic bonding technique to facilitate handling of the

diamond probe. Figure 6 shows an AFM image obtained with the fabricated diamond probe with a spring constant of

about 1 N/m (ref. 45). This image proves the ability of the fabricated diamond probe to measure AFM images.

The anodic bonding could become a key technology in the utilization of CVD diamond film as a MEMS

material. However, few major developments for the bonding of diamond film are currently under way. More research

needs to be carried out into bonding techniques for a diamond film if its potential as a material for MEMS is to be

exploited.

i ;i ! ::!?sliiiiiiaisiaiigiii::i::!isi a! :"

Figure 5. Diamond AFM probe. Figure 6. AFM image taken with

diamond probe.

OTHER TECHNIQUES

In order to fully realize diamond MEMS, other techniques such as polishing and planarizing, heteroepitaxial or

textured growth, and low temperature growth may be required.

Polishing and Planarizing

The surface roughness of diamond film poses a serious problem if the design specifications of a MEMS device

dimensions are of less than a few micrometers. Furthermore, diamond film with a smooth surface can be used in

optical applications. In addition to conventional mechanical polishing, thermochemical polishing using hot metallic

plates film, laser polishing, ion beam polishing (see, e.g., a review paper reference 46 and related references therein),

and gas cluster ion beam etching (ref. 47) have all been proposed. The planarizing technique based on RIE has also

been proposed (ref. 48). A bilayer made up of diamond and SiO2 was etched in SF6/O2 plasma, and the surface

roughness was reduced from 40 nm (rms) to 14 nm (rms).

Heteroepitaxial and Textured Growth

To date many efforts have been made to realize the heteroepitaxial growth of diamond on non-diamond

substrates for electronic applications. This also promises well for application in MEMS devices. This is because

heteroepitaxial or highly oriented diamond film offers a smoothly textured surface without any additional polishing

process. This would result in diamond MEMS devices of higher accuracy, smaller dimensions, and lower cost.

Furthermore, such highly oriented diamond film was found to have not only better electrical properties (ref. 49), but

also higher mechanical strength (a Young's modulus of approximately 850 GPa) (ref. 50), better piezoresistivity (ref.

51) and higher thermal conductivity (ref. 52) than randomly oriented polycrystalline diamond film. These excellent

properties of highly oriented diamond film will offer strong advantages for its use in MEMS.

NASA/CP--2001-210948 101 August 6, 2001
Theheteroepitaxial nucleationonSiwasgenerallyinducedbytheBENprocedureinmicrowaveplasmaCVD

(ref.53)andhotfilamentCVD(ref.54),sincethefirstreportofhighlyorienteddiamondfilmong-SiC(ref.55).In

additiontotheconventional BEN,todateavarietyofmodifiedprocesses havealsobeenreportedwhichaimto

improvethequalityoftheheteroepitaxial films(refs.5658).Althoughhighlyorienteddiamondfilmshavebeen

depositedon2in.Siwafers(ref.59),therestillneedstobefurtherdevelopment forrealizingindustrialpplications,

e.g.,morehomogeneous filmsovera largearea,highlyorientedfilmswithsmallthickness(lessthan5 _tm),

depositionfsuchfilmsinahotfilamentCVDprocess, andhighlyorientedfilmsonothermaterials.

LowTemperature Growth

Lowtemperaturegrowthof diamondfilms(LTGD)is alsoof importanceif therangeof applicationsfor

diamondMEMSisto beextended. Thistechniquecanrealizethegrowthof diamondonvariousubstrates or

depositedthinfilmswithlowmeltingpointssuchasglasses, plasticsandA1.Inthelastdecade, thetechniquefor

LTGDatlessthan500°ChasbeenrealizedusingavarietyofCVDprocesses andgassystems; see,e.g.,areview

paperreference60andrelatedreferences therein.InaCH4/H2systemoroxygen-containing gassystems, it seems

thatheCH4concentrations fordiamondformationcouldbeextendedtoashighas5%atlowtemperatures ina

microwaveplasmaCVD,whereasthemaximumconcentration ofCH4torealizethegrowthofdiamonddecreased

(tolessthan1%)withdecreasing temperature inahot-filament CVD.Theadditionof oxygentothegasmixture

suppressed thedeteriorationfcrystallinenatureof diamondfilms.Sincesuccessful growthof diamondatlow

temperatures (250750°C)withhalogenorhalogen-containing gasesin a thermalCVDmethod(withoutthe

productionofatomichydrogen) wasreported(ref.61),severalttemptsforLTGDhavealsobeenmadeusingsuch

gassystems(ref.62).However, therestillneedstobefurtherimprovement indepositionrateandquality.

SUMMARY

Thereisnodoubtthatheexcellent propertiesofdiamondfilm will offer significant advantages for its use in

MEMS. It is therefore to be expected that diamond will become a major material of MEMS devices in the future. In

this paper, a number of current and developing micromachining techniques for CVD diamond used in the fabrication

of diamond MEMS devices have been described. These include the patterning of diamond film, surface

micromachining technique for the fabrication of movable microstructures, a molding technique for the fabrication of

three-dimensional microstructures, a bonding technique for assembly, and other necessary techniques such as

polishing and planarizing, heteroepitaxial or textured growth, and low temperature growth. However, the use of these

diamond MEMS technologies has both advantages and disadvantages for the fabrication of diamond MEMS devices.

We therefore have to choose the most suitable technique for each device. Furthermore, in order to fully realize the

potential of diamond as a material in MEMS, further improvement of conventional micromachining and deposition

techniques of CVD diamond, and the development of novel diamond MEMS technologies will be required.

ACKNOWLEDGEMENTS

The author would like to express his appreciation of the helpful suggestions and discussions on this work offered

by Dr. E. Makino of Hokkaido University. The author would also like to thank Mr. K. Unno, Mr. Y. Kitamoto, Miss

H. Takahashi and Miss T. Nakatsuji for their contribution in realizing various diamond devices. This work was partly

supported by the Japanese Ministry of Education, Science, Sports and Culture, Grant-in-Aid for Scientific Research.

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NASA/CP--2001-210948 103 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

FABRICATION AND BEHAVIOR OF DIAMOND FIELD EMITTER TRIODE

UTILIZING SILICON-ON-INSULATOR TECHNOLOGY AND CVD DIAMOND

W. E Kang, A. Wisitsora-at, and J. L. Davidson

Dept. of Electrical and Computer Engineering, Vanderbilt Univ., Nashville, TN 37235

wkang@v-dse.vanderbilt.edu, phone: (615) 322-0952, fax: (615) 343-6614

D. V. Kerns

Olin College of Engineering, 1735 Great Plain Avenue, Needham, MA 02492-1245

ABSTRACT

This paper reports a new fabrication approach that utilizes silicon-on-insulator (SOI) micro-electromechanical-system (MEMS) processing technology and chemical-vapor-deposition (CVD) diamond to achieve

totally monolithic diamond field emitter devices. The approach allows the core structure of field emitter device

such as the gate, anode, field emitter geometry, emitter array configuration and the final physical location of the

diamond cathode to be designed on the SOI wafer where precise self-alignment of diamond field emitter structures

can be achieved on an integral substrate. Moreover, this technique allows the fabrication of high-density gated

diamond field emitter arrays with precise uniformity and planarity over a large area. A uniform array with millions

of gated diamond micro-emitters per cm 2 can be reproducibly achieved. The emission characteristics in triode

configuration indicate that there are the three distinct regions associated with transistor operation: the cutoff

region, the linear region, and the saturation region. The emission characteristics demonstrate excellent transistor

behavior for vacuum microelectronic applications.

Keywords: diamond, vacuum microelectronics

INTRODUCTION

Diamond has been experimentally demonstrated by many to be an efficient electron field emitter (refs. 1 to

7). In addition, the robustness, chemical inertness, and high thermal conductivity of diamond promise diamond

emitter devices with good emission stability and reliability. However, for diamond emitter device technology to be

applicable, many obstacles such as performance, yield, and reliability have yet to be overcome. Most of present

diamond emitter development examines non-uniformly diamond coated silicon tips, planar as-deposited diamond

film, or irregular ion etched conical structure. These fabrication methods are nominally simple, but produce nonuniform emitter microstructures that result in inconsistent emission behavior and poor stability.

Practical fabrication techniques for diamond field emitter devices are needed. A molding technique is an

advantageous method for fabrication of diamond field emitters with well-controlled and uniform micro-emitter

structures. We have developed micro-patterned diamond microtips on diamond films by molding technique (refs. 7

to 8) and achieved self-align gated diamond field emitter (refs. 9 to 10). However, there are limitations. Since a

uniform array of self-align gated diamond emitters over a large area is difficult to achieve while keeping gate

height and opening within specifications, we recently developed a new molding fabrication process utilizing

silicon-on-insulator (SOI) wafer as the base substrate for the fabrication of large uniform arrays. Emission

characteristics in triode configuration have been achieved.

DEVICE FABRICATION

The fabrication flow chart of the self-align gated diamond field emitter devices utilizing SOI based wafer

is shown in Fig. 1. The fabrication process begins with wafer bonding of two pieces of oxidized (100) silicon

wafers followed by etchback and electropolishing. The resulting SOI wafer is comprised of 2.2 gm-thick silicon

active layer, 1 gm-thick SiO2 layer, and 525 gm-thick silicon substrate. A 0.1 gm-thick SiO2 layer was then grown

on the wafer surfaces. The inverted pyramidal cavities were then formed on the silicon active layer (as gate layer)

by photolithographic patterning and anisotropic etching of silicon using KOH solution. The square patterns are

sized such that complete inverted pyramidal cavities are formed within the silicon active layer. Next, a silicon

dioxide layer was grown on the active silicon layer to form the gate dielectric, which concurrently produces a wellNASA/CP--2001-210948 104 August 6, 2001
sharpened apexontheinvertedpyramidal SiO2layer.Thegrowthofthegateoxideproceededuntilit touchedthe

embedded1gm-thickSiO2layer(astheetch-stoplayer).Diamondwasthendepositedonthemoldbyplasma

enhancedchemical vapordeposition(PECVD). ThePECVDfabricationparameters werecontrolledtoachievea

smallbutdeliberatesp2content inthediamondfilm.Next,thebacksideofthesiliconwasetchedawayandstopped

attheembeddedSiO2layer.Finally,theSiO2layerwasetchedandthesharpeneddiamondpyramidalpexes

exposed. TheremainingSiO2andsiliconformthedielectricspacer andthegate,respectively.

(a) SOIwaferbonding, etchback, andoxidation

, ,_ _S.iO2

, , "_"_ g]O 2

(b) Photolithographic patterning and Si anisotropic etching

(c) Dry thermal oxidation for gate oxide

i02

iO2

(d) PECVD diamond deposition

I'_------Diamond

]"o-_ Si02 '.:::::::::::::::::::::_:':':':':':':':':':':':I "_'-'-'-'- Si

:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:_:°_------ _S.iO2

(e) Back etching of silicon substrate

......................... ",4------ Si02

::::::::::::::::::::::::::::::::::::::::::::::::::: si

[ ['-'-'-J ['-'-'-J _ ] I SiO2

]'_''-Diamond

(f) SiO2 etching to expose diamond tips

G2

_Diamond

Fig. 1: The fabrication flow chart of the self-align gated diamond emitter utilizing silicon-on-insulator wafer.

RESULTS AND DISCUSSION

Figure 2 demonstrates a large array of uniform gated diamond emitters achieved by this technique. The

insert of Fig. 2 shows a high magnification of a gated pyramidal emitter with very sharp apex located below a selfaligned silicon gate. The sharpened pyramidal diamond microtip is surrounded by a silicon gate separated by the 1

gm SiO2 layer. The quality and composition of the diamond tips were characterized using Raman spectroscopy.

The Raman spectra of diamond tips show asp 2 shoulder and a sp 3 peak at 1332 cm q, which indicates there is sp 2

content in the diamond tip.

This new self-aligning, height limiting fabrication technique utilizing SOI substrates is intended to

improve the fabrication yield and uniformity of self-aligned gate arrays over larger area than other etching

techniques. The etch-stop by SiO2 layer can be easily achieved by conventional chemical etching, in contrast to

another technique, which requires a more complicated electrochemical etching of silicon p-n junction (Ref. 11).

Another advantage is that the gate opening and position can be controlled and optimized. Thus, this fabrication

NASA/CP--2001-210948 105 August 6, 2001
techniqueispractical forverylargeareamicroemitterfabrication. Moreover, completemonolithicdiamondfield

emitterstructureswithself-alignedgateandintegratedanodecanbebuiltusingtheinherent layersassociated with

theSOIbasedwafer.

,°°°%°°%%%°°°°°°°°°°°°%°°°°%°°%°°%°°°°%%%°°°°°°%°°%°°°°°°.°°°%%,

'°%°°°°°°%°°°°°o°o°o%°°°°%%°°%°°°°°°%°°°o%%°°°°°°%°°%%%°°°°°°°°%°° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° ° °, ! i {i !i{ii i i i iiiii !ii!i iii{ i i { ! 

i iii ii ! !i iii ii!! {i % iiiiii i iiiiiii!iiii iii !iii: iii iiiii ii iiiiiiiiiiiiiii iiiii iii 

{{i{{ i iiiii!! iii!!!{iiiiiiiii{ii !! iiii{iiiii!{{ii!{

i i i i!!i i i i i i i?! ! !!!ii !!! ! !:! i i i i i i ii i!ii ii iii iiiii iiiiiiiii iiii iiii

'_%_%_%_%%_%%_%%%_%_%_%_%_%%°_°_%%%%_°%%'°_'_°%_%°t_%%%_%%%°_%%%%°°°°%°*%°°°_%°d

iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii___

Fig. 2: Scanning electron micrograph of self-align gated diamond emitters over a large array. Insert shows a high

magnification of a single self-align gated diamond emitter structure

Emission testing was performed in vacuum, 10 -6 Torr. The self-align gated diamond emitter was operated

in triode configuration with an external anode placed 50 gm above the gate. The anode emission current vs. anode

voltage (Ia-Va plots) of a self-align gated diamond emitter normalized to a single tip for various gate voltages (Vg)

is shown in Fig. 3. The plots demonstrate the desired saturation behavior of a field emission transistor. The

saturation is seen for various gate voltages at anode voltage above - 200 V. The anode emission currents vs. gate

voltage of a self-align gated diamond emitter normalized to a single tip for various anode voltages is shown in Fig.

4. It shows a low gate turn-on voltage of 26 V and a high emission current of 1 BA at an applied voltage of 60 V at

anode voltage ranging from 150 to 400 V. The threshold current used to determine the turn-on voltage is 10 nA per

tip. The turn-on voltage is considerably lower than silicon field emitter triode with the same gate-cathode spacing

of 1 Bm, whose turn-on voltage is typically >80 V (Refs. 12 to 15). Furthermore, the self-align gated diamond

emitter has relatively stable emission current at a given anode voltage. The fluctuation of emission current is found

to be less than 10% for high emission current of 1 BA per tips. This fluctuation is considerably smaller than the

fluctuation of 50% at the emission current of approximately 1 BA from single silicon tip with the triode structure

(ref. 16). The tip sharpening and sp 2 content are believed to contribute to the low operating voltage and high

emission current of the diamond emitter. A detailed discussion of the effect of sp 2 content on the emission

characteristics of diamond tip has been presented (ref. 17).

The simplest way to estimate the amplification factor, B, from emission data is to estimate graphically

from the basic definition

= (dVa/dV_) at Ia= constant (1)

From Fig. 4, at a constant Ia of 1 gA, ifVa changes from 350 V to 400 V, Vg is required to change from

approximately 59 V to 59.85. Thus, g _ 50/0.85 _ 58. This amplification factor is quite high for field emission

transistor. This high amplification factor value indicates that the diamond field emission transistor has good

performance for small signal amplification applications.

NASA/CP--2001-210948 106 August 6, 2001
<

=k

i

1.4

1.2

0.8

0.6

0.4

0.2

. Vg=62 V

0 _

0 100 200 300 400 500

% (v)

Fig. 3: The anode emission current vs. anode voltage (Ia-Va), normalized to a single tip, of a self-align gated

diamond emitter triode

<

1.4

1.2

1

0.8

0.6

0.4

0.2

0

0 10 20 30 40 50 60 70

vg(v)

_Va=150

...... Va=250

....... Va=350

_Va=400

o. Va=400

•", Va=350

Va=250

,o Va=150

Fig.4: The anode emission current vs. gate voltage (Ia-Vg), normalized to a single tip, of a self-align gated diamond

emitter triode.

NASA/CP--2001-210948 107 August 6, 2001
CONCLUSION

Wehavedevelopedanoveltechniqueforfabricationoflargeuniformarraysofself-aligngatediamond

micro-emitters onSOIwafer.Thenewfabricationtechniqueisefficientandproduceslargearrayofultrasharp

self-aligngatediamondemitters. Thus,thisfabricationtechniqueispractical forverylarge-scalemicroemitter

fabrication. Thefabricateddevicesdemonstrate promisingfieldemissioncharacteristics withlowturn-onvoltage.

Thisexcellent fieldemissionperformance oftheself-aligndiamondemittertiodeisattributedtothespecifically

designedgatedultra-sharpened pyramidal diamondtip structure. Therefore, thisnovelfabricationtechniqueand

associated deviceshavepromisingpotential forverylarge-scale integratedvacuumicroelectronic applications.

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NASA/CP--2001-210948 108 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

Electronic properties of phosphorus-doped triode-type diamond field emission arrays

Chia-Fu Chen, Chia-Lun Tsai, Chien-Liang Lin

Department of Materials Science and Engineering, National Chiao Tung University

I001 Ta-Hsueh Road, Hsinchu 300, Taiwan

Corresponding Author: Chia-Fu Chen

e-mail: cfchen_cc.nctu.edu.tw

Fax: 886-3-5723811

Phone: 886-3-5731898

Abstract

Field emission display (FED) is evolving as a promising technique of fiat panel displays in the future.

Many factors affect the performance of field emitter arrays including the shape and work function of

emission materials, distance between tip and gate, and the environmental vacuum condition.

In this work, we present a novel scheme that involves a new fabrication process of gate structure metal-insulatorsemiconductor (MIS) diode using IC technology. Using a bias assisted microwave plasma chemical vapor deposition

(BAMPCVD) system synthesized P-doped or B-doped completes this process. Based on our experimental results, it

showed dendrite-like diamond with non-doped and nanotube-like diamond with B or P doping. Doping phosphorus

or boron can enhance its electric characteristic by reducing the turn-on voltage and enhancing the emission current

density. The turn-on voltage of non-doped, B-doped and P-doped is 15V, 8V and 5V, respectively. The field

emission current (Ia) of non-doped, B-doped and P-doped is 4gA (at 45V), 76gA (at 77V) and 322gA (at 120V).

Kcywords: Field emission, Diamond, CVD, Doping

NASA/CP--2001-210948 109 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

MICROFABRICATION OF VARIOUS ELECTRON EMITTERS

ON SINGLE CRYSTAL DIAMOND

Y. Nishibayashi, Y. Ando and K. Kobashi

FCT project/JFCC, Center for Advanced Research Projects, 6F, Osaka University,

2-1 Yamada-oka, Suita, Osaka 565-0871, Japan

K. Meguro and T. Imai

Itami R&D Laboratories, Sumitomo Electric Industries, Ltd., 1-1-1 Koya-kita, Itami, Hyogo 664-0016, Japan

T. Hirao and K. Oura

Faculty of Engineering, Osaka University, 2-1 Yamada-oka, Suita, Osaka 565-0871, Japan

ABSTRACT

We have developed a new process technology for emitter tips of single crystal diamond. The technology has a

process forming fine columns on single crystal diamond and a process reforming fine columns to sharp tips. Fine

columns were formed after diamond substrate with A1 dots mask was dry etched. Although it is generally difficult to

obtain both a high aspect ratio (height/diameter) column and smooth etched surface, we have fabricated tall fine

columns (aspect ratio of 8) using a novel reactive ion etching technique. It has been found that it is important to

mix a little CF4 gas in 02 gas. Using the high aspect ratio column, novel sharp tips of needle type have been

fabricated by anisotropic etching in microwave plasma at the optimized condition and novel pyramidal shape tips

have been fabricated by anisotropic homoepitaxial growth. It has been confirmed that a diameter of emitter top is less

than 40nm for a needle type tip and less than 80nm for a new pyramidal tip.

Keywords: homoepitaxial growth, anisotropic etching, reactive ion etching, sharp tips, emitter array

INTRODUCTION

Diamond is a useful material for highly efficient electron emitter devices because of its negative electron affinity,

so there are many reports on field emission from diamond. Sharp emitter tips of diamond have been fabricated on

aligned Si tips or by the transfer mold method in some paper (refs. 1 to 4), because it is important to controlled

emission site for electron devices and to make a threshold voltage lower. However, an emitter tip consists of

polycrystalline diamond and contains grain boundary in these papers. Therefore, it is difficult to investigate an

emission mechanism from polycrystalline diamond. On the other hand, we have fabricated a tip of single crystal

diamond. It has no grain boundary and an each tip has same crystal orientation. As an electron emission efficiency of

one crystal face is generally different from that of another crystal face, it is useful to select better crystal face or

orientation. In electron emission property, it is also important to transport an electron in diamond without a

collision. Therefore, an electron emitter of single crystal tip would be higher efficient than that of polycrystalline tip.

In this paper, we report a fabrication of new type emitter of single crystal diamond.

FABRICATION PROCESS AND EXPERIMENTAL

Sharp tips were fabricated on single crystal diamond after following two processes. First process is to form fine

columns on diamond substrate and second process is to reform fine columns to sharp tips. Formation method of fine

columns is as follows. After an aluminum (A1) of l#m thickness was evaporated on a single crystal Ib diamond by

sputtering method, A1 dots pattern of 1 to 15 #m diameter was formed by photolithography. A diamond surface was

etched by 3M0#m by a RIE using CF4/O2 (0-25%) gas, and an A1 was removed in a HF acid. Then, fine columns

NASA/CP--2001-210948 110 August 6, 2001
ondiamondsubstratewerefabricated. Reformation offinecolumnstosharptipswascarriedoutusingananisotropic

growthtechniqueorananisotropic etchingtechnique.

AnetchingofafinecolumnwascarriedoutinanNIRIM-typemicrowaveplasmareactor. H2andCO2gaswere

introducedtothereactor. Etchingconditionswereasfollows:Asubstratetemperature wasabout1100°C,aCOff

H2concentration 0- 5%,atotalgaspressure100torr,andanetchingtime0.1- 4hours.Asubstrates was(100)

singlecrystaldiamondwithfinecolumns. Fabricationoffinecolumnsonsubstratewasdescribedabove.

Anepitaxial growthonafinecolumnofsinglecrystaldiamondwascarriedoutinmicrowaveplasmareactorof

stainlesssteel.H2andCH4 gas were introduced to the reactor. Growth conditions were as follows. A CH4/H2

concentration was 0.25 to 9%. A substrate temperature 580 _ 1050 °C. A total gas pressure 50 _ 160 torr. A

microwave power 0.75 _ 1.4 kW. A growth time 20 _ 300 min. A substrate was (100) single crystal diamond with

fine columns. It have been known that growth rate ratio of V(100)/V(lll) is higher at a higher methane

concentration and at a lower substrate temperature, where V(100) and V(111) are growth rate of <100>- and <11 l>-

direction, respectively (ref. 5). Therefore, sharp pyramidal tips on various orientation substrates are able to fabricate

controlling the growth rate ratio.

A figure of column and tip on single crystal diamond was observed using a scanning electron microscope

(SEM) and a field emission SEM (FESEM). We also evaluated surface roughness by an optical step meter attached

on optical microscope and an atomic force microscopy (AFM).

RESULTS AND DISCUSSION

Formation of fine columns using novel dry etching technique

As well known, although diamond is etched easily and deeply in 02 plasma, the etched surface has generally

very large roughness. In order to obtain a smooth surface, it is effective to use CF4 gas (ref. 6). Figure 1 shows a

dependence of surface morphology by SEM on CF4 concentration in 02 gas. One dot mask of A1 has existed at a

center of each photograph, but the diameter is different. After dry etching, a column was formed on diamond surface

except for one photograph. As shown in this Fig. 1, the etched surface had a lot of needles at the condition of

CF4/O2=0 and 0.4%. On the other hand, an etched diamond surface had no needle and is smooth in the condition of

CF4/O2 >2%. However, A1 mask was etched and almost disappeared in the condition of CF4/O2 >5%, though a

diamond was not etched. Other conditions were examined in more detail. Thus, it was found that the best condition

of CF4/02 is 1 to 3%. At the best condition, we tried to fabricate a column of very small diameter and obtained a

fine column of a very high aspect ratio (height/diameter) and its array as shown in Fig. 2. The highest aspect ratio is

8 for a column of 1.2#m diameter, the height of the highest column is 22.5#m using A1 mask of l#m, etching rate

is 9.6 #m/h and surface roughness is less than lnm by measuring of AFM. It is found that our dry etching technique

has an excellent performance. We think that a little CF4 gas reacts very effectively in exceed 02 gas and that

diamond was vertically etched without side-etching, because an etching surface was covered by non-diamond film.

NASA/CP--2001-210948 111 August 6, 2001
Reformationffinecolumnstosharptipsbyananisotropicetching

AfinecolumnwasetchedinmicrowaveplasmausingCO2/H2gas.Althoughacolumndidnotchangedinonly

H2plasma,columnchangeddrasticallyinCO2/H2plasma. ThebestCO2/H2conditionwasapproximately 0.5%.

Atthiscondition, lowheightcolumnon(100)substratewasreformedto acubicshapeandthaton(111)substrate

wasreformedtoapyramidal shape. Therefore, weguessthatanetchingatthisconditionisananisotropicetchingof

V(100)>V(111), whereV(100)andV(111)areanetchingrateof<100>-and<11l>-direction. Figure3 showsthe

shapeofalowcolumn(aspect ratio=0.7)(a)beforeand(b)afteretchingandahighcolumn(aspect ratio=2.4)(c)

beforeand(d)afteretching. Theshapeof(b)isalmostcubic.However, topfaceisnotperfect squareplane.In the

caseofahighcolumn, thatisclearandacolumntopbecomesaharptip.It isseemedthattopofhighcolumnwas

etchedisotropically. It wasthoughtthatamethaneconcentration, a distributionofradicalincidentdirectionand

electricfieldaredifferent partially.Wehaveobtainednewtypesharptip fromhighaspectratiocolumn.Thetop

shapewasobservedbyFESEMandit wasconfirmedthatitsdiameter islessthan40nmasshowninFig.4.

Fig.3Ashapeoflowaspect ratio(0.7)column(a)beforeand

(b)afteretching, andhighaspect ratio(2.4)column

(c)beforeand(d)afteretchingon(100)substrate.

Fig.4Topshapeofneedletypesharptip

inFig.3(d).

iVaa'-OT;'ti:i i:i"i i:'i:i:i:i :: 7 a'T'iNNiiiiii 

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.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.. •. : i.:.;.:.:.:.:.:.:.:.:.:.:.:_...._.._..:._'_:°._.:_.

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::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::.:.:.:.:.'.:.:.X_..:.:.:.:.:.:.;.:.:.:.:.:.:.:.:.:.'-.:-:.:-:

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::: ::::::::::::::::::::::::::::::::::::::::::::::::::::::::

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

(a) (b)

......_......:...:........._._-._..._.........,..,,.......,.........

x "_.'......'._'.'.'.'..'.'.:.:.:.:.:.:.'.'.'.'.'.'.'.'.×.._e_.: _.'..._'_._"×.'.'.::.:.:.:...:...:.:.:.'-'.'.'.'.'.',..'.'.x_.'.

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::: x.'...'_...",'.'.'.'.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.'.'.'.'.'._'_.:

::::::::::::::::::::::::::::::::::::::

li'ii_!:_iiiii!iiE!!iiiiiiiiiiiiiiiiiiiiii_:i:_.'_:?!:_'_:_'i'_'i:_iiiiiii

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

Fig. 5 A shape of (a) low aspect ratio (0.4) column and

(b) high aspect ratio (1.0) column after growth

on (100) substrate. A column shape before etching

insert each photograph.

Fig. 6 Top shape of pyramidal sharp tip

in Fig. 5 (b).

Reformation of fine columns to sharp tips by an anisotropic growth

Diamond was grown on a fine column by microwave plasma CVD method using CH4 and H2 gas. Controlling

growth condition, a pyramidal tip has been fabricated for substrates of (100) and (110) (ref. 5). The shape of

NASA/CP--2001-210948 112 August 6, 2001
pyramidal tipon(100)substrateisshowninFig.5(a).Inthisfigure,smallpyramidal tipandfourslightslopes,

whichsurroundit,areobserved. Whengrowthtimeismuchlonger,slightslopesgrowupoversmallpyramidal

tip.Fig.5 (b) shows a pyramidal tip grown from a high aspect ratio column. The growth rate ratio (V(100)/V(111))

of (b) was higher than that of (a). For tip of (b), slight slopes do not exist. Even though growth condition of (b) is

same to that of (a), slight slopes can not easily grow up over sharp pyramidal tip because starting column is high. A

top shape of a new type pyramidal tip was also observed by FESEM. Top diameter was estimated to be less than

80 nm.

We have fabricated a needle type sharp tip array using both etching technique and high aspect ratio column, and

a new pyramidal type sharp tip array using both growth technique and high aspect ratio column as shown in Fig. 7.

;""'._'_'_]:::::::::::::::::::::: ....:.::_......:.

' "-lI'I'" "lllI_II llIillI." ..lllllli "l_'. _._::. .'')':" "g::::" ._:;':" :::::"

,..:.:.:i:. .j

:.._._ • .:..

(b)

Fig. 7 (a) an emitter array with needle type sharp tips and

(b) an emitter array with pyramidal sharp tips

CONCLUSIONS

We have developed the formation technique of a fine column on diamond and the reformation technique of it to

a sharp tip. In the reformation method, we can use the etching technique or the growth technique. The etching

technique is anisotropic etching, which reform a low aspect ratio column to cubic shape. A new needle type sharp

tip has been obtained using both high aspect ratio column and the etching technique. A new pyramidal type sharp

tip has been obtained using both high aspect ratio column and the growth technique. A needle type tip and a

pyramidal type tip have top diameter of less than 40 nm and 80 nm, respectively. These new type tip arrays also

have been fabricated on diamond substrate by our new technique.

ACKNOWLEDGEMENTS

This work was supported by the Frontier Carbon Technology (FCT) project, which was consigned to Japan

Fine Ceramics Center (JFCC) by New Energy and Industrial Technology Development Organization (NEDO),

REFERENCES

1. J. Liu, et al.: Electron emission from diamond coated silicon field emitters, Appl. Phys. Lett., 65 (22),

(1994) 2842.

2. E.I. Givargizov: Silicon tips with diamond particles on them: New field emitters?, J. Vac. Sci. Technol.

B 13 (2), (1995) 414.

3. K. Okano, et al.: Fabrication of a diamond field emitter array, Appl. Phys. Lett., 64, (1994) 2742.

4. W.P. Kang, et al.: Micropatterned polycrystalline diamond field emitter vacuum diode arrays, J. Vac. Sci.

Technol., B, 14 (3), (1996) 2068.

5. Y. Nishibayashi, et al.: Homoepitaxial growth on fine columns of single crystal diamond for a field emitter,

Diamond and Related Materials, 9, (2000) 290.

6. H. Shiomi: Reactive Ion Etching of Diamond in 02 and CF4 Plasma, and Fabrication of Porous Diamond

for Field Emitter Cathodes, Jpn. J. Appl. Phys. 36, (1997) 7745.

NASA/CP--2001-210948 113 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Field emission from graphite-diamond composites

L. Ley, J.B. Cui*, M. Stammler, and J. Ristein

Institute of Technical Physics, University of Erlangen

Erwin-Rommel-Strasse 1, D-91058 Erlangen, Germany

ABSTRACT

The field emission properties of diamond-graphite composites were investigated as a function of composition both

for oxidized and hydrogen covered diamond. The composites consist of mixtures of nano-crystalline diamond and

graphite particles. In this way their composition could be varied at will while the field enhancement factor of the

individual crystallites remained unchanged. The measurements prove that graphite is the phase responsible for low

threshold field emission. The apparent emission threshold is strongly influenced by the conductivity of the

composites. Hydrogenation has two beneficial effects. It provides a conducting path to the emission sites via the

hydrogen induced surface conductivity of diamond. It also lowers the effective emission threshold of graphite in

contact with diamond that exhibits negative electron affinity after hydrogenation. The latter effect was

experimentally verified by photoelectron yield spectroscopy.

Keywords: field emission, diamond, graphite, negative electron affinity, percolation

INTRODUCTION

Thin film diamond synthesized from the gas phase by chemical vapor deposition (CVD - diamond) is an

excellent cold electron emitter that yields current densities of lmA/cm 2 for applied fields of 10 to 100 V/gm (see Fig.

1) (refs. 1 to 5). Even lower fields of 2 - 5 V/gm are needed to draw the same current density from nano size

diamond particles (diamond powder) (ref. 6). The emission properties of CVD-diamond films are markedly better for

hydrogenated films compared to unhydrogenated CVD diamond (ref 7, 8) as illustrated in Fig. 2. and hydrogenation

is a necessary requirement to draw any current at all from diamond nano powder (ref. 6). The field required to extract

comparable current densities from a flat metal surface with a typical work function of 4eV is of the order of 2 x 103

V/gm if one assumes field assisted tunnelling (Fowler - Nordheim tunnelling) as the electron emission mechanism

(ref. 9). Thus, the diamond based materials must either have a vanishingly small emission barrier or the local field at

the emission sites has to be enhanced by factors of the order of 103 in order to account for the low emission fields.

E 10 -2

= 10 -4

10 -s

c

= 10 -s

O

/llfll

CVD diamond

_ano crystalline /////

,.diamond powder

decreasing film quality

I I I I I

3 10 30 100

Electric Field (V/gm)

Figure 1. Current-Field characteristics of

CVD diamond films of different quality in

terms of sp3/sp2 ratio (right hand side) and

of nano-crystalline diamond powder after

hydrogen plasma treatment (adapted from

ref. 6).

6xl 0 6

E

o

-<

4xl 06

$

o 2xl Os

O

0

I

i non-hydrogenated

hydrogenated ."

|

J

40 80 120 160

Field (V/#m)

Figure 2. Field emission characteristics of the same

CVD diamond film before and after plasma

hydrogenation.

NASA/CP--2001-210948 114 August 6, 2001
Forsometimeit hasbeenthoughtthattheuniquepropertyof thediamondsurfacetoexhibitnegative

electronaffinity(NEA)whenfullyorpartiallyhydrogenated wascrucialforthefavorablefieldemissionproperties

of diamondbasedmaterials. Theelectronaffinity)_ is the energy difference between the vacuum level and the

conduction band. It thus constitutes the barrier for electron emission from the conduction band. For the (111) surface

of diamond )_ varies from +0.35 eV for the clean (no hydrogen or other adsorbates) surface to -1.27 eV for the fully

hydrogenated surface (ref. 10) and similar values have been measured for the (100) face (ref. 11). Thus, for a fully or

even a partially hydrogenated diamond surface there exists no barrier for electron emission from the conduction band

edge provided one has electrons in the conduction band in the first place. However, diamond is a wide band gap

semiconductor and in the absence of n-type doping - this is the situation we are dealing with in CVD diamond and in

the diamond powders of ref. 6 - the concentration of electrons in the conduction band is negligible. Therefore,

electrons have to be injected over a barrier at the back contact which is of the order of 4 eV. The surface barrier of a

metal emitter is thus merely replaced by a Schottky barrier of similar height at the back contact in the case of

diamond and nothing has been gained. Also, this scenario cannot account for the improved field emission properties

of diamond-based materials after hydrogenation. Finally, there exists a wealth of data - some of them shown in Fig. 1

- that indicates that the field emission properties of CVD diamond actually improve with decreasing film quality

(refs. 12 to 16), i. e. as the relative concentration of sp 2 bonded graphitic material increases. In fact, in some cases,

purely sp 2 bonded graphitic samples have emission properties almost as good as diamond samples (ref 4). On the

other hand, if the graphitic phase is responsible for field emission as has been suggested in the past (refs. 2, 4, 16,

17), the role of hydrogen in improving the emission properties is hard to understand. Hydrogenation lowers the

electron affinity of diamond but does not affect the work function of graphite.

Here we report systematic investigations of the field emission properties of diamond-graphite composites

that were specifically designed to answer two main questions. (i) What are the relative contributions of diamond and

graphite to the field emission current; and (ii) what is the role of hydrogen in lowering the field emission threshold?

The use of diamond-graphite composites consisting of mixtures of nano-crystalline diamond and nano-crystalline

graphite powders is crucial for these investigations because in this way the relative volume concentration of the nondiamond phase can be controlled over the full concentration range from 0 to 100% without changing the morphology

of the emitting sites which constitute in each case the same nanocrystals of diamond or graphite. This last point

cannot be overemphasized because the morphology determines the field enhancement factor and is thus one of the

crucial factors affecting field emission. In any attempt to elucidate the microscopic mechanism of field emission in a

many phase system with varying surface morphologies such as CVD diamond it is thus mandatory to have control

over the field enhancement factor _3. The results to be presented below confirm the dominant contribution of the nondiamond phase to the field emission current. They also elucidate the subtle but important role hydrogen plays for the

field emission process. A detailed account of this work has been presented earlier (ref 18).

EXPERIMENT

For our experiment we used commercially available diamond and graphite powders with particle sizes

around 500nm that were thoroughly mixed and sprayed as a thin homogeneous layer on highly doped silicon. The

diamond crystallites are undoped and were originally oxidized as measured by x-ray induced photoelectron

spectroscopy (XPS). Hydrogenation of the composites was achieved by exposing them to a microwave-generated

hydrogen plasma for 2 min at 800°C under conditions that are used in diamond CVD. The field emission was

measured before and after hydrogenation with a gold coated metal anode of 80gm diameter in a vacuum of 10 .4 Pa.

The anode-cathode distance was determined by SEM with an accuracy of lgm.

RESULTS

In Fig. 3 current - field (I - F) characteristics are shown for a series of composites with a graphite content x

(defined as the relative volume fraction of graphite) varying from 0 to 100%. On a semilogarithmic plot the I - F

curves show the exponential increase in emission current with applied field (calculated from the applied voltage and

the anode - cathode distance) expected from the Fowler - Nordheim formalism for field emission. The exception is

the pure diamond sample which has not only a markedly different I - F characteristics but also requires a

significantly higher field than the sample with only 6.5% graphite for the same current. The compilation of threshold

fields Ft (defined as the field necessary to draw a current of 10 -12 A) in Fig. 4 shows that Ft drops with increasing

NASA/CP--2001-210948 115 August 6, 2001
graphitecontent from230V/pmforpurediamondtoaminimumvalueof23V/pmforthesampleswith87and100%

graphite, respectively• Afirstconclusiontobedrawnfromthesemeasurements isthatbeforehydrogenation pure

graphiteisthebestfieldemitterinthisseriesandthathecontributionfdiamondtoflaefieldcurrent isnegligibleby

comparison•

£1°-°_ ill//

,,.- 10-8

8

._'-_10-1° J

_ 1°-'2 ',wf r r

• . _. .... , . .

10 100

Field (V/pm)

Figure 3. Field emission current measured as a function of the

macroscopic field F for a number of diamond-graphite composites.

The graphite content by volume is 0% (bullet), 6.5% (circle), 24%

(full triangle), 57% (open triangle), 87% (open square), and 100%

(full square), respectively for the different curves.

A

E

¢

_'10(]

LI.

_- 10 ' 2'o' 4'o'6'o' s'o'160-

Graphite Concentration x(%)

Figure. 4. The threshold fields Ft, as a function of graphite content

x before (triangle) and after (circle) plasma hydrogenation. The

threshold fields for composites have been corrected for the reduced

surface fraction of graphite. The error bars on the data points

indicate the combined uncertainty in Ft due to the distance

measurement and the scatter in the I - F data.

After plasma hydrogenation the I - F curves retain their characteristic shape but the threshold fields change

in the way indicated in Fig. 4. For the films with high graphite concentration (87 and 100%) the threshold field is

unchanged within the experimental uncertainties• This important observation is proof that the morphology and thus

the field enhancement factor of flae crystallites has not been affected by the plasma treatment• Ft of the pure diamond

sample is lowered considerably to 65 V/pm but remains higher than any of the other threshold fields• The lowest

thresholds of about 17 V/pm are attained by the samples with average carbon concentrations of 24 and 56%. The

results for films with graphite concentrations below -60% reproduce thus the generally observed trend for CVD

diamond, that flae emission properties improve as the sample quality deteriorates and after hydrogenation.

0 20 40 60 80 100

Graphite Content x (%)

Figure 5. Measured resistivities of the diamond-graphite

composites as a function of graphite content (dots). The line

is the calculated resistivity of a two phase system in which

one phase is well conducting and the other insulating. The

curve is adapted from ref. 19 by scaling the end points to the

actually measured resistivities

_-'IE

E

o 14

:>_-lC

ad 6

q4

For an explanation of these results one has to consider two aspects, the transport of electrons from the back

contact to the surface and the field emission from the surface and the way both factors are influenced by

hydrogenation• The composites constitute macroscopically inhomogeneous media consisting of a well conducting

phase (graphite) and an insulating phase (diamond). The effective conductivity of such a medium is highly non linear

as a function of composition• For an ideal homogeneous mixture in which ohmic conductivity is the only conduction

mechanism flae effective conductivity exhibits a percolation threshold at about 33% graphite content as shown in Fig.

5. (ref. 19). This is indeed born out by the resistivities of the composites which are also plotted in Fig. 5. It is thus not

unreasonable to expect that a sizeable fraction of the applied voltage drops across the sample for any given current

and the Ft values plotted in Fig. 4 are only nominal threshold fields• According to the percolation model the voltage

drop is further expected to increase with decreasing graphite fraction which would explain the trend in Ft with x

NASA/CP--2001-210948 116 August 6, 2001
beforehydrogenation. Theseconsiderations aresupportedbytheoutcomeof anexperiment forthesamplewith

x=0.65.Wecomparethefieldemissioncurrent voltage(I-U)characteristic wththecasewheretheanodeisindirect

contact withthesamplesurface. Foranygivencurrentdensitymorethanhalfoftheappliedvoltagedropsoverthe

sampleandthusdoesnotcontributetothefieldinvacuo. Infact,convertingthevoltagedropoverfreespaceintoan

electricfieldusingtheknowncathode- anodedistanceyieldsfieldsbetween32and50V/gmthatareingood

agreement wihthefieldsrequiredtodrawcomparable fieldcurrentsfrompuregraphite. Thetwosampleswith100

and87%graphitearewellconducting andtheconductivity isnotinfluencedbyhydrogenation.

Afterhydrogenation allsamples, includingthepurediamondone,weremeasuredtobehighlyconductive.

Thisisdueto thesurfaceconductivitythatisgenerallyobservedonhydrogenated diamond(ref.20) Thus,for

hydrogenated composites thehighlyresistivepercolationtransport of electronstothesurfaceisbypassedbythe

surfaceconductivitywhichsuppliestheemittinggraphitecrystalliteswithelectrons. Followingourlineofargument

wewouldthusexpect thathethresholdfieldsofallsamplescontaininggraphitefallonthedashedhorizontal linein

Fig.4whichcorresponds totheFtappropriateforgraphitecrystallites. Instead, wefindthresholdfieldsthatareeven

lowerforall composites exceptthosewiththehighestgraphitecontent.A specialfieldemissionmechanism

involvingdirectlythesurfaceofdiamondcanbeexcludedbythedatafromthepurediamondsamplewhichexhibits

afarhigherthresholdfieldthanallgraphite- diamondcomposites evenafterhydrogenation. Theexplanation forthe

loweredthresholdfieldsisgiveninFig.6whereweshowphotoelectron yieldspectraofdifferent samplesbeforeand

afterhydrogenation. Inphotoelectron yieldthenumberofphotoemitted electronsperincident photonismeasured

andexperimental detailsaregivenin(ref.21).Themeasurements areperformedunderlowfieldconditionsandthe

photonenergyhVonwherephotoemission setsinisdetermined bythatpartofthesurfacewiththelowestthreshold

forphotoelectron emission. FordiamondhVonhastobegreaterorequaltothegapenergyEg= 5.5eV.ThecasehVon

=EgholdsforasurfacewithnegativeelectronaffinityandhVon=Eg+)_ holds for positive electron affinity. An

emission from the Fermi level, i.e. with a threshold corresponding to the work function of diamond is not possible

for a perfect diamond crystal because the density of states at E_ is zero. In Fig. 6a) and b) yield spectra for the pure

graphite sample and two composites with x = 0.57 and 0.24, respectively, are shown both on a logarithmic and on a

linear scale. All three spectra are identical with a common threshold photon energy hVon = 5.0+ 0. l eV which was

determined from a Fowler analysis of the spectra and which corresponds to the work function of graphite. Emission

intensity below hVon is due to thermal broadening of the Fermi-Dirac occupation function of the electrons in the

semi-metal graphite. These spectra confirm thus our initial conclusion that the graphitic part of the composites are

responsible for the field emission. The reason is that their emission barrier of 5.0 eV is smaller than that of oxidized

diamond which is at least 6 eV (ref. 11).

_101°11_ a

=

o

- hv

a. 102 4'0 45 50 55 60

Photon energy (eV)

a. 10 _

4;0 4;5 5;0 5,5 6',(

Photon Energy (eV)

I ex_

j 14x_

Photon energy (eV)

ail _ t _x_

tt I

Photon Energy (eV)

Figure 6. The Photoelectron yield spectra of three

composites with x=l (open squares), 0.57 (circles), and 0.24

(triangles) before (a and b) and after (c and d) plasma

hydrogenation. Both logarithmic (a and c) and linear (b

and d) plots are shown. Comparable onset photon energies

hVon are derived from a Fowler analysis of the spectra a

and c and a linear extrapolation of the yield spectra in b

and d, respectively. The emission below hVon is due to the

thermal broadening of the Fermi - Dirac occupation

function of the electrons in graphite.

Yield spectra recorded for the same samples after plasma hydrogenation are plotted in Fig. 6c) and d). Now

there is a remarkable difference in the photoemission thresholds. For the pure graphite sample hVon remains virtually

unchanged at 4.9 eV. However, for the mixed diamond-graphite composites hVon is lowered by about 0.6 eV to

values around 4.4 eV. Again, such low values cannot be due to electrons emitted from diamond for the reasons

mentioned above and the lowered thresholds must be ascribed to a special effect that affects graphite in close contact

NASA/CP--2001-210948 117 August 6, 2001
withdiamondandnotgraphitealone.It isalsoevidentfromtheyieldspectraofthecompositeshathesurfacesof

thediamondcrystalliteshaveturnedfrompositiveelectronaffinitybeforehydrogenationto NEAafter

hydrogenation. ThesignatureofNEAistherise(byoneorderofmagnitude) intheyieldspectraofFig.6c)at5.5eV

thatisabsent forPEAsurfaces(Fig.6a)ashasbeendiscussed elsewhere (refs.22,23).

DISCUSSION

Irrespective of the reasons that affect the photoemission threshold it is correct to state that the lowest

threshold for photoemission constitutes also the current limiting barrier for field emission: • = hVon. A reduction of

this barrier from 5.0 to 4.4eV corresponds, according to the Fowler - Nordheim theory to a reduction in threshold

field by 20%. This is exactly what we measure for the x = 0.56% and x = 0.24% composites after hydrogenation for

which the threshold fields are 18 V/gm compared to 23 V/gm for the graphitic reference samples. All experimental

evidence points towards the graphite grains as the electron emitting phase despite the fact that hydrogenation affects

solely the diamond phase. One aspect of hydrogenation, i.e. the surface conductivity of diamond, and its role in the

field emission process has already been discussed. That leaves the question how hydrogenation might affect the

emission threshold of graphite as demonstrated by the yield spectra and the reduced field emission threshold. A clue

comes from earlier investigations of photoelectron emission from graphitic patches on single crystal diamond (ref.

24). There we had shown that the photoemission threshold from the graphitic patches follows exactly the work

function _D of the diamond surface that was changed between 3.4 and 5.0 eV. The mechanism that allows for this

correlation between the emission threshold of graphite and the work function of the surrounding diamond is

operative here as well. For a discussion we refer the reader to refs. 18 and 24. Briefly, at the junction between

diamond and graphite electrons coming from the graphite are able to escape into vacuum by "sneaking by" the high

barrier of graphite and using the lower barrier of diamond. Naturally, this mechanism is only operative for electrons

escaping within a few nm of the diamond-graphite barrier (ref. 25). If the work function of diamond is higher than

that of graphite which is always the case for unhydrogenated or even oxidized diamond, nothing has been gained and

electrons have to overcome the full 5.0 eV work function of graphite.

SUMMARY AND CONCLUSIONS

From a systematic investigation of the field and photoelectron emission thresholds of diamond-graphite

composites we could identify key factors in the field emission process and in particular the role hydrogenation plays

for diamond based emitters. Our results shed some new light and substantiate earlier conjectures concerning the field

emission from CVD diamond and nano-crystalline diamond powder. Our results may be summarized as follows.

(i) The graphitic phase is a necessary requirement because it provides the emission sites. Emission from the Fermi

level is indeed observed in investigations that measure the energy distribution of field emitted electrons (ref. 26).

(ii) Hydrogenation is beneficial because it improves the conductivity of the diamond matrix. This is the reason why

hydrogenation was a necessary requirement to observe any field emission in the work of Zhu et al. (ref. 6). In the

m-size polycrystalline diamond used by Zhu et al. the graphitic phase can only be present in the form of grain

boundaries or graphitic patches at the surface. Thus, the percolation threshold for a conducting path between the

graphitic phases is apparently not attained and only after plasma hydrogenation does the diamond powder exhibit

sufficient conductivity so as to provide electrons for the emission sites.

(iii) Hydrogenation lowers the effective emission barrier for the graphitic part of any composite or CVD-diamond.

This lowering is, however, limited to those parts of the emitting graphitic surfaces that are in immediate contact with

the diamond (ref. 25). The lowering sets in as soon as the work function of the pure diamond phase drops below the

5.0 eV of graphite and that requires at least a partial hydrogenation.

Based on these results we are in a position to formulate requirements for an optimum, diamond based field

emission material. It would consist of a fine grained diamond layer with a columnar structure. Narrow (a few nm)

graphitic grain boundaries provide the conducting channels to the surface while the surrounding diamond serves after

hydrogenation as a means to lower the effective emission barrier. The lateral dimensions of the graphitic and the

diamond surface area should be adjusted so as to provide the highest possible fraction of graphitic surface area that is

compatible with the lowering of the effective work function. Such an emitter could in principle be realized, for

example, by electrically conducting ion tracks in diamond as was suggested recently (ref. 26). The main contribution

to an efficient field emission is, however, the local field enhancement at the emission sites.

NASA/CP--2001-210948 118 August 6, 2001
ACKNOWLEDGEMENT

Thisworkwasperformedunderthe auspices of the trinational D-A-C-H cooperation between Germany, Austria, and

Switzerland on the "" Synthesis of Superhard Materials".

* present address: Max-Planck-Institute for Solid State Research, Stuttgart, Germany

REFERENCES

1. C. Wang, A. Garcia, D.C. Ingram, M. Lake, M.E. Kordesch, Electron. Lett. vol.27, 1459 (1991)

2. N.S. Xu, R.V. Latham, Y. Tzeng, Electron. Lett. Vol. 29, 1596 (1993)

3. K. Okano, K. Hoshina, M. Iida, S. Koizumi, T. Inuzuka, Appl. Phys. Lett. vol.64, 2742 (1994)

4. M.W. Geis, J.C. Twichell, J. Macaulay, K. Okano, Appl. Phys. Lett. Vol.67, 1328 (1995);

M.W. Geis, J.C. Twichell, T.M. Lyszczarz, J. Vac. Sci. Technol. vol.B14, 2060 (1996)

5. K. Okano, S. Koizumi, S. Ravi P. Silva, G.A.J. Amaratunga, vol.381,140 (1996)

6. W. Zhu, G.P. Kochanski, S. Jin, Science vol.282, 1471 (1998)

7. T. Sugino, S. Kawasaki, Y. Yokota, Y. Iwasaki, J. Shirafuji, J. Vac. Sci. Technol. vol. B16, 720 (1998)

8. J.B. Cui, J. Ristein, M. Stammler, K. Janischowsky,G. Kleber, L. Ley, Diam. Rel. Mat. (to be published)

9. R.H. Fowler and L.W. Nordheim, Proc. R. Soc. London vol. A119, 173 (1928);

E.L. Murphy and R.H. Good, Phys. Rev. vol.102, 1464 (1956)

10. J.B. Cui, J. Ristein, L. Ley, Phys. Rev. Lett. vol. 81,429 (1998)

11. F. Maier, J. Ristein, and L. Ley, Phys. Rev. B (submitted)

M. J. Rutter and J. Robertson, Phys. Rev. vol. B60, 16135 (1999)

12. N.A. Fox et al., Appl. Phys. Lett. vol.71, 2337 (1997)

13. A.A. Talin et al., Appl. Phys. Lett. Vol.69, 3842 (1996)

14. I.T. Han et al., J. Vac. Sci. Technol. vol.B16, 2052 (1998)

15. K.H. Park et al., J. Vac. Sci. Technol. Vol.B16, 724 (1998)

16. W. Zhu, G.P. Kochanski, S. Jin, L. Seibles, J. Appl. Phys. Vol.78, 2707 (1995)

17. F. Lacher, C. Wild, D. Behr, and P. Koidl, Diam. Relat. Mater. Vol.6, 1111 (1997)

18. J. B. Cui, M. Stammler, J. Ristein, and L. Ley, J. Appl. Phys. Vol. 88, 3667 (2000)

19. D.J. Bergman and D. Stroud, in Solid State Physics, vol. 46, H. Ehrenreich, D. Turnbull, editors (Academic

Press, Boston 1992) p. 148

20. F. Maier, M. Riedel, B. Mantel, J. Ristein, and L. Ley, Phys. Rev. Lett. vol. 85, 3472 (2000) and references

tlaerein

21. J. Schfifer, J. Ristein, L. Ley, andH. Ibach, Rev. Sci. Instr. vol. 64, 653 (1993)

22. C. Bandis and B.B. Pate, Phys. Rev. Lett. vol.74, 777 (1995)

23. J. Ristein, W. Stein, and L. Ley, Phys. Rev. Lett. vol.78, 1803 (1997)

24. J.B. Cui, J. Ristein, and L. Ley, Phys. Rev. vol.B60, 16135 (1999)

25. For this reason tlae sample with $x=0.865 does not exhibit tlae lowering of the threshold field. On account of

tlae high graphite concentration most graphitic grains are not surrounded by diamond and therefore do not

encounter the lowering of tlae effective emission barrier.

26. O. Gr0ning, O.M. Kattel, P. Gr0ning, and L. Schlapbach, Appl. Phys. Lett. vol.71, 2253 (1997)

27. M. Waiblinger et al., Appl. Phys. vol.A69, 239 (1999)

Prof. Dr. Lothar Ley

Institute of Technical Physics

University of Erlangen

Erwin-Rommel-Str. 1

D-91058 Erlangen

Germany

Phone: ++49 9131 8527090

Fax: ++49 9131 8527889

E-mail: lothar, ley@physik.uni-erlangen, de

NASA/CP--2001-210948 119 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Field Emission from diamond films containing ion-implanted graphitic regions

N. Koenigsfeld and R. Kalish

Physics department and Solid State Institute, Technion, Haifa 32000, Israel

1st. Hoffman and R. Akhvlediani

Chemistry department, Technion, Haifa 32000, Israel

ABSTRACT

Electron field emission properties of different types of carbon films are promising for the realization of

cold cathode devices. Exceptionally good emission properties were obtained for carbon films which seem to contain

some kind of graphitic phase, introduced either during film deposition or in a conditioning stage (vacuum arc

discharge) during the field emission measurement.

In order to investigate the role of graphitic inclusions in CVD diamohd and their field emission properties

we have studied the field emission of diamond films deliberately patterned with Graphitic columns. Very thin

(200nm thick) insulating undoped diamond or Boron doped conducting films with sub-micron size grains were used.

Graphitic regions which penetrate all the way through the diamond layers to the conducting substrate were realized

by selective area ion-implantation induced damage followed by thermal annealing. Field emission current-voltage

curves are measured at different anode-sample distances and the turn of field of each sample is evaluated. Special

measures are taken to avoid any conditioning processes such as vacuum arc discharge that can modify material

properties.

The results illuminate the contribution of the presence of the Graphitic regions embedded in the diamond

layer to the FE properties, while eliminating any geometric field enhancement effects or creation of conductive pathways due to uncontrolled discharge, which were previously considered as possible mechanisms for improved FE

from diamond films.

Keywords: Field emission, Diamond properties and applications, Fowler-Nordheim injection, Diamond films.

Presenting author:

N. Koenigsfeld

Solid State institute

Technion, Haffa 32000, ISRAEL

Fax: (972) 4-8235107

Tel: (972) 4-8293907

noamk@ssrc.technion.ac.il

NASA/CP--2001-210948 120 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FIELD EMISSION CHARACTERISTICS OF ION-IMPLANTED CVD DIAMOND

T. Ono, T. Sakai, N. Sakuma, L. Zhang and H. Yoshida

FCT/JFCC and Toshiba Corporation,

1, Komukai Toshiba-cho, Saiwai-ku, Kawasaki 212-8582, Japan

M. Hasegawa and H. Okushi

Electrotechnical Laboratory, 1-1-4 Umezono, Tsukuba, Ibaraki 305-8568, Japan

ABSTRACT

Field emission from sulfur-implanted homoepitaxial diamond film, which showed n-type character at above 300

°C from Hall-effect measurement, was studied. As a result, the threshold field for field emission of about 12 V/gm,

the maximum current density of about 150 mA/cm 2 and stable emission current were obtained. Emission current

tended to become larger as the lateral distance between the measurement position and the electrode for electrical

contact became shorter. On the other hand, tlae threshold field tended to be unchanged. This tendency is thought to

be partly due to the large electrical resistance of the film.

Keywords: field emission, diamond, ion implantation, sulfur

INTRODUCTION

There is a possibility that electron field emitter devices may break through the limitations to existing Si power

switching devices, because a vacuum can act as a good insulator and emitted electrons can move freely (ref.1). Such

field emitters should have the capability to produce large and stable current at low operating voltage. Materials of

conventional field emitters are usually metals (Mo, W) and semiconductors (mainly Si), which have relatively high

work function (4 - 5 eV). This means that it is necessary to enhance electric field by using very sharp tips and

electrode with small holes to extract electrons. As a result, the fabrication process becomes complicated. However,

the operating voltage of conventional field emitter is still too high (10 - 100 V) for actual applications, and it is

necessary to lower the operating voltage.

Diamond is very attractive for application as an emitter material because of its negative electron affinity (NEA),

which allows electrons to escape easily from tlae surface. The NEA feature has the possibility to lower the operating

voltage drastically and realize planar structured emitter which would enable the fabrication process to be simplified.

In this paper, we report field emission characteristics of sulfur-implanted homoepitaxial diamond film, which

showed n-type character at above 300 °C from Hall-effect measurement.

EXPERIMENTAL

Diamond film was homoepitaxially grown on high-pressure high-temperature (HPHT) Ib (100) diamond

substrate by chemical vapor deposition (CVD). The dimensions of the substrate were 4 mm x 4 mm x 0.3 mm.

Sulfur ions were implanted with energy of up to 400 keV at 400 °C, and then heat treatment was carried out at 700

°C in nitrogen atmosphere. Electrodes were formed at the four comers of the sample surface by Ar-ion irradiation at

room temperature, which leads to the graphitization of the implanted region. Au/Ti electrodes for electrical contact

were formed by vacuum evaporation on tlae graphitized area. This sample showed n-type character at above 300 °C

from Hall-effect measurements (ref.2).

Field emission characteristics were measured using a diode configuration consisting of an emitter (sample) and a

metal rod anode as shown in Fig. 1. Metal wires were used to mount the sample on a metal holder and contact tlae

Au/Ti electrodes on the sample. The distance between the sample and rod anode was adjusted using a micrometer

and determined by a telescope. This setup was installed in a vacuum chamber at a pressure of 10 .7 Torr. The sample

and anode were connected to a DC voltage source. The emission current was measured using an electrometer and a

series resistance of 10 M_ was inserted for protection.

NASA/CP--2001-210948 121 August 6, 2001
__ Metalrodanode Resistor

DCvoltageso/rce_

Electrometer_///

Au/Tielectrode

Figure1.Experimental setuptomeasurefieldemissioncharacteristics

Metalwire

RESULTANDDISCUSSION

First,measurements werecarriedoutin flaemiddleofthesampleusingametalrodoflmmindiameter.

Theappliedvoltagewasincreasedfromzeroto themaximumandthendecreased to zero.Several voltagecycles

wererepeated. Althoughsurfaceofthefilmisnothydrogenised becauseoftheheattreatment, it showedclearand

reproducible fieldemission. Figure2showstypicalfieldemissioncurrent-voltage (I-V)characteristics. Thedistance

betweenthesampleandanoderodwas33_tm.Asshowninthefigure,theI-Vcurvehaddistincthysteresis during

firstvoltagecycle,whereasit disappeared aftersecondvoltagecycle.Apparently, someactivationprocessof

emissionoccurredduringtheperiodofincreasingoffirstvoltagecycle.After2.5hours,thesamemeasurement was

carriedout.Thesamehysteresis curvewasobservedagainduringfirstvoltagecycle.Thisindicatesthatactivationis

notadestructiveprocess. Figure3 showsemissioncurrent-voltage characteristics oflowappliedvoltageregion

beforeandaftertheabove-mentioned activation. Asshowniflaefigure,thethresholdfieldforelectronemissionwas

about12V_tm

1E-6• | • i • | • | • | •

D First cycle

• Second cycle

1E-7 .._'_ .

1E-8

1E-_

1E-l( '

500 1000 1500 2000 2500 3000

Voltage (V)

Figure 2. Typical field emission current-voltage

characteristics

2.0

fl Before activation

1.5 • After activation

° /

1.0

29

0.5

0.0

0 250 500 750 1000 1250

Voltage (V)

Figure 3. Emission current-voltage characteristics

of low applied voltage region before and after

activation

NASA/CP--2001-210948 122 August 6, 2001
Figure4showsemissioncurrent-voltage characteristics forthreedifferent distances (d)betweenthesampleand

anoderod.Theemissioncurrentwasdecreased withincreasingthedistance. Figure5 showsthecorresponding

Fowler-Nordheim(FN)plotsderivedfromdatashowninFig.4.ThefeatureoftheFNplotsforthedistanceof20

and33_tmshowedtworegions. In lowappliedvoltageregion,emissioncurrentsobeyedtheFNtheory.In high

appliedvoltageregion,asaturationtendencywasobserved. Thebehavior ispossiblydueto thelargeresistanceof

thediamondfilm.Bycontactingtheanoderodtothefilm,theresistancewasmeasured. Theresultwasabout10

G_.Thisindicatesthathevoltagedropof1000Vinthediamondfilmoccursfortheemissioncurrent of100nA.

Thestraight lineinFig.5corresponds tofitofthedataforthedistanceof53_tm.Fieldemissionisgoverned

bythewell-knownFNequation(ref.3)asfollows

log(I/V2)= log(1.54x10 -6o__2/_)_2.98x107_3/2/(_V) (1)

whereI isthecurrent inA,VtheappliedvoltageinV,o_theemissionareaincm2,[3thefieldenhancement factorin

cm-1and• theworkfunctionineV.Assumingthatemissionoccursuniformlyundertheanode, o_is equalto flae

sectional areaoftheanoderod(o_=0.00785cm2).[3and• wereobtainedas[3= 5.97x10-7cm -1and• = 1.46x10-7

eV.Ontheotherhand,assumingthattheelectricfieldis uniformundertheanode,[3is equalto the1/d([3=

188cm-1), o_and• wereobtainedaso_=3.65x10-Ucm2and• =0.0678eV.Thevalueof[_= 5.97x10-7cm -1based

onthefirstassumptionseemsunrealisticbecauseof[3< 1/d.Thesecondassumptionseemstobeadequate. The

valueofo_= 3.65x10-Ucm2basedonflaesecondassumptionisverysmallcomparedto thesectionalreaof flae

anoderod.Thissuggests thatfieldemissionarisesfromrestrictedpoints.

400 , , ,

t tl d=20 gm

A d=33 gm

_ 300[ " d=53gm

i 2001 __,,_ f :-

0"

0 1000 2000 3000 4000

Voltage (V)

Figure 4. Emission current-voltage characteristics

for three different distances between the sample

and anode rod

1E-13 , , ,

.__ a d = 20btm

A d= 33btm

1E-14 n° 0 [] = ;5

_A A & &

a,

1E-15 •

1E-16

i | | | |

d = 53btm,

0.000 0.001 0.002 0.003 0.004

1/V (V_)

Figure 5. Fowler-Nordheim plots derived from

data shown in Fig. 4

The emission current stability of the diamond film was also investigated. Figure 6 shows the typical example of

emission current fluctuations. The distance between the sample and anode rod was 33_tm. During the test, applied

voltage was fixed to 2250 V. At an average current of about 100 nA, the fluctuation of emission current for a period

of 15 minutes was less than 5%. It was observed that the emission current fluctuations decreased as the emission

current increased.

Next, measurements were carried out using a metal rod of 0.3 mm in diameter in order to investigate the

dependence of emission current on measurement positions as shown in Fig. 7. The distance between the sample and

anode rod was 33 _tm. It was observed that emission current tended to become larger as a lateral distance (x) between

the measurement position and the electrode for electrical contact became shorter. On the other hand, the threshold

field tended to be unchanged. This tendency is thought to be due to the above-mentioned large electrical resistance

of the diamond film. The dependence of emission current stability was also investigated. It was observed that

emission current became unstable as a lateral distance became shorter.

NASA/CP--2001-210948 123 August 6, 2001
©

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".'.* ".I".'.*" "o'. ". ". ". %', ". ". "." ,'.'.'.'." .°.'.'.'." ." .'.','." .° .'. ".'. ". °. "." .'. ". ". ". ".°.'. %'. ".°. ". "_'.'.'.'." ,'.'. ". ". '.,'.'.'.'.' .',','.. .'." ", "."

:i:i:!:if:i:!:i:i:i:i:i:i:i:i:i:i:i:!:i:i:i:i:i:i:i:!:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:ii:i:i:i:i:i:i:i:i:i:!:i:i:i:i:i:i:i:i:i:_!:i:i:i:i:i:i:i:i:i:i:i:i:i:i:

•.-..g......°.......°.......o.°,o.°.....°..°.°o .°.°.°.°°°..°...,..°.°.. ,° ...,...................°. • -.... - • .°..°° ............°....

• .......g...,.,., .°.,.,., ...°., ......, .,.., .., ..., .,.,.,., ...,°.....,.. ,, .....,., ,o., .,_, ....., .,°,., ..,.,.,., ..,., ..,. ,;, ", ..., .°..,... ,., .....,.

Figure 6. Typical emission current fluctuations of the diamond film

I (nA)

soU/J.//Y// jfioo

V(V)

0.75 - --_""0

I

x (ram)

Figure 7. The dependence of emission current on measurement position

Finally, we tried to bring the anode rod as close as possible to the electrode. The result is shown in Fig. 8.

The lateral distance between the measurement position and the electrode was estimated to be less than 50 gm. As

shown in the figure, the maximum current reached 100 gA at 3500 V. The calculated current density was very large,

2 ........

about 150 mA/cm. It seems difficult to explain this rapid increase of current by reduction of the resistance due to

short current path. Further investigation is necessary to understand this behavior.

NASA/CP--2001-210948 124 August 6, 2001
CONCLUSION

<" 150 t .......

:::L

dt0

I . ! .

0 500 1000 1500 2000 2500 3000 3500 4000

Voltage (V)

Figure 8. Emission current-voltage characteristics for the lateral distance of less than 50gm

Field emission from sulfur-implanted homoepitaxial diamond film, which showed n-type character at above 300

°C from Hall-effect measurement, was studied. Alflaough, surface offlae film is not hydrogenised because of the heat

treatment, it showed clear and reproducible field emission. The threshold field for field emission of about 12 V/_tm,

the maximum current density of about 150 mA/cm 2 and stable emission current were obtained. Emission current

tends to become larger as flae lateral distance between the measurement position and the electrode for electrical

contact become shorter. On the other hand, the threshold field tends to be unchanged. This tendency is thought to

be partly due to the large electrical resistance of the film. These results indicate that sulfur-ion implantation is a

promising method to enhance field emission from the diamond surface.

REFERENCES

1. T. Ono, T. Sakai, et al.: "Field Emitter Triode for Power Switching - ANew Power Device Candidate -,"

Proc. ISPSD'98, 1998, pp.151-154.

2. M. Hasegawa, D. Takeuchi, et al.: "n-Type Control by Sulfur Ion Implantation in Homoepitaxial

Diamond Films Grown by Chemical Vapor Deposition," Jpn. J. Appl. Phys., vol.38, 1999, pp.L1519-

L1522.

3. H.H. Busta: "Review: Vacuum Microelectronics - 1992," J. Micromech. Microeng., vol.2, 1992, pp.43-

74.

NASA/CP--2001-210948 125 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

THERMIONIC FEEM, PEEM AND I/V MEASUREMENTS OF HYDROGEN

TERMINATED N-DOPED CVD DIAMOND SURFACES

F.A.M.Kiick, J.M. Garguilo, B. Brown, R.J. Nemanich

Department of Physics

North Carolina State University

Raleigh, NC 27695-8202

USA

ABSTRACT

Imaging of field emission and photoemission from hydrogen terminated diamond surfaces is accomplished

with a high resolution photo-electron emission microscope (PEEM). Measurements obtained as a function of sample

temperature up to 1000°C display thermionic field emission images (T-FEEM). The system can also record the

emission current versus applied voltage. N-doped hydrogen terminated diamond films have been produced by

MPCVD with a N/C gas phase ratio of 48. The hydrogen terminated surfaces display uniform emission in PEEM at

all temperatures. No FEEM images are detectable below 500°C. At -680°C the T-FEEM and PEEM images are

nearly identical in intensity and uniformity. At higher temperatures the hydrogen passivation layer desorbes

continuously leading to a reduced electron emission current, with an increasing loss of the negative electron affinity

(NEA) property of the surface. The uniform electron emission of these films is to be contrasted with other carbon

based cold cathodes in which the emission is observed from only a low density of highly emitting sites. The I/V

measurements obtained from the hydrogen terminated N-doped films in the T-FEEM configuration show a

component that depends linearly on voltage at low fields. At higher fields, an approximately exponential dependence

is observed. At low temperatures employed (<700°C), the results indicate a thermionic component to the emitted

current.

I. INTRODUCTION

One of the most unusual properties of diamond is the possibility for its (100), (110) and (111) surfaces to

gain a negative electron affinity (NEA) by undergoing special surface treatments. The particular applications using

the NEA property of surface treated diamond films are electron emission sources and photocathodes. One way to

achieve a negative electron affinity on diamond surfaces is through hydrogen passivation. The NEA enables

electrons from the conduction band minimum to escape the diamond without an energy barrier at the surface. To

determine whether the surface exhibits a NEA, one can employ photoexcitation of the electrons into the conduction

band minimum and observe if electron emission occurs [1,2,3].

The technique of photo electron emission microscopy (PEEM) can be used to image the electron emission

properties of a surface in a controlled UHV environment. In addition, the same apparatus can be employed to image

field emission from a surface. This technique is termed field electron emission microscopy or FEEM.

In this paper, we investigate the electron emission properties of hydrogen terminated diamond in both

PEEM and FEEM modes as a function of sample temperature up to 900°C. The FEEM mode at high temperature

involves thermionic field emission processes, and we have termed this imaging condition as thermionic-field

electron emission microscopy (T-FEEM).

It is well established the diamond can exhibit p-type conductivity with substitutional boron doping, and

Hall measurements indicate an activation energy of 0.38eV for B-doped diamond. This value was reported to be in

good agreement with the ionization energy of 0.4eV calculated from the hydrogen-like model using 5.7 for the

dielectric constant of diamond, when assuming the hole effective mass equals the electron mass [4]. In contrast, ntype doping of diamond is difficult to obtain. The most common approach is to employ nitrogen doping, which

exhibffs a conductivity activation energy of 1.7eV for substitutional nitrogen, suggesting that the energy level caused

by the substitutional nitrogen lies 1.7eV below the conduction band minimum [5].

We have found significant differences in the thermal emission properties of B- and N-doped diamond

films. With increasing temperature highly N-doped diamond films exhibit strongly enhanced field emission, whereas

B-doped diamond shows very low field emission that remains nearly constant over a wide temperature range. For Ndoped diamond films with low N/C ratios the temperature dependent electron emission can not be observed. At

NASA/CP--2001-210948 126 August 6, 2001
highertemperatures wefindareductionintheelectronemission, whichweattributetothedegradationofthe

hydrogenpassivation layer.

II. EXPERIMENTAL DETAILS

N-doped diamond films were grown on 25mm diameter, 1 _cm Si <100> substrates. Sample preparation

included: 30min. ultrasonic abrasion in a diamond/titanium/ethanol suspension, rinse with methanol and drying with

nitrogen gas. The CVD reactor is a 1500W ASTeX IPX3750 microwave assisted CVD system with an RF induction

heated graphite susceptor. The temperature was calibrated with an optical pyrometer. The process gases were zero

grade N2, H2 and CH4 . The gas flows were monitored and controlled with mass flow controllers. During film

growth laser reflectance interferometry (LRI) was used to in situ monitor the thickness of the diamond layer.

The growth of N-doped diamond film was divided into three steps: (1) establishment of the nucleation

layer, (2) growth of the N-doped film, and (3) post growth surface treatment. The conditions for the nucleation layer

were 400 sccm H2, 8 sccm CH4, chamber pressure of 20 Torr, substrate temperature of 780°C and a microwave

power of 600 W. After formation of the nucleation layer the conditions were changed to 437sccm H2, 2.5sccm CH4,

60 sccm N2, pressure of 50 Torr, substrate temperature of 910°C and microwave power of 1300 W for a nitrogen

doped diamond film with N/C-48. For nitrogen doped films with N/C-0.5 the flow rates were adjusted accordingly.

By growing for 4 to 8 hours, films with thicknesses of 0.5gm and 1 gm were fabricated. Growth was terminated by

shutting off the gas flows except H2, reducing the microwave power to 600W and reducing the substrate temperature

to 780°C. The diamond film was then treated with H2 plasma for 5 min at a pressure of 20 Torr.

The major focus of this study is the emission from N-doped films. For comparison, PEEM and FEEM were

measured on B-doped CVD diamond films, which were obtained from Dr. K. Kobashi of Kobe Steel.

The PEEM and FEEM measurements were completed in an Elmitec UHV-Photo Electron Emission

Microscope [6]. Measurements were obtained at a base pressure less than lxl0 -1° Torr. The field of view can be

changed from 150gm to 1.5gm with a resolution better than 15 nm. For all measurements a high voltage of 20 kV is

applied between the anode and the sample surface, and the anode is a distance of 2 to 4mm from the surface

(resulting in an applied field of 10 to 5 V/gm). In the PEEM measurements a mercury arc lamp was used as the UVlight source. In the FEEM measurements no UV excitation was employed and the emission was due to the high

applied field.

The system has sample heating up to 1200°C to obtain T-FEEM images. The electron emission current

from the sample surface can be monitored and recorded to obtain current-voltage dependence. Electrons emitted

from the sample pass through a perforated anode and are imaged using electron optics. The image is intensified with

a double microchannel plate and a fluorescent screen. A CCD camera is used for image capturing. The gain of the

system is dependent on the voltage on the image intensifier. The images reported here have been digitally processed

to remove dark regions of the intensifier.

III. RESULTS AND DISCUSSION

The influence of nitrogen on the growth and properties of diamond films has been reported in several

articles [7-11]. With increasing nitrogen content in the growth chamber, increased sp 2 bonding signatures are

observed and the increased FWHM of the 1340 cm -1 diamond peak also increases significantly. Displayed in

Figure 1, the Raman scattering from the 0.5gm diamond film shows a broad spectral background extending from

1100 to 1600cm -1. The spectral components attributed to sp 2 bonded carbon include the 1355 1580 cm -1 doublet

due to microcrystalline graphite, and the broad peak centered near 1500 attributed to sp 2 bonding in or around

diamond crystals. A feature at -1140 cm -1 has been attributed to disordered sp 3 bonded carbon, and the broadening

of the 1340 cm -1 diamond peak indicates small grain size and the presence of defects or impurities [12, 13]. It should

be noted that Raman spectroscopy is about 75 times more sensitive to sp 2 bonded carbon sites than it is to crystalline

diamond when 514.5nm excitation is used [14]. It is evident that the film has a significant diamond crystal

component but it also exhibits a distribution of sp 2, and the variations in bonding are on a nanometer scale. When

using nitrogen gas as the nitrogen dopant source, it is found that the nitrogen concentration in the diamond film is

low [15,16]. However, the Raman spectrum shows that the diamond films change significantly with the addition of

nitrogen. With increasing nitrogen concentration in the growth chamber the diamond peaks FWHM at 1340 cm -1

increases and the peak corresponding to the sp 2 bonded carbon becomes more significant.

NASA/CP--2001-210948 127 August 6, 2001
1.1-

0.9

0_ 0.8

c 0.7

ill

E

0.6

0.5

0.4

N/C=4

Diamond 1340 Graphite 1554

1000 1100 1200 1300 1400 1500 1600 1700 1800

Raman shift [l/cm]

Figure 1. Raman Spectroscopy of nitrogen doped diamond films with N/C=48 and N/C=0.5 showing a broadened

diamond peak and a spectral components attributed to sp 2 bonding for the diamond film with higher nitrogen

concentration. The spectrum was taken at 514.5nm at 300K.

Shown in Fig. 2 are PEEM images of different magnification of the N-doped diamond film at room

temperature. The images show that the surface exhibits a uniform electron emission over the whole sample surface.

The images show a fine textured grainy structure that corresponds to the morphology observed by SEM. The

brighter spots in the image of the highest magnification are attributed to emission from the most raised points on the

surface. These regions will have the highest field. Images resulting from PEEM measurements can be understood in

terms of photoelectron emission from NEA diamond [17,18]. The photo excited electrons in the conduction band of

the diamond are emitted into vacuum following the theory of photoemission. Note that the emission is uniform over

the surface with no evidence of spot emission often observed in other carbon based cold cathode materials.

This uniform emission is indicative of a relatively uniform surface barrier and of a uniform carrier

distribution near the surface. In the following we suggest that observation of uniform emission may be indicative of

emission from the conduction band.

(a) (b) (c)

Figure 2. PEEM images obtained at room temperature with Hg-arc lamp excitation using a field of view and channel

plate voltage of: (a) 150gm, 1.476kV, (b) 50gm, 1.596kV and (c) 20gm, 1.627kV. A higher channel plate voltage is

required to image at higher magnification.

By turning off the Hg-arc lamp and increasing the sample temperature T-FEEM of the nitrogen doped

diamond film can be observed. T-FEEM images were obtained at various temperatures using a constant channel

NASA/CP--2001-210948 128 August 6, 2001
platevoltage. Figure3(a),(b),(c),(d)showsT-FEEMimageswitha20gmfieldofviewatachannel platevoltage

of 1.55kV,corresponding tosubstratetemperatures of640°C,680°C, 700°Cand720°C, respectively. FromtheTFEEMimages,averyuniformelectronemissioncanbeobservedthatincreasesin intensitywithincreasing

temperature. It isevident thatheFEEMisverytemperature dependent withastrongincreaseabove640°C. Similar

measurements onB-doped, p-typediamondfilmsdonotdisplayobservableFEEMimagesbelowtemperatures of

800°C. Wehavealsorepeatedthemeasurements oaMoplatewithsimilarsurfaceroughness tothediamond. Here

againFEEMwasobservable onlyfortemperatures greaterthan900°C. It isevident thatheN-dopeddiamondfilms

exhibitsignificant emissionatrelativelyowtemperatures.

Theuniformityoftheemissionsuggeststhatheelectronsmayoriginatefromtheconductionbandofthe

diamond.Alternatively, it ispossiblethatheemissionoriginatesfromdefectstatesinthebandgap.If thisisthe

casethentheuniformityoftheemissionwouldrequirethathedefectsareuniformlydistributedthroughout thefilm.

Fromtheemissionuniformityin theT-FEEMimagesit isnotobviousthatgraphiticregionshowenhanced

thermionicfeldemissioncontributiontotheelectronemissioncurrent.Thefactthatfieldemissionfromdiamond

filmsappearstobeanti-correlated withdiamondquality,thusshowingenhancedemissionforfilmswithhighernondiamondcarbonhastobeinvestigated forthethermioniccase[19].

(a) (b) (c) (d)

Figure3.T-FEEMimagesofN-dopeddiamondat640°C(a),680°C(b),700°C(c)and720°C(d).All imageswere

obtainedatachannel platevoltageof1.55kV, andthefieldofviewis20gm.

Foreachsampletemperature of 640°C,680°C,700°Cand720°CanI/Vcurvewasrecorded. Figure4

showsasetofI/Vcurvesobtainedatvarioustemperatures. Atlowfieldswe

6O

5O

4o

E

•'- 30

o

20

10

1,

....... 720°C ,t,o

,,,'

....... 700oc ,_'

,,,,

......... 680oc z,:°

I'; z:

....... 640°C . ,/'[.2 ,'

y*," ,^

........ ., • -.--''_._,,2_._:,._ o_, /

y._ .o__ _ - _-

__-o5 10 15 20

Voltage [kV]

Figure 4. The emission current vs. applied voltage obtained in the electron emission microscope for the N-doped

diamond film. Note that the baseline signal increases with temperature.

NASA/CP--2001-210948 129 August 6, 2001
findanearlyconstant valueoftheelectronemissioncurrent. Increasingtheanodevoltagebeyond15kVresultsinan

exponential behavior, whichdominatestheemissioncurrentcharacteristics. Electrical conductivitymeasurements

ontypeIIanaturaldiamond, thathasnitrogenasthedominant impurity,showsanexponential temperature

dependence ofthecarrierdensity, andabouttwoordersofmagnitudeofchangethatisattributedtochangesinthe

mobility.FortypeIIanaturaldiamondS.Hanet al. report an activation energy of 1.4eV that likely reflects the

excitation of electrons from nitrogen levels [20].

To show how the hydrogen passivation layer stability at higher temperatures affects the electron emission,

we have imaged the hydrogen terminated diamond surface before and after the anneal process. Figure 5 shows

PEEM images of a nitrogen doped diamond film with N/C-60. The left part of Figure 5 shows the PEEM image

before annealing the sample to 800°C and the right side shows the image obtained after annealing. It is evident that

the emission is highly reduced after annealing. We suggest that the change is due to the evolution of the hydrogen

passivation from the diamond surface, which in term increases the electron affinity from a negative value to a

positive value [21, 22].

Figure 5. The electron emission as imaged in PEEM mode before (left part) and after (right part) annealing a

nitrogen doped diamond film to 800°C showing the degradation of the hydrogen passivation that induces the NEA

and thus reduces the electron emission from the diamond film. The field of view is 50gm, the channel plate voltage

1.563kV.

The emission characteristics presented here suggest a thermionic contribution to the electron emission of

N-doped diamond films at elevated temperatures. It appears that the emission at low fields could be attributed to

thermionic emission of electrons in the conduction band of the diamond. The exponential increase of the current at

higher fields indicates a tunneling process. This process may be related to electron emission at the film surface or to

supply of electrons at the film substrate interface. Future research will explore these issues.

IV. CONCLUSIONS

We have grown nitrogen doped diamond films by plasma assisted CVD. Raman spectroscopy shows a

broadened diamond peak and the presence of sp2 bonding in the film. PEEM investigation of the NEA nitrogen

doped diamond films shows uniform electron emission from the surface at room temperature. At temperatures below

500°C FEEM does not resolve electron emission from the specimen. Increasing the sample temperature up to 720°C

we find a strong increase in the electron emission suggesting that electrons from the nitrogen donor levels are

excited into the CB. It is also noted that defects could play a role in the emission process. Further investigation will

explore the contribution of defect states and attempt to understand the thermionic emission properties.

Acknowledgements: We gratefully acknowledge the Duke University Free Electron Laser Laboratory where all the

Electron Emission Microscopy Experiments were performed. This research is supported through the Frontier

Carbon Technology Project, the Japan Fine Ceramics Center, ONR and PTS Company - John Driscoll.

NASA/CP--2001-210948 130 August 6, 2001
REFERENCES

1. R.J. Nemanich, P.K. Baumann, M.C. Benjamin, O.-H. Nam, A.T. Sowers, B.L. Ward, H. Ade and R.F. Davis,

Appl. Surface Sci. 130-132 (1998) 694-703

2. R.J. Nemanich, P.K. Baumann, M.C. Benjamin, S.P. Bozeman, B.L. Ward, Proceedings of the Int. School

Physics Enrico Fermi, IOS Press, Amsterdam, 1977, p. 537.

3. J.B. Cui, J. Ristein, and L. Ley, Physical Review B, PRB60, 23, 1999, p.16135

4. K. Okano, Doped Diamond in Diamond: Electronic Properties and Applications, L.S. Pan and D.R. Kania edt.

Kluwer Academic Publishers 1995

5. R.G. Farrer, On the substitutional nitrogen donor in diamond, Solid State Comm. 7 (1969) pp. 685-687

6. H. Ade, W. Yang, S.L. English, J. Hartman, R.F. Davis and R.J. Nemanich, Surface Rev. and Lett. Vol. 5, No.

6, 1257-1269 (1998)

7. V. Baranauskas, B. B. Li, A. Perlevitz, M. C. Tosin, and S. F. Durant, J. Appl. Phys. 85, 7455 (1999)

8. N. Jiang, A. Hatta and T. Ito, Jpn. J. Appl. Phys. 37, 1175 (1998)

9. I.T. Han and N. Lee, J. Vac. Sci. Technol. B 16(4), 2052 (1998)

10. W. Mfiller-Sebert, E. W6rner, F. Fuchs, C. Wild and P. Koidl, Appl. Phys. Lett. 68 (6), 759, (1996)

11. A. T. Sowers, B. L. Ward, S. L. English and R. J. Nemanich, J. Appl. Phys. 86 (7), 3973, (1999)

12. R. J. Nemanich, J. T. Glass, G. Lucovsky, and R. E. Shroder, Raman scattering characterization of carbon

bonding in diamond and diamondlike thin films, Journal of Vacuum Science & Technology A: Vacuum,

Surfaces, and Films May 1988 Volume 6, Issue 3 pp. 1783-1787

13. M. Park, A.T. Sowers, C. Lizzul Rinne, R. Schlesser, L. Bergman, R.J. Nemanich, Z. Sitar, J.J. Hren and J.J.

Cuomo, J. Vac. Sci. Technol. B 17(2), 1999

14. R. E. Shroder and R. J. Nemanich, Analysis of the composite structures in diamond thin films by Raman

spectroscopy Phys. Rev. B 41, 3738 3745 (1990)

15. S. Jin, T.D. Moustakes, Appl. Phys. Lett. 65 (1994) 403

16. E. Boettger, A. Bluhm, X. Jiang, L. Schafer, C.P. Klages, J. Appl. Phys. 77 (1995) 6332

17. F.J. Himpsel, J.A. Knapp, J.A. van Vechten and D.E. Eastman, Phys. Rev. B 20, 624 (1979)

18. J. van der Weide, Z. Zhang, P.K. Baumann, M.G. Wensell, J. Bernholc and R.J. Nemanich, Phys. Rev. B 50,

5803 (1986)

19. J.B. Cui, J. Ristein, M. Stammler, K. Janischowsky, G. Kleber & L. Ley, Hydrogen termination and electron

emission from CVD diamond surfaces: a combined secondary electron emission, photoelectron emission

microscopy, photoelectron yield, and field emission study, Diamond and Related Materials 2000, 9:3-6:1143-

1147

20. S. Hart, L. S. Pan, D. R. Kania, Dynamics of Free Carriers in Diamond, Diamond, in Electronic Properties and

Applications, L. S. Pan, D. R. Kania Eds., Kluwer Academic Publishers 1995

21. J.B. Cui, J. Ristein, M. Stammler, K. Janischowsky, G. Kleber, L.Ley, Diamond and Related Materials 9 (2000)

1143-1147

22. P.K. Baumann and R.J. Nemanich, Surface cleaning, electronic states and electron affinity of diamond (100),

(111) and (110) surfaces, Surface Science, 409 (2) (1998) pp. 320-335

NASA/CP--2001-210948 131 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

N-DOPED CVD DIAMOND FILMS AS A LOW TEMPERATURE THERMIONICFIELD ELECTRON SOURCE

F. A. M. Kiick, J. M. Garguilo, B. Brown, R.J. Nemanich

Department of Physics

North Carolina State University

Raleigh, NC 27695-8202

Abstract

Microwave Plasma assisted Chemical Vapor Deposition (MPCVD) has been utilized to synthesize nitrogen doped

diamond films for application as a low temperature thermionic-field emission cathode. The critical result of this

study is the imaging of electron emission from UV photo excitation and from thermionic-field emission. The

samples were imaged in UHV by photo electron emission microscopy (PEEM) using a UV Hg lamp for

photoemission excitation. The same instrument was used to obtain the thermionic-field emission electron

microscopy images (T-FEEM) at temperatures up to 900°C and to record the electron emission current in

dependence of the applied voltage and substrate temperature. The microscope has 15nm resolution. Nitrogen doped

diamond films were grown at substrate temperatures from 850°C to 900°C, and the Raman spectra of the films

showed a strong diamond peak at 1332cm -1 and weak signal from the graphitic regions in the sample. Field emission

could not be measured at room temperature, but the PEEM images showed relatively uniform emission. The PEEM

images showed little change as the temperature is increased. At -680°C the T-FEEM and PEEM images are nearly

identical in intensity and uniformity. This is to be contrasted with other carbon based cold cathodes in which the

emission is observed from only a low density of highly emitting sites. The I/V measurements obtained from the Ndoped films in the T-FEEM configuration show a component that depends linearly on voltage at low fields. At

higher fields, an approximately exponential dependence is observed. It appears that the emission at low fields could

be attributed to thermionic emission of electrons in the conduction band of the diamond. The exponential increase of

the current at higher fields indicates a tunneling process. This process may be related to electron emission at the film

surface or the supply of electrons at the film substrate interface. These results indicate a promising new material for

the production of low temperature, high brightness electron sources.

Keywords: diamond, peem, thermionic field emission

NASA/CP--2001-210948 132 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Field Emission, PEEM and FEEM Measurements of Emitting Sites of MPCVD grown NanoCrystalline

diamond Films

J. M. Garguilo, F. A. M. Kiick, Billyde Brown, R. J. Nemanich

Department of Physics, North Carolina State University, Raleigh, NC 27695-8202

Tel: 1-919-515-3225; Fax: 1-919-515-7331; Email: Robert Nemanich@ncsu.edu

ABSTRACT

Field emission characteristics of nanocrystalline diamond thin films grown by microwave plasma assisted

chemical vapor deposition (MPCVD) were examined in this study. Raman spectroscopy shows a broad peak at

1350 cm -1 associated with sp 2 bonding. Field emission characteristics were measured in ultrahigh vacuum with a

variable distance anode technique. Variable distance field emission measurements of the nanocrystalline diamond

films resulted in relatively low threshold fields of 4-6 V/gm for a current of 0.5 nA. In addition, Photo Electron

Emission Microscopy (PEEM) and Field Electron Emission Microscopy (FEEM) measurements were conducted to

image the emitting sites of the nanocrystalline diamond films. Electron emission current measurements of the

emitting sites were recorded. PEEM images show a smooth surface with very strong emitting spots that retain the

same high electron emission during FEEM. FEEM was employed to characterize the spot size of the emitting sites.

The results show that the emission originates from a localized region smaller than 0.1cm. The emission site density

of the nanocrystalline diamond films is - 104 cm -2. Thermionic Field Emission Electron Microscopy (T-FEEM) does

not result in a change in the electron emission for temperatures up to 1000 ° C. These results suggest that the emitting

sites do not originate solely from grain boundaries of the nanocrystalline diamond film but involve additional

emission mechanisms, which is a topic of ongoing research.

INTRODUCTION

Field emission from highly sp 2 bonded diamond films has been characterized as non-uniform spotty

emission (1). This non-uniform emission is indicative of field enhancement effects. Since the film surfaces are

relatively smooth, it has been suggested that the origin of the emission is from the nanocrystalline grain boundaries

(2). However, the low emission site density suggests that only specialized structures act as emission sites.

This study further investigates the electron emission properties of nanocrystalline diamond films through a

variety of methods in order to better understand the mechanism of electron emission.

The technique of photo electron emission microscopy (PEEM) can be used to directly image the electron

emission properties of a surface in a controlled UHV environment. The same apparatus can be used to image field

emission from a surface. This is known as field electron emission microscopy (FEEM). With these measurements,

we are able to image the individual electron emission sites of the nanocrystalline diamond film and obtain a

maximum limit to the size of an individual site. This technique also gives us the density of emitting sites of a

sample.

Atomic force microscopy is used to measure the surface roughness and density of grain boundaries on the

surface. By comparing the emission site density obtained by FEEM to the grain boundary density obtained using

AFM, we can related the emission in a general way to the film structure. Our results indicate that few grain

boundaries emit suggesting that an additional mechanism for electron emission from the grain boundaries must be

adopted or an entirely new explanation needs to be generated.

EXPERIMENTAL DETAILS

Nanocrystalline diamond films were grown on 25mm diameter 1 g-cm Si <100> substrates. Sample

preparation consisted of a thirty minute ultrasonic abrasion in a diamond/titanium/ethanol suspension (3), rinse with

methanol and drying with nitrogen gas. The CVD reactor is a 1500W ASTeX IPX3750 microwave assisted system

with an RF induction heated graphite susceptor. The temperature was calibrated with an optical pyrometer. The

process gases were zero grade hydrogen and methane. The gas flows were monitored and controlled with mass flow

controllers. During film growth, laser reflectance interferometry (LRI) was used to monitor the thickness of the film.

The growth of the nanocrystalline film was obtained using gas flows of 180 sccm of hydrogen and 20 sccm

of methane and a chamber pressure of 20 Torr. The substrate was held at a temperature of 900 ° C and a microwave

NASA/CP--2001-210948 133 August 6, 2001
powerof900wattswasused.LRImeasuredadepositionrateofapproximately 0.16gmperhour.Filmsweregrown

for2to6hours.

Thresholdfieldswereobtainedwithvariabledistancefieldemissionmeasurements. Thesamplewasplaced

anarbitrarydistancefromapolishedmolybdenumanode. Anodevoltagesbetween0and1100Vwereprovidedbya

Keithley237SourceMeasureUnitwhichalsomeasuredthecurrent generated. Thevoltageneededtogeneratethe

thresholdcurrentwasrecorded, andtheanodewasmovedtowardsthesamplebyaknownstepsize.Byrepeating

thisprocessseveral times,athresholdfieldforaselectedthresholdcurrentcanbecalculated. Thestepsizeforthis

systemis55nm.(4,5)

Ramanspectroscopy wasusedtoidentifythebondingcharacteristics ofthenanocrystalline diamondfilm

andatomicforcemicroscopy wasemployedtoimagethesurface.

PEEMandFEEMmeasurements providedimagingof electronemissionsitesforthisstudy.These

measurements wereperformedusinganElmitecUHV-PhotoElectronEmissionMicroscope(6).Measurements

wereobtainedinabasepressureoflessthan1x 10-l°Torr.Thefieldofviewcanbechangedfrom150to1.5gm

witharesolutionofbetterthan15nm.Forallmeasurements, ahighvoltageof20kVwasappliedbetweenthe

anodeandthesamplesurface. Theanodeisspacedadistanceof2to4mmfromthesurfacersultinginanapplied

fieldof 10to5V/gm.ForthesePEEMmeasurements, amercuryarclampwasusedastheUVlightsource. In

FEEMmeasurements noUVexcitationwasemployed, andtheemissionwasduetothehighappliedfield.

Thesystemhassampleheatingcapabilities upto1200°CtoobtainT-FEEMimages. Theelectronemission

currentfromthesamplesurfacecanbemonitoredandrecordedtoobtaincurrent-voltage dependence. Electrons

emittedfromthesamplepassthroughaperforatedanodeandareimagedusingelectronoptics.Theimageis

intensifiedwithadoublemicrochannel plateandafluorescent screen. ACCDcameraisusedtocapturetheimages.

Thegainofthesystemisdependent onthevoltageontheimageintensifier. Theimagesreportedherehavebeen

digitallyprocessed toremovedarkregionsoftheintensifier.

RESULTS AND DISCUSSION

Raman spectroscopy was used to identify the predominant bonding in the diamond films. Figure 1 shows a

typical Raman spectrum obtained from a nanocrystalline diamond film. The characteristic 1355-1580 cm -1 peaks for

sp 2 bonded carbon are present. These peaks are broad indicating a small grain size (7, 8). A small shoulder at

approximately 1140 cm -1 is sometimes present and has been attributed to disordered sp 3 bonded carbon (7). Raman

spectroscopy is about 75 times more sensitive to sp 2 bonded carbon sites than it is to crystalline diamond when the

514.5 nm excitation is used (9).

1:34u

1582

/\

i

O00 1100 1200 1300 1400 1500 1600 1700 1800

Raman Shift (Rcm-1)

Figure 1. Raman spectrum of a nanocrystalline diamond film showing the characteristic peaks associated with sp 2

bonding. The spectrum was obtained using 514.5 nm excitation at -300 K.

NASA/CP--2001-210948 134 August 6, 2001
Variabledistancefieldemissionmeasurements wereemployedtomeasureI-Vplotsataseriesofdifferent

anode-to-sample distances. Thevoltagetoobtainacurrentof0.5nAwasdeterminedforeachrelativedistance.

Thevalueswereplottedversusrelativedistanceandfit withastraightline. Theslopeofthelinefit givesthe

appliedfieldforthresholdemission.Thresholdfieldsforthenanocrystalline diamondfilmswasfoundtovary

between4to6Vpergm(foracurrent of0.5nA).Thiscorresponds tothresholdfieldspreviouslyobserved(10,11).

Insomeinstances, thresholdfieldsofapproximately 1V/gmwereobservedinseveral samplespreparedinthesame

manner. Thesesampleshaveshownimprovedemissioncharacteristics inFEEMmeasurement tchniques although

theoriginofthisimprovedemissionisasyetunknown.

Theemissionsitedensityofthenanocrystalline diamondfilmswasfoundtobeabout1siteper10 -4cm2in

FEEMmeasurements. Withthesetechniques, bothemissionsitesandthenonemittingareasofthefilmscanbe

imagedandcompared. Figure2showscomparisons betweenPEEMandFEEMofareasofthefilmwherean

emissionsiteisbothpresent andabsent. Whenanemissionsiteisabsent fromthefieldofview,thereisnoFEEM

imagedetectable, andthePEEMimagedisplaysauniformemission. Thisindicatesasmoothsurfaceofthefilm

whichwasconfirmedwithAFM.Whenanemissionsiteispresent, theuniformbackgroundisstillobservable,

however, theimageisdominatedbytheintenseemissionfromalocalizedarea.WhentheUVexcitationisremoved,

thissitecontinuestoemitasaresultofthefieldapplied.It isevidentthathecurrentobservedinthevariable

distancefieldemissionmeasurements areduetothesebrighthighlyemittingsites.

Figure2.150gmfieldofviewPEEMandFEEMcomparison ofananocrystalline diamondfilm:(a)PEEMimage

of uniformlyemittingfilmwithnoemissionsitepresent, (b)FEEMimageof sameregionshowingnoemission

detectable, (c)PEEMimageofaregionwithabrightemissionsiteandbackground, and(d)FEEMimageofthe

sameregionshowingthatheemissionisdetectablefromalocalizedregionoftheemissionsite.All imageswere

obtainedwiththesampleatroomtemperature. ThePEEMimagesemployedaHg-arclampexcitationsource.

Afieldofviewof1.5gmcanbeobtainedwiththephotoelectronemissionmicroscope. Figure3showsa

comparison ofseveral emittingsitesonasinglenanocrystalline sample. Sixbrightemissionsiteswereimagedwith

thehighest magnificationofthesystematafieldofviewof 1.5gm.Thechannel patevoltageforeachspotwas

adjustedtoimagetheemissionsiteclearly.Theintensityofemissionisnotuniformforthedifferent emissionsites.

Arangeofemissionintensityisobtainedfromthedifferentsiteswithsite3beingthebrightest andsite1beingthe

dimmest. Thesizeofeachspotislessthan0.1gm.Whilewehavenotapplieddetailedanalysisofthesites,it

appearsthathesmallest siteshavecharacteristic dimensions of-20nm.

NASA/CP--2001-210948 135 August6,2001
Spot 1,1.150kV Spot 2,1.075kV Spot 3,1.050kV

Spot 4,1.100kV Spot 5,1.050kV Spot 6,1.100kV

Figure3.Comparison ofFEEMimagesofsixdifferent emissionsitesonthesamenanocrystalline flmsurface. All

imageswereobtainedatroomtemperatureanda 1.5gmfieldof view,themaximummagnificationf the

microscope. Thevoltageappliedtotheimageintensifier channel platewasvariedtoobtainthebestresolutionandis

listedforeachimage.Ahigherintensifiervoltageindicatesweakeremission. Thesixsitesarenotequalintheir

electronemissioncharacteristics. All sitesarelessthan0.1gminsize.

Weexploredtheeffectoftemperature ontheemissioncharacteristics ofasinglemissionsite.Imageswere

obtainedfromthesamesiteattemperatures upto1000°C.Figure4showsanemissionsiteat600,800,and1000°

C.Theintensityoftheemissiondoesnotchangeoverthistemperature range.Thermionicfeldelectronemission

microscopy(T-FEEM)of thenanocrystalline diamondfilmsshowednotemperature dependence oftheelectron

emissionfortemperatures upto1000°C.Thisresultssuggests thatheelectronemissiondoesnotoriginatefromthe

conductionbandandthatheworkfunctionislarge(>4eV).

Figure4.Thermionicfieldelectronemissionmicroscopy (T-FEEM) seriesofasinglenanocrystalline emissionsite.

Theintensityoftheelectronemissiondoesnotchangeupto1000°C.Imagesweretakenatafieldofviewof20gm

andanintensifier channel platevoltageof1.3kV.

Figure5showsacomparison betweena5gmAFMscanofananocrystalline diamondsampleanda5gm

fieldofviewFEEMimage.It isevidentthatherearemanymoregrainboundariesthanemissionsitesatthese

NASA/CP--2001-210948 136 August6,2001
dimensions. Thedensityofemissionsitesascomparedtothedensityofgrainboundaries ladstotheconclusionthat

additional emissionmechanisms mustbeconsidered.

•..:..... ".:.- :.. .:.

...

.,.' ,, ".

Figure 5. Comparison of AFM (on left) and FEEM (on right) images from a nanocrystalline film. Each image has a

5 gm field of view. The AFM image shows many small grains and grain boundaries while the FEEM image shows

only one emission site over a similar area.

CONCLUSIONS

Nanocrystalline diamond films were prepared via plasma assisted CVD. Raman spectroscopy of these films

identifies peaks due to sp 2 bonded carbon, and the films contain nanocrystalline diamond along with significant sp 2

bonding.. The thermionic independence of the films suggests that the electron emission does not originate from the

conduction band and that the work function of the film is high. Atomic force microscopy gives average surface

roughness on the order of 100 angstroms correlating to a relatively smooth surface. Field emitted electron energy

distribution measurements of nanocrystalline diamond indicate electrons originate from the Fermi level of the

material (11). A comparison between AFM and FEEM images of nanocrystalline samples shows a disparity between

the number of grain boundaries and emission sites as displayed in Figure 5. This leads to the conclusion that there

are additional or unexplained mechanisms at work in the electron emission process. This is a topic of ongoing

research.

Acknowledgements: All electron emission microscopy measurements were performed at the Duke University Free

Electron Laser Laboratory. This research is supported through the Frontier Carbon Technology Project, the Japan

Fine Ceramics Center, ONR and PTS Company-John Driscoll.

REFERENCES

1. R.L. Fink, Z. Tolt, and Z. Yaniv, Surf. Coat.Technol. 108-109, 570 (1998).

2. F. Lacher, C. Wild, D. Behr, P. Koidl, Dia. and Rel. Mat. 6 (1997) 1111-1116.

3. Y. Avigal, A. Hoffman, Dia. and Rel. Mat. 8 (1999) 127-131.

4. B.L. Ward, O. H. Nam, J. D. Hartman, S. L. English, B. L. McCarson, R. Schlesser, Z. Sitar, R. F. Davis,

and R. J. Nemanich, J. Appl. Phys. 84 (1998) 5238-5242.

5. A.T. Sowers, B. L. Ward, S. L. English and R. J. Nemanich, J. Appl. Phys. 86 (1999), 3973.

6. H. Ade, W. Yang, S. L. English, J. D. Hartman, R. F. Davis, and R. J. Nemanich, Surface Rev. and Lett. 5

(1998) 1257-1269.

7. R.J. Nemanich, J. T. Glass, G. Lucovsky, and R. E. Schroder, Raman scattering characterization of carbon

bonding in diamond and diamondlike thin films, J. Vac. Sci. Technol. A 6(3) 1998.

8. M. Park, A. T. Sowers, C. Lizzul Rinne, R. Schlesser, L. Bergman, R. J. Nemanich, Z. Sitar, J. J. Hren and

J. J. Cuomo, J. Vac. Sci Technol. B 17(2) 1999.

NASA/CP--2001-210948 137 August 6, 2001
9. R.E.Schroder andR.J.Nemanich, Analysisofthecompositestructuresindiamondofthinfilmsby

Ramanspectroscopy Phys.Rev.B41,3738-3745 (1990).

10.D.Zhou,A.R.Krauss, L.C.Qin,T.G.McCauley, D.M.Gruen, T.D.Corrigan, R.P.H.Chang, andH.

Gnaser, J.Appl.Phys.82(1997)4546-4550.

11.O.Groning, O.M.Kuttel,P.Groning, L.Schlapbach, J.Vac.Sci.Technol. B17(1999)1970-1986.

NASA/CP--2001-210948 138 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

INVESTIGATION ON FIELD EMISSION OF SI TIPS WITH NANOCRYSTALLINE

DIAMOND COATINGS

W.L.Wang, K.J.Liao C.G.Hu, Z.Li

Department of Applied Physics, Chongqing University, Chongqing 400044, P.R.China

ABSTRACT

The electron field emission from Si tips coated nanocrystalline diamond films was investigated. Nano-diamond films

were deposited on Si tips by hot filament chemical vapor deposition, and Si tips were formed by plasma etching. The

average radius of curvature for the Si tips was about 50 nm. The microstructure of the diamond films was examined by

scanning electron microscopy, Raman spectroscopy etc. The field emission properties of the samples were measured in

an ion-pumped vacuum chamber at the pressure of 10 .8 Torr. The experimental results showed that the nanostructured

films on Si tips exhibited a lower value of the turn-on electric field than those on flat Si substrates as compared with the

diamond films on flat Si. It were found that the tip shape and non-diamond phase in the films have a significant effect on

the field emission properties of the films.

Keywords: field emission; diamond films; CVD; Si tips; nanocrystalline diamond

INTRODUCTION

Recently, carbon nanotubes have attracted much interest because of outstanding electrical properties and their potential

application as a possible material for fabrication of cold cathodes (1 to 4). It was found that the field enhancement from

carbon nanotubes exhibited large local field enhancement and considerable field emission currents at relatively low

applied voltage. Carbon nanotube-based field emission displays have been fabricated using well-aligned nanotubes on

glass substrates. The displays showed a high brightness of 1800cd/m 2 at 3.7V/_tm from the green phosphor, and the

threshold field of 1V/_tm and emission current of 1.5mA at 3V/_tm were observed (5). However, it is difficult for

nanotubes to be vertically aligned after the growth. The arrangement of large area synthesis is still not easily accessible

although vertically aligned carbon nanotubes have been grown by plasma-enhanced chemical vapor deposition (CVD)

recently. To explain the field emission properties of the carbon nanotubes, Obraztsov et al (6) suggested a model for an

energy band diagram of nanostructured graphite-like materials. According to the model, the electron emission from the

nanostructured carbon was a result of tunneling of electrons through a thin modified surface layer with a low value of

efficient work function.

In this paper, the field emission from silicon tips with nanostructured diamond thin layers was investigated. The Si tips

and the nanostructured diamond films were obtained under various conditions in the same reactive chamber rather than

by a complicated photolithography. The orientated growth Si tips coated nanostructured diamond was similar to the

carbon nanotubes, i.e. carbon nanotube-like materials. Different substrate temperatures were employed to study the

effect of film structures on their field emission properties. The mechanism of the field emission from nanostructured

diamond was also analyzed.

EXPERIMENTAL DETAILS

The nanostructured diamond thin films for this study were grown on Si tips using a conventional hot filament

chemical vapor deposition (CVD) apparatus as reported in Ref 7. Mirror-polished n-type silicon wafers were used as

substrates. The substrates were ultrasonically cleaned in acetone and methanol bath for 10 min, and rinsed in 30% HF for

1 min, before loading into the deposition chamber in order to remove the oxide layer of Si surface. After a short rinse in

de-ionized water, the substrate was put on a molybdenum substrate holder. The reactive gases were a mixture of CH4 and

H2. The CH4 concentration in H2 was 1.5-15%, and total gas pressure was 4xl03pa.The gas flow was about 200 sccm.

The chamber was pre-evacuated to 10 -3 Pa. The reactive gases were directed onto the substrates. The gases were

activated by tungsten filament heated at 1800-2200 °C as measured by a pyrometer. The substrate temperature was set

NASA/CP--2001-210948 139 August 6, 2001
between550and850°CasmeasuredbyaPtRhthermocouple mountedonthesubstrateholder.Thefilament wasfirst

heatedinamethane-hydrogen atmosphere tonhancetheformationofatungstencarbidelayercoveringthefilament

surfaceinordertoreducethetungstenevaporation duringthedepositionprocess. Thefilament-to-substrate distancewas

fixedat8mm.

Anegativebiasrelativetothefilamentwasappliedtothesubstrates whenalldepositionconditionswereadjusted

to thedesiredvalues.Biasvoltageandcurrentwere-300Vand250mA,respectively. Thebiascurrentonsetwas

accompanied bytheplasmaformationonthesubstratesurfacewhenbiasvoltageexceeded thethresholdvoltage. Sitips

of several micrometers highwereproducedonSisubstrates byplasmaetchingafter15mindeposition. Thetipsize

dependedonplasmaintensity, andtheplasmaintensitywasincreasedwithincreasingbias(orcurrent).Aftertheplasma

treatment, biasvoltagewasturnedoff,CH4concentration andsubstratetemperature wereadjustedto3.5%and650°C,

respectively.

RESULTS AND DISCUSSION

Figure 1 shows SEM images of the Si tips after 20min of deposition. Fig. 1a and b show the SEM morphology of Si

tips and nanostructured diamond films on Si tips, with an orientated growth, respectively. As can be seen, the films on Si

tips consist of many micrograins without any crystallographic facets. The average radius of curvatures for Si tips is about

20nm, which are similar to the carbon nanotubes, i.e. carbon nanotube-like.

%

Li•N 11

_.__'_:,:,_-._... N

Figure 1. SEM images of the Si tips; (a) Si tips; (b) nanostructured

diamond films on Si tips;(c) diamond films on flat Si substrate.

The field emission characteristics of the films were measured at different substrate temperatures. The experiments

were performed in an ion-pumped vacuum chamber at a pressure of 10 s torr. ITO-coated glass of 3 mm 2 in area was

used as an anode electrode. Anode-to tip separation was about 100gm. Cathodes were heated by a small furnace. Ag and

Au were used as ohmic contacts to anode and cathode, and a resistor of 5M_ was connected to the cathode to protect the

measurement equipment.

Typical emission characteristics from the nanostructured carbon films on Si tips are shown in Figure 2. Fig.2a-c

represents three samples with different annealing temperatures, respectively. In Fig.2c a turn-on field of 1.2V/gm and a

maximum current of 500gA at 2.5V/gm for annealing temperature of 350 °C in nitrogen atmosphere is observed,

whereas, a lower emission current of 160gA at 2.5V/gm with a turn-on electric field of 1.8V/gm without annealing

treatment is shown in Fig.2a. Compared to Fig.2c, the emission current is rather low.

NASA/CP--2001-210948 140 August 6, 2001
40_

2o_

2

r_ ._ (a)

200 400 600 800

Applied Voltage (V)

1000

Figure 2. Field emission characteristics with different

substrate temperatures 25 °C (b) 150 °C (c) 350 °C

Fig.3 shows the comparison in the emission characteristics of the nanostructured diamond films on Si tips with

those on the flat Si substrate. Fig.3a is the field emission from Si tips coated crystalline diamond films without annealing

treatment, and Fig.3b denotes that of nanocrystalline diamond films on the flat Si substrate. It is quite apparent that the

field emission was enhanced by Si tips. The films on the flat Si substrate exhibit a low emission current and high

threshold field.

400

200

=

2

• .

0 200 400 600 800

Applied Voltage (V)

Figure 3. Comparison in the emission characteristics of the nanostructured diamond

on Si tips with those on flat Si substrate. (a) on Si tips; (b) on flat Si

Figure 4 shows Raman spectra of nanostructured diamond films on the Si tips at different annealing temperatures.

The strong sharp peaks at 1580cm -1 (G line) correspond to the high-frequency E2g first-order mode, which are related to

various forms of graphite materials. The peaks at 1345cm -1 (D line) have also been assigned to graphite materials. A

weak peak centered at 1200cm -_ represents nanocrystalline diamond as shown in Fig.4a-c. Fig.4d shows the diamond

films on the flat Si substrate, which revealed that the films were main diamond crystals as evidenced by the sharp

1332 cm-_ peak. It can be seen that the peaks of nanocrystalline diamond were enhanced with annealing temperature.

NASA/CP--2001-210948 141 August 6, 2001
=.

_=

1000 1200 1400 1600 1800

Raman Shift (cm "_)

Figure 4. Raman spectra of nanostructured carbon

films on Si tips at different substrate

temperatures; (a) 25 °C; (b) 150 °C

(c) 350 °C; (d) diamond films on flat Si.

As mentioned above the results, the enhancement in emission properties with the annealing temperature can be

attributed to the changes in the structure of the films on Si tips and barrier height as well as the geometric form of Si tips.

It is well known that diamond is an initial reason for the negative electron affinity on its surface, which can provide a

very low energy threshold for electron escaping into the vacuum from the diamond conduction band. On the other hand,

nanocrystalline diamond has the tip-like shape of its surface morphology. The tip shape of the surface should provide an

additional local increase in electric field at the tip edge. The film composition on Si tips mainly consisted of the graphite

material as shown in Fig.4. The better conductivity of graphite along a based crystallographic plane may lead to the

improvement of the field emission properties as compared with pure diamond films.. The thin graphite layers on Si tips

are similar to carbon nanotubes, i.e. carbon nanotube-like. The Fowler-Nordheim description of the field emission can be

expressed as

J= A _exp(_) (1)

where A-0.014 and B-6.8×109 are constants, J(A/m 2) is current density, F-t_V/d is electric field, d is distance between

anode and cathode, d?is the effective barrier height, and ]3 the field enhancement factor depends on the emitter geometry.

]3 can be estimated by the following expression,

h

]_ = -- (2)

F

where h is the tip height, and r is the tip end radius. Combining eq.(1) with eq.(2), it was found the _ is about

0.30eV-0.80eV for sample c (see Fig.2c) , but _? lies in the range of 1.0eV-2.60eV for sample a as shown in Fig.2a As

can be seen, the energy barrier height with graphite-like structure was decreased by thirty percent as compared with that

ofnano-diamond structure and less than the work function for graphite (5eV).

The decrease in the energy barrier may be due to the modification of electronic configuration of carbon atoms

located at the graphite-like sheet surface for carbon nanotubes. However, the graphite-like thin films on the Si tips are

quite different from that the structure of carbon nanotubes. Here we propose a modified model for electron emission

NASA/CP--2001-210948 142 August 6, 2001
fromnanocrystalline diamondfilmsonSitips.Thegraphite-likefilmsareassumedto bea conductivemetal.A

Si/graphite-like intrfaceispresumedtobeametal/Si schottkyjunction.Asaresultofthemirror-imageforceandtunnel

effect,schottkybarrierheightisreducedwithappliedvoltage.Electronstunnelthroughthejunction,anddrifttothe

surfaceofthenanostructured carbonfilmsandescapeforvacuum. Thewidthofthepotential barrieratthetipvicinity

willbedecreased, furtherincreasingthetunnelingprobabilitythroughit.Ontheotherhand,it canbeseenfromFig.4

thatsampleacontainsahighersp2content inthefilmsthanthatofdiamondfilms.Itwasconfirmedthatheconducting

sp2particlesinthediamondfilmscanincreasethefieldenhancement facor.Therefore, thetotalfieldenhancement

factor_3canbeexpressed intermsoftheproduct ofeachfieldenhancement factoras

t_ = t_g t_ r t_ p (3)

where ]_x, ]_r and ]_p are the field enhancement factor for the tip geometry, tunnel effect and sp 2 content. Taking into

account the graphite work function of 5eV, _3 can be estimated as 3500 using the slope of FN plot. _3gis found to be 120

from eq.(2). Thus _3T_3pis about 29. This means that the geometrical enhancement is greater that of the tunnel and sp 2

content effects.

CONCLUSION

Nanostructured diamond films on Si tips were prepared by hot filament CVD. The Si tips and films can be

obtained under different deposition conditions in the same reactive chamber. It was found that the field emission

properties from nanocrystalline diamond coatings on Si tips were greatly improved as compared with those of the films

on the flat Si substrate. This is due to the tip shape of the surface that provides an additional local increase in electric

field at the tip end, which led to the enhancement of the field emission.

ACKNOWLEDGMENT

The work was financially supported by the National Natural Science Foundation of China under Grant No. 19904016.

REFERENCES

1. Iijima S and Ichihashi T: Single-shell carbon nanotubes ofI nm diameter. Nature, 363, 1993, 603-605.

2. Heer W.A.der: Recent developments in carbon nanotubes. Current Opinion in Solid States & Mater,Sci. 4,

1999, 355-359.

3. Schlesser R, Collazo R, Bower C, et al.: Energy distribution of field emitted electron from carbon nanotubes.

Diamond Relat. Mater. 9, 200, 1201-1204.

4. Wang Q.H., Seelur A.A, Lauerhaas J.M, et al.: A nanotube-based field emission flat panel display. Appl. Phys.

Lett., 72(22), 1998, 2912-2914.

5. Kim J. M, Chai W.B, Lee N.S, et al.: Field emission from carbon nanotubes for displays. Diamond Relat.

Mater. 9, 2000, 1184-1189.

6. Obraztosov A.N, Volkov AP, Pavlovsky I, Field emission from nanostructured carbon materials. Diamond

Relat. Mater. 9, 2000, 1190-1195.

7. Wang W.L, Polo M.C, Sanchez G, et al. Nucleation and initial growth of diamond by biased hot filament

chemical vapor deposition. Appl. Phys. Lett. A65, 1997, 241-249.

8. Wang W.L, Polo M.C. Sanchez G, et al. internal stress and strain in heavily boron-doped diamond films by

microwave plasma and hot filament chemical vapor deposition. J.Appl. Phys. 80(3), 1996, 1840-1850.

NASA/CP--2001-210948 143 August 6, 2001
Proceedings of the SLxth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

PARAMETER ESTIMATION FOR ELECTRON FIELD EMISSION

D.G. Walker and T.S. Fisher

Department of Mechanical Engineering

Vanderbilt University

Box 1592, Station B

Nashville, TN 37235

greg.walker _vanderbilt.edu

fax: (615) 343-6687; ph: (615) 343-6959

J.L. Davidson, W.P. Kang

Department of Electrical and Computer Engineering

Vanderbilt University

Box 99, Station B

Nashville, TN 37235

ABSTRACT

Electron field emission is often described by the basic Fowler-Nordheim model for electron tunneling

currents under an applied field. A convenient expression of this model is a linear fit of the log of measured data

as a function of the inverse field. However, this approach eliminates the possibility of extracting more then

two inherent model parameters simultaneously. Furthermore, Fowler-Nordheim emission from non-planar

surfaces is based on non-physical parameters that offer little insight into the mechanisms of electron field

emission. Using a new model with physically-based parameters to describe the emission process, estimates

can be made from measured data providing more rigorous insight into the physics. The objective of this

work is to evaluate two models used to describe field emission from diamond tips. Model parameters are

extracted through a rigorous application of parameter estimation methods. Conclusions on the nature of

electron field emission are gleaned from examination of the resulting parameters.

Keywords: parameter estimation, field emission

INTRODUCTION

Electron field emission occurs when the potential barrier between a surface and vacuum (or two dissimilar

materials) is thin enough to allow an appreciable probability of electron tunneling. When a wide band gap

semiconductor such as diamond is used as the emitting material, highly energetic electrons are more likely

to tunnel [7]. The energy difference between the emitted electrons and replacement electrons can result in

a net cooling of the cathode. This selective emission of energetic electrons process also suggests that wide

band gap semiconductors can be used in power generation applications as well. At high temperatures, the

potential barrier becomes thinner increasing the probability of tunneling. The thermal energy is thereby

converted to electrical current.

To evaluate the thermodynamics of field emission the electrical current and average energy of the emitted

electrons must be described. Accurate description requires models to include nanoscale effects which govern

the probability of electron emission. Several central f_atures that influence the emission of electrons in

semiconductors are topology of the emitting surface [10], grain structure of the emitting material, work

function of the semiconductor and doping level of the prepared emitter. Although cathodes can be prepared

with a great deal of geometric precision, their effects on the emission process and are not well understood.

Furthermore, the thermodynamics are also governed by several macroscopic phenomena such as temperature,

area available for emission and choice of material. A good model, therefore, should incorporate many of these

f_atures.

Electron field emission is often described using a basic Fowler-Nordheim model for tunneling currents

under an applied field. With this model, whose derivation can be found in the literature [4, 2], the tunneling

NASA/CP--2001-210948 144 August 6, 2001
current (I) varies exponentially with applied field (F).

I 1.56x 101.1 °A/32F2 exp \f"10.4 /exp( -6.44 /

anlp/cnl 2 (1)

The model parameters include emission area (A), the effective work function of the material (g;) and an

"enhancement factor" (/3). All these parameters are interrelated and estimation can prove difficult.

Emission area which is a geometric quantity is typically unknown and will change with the topology

of the sample, applied field and crystalline structure of the cathode. With specifically prepared samples

such as pyramidal "tips," this area can be controlled and is conceivably related to the number of emission

sites (or tips) and the radius of curvature at the end of the tip. However, the area is usually taken to be

the macroscopic surface area available for emission. The effective work function can be approximated from

the elemental composition of the cathode. However, slight variations in this value will occur with surface

contamination and changes in crystalline structure. Therefore, the height of the solid-vacuum barrier is not

consistent for different emission sites.

The enhancement factor represents the amount of band bending created by any number of influences

such as gated structures or geometric surface enhancement. In the Fowler-Nordheim model, the potential

near the barrier is assumed to be linear with a slope of a perfbctly fiat uniform emitter. The enhancement

factor, therefore, represents the change in slope at the barrier due to band bending. Because of the nonphysical nature of this parameter, it has been used to quantify any number of non-geometric effects such

as grain size, and doping level [11] which are not described explicitly in the Fowler-Nordheim model. As a

result, the Fowler-Nordheim model has been relatively successful in providing comparative analysis between

similar samples.

A unique evaluation of all three model parameters in the simplified Fowler-Nordheim model is mathematically difficult because the parameters are highly correlated. Normally, one of the three is assumed to

be known (either emission area A or effective work function g;) and the remaining two can be determined

uniquely. This %ature of the Fowler-Nordheim model becomes apparent when its reduced form is examined.

The Fowler-Nordheim model is conveniently cast as a linear relationship.

in(/) Co(A,/3,0)+C1(/3,_) F (2)

where Co and C1 are functions of the model parameters. It is immediately clear that the three model

parameters can not be determined uniquely fbr any given set of data.

To overcome the limitations of basic Fowler-Nordheim, Fisher [3] suggests using a model which describes

the field enhancement due to band bending through a physically-based geometric radius. We note that other

similar models have been proposed [8, 6, 5]. In this model, the radius of curvature of the emitting tip is used

to describe the field enhancement. Note that for a flat emitter with no surface topology, the tip can be taken

as infinity without loss of generality. Also note that the new model is not restricted by the assumption of a

linear potential at the barrier. The details of the model can be fbund in the literature.

In the present work, parameters from each model are estimated using inverse techniques and experimental

electron-field-emission data. The emission data was collected from a boron-doped diamond cathode with

2#m (at the base) tips. Details of the cathode preparation and experimental procedure can be found in

the literature [11]. Sensitivities of each model to the estimates will be examined to determine confidence

intervals for the estimates. Further, a comparison of non-physical models and parameters to physical model

and parameters will be performed to gain additional insight into the emission process and to assess the utility

of each model.

PARAMETER ESTIMATION

Inverse techniques can loosely be described as approaches for solving ill-posed problems [9]. The problem

at hand can be described as ill-posed because experimental data inherently contain measurement errors.

Therefore, the estimates will also contain errors. However, the error in the estimates can be characterized

through confidence intervals [1]. Inverse techniques are related to curve fits (and actually reduce to curve

fits for polynomial models) but can incorporate any forward model and measurement errors. If we define

NASA/CP--2001-210948 145 August 6, 2001
the forward model as a function of the parameter vector [r//3)], then the estimates can be found from a

minimization of a least squares fit of the experimental data

S [y _ f/(/3)]T_ l[y _ f/(/3)] (s)

where S is the objective to be minimized, Y represents the measurement vector and 6 is the covariance

matrix of measurement errors. For simplicity, the covariance matrix is given a square matrix with relative

measurement errors for each measurement along the diagonal. If the error is constant the matrix can be

defined as the variance times the identity matrix. In this special case, the measurement errors can be

eliminated from the formulation as they do not affect the value of the estimates.

To extract the parameter vector _, the derivative of the objective with respect to the parameter vector

can be set equal to zero. Because the model will in general be non-linear, we use a Taylor series expansion,

neglecting higher order terms to obtain an expression for the parameter vector correction.

/_/+s /_/+ A_; A_ xT¢) s[y_ r/(_)] (4)

xT(_ 1 x

where the tilde represents the fact that the parameter vector is an estimate and the subscript is an iteration

index. The sensitivity matrix represents how the model changes with respect to the parameters, and is

central to determining the parameter extraction process.

x or(#) v(#+_#)-v(#) (5) oH _#

This matrix can be determined analytically in some rare instances; normally, it must be determined numerically. For example, the Fowler-Nordheim model has analytic derivatives in _. However, the derivatives of

the new emission model are determined numerically by perturbing a single parameter (5/3).

RESULTS AND DISCUSSION

The linear dependence of the sensitivity coefficients indicates whether the parameters are correlated.

High correlation (or confounding) means that the parameters can not be estimated independently. It has

already been suggested that the parameters are highly correlated for the Fowler-Nordheim model because

of the ability to cast the formulation into an expression for a line. The line contains at most two independent parameters. Figure la shows how the sensitivities grow with applied field. The actual sensitivity was

g

c

E

0.06

0.04

0.02

0

-0.02

-0.04

-0.06

-0.08

1.5

//

,/

//

,//

/

..... ....,..

Emission Area k,

Field Enhancement Factor "

Work Function

0.01

0

-0.01

-0.02

c -0.03

_N_ -0.04

-o.o5

E

-0.06

-0.07

-0.08

1.5

/

Emission Area

Tip Radius

Work Function

2 2.5 3 3.5 4 4.5 2 2.5 3 3.5

Field (V/gm) Field (V/gm)

(a) Fowler Nordheim (b) New Model

4 4.5

Figure 1: Normalized sensitivity coefficients for the two emission models.

normalized by the nominal value of each parameter. Prom visual inspection, notice that all sensitivity coefficients exhibit possible linear dependent behavior. This fact can be verified further through the identifiability

factor which can be found from the the sensitivity matrix (M IxTxI). An identifiability of zero indicates

NASA/CP--2001-210948 146 August 6, 2001
Table1:Estimatedparameters(with 95%confidenceintervals)andidentifiabilityfortheemissiondatasetandboth models

Parameter NewModel Fowler-Nordheim

Emission Area (#m 2) 2.94x1034_1.20x103 6.81x10n4_1.78x 10 n

Work Function (eV)

Enhancement Factor

Tip Radius (#m)

Identifiability

2.678 4- 5.33 x 10 '_

9.83x 10 _4-3.18X 10 4

2.44 4- 3.79 x 10 '_

905 4- 23

2.6x 10 x9 3.8x 10 27

perfect linear dependence, therefore any value greater than zero indicates a relative ability to extract unique

parameters. Although the identifiably is arguably small in both cases, the new model clearly indicates an

advantage over the Fowler-Nordheim model. These data are reported in table 1.

The sensitivities of the new emission model are found in figure lb. For this model, the parameters are

also relatively correlated and simultaneous estimation is difficult. Mathematically this results in a fairly

flat design space and a global optimization scheme such as a genetic algorithm must be used. Despite the

apparent linear dependence between parameters, significant non-linearity in the design space was found. In

fact, a secondary global minimum representing a possible alternate set of estimates was located. This result

was deemed physically unrealistic so the design space was limited to exclude this region.

In the two models, the emission area and work function represent essentially the same quantity. The tip

radius from the new model and the enhancement factor from Fowler Nordheim, on the other hand, can be

loosely described as the inverse of the other. As the radius of the tip becomes smaller (becomes %harper"),

we expect the emission to increase because the potential barrier at the cathode surface because thinner

allowing more tunneling. As the radius decreases, we expect the field enhancement factor to increase. This

relationship is reflected in the sensitivity coefficients. The signs of the two parameters are opposite.

The validity of the two models can be deduced from the parameters that are similar in each model.

The emission area from each case should be approximately equal. The two models predict slightly different

values. However, the large confidence intervals indicate that this parameter is extremely sensitive to the

measurement errors. Therefore accurate estimation is difficult. Note that the work function estimated from

each model is very similar. Realize that the two models are different, and therefore the parameter estimates

are expected to differ somewhat.

Since the tip radius, which is extracted from the new model, is a physical quantity, we can obtain a

reasonable verification from SEM photographs of the tips. Through visual examination of an SEM image,

the tip of one pyramid was estimated to be 5nm compared to the estimated 9nm. Realize that the image

does not provide a high resolution of the tip and we expect tips to vary to some degree. We believe this is

remarkable agreement between the estimate and the "measurement".

An additional piece of information available from the new model is the number of emitting tips. The

emission area can be assumed to be related to the tip radius by the number of emitters (A Nrcr2). Prom

the estimates, we find that there are N 9.68 tips emitting. This value is not necessarily an integer because

we do not expect all tips to be created equal. On the microscopic level, there will be differences in the

grain structure as well as doping levels etc. Therefbre, this quantity represents some average number of tips

emitting.

The models fitted to the data are shown in figure 2. Although the both match the data reasonably

well fro large fields, we notice that the the new model tends to match better at lower fields. This artifact

is attributed to the fact that the new model does not assume a linear potential profile at the surface like

the Fowler-Nordheim does. This effect will be more important at lower fields where the curvature of the

potential is near the average emission energy.

Finally, the traditional Fowler-Nordheim plot is shown in figure 3 to illustrate some differences between

the two models. Realize that the Fowler-Nordheim model is exactly a straight line. Therefore, the new

model contains some curvature due to the non-linear potential profile. At initial inspection, it may appear

that the Fowler-Nordheim expression matches the data more closely. However, realize that the vertical axis

NASA/CP--2001-210948 147 August 6, 2001
<v

E

o

0.006

0.005

0.004

0.003

0.002

0.001

Measured ' Da.ta '

New Model

Fowler-Nordheim

I

1.5 4 4.5

iI ....-..... k "_...... =_JS:_I- I

2 2.5 3 3.5

Field (V/pm)

Figure 2: Fit of the data with both models using the parameters reported in table 1.

is a log scale. Therefore, at small fields (large 1/F), the difference between the models is negligible. For

moderate fields (as also demonstrated in figure 2), the new model accounts for the identifiable bias in the

data compared to the linear fit.

CONCLUSIONS

Through inverse techniques we were able to extract model parameters from two different electron field

emission models. For both models, it was found that appropriate treatment of the experimental errors in

the estimation process can lead to smaller confidence intervals and more accurate estimates. As identified

by the sensitivity coefficients, it was found that neither model uncorrelates their corresponding parameters,

and estimation is difficult.

Even though the sensitivity coefficients do not identify a clearly prefbrred model for the estimation of

parameters, the new model uncorrelates the parameters slightly (as seen from the identifiability). This is

thought to be because the field in the new model is not assumed to be linear as in the Fowler-Nordheim. Furthermore, a model which employs physically meaningful parameters (tip radius as opposed to enhancement

factor) is generally more desirable.

References

[1] James V. Beck and Kenneth J. Arnold. Parameter Estimation in Engineering and Science. John Wiley

&; Sons, 1977.

[2] Ivor Brodie and Paul Richard Schwoebel. Vacuum microelectronic devices. Proceedings of the IEEE,

82(7):1006 1034, July 1994.

[3] Timothy S. Fisher. Influence of nanoscale geometry on the thermodynamics of electron field emission.

in review Applied Physics Letters, February 2001.

[4] R. H. Fowler and L. W. Nordheim. Electron emission in intense electrical fields. Proceedings of the

Royal Society A, 119:173 181, 1928.

NASA/CP--2001-210948 148 August 6, 2001
U_

m

v

c-

-8

-8.5

-9

-9.5

-10

-10.5

-11

-11.5

-12

-12.5

0.24

;%,,

4"-.'-..

Measured bata +

New Model

Fowler-Nordheim

0.26 0.28 0.3 0.32 0.34 0.36 0.38 0.4

1/F

Figure 3: Fit of the data with both models using the parameters reported in table 1 cast into

traditional Fowler-Nordheim axes.

[5] K. L. Jensen. Semianalytical model of electron source potential barriers. Journal of Vacuum Science

and Technology B, 17:515 519, 1999.

[6] K. L. Jensen and E. G. Zaidman. Field emission form an elliptical boss: Exact versus approximate

treatments. Applied Physics Letters, 63:701 704, 1993.

[7] N. M. Miskovsky and P. H. Cutler. Microelectronics cooling using nottingham effect and internal field

emission in a diamond (wide-band gap material) thin-film device. Applied Physics Letters, 75(14):2147

2149, October 1999.

[8] A. Modinos. Field, Thermionic and Secondary Electron Emission Spectroscopy. Plenum Press, New

York, 1984.

[9] Andrey N. Tikhonov and V. Y. Arsenin. Solutions of Ill-Posed Problems. V. H. Winston & Sons,

Washington, D. C., 1977.

[10] Takao Utsumi. Keynote address vacuum electronics: What's new and exciting. IEEE Transactions on

ELectron Devices, 38(10):2276 2283, October 1991.

[11] Anurat Wisitsora-at, W. P. Kang, J. L. Davidson, and D. V. Kerns. A study of diamond field emission

using micro-patterned monolithic diamond tips with different sp 2 contents. Applied Physics Letters,

71(23):3394 3396, December 1997.

NASA/CP--2001-210948 149 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CONDUCTIVE DIAMOND: A UNIQUE ELECTRODE MATERIAL FORANALYT[CAL

APPLICATIONS

Akira Fujishlma

Department of Applied Chemistry, School of Engineering,

• The University of Tokyo, 7-3-1 Hongo, Bunkyo-ku

Tokyo 113-8656, Japan

ABSTRACT

Electrochemical analysis has assumed an important place in the arsenal of analytical techniques for the detection

of trace amounts of many organic and inorganic molecules of biological significance. Various electroanalytical

techniques such as differential-pulse voltammetry (DPV) and stripping voltammetry have been developed for

electroanalysis. The selection of electrode material is very crucial for the selective analysis of specific group of

compounds. Although wide range of organic species are electroactive at glassy carbon and pyrolytic graphite

electrodes, these electrodes suffer from practical problems such as electrode fouling due to binding of reaction

products and non-electrochemically active biochemical species. Boron-doped diamond, while acting as conducting

electrode, has overcome many of these deactivation problems. First, the wide electrochemical potential window of

the diamond electrode allows the sensitive electroanalytical detection of chemical species that react at relatively high

potentials. Secondly, the low background ettrrent of the diamond electrode has facilitated the detection of several

eleetroactive species with sensitivity at least on_ order of Magnitude greater than that of the GC electrode even at

moderate potentials. Besides, this electrode is its long-term stability due to lack of adsorption of chemical species

on the inert electrode surface. This article gives an overview of recent developments in the eleetroanalytieal use of

diamond electrodes.

Keywords: conductive diamond, electroanalysis, sensor, phenol, and electrode-modification

INTRODUCTION

Diamond possesses various types of superior properties, including high mechanical stability, high thermal

conductivity and chemical inertness. These properties made diamond very promising for electronic device

applications. Recently, conductive boron-doped chemical vapor-deposited diamond thin films have emerged as

unique electrode materials in electrochemistry due to their attractive properties, including very low background

current, wide electrochemical potential window in aqueous media, high resistance to corrosion and mechanical

stability. Boron-doped diamond (BDD) electrodes are expected to have distinct advantages over other conventional

carbon-based electrodes such as glassy carbon (GC) and highly oriented pyrolytic graphite (HOPG) in some

electroanalytieal applications (ref. 1) (Fig. 1). First, the wide electrochemical potential window of the diamond

electrode allows the sensitive electroanalytical detection of chemical species that react at relatively high potentials

(ref. 2). We have recently demonstrated the detection of histamine at a potential of 1.28 V vs. Ag/AgCI, with an

experimental detection limit of 500 nM (S/N = 13) (ref. 2). Another promising feature of this electrode material is

the long-term electrode stability. Although HOPG has been reported to be useful for detection at high potentials, it

fails to exhibit long-term stability due to problems such as intercalation and electrode fouling. Xu et at. have

previously reported such results for azide, which also oxidizes at high potentials (ref. 3). One of the main factors

contributing to the long-term stability of the diamond electrode is the lack of adsorption of most types of chemical

species on the inert electrode surface. We have presented evidence for the inert properties of diamond during the

electroanalysis of NADH and serotonin (refs. 4 and 5). This property has been attributed mainly to the hydrogen

termination of the as-deposited diamond surface, which makes the surface non-polar. It has been shown that there is

a lack of adsorption of polar molecules such as anthraquinone disulfonate on the as-deposited diamond surface, this

behavior eontrasting with that observed at GC electrodes (ref. 6). Reeently, we have observed similar behavior for

the quinone formed during the electrochemical oxidation of serotonin (ref. 5). Interestingly, even oxidatively treated

diamond electrodes did not show any tendency towards adsorption of the quinone, indicating that the oxygen

functional groups are probably of a different type from those present on GC. Thus, both as-deposited and

oxidatively treated diamond electrodes exhibit high stability for electroanalysis (refs. 2 and 4).

NASA/CP--2001-210948 150 August 6, 2001
O'2

L)

/

Diamond__._

Glassy carbon

m_ Pt

2 mAc -

Au

•

• -V

I I I

-1 0 1

Potential (V vs. Ag/AgC1)

I

2

Figure 1. Cyclic voltammograms obtained for Various electrodes in 0.5 M H2804, scan rate, 200 mV/s

Secondly, the low background current (low double-layer capacitance) of the diamond electrode has faeilitatedthe

detection of several electroactive species with sensitivity at least one order of magnitude greater than that of the GC

electrode even at moderate potentials (Refs. 4 and 5). We have achieved amperometric detection limits as low as 10

nM for NADH and serotonin at operating potentials of 0.58 V and 0.43 V vs. Ag/AgC1, respectively (refs. 4 and 5).

We were not able to observe any analytically useful signal at GC for these concentrations, due to large background

current and noise. Although the sensitivity of the diamond electrode is somewhat low at higher operating potentials,

as indicated for histamine (detection limit, 500 nM), the electrode performance is much superior to GC. An

additional advantage with diamond for electroanalysis is that it is relatively insensitive to dissolved oxygen over a

wide range of potentials in both alkaline and acidic aqueous electrolytes (ref. 7). Hence, the electrolyte does not

require any deaeration.

EXPERIMENTAL SECTION

Highly boron-doped diamond thin films were deposited by use of a microwave plasma chemical vapor deposition

system (ASTeX, Woburn, MA), on silicon (100) wafers, as previously reported (ref. 7). Prior to deposition, the Si

substrates were hand-polished with diamond powder (0.5 gin) for nucleation, followed by rinsing with acetone. A

mixture of acetone and methanol (9:1 v/v) was used as the carbon source, with dissolved B203 as the boron source.

High purity (99.99 %) hydrogen gas was used as the carrier gas. The deposition of the film was carried out at a

microwave power of 5 kW. A film thickness of approximately 40 gm was achieved after 10 hours deposition. The

film resistivity, measured by the four-point probe method, was on the order of 10.3 _ cm. The Raman spectra of

these films showed them to be of high quality, as evident from the strong characteristic peak at 1332 em "_. In

addition, a broad peak centered at approximately 1200 cm -t was observed, which is cb.araeteristic of highly borondoped samples. The GC electrode (5 mm diameter, GC-30, Tokai Carbon Co., Ltd., Japan) was pretreated by

polishing with alumina (1 gm), followed by rinsing with high-purity water.

The flow injection analysis (FIA) and liquid ehromatographie experiments were carried out with a miero-LC

pump (BAS, LC-100), an injector (Rheodyne) with a 20-gL loop, a thin layer flow-eelI (BAS), an amperometrie

detector (BAS LC-4C) and an X-Y recorder (Graphtec, WX4000). The wall-jet-type flow-cell consisted of the

Ag/AgC1 referanee electrode and a stainless steel tube as the counter electrode, which also served as the tube for the

solution outlet. The geometric area of the diamond electrode in the cell was estimated to be 0.64 em 2. A CN-3

column (GL Science, 150 × 4.6 mm I. D.; particle size, 5 l/m) and ODS columns were used for the separation of

the pesticides. Phosphate buffer (pH, 2.25)/acetonitrile was used as mobile phase for pesticide detection. A

mixture of acetic aeid/water/aeetoniile was used as the mobile phase for the detection of prehydrolysed pesticides

(phenolic derivatives). All other reagents were analytical reagent grade. Glucose oxidation at metal-modified

diamond was carried out in 0.1 M NaOH solution. All solutions were prepared in Milli-Q water (Millipore).

NASA/CP--2001-210948 151 August 6, 2001
RESULTS AND DISCUSSION

Detection of Carbamate Pesticides

As mentioned above the wide potential window of diamond electrode, as shown in the Figure 1, allows its use

for the detection of molecules, which oxidize at high potentials. Figure 2 shows an example for oxidation of such

molecule, namely carbaryl that is one of the several carbamate pesticides, which is widely used as an insecticide.

The allowed levels of residues of these pesticide compounds in crops and water samples are regulated by

international organizations; for example, the European Drinking Water Directive imposed a limit of 0.1 ng/mL for a

single pesticide and 0.5 ng/mL for the sum of all pesticides in environmental waters. These pesticides are heat

sensitive, decomposing to methylisocyanate and phenol. This tendency to decompose thermally makes it difficult to

analyze many carbamate pesticides directly by gas chromatography. For this reason, numerous HPLC methods have

been developed to analyze carbamates by UV spectrometry or by fluorimetry, and only in some instances are

electrochemical methods used. However, electrochemical methods can offer some advantages as greater sensitivity

and selectivity without the need of various-step derivatization procedures as in fluorescence. However, no suitable

electrode is available for direct detection of these pesticides due to high oxidation potentials. Diamond due to its

wide potential is an excellent choice for such detection.

Figure 2 shows cyclic vottammograms for 100 pM carbaryl together with the corresponding background

voltammograms in phosphate buffer (pH, 7.2) at diamond and glassy carbon electrodes. At the diamond electrode,

an irreversible oxidation voltammetric wave was observed. However, at the glassy carbon electrode, the

voltammogram appears featureless with a broad shoulder along the rising background current, which is due to a

combination of oxygen evolution and electrode corrosion. However, at the diamond electrode, a well defined,

sweep rate-dependent oxidation peak was observed at 1.4 V vs. SCE, the peak potential being slightly more

positive than that for glassy carbon. Due to the low background current, a signal/background (S/B) value of 10 was

obtained for the diamond electrode, indicating the possibility of obtaining high sensitivity at this electrode.

Furthermore the voltammograms were reproducible even after several sweeps, indicating the high stability of the

electrode.

/

I I i

0 0.4

Diamond

5 _A

Glassy carbon

I I I

0.8 1.2 1.6

Potential (V vs SCE)

Figure 2. Cyclic voltammograms for 100 gM carbaryl pesticide at boron-doped diamond electrode and

glassy carbon electrodes in phosphate buffer (pH, 7.2).

The low background current of the as-deposited diamond electrode is mainly attributed to the non-polar nature

of the surface due to hydrogen termination and a depressed density of electronic states at the Fermi level. However,

NASA/CP--2001-210948 152 August 6, 2001
wehaveearlierdemonstrated thathediamondelectrodertainsitshighsensitivityfordopaminedetectionevenafter

severeoxidation in alkaline conditions (ref. 8). In another recent study, we have shown that anodic oxidation of

diamond does not alter its background voltammetrie behavior (ref. 4). Although the presence of oxygen-containing

functional groups on the glassy carbon surface is believed to be partly responsible for its high background current,

surface oxygen does not seem to play a major role in the case of diamond in determining the background current.

We believe that the lack of porosity, even at the sub-nanometer level, rather than the oxygen content, is the main

factor contributing to the tow background current for diamond. Diamond being highly resistive to electrochemical or

chemical oxidative attack does not acquire significant porosity after oxidative treatment. These treatments only

appear to change the surface termination. This is a unique property of the diamond electrode, which allows it to

exhibit very low background current.

Figure 3 demonstrates the utility of diamond as an electrochemical detector for pesticide analysis in high

performance liquid chromatography (HPLC) system. A mixture of five pesticides (100 [.tM each except dichloron,

which is 25 gM) was detected at diamond electrode alter HPLC separation. An experimental detection limit of 200

nM was obtained using this electrode. The electrode exhibited excellent stability with low background currents.

.g

r..)

Met

I

I

ayl benzimidazole

Carbofuran

Time (min)

Figure 3 Chromatogram obtained using diamond detector after HPLC seperation of various carbamate

pesticides (100 gM) each and dichloron pesticide (25 gM).

Detection of Nitrite

We have also shown the analytical utility of diamond electrode for the detection of nitrite and nitrogen

oxides (ref. 9). The detection of nitrite in environmental and biological samples is important. In eleetrochernical

analysis, modified glassy carbon electrodes are usually used to increase the sensitivity and improve the stability. In

case of diamond, due to its inert nature, it does not require any surface modification for the nitrite detection. Figure

4 shows the flow injection analysis data obtained using diamond electrode. The hydrodynamic voltammogram

shown in Fig. 4a shows a peak at 1.15 V vs Ag/AgC1 with the maximum signal to background ratio. Therefore the

flow injection analysis was carried out by holding the electrode potential at 1.2 V vs ag/AgC1 with phosphate buffer

as the mobile phase. Repetitive injection of nitrite stock solutions with various concentrations yielded the data

shown in Fig. 4b. The response for 1 gM was repeatedly tested to check the reproducibility. As shown in Fig. 4b,

the traces are very reproducible with a relative standard deviation of 1.65 %. This figure also shows the response for

50 nM nitrite. The high quality of signals obtained for 50 nM suggests that significantly lower concentrations could

be measured using diamond electrode.

NASA/CP--2001-210948 153 August 6, 2001
300 I_M nitri_

....... __.

._ 200 20

150

_o100

50

o

0.4 0.6 0.8 1 1.2 1.4 1.6 6 ' _ ' 2_ ' _5 ' _ b ' _ ' 3, ' _i '

Potential (V vs Ag/AgCI) Time (min)

Figure 4. (a) S/N variation as a function of the applied potential (data obtained from hydrodynamic

voltammogram) for nitrite detection by flow injection analysis. (b) Electrochemical response for 20 _tL

injections of I gM and 50 nM nitrite in 0.1 M phosphate buffer, at an applied potential of 1.15 V vs

Ag/AgCI.

Detection of Chlorophenol

Chlorophenot (C1P) detection using diamond electrodes is another interesting study. Phenol and

substituted phenols such as chlorinated phenols and related aromatic compounds are known to be widespread as

components of industrial waste. Most of them are carcinogenic and their presence can be harmful for life in general.

Therefore, phenols are considered to be priority pollutants by the US EPA (Environmental Protection Agency).

CIPs also in flue gas from waste incinerators are noted as the precursors of polychlorinated dibenzo-p-dioxins and

dibenzofurans (PCDD/F).

Although organic compounds related to phenols in general can be oxidized at numerous electrode materials, the

oxidation of phenolic compounds at solid electrodes produces phenoxy radicals, which couple to form a passivating

polymeric film on the surface of the electrodes (re£ 9). However, highly boron-doped diamond electrodes have

recently received a great deal of attention, particularly for electroanalysis, owing to their unique electrochemical

properties. In this present study, anodically oxidized diamond surface has been shown to be very stable and

sensitive for ehlorophenol detection. Electrochemical anodic pretreatment is known to change the diamond sin-face

from hydrophobic to hydrophilc by introducing oxygen functional groups on the surface. Earlier we have shown

that anodically pretreated diamond is suitable for selective detection of dopamine in the presence of large amount of

aseorbic acid, as the oxidation potential shifts to more positive potentials relative to dopamine oxidation potential

at the pretreated electrode (ref. 8).

Figure 5a shows the cyclic voltammograms for the first and fifth scans for 5 mM 2,4 dichlorophenol at

anodically oxidized diamond. In this ease the as-grown diamond electrode was treated anodieally at 2.64 V vs. SCE

for 4 rain. The anodieally treated diamond has become hydrophilie fi'om hydrophobic. It is interesting to note that

the voltammetric peak current does not decrease significantly after 5thscan even at high chlorophenol concentration.

This result is in contrast to that at as-grown diamond or glassy carbon where the electrode fouled just after second

scan. The unique stability of diamond electrode is not clear yet, however, the polar oxygen functional groups

generate by oxidation probably repel the phenoxy radicals generated during phenol oxidation.

Figure 5b shows the FIA results for 100 repetitive injections of 5 mM 2,4 dichlorophenol at glassy carbon

elecla'ode and anodized diamond electrode. The detection potential was 1.2 V for glassy carbon while at anodized

diamond 1.4 V was chosen as an optimum faradaic response from hydrodynamic voltammetry (not shown). A

highly stable detection peak is observed when anodized diamond is used, the relative standard deviation (R.S.D.) of

the peak heights is 2.3 % and the decrement is 10 %. In contrast, repetitive injections of coneentrated 2,4

dichlorophenol at glassy carbon results in a 70 % reduction of the peak height. The value of R.S.D. in the peak

height ranges from 30 % to 40 %. After this durability experiments, we could be clearly visible to the naked eye the

passivating layer on the surface of glassy carbon. These results clearly demonstrate the outstanding stability of

diamond electrode for ehlorophenol detection.

NASA/CP--2001-210948 154 August 6, 2001
_, 1 . ........ I'siscah'- /_-

'_ 0.8

_0.6

;_ 0.4 5th scan\ _,/_rf

"_ 0.20 _°

_-0.4 I, .... _ ....................

_) -0.5 0 0.5 1 1.5 2

Potential (V vs SCE)

o

_100

o

880

6o

o 40

o

.;_

z0

o

0 '

o

Diamond _

_qe_= o Glassy carbon

I . , I , , , I , , , I , , ,

20 40 60 80 100

Injection number

Figure 5 cyclic voltammograms for the first and fifth scans for 5 mM 2,4 dichlorophenol at anodically

oxidized diamond.

Metal-modified Diamond for Glucose Detection

Diamond electrodes also act as inert supports for metal catalysts. We have succeeded in modifying diamond

electrode with metal particles such as Cu and Ni in view of its applications to glucose (carbohydrate) detection.

Electrochemical detection of carbohydrates is very attractive due to the possibility of high sensitivity and wide

dynamic range. Metal electrodes such as nickel and copper are known to oxidize carbohydrates in alkaline solution.

The advantage of Ni and Cu electrodes is that they produce quite stable responses. These electrodes have been

widely used in liquid chromatography and capillary electrophoresis. However, dispersion of metallic particles

within an organic polymer or simply on an inert surface results in a drastic increase in the catalytic activity and

sensitivity of the electrode (ref. 10). A stable, inert electrode with low background current would be the best choice

for the deposition of metal electroeatalysts. Glassy carbon (GC) electrodes lack some of these requirements, with

additional problems such as weak adherence of the metal particles to the substrate. Conductive diamond films

appear to be well suited to overcome these problems (ref. 11).

Ni-diamond was prepared by depositing 100 I.tL of 10 mM Ni(NO3)2 solution on the surface of the

diamond, drying at 30 °C for more than 8 hours and washing with purified water. The electrode was

electrochemically conditioned in 0.2 M NaOH by cycling the potential between 0.0 and +0.8 V (vs SCE) until no

further increase in the anodic peak current at +0.48 V vs SCE was observed. The significance of this peak is

discussed below. Cu-diamondwas prepared by single step potentiostatic deposition of Cu on the diamond surface in

50 mM H2SO4 solution containing 1 mM CuSO, The deposition was carried out at -0.12 V vs SCE for 22 mill.

Prior to copper modification, the diamond surface was pretreated electrochemically at +2.0 V (vs SCE) for 1 hour in

0.1 M NaOH solution. Such an oxidative pretreatment was necessary to achieve strong adherence of Cu particles.

Oxidative treatment is known to improve the adherence of the diamond surface by making it to hydrophilie. In the

case of Ni-diamond, this treatment showed no effect. The amount of metal deposited on the diamond surface was

estimated to be 25 Ilg/em 2 for Ni and 30 I.tg/em 2for Cu.

Figure 6a shows a linear sweep voltammogram obtained for a bare diamond electrode in 0.2 M NaOH

solution containing 1 mM glucose. No Faradaic response was observed within the potential window. Furthermore,

the background current was very low, as mentioned previously. However, Ni-diamond produced a peak-shaped

voltammogram even in the absence of glucose, as shown in Figure 6a. As previously reported, the anodie and

cathodic peaks at +0.48 and +0.36 V vs SCE, respectively, earl be attributed to the Ni (II)/Ni (III) couple. A

significant increase in the anodic peak current at +0.48 V is observed in the presence of 3 mM glucose. The

voltammograms obtained in the presence of glucose were very reproducible. The voltamrnetrie response of Cudiamond is very similar to that of Ni-diamond, except that the oxidation occurs at 0.58 V vs SCE.

NASA/CP--2001-210948 155 August 6, 2001
_9

20

0.2

A

a

,-,2.

v

El.

B

I I r " " "

0.3 o.4 0.5 0.6 o io 4"0 60

Potential (V vs. SCE) Time (min)

Figure 6 (A) Cyclic voltammograms for 3 mM glucose at bare (a) and Ni modified (c) diamond electrodes in

0.2 M NaOH. Curve b is voltammogram for blank solution for Ni-diamond electrode. (B) Amperometric

response of a Ni-diamond for repetitive injections of I mM glucose in flow injection analysis system at 0.45

V vs Ag/AgCI.

Figure 6B shows the amperometric response of Ni-diamond for 20-gL injections of 1 mM glucose, with

30 mM NaOH as the mobile phase. A highly reproducible response, with a peak variability less than 3% (n =11)

was observed. The operational potential of +0.45 V vs Ag/AgCI was selected from the hydrodynamic

voltammogram for these measurements. In the ease of Cu-diamond the operational potential was +0.625 V. The

background current for Ni-diamond in Figure 6 is as low as 100 nM, which enables the detection of glucose to be

carded out with high sensitivity. This value is at least one order of magnitude lower than that for the bulk Ni

electrode, with the response for glucose also being higher for the modified diamond electrode. Both Ni-diamond and

Cu-diamond showed excellent stability for one week with regular use. Experiments to examine long-term stability

are in progress. Under similar conditions, metal-modified glassy carbon electrodes were stable only for one day due

to adherence problems. Thus, diamond appears to provide a highly suitable substrate for metal modification.

However, it should be noted that highly alkaline conditions (pH _>12) are necessary to achieve high sensitivity.

REFERENCES

1 Rao, T.N., Fujishima, A.: Recent advances in electrochemistry of diamond. Diamond and Relat. Mater., 9,

2000. 384.

2 Sarada, B.V., et al.: Electrochemical oxidation of histamine and serotonin at highly boron-doped diamond

electrodes. Anal. Chem., 72, 2000, 1632.

3 Xu,J.Z., Swain, G.M: Oxidation of azide anion at boron-doped diamond thin-film electrodes, Anal. Chem.,

70, 1998, 1502

4 Rao, T.N., et al.: Electrochemical oxidation of NADH at highly boron-doped diamond electrodes. Anal.

Chem., 71, 1999, 2506.

5 Sarada, B.V., et al.: Electrochemical detection of serotoninat diamond electrode. Chem. Lea., 1999, 1213.

6 Xu, J.S., Chen, Q.Y., Swain, G.M.: Anthraqunionedisulfonate electrochemistry: Acomparison of glassy

carbon, hydrogenated glassy carbon, highly oriented pyrolytic graphite, and diamond electrodes, Anal. Chem.

70, 1998, 3146.

7 Yano, T., et al.: Electrochemical behavior of highly conductive boron-doped diamond electrodes for oxygen

reduction in alkaline solution. J. Eleetroehem. Sot., 145, 1998, 1870.

8 Popa, E., et al.: Selective electrochemical detection of dopamine in the presence of ascorbie acid at anodized

diamond thin film electrodes, Electrochem. Solid St. Lea., 2, 1999, 49.

9 Spataru, N., et al.: Determination of nitrite and nitrogen oxides by anodic voltammetry at conductive diamond

electrodes. J. Eleetroehem. Soc., t48, 2001, E112.

NASA/CP--2001-210948 156 August 6, 2001
10 Rodgers, J.D.,Jedral,W.,Bunce, N.J.: Electrochemical oxidation of chlorinated phenols. Environ. Sci.

Technol., 33, 1999, 1453.

11 Prabhu, S.V., Baldwin, R.P.: Eleetroeatalysis and detection of amino sugars, alditols, and acidic sugars at a

copper-containing chemically modified electrode.Anal. Chem., 61, 1989, 2258.

12 Uchikado, R., et al.: Metal-modified diamond electrode as an electrochemical detector for glucose. Chem.

Lett., 2001, 144.

NASA/CP--2001-210948 157 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ADVANCED ELECTROCHEMISTRY WITH DIACHEM ® ELECTRODES

L. Schiller, M. Fryda, D. Herrmann, I. Triister

Fraunhofer Institute for Surface Engineering and Thin Films, IST

Bienroder Weg 54 E, 38108 Braunschweig, Germany

E-mail: schaefer@ist._hg.de,: Fax: ++49-(0)531-2155-901, Phone: ++49-(0)531-2155-520

W. Hiinni, A. Perret

Centre Suisse d'Electronique et de Microtechnique SA CSEM

Jaquet-Droz 1, CH-2007 Neuchfitel, Switzerland

ABSTRACT

Boron-doped conductive diamond films applied as electrode material exhibit unique properties in

electrochemical processes. Beside the inherent extreme chemical stability of diamond, even at high doping levels, in

particular the high overpotential for water decomposition opens new opportunities for improved and also new

electrochemical processes. To investigate conductive CVD diamond for electrochemical processes under laboratory

as well as industrial conditions we developed large-area conductive DiaChem ® deposition on a variety of electrode

substrate materials. DiaChem ®electrodes with different geometries and dimensions have been fabricated using hotfilament diamond CVD with deposition areas up to (40 x 50) cm 2. Electrical resistances are in the range of (5-100)

m_cm using in-situ doping with diborane or trimethylboron, respectively.

The extreme chemical stability of DiaChem ® electrodes is demonstrated by high current density testing. In

these experiments the electrodes are loaded with increasing current densities up to 10 A/cm 2 in sulfiaric acid over

several months without degradation of the electrode performance. Electrochemical cells with DiaChem ® electrodes

have been used for treatment of industrial wastewater, synthesis of inorganic chemical products and in electroplating

processes. In all tests conductive diamond exhibits superior effectiveness and energy efficiency compared to

conventional electrode materials. The results are promising for the transfer of DiaChem ® electrodes into the market.

Keywords: large-area hot-filament CVD, diamond electrodes, electrochemistry, applications

INTRODUCTION

Diamond is known for its extreme chemical stability. The considerable increase of availability of diamond due

to the activated chemical vapour deposition (CVD) techniques for polycrystalline films (ref. 1) thus made CVD

diamond attractive for applications in electrochemical processes with severe demands on electrode stability. To

apply CVD diamond as an electrode material for electrochemical processes, it has to be doped to achieve the

necessary electrical conductivity. In the diamond CVD process sufficient conductivity can easily be realised by insitu doping with a boron precursor in the activated gas phase. Despite high doping levels of up to a few thousand

ppm the diamond coatings retain the extreme chemical stability ofundoped CVD diamond. It has been

demonstrated, that diamond coated electrodes show no significant corrosion even under high electrochemical load

(refs. 2 and 3). Also, no degradation of the electrochemical performance could be observed for operation in solutions

of NaF and concentrated nitric (ref. 3) or concentrated sulphuric acid (ref. 4), respectively. Even loading with high

• . -2

current densmes of 2 to 10 A cm for several months, the electrochemical activity of these diamond electrodes

remains unaffected in contrast to conventionally used carbon electrodes (ref. 5).

Besides the extreme electrochemical stability the extraordinary high overpotentials for water electrolysis (ref. 6)

is the most striking property of diamond electrodes. It opens new opportunities for electrochemical processes in

water treatment, electrochemical synthesis and electroplating. However, in order to exploit these opportunities,

conductive diamond coatings have to be available for large-area electrodes and also on various geometries. We

therefore developed large-area hot-filament chemical vapour deposition (CVD) of doped polycrystalline diamond

• • 2 .....

films on planar deposmon areas of up to 0.5 m and on flexible geomemes for coating metalhc and ceramic

• ® ....

substrates to produce DmChem electrodes. These have been investigated on laboratory scale using specific model

systems and in electrochemical test systems on industrial sides.

NASA/CP--2001-210948 158 August 6, 2001
PRODUCTIONOF DIACHEM ® ELECTRODES

The technology for the production of DiaChem ®electrodes has been developed in a co-operation between

Fraunhofer IST, Germany, and CSEM, Switzerland. DiaChem ® electrodes with reproducible electrochemical

properties are fabricated on metals (e.g. Nb, Ta, Ti, Mo, W, Zr), silicon, ceramic and graphite base materials by a

large-area hot-filament chemical vapour (HFCVD) process (see figure 1)

 ::!iii ii i;ii' i i i iiiiiiiiii ili iiiiiiiiiiii

i::::::::i:! ...... :

Figure 1: Hot-filament diamond CVD system for fabrication of large-area DiaChem ® electrodes (left) and

selection of DiaChem ® electrodes demonstrating the variety of forms and dimensions used.

Technical electrode materials are used for the preparation of metal-based DiaChem ® electrodes. The electrodes

are preconditioned by sand blasting for increasing surface roughness. The increased roughness promotes diamond

adhesion and increases the electrochemical active surface area. Awet cleaning process is used to remove debris from

the sand blasting process. Subsequent ultrasonic seeding is performed in a suspension of diamond powder in

alcohol. After seeding the metal substrates are cleaned with ultra pure water before diamond deposition. In the case of

Ti-based electrodes, which are very common in industrial electrochemistry, additionally a conducting carbide layer

at the interface is formed prior to diamond deposition to avoid embrittlement and phase transformation by hydrogen

diffusion and incorporation during the diamond deposition process (ref. 7).

Because of porosity graphite substrates used for electrodes have a high surface area, which interacts with atomic

hydrogen in the activated gas phase. The resulting graphite etching might deteriorate the substrate surface and also

contribute to the carbon content in the gas phase in an undefined way. Therefore, graphite electrode substrates were

electroplated with a thin (10 50 nm) Au or Pt interlayer prior to seeding and cleaning like the metallic substrates.

The technical metal as well as graphite electrode substrates were coated with 1 10 gm thick conductive diamond

films.

For silicon-based diamond electrodes the preconditioning and diamond deposition differ from that of the

technical metal and graphite substrates. No surface roughening is performed and an etching step with diluted HF

solution is added. The smooth Si surfaces of the conductive silicon wafers (0.1 _cm and 1 m_cm, Siltronix) used,

are coated with diamond films in the thickness range of 0.1 gm to 1.5 gin. The Si-based DiaChem ® electrodes are

preferable for sensor applications and in electrochemical cells with electrode stacks which are sensitive to the distance

between the electrodes.

The diamond deposition process is consisting of the controlled steps of heating, diamond deposition and

cooling. Typical deposition parameters for the fabrication of DiaChem ® electrodes are:

• filament temperature (2200 °C 2600 °C measured with an optical pyrometer),

• substrate temperature (700 °C 925 °C measured by a thermocouple in a dummy substrate)

• total pressure (10 50 mbar),

• gas phase composition (0.5% - 2.5% cn4 and 10 200 ppm diborane in H2)

Typical growth rates for the boron-doped diamond coatings are between 0.2 gm/h and 1.2 gm/h. Differences in

growth rates are predominantly influenced by the substrate temperatures, which differ depending on substrate

materials. After diamond growth the DiaChem ® electrodes were cooled to room temperature under pure hydrogen gas

flow. All deposition-related steps are performed with a programmable control unit for automatic process control to

ensure reproducibility. Examples for DiaChem ® electrode types are presented in figure 1, demonstrating the variety

of electrodes used for electrochemical testing.

NASA/CP--2001-210948 159 August 6, 2001
Forthetechnical electrodesubstrates thesurfaceroughness ofseveral gmrepresents theroughening

preconditioning. Thisroughness iuperimposed bythediamondfilmroughness whichdependsonthefilm

thicknessandtheboronconcentration inthefilm.Withincreasingborondopinglevelsmallerdiamondgrainsizes

andhighergraindensitiesareobserved. Forhighborondopinglevelsthediamondgrainsizeisbelow1gmeven

forseveral gmfilmthicknessduetosecondary nucleationpromotedbyboronincorporation.

Secondary ionmasspectrometry (SIMS)characterisation revealsnegligibleamounts(below1ppm)offilament

material inthediamondcoatings. Additional hydrogenincorporation ofabout1%istypicalforCVDdiamond

films.Boronconcentrations inthefilmarebetween500ppm 8000ppmresultinginelectrical resistances between

0.1_cmand0.05_cm.Theborondopingaswellasimpuritylevelsareconstant throughthewholediamondfilm,

thusdemonstratingheaccuracyofthedepositionviaautomatedprocesscontrol.Ramanspectraaredominatedby

thesharpdiamondpeakintherangefrom1332cm-1to1340cm-1.NoRamanfeaturesofnon-diamond carbonithe

regionof1450cm-1to1600cm-1arevisible.Forfine-grained coatingsathighdopinglevelsaweakpeakfor

submicrondiamondgrainsat1140cm-1isobserved. Additionallyabroadpeakatapproximately 1200cm-1becomes

visible.Accordingtoliterature(see.g.refs.8to9)thispeakiscausedbyincreasingdisorderinthediamond

crystallitestructureduetotheincorporation ofboron.

ELECTROCHEMICAL PROPERTIES OF DIAMOND ELECTRODES

The electrochemical behaviour of diamond electrodes is influenced by the type and concentration of the dopant,

the primary crystallographic orientation, the presence of non-diamond amorphous carbon phases and the surface

termination (H or O) (ref. 10). Cyclic voltammograms are used as a standard electrochemical characterisation for

diamond electrodes. These voltammograms are performed in a conventional EG&G Parc three electrode flat cell

using a ZAHNER electric IM5d potentiostat, where the diamond electrode was used as working electrode, Pt or Zr

as counter andAg/AgC1 (sat. NaC1) as reference electrode. In 0.5 M H2SO4 metal-based DiaChem ® electrodes

exhibit electrochemical overpotentials of up to 2.8 V for oxygen generation (anodic operation modus) and about

.... ®

1.0 V for hydrogen evolution (cathodic operation modus). Thus DlaChem electrodes are able to produce strong

oxidants like ozone (oxidation potential 2.07 V) and even hydroxyl radicals (oxidation potential 2.80 V) without

unacceptable energy loss by the generation of oxygen. However, cyclic voltammogram investigations showed, that

for anodic operation a preconditioning by anodic polarisation in 1 M H2SO4 with current densities of 50 mA/cm 2 for

at least 30 min is necessary to ensure stable and reproducible operation (ref. 11).

The extreme electrochemical stability of preconditioned electrodes is demonstrated by operation in 1 M H2SO4

-2 ..... ®

at 10 A cm for more than 750 h. During these severe operation conditions, Nb-based DlaChem electrodes showed

no degradation in their electrochemical performance for several weeks or thousands ofAh cm -2.

ELECTROCHEMISTRY WITH DIACHEM ® ELECTRODES

DiaChem ® electrodes have been investigated applying well defined electrochemical model systems in the

laboratory and in industrial testing at end-users to identify relevant fields for technology transfer to industrial

applications. Promising areas are water treatment, electrochemical synthesis and electroplating. In the sections

below, a selected representative example for each field of application is discussed.

Water treatment

The preservation of resources, especially water, is one of the major environmental issues for furore technological

development. Biological treatment of water for the removal of organic load - characterised by COD (chemical oxygen

demand) - is well established and relatively cheap. However, biological treatment techniques are slow and not

suitable for all pollutants. For polluted waters, which are difficult to treat with biological methods, Advanced

Oxidation Processes (AOP) have been developed (ref. 12). The basis of AOP is the total mineralisation and

decomposition of organic pollutants by hydroxyl radicals with the highest oxidation potential for an effective

oxidation process. Economical and efficient oxidation rates for industrial water treatment, however, require

appropriate anode materials. As the electrochemical overpotential of DiaChem ® anodes and the chemical stability are

higher than for any other electrode material, the production of hydroxyl radicals is extremely energy efficient and

reliable. The high current efficiency and the effectiveness with respect to the kind and amount of organic pollutant

has been demonstrated for a variety ofwastewaters from different industries. One representative example for this

electrochemical advanced oxidation process (EAOP) with DiaChem ® electrodes is shown in figure 2.

NASA/CP--2001-210948 160 August 6, 2001
8000 100

7000 .%,, .'. ".' ".. ', . . -. ,: . .

6000 ._

:.'. ",'.'.:::::'5"-'.'.:.'-..:.'," ".'-:.:. ,_

0_5°°°_ " 60 ,&

E 4000

::_: :::::::::::::::::::::::::::::::::::::: 40

oooo 2O00

1000 _ 20

0 ' , i , I , ' ' ' ' _ 0

10 20 30 40 50

Q I Ahll

Figure 2: Reduction of the organic load, represented by the chemical oxygen demand (COD), of a dye

• • ®

containing wastewater applying an electrochemical advanced oxidation process with DlaChem electrodes•

For the investigation of EAOP water treatment on a larger scale, a pilot system with a monopolar flow cell has

been constructed. The system includes a stack of diamond grid electrodes at a distance of 4 mm and an active

diamond area of 1 m 2. First investigations on 500 1 ofwastewater from the automotive industry have been

performed. The system was operated at flow rates from 500 1/h to 2400 1/h. The EAOP yields a maximum average

current efficiency of 65% for the COD reduction from 3200 mg O2/1 to 400 mg O2/1 at a flow rate of 2000 1/h. During

several weeks of oxidation the properties of the DiaChem ®electrodes have not been affected.

Besides the high energy efficiency and effectiveness there are other features of EAOP for water treatment, which

have to be considered. Even for DiaChem ® electrodes a sufficient production of hydroxyl radicals occurs only under

conditions of simultaneous oxygen evolution, resulting in effective, but non selective oxidation. Furthermore, the

extreme high oxidation potential of hydroxyl radicals yields unspecific side reactions for example formation of

nitrate, chlorate or AOX, which have to be considered in the construction of EAOP systems for water treatment. A

decrease of current efficiency for low COD values due to mass transfer limitations is often observed. This effect may

be solved by the optimisation of the cell design in cases where a complete destruction of the pollutant is necessary.

Electrochemical synthesis

• • • ®

Because of the high overpotential for water decomposition DiaChem electrodes, operated in the anodic mode,

are very suitable for the electrochemical synthesis of oxidants like ozone and peroxide compounds. These oxidants

are widely used as initial products in chemical industry. In figure 3 a comparison of the economic performance

• • ® • • •

between conventionally used Pt electrodes and DiaChem electrodes is shown. For the production of ammonium

.... ® •

peroxodisulfate, performed in a test system at an end-user, the efficiency of DiaChem electrodes with respect to

o

production rate per electrode area is increased by roughly 25 _ for high current densities. Furthermore, in contrast to

• ® .....

Pt electrodes, DiaChem electrodes can be operated at lower current densities. This operation results in lower energy

consumption levels and even higher production efficiency. This considerable increase in electrode performance will

reduce production costs and also may contribute to the environmentally beneficial energy saving.

Additional advantages of DiaChem ®electrodes have been observed for the production ofperoxodisulphuric acid

from sulphuric acid (ref. 13). Compared to conventional Pt electrodes the production process can be operated with

much lower concentrations of sulphuric acid at acceptable efficiencies. This will reduce handling cost for sulphuric

acid on production level. Also, the purity of the initial products is higher in the case of DiaChem ® electrodes due to

the reduction of by-products by avoiding additives and reducing the oxygen production. DiaChem ® electrodes are

now investigated for long-term performance in technical systems at different industrial sides for the production of

persulfate, perchlorate and periodate, respectively.

NASA/CP--2001-210948 161 August 6, 2001
120

110

_1oo

g 90

8o

u_ 70

LU

6O

5O

4O

--Platinum

with 494 mA/cm 2

Diamond

with 494 mA/cm 2

Diamond

with 156 mA/cm 2

i i i i

0 1 2 3 4

Es/(kWh/kg)

Figure 3: Production rates per electrode area (efficiency EFZA) for different energy consumption levels per

• ® . .

produced mass (Es) compared for DlaChem electrodes (Diamond) and Platinum electrodes•

Electroplating

In a variety of processes DiaChem ® electrodes have been tested against conventionally used lead or mixed metal

oxide (MMO) electrodes in the production environment of electroplating industry. DiaChem ®electrodes have been

VI.

used for example for the oxidation of Cr III to Cr in chromic sulphuric acid applied for etching ABS plastics.

Compared to MMO electrodes an increase of efficiency with respect to the electrode area by a factor of about 5 at a

simultaneous reduction of processing time by a factor of 2.5 is observed (figure 4). This considerable improvement

of the economics is especially interesting when considering the increased use of ABS plastics.

100

80

6O

4O

2O

0

z conventionalelectrode

; diamondelectrode

_ t

_ t

:-:-:-:-:. :-:-:-:-:. [

::::

i

electrode area processingtime

...... ]]]

Figure 4: Comparison of electrode area efficiency and processing time for the oxidation of Cr to Cr vI in

.... ®

chromic sulphurlc acid for DlaChem electrodes and conventional MMO electrodes•

DiaChem ®oxidation anodes are also investigated in combination with MMO working electrodes for the

development of lead-free chromium electroplating processes In conventional electroplating chromium oxidation is

achieved by introducing ecological critical Pb into the process An additional aspect is the recycling of electroplating

baths where the decomposition of organic or cyanide additives is required DiaChem ®electrodes are highly efficient

for destruction of cyanides in different wastewater Cyanide has also been decomposed when bound in a complex

molecules, as demonstrated by the complete cyanide reduction in solutions containing K4[Fe(CN)6]

CONCLUSIONS

The development of large-area hot-filament activated CVD processes with in-situ boron doping via gas phase

precursors and automated process control resulted in reproducible deposition of conductive diamond films on areas of

2 ........

up to 0.5 m and also on three dimensional shapes. The reliability of these coatings on conventional electrochemical

electrode materials is sufficient to transfer the CVD diamond technology into electrochemical industry. The

NASA/CP--2001-210948 162 August 6, 2001
• ®

electrochemical performance o f DlaChem electrodes is highly promising with respect to effectiveness and energy

efficiency. Laboratory as well as industrial testing revealed electrochemical wastewater treatment, electrochemical

..... ® . . .

synthesis and electroplatlng as application fields for DlaChem electrodes. To exploit the opportunities of the new

• ® .... ® ....

DlaChem electrodes in these fields, industrial DlaChem -based systems for the various specific application

requirements have to be developed in co-operation between electrode producers, system manufactures and end-users.

ACKNOWLEDGEMENTS

• • ®

We gratefully acknowledge electrochemical DlaChem investigations and discussions by Prof. Comninellis and

his co-workers, EPFL Lausanne, and by Dr. M. Blaschke, M. Stadelmann and Dr. A. Kraft G.E.R.U.S. mbH,

Berlin. We also like thank Dr. T. Lehmann, Infracor Hanau, Dr. J. Wurm, METAKEM GmbH Usingen, and our

co-operation partners Dr. W. Haenni and Dr. A. Perret at CSEM, Neuchatel, for fruitful discussions. For SIMS

measurement we thank Dr. Peter Willich at Fraunhofer IST, Braunschweig.

REFERENCES

1. Dischlerw, B., Wild, C.: Low-Pressure Synthetic Diamond, Springer Series in Materials Processing,

Springer Verlag, Berlin Heidelberg, 1998

2. Perret, A., et al.: Diamond electrodes and microelectrodes, Electrochemical Society Proceedings, vol. 97-

32, 1997, 275-283

3. Swain, G..M.: The susceptibility to surface corrosion in acidic fluoride media: a comparison of diamond,

HOPG, and glassy carbon electrodes, J. Electrochem. Soc., vol. 141, 1994, 3382-3393

4. Fryda, M., et al.: Diamantschichten fiir die chemische Industrie, Nachr. Chem. Tech. Lab., vol. 45, 1997,

379-382

5. Swain, G.M.: The use of CVD diamond thin films in electrochemical systems, Adv. Mater., vol. 6, 1994,

388-392

6. Martin, H.B., et al.: Hydrogen and oxygen evolution on boron-doped diamond electrodes, J. Electrochem.

Soc., vol. 143, 1996, L133-L136

7. Heinrich, G.., et al.: The influence of diamond chemical vapour deposition coating parameters on the

microstructure and properties of titanium substrates, Diamond Relat. Mater., vol. 5, 1996, 304-307

8. Ager III, J. W., et al.: Fano interference of the Raman phonon in heavily boron-doped diamond films

grown by chemical vapour deposition, Appl. Phys. Lett., vol. 66, 1995, 616-618

9. Yano, T., et al.: Electrochemical behaviour of highly conductive boron-doped diamond electrodes for

oxygen reduction in alkaline solution, J. Electrochem. Soc., vol. 145, 1998,1870-1876

10. Granger,M. C., Swain, G. M.: The influence of surface interactions on the reversibility of ferri/ferrocyanide

at boron-doped diamond thin film electrodes, J. Electrochem. Soc., vol. 146, 1999, 4551-4558

11. Fryda, M., et al.: Wastewater treatment with diamond electrodes, Electrochemical Society Proceedings,

vol. 99-32, 1999, 473-483

12. Andreozzi, R., et al.: Advanced oxidation processes (AOP) for water purification and recovery, Catalysis

Today, vol. 53, 1999, 51-59

13. Michaud, P. _., et al.: Preparation of peroxodisulphuric acid using boron-doped diamond thin film

electrodes, Electrochemical and Solid-State Letters, vol. 3, 2000, 77-79

NASA/CP--2001-210948 163 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ELECTROCHEMICAL PROPERTIES OF SULFUR-TREATED DIAMOND

Sally C. Eaton,* Alfred B. Anderson, John C. Angus

Case Western Reserve University

Cleveland, OH 44106 USA

Yulia E. Evstefeeva, Yuri V. Pleskov

Frumkin Institute of Electrochemistry

Moscow, 117071 Russia

ABSTRACT

The electrochemical properties of sulfur-treated diamond were explored. Growth of diamond with H2S present

in the source gases and treatment of diamond in the absence of growth with a sulfur-containing plasma were

performed. In both cases, Mott-Schottky analysis showed evidence of donor centers. Measurements of open-circuit

potentials in the presence of UV irradiation and measurements of the thermoelectric effect also showed n-type

conductivity. The presence of sulfur was confirmed by particle induced x-ray emission (PIXE). The donor activity

of the sulfur may arise from surface states or grain boundaries rather than sulfur incorporated substitutionally into

the bulk of the diamond. Preliminary evidence indicates that boron may facilitate the attachment of sulfur onto the

growing diamond surface and hence aid its incorporation into diamond.

Keywords: Diamond films; n-type conductivity; Sulfur-doping; Mott-Schottky plots

INTRODUCTION

Boron has long been used as a p-type dopant in diamond, but a viable n-type dopant remains elusive. For ntype conductivity, several substitutional impurities have been investigated. Nitrogen is a deep thermal donor at 1.6

eV below the conduction band; phosphorus has a donor level about 0.6 eV below the conduction band and a

mobility of 100 cm2/V-s (ref. 1). Recently, sulfur has been reported to give n-type conductivity (refs. 2,3).

However, other measurements have indicated that the samples contained boron and were p-type (ref. 4). The

experiments described here were motivated by the possibility of co-doping diamond with boron and sulfur to

produce n-type conductivity.

EXPERIMENTAL

Growth experiments

In this work, diamond films were grown in an ASTeX microwave reactor using H2S as the sulfur source. The

methane concentration ranged from 0.1 - 0.4%; the S/C atomic ratio from 15 to 20,000 ppm. In some experiments,

co-doping of sulfur and boron was attempted by including trimethylboron (TMB) in the source gases. Single crystal

diamond with { 111 }, { 110}, { 100} faceting and n-type silicon wafers were used as substrates. The gas flow was

200 sccm, the pressure was 25 torr, and the microwave power was 1020 W. Substrate temperatures were measured

using an optical pyrometer (Williamson 8220C) and ranged from 700 to 750°C. The samples were analyzed for

sulfur by Secondary Ion Mass Spectroscopy (SIMS) and Particle-Induced X-ray Emission (PIXE). Gold-capped

titanium contacts were annealed under vacuum up to 400°C for two hours to create ohmic junctions when required.

Electrochemical measurements were conducted in a cell filled with 0.5 M H2SO4 electrolyte solution, using a

Ag/AgC1 reference electrode. Differential capacitance measurements were conducted with an EG&G Instruments

Potentiostat, Model No. 283, and EG&G Instruments Frequency Response Detector, Model No. 1025.

The growth experiments proceeded in two distinct stages. Prior to the first set of growth experiments, the

chamber was dismantled, cleaned, and reassembled. In the first set of experiments, no boron was used in the feed

gas, and it is believed that there was no significant amount of boron in the reactor. Analysis by SIMS of diamond

grown in the reactor showed no increase in boron concentration above background levels. PIXE analysis of

* Email: sce@pc,cwru.edu, Phone: 216-368-2648, Fax: 216-368-3016

NASA/CP--2001-210948 164 August 6, 2001
diamondsamplesgrownundertheseconditionsgavenodetectablesulfur-incorporation whenHzSwasincludedin

thesourcegas.Additionally, mostsamplesweretooresistiveforelectrochemical measurements.

Thesecondsetofexperiments beganafterboron(asTMB)wasintroducedtothereactor.Thereactorwasnot

disassembled ancleanedbetweentheserunssotherewerelowlevelsofboronpresent evenwhenTMBwasnot

addedtothesourcegases.Residual gasanalysisshowedthepresenceofBH3inrunsdoneintheabsenceofTMB.

Inthissecondsetofruns,i.e.,afterboronwasintroducedtothereactor, sulfurcouldbeobservedbyPIXE.SulfuJc

wasobservedindiamondgrownon{111},{110},and{100}surfaces. Figure1showsaPIXEanalysisfora{110}

sample.

WhenTMBandHzSarebothusedinthefeedgases, theresultingdiamondfilmsarep-type.TheMottSchottkyplotsshowanegativeslope,indicativeofp-typecarrierswithanon-compensated acceptor concentration

rangingfrom1018to1021cm-3.Arepresentative exampleisshowninFigure2,whichhasaflat-bandpotentialt

0.43Vvs.Ag/AgC1. Thisflatbandpotential isabout1Vmorenegativethanisobservedondiamondgrownwithno

sulfuJcpresent inthesourcegas.(SeeFigure5.)If TMBwasnotaddedtothefeedgas,thenthesamplesgrown

usingHzSaregenerallyn-type.TheMott-Schottky plotsshowapositiveslope,whichdenotesn-typeconductivity.

Theslopesindicateddonorconcentrations thatrangedfrom1014to1021cm-3.Arepresentative Mott-Schottky plot

ofoneofthesesamplesisshowninFigure3.

_.(_:E'_.:.::! _ ..................................................................

_._

• .: ,2

•' n . i .%_ :.N

._';_._._._._._._._._._._......_.._._._.5_._._._._._'._._._.._._._._._._._._.._._._._._._._._._._._':''_.._._._..._.._

:2;71,_?_-_ 4, ¢.'_,_'.':

Energy in keV

Figure 1. PIXE results:

counts in arbitrary units

versus energy in keV. H-ion

energy was 3 MeV with a

22.5 ° incident angle. The

sulfur peak occurs at 2.31

keV. The substrate was a

(110) diamond. Growth

conditions were S/C atomic

ratio in the gas phase of

1250 ppm, a methane

concentration of 0.2% and

residual boron in the

reactor. Film thickness was

14 microns.

NASA/CP--2001-210948 165 August 6, 2001
? ..?:¢.;-

),;.

_iii

Jq.

Figure 2. Mott-Schottky plot,

C 2 vs. E, from a (111) diamond

surface. Growth conditions

were S/C atomic ratio in the gas

phase of 2500 ppm, a B/S atomic

ratio of 0.4, and methane

concentration of 0.2%. The

negative slope indicates p-type

conductivity. The number of

l acceptors, NA, is 5 x 1021 cm3.

i The electrolyte is 0.5 M H2SO4.

-5_ -':.: ,':5 ":....., ,:" ',5_, :','. _5

.... '.&,m.._I_.m_.

9

Figure 3. Mott-Schottky plot,

, C 2 vs. E, from (110) diamond

surface. Growth conditions

;: : :' were SIC atomic ratio of 1250

,-,,_., :_::,:; : _. /; ::: : ppm, methane concentration

of 0.2% and residual boron in

_ ::: :; , the reactor. The positive slope

• i ...... _ : ', indicates n-type conductivity.

..t :::,,.::;::i:!ii : The number of donors, ND, is 17 3

: : . 4 x 10 cm. The electrolyte is

d 0.5 M H2SO0.

•, o

,+},,..,. -Ty.--¢-7.-.T.-,?.-.-.-.T77,-.-.--.- ,_ .-,.-.-.-..7-.-,?..-7.--.-,T.-.-,--T.T--Ty,-( :

NASA/CP--2001-210948 166 August 6, 2001
Plasma treatment experiments

Other experiments were performed in which diamond substrates were treated in an H2/H2S plasma at 25 torr for

12 hours without CH4 in the feed gas. There were some residual hydrocarbons in the growth chamber however.

These samples exhibited some of the same electrical properties as samples in which diamond growth occurred in the

presence of H2S. For example, a virgin diamond macle treated for 12 hours with a hydrogen plasma containing 13

ppm H2S gave an n-type Mott Schottky plot with a donor concentration of about 1021 cm -3. This is shown in Figure

4.

2.0E+09

4 _

E

O

¢q

1.5E+09

1.0E+09

5.0E+08

O.OE+O0 i i

0 0.2 0.4 0.6 0.8 1 1.2

E [V vs. Ag/AgCI]

Figure 4. Mott Schottky plot, C 2 vs. E, from a virgin diamond macle (111) surface treated with a H2/H2S

plasma containing 13 ppm H2S for 12 hours.

Treatment of boron-doped diamonds with a sulfur containing plasma was also performed. A Mott-Schottky plot

from a typical heavily boron-doped diamond is shown in Figure 5. The slope indicates p-type conduction with an

acceptor concentration of 2x1019 cm-3; the intercept gives a flat band potential of about 1.3 V vs. Ag/AgC1. The

macle of Figure 5 was then treated under a H2/H2S plasma for 12 hours; the results are shown in Figure 6. The

maximum at _ 0.6 V vs. Ag/AgC1 was reproducible and has been seen in other samples. This type of behavior has

been observed in other systems by others and can arise from several causes: the presence of rapid surface states

located in the band gap (ref. 5) or the presence of more than one donor or acceptor level.

Other Measurements

The response of the open-circuit photo-potential under ultraviolet illumination at 254 nm was also tested and

these experiments confirmed the Mott-Schottky results. Samples that showed n-type Mott-Schottky behavior had a

negative shift in the open circuit photo-potential and p-type samples had a positive shift as expected.

NASA/CP--2001-210948 167 August 6, 2001
Figure 5. Mott-Schottky plot, C 2 vs. E, from a heavily B-doped diamond film grown on a (111) diamond

macle surface. The number of acceptors, NA, is 2.3 x 1019 cm 3. The electrolyte is 0.5 M H2SO4.

C:2, F!2:orn_

8.0×10_

.4..OX'IQla

2.:0x10'_

0,O

I

;(

[]

[]

I • I } " I I I T "I I ' .I ' } I I ' I "" I

-0,2 0.0 0.2 0:4 0.6 :0:8 i.0 i.2 1.4 1.6

E.V {:Ag/AgCi

Figure 6. Mott-Schottky plot of macle from Figure 5 after treatment in a H2/H2S plasma containing 13 ppm

H2S for 12 hours. The electrolyte is 0.5 M H2SO4.

NASA/CP--2001-210948 168 August 6, 2001
Measurements of the sign of the thermoelectric effect were also performed. These measurements were done

using a two-point probe in which one of the probes was heated several tens of degrees above room temperatuace by a

very small resistance heater wound around the probe. Both probes were tungsten. The sign of the thermoelectric

effect was determined by measuring the voltage deflection when the heater was turned on. The system was tested

using known samples of n-type and p-type silicon and with pure copper metal. For the n-type silicon samples, the

cool probe was negative relative to the warm probe; for p-type samples the cool probe was positive relative to the

warm probe. Copper gave no detectable signal. All diamond samples that showed n-type Mott-Schottky plots also

showed n-type thermoelectric effect; all p-type diamond samples showed a p-type thermoelectric effect.

DISCUSSION OF RESULTS

The results can be summarized as follows. Sulfur incorporation, in or on diamond, was enhanced by the

presence of boron. Secondly, n-type behavior was only observed when sulfur was used, either during growth or in a

post-growth plasma treatment, and when only residual amounts of boron were present in the reactor. The

observation of n-type conductivity following plasma treatment with sulfur in the absence of appreciable growth is

indication that surface states or near-surface structures play a role in the observed n-type behavior. This does not

preclude the possibility that donor centers may also arise from substitutional sulfur or sulfuac at grain boundaries.

Recent calculations by Albu (ref. 6) indicate that substitutional S and substitutional BS centers are deep donors, with

levels that each lie about 1.5 eV below the conduction band.

The co-doping experiments with boron and sulfur were motivated by elementary strain energy considerations

that suggest that the smaller boron atom may facilitate the incorporation of the larger sulfur atom into diamond.

Also, gas-phase equilibrium calculations and Langmuir adsorption calculations show that the sulfuac concentration on

the growing surface is increased by the presence of boron. These calculations indicate that gas phase radical species

such as BS and BS2 can increase the sulfur concentration on the surface by up to two orders of magnitude by

bonding to diamond surface radicals.

In the second group of experiments, residual boron was present in the reactor even when TMB was not added to

the source gases. The wide variation in the observed number of donors may be due to different levels of

compensation of the sulfur by the residual boron. Two boron atoms are required to compensate doubly ionized

sulfur; however, boron is more easily incorporated into diamond than sulfur.

It should be noted that in some cases the doping levels in the diamond may have been too high to meet the

assumptions inherent in the Mott-Schottky analysis. Also, as noted, surface states can influence the results.

Therefore, the values of the carrier concentrations obtained from the plots must be treated as approximations.

ACKNOWLEDGEMENTS

The support of the National Science Foundation (Grant CHE98-16345) and the Civilian Research and

Development Foundation (Grant RC1-2053) is gratefully acknowledged. Mr. Cliff Hayman designed and

constructed the thermoelectric device.

REFERENCES

1. Koizumi, S; Teraji, T.; and Kanda, H.: Phosphorus-doped chemical vapor deposition of diamond. Diam. and

Relat. Mater. 9, 3-6, April-May 2000, 935-940.

2. Sakaguchi, I., et al.: Sulfur: A Donor Dopant for n-type Diamond Semiconductors. Phys. Rev. B. 60, 4, July

1999, 2139-2141.

3. Nishitani-Gamo, M., et al.: Homoepitaxial Diamond Growth with Sulfur-Doping by Microwave Plasma-Assisted

Chemical Vapor Deposition. Thin Solid Films, 382, 1-2, Feb. 2001, 113-123.

4. Kalish, R.,et al.: Is Sulfur a Donor in Diamond? App. Phys. Lett., 76, 6, Feb. 2000, 757-759.

5. Barrai, G.; Njanjo-Eyoki, F.; and Maximovitch, S.: Characterization of the Passive Layer and of Hydroxide

Deposits of Nickel by Impedance Spectroscopy. Electrochim. Acta 40, 17, 1995, 2815-2828.

6. Albu, T.; Anderson, A. B.; and Angus, J. C.: Extended Abstract, 2001 Joint International Meeting of the

Electrochemical Society and the International Society of Electrochemistry, San Francisco, September 2-7, 2001

NASA/CP--2001-210948 169 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

STM AND ELECTROCHEMICAL INVESTIGATION OF HOMOEPITAXIAL BORONDOPED CVD DIAMOND FILMS

John B. Cooper, Jason A. Moulton

Department of Chemistry, Old Dominion University, Norfolk, Virginia 23529

Sacharia Albin, and Bing Xiao

Department of Electrical and Computer Engineering, Old Dominion University, Norfolk, Virginia 23529

ABSTRACT

Homoepitaxial growth of boron-doped CVD diamond films was carried out on (100) and (111) oriented

substrates. Atomic resolution images were obtained for both (100) and (111) surfaces using scanning tunneling

microscopy. STM images reveal the presence of a 2xl-monohydride reconstruction for the untreated (100) surface

and a lxl reconstruction for the untreated (111) surface. No other atomically resolved reconstructions were

observed under a wide range of growth conditions. Non-aqueous electrochemical investigations were carried out on

the films exhibiting atomically resolved reconstructions. Evidence for potential-induced surface-reconstruction and

surface chemical modification of the (100) 2xl-monohydride surface has been observed.

Keywords: Homoepitaxial, STM, Electrochemical Modification, Diamond, Surface Reconstruction

John B. Cooper

Department of Chemistry

Old Dominion University

Norfolk, VA 23529 USA

Fax: (757) 683..-4628

Tel: (757) 683-4096

NASA/CP--2001-210948 170 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

HTHP DIAMOND ELECTRODES: ELECTROCHEMICAL IMPEDANCE

CHARACTERISTICS OF INDIVIDUAL CRYSTAL FACES

Yu. V. Pleskov 1, M. D. Krotova, V. Ya. Mishuk, Yu. E. Evstefeeva

Frumkin Institute of Electrochemistry, Russian Academy of Sciences, 117071 Moscow, Russia

V. A. Laptev

All-Russian Research Institute for Synthesis of Mineral Raw Materials, Aleksandrov, 601600 Vladimir district,

Russia

Yu. N. Palyanov, Yu. M. Borzdov

Institute for Mineralogy and Petrography, Siberian Division, Russian Academy of Sciences, 630090

Novosibirsk, Russia

ABSTRACT

Effects of crystal structure on the electrochemistry of boron-doped high-temperature-high-pressure diamond

crystals grown from a Ni Fe C melt are studied. The boron concentration in the batch was 0.01 to 0.1 wt. %. On

the { 111 }, {100}, and {311 } faces, the impedance-spectroscopy measurements were performed and the kinetics of

• . 3-/4-

redox reactions in the Fe(CN)6 system were evaluated. The HTHP single crystals closely resemble

polycrystalline CVD films in their electrochemical properties. In particular, the electrodes' equivalent circuit

includes a constant phase element. The non-compensated acceptor concentration in the semiconductor diamond was

determined from Mott--Schottky plots and by the nonlinear impedance measurements; it varied in the 1016 to 102°

-3

cm range. The difference in the electrochemical behavior of individual faces can be tentatively put down to different

boron concentration in the growth sectors associated with the faces, resulting from the different ability of the

diamond crystal faces to entrap the boron dopant during growth.

Keywords: diamond single crystals, boron dopant, crystal faces, impedance spectroscopy, electrochemical

kinetics.

INTRODUCTION

Owing to its extraordinary chemical stability, diamond is an attractive electrode material for use in theoretical

and applied electrochemistry (refs. 1 3). Comparative studies of chemical-vapor-deposited single-crystal

(homoepitaxial) and polycrystalline diamond thin-film electrodes, grown by the hot-filament techniques, as well as

amorphous diamond-like carbon electrodes, revealed the role of crystal structure in the electrochemical behavior of

diamond. By comparing the measured values of kinetic characteristics of various redox reactions, it was concluded

that the single crystal and polycrystalline diamond electrodes are similar in their kinetic behavior. Because the widegap diamond-like carbon (assumed to be a model material of intercrystallite boundaries) turned to be inactive as an

electrode, we concluded that the electrode behavior of polycrystalline diamond films is entirely determined by the

diamond crystallites, rather than the intercrystallite boundaries, at least, over a moderate electrode polarization range

(ref. 4). However, at high anodic potentials (just prior to the onset of oxygen evolution) a minor current peak is

observed at potentiodynamic curves taken with polycrystalline electrodes, which cannot be seen with single-crystal

films. The peak was ascribed to the oxidation of sp_-carbon of intercrystallite boundaries (ref. 5).

Our preliminary studies showed that high-temperature-high-pressure single crystals, on the whole, are similar to

the CVD polycrystalline films in their electrode behavior. In particular, both the polycrystalline and single crystal

diamond electrodes are equally characterized by a special type of frequency-dependent capacitance (refs. 6, 7).

In this work our principal attention was focused on studying electrochemical properties of individual crystal

faces of the HTHP diamond single crystals. The acceptor concentration was evaluated by differential capacitance and

nonlinear impedance measurements; the electrochemical activity of the faces, by taking cyclic voltammograms in the

Fe(CN)63-/4- system.

EXPERIMENTAL

1 E-mail: pleskov@electrochem.msk.ru ; Fax: 7 095 9520846; Tel.: 7 095 9554549

NASA/CP--2001-210948 171 August 6, 2001
Theboron-doped diamondsinglecrystalsweregrownfromaboron-containing Ni_e Cmelt,atp,Tconditionswithinthediamondthermodynamic stabilityrange.Thegrowthwasperformed, usingaseed, bythe

temperature gadient techniques, underconditionspreventingitrogencapturebygrowingcrystals(eerefs.8,9

fordetails). Attheconclusionfsynthesis, thesolidifiedmetalwasdissolvedinaCr207+H2SO4mixtureat

80 100°C,anddiamondcrystalsextracted. Thenear-seed regionofthecrystalsthusobtainedwasthenground-off

mechanically.

Fig. 1. Crystal #168/1 (top view).

The two samples studied are presented schematically in Figs. 1 and 2. The crystal #168/1 (Fig. l) is a cubooctahedron depressed along the 3r° order-axis. Its basal plane is a { 111 } growth face. The opposite face was ground

and polished, as mentioned above; its resulting orientation approached { 111 }. The sample's lateral faces make a

combination of {100 } and { 111 } facets of smaller size. Still smaller are the rhombododecahedral [{ 110 }, truncating

the edge between two { 111 } faces] and tetragontrioctahedral [ {311 }, truncating the edge between { 111 } and { 100}

faces]. During the growth process, the batch contained relatively small amount (0.01 wt. %) of boron, which resulted

in a moderate doping level. The growth sectors associated with the octahedral faces, which dominate the crystal

bulk, are colored light-blue (shown gray in Fig. 1).

Crystal #360 (Fig. 2) is a cubo-octahedron with unevenly developed {111 } and {100} faces, which is due to

Fig. 2. Crystal # 360 (top view).

the complicated character of symmetry in the source melt. The boron concentration in the batch (0.1 wt. %) was an

order higher compared to the crystal #168/1. This resulted in a higher doping level; a nearly nontransparent crystal

of dark indigo_lack color was obtained. The coloration is also enhanced by a larger relative volume of the growth

sectors associated with the cube faces, which are known (ref. 10) to be characterized by much higher boron-admixture

concentration as compared to those associated with the octahedral faces.

Ohmic contacts to the rear sides of the crystals were made using silver epoxy. The octahedral, cube, and

tetragontrioctahedral faces were consecutively exposed to the electrolyte, the rest of the crystals' surface being

insulated by high-purity paraffin. Prior to measurements, the working surface was "refreshed" in concentrated HF

solution.

NASA/CP--2001-210948 172 August 6, 2001
Thespectraweremeasuredoverthefrequency rangefrom20Hzto200kHzusinganR-5021acbridge;insome

experiments, withanimpedance spectraanalyzer Solartronmodel1250,equippedwithanelectrochemical interface

1286.Theequivalent circuitcomponents werecalculatedusingaproceduredscribedinrefs.6,7.Thenonlinear

impedance oftheelectrodes wasstudiedbytheamplitude-demodulation method(ADM).Themeasuringprocedure

consistedinapplyingahigh-frequency current signal,modulatedinitsamplitudeatalowfrequency, tothecelland

measuringavoltageresponseatthislowerfrequency (ref.4).Thecyclicvoltammograms weretakenunderlinear

potential scanning(atascanrateof0.005Vs-l),usingaPI-50-1potentiostat equippedwithaPR-8programming

unitandaPDA-1X--Yrecorder.

Electrochemical measurements wereperformedinathree-electrode cell,withaplatinumauxiliaryelectrodeanda

normalsilver/silver chloriderferenceelectrode(inwhatfollows,allpotential valuesaregivenagainst this

electrode). Thesupportingelectrolytewasa0.5MH2SO4 solution. The measurements were performed at room

temperature.

RESULTS AND DISCUSSION

From the samples' resistivity and the sign of photocurrent and photopotential we concluded that the diamond

under study is a p-type semiconductor. Therefore, we used our experience in the electrochemistry of semiconductors

(ref. 11) in the characterization of the samples with electrochemical techniques.

Shown in Fig. 3 is the complex-plane presentation of an impedance spectrum. It is a semicircle somewhat

distorted at lower frequencies, probably due to a slow process at the interface or in electrolyte. A closer inspection of

the spectrum revealed a specific type of frequency dependence represented by a constant phase element (CPE) in the

equivalent circuit. The CPE impedance is known to equal Zcp_ = 6 -1 (i0)) -a, where i = (- l) 1/2, 0) is the ac angular

frequency, the power a determines the character of the frequency dependence, and the quantity 6 is measured in

Fa_a-lcm-2 units.

,_,

5

/

.e

•l ...... _------_,------_-------_ ............... T.............. _ ....... L._._.___._._.__._._.__

Fig. 3. Impedance spectrum for sample #168, {111} basal face. The ac frequencies are shown in kHz.

In this respect, the HTHP crystals are full analogy to the CVD films we studied earlier. The frequency

dependence of the differential capacitance of diamond electrodes was tentatively ascribed to a slow process in the

space-charge layer, probably, involving a deep energy level (refs. 4, 7).

Because a usually approached 1 (see Table 1 below), in our further analysis we approximated 6 with a

frequency-independent capacitance C, which we used in plotting Mott--Schottky (that is, the reciprocal of C 2 vs.

potential E) dependences (Fig. 4). The plots are reasonably linear; the intercepts at the abscissa axis give the flatband potentials. The very positive flat-band potentials are, probably, due to the oxidized state of the HTHP crystal

surfaces, caused by the oxidative pre-treatment, see the Experimental section. And for the heavily doped { 100} face,

an apparent positive shift of the flat-band potential might be partly caused by a transition to degeneracy (NA

1020 cm -3, see Table 1).

NASA/CP--2001-210948 173 August 6, 2001
_V

Fig. 4. Mott--Schottky plots for two faces of sample #360.

From the slope of the lines d(C-2)/dE the uncompensated acceptor concentration was calculated:

-2

NA = _ [d(C-2)/dE] -1 , (1)

where _ and _0 are the permittivities of diamond and free space, respectively, e is the electron charge. The NA values

thus obtained are given in Table 1.

Table 1. Properties of Individual Faces.

No. of Crystal Face NA, cm -_ a (y AEp, mV

Mott Schottk_ ADM

1.2 1019 2.3 1019 0.95 0.9 1.0 130

360 4.1 10 Is 0.95 1.3 90

168/1

{111} basal

{ 111 } lateral

{111} rear

_1oo1

t1111 basal

{ 1111 lateral

_311}

1 1019

5 1019

2.3 101°

1 020

2 101°

1.4 101/

0.9 0.94

0.93

0.92

-0.9

1.3 2.2

0.016

0.06

0.04 0.09

*The redox current maximums superimposed on the H2 and 02 evolution current and cannot be se

1700"

730

- 1400"

_arated.

Another way to find the acceptor concentration we used is the amplitude demodulation method (ADM).

..... 2

According to the theory of the method, the measured cell response is inversely proportional to the d(C )/dE value.

This opens the way to directly determining NA; the values found are also given in Table 1. We see that the two

methods are in reasonable agreement with each other. Some deviation can be due to the fact that the Mott Schottky

data relate to the measuring frequency of 61 kHz; the ADM data, to 500 kHz.

Despite some scatter in the data, we can draw the following conclusions:

(1) Crystal #360 is much heavily doped than crystal #168/1. This is no surprise, since the boron-dopant

concentrations in the batch differed by order of magnitude. The higher capacitance (more precisely, 6) values

corroborate qualitatively the conclusion drawn from the Mott Schottky and ADM measurements.

NASA/CP--2001-210948 174 August 6, 2001
(2)The{100} faceismoreheavilydopedthan{111} faceinthecrystal#360.Thisdisagrees withtheknownability

ofa{111} facetoentrapboronmoreintenselyascomparedwith{100}faceduringrowthprocess. Thispoint

deserves furtherinvestigation.

(3)Surfacepolishingdoesnotaffectsignificantly thecapacitance: nodifferenceinNAwasfoundbetweenthebasal

(growth) andtherear(polished) {111}facesofthecrystal#360.

The{311}faceinthecrystal #168/1appearedtoberatherlightlydoped.

Fromthe<<series_ resistance in the equivalent circuit, the resistivity of the crystal #360 can be roughly

estimated as 10 to 20 _ cm, a reasonable value for the heavily doped single crystal.

We now turn our attention briefly to the kinetic data. The degree of reversibility of an electrochemical reaction

can be qualitatively judged from the potential difference for the anodic and cathodic current peaks AEp on cyclic

voltammograms. The higher AEp (the wider the current peaks separate along the potential axis), the more

irreversible (slow) is the reaction. For reversible reactions, AEp = 59 mV (ref. 12). A typical voltammogram taken in

the 0.5 M H2SO4 + 0.01 M Fe(CN)63- + 0.01 M Fe(CN)64- solution is shown in Fig. 5.

"7'

;i.

[\,

/

j i

_%% ,'

Fig. 5 Cyclic voltammogram for the {111} face of crystal #360.

Comparing the AEp values obtained for the crystals under study (Table 1), we see that on the heavily doped

sample #360 the basal {111 } face is very active: the reactions proceed in nearly reversible mode (the AEp values

approach the theoretical 59 mV value). The moderately doped crystal # 168/1 demonstrates much less degree of

reversibility for the {111 } faces, which falls into general rules of electrochemical kinetics on semiconductor diamond

(ref. 13). Unexpected is the behavior of the {100} faces which are more heavily doped than the {111} faces: the

electrode reactions are markedly hindered. This point deserves further investigation.

CONCLUDING REMARKS

The difference in the electrochemical behavior of individual faces can be tentatively ascribed to different boron

concentration in the growth sectors associated with the faces, resulting from the different ability of the diamond

crystal faces to entrap the boron dopant during the growth process, rather than to the surface atomic density or other

purely surface properties. To reveal these fine effects, individual faces of different crystals must be compared under

conditions of equal doping level.

NASA/CP--2001-210948 175 August 6, 2001
ACKNOWLEDGMENT

Thisworkwassupportedby theCivilianResearchandDevelopment Foundation, grantno.RC1-2053.

Authorsaregrateful toProf.J.C.Angusforhelpfuldiscussion ofthemanuscript.

REFERENCES

1. Pleskov, Yu.V.: Synthetic Diamond in Electrochemistry. Russ. Chem. Rev., vol. 68, no. 5, May 1999,

381 392.

2. Swain, G.M., Anderson, A.B., and Angus, J.C.: Applications of Diamond Thin Films in

Electrochemistry. MRS Bulletin, vol. 23, Aug 1998, 5642.

3. Tenne, R., and Levy-Clement, C.: Diamond Electrodes. Israel J. Chem., vol. 38, 1998, 5743.

4. Pleskov, Yu.V., Evstefeeva, Yu.E., Krotova, M.D., et al.: Synthetic Semiconductor Diamond Electrodes.

The Comparative Study of the Electrochemical Behaviour of Polycrystalline and Single Crystal Borondoped Films. J. Electroanal. Chem., vol. 455, no. 2, 1998, 139 146.

5. Martin, H.B., Argoitia, A., Landau, U., et. al.: Hydrogen and Oxygen Evolution on Boron-doped

Diamond Electrodes. J. Electrochem. Soc., vol. 143, no. 6, June 1996, L133 136.

6. Pleskov, Yu.V., Evstefeeva, Yu.E., Krotova, M.D., et al.: Synthetic Semiconductor Diamond Electrodes.

A Study of Electrochemical Behavior of Boron-doped Single Crystals Grown at High Temperature and

High Pressure. Electrochim. Acta, vol. 44, 1999, 3361 3366.

7. Pleskov, Yu.V., Elkin, V.V., Abaturov, M.A., et al.: Synthetic Semiconductor Diamond Electrodes.

Elucidation of the Equivalent Circuit for the Case of Frequency-dependent Impedance. J. Electroanal.

Chem., vol. 413, no. 2, 1996, 105 110.

8. Pal'yanov Yu.N., Khokhryakov AJ_., Borzdov Yu.M., Sokol A.G., Gusev V.A., Rulov G.M., Sobolev

N.V.: Growth conditions and real structure of synthetic diamond crystals. Russ. Geology and Geophysics,

vol. 38, no. 5, 1997, 920-945.

9. Butuzov, VP., Laptev, V.A., Presnov, V.A., et al.: Growth and Study of Synthetic Semiconductor

Diamond of Different Conductance Types (in Russian). DokladyAkad. Nauk SSSR, vol. 226, no. 2, 1976,

328 332.

10. Laptev, V.A., Pomchalov, A.V., Palyanov, Yu.N., et al.: Growth of Diamond Crystals (in Russian). Proc.

of 3rd Int. Conf. "Crystals: Growth, Properties, Real Structures, and Applications", Aleksandrov, Vladimir

district, vol. 2, 1997, 119 123.

11. Pleskov, Yu.V. and Gurevich, Yu.Ya.: Semiconductor Photoelectrochemistry. Consultants Bureau, New

York, NY, 1986.

12. Delahay, P.: New Instrumental Methods in Electrochemistry. Interscience Publishers, New York, NY,

1954.

13. Modestov, A.D., Evstefeeva, Yu.E., Pleskov, Yu.V., et al.: Synthetic Semiconductor Diamond Electrodes.

Kinetics of some redox Reactions. J. Electroanal. Chem., vol. 431, no. 2, 1997, 211418.

NASA/CP--2001-210948 176 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FAR INFRARED CVD DIAMOND OPTICS: PROPERTIES AND APPLICATIONS

T P Mollart and K L Lewis

Defence Research and Evaluation Agency

St. Andrews Road, Great Malvern, Worcs. WRI4 3AX UK

ABSTRACT

Diamond is a unique optical material having the combination of high strength and transparency in file visible,

near and far infrared regions. As a far infrared window it has a wide range of potential uses in demanding

environments, e.g. high power CO: lasers exit windows and airborne optics. In this paper, the general optical

properties of diamond will be briefly reviewed and why it is so useful will be highlighted by a comparison to other

infrared materials. The fax infrared optical properties, such as absorption and scatter and how they are linked to the

material characteristics will be described.

Recent progress in synthesis technology enables large area three dimensional diamond components to be

fabricated. The optical properties of a polished 70 mm diameter hemispherical diamond dome will be described.

When considering diamond for large area applications the mechanical properties are importhm. The strength of

CVD diamond is determined by its polycrystalline nature, how it is one of the limits to potential applications will be

discussed.

High temperature transmission spectra from 5000 cm"1(2 gin) to 200 cm"l (50 gm) for CVD diamond will be

presented. In the far infrared region the transmittance drops with temperature and a broad peak approximately

centred at -1275 cm * (7.8 gin) emerges. In this experiment the temperature dependency of the absorption

coefficient has been deduced by assuming there are no additional losses from the surface and calculating the

reflectance using a temperature dependent value for the refractive index. Direct emissivity and photo-thermal

deflection measurements have verified these finding and indicate that it is art intrinsic property of diamond.

Work is still necessary on optical coatings for diamond, since in many of these applications coatings would

enhance the functionality of the optic. Good adhesion to diamond can be achieved by exploiting the surface

termination of the surface. How durable antireflection oxide coatings, such as yttria (Y203) can be deposited on

diamond and their properties will be described.

Keywords: Diamond, optical properties, domes, AR coatings

T.R Mollart

M210

DERA Malvern

St. Andrews Road

Malvern

Worcs.

WR14 3AX

Tel: + 44 (0) 1684896649

Fax:+44 (0) 1684894311

tpmollart(&,dera, gov.uk

NASA/CP--2001-210948 177 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Thin Film Diamond Detectors for Industrial Applications in Next Generation 157nm

Lithography and High power Excimer Laser Beam Diagnostics

Michael D Whitfield 1,Nadeem Rizvi 1, Robert D McKeag 2, Stuart P Lansley 3, Olivier

Gaudin 3 and Richard B Jackman 3

1Exitech Ltd. Hanborough Business Park, Long Hanborough, Witney OX29 8SL. UK.

2Centronic, Croydon. U.K

3Departement of Electronics & Electrical Engineering, University College London,

Torrington Place, London WC1E 7JE. U.K.

ABSTRACT

The next generation of photolithography stepper tools will operate at 157nm allowing the

routine fabrication of sub-100 nm device features. For efficient commercial operation

these machines will require robust solid-state UV detectors for direct beam monitoring

and photoresist exposure dosimetry. Currently, no commercially available detectors can

fully meet the all the specifications required for these applications. Silicon 'detectors, for

example, suffer from rapid performance degradation and unacceptably high attenuation in

the surface metallisation and passivation layers. This paper examines the role of CVD

diamond as a detector material solution for use in 157nm research and future production

steppers and in other high power excimer laser micro-machining applications. The

authors discuss the fabrication of simple inter-digital surface point source detectors

including the passivation and smoothing steps needed to realies high performance

detector structures. The results of a detailed study of detector performance under

industrial conditions on the world's first commercially available 157nm micro-stepper

tool are presented. To improve device sensitivity and transient response times it is

necessary to carry out gaseous thermal passivation treatments to modify the defect

structure of the near surface layer forming the active layer in these detectors. Spectral

and pulse response characteristics have been evaluated by steady state and transient

photoconductivity methods. Device sensitivity was studied over the laser fluence range

0.05-1.5 mJcm a. Device gain was found to be liner with applied bias over a 3-order

range, Long-term operational performance was studied up to a total of I07 laser pulses.

Charge sensitive deep level transient spectroscopy and transient photoconductivity has

been used to relate the observed device performance to the electrical defect structure of

the CVD diamond material.

Many laser micro-machining applications would benefit from a device structure that

would enable direct beam imaging for beam shaping and monitoring purposes at the

work-piece. The paper will present results from performance studies on a 1-D array

CVD diamond deep UV detector that has been used to image 193nm excimer laser beams

and will discuss the design, fabrication and testing of a prototype 2-D array that is

currently being built by the authors.

NASA/CP--2001-210948 178 August 6, 2001
Proceedingsof the SixthApplied DiamondConference/Second

FrontierCarbonTechnologyJoint Conference(ADC/FCT2001)

Y. Tzeng, K.Miyoshi,M. Yoshikawa,M. Murakawa,Y. Koga,

K. Kobashi,G.A.J.Amaratunga(Editors)

SINGLE-PIXEL DIAMOND-BASED VUV PHOTODETECTORS

FOR SPACE APPLICATIONS

A. De Sio, E. Pace ", A. Pini

Dip. Astronomia e Scienza dello Spazio, Univ. Firenze, L.go E. Fermi 2, 50125 Firenze (Italy)

S. Scuderi

Osservatorio Astrofisico di Catania, Catania (Italy)

ABSTRACT

During the last decade, many studies have addressed the use of synthetic diamond as sensing material for deep UV

photon detection. Experiments on the photoelectrical properties of diamond at wavelengths down to 190 nm have been

reported in literature, while a few papers concern the same properties at shorter wavelengths (_ > 120 nm). Some future,

but already under way, space missions will require UV photon detectors with unprecedented sensitivity at wavelengths

_,> 30 nm and radiation hardness. Diamond appears to be the ideal photosensitive material for such applications, but an

experimental and technological effort must be faced to investigate its properties in the unexpIored region between 30

nm and 120 nm. Up 1o now, this has been an unexplored spectral region for the electro-optical properties of diamond

and therefore this paper is aimed at filling the gap, reporting a comparative analysis and characterization of the

performances of different diamond-based photodetectors in the vacuum ultraviolet, namely 50-200 nm. In partieular

four different types of diamond have been tested: polycrystalline CVD (pCVD), homoepitaxial CVD type IIa (sCVD),

HPHT type Ib single crystal (IbHPHT) and HPHT type IIa single crystal (IIaHPHT). Photoconductive devices have

been fabricated evaporating interdigitated planar gold contacts on each sample by photolithographic technique and

annealed to ensure a good adhesion of the metal on the diamond surface.

A full characterization has been performed including I-V characteristics in dark and under UV steady light to

determine the gain curve versus the applied voltage for each device. The contribution of material quality and

morphology to the photo-induced currents has been also investigated. Moreover the spectral responsivity has been

measured for each device operating at the best working point previously determined. Finally time responses have been

analysed using pulsed radiation.

Keywords: Polycrystalline diamond, Single crystal diamond, photoconduetivity, UV photodetectors

* Presenting and corresponding author:

Dr. Emanuele Pace

Dip. di Astronomia e Scienza dello Spazio

Universit_ di Firenze

L.go E. Fermi, 5, 50125 Firenze (Italy)

E-mail paee@areetri, astro.it

Off. Tel. +39 (55) 27 52 281

Lab. Tel. +39 (55) 27 52 222

Fax +39 (55) 27 52 222

NASA/CP--2001-210948 179 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

DIAMOND THIN-FILM FABRY-PEROT OPTICAL RESONATORS

R. A. Booth and D. K. Reinhard

Department of Electrical and Computer Engineering

2120 Engineering Building

Michigan State University

East Lansing, MI 48824-1226

Email: "_o_o__! Lh_r_o_ g _ _e_L__s_tl &d__

Phone: 517/355-5214

Fax: 517/353-1980

ABSTRACT

Diamond thin-film based Fabry-Perot optical resonators are fabricated on a silicon substrate, based upon

established integrated-circuit fabrication technology. Processing procedures include thermal oxidation,

photolithography, wet etching, plasma enhanced chemical vapor deposition, and sputtering. The resonant structure

consists of a diamond membrane with partially-reflecting gold films on both sides. The through-transmission of the

resulting structures is investigated both experimentally and theoretically. Simulations of device performance are

based upon standard optics theory, using the matrix approach for multi-layer optical structures. The simulations

account for absorption and phase shifts in the front and back gold films and use published empirical data for

constants such as the gold absorption coefficient and index of refraction for gold and diamond. A limitation of the

theoretical calculations is that the matrix approach used to determine percent transmission assumes smooth, scatterfree interfaces. The experimental performance of the diamond-film resonators has been investigated from the

ultraviolet, through the visible, and into the infrared region of the optical spectrum (200 nm 2500 nm).

Experimental results will primarily be presented for the range of 1050nm to 1900nm. Simulations accurately predict

the wavelengths at which peak transmission is observed and the general shape of the transmission peaks. However,

the simulations predict higher percent transmission and higher Q values than are actually observed. Simulation

results suggest that the difference in theory and experiment is not due to absorption in the diamond. Therefore the

role of scattering at the cavity surfaces is currently being studied.

Keywords: Fabry-Perot, thin-film, diamond, optical, fiker

INTRODUCTION

In a world which has seen everything from transistors to gyroscopes to lasers integrated onto silicon wafers, it

should come as no surprise that various optical devices are being constructed in small scale on silicon wafers. CVD

diamond is of interest as an optical material for structures on silicon because of its wide optical transmission

window and because it is compatible with silicon substrates. Diamond films have previously been reported to be

effective in enabling photon entry into silicon as anti-reflection layers over a wide range of wavelengths (ref. 1). In

this study, the performance of diamond as an optical medium for optical resonators is investigated in structures

formed by photolithography on silicon. Experimental results are compared with model calculations.

BACKGROUND

The Fabry-Perot resonant cavity is an important device for a variety of optical applications, including highly

selective optical band-pass filters. It is of interest to integrate such devices into on-wafer structures by incorporating

optical materials into the processing sequence (ref. 2,3). An ideal Fabry-Perot cavity consists of a non-absorbing

optical medium with a partially reflective, non-absorbing coating on either side with perfectly smooth interfaces

between the different materials. The ideal cavity will allow 100% transmission of certain wavelengths of light,

while heavily attenuating other wavelengths.

When constructing an optical resonator, absorption in the optical medium between reflectors can become a

concern. For example, SiO2, which absorbs heavily in portions of the infrared (IR), would not be a good FabryPerot medium if the device had to operate at IR wavelengths. Likewise, silicon would not work well in the visible

portion of the spectrum, due to its small band gap. Diamond is a particularly attractive material for broad-based

NASA/CP--2001-210948 180 August 6, 2001
Fabry-Perot applications, becauseit istransparent fromtheultraviolet(UV)intothemmwaveportionofthe

spectrumwithveryslightabsorptionintheIR.

Anotherpractical concernisthathematerial usedtocreatehereflectivesurfacemaycauseabsorptionlosses.

Thisisthecasewithmostmetals, especially intheIRregime.Additionally, resonator performance maybelimited

if surfaceroughness isnon-negligible. Theimpact ofthesefactorsiinvestigated inthiswork.

OPTICAL MODELING

The theoretical calculations in this paper are based primarily on the matrix approach to multi-layered optical

structures (ref. 4,5). The matrix method assigns a characteristic matrix to each layer of the optical structure. The

matrix contains information about the incident wavelength, film thickness, and complex index of refraction. This

matrix is given by:

I cosfj i sin 5j

Mj

= L nj in j sin 5 j cos 5 j

(1)

where nj is the index of refraction for the jth layer, and 8j is defined for vertical incidence as:

(2)

with _ being the incident wavelength, and tj being the thickness of the jth layer.

A special matrix is assigned for the final medium in the beams path. If the index of refraction for the final

medium is nf, the matrix for the final medium is given by:

(3)

The fraction of the incident beam that is transmitted, T, is given by:

T- nf_12

n o

(4)

where n o is the index of refraction of the medium in which the beam originates. For purposes of this paper, n oand nf

can be considered equal to unity, since the beam originates and terminates in air. The quantity t is defined by:

2

t-

(a + b)

(5)

where for an m-layered optical structure, a and b are obtained from the characteristic matrices in the following

manner:

(6)

This matrix method is valid for optical structures with arbitrarily large numbers of layers down to single layered

optical structures, which is an appealing attribute in the analysis and design of optical resonators.

As example, a modeling calculation for a gold-diamond-gold Fabry-Perot resonator is shown in Figure 1.

Calculations are based upon a diamond membrane thickness of 1.59 gm, and partially transparent gold film coatings

of 15 nm and 25 nm. Optical data for gold is taken from Kingslake (ref. 6). The index of refraction for diamond is

NASA/CP--2001-210948 181 August 6, 2001
takenfromtheSellmeier equation(ref.7).Thegoldfilmsarenotofthesamethicknessduetoexperimental

considerations explainedinafollowingsection.Thesemodelingresultsaccount forabsorption inthegoldfilmsbut

donotaccount forabsorption inthediamondorforscatteringlosses.

O

O

_7

•-- 0.4

¢-

O

0.3

E

c 0.2

t-

¢-

o 0.1

" 0

0.5 ........................................................................

J /

J-" <-.____j// L :_.

1050 1250 1450 1650 1850

Wavelength (nm)

Figure 1 Theoretical calculation of transmission versus wavelength for a gold-diamond-gold optical

resonator.

One measure of resonator performance is the cavity Q (ref. 8), where Q is defined by:

Q= v0

AV1/2

(7)

where v0 is the resonance frequency and AVl/2 is the full width half maximum. The modeling results indicate Q

values in the range of 59-67 for the simulation shown in Figure 1.

DEVICE FABRICATION

Figure 2 shows a cross-sectional view of the device constructed for this study. The Fabry-Perot device is

fabricated on a double-side polished, 2" diameter p-type silicon wafer. After an RCA cleaning process (ref. 9), the

wafer is thermally oxidized to an approximate oxide thickness of 1.5 gm on both sides. One side of the wafer is

then coated with photo resist and patterned with an array of small squares, approximately 2mm X 2mm. The

squares are aligned with the flat edge of the wafer. An oxide etch is performed after the photolithography. One side

of the wafer is not covered with photo resist at all, while the other side has a pattern. The oxide is etched all the way

to the silicon, so after etching, one side of the wafer is bare silicon, while the other side has a patterned oxide. The

photo resist is then stripped.

The wafer is then polish seeded for diamond film growth on the non-patterned side, and a diamond film is then

deposited on the wafer. Diamond deposition parameters proved to be very critical, as intrinsic stress in the film due

to thermal mismatch is a large factor in the flatness and durability of the film. The deposition was performed at 35

Torr, at a temperature of approximately 660°C. This resulted in a flat film membrane. Depositing at substantially

lower temperatures would usually result in films that would wrinkle and in some cases break at the completion of

the silicon through-etch.

After the diamond film deposition, the silicon wafer is through-etched to create the diamond membranes. The

patterned oxide is used as a mask, while the diamond film has proven to be relatively resilient to the KOH etch, thus

NASA/CP--2001-210948 182 August 6, 2001
servingasamaskonthesecondsideofthewafer.Thethrough-etch isperformedatatemperature of60°C,ina44

/56 weight percentage KOH/H20 solution. When the diamond membrane is created, the through transmission of the

membrane is measured as described in the Experimental Results and Discussion section.

_NNN'xNN\\\NNN\\NNNN\\J t t t \X,.\\\'xN%\\\'xN%\\\',Xx,Jt

i:i:i:i:i:__:i:i:i:i i:i:i:i:i:ii:i:i:i:i:i i:i:i__ i:i:i:i:

[] Silicon _ Gold

.. Si02 :t:i: Diamond

Figure 2 Cross-sectional view of two Fabry-Perot resonators. Drawing not to scale.

Once the diamond membranes are formed and measured, the final step is to sputter coat gold on both sides of

the membrane. A simple shadow mask is used to selectively coat the windows. The diamond film tends to have a

slightly rougher surface on the growth side as opposed to the silicon side of the film. The result is that the gold

deposition on the growth side of the membrane has to be slightly thicker to form a continuous film. Experimentally,

a ratio of 5/3 was determined for the thickness of the gold film on the rough/smooth surfaces of the membrane for

minimum gold thickness.

EXPERIMENTAL RESULTS AND DISCUSSION

The device is characterized in two steps. First, before the diamond membrane is sputter coated with gold, the

through-transmission vs. wavelength is obtained for the diamond membrane only. This aids in determining the

thickness of the particular membrane window. Secondly, the transmission through the gold-coated resonator is

measured.

The through-transmission is measured by passing a monochromatic beam of light through the sample, which is

masked to block any light from passing through the silicon or other diamond membranes. The wavelength is varied

and transmitted power at each wavelength is noted. The sample is then removed from the optical system, and a

mask with the same size opening as was used on the sample is placed in the path of the beam. The power of light

passing through the mask with no sample is then recorded for each wavelength. The "sample" data is then

normalized against the "air" data to obtain an accurate value for percent transmission.

Figure 3 shows the transmission versus wavelength obtained for an uncoated diamond membrane using the

method described above from the UV to IR. For a lossless diamond film, the transmission would oscillate vs.

wavelength with peaks of 100% and minimum values of approximately 50%. The separation of the measured peaks

provides an accurate estimate of the particular diamond membrane's thickness. In the case of the data shown in

Figure 3, the thickness of the film is found to be approximately 0.7 btm. In the measured transmission strong

absorption is seen in the UV beyond the bandgap (approximately 225nm). Previous studies indicate the nonbandgap absorption losses are mainly due to scattering at the rougher surface (ref. 10). AFM results for our films

indicate a surface roughness of approximately 15nm.

Figure 4 shows the results of the measured transmission versus wavelength for a fabricated Fabry-Perot

resonator. The thickness of the diamond film for this set of data is approximately 1.59 _tm, with gold coatings of

approximately 15 nm and 25 nm. This device served as the basis for the simulation shown in Figure 1. Location

and general shape of the peaks correspond well with theoretical calculations, however the Q values are less than the

NASA/CP--2001-210948 183 August 6, 2001
modelingwouldpredict.Figure5showsanexpanded viewofthepeaklocatedatapproximately 1600nmseenin

Figure4. ThepeakshowninFigure5hasaQvalueof18.

Thefacthathemeasured QislessthanthemodeledQislikelyaresultofthefacthatscatteringisnot

includedinthemodeling. Itisanticipated thatsmoother filmswouldresultinahighermeasured Q.

0.9

O

]i 0.8

,- 0.7

O

"_ 0.6

.__

_= 0.5

0.4

_- 0.3

tO

:= 0.2

0.1

1,1.

01

200 700 1200 1700 2200

Wavelength (nm)

Figure 3 Representative data plot showing measured fraction of incident light transmitted through an

uncoated diamond membrane versus wavelength.

¢-

O

.m

E

¢-

O

.m

U

IJ0.1

0.08

0.06

0.04

0.02

/ /' Is 1

\\, J,

"\ jJ i

1050 1250 1450 1650 1850

Wavelength (nm)

Figure 4 Measured transmission characteristics for a fabricated Fabry-Perot resonator.

NASA/CP--2001-210948 184 August 6, 2001
co

.m

t_

t_

E

t_

eb-

,g

1.1.

0.08

0.06

0.04

0.02

\,

1450 1500 1550 1600 1650 1700 1750 1800

Wavelength (nm)

Figure 5 Expanded view of a peak shown in Figure 4.

CONCLUSIONS

The measured transmission characteristics of diamond film Fabry-Perot resonators correspond to theoretical

calculations in a promising manner, but measured Q factors are less than model prediction. This difference may be

due to scattering losses and absorption losses not incorporated into the model. Presently, work is proceeding in

parallel to investigate both possibilities. The modeling calculations will be expanded to include effects of surface

roughness at the interfaces. Also, new device fabrication sequences are being investigated in an attempt to produce

smoother diamond membranes.

REFERENCES

1. D.K. Reinhard and I. Kleindienst: Diamond Films as Optical Coatings for Silicon, New Diamond And

Frontier Carbon Tech., 13, 2000.

2. J.H. Correia, M. Bartek, and R.F. Wolffenbutteh High-Selectivity Single-Chip Spectrometer for Operation

at Visible Wavelengths, IEDM Tech. Dig., IEEE, 1998, 467-470

3. J.H. Correia, M. Bartek, and R.F. Wolffenbutteh High-Selectivity Single-Chip Spectrometer in Silicon for

Operation at Visible Part of the Spectrum, IEEE Trans. On Elect. Dev., Vol. 47, No. 3, MAR 2000, 553-

559

4. W.H. Southwelh Modeling of Optical Thin Films, Vacuum & Thinfilm, May 1999, 17-23

5. H.A. Macleod: Thin-Film Optical Filters, Macmillan, NewYork, NY, 1986

6. R. Kingslake: Applied Optics and Optical Engineering, Academic Press, New York, NY, 1965

7. M. Thomas and W.J. Tropf: Johns Hopkins APL Tech. Dig. 14, p.16, (1993)

8. J. Verdeyor: Laser Electronics, p. 131, Prentice Hall, Englewood Cliffs, NJ, 1989

9. W. Kern and D.A. Puotinen: Cleaning Solutions based on Hydrogen Peroxide for use in Silicon

Semiconductor Technology, RCA Review, 187-206, June 1970

10. M.J. Ulczynski, B. Wright, D.K. Reinhard: Diamond-coated glass substrates, Diamond and Related

Materials 7, 1639-1646, 1998

NASA/CP--2001-210948 185 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

SURFACE FLASHOVER CHARACTERISTICS OF

POLYCHRISTALLINE DIAMOND AND DLC DIELECTRIC THIN

FILMS IN VACUUM: A REVIEW

Hulya Kirkici

Auburn University

Electrical and Computer Engineering Department, 200 Broun Hall

Auburn AL 36849, USA

ABSTRACT

This work presents a discussion on surface flashover phenomena in vacuum and electrical

insulation in space. Experimental results of vacuum surface flashover characteristics of new dielectric

materials such as polycrystalline diamond and diamond like carbon (DLC) to be used in space

environment as protective coating or as dielectric materials. Specifically, experimental results of

voltage and current characteristics and optical emission during the surface flashover of diamond like

carbon (DLC) on silicon and polycrystalline diamond on silicon are discussed. The thin film samples

used in the experiments were polycrystalline diamond on silicon (D/Si), diamond like carbon on silicon

(DLC/Si), and silicon dioxide on silicon (SiO2/Si) substrate, free standing polycrystalline diamond

(PD), single crystal diamond (SD), quartz, and Teflon TM. We also used DLC film deposited on silicon

dioxide, which acted as the interface between the DLC and silicon substrate (DLC/SiO2/Si) . The

results are compared with the surface flashover characteristics of other materials such as quartz and

Teflon TM.

INTRODUCTION

Electrical insulation of high voltage (HV) power systems to be used in space environment is

currently a major research area because of the interest in HV, high-power systems intended to support

advanced space vehicles and modules such as earth orbiting satellites, inter-satellite modules, and

interplanetary or intergalactic vehicles. As technology advances and new materials developed, there is

a growing need to study the electrical integrity of these materials in space environment. These new

materials can be either used in the construction of electronic components as dielectric material, or can

be used as protective coating of the sub systems. In either case they are exposed to harsh space

environment such as ionospheric plasmas, high-energy particle flux, solar flare, galactic or intergalactic

micro-meteorites, atomic-oxygen, thermal variation, and high energy UV or other radiation effects.

NASA/CP--2001-210948 186 August 6, 2001
Furthermore, thespaceenvironmentis in realitya vacuumenvironment.Therefore,the

electricalinsulationin spacerequiresfurtherconsiderationthantheearth-boundsystemsoperatedin

atmosphericpressures. Onemajorproblemwiththeelectricalsystemsoperatingin spaceorvacuumis

thedifferentialchargingoftheindividualcomponentsduetotheinteractionof spacecraft withthelocal

spacenvironment anddueto thedifferencein dielectricpropertiesof thematerialsusedto construct

thesesystems.Thisdifferentialchargingeventuallyresultsin surfaceflashoverorsurfacebreakdown

of thematerials,thuslossof highvoltageisolationandpermanent damagetotheelectricalsystemsof

subsystemsdueelectricaloverloading[1]. Consideringthespaceenvironmental effectsandnew

advancedmaterialcandidatefor usein space,wehavetestedelectricalcharacteristicsandstudied

surfaceflashovercharacteristics of severalsynthetic[2-4]materialsconsideredtobeusedasinsulating

and/orprotectivecoatinginspacepowersystems.

SURFACEFLASItOVER

Surfaceflashoverorsurfacedischargeacrosssolidinsulatorinvacuumhasbeeninvestigatedfor

manyyearsbyvariousresearchers[5,6]. It iswidelyacceptedthatthesurfaceflashoverphenomena

acrossolidinsulatorssupportingthehighvoltageelectrodesisduetothefieldemissionof electronsat

thecathode"triplejunction"whichistheintersectionof electrode, insulatorandvacuuminterfacewhen

theappliedfieldexceedssomethresholdvalue.

Besidesthemaincauseof surfaceflashover,othercontributingfactorsarelistedasthetypeand

surfaceconditionof thedielectricmaterial[2],thetypeandsurfaceconditionof theelectrodematerial,

thegeometryof theelectrodes, andoperatingenvironment [7]. Thebackgroundandlocalgaseous

speciesandpressuregradient[8]in thevicinityof thehighvoltagesystemarealsoimportantfactors

contributetothesurfaceftashovercharacteristics.

Regardlessof thecontributingfactors,in general,surfaceflashoverresultsfromtherandom

emissionof arelativelysmallnumberof fieldemittedelectronswhicharetheseedelectrons,atthe

cathode/dielectric/vacuum interfaceduetothehighelectricfieldestablishedatthisinterface.Theseseed

electronsincombinationwiththepositivechargebuiltuponthedielectricsurfaceactasaprecursorto

anelectronavalanchetowardstheanodeacrossthesurfaceofthedielectric.Thisenhancedelectricfield

atthetriplejunctionfurtherhelpssustaintheelectronemission.Thusasurfaceflashoveris initiated

whichisfollowedbyasurfacebreakdowninmostcases.Thethresholdpotentialfortheinitiationofthe

surfaceflashoverismainlydependsonthefactorsthatarelistedabove.

NASA/CP--2001-210948 187 August6,2001
EXPERIMENTAL SET UP

Experimental set up to study the surface flashover phenomena consisted of a pump station with a

ultra-high vacuum chamber, a mass spectrometer and an optical multi-channel analyzer with each

computer controlled data acquisition systems, a high voltage dc power supply, and electrical

(current/voltage probes) diagnostics instrumentation. A pair of copper electrodes was placed on the flat

surface of the test sample as shown in Figure 1.

 ath°d iiiii n°de

Figure 1. Top view of Electrodes and sample assembly.

The electrodes were polished to reduce surface irregularities; and the distance between them was

either 1 mm or 3 mm depending on the experimental set. This electrode/sample assembly was places on

a Teflon TM sample holder and was placed in the vacuum chamber in a way that it was seen through a

glass window to allow visual observation and optical diagnostics during the surface flashover

phenomena.

SAMPLES AND SAMPLE PREPERATION

The thin film samples used in the experiments were polycrystalline diamond on silicon (D/Si),

diamond like carbon on silicon (DLC/Si), and silicon dioxide on silicon (SiO2/Si) substrate, free standing

polycrystalline diamond (PD), single crystal diamond (SD), quartz, and Teflon TM. We also used DLC

film deposited on silicon dioxide, which acted as the interface between the DLC and silicon substrate

(DLC/SiO2/Si). All the thin film samples were produced in house by microwave plasma assisted (CVD)

system using proper gas mixtures depending on film type. The surface resistivity of the samples was

measured by a surface probe and determined to be larger than 20 Mf_-mm. The DLC film thickness was

5 _tm and 30 _tm. The film thickness of SiO2 was 2 _tm for both cases, as the interface or as the thin

film. The polycrystalline diamond film was 5 _tm. Both the quartz and free-standing polycrystalline

samples were 1 mm in thickness. All the samples, except the single crystal diamond, were 1 cm by 1 cm

in surface area. All the samples had optical quality smooth surfaces except the polycrystalline diamond

on silicon substrate. One sample at a time was placed on the sample holder and tests were conducted.

NASA/CP--2001-210948 188 August 6, 2001
Thesurfaceconditionof eachsamplewasexaminedbothbyanopticalmicroscopeorscanning

electronmicroscopebeforeandafterthesurfaceflashoverexperiments. Somesamplesseemedtohave

experiencedpermanentsurfacedamage,andthe pictureof damagedregionwasobtainedfor

presentation.

EXPERIMENTALPROCEDURE

Thevacuumchambercontainingthesampleassemblywaspumpedtoahighvacuumlevel,and

waskeptundervacuumfor severalhoursto reducetheoutgassingeffectsonthesurfaceflashover

events,andallthechamberpressurewasontheorderof 10-5Pathroughout experiments. Adcvoltage

wasappliedbetweentheelectrodes, andanoscilloscopewasusedtorecordthesurfaceflashoverevents.

Thebreakdowncurrentandvoltagecharacteristicsof thesurfaceflashovereventswererecordedby

usingahighfrequencyhighvoltageprobeand0.1f_ currentviewingresistor(CVR). Theapplied

voltagewasrampedupsteadily(approximately100V/s)untilaflashovereventinitiatedatthecathode,

andthevoltagewasturnedoff immediatelyafterthesurfaceftashovereventoccurred.Thesame

procedurewasrepeatedseveraltimesuntil asteadysurfaceftashovervoltagewasreached.In some

cases,however,afterafewshots,thebreakdownvoltagebecameunpredictable. Forthesesamples, the

experimentswereterminated, andthesamplewastakenoutofthechamberforexaminationandsurface

analysis.

EXPERIMENTAL RESULTS

It should be noted that the present experiments were conducted using a non-uniform electric

field since the electrodes were rounded at the tips, but flat on the surface contacting the sample's surface

(see Figure 1) to avoid any sharp structure that may enhance the field at the vacuum electrode interface.

With this electrode geometry, the highest electric field was on the surface of the dielectric where the

electrode separation was the smallest. Under this condition it was expected that the filed emitted

electrons from the cathode would follow the field lines while propagating toward the anode, and initiate

a surface ftashover discharge. This behavior was verified by observing the light emission during the

surface ftashover events, and the voltage collapse and the current increase were recorded as the evidence

of the events.

Surface Flashover Voltage Waveform

The waveform data for all the samples were normalized to compare the exact shape and time

evolution of the signals obtained during the surface flashover events. This comparison of voltage wave

NASA/CP--2001-210948 189 August 6, 2001
formsasa functionof time for all the samplesshowedthatthe TeflonTM, quarts, free-standing

polycrystalline and single crystal diamond samples exhibited the same characteristics. The breakdown

event (the voltage collapse) occurred in few (1 to 2 ns) ns. However, the duration of the surface

flashover was considerable long (in the order of 7 to 8 _ts). On the contrary, although the voltage

collapse time was again in few ns, the voltage recovery time (or the breakdown duration) was in the

order of 100 ns for the thin films (DLC/Si, DLC/SiO2/Si, D/Si, SiO2/Si) on silicon samples regardless of

the film thickness.

This result indicated that the film thickness was not a major effect on the surface flashover

characteristics of these materials. However, having the substrate, which is a semi-conducting material,

had played an important role in the duration of the breakdown. Dielectric film on semiconductor

material structure seemed to have appeared to the field as a metal-insulator-semiconductor (MIS) diode,

and depending on the biasing level and/or direction, a mechanism similar to an avalanche breakdown on

bulk breakdown may have occurred. Also, having a dielectric interface, namely SiO2 between DLC and

Si seemed to have increased the integrity of the film to the applied field.

All these thin films including the diamond film, exhibited "conductive" nature after few surface

ftashover events. Furthermore, it was observed that the DLC samples were more vulnerable to the high

electric fields. After few shots, only these DLC samples experienced visible permanent damage on them

at the vicinity of the cathode "foot" print, where the flashover was initiated and where the field was the

highest [3]. These samples were taken out of the vacuum chamber for surface analysis.

It was reported above that the polycrystalline diamond sample did not have an optical quality

surface smoothness. And yet behaved the same as the DLC films in terms of vulnerability to high

voltages, i.e., lost high dielectric properties, and became conducting after a few shots. However, there

was no visible permanent damage on the surface of the sample. Furthermore, the breakdown voltage

waveform and order of magnitude of the breakdown voltage of D/Si were same as the other thin film

samples having optical quality surfaces [2]. Therefore it may be concluded that the surface condition

didn't seem to be an important factor in the initiation or time evolution of the surface flashover events.

Surface Flashover Voltage

Several sets of experiments were conducted with all the samples. Figure 2 shows the breakdown

voltage of thin films (DLC/SiO2/Si, DLC/Si, SiOjSi, and D/Si) for different event number. It is seen in

this data that DLC/Si seems to have the highest breakdown voltage approaching 5.5 kV, in comparison

SiO2/Si has the lowest breakdown voltage (approximately 1. KV). Further, the electrode conditioning

NASA/CP--2001-210948 190 August 6, 2001
phenomena, wherethebreakdownvoltageincreasewiththeeventnumberandreachesto a steady

breakdownvoltagelevelwasobservedforeachsampleunderinvestigation.However,afterafewshots,

asseeninFigure2,(inmostcasesthiswasafterthe3raor4thevent),thesamplesseemedtohavelostthe

dielectric,orhighresistivityproperties, andthebreakdownvoltagesuddenlybecameverylow. It was

alsoobservedthattherewasarelativelylow intensityof lightemission,andthebreakdownvoltage

waveformseemedtohaveexhibited"micro"surfaceflashoverevents.Themaximumcurrenthowever

wasrelativelylow,ontheorderofmA. Thenextandthefollowingsurfaceflashovereventsoccurredat

muchlowervoltagelevels,andthetestswereterminatedafter7 th or 8 th shot for each sample. On the

contrary, although the surface flashover breakdown voltage was much lower compared to the other thin

film samples, the SiO2/Si sample did not experienced similar catastrophic breakdown event, and it

sustained its dielectric characteristics longer than the others.

• DLC/Si:_':_" D/Si .: DLC/SiO2/Si.,'" SiO2/Si]

_o

©

>

6

5

4

3

2

1

0

____ .... ,.... :.... :........ ____-____ ......... 4_

---, .... :.... :.... :.... :.... :....:.....: .... :.....

,,_,,

.............................................. i

, , ' , , , , , ,

-- -:;_c--:!i" - - -:.... :.... :.... ;- - - "-- - : .... :.....

, , , , , , , ,

.i. _i.... .!.... _ ;!: !_ "i.... !........

,, ,, ,, ,, _', ,* ,_, _+_' ,

0 1 2 3 4 5 6 7 8 9 10

event #

Figure 2. Breakdown voltage versus event number for four different thin-film samples. The electrode

separation was 3 mm throughout these experiments.

Figure 3 shows the same plot for free-standing polycrystalline diamond and single crystal

diamond samples. It should be noted that the surface of the free standing polycrystalline diamond was

polished to an optical quality smooth surface.

In Figure 3, the surface flashover voltage seems to remain at approximate value of 9.5 kV,

whereas the single crystal sample's breakdown voltage seems to oscillate around 5.8 kV range with a

NASA/CP--2001-210948 191 August 6, 2001
variationof +3 kV. This was an unexpected result because the surface quality of this sample was same

as the polycrystalline and quartz sample. In a previous work, surface flashover characteristics of a

quartz sample was studies and, it was found that the breakdown voltage increased for the firs few shots

and stabilized at approximately 15 kV range [7]. The polycrystalline sample's surface flashover

characteristics are similar to the quartz's case.

,-. 12

g'_ 8

6

4

2

0

single-xtal :;:;poly-xtal

"-'-'-" " "!-'_'-" ";-'-"" _-" " .... i

;.;, _'__" ;.;,. ,.;'; ;';,. _-'_', ;';'. --

....................... "-'--0- - [ ...............

;"--: 0 •

- -0- • ........ :--0- .... 0- '

0 ° 0000

.............. :.............. ; .......... _;--

...........................................

0 5 10 15

event #

Figure 3. Breakdown voltage versus event number for single crystal and polycrystalline diamond

samples. The electrode separation for these experiments was 1 mm.

Optical Emission Characteristics

Optical diagnostics light emission during the surface flashover events were recorded by several

instruments namely, optical multi-channel analyzer (OMA) to record optical spectroscopy, CCD camera

to record optical image, and photo-multiplier tube (PMT) to record the time evolution of the light

emission. The CCD camera images revealed that the surface flashover starts at the cathode-dielectricvacuum interface as a small "hot" spot and then develops to a full arc propagating across the surface of

the dielectric to the anode. One needs to note that the intensity of the light emission during the surface

flashover of thin-film samples was relatively low for the CCD camera to detect. Therefore, only the

light emission images of the surface flashover of quartz, free-standing and single crystal samples were

recorded.

NASA/CP--2001-210948 192 August 6, 2001
AnOMAwasusedtoinvestigatehenatureof theemissionspectrum.Opticalspectrumof the

light emissionduringthe surfaceflashoverwasrecordedandthe wavelengthidentificationwas

conducted.Theemissionspectrumof theplasmaproducedduringthesurfaceflashoverofquartzsample

showedthattheemissionwasmainlywasdueto thenitrogenspecies.Thefirstnegativeandsecond

positivemolecularemissionbandswerethestrongest emissionbands.In additionto these,therewere

linesbelongingtoSiO2molecule[2]. Theseresultsindicatesthatheresidualgaspresentinthevacuum

chamberplaysanimportantroleintheinitiationof asurfaceflashoverandinthegenerationofaplasma

environment inthevicinityof theelectrodes.Also,thelinesbelongingtoSiO2moleculeindicatedthat

theappliedelectricfiledhavesufficientenergyto excitetheboundmoleculesin thesoliddielectric

material.

TheintensityofthelightemissionfromthepolycrystallinewastoolowforOMAsensitivityto

obtainanyvaluabledata.However, visualobservationoftheemissionindicatedthattheemissionwasin

theblue,blue-violetregionof thespectrum.Thisisin agreement withthedataintheliteraturethatthe

emissionfrompolycrystallinediamondwasreportedtobeabroadbandspectrumwithhighestintensity

centeredat2.85eV(435nm)withtwosidebandscenteredat2.155eV(575nm),and1.68eV(738nm)

[10]. In thatwork,theseemissionlineswereattributedtothedeepdonoracceptorpairrecombination,

andbroadvibronicbandwithazerophononline. Incomparison, wemayconcludethattheemissionin

ourobservationisasimilartotheonesgivenaboveandisaresultofazero-phononlie,whichinturnis

dueto anitrogenvacancycomplex.Consideringthatthefield-emittedelectronsmaypenetrateintothe

material,theeffectmaybeinterpretedaselectroninjectionintothedielectricmaterial,andwouldbe

responsibleforthelightemission.

Opticalspectrumof theemissionfromthesinglecrystaldiamondwasrecordedbytheOMA.

Thisspectrumshowedthatit wasabroadbandemissionwithapeakat550nmandtailtowardsthe

infrared. Themechanismfor this emissionwasconcludedto be assimilarto theoneseenin

polycrystallinecase,wherethefieldemittedelectronsareinfectedintothematerialandexcitedthe

molecularlevelsinthedielectric.However,studiesarecontinuingtoinvestigatehenatureof thelight

emissionfromthesematerialsfurther.

ThePMTdatashovedthatthelightemissionduringthesurfaceflashoverforquartzsamplewas

ontheorderof severalmicro-seconds, almostsamedurationasthevoltagerecoverytimeforthissample.

However,thedurationof thelight emissionfromthefreestandingpolycrystallineandsinglecrystal

diamondfilmswasontheorderof severalseconds, whichwasmuchlongerthanthevoltagerecovery

time. Thisis anotherreasonthatwebelievethatthelight emissionwasresultof excitationaddeexcitationofthemolecules/atoms withinthecrystallinestructure.

NASA/CP--2001-210948 193 August6,2001
Outgassing Effects:

Desorbing or outgassing gas molecules in vacuum can also contribute to the surface

flashover initiation [2, 9]. Any surface of solid (dielectric or other) exhibit forces of attraction

normal to the surface and gas is normally adsorbed on all the surfaces, when left in atmospheric

environment long enough. Under certain conditions of temperature and pressure, this gas is

desorbed and constitutes one of the main sources of gas in vacuum. The material outgassing is

another source of increased gas density in the vicinity of the surfaces. This molecular or atomic

gas density gradient near the dielectric lowers the surface breakdown potential threshold further,

because the breakdown potential depends on the operating pressure.

In search of effect of outgassing and desorbtion, we used a mass spectrometer to

determine the gas species before and after each experimental set of each sample being

investigated. These experiments showed that before the surface flashover events, the gas

composition was mainly of water, nitrogen, hydrogen, oxygen, and carbon-dioxide, which are

concluded to be result of adsorption of atmospheric species in the vacuum chamber. However,

mass spectroscopy experiments showed no traces of gas species that might be by product of the

synthetic thin film dielectric materials (DLC or polycrystalline diamond) after the surface

flashover events. This may be due to the fact that the surfaces area of the samples is much

smaller compared to the surface area of the vacuum chamber, and the amount of outgassing or

desorbing gas from the samples is much less than the detectable level for the mass spectrometer.

On the contrary, mass spectrometer analysis of Teflon TM sample after the flashover events

showed traces of perfluorotibutylamine (PFTBA) and trifluoroethane (C2H3F3) molecules,

which are most likely to be by-products of Teflon TM sample. Also, optical diagnostics of the

plasma generated during the flashover showed similar results as the mass spectroscopy analysis.

CONCLUSIONS

Experiments to investigate surface flashover characteristics of material intended to be used in

space environment conducted. The results showed that polycrystalline and DLC thin films might not be

suitable materials for high voltage application as dielectric material in vacuum when coated onto

semiconductor material. Having a dielectric interface between the DLC and silicon seemed to improve

NASA/CP--2001-210948 194 August 6, 2001
the breakdown characteristics, however it was determined that this also was not sufficient to have a

reliable film for HV applications. All the thin films experienced permanent damage; therefore they were

vulnerable for HV applications. The duration of the surface flashover was on the order of 10s of ns for

thin films. However it was on the order of several micro-seconds for the other free standing samples.

Also, there was almost no light emission during the surface flashover for thin films, where as the light

emission was a long duration event for the other, flee standing films. It was concluded that the light

emission was due to the excitation and emission of impurities within the crystal structure for the

diamond samples, whereas the light emission was due to the background gas for the case of quartz

sample.

ACKNOWLEGEMENT:

Diamond samples used in this experiments were obtained from Dr. Tommy Tzeng, and the DLC and

SiO2 samples were obtained form Dr. R. Ramesham.

REFERENCES

[1] M. Gollor and K. Rogalla, "HV Design of Vacuum-Insulated Power Supplies for Space

Applications," IEEE Trans. Elect. Insul., vol. 28, pp 667-80, 1993

[2] H. Kirkici, M. F. Rose, R. Ramesham, and R. F. Askew, "Surface Breakdown Characteristics of

Polycrystalline Diamond Thin Films in Vacuum," 1994 CEIDP Annual Report, 94-CH3456-1, pp

699-704, October, 1994.

[3] H. Kirkici, "Surface Flashover Characteristics of Diamond-Like-Carbon (DLC) in Vacuum," IEEE

Trans. on Dielectrics and Electrical Insulation, vol. 4, pp 71-78, 1997

[4] H. Kirkici, "Optical mission Characteristics of Polycrystalline Diamond during Surface Flashover

in Vacuum," IEEE Trans. Dielectrics and Electrical Insulation, vol. 4, pp822-26, 1997.

[5] P.H. Gleichauf, "Electrical Breakdown over Insulators in High Vacuum," J. Appl. Phys., vol. 22,

pp 535-41, 1951

[6] H.C. Miller, "Surface Flashover of Insulators," IEEE Trans. Electr. Insl., vol. 24, pp 765-86,

1989.

[7] H. Kirkici, M. F. Rose, R. R. Criss, R. Ramesham, and R. F. Askew, "Optical Spectroscopy

Studies of Surface Breakdown of Polycrystalline Diamond Thin Films in Vacuum," 1994 CEIDP

Annual Report, 94-CH3456-1, pp 606-611, October, 1994.

NASA/CP--2001-210948 195 August 6, 2001
[8] G.Masten,T.Muller,F.Hegeler,andH.Krompholz,"PlasmaDevelopment itheEarlyPhaseof

VacuumSurfaceFlashover,"IEEETrans.PlasmaScience, vol.22,pp1034-42, 1994.

[9] R.A. AndersonandJ.P.Brainard,"MechanismofPulsedSurfaceFlashoverInvolvingElectronStimulatedDesorbtion,"J.Appl.Phys.,vol.51(3),pp1414-21, 1980.

[10] L.H. Robins,L.P.Cook,E.N.Farabaugh, andA. Feldman, "Cathodoluminescence ofDefectsin

DiamondFilmsandParticlesGrownbyHot-Filament Chemical-Vapor Deposition,"Phys.Rew.

B,vol.39,pp13367-77, 1989.

NASA/CP--2001-210948 196 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

OPTICAL PROPERTIES AND RESIDUAL STRESS OF DIAMOND FILMS

Y.N.Sun, J.H.Zhang, J.Q.Li, W.T.Guo, and X.Gao

Lanzhou Institute of Physics, RO.Box 94, Lanzhou 730000, CHINA

ABSTRACT

Diamond is an excellent protective material for optical windows. With large residual stress, crack

through the film or even delamination from the windows can occur, while diamond film is brazed or

coated on the windows. Therefore, it is important to reduce stresses during deposition of diamond film.

This paper describes optical properties and residual stresses of diamond film deposited by

different methods and processing parameters. Diamond films eoated on silicon wafers were produced

by separate hot filament CVD and microwave plasma CVD. SEM micrograph, and Raman spectra of

the films, and optical transmission of free-standing films were measured. The transmission model of

diamond film was improved on the basic of the one of single layer dielectric film, considering the

effects of absorption, dispersion, uniformity, and surface roughness on the transmission. The change of

index of refraction with frequency, free-carrier absorption of graphic component from complex

permittivity in Drude theory and the absorption of C-H stretch were simultaneously involved in the

model. The Levenberg-Marquardt algorithm was adopted to produce nonlinear least square fit. Results

show that the sample produeecl by HFCVD contains refractive index n = 2.35 which is few change with

frequency, surface roughness 0.36 microns, absorption coefficient is lower than 0.1 cm q , and it means

the film has higher purity and small graphic components. The sample by MPCVD contains refractive

index that varies from 2.36 to 2.47 with frequency, surface roughness 0.12 microns, and absorption

coefficient is lower than 0.3 cm q. The Stress State of films was estimated by observing the shift of

Raman peak. The stresses are composed of two sources, intrinsic stress and thermal stress. The thermal

stress of diamond coated on silicon substrate displays compressive. The intrinsic stress in films comes

from deposition method and operation conditions and relates to microstructure, impurity, and relaxation

model. The intrinsic stress usually displays tensile. Observation indicates that residual stress increases

with increasing methane concentration. It is expected to obtain higher optical quality diamond films

with lower stress through controlling processing methods and condition.

Keywords: Optieal properties; Residual stress; Modeling; HFCVD; MPCVD

Corresponding author: Sun Yi-Ning (Y.N.Sun)

Lanzhou Institute of Physics,

P.O.Box 94, Lanzhou 730000, China;

E-mail address: sunyn@lzu.edu.cn;

Phone: +86-931-8267121-5506

NASA/CP--2001-210948 197 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

STATE-OF-THE-ART AND FUTURE TRENDS FOR CVD-DIAMOND COATED TOOLS

FOR CUTTING OPERATIONS

Dr. O. Lemmer, D. Breidt, M. Frank, Dr. T. Leyendecker

CemeCon GmbH, Adenauerstr.20B1, 521146 Wuerselen, Germany

ABSTRACT / INTRODUCTION

In the beginning of the nineties, a world-wide turnover of 450 million dollar was predicted for the year 2000 in the

market " CVD diamond coated tools". Up-to-date, these number is not reached by far. Careful estimations at present

amount to only 10% of this value. The reasons for this are of various nature:

• Underestimation of the complexity in the development of CVD-diamond-coatings on cutting tools.

• The systematic exploitation of new markets for diamond coated tools was difficult and proceeded substantially

slower than expected.

• The understanding, that is necessary for the successful use of CVD diamond coated tools, had to be generated and

communicated to the tool manufacturer, the trade and the end-users

Today the status of the development presents itself as follows: CVD diamond coated tools for cutting operations are

used for machining graphite and green compacts with great success. In this area of applications these tools are

considered as state-of-the-art. End users in this sector normally expect their tool manufacturer to offers CVD-diamond

coated tools in his catalog.

The exploitation of these first fields of applications with economic relevance was possible by a systematic approach

during the product development. A close cooperation of the tool manufacturer, the end-user and the tool coater was

necessary. The knowledge and the know-how in technology and strategy, that were achieved during this product

development, will be transferred and adapted to other, larger sectors of the tooling market and will lead to successful,

innovative CVD-diamond coated tools.

APPLICATION AREAS FOR CVD DIAMOND COATINGS

Conventional diamond tools (PCD, CVD thick coatings) are applied preferentially for cut machining operations of

abrasive materials, e.g. A1Si with a Si-content of >10%, or MMC (__Metal Matrix Composites). Tools having cutting

edges "coated" with a CVD thick coating, may be compared to the property profile of a cutting tool with PCD inserts.

Also for the application of thick CVD diamond coatings only 2-dimensional geometries can be produced with a

justifiable economic expense. If there are very high requirements regarding surface quality of the manufactured

components, PCD tools are successfully applied for machining of soft materials. Here they are in direct competition to

uncoated and conventionally coated tools (CVD, PVD).

Extremely wide fields of application are expected in the future, i.e. CVD diamond coated tools are expected to be

applied very generally. This is exactly the biggest chance, however, also the biggest challenge.

Breaking down the requirements for tools with respect to the different areas of application and commercial aspects

to the coating, one recognizes the qualities that CVD diamond coated cutting tools must have for successful use within

specific application areas. One realizes, that individual aspects of the quality profile may have different priorities

according to the area of application.

Adhesion and wear resistance of the coating play an important role for cut machining of abrasive materials. For cut

machining of softer materials, the coating surface has to be "anti-adhesive", i.e. it must show a low tendency towards the

formation of built-up cutting edges. Furthermore the cutting edge should be as sharp as possible, i.e. the coated cutting

edge should show as low as possible rounding of edges. The machining of soft non-ferrous metals is a typical area of

application where mass production takes place, i.e. the economic efficiency of the tool is of major importance. The

economic pressure is not so high for machining of expensive exotic materials like MMC.

It is evident, that a single CVD diamond coating cannot cover all areas of application in the same manner, since

often completely opposite qualities are required. It is obvious that CVD diamond coatings optimized for specific

applications are preferred if the maximum performance is expected from tools. In plain English:

The user cannot expect one "type" of diamond coating for cutting tools covering all possible areas of application; or

as reverse conclusion: If different diamond coatings are tested against each other, the application and the quality criteria,

e.g. tool life, have to be considered.

NASA/CP--2001-210948 198 August 6, 2001
CVD DIAMOND COATINGS WITH SPECIFICALLY TAILORED PROPERTY PROFILE

If the aimed for property profile has been determined in cooperation with the customer, it can be "modeled" through

the following factors:

"Wear resistance" : Especially for extreme abrasive wear, life time of the tool is proportional to the wear

resistance. This resistance against wear is again influenced through coating thickness, purity and morphology.

"Anti-adhesive" surface : The machining of soft non-ferrous metals requires a coating surface, that counteracts

possible spread-on's or that at least restrains this phenomenon. Therefore a smooth surface is favorable. In this context

especially the crystal size, the coating texture and the pattern of the interface have an influence. Furthermore, the

chemical affinity between coating and workpiece material plays a role. Here, diamond has distinct advantages compared

to most other cutting materials.

"Sharpness" of the cutting edge, cutting edge round-off: The quality of the machined surface and the machining

forces are influenced through the sharpness of the cutting edge of the tool. Coated cutting edges in principle always

show a stronger cutting edge round-off than uncoated ones, however, the coated ones are more resistant against wear.

The coating thickness and the grain size of the diamond coating as well as the interface have an influence on cutting

edge round-off.

"Adhesion" : Sufficient adhesion of the coating is of central importance. This essentially depends on the applied

"Adhesion Technique", on coating thickness and on the coating morphology (crystal arrangement, crystal size, texture)

as well as on the purity of the diamond coating.

These tool characteristics can be tailor by certain techniques of the coater. These are mainly the

• "Coating thickness" (2-60 gm)

• The "coating morphology" (crystal shape, density and texture) that also determines the surface roughness

• The "edge rounding" caused inevitably by the coating itself or by pretreatment techniques on purpose

• The "adhesion techniques"

• Others: post-treatments, multilayers etc.

The Adhesion Techniques

Especially this point is the key technology of a successful application of CVD-diamond coated tools. It describes the

technologies, mechanisms and processes, which guarantee the adhesion of CVD diamond coatings. The following

processes are known (separately or in combination):

• Etching of Cobalt

• Diffusion barriers

• Passivation of Cobalt

• "Rough" interface

As an example for the importance of the last point Figure 1 impressively shows the effectiveness of a rough

interface that is formed by using coarse gained cemented carbide. An outstanding adhesion of the coating could be

reached even with a high coating thickness of 25 gm. Since diamond grows in the gaps of these carbides.

.c_

>

600-

5OO

400-

300-

200-

IO0

0

_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'_'

Fein grained Coarse grained

cemented carbide cemented carbide

Figure 1: Diamond coating on fine / coarse grained cemented carbides

NASA/CP--2001-210948 199 August 6, 2001
TheMMCmaterial AA6082with18Vol.-%SiC-particles wasmilledwithindexableinsertsofthegeometry

M680,withthefollowingmachiningparameters: vc:500m/min.; fz:0,1mm/revolution; ap.:1,5mm.Thedifferenceof

thetoolifedepending ontheusedHMtypeisobvious(ref.3).

TheCoatingThickness

Formachiningof extremelyabrasivematerials, thetoollifeisproportional tothecoatingthickness. In suchcut

machiningprocesses, e.g.machiningofgraphite, it isimportant thatduringthemachiningprocessthereisaratherhigh

coatingvolumeavailableforwear.InFigure2and3thisbehavior isshownfordifferent cutmachiningoperations.

35

30

25

,o 20

-u 15

F: lo

5

0

i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:_:i;i:::!:!_i:i:i:i:i:i:i:i:i:i:i:

iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiNiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii

0 2 4 6 8 10 12

coating thickness [pm]

Figure 2: Dependence of the tool life on the diamond coating thickness for machining of CFC material

12-

10-

t_

t_

8-

0

6-

E 4-

7 2-

0

_,_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_'

......... iiiiiiiiiiiiiiii

 iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii i i i i i i i i i i!iiiiiiiiiiiiiiiiiiii

uncoated CCDia®08 CCDia®15

Figure 3: Dependence of the tool life on the diamond coating thickness for milling of AISi 20 castings

(SDHT1204AEF Vc = 785 m/min, fz = 0,2mm/revolution, ap=0,5mm)

The Coating Morphology

As already indicated, CVD diamond coatings can be deposited in different morphologies and in different roughness.

For machining of strongly abrasive materials, the surface of a crystalline rough diamond coating is smoothened during

the application. If the surface quality of the workpiece plays only an insignificant role, a CVD diamond coating with

distinctive crystalline surface structure is no disadvantage. For some machining applications, in which the surface

NASA/CP--2001-210948 200 August 6, 2001
qualityof theworkpieceandthepreventionof formationof built-upcuttingedgesareof greatimportance, nanocrystallinediamondcoatingsareused.Figure4and5showtwodifferent applications iwhichnano-crystalline diamond

coatingswithaccordingly verysmoothsurfacesareused.

3,5

-_ 3

'_' 2,5

2

1,5

....--

.-°

0,5

length of cut [m]

I I I I

1 20 40 60 80

Figure 4: Surface qualities for turning of Al 99,5 soft, vc=1000 m/min, f = 0,2mm,

ap = 0,5 mm, minimal lubrication technology with nano-crystalline diamonds (CCDia08s)

:.i:!:i:!:i:i:i:i:i:i:i.:

Figure 5: Dry drilling in magnesium-alloy AZ91, vc = 250 m/min, f = 0,2 mm, left uncoated after 30 holes, right

nano-crystalline diamond (CCDial2s) after 300 holes (University of Hanover, Germany, IFW).

Figure 5 clearly shows that especially with dry cut machining the advantages of smooth, nano-crystalline coatings

take effect. While the uncoated drill after 30 holes was no longer usable due to the spread-on's, the diamond coated drill

after 300 holes is still fully operational (ref. 4).

The Edge Rounding

Often a "sharp" cutting edge is demanded by the user. In the least of these applications, the fast wear of the cutting

edge is taken into account. With the machining of extremely abrasive materials the "sharp" cutting edge offers no

considerable wear volume, that could be directly used as advantage for tool life. Figure 6 shows a typical application

case, the treatment of CFRE. It is obvious that the uncoated tool begins the machining process with a "sharper" cutting

edge, however, already after one meter it shows the same round-off as the coated tool. This wear continues rapidly,

while the diamond coated cutting edge is not considerably further worn off in the time period under consideration.

NASA/CP--2001-210948 201 August 6, 2001
0,16

0,14

,o 0,12

"O

E

 g0,1

g_

-o : 0,08

:

._ o_ 0,06

O 0,04

0,02

..... Uncoated

°°°°*

-- CCDia®08s ,,'" :

°°°°

°,"

°°*

°°°

,"

°,"

°°°

,..°°

i

0 Cutting distance [m] 8

Figure 6: Milling of CFRE, vc = 800 m/min, ap=3 mm, ae = 2 mm (IPT)

The formation of burrs at the workpiece is caused through a too large cutting edge round-off. Again, the end-user

would like to use a cutting edge as "sharp" as possible. The reasons are evident: as much wear volume as possible until

wear has caused such a high cutting edge round-off and consequently due to that high cut machining forces and

"pressing" of the cutting edge will result in the formation of burrs.

In the example shown in the next figure, again, a "sharp" cutting edge was originally demanded for the machining of

an A1Si alloy. A cutting edge round-off of approximately 30 gm, that resulted from a round-off of 10 gm due to the

preparation before the coating process and the coating thickness of 20 gm, was considered as unsuitably high.

Machining tests in the framework of the improvement process finally showed, that a minimum thickness of the diamond

coating of 12 gm (CCDial 2s) was necessary to achieve good results.

if)

"l:J

¢-

.m

cO

60-

40-

20-

_`:______________________________________________________________________________________________________

/:_______________________________________________________________________________________________

 iiiiiiiiiiiiiiiiiiiiiiiiiiiiii............ iiiiiiiiiiiiiii

:::::::::::::::::::::::::

uncoated CCDia®20

Figure 7: Milling: G-AISi07Mg, APHT15T3PDR-K88 WK10, vf = 1.500 mm/min, ap = 0,7 mm, n = 6000

revolutions/min, tool life criterion: formation of burrs

NASA/CP--2001-210948 202 August 6, 2001
Evenformachiningof ductileandself-hardening electrolytecopper,adiamondcoatedtoolwitharatherhigh

cuttingedgeround-off ofapproximately 15-20gmcanbeusedwithsuccess(Figure8).Heresurfacequalitywasthetool

lifecriterion.

o0 1Z

Q.

r--

.m

ro

60-

50-

40-

30-

20-

10-

0

`,_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_-_'

 ii i i i i i i i i i i i i i i i i i i i i ...... iiiiiiiiiiiiiiiiii

iiii i iiiiiiiii iiiiiiiiiiiiiiiiiii i

uncoated PCD/DP CCDia®08s

Figure 8: Milling of electrolyte copper (n=20.000 revolutions/min, vc=800m/min, ap: small, ae: small)

SUMMARY

The development of CVD diamond coated high performance tools (brand-name tools) requires a systematic

approach. It is important to go through the partial processes: Marketing, Design, Prototyping and Improvement in order

to reach the requirements for a successful production and market introduction of the new product. During this

development process, the tool manufacturers, coater and if necessary also end-users (pilot users) must work together in

the sense of a close partnership. On this occasion the coater plays the role of a " source of technology " for all concerns

regarding the coating for the new product.

Based on the examples from practical applications it was shown, that especially for the extreme coating material

"CVD diamond" a distinct thorough procedure is necessary in order to use the high potential of this material.

The CVD diamond coated cutting tools just stand at the beginning of their development and their commercial

transformation. With growing know-how on the part of coater and tool manufacturer, a strong increase of the

applications and with it an expansion of the market for CVD diamond coated cutting tools can be expected.

REFERENCES

1. K6nig,W.; Klocke,F.; K6nig,M., Hochleistungzerspanung von Graphit, wt-Produktion und Management 85(1995)

503-509

2. Lemmer; O. Revolution im Formenbau - Graphit kontra Kupfer - HSC-Bearbeitung mit CVDdiamantbeschichteten Werkzeugen statt konventionellem Frfisen"Stahl, 59, 1997

3. Jaspers,SPl_.C.; Dauzenberg,J.H.; Oosterling,J.A.J. Machinability of MMCs using diamond coated tools

Lecture at the EUROMAT 1997 Conference, Symposium AI: Maastricht, Netherlands, 1997

4. Brite Euram project BE97-4586, SAMMI "Safe and Economic Machining of Magnesium castings".

Consortium: Universitfit Hannover IFW (D), Fiat (I), Renault Automation (F), Natan (E), Elf Minearl61werke (D),

Fuchs Lubricants (E), Plansee Tizit (A), CemeCon (D)

NASA/CP--2001-210948 203 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ULTRASONIC END MILLING OF HARD FERROUS MATERIALS

USING DIAMOND TOOL

Masao Murakawa, Hiroyuki Noguchi and Masahiko Jin

Nippon Institute of Technology

4-1 Gakuendai, Miyashiro-machi, Minami-saitamagun, Saitama-ken 345-8501, Japan

E-mail ad&cess: mura@nit.ac.jp

Fax: 81-480-33-7645

Phone: 81-480-33-7618

ABSTRACT

Diamond is known to be the hardest material on the earth. Thus, it shows a superb anti-wear capability when used

as a tool material for cutting almost all work materials except ferrous ones. This is because in the case of ferrous

work materials diamond can be decomposed and diffused very easily into the work materials because of the cutting

heat generated when diamond contacts ferrous work materials.

To solve this problem and fully utilize the hardness of diamond even for ferrous work materials we have attempted

to utilize the effects of ultrasonic vibration applied vertically to the rotating shaft of an end milling tool so as to

increase the lubricity and/or cooling capability of the diamond tool and thereby reduce the involved temperature over

the rake face of the tool cutting edge. By this means it was unexpectedly found that the life ofa sintered diamond end

mill tool used to cut hard ferrous materials could be more than ten times that of a WC-Co alloy tool coated with

TiA1N which is believed to be the best coating material at present for cutting hard ferrous work materials.

Keywords: diamond tool, DLC, ultrasonic end milling, hard ferrous materials

INTRODUCTION

Recently, the requirements for machining tools such as those in ball end milling used, for example, in die making,

have become more severe, particularly in terms of their lifetime, because an end mill tool having a shorter lifetime

would result in the necessity for more frequent tool changes which will inevitably accompany degrading of the work

surface or dimensional accuracy. In an effort to solve this problem, the tool industry has provided various anti-wear

tools such as diamond tools for aluminum work materials and coated WC-Co alloy tools such as TiA1N ones for

ferrous work materials. However, even the best solution of the TiA1N-coated WC-Co alloy end mill tool is not

sufficient when cutting hard ferrous materials such as pre-hardened or hardened steel. In view of this situation it will

be very advantageous if we can utilize the hardness of diamond when we approach the cutting or ball end milling of

hard ferrous materials, although it has been generally believed to be impossible to cut ferrous materials with diamond

tools due to the ease with which diamond composed of carbon is decomposed or diffused into the work material with

the heat generated by its contact with ferrous material. The only countermeasure against this decomposition may be

the use of ultrasonic vibration applied in such a way that there is a gap or a rest period during which the tool is not in

contact with the work material and a cooling medium such as air or lubricant can be introduced into the gap, thereby

preventing the harmful increase in temperature of the tool cutting edge. This countermeasure has been applied in the

field of simple lathe cutting (refs. 1 to 2). However, it has never been applied in the end milling operation.

EXPERIMENTALS AND RESULTS ON ULTRASONIC FRICTION TEST

When ultrasonic machining or cutting is performed, ultrasonic vibration is applied along the cutting direction since

this direction is the most efficient one, in that the machining force can be lowered the most. However, this inevitably

limits the practical maximum cutting speed available for ultrasonic cutting to about 30 m/min. Also, in the case of

end milling, if we apply the vibration along the cutter rotation direction, the maximum peripheral speed available for

NASA/CP--2001-210948 204 August 6, 2001
availableforultrasonicmillingwillbeabout30m/min.Thismeansthatif weusea3-mm-diameter ballendmill,

theallowablerotational speedwillbeonly3200rpm,whichissignificantlylessthanthedesiredrotational speedfor

high-speedmilling( atleast10000rpmbutpreferable morethan40000rpm).In viewofthislimitationinthe

conventional "rotational"ultrasonicendmillingoperation, in thisstudywehaveattemptedto applyultrasonic

vibration(u.v.)totheendmill inaverticaldirection, whichwill eliminatethelimit ontheavailablerotational

speedofendmillingoperation. Sincenoonehaspreviouslyinvestigated theeffectsofultrasonicvibrationapplied

verticallytotheendmilltool,wehavefirstinvestigatedheseffectsusinganultrasonictypefrictiontesterspecially

designedtoinvestigatethetribological propertiesofmaterialssubjectedtovertical u.v.(ref.3).

Figure1showsaschematic oftheultrasonicfrictiontester.WC-Coalloyballs(coatedoruncoated) were

pressedagainst arotatingdiskplateattachedto theendofanultrasonichornvibratedverticallyat20kHz.The

rotatingdiskchosenwasaustenitestainlessteel(JIS:SUS304), knowntoeasilyadheretomosttoolmaterials. The

lubricant chosenwasanonrinsingoil whichhasaverypoorlubricatingcapability, butit isveryenvironmentally

friendlysincethereisnoneedtowashawaythelubricant aftertheformingormachiningprocedure. TheWC-Co

alloyballstestedwereanon-coated one,aDLC-coated one[bymeansofthePSIImethod(ref.4)], anda TiA1Ncoatedone.

Figure2showsanexampleofthefrictiontest.Ascanbeseen,theballcoatedwithDLCfilmshowedavery

lowandstablefrictioncoefficient valueevenwithoutheadditionofu.v.,but showedanevenlowerfriction

coefficient valuewithoutanydelamination ofthecoatingwhentheu.v.wasapplied.Thisresultledustoanticipate

thatheadditionofvertical u.v.asacountermeasure against theharmful effectsinthecaseofferrousmaterial might

beeffectiveevenforcuttingtoolscoatedwithDLCorthosefabricatedfromsimilarmaterialssuchasdiamond.

"---.___;__) .1.)l!rasor_t,: vibrot;,,Jt ht)J:t

J i W : J-R_tatb_,e, d_sk / _,

I| i lL.-r .u / / \

,____Lt r17.-.-.7.7=7-.-.-.-.-.-._- r1-1-=-.-.-.- -lZ_C-t:L_. , ,_ Y ._.K

A ti:!! ; [lie a , :

iii .'.'. " i......... _? ??..i _, , '

:+> i . +:.: ', / ,

._a) V_roole viev¢ (b; A-A v;,Jv.

Figure 1. Schematic view of the ultrasonic vibration type friction tester

With u.v. (ultrasonic vibration)

Without u.v.

,_= 0.25 /

"u 0.2 .._. _,

0.15

eo

"_ 0.1 ...............................

005 :::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

.... ., • _.. •• .,. • .,.......,,....!.!..,....!._...!.......!._...!,?, •

0 I I" ' ' ' I "'" ' ' I '

0 20 40 60 80

Sliding distance (m)

Figure 2. Tribology test results [Balls coated with DLC(PSII) against austenite stainless steel]

NASA/CP--2001-210948 205 August 6, 2001
EXPERIMENTALSAND RESULTS ON END MILLING WITH DIAMOND TOOLS

Figure 3 shows a schematic of the ultrasonic end milling apparatus. The u.v. units (2) and (3) were attached to

the spindle of the machining center. The end mill was mounted to the horn (3) and used to cut the pre-hardened steel

work material. The end mill tools chosen were a DLC-coated WC-Co ally tool, a CVD diamond-coated one, TiA1Ncoated one and sintered diamond tool. Table 1 shows details of the tool, work material and the cutting conditions. As

is evident from the Table, the conditions are those for the finishing end milling operation.

Referring to the results, the DLC-coated tool showed a premature wear out and/or delamination over the cutting

edge, regardless of the addition of u.v., although there was no adherence observed onto the work material. This is

probably due to a severe cutting load or stress which the DLC coating could not withstand. The CVD diamondcoated tool showed similar results in that the coating delaminated prematurely.

Figure 4 shows a graph plotting the cutting distance versus wear amount over the flank faces of the WC-Co

alloy ball end mill tool coated with TiA1N and sintered diamond tool with the addition of u.v. As is clearly evident in

the figure, the TiA1N-coated ball end mill tool has gradually increased its flank wear, reaching the end of its lifetime

at a cutting distance of 25 m where the wear amount of the tool flank reached 40 gm, resulting in an unacceptable

surface roughness of the work material. In contrast with this, in the case of the sintered diamond end mill tool, the

flank wear shows a linear wear increase up to about 40 gm until the cutting length of 250 m is attained. In other

words, the lifetime of the sintered diamond end mill tool with u.v. was more than ten times that of the TiAlN-coated

WC-Co alloy end mill tool, whereas without u.v. the same sintered diamond end mill tool showed a very poor

lifetime: even after a small cutting distance of about 2.5 m the flank wear reached about 25 gin, confirming the

widely held belief that diamond is not suitable for cutting ferrous materials.

Figure 5 shows a micrograph of the work material surface and the corresponding surface roughness respectively

for the operation of TiA1N-coated WC-Co alloy end mill and the sintered diamond one at a cutting distance of 25 m

with the addition of u.v.. In the case of TiA1N-coated end mill, the tool cutter marks observed were obscure due to

the severe wear of the cutter, while the diamond tool sustained the initial clear cutter marks accompanied with a

rainbow-colored surface peculiarly observed for an ultrasonically cut surface.

Figure 6 shows a micrograph of the TiAlN-coated WC-Co alloy end mill and the sintered diamond one.

,¢f-%...-

" I '|,

......I.... ....

[,.7 ooi "'--7 "_I) (1) Machining center spindle

" (2) Ultrasonic vibration unit ! ..... ----(2) (3) Ultrasonic vibration horn

L'..- - - -- - - - - -_@_. .._._-_3) (4) End mill

"_g' (._[_ _ - _ _.//-(,4) (5) Vise

_.._ _-i_/f_: I (6) Generator

(7) Work material

d.J ........... L3

Figure 3. Schematic of the ultrasonic vibration end milling apparatus

Table 1. Cutting condition

Tool

Number of teeth

Work material (Hardness)

Frequency f kHz

Rotation speed N min -1

Feed per teeth mm/rev.

(a) DLC-coated WC-Co alloy end mill

(b) CVD diamond-coated WC-Co alloy end mill

(c) TiA1N-coated WC-Co alloy end mill

(d) Sintered diamond end mill

(a):2, (b):2, (c):2, (d): 1

Pre-hardened steel (HRc:40)

60

3OOO

0.05

Axial depth of cut mm 0.2

Radial depth of cut mm 0.1

Type of milling

Lubricant

Down cut

Water-immiscible cutting fluid

NASA/CP--2001-210948 206 August6,2001
45 (C) Sintered diamond end mill (without u.v.)

40 _coated end mill --

3o

25 __ _)Sint reedd!a?ond--

I 20_ 1015 5 1 end mill (with u'v') 2

0 I° r i i i i

0 50 100 150 200 250

Cutting distance/m

Figure 4. Cutting distance versus wear amount over the flank face of (a)

TiAIN-coated end mill without addition of u.v., (b) sintered diamond one with

addition of u.v. and (c) sintered diamond one without addition of u.v.

lgm

Ry: 2.52, Rz: 1.74 Ry: 1.47, Rz: 1.27

(a) TiA1N-coated WC-Co (b) Sintered diamond end

alloy end mill mill (with u.v.)

(without u.v.)

Figure 5. Micrographs of the work

corresponding chart of surface roughness

(Cutting distance of 25 m)

material surface and the

NASA/CP--2001-210948 207 August 6, 2001
•,o,o,oo%..oo.e=w= • _.•.°.o.O..o.,OoO.OoO°..Oo

,°o o, °.oy..,-•-.°, °°°oO°OoO°°,%%.,¢%o_%o_,_%,_,o°%%.

.,,°°°°, %°,,,, .°.,oOo°°O.°°_._°O_OoOo.,,;.,,.°_oo°.%oo°

•.... lmm iiiiiiiii iii ili i iii iiiiiiiii!i 

(arrow view)

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::: ::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

,°°O°°oO°Ot°°°oOo°o°o°o6oO.%%°°OoO°°o°o% °oo.oo°°o..°o°.oo°o° ......o ° °OoOoOo,°,°%%, ..°,°, ,OoOo°°,oOoO°°°O.Oo .......... oo.,o,%°o, o°oOoO,, oo°,°o°,o°o,°_%Oo ............,O°.oOoOo

i'i'?i:?i.'i:i'_:_:_:!:!:???!:i:i:i.i.i'i:i:i'i:i'i'i'i.'i:i: !:i:!:i:!:i_:'.':i:.":i:_,i:!._:?i:_:!:i:i:i:i:!:!:i:i:??_:i:i

,°o°oOoO°%o.o, Oo•,• o•, •, *, O°°o°oOoOo°oOo°o%Oo%Oo%OoO°OoOo%, / .%OoOoOoO°°°°°ooOdoOoO,Oo.,-..°OoO.O.°°..°,*. %*_,°.-,°. •.

•,'°',','°'.'.'.°.°,°o°°',°°°,'°",','°°o°,°°'," • *,',',°°°,'.' ,'o°.'.'.'+'°'.','.'g '," %,,..... .?,°°"",'o" °',

::::::::::::::::::::::::::::::::::::::::: ::::::::::::::::::::::: (with u.v.)

,.%%%%%%%%°,o,,.=°o_ o._ OoOo,o°°°.. ,.°, .°oOd.OoO°O°,o_OoO°Oo•

,.°.%%o°%%%%%o°°_°o. o°.°°.°°.°°o.°... °°°o, °°,o,o°°-°O.OoO°O° .Oo°oOo° iii ii!!iiiiii ii i

l'mm'"n': __;::':'-'" , 50ttm ,

................. I I (arrow view)

(a) TiA1N-coated WC-Co alloy end mill

(without u.v.)

(b) Sintered diamond end mill

Figure 6. Micrographs of the TiAIN-coated WC-Co alloy end mill and the sintered diamond one.

[(a) cutting distance of 50 m, (b) cutting distance of 250 m]

ULTRASONIC END MILLING OF THREE DIMENSIONALLY CONTOURED SURFACES

From the above results it is clear that ultrasonic end milling using a diamond tool is applicable in the cutting of

hard ferrous work material and that it has an adequate tool lifetime. Finally, we checked the adaptability of the u.v.

diamond end mill tool when used to finish cut a pre-hardened steel workpiece having the three-dimensional contour.

Figure 7 shows the configuration of the contoured surface in question. The wear amount in such cases, where cutting

points over the tool edge change constantly, was substantially the same as that in the case of plane cutting operation

where the cutting points remain the same.

Tool feed direction :.!.:iiiiiiiiiiiiiiiiiiii!!iiiiiiiiiiiiiiiii:!........ ;.;!ii::iiii:i2.0. :i:i:i:i:i:i:

.... ! Zoo!feed.direction !ii!ii iiiiiiiiiiiiiiii

iii:!:!ii;ii;

(a) Tool feed: straight line

(cutting distance : 15 m) (b) Tool feed: curved line

(cutting distance : 30 m)

Figure 7. Examples of finished surface for work materials with three dimensional contours

(Work material: pre-hardened steel, Tool : sintered diamond end mil)

NASA/CP--2001-210948 208 August 6, 2001
CONCLUSIONS

In an effort to enable the utilization of the diamond tool for ferrous work materials, we have attempted to apply

ultrasonic vibration to the end milling tool along the vertical direction so as to obtain a periodic gap between the tool

tip and the work material surface, thereby enabling the introduction of a lubricant or air into the gap to prevent the

harmful temperature rise of the tool cutting edge. The results have shown that a sintered diamond end mill tool with

the addition of ultrasonic vibration used to cut hard ferrous material (pre-hardened steel) could last ten times longer

than the TiA1N-coated WC-Co alloy tool. Furthermore, the diamond tool assisted by the ultrasonic vibration

demonstrated an equally good performance in terms of tool wear against a pre-hardened steel workpieces with threedimensional contours to be finished. Future research will involve further improvement of lifetime of the

ultrasonically vibrated diamond tool so as to enable the cutting of even harder materials such as hardened steel.

REFERENCES

1. M. Murakawa, M. Jin: Turning of beta-titanium alloys by means of ultrasonic vibration, Transaction of

NAMRI/SME, vol.26, 1998, 153.

2. M. Jin, T. Watanabe, M. Murakawa: Study on prevention of chipping of tool edge in ultrasonic vibration cutting

(in Japanese), Journal of JSPE, vol.65, no.12, 1999, 1813.

3. M. Murakawa, H. Noguchi, M. Jin, S. Kataoka, K. Kato, A. Motoi: Tribological properties of hard coatings

subjected to ultrasonic vibration, Proc. ICMCTF, 2001.

4. J. R. Conrad, J. L. Radtke: Plasma source ion-implantation technique for surface modification of materials, J. Appl.

Phys., vol.62, 1987, 4591.

NASA/CP--2001-210948 209 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Diamond Cutting Tools for Dry Machining Application

Prabhu Ammugam, Ajay P. Malshe* and Steve Batzer

Materials and Manufacturing Research Laboratories (MRL), Department of Mechanical

Engineering, University of Arkansas, Fayetteville, AR

* Contact Information- Tel (501) 5 75-6561 E-mail <apm2@engr. uark edu>

Fine mist of cutting fluid during machining, and associated environmental and operator's

health concerns make environmentally benign machining and manufacturing a major

thrust in US, Europe, Japan and other parts of the world. Further, continued demand for

higher machining yield and hence, high speed machining in combination with

environmentally benign machining makes diamond coating a unique candidate for dry

machining. Diamond research in the past few decades has resulted in few products and a

diamond tool is one of them. Diamond in PCD, thick film and CVD coating (polished

and unpolished) forms offers unique advantages for dry machining. This paper discusses

correlation among diamond tool morphology, machining parameters, nonferrous work

piece properties and particulate emission in dry machining. These findings are important

to gauge the "true" performance of diamond tools for dry machining.

NASA/CP--2001-210948 210 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CMP Pad Dresser: A Diamond Grid Solution

Abstract

James Sung

KINIK Company, Taipei, Taiwan, ROC

Chemical mechanical planarization (CMP) is the enabling technology for densification of interconnected

circuitry (e.g., VLSI) or memory storage (e.g., DRAM) in semiconducting chips. CMP flattens the wafer by

polishing it against a rotating pad. The pad is immersed in slurry that contains microscopic (e.g., 0.15 micron)

abrasive particles suspended in a solution of chemical (e.g., H202). During the polishing process, a diamond

dresser must be frequently employed to remove the polishing debris to prevent the pad top from glazing.

Moreover, the diamond dresser must refurbish continually the pad surface in order to sustain the polishing rate of

the wafer.

A diamond grid that contains an array of diamond grits is an ideal CMP dresser. The diamond grid may

double the life in dressing. It can also slow down the decay rate of the polishing pad. The process stability is also

reflected in lower machine-to-machine variation, and also lot-to-lot fluctuations.

Furthermore, the open structure of the diamond grid could allow the table speed or down load to increase, so

the wafer can be polished even faster than current technology allows. Because the pad is evenly dressed by the

diamond grid, its life can be extended by at least 50%. In addition, because the slurry can be pushed around to

coat the pad surface uniformly, its abrasive concentration may be reduced (e.g., by using a diluted slurry) to

further save the cost. Thus, the overall reduction of CoO (cost of ownership) for the CMP process may be more

than 20%.

Because the load of a diamond grit is proportional to its separation from the surrounding grits. The diamond

grid can avoid excessive loading to any of its grits. Hence, the likelihood for grit to chip or pullout is negligible.

Thus, the diamond grid has reduced the occurrence of macroscratches on wafers from a typically about 100 PPM

to about 10 PPM (oxide CMP on Mirra machine). Moreover, the diamond grid can assure an uniform pad surface

for polishing, hence, the thickness uniformity of polished wafers is greatly improved (standard deviation reduced

from about 5% to 2%), This is true for wafer-to-wafer variations too. Furthermore, the diamond grid can produce

a flatter pad surface with uniform texture. As a result, the surface of polished wafer is also flatter. Thus, the yield

as well as the quality of the wafer is both increased by using the diamond grid dresser.

The diamond grid may also be coated with a thin layer of nano-crystalline diamond to shield it from the

chemical attack by the slurry. This is critical for polishing metal (e.g., W, AI) layers when acidic slurries (pH.

value may be 3 or less) are used, in particular for Cu recipe. By making the diamond grid acid-proof, it can dress

the pad while the wafer is being polished at the same time. As the result of this in-situ processing, the throughput

of the wafer production may be boosted by 20%. Moreover, with the continuous in-situ dressing, the pad does not

have to be overdressed as in the case of ex-situ dressing. As a result, the pad life will increase further. The mild

NASA/CP--2001-210948 211 August 6, 2001
butuniform dressing of the pad will also improve the quality of the polished wafers.

Diamond griddressershave now been used for polishingoxidesand metals in productionby major

semiconductorfabsworldwide. The use of thisrevolutionarydressercan improve significantlythe CMP

technologywiththebenefitsincostofownership,productionthroughput,and waferquality.

Key Words: Pad Dresser, Pad Conditioner, CMP, Diamond Disk, Diamond Brazing, Diamond Coating

NASA/CP--2001-210948 212 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

The Reeiprocative Sawing of Granite

James C. Sung

KINIK Company, Taipei, Taiwan, R.O.C.

Taipei University of Technology, Taipei, Taiwan, R.O.C.

Abstract

Diamond is notoriously difficult to adhere to a metal matrix, hence, diamond grits in saw

segments axe commonly locked in place mechanically. In order to prevent such weakly held diamond

from falling out during the cutting action, diamond saws axe typically used to cut in the same direction

so the metal trailing behind diamond can provide the support against the impact force, such as in the

case of circular saws and wire saws. Multiple blades frame saws with reciprocative motion are

advantageous in cutting large area thin slabs at lost cost. Because poorly adhered diamond cannot cut

both directions under strong impact force, frame saws equipped with diamond segmengs axe used only

to cut soft marble. However, steel blades immersed in the slurry of iron particles are used to drag

along the cutting path to grind away granite slowly (about four days to cut through a block).

Diamond can now be brazed in place by a hard alloy. The strong chemical bond can allow

diamond to cut granite both directions without falling out. Moreover, the protrusion of diamond can

be higher so the cutting rate is increased. Frame saws equipped with brazed diamond segments have

been used successfully to cut hard granite at a speed three times faster than conventional steel shot

frame saws.

The world now employs more than 2000 frames saws that slide granite with iron grits. The

sawing of one square meter using this primitive technology will consume about 6.5 Kg of materials

including 3 Kg of iron grits, 1.5 Kg of steel blade and 2 Kg of lime. If the brazed diamond segment

is used, only 10 grams of diamond and segment would be spent. Hence the latter solution can not

only improve the productivity of the sawing (e.g., by cutting three times faster) and the quality of

granite slabs (e.g. no wavy surfaces with rusty stains), but also do away with massive waste by not

generating it in the first place.

Key Words: Diamond Saw, .Frame Saw, Granite

NASA/CP--2001-210948 213 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

The Casting of CVD Diamond

James Sung & Shiao-Kuo Chang

Kinik Company & National Taipei University of Technology, Taiwan, R.O.C.

Abstract

CVD diamond films have many unsurpassed properties including mechanical hardness, thermal

conductivity, acoustic speed, optical transparency, and chemical inertness. Equipped with such unique

attributes, CVD diamond films have invaded in the market of cutting tools, heat spreaders, tweeter

diaphragms, SAW filters, optical lenses, and electrodes.

CVD diamond films are typically deposited on a flat substrate. The growing surface contains

faceted diamond grains that tend to enlarge with increasing thickness of the film. " Hence, CVD

diamond is intrinsically rough on the growth side of the surface. Sometimes the surface relief may

reach over I00 microns. Because diamond is extremely difficult to machine or grind, the flattening of

CVD diamond is often very hard, if not impractical. In many applications, the post deposition cost is

higher than that for growing CVD diamond.

Many CVD diamond applications rely on using a smooth surface that may be either flat or curved.

Moreover, certain applications calls for CVD diamond formed in complicated shapes. Such surface

quality and/or shape requirements are extremely difficult to achieve, if not impossible .to make.

Hence, no such products are currently available.

As CVD process is derived from a fluid of gas, instead of using a substrate, the diamond film can

be deposite£1 inside a mold. The CVD diamond so formed is similar to a cast. Alternatively, the

CVD diamond can be deposited around a wire, so it resembles a coating process. In either ease, the

metal part (mold or wire) c_ be dissolved in a pool of warm acid (e.g., aqua regia) or by a flux of hot

corrosive gas (e.g., chlorine). The remaining diamond will show the surface that is a replica of the

mold, or a wire (i.e., a tube). In this case, the surface and the shape of the CVD diamond would

duplicate the conditions of the metal part. Hence, no further machining or polishing work may be

necessary.

Kinik Company of Taiwan has developed several forms of CVD diamond products based on this

reverse casting process. Among them are cutting tools with chip breaker, water jet nozzles with long

pinhole, dressers with identical pyramids, etc. This paper will unveil such exotic products currently

under development.

NASA/CP--2001-210948 214 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

AN ANALYSIS OF THE CATALYSIS OF FE, NI OR CO ON THE WEAR OF

DIAMONDS

M. Uemura

Toyohashi University of Technology, 1-1 Tempaku-cho, Toyohashi, 441-8580 JAPAN

ABSTRACT

The oxidation probability of hydrogen of methane with metal oxides of Cu, Ni, Co, Fe, Ti and A1 was discussed

by the calculation of equilibrium partial H20 pressure, and diffusional wear rate of diamond was computed and

compared with experimental results. The calculated partial pressure of H20 suggested that the surface layer of

diamond cutting of Cu, Ni, Co or Fe transformed to graphite. In case of Cu, the graphite film formed on the diamond

tool was supposed to act solely as a solid lubricant reducing the wear of the diamond, because the carbon atoms are

not soluble in copper and do not combine with copper. The calculated wear rate of diamond explained experimental

data very well, when the average temperature of cutting interface was supposed to be 560°C for cutting Fe and

600°C for cutting Ni. These results supported the hypothesis that a dissociation of hydrogen chemisorbed on the

diamond surface was the trigger of the graphitization of diamond and the dissociation of hydrogen was caused by the

deoxidization of the metal oxide and the ability of metal oxide as the oxidizing agent was the origin of the metal

catalysis.

Keywords: diamond, metal catalysis, dissociation of hydrogen, diffusional wear

INTRODUCTION

Though diamonds are very hard and highly wear resistant material, it is known (ref. 1) that the diamonds easily

wear during cutting processes of Fe or Ni. This is called as the metal catalysis on the wear of diamonds. Tanaka and

Ikawa (ref. 2) suggested that this phenomenon is concerned with the transformation of diamond to graphite and the

diffusion of carbon into the metal.

Phaal (ref. 3) reported that diamond surfaces became coated with a layer of amorphous carbon throughout the

temperature range of 650-1350°C at relatively low oxygen gas pressures of 0.4 torr and that the presence of the

surface carbon layer at relatively low temperatures of 650°C-750°C could not be accounted for by a purely physical

phase transformation from diamond to the graphite structure because no surface carbon was formed and no surface

attack was found to take place by the heating of a clean diamond to 1350°C in a vacuum of better than 10 -_ torr for

long periods. Matsumoto, Sato and Setaka (ref. 4) found that graphitic carbon was formed with the evolution of

hydrogen from diamond powder on the heating in vacuum up to 1250°C. Pate (ref. 5) showed that the experimental

evidence of the as polished (111) diamond surface suggested hydrogen termination of the surface and that very little,

if any, hydrogen-bonded carbon existed on the reconstructed diamond (111) surface prepared by annealing up to

1000°C in vacuum.

These results look like suggesting that the dissociation of hydrogen chemisorbed on diamond surface would be

the trigger of the graphitization of diamonds. If that is the mechanism, the dissociation of hydrogen should be caused

by the oxidation of hydrogen, because the graphitization of diamond occurred at relatively low oxygen gas pressures

and at relatively low temperatures of 650°C-750°C. In case of cutting metal in air, it is possible that the

deoxidization of the metal oxide, produced by the oxidation of nascent surface formed by cutting, is acting as the

oxidizing agent and that the oxidizing ability of metal oxide is the origin of the metal catalysis.

The present work treats the estimation of the oxidation probability of C-H bond with the oxides of Fe, Co, Ni,

Cu, Ti and A1, and the calculation of the wear rate of diamonds.

METAL CATALYSIS FOR THE GRAHITIZATION OF DIAMONDS

If a diamond surface is terminated with hydrogen (ref. 4 to 5) and the dissociation of the hydrogen is the origin

of the graphitization of diamond, it may be reasonable to deduce that the essence of metal catalysis is the oxidation

of hydrogen atoms chemisorbed on the diamond surface. Therefore, it must be the key point of the metal catalysis

NASA/CP--2001-210948 215 August 6, 2001
whethertheoxygen, diffusingintotheregionaroundalipofadiamondtool,isconsumed fortheoxidationofthe

nascent surfaceofworkmaterial orfortheoxidationofhydrogenterminatingthediamondsurface.

AssumingthathestandardfreeenergyforformationofmethaneissimilarleveltothatofC-Hbondondiamond

surface, asthevalueofastandardfreeenergyforchemical adsorptionfhydrogenoadiamondsurfaceis

unknown, thepossibleractionprocesses, namelytheoxidationofhydrogenbythedeoxidization ofmetaloxide,can

beexpressed asfollows:

CH4(g)+(2/y)MxOy(s) =C(s)+(2x/y)M(s) +2H20(g) (1)

wherethe(s)and(g)meansolidstateandgaseous matter,respectively. Thechangeofstandardfreeenergyforthis

reactionprocessisshownasfollow:

AG ° = 2 A G° (H20) - A G ° (CH4)- 2 A G° (MxOy)/y (2)

where AG ° (CH4), AG ° (H20), AG ° (MxOy) are the changes of standard free energy at T K for the formation of

CH4, H20 and a metal oxide of MxOy, respectively. Then the ratio between water vapor pressures (arm) and methane

gas pressures (arm) is given as follow:

{P(H20)2/P(CH4)} = 10 -AG°/19"I4T (3)

The water vapor pressures, existing in equilibrium with methane at any pressures, can be calculated by

substituting the standard free energy (ref. 6) into the equation (3). Here there is a problem that what pressures of

methane would be equivalent to the number of C-H bonds on the diamond surface. In order to simplify the matter, we

consider { 100 } crystal face of a diamond single crystal as the lip surface of a diamond tool. Then four hydrogen

atoms would be adsorbed on each {100 } face of a unit crystal cell. As the lattice constant of diamond structure is

0.3567 nm, we can say that about 3x107 hydrogen atoms per lbtm 2 are being adsorbed on the { 100 } face. If we give,

for example, 300 nm for an average gap between the lip surface of a diamond tool and the work surface, namely, the

gap related with the surface roughness of the tool surface and the work surface, we can say that about 3x107

hydrogen atoms are contained in a space of 0.3gmxl_tm 2. Therefore there are about l0 s hydrogen atoms per l_tm 3.

On the other hand, gaseous matter at pressure of 105 Pa contains about 3x107 atoms per lgm 3 at 0°C and about

8x10 _ atoms per lgm 3 at 700°C. Then methane gas at pressure of 105 Pa contains about 1.2x10 s hydrogen atoms per

l_tm 3 at 0°C and about 3x107 atoms per l_m 3 at 700°C, each methane molecule contains four hydrogen atoms.

Though the actual cutting temperature is unknown, the number of hydrogen atoms contained in a unit volume of

methane gas does not change remarkably with temperature. Then we estimated the corresponding methane pressure

to be about 105 Pa.

Table 1 gives the water vapor pressures existing in equilibrium with methane at pressures of 105 Pa. The given

pressures mean that when the water vapor pressures rise above the given one, the oxidation of hydrogen will be

stopped. In case of Cu20, NiO, CoO and FeO, very high equilibrium water vapor pressures are given. Then we can

say that, in case of cutting cupreous, nickelic, cobaltiferous or ferrous material, the chemisorbed hydrogen on

diamond surface will be oxidized and then the diamond surface will be transformed into graphite.

Table 1. E_ uilibrium water vapor pressure for the reduction of CH4 gas of 10 5 Pa by metal oxide, Pa.

300°C 400°C 500°C 600°C 700°C

CuO2 1.2x1012 3.6x1011 1.5x10 n 7.4x101° 4.2x101°

NiO

CoO

FeO

3.5x10 _

6.7x105

5.3x102

2.3x10 -15

9.4x106

1.8x106

3.6x103

9.0X10 -12

1.9x 10 7

.6x106

1.5x 104

4. lx 10 -9

3.4x107

6.3x10 _

4.6x104

TiO2* 4.6x10 -7

A1203 1.2x10 -22 7.5x10 -lg 2.7x10 -14 1.5x10 -n

*Standard free energy is obtained in the temperature

5.4x107

9.9x10 _

1.1xl05

2.0x10 -5

region of 1500K-1940K

In case of TiO2, a fairly low equilibrium water vapor pressure of 4.6x10 -7 Pa is given at 600°C. When we

suppose the area of diamond tool's rake or flank face to be lmm 2 and the gap between the diamond tool and work

material to be 300nm, the equilibrium water vapor pressure of 4.6x10 -7 Pa at 600°C is equivalent to being about ten

pieces of H20 molecules in a space of 3xl0Sjam 3 (300nmxlmmxlmm), because gaseous matter contains about 10 7

molecules per lbtm 3 at 600°C and 105 Pa. Also the equilibrium water vapor pressure obtained for TiO2 at 500°C is

NASA/CP--2001-210948 216 August 6, 2001
equivalent tobeingabout0.1piecesofH20molecules. Thismeansthatevenif noH20moleculeexistedinaspace

of300nmx1minxlmm,theprobabilityofoxidationofchemisorbed hyrogenwillbeverylow.

IncaseofA1203, theequilibriumwatervaporpressures obtainedshowsthathechemisorbed hydrogenothe

diamondsurfacewillnotbeoxidizedduringcuttingprocessofaluminous materialndthediamondwillnotbe

transformed intographite.

CAPABILITY OF COPPER OXIDE OF GRAPHITIZATION OF DIAMOND AND MATAL CATALYSIS

The calculated equilibrium water vapor pressures show that Cu have the most remarkable activity of the

graphitization of diamond. Though there is no report supporting metal catalysis activity of copper, the wear of

diamond is said to become serious when a diamond tool is cutting cupreous material containing hard particles like

Cu-Si alloy (ref. 7). Phaal (ref. 8) suggested that bulk diffusion processes of oxygen through the thick carbon barrier

were controlling the kinetics of etch rate of diamond. Cu is characterized by the property of having no solubility of

carbon.

Considering these data, it become known that the characteristic of copper for metal catalysis can be explained by

accepting the assumption that the continuous or macroscopic graphitization of diamond is caused by the continuous

oxidation of a few carbon atoms on the diamond surface at the diamond-graphite interface and the oxidizing agent is

the oxygen diffused through the graphite film. Then, on the cutting of Cu, even if a thin graphite film were formed on

a lip surface of diamond tool, the thin film will not grow to thick, because copper has no solubility of carbon and

then does not promote the diffusion of oxygen through the graphite film. Therefore the thin graphite film would

solely act as a solid lubricant reducing the wear of diamonds.

In case of cutting of hard particles contained Cu alloy, the surface graphite film will be scraped off. This means

that oxygen can attack the diamond surface. Then the diamond surface will transform in to graphite and the surface

graphite film will wear again by the abrasion of hard particle. Then we can also say that the metal catalysis would be

the promotional ability of the metal atoms diffused in the graphite film for the oxygen diffusion in the graphite film.

MODEL OF DIFFUSIONAL WEAR RATE OF DIAMOND TOOL

When we apply the above hypothesis to the cutting of Ni, Co or Fe with a diamond tool, the graphite film

formed on the diamond surface will be removed by diffusion of carbon in to the work metal or the ejecting chip. The

removing speed would be very fast, because the diamond surface is being always in contacting with a nascent surface

of metal. That is to say, the diffusion of carbon atoms in metal is controlling the wear of diamond.

In case of cutting, it would be a very short period while the diamond surface was in contact with a machined

surface or a chip surface. Then we assume one-dimensional diffusion. In this case, Fick's law is represented by

3C(x, t)/Ot = D 32C(x, t)/Ox 2 (4)

where D is the diffusion coefficient, x is the diffusion length, t is the diffusion time and C(x,t) is the carbon

concentration at a distance x and at a time t. Supposing carbon concentration at the contacting interface was a

solubility limit of carbon in the metal, Co, equation (4) can be solved under the following boundary conditions,

C(x,t) = 0, at t = 0 and C(x,t) = Co, at x = 0. Then we have

C(x,t):Co{1-2/_/-_I[/2US;e-y2dy} (5)

Figure 1 shows the contact model between a lip surface of diamond tool and work surface or chip surface. In this

figure, the diffusion time, t, corresponds to the contacting period between the diamond surface and a certain point of

the work surface. Then the diffusion time is given by

t = 1/V (6)

where 1 is the contact length of the diamond tool along the cutting direction and V is the cutting speed. Then the total

number of carbon atoms, diffusing into the machined metal or chip, can be shown as follow:

NdifC : LBA_73 So C(x,1/V)dx (7)

where B is the contact width, L is the total cutting length and AcN is the neighboring atomic distance of the metal.

Then, LB is the whole superficies of machined surface or chip and LBAd_ is the total number of the metal atoms on

the area of LB.

NASA/CP--2001-210948 217 August 6, 2001
Diamond _ _

.......i........... ",

Figure 1. Contact model between diamond tool and work.

On the other hand, the total number of carbon atoms, which composed the wear volume of diamond tool, can be

written as follow:

Ndia = B1ZA_3 (8) C

where Z is the wear depth, B1Z is the wear volume of diamond and AdD is the neighboring atomic distance of

diamond. If the total wear volume of the diamond were caused by the diffusion of carbon atoms into the work

material, we obtain:

N dif _V dia

c = "*C (9)

Then the wear depth of diamond can be shown as follow:

Z=TtYgd )IoC° L(AdD) 3 {--2_f2_e-y2_x (10)

Here the diffusion coefficient is given by

D = Aexp(- Q/1.99T xl0 -3 ) (11)

where A is a coefficient, Q is the activation energy and T (K) is the average temperature at the cutting interface.

COMPARISON OF CALCULATED WEAR RATE WITH EXPERIMENTAL RESULTS

Ikaw and Tanaka (ref. 1) performed cutting tests using conical diamonds as cutting tools. As their results are useful

to verify our calculation, we calculated the wear rates of conical diamond tools. When a conical diamond was worn

into a truncated cone, the wear depth of cone is given by

Z = d/2tana (12)

where d is the diameter of the upper end face of the truncated cone and _ is the semi-apex angle of the cone. In this

case, 1in equation (10) corresponds to d, namely d = 1. Then

1 = 2Z tang (13)

We call 1 wear width after Ikaw and Tanaka's definition (ref. 1), hereafter. If T, V and _ are given and an initial value

of wear width, 10, or wear depth of the truncated cone, Z0, is chosen, we can calculate a very short wear depth

increase, AZ1, by putting l0 and AL, a short cutting distance, in equation (10). Then putting AZ1 in equation (13), we

get the wear width, 11, after a short distance cutting of AL:

11 = 2(Z 0 + AN1) tang (14)

Then putting 11 in equation (10), we get AZ2. Repeating these calculations, we get the wear width as a function of

the cutting distance. Table 2 gives experimental and calculative conditions. Table 3 gives solubility limits of carbon,

Co, in the metal obtained from phase diagram (ref. 9). Table 4 gives coefficients, A, activation energies, Q, and

neighboring atomic distances, AdM and AdD.

Figure 2 contains the plot of theoretical expression of wear length of the conical diamond tools on the cutting of

Fe and Ni, and experimental results (ref. 1). In this calculation, the temperatures of cutting interfaces are chosen

560°C for Fe and 600°C for Ni to give the best fit results.

NASA/CP--2001-210948 218 August 6, 2001
Table2.Experimental and calculativ conditions.

Experiment fief. 1) Calculation

Conical diamond tool cutting Truncated cone diamond tool cutting

Apex 110 ° 110 ° angle

Nose radius

Cone axis

Cutting <100> direction

Cutting speed

Depth of cut

About 5 lam

<111>

1800 m/min

Apex angle

l0

Cutting speed

AL

1 _tm

1800 m/min

2 gm 0. lm

Coolant No

Table 3. Solubility limit of carbon (ref. 9), at.%.

300o C 400°C 500°C 600°C 700°C (573K) (673K) (773K) (873K) (973K)

Ni 0.40 0.49 0.57 0.65 0.74

Fe 0.086 0.102 0.121 0.139 0.163

Table 4. Diffusion coefficient and neighboring atomic distance.

Diffusion coefficient (ref. 9)

D = Aexp{-Q/(1.99T/lO00) }

A, cm2/s Q, kcal/mol

Ni 0.1 33.0

Fe 6.2x10 -3 19.2

Neighboring atomic

distance, nm

Diamond 0.178

Ni 0.2491

Fe 0.2481

80

THEORY 0

7o

O Fe X Ni

60 -k9 560oc for Fe IKAWA and TANAKA DATA 1_,-"'_ o

50

C for Ni

30 _ v

0 I I I I I

0 3_0 20 30 40 50 60

Cutting distance L, m

Figure 2. Comparison between theoretical and experimental (ref. 1) wear width of conical diamond tool as a

function of cutting distance.

Fig. 2 demonstrates a quite satisfactory agreement between calculated wear rate and experimental data. But the

temperatures of the cutting interfaces seem to be relatively low for the cutting done under the dry condition. These

results can be explained by the high thermal conductivity of diamond. Furthermore, though Ni shows higher thermal

conductivity than Fe, the interface temperature chosen for Ni is higher than that for Fe. Miyoshi and Buckley (ref.

10) reported that Fe shows lower friction coefficient than Ni, when in contact with diamond in vacuum. Therefore,

we can say that the thermal conductivity of diamond should be controlling the cutting interface temperature and the

higher temperature of cutting interface for the cutting of Ni would be caused by higher calorific value due to higher

friction coefficient.

NASA/CP--2001-210948 219 August6,2001
DISCUSSION

The proposed metal catalysis mechanism is based on the following five assumptions:

(1) The trigger of the graphitization is the dissociation of the hydrogen chemisorbed on the diamond surface.

(2) The dissociation of hydrogen was caused by oxidation.

(3) The oxidation of hydrogen is controlled by the standard free energy for the formation of metal oxide and

consequently by the water vapor pressures around the diamond.

(4) The attractive forth of positive metal ion diffusing in the graphite film promotes the diffusion of oxygen.

(5) The macroscopic graphitization results from continuous partial oxidation of diamond surface.

Among these assumptions, the assumption (1) is supported by the experimental results that hydrogen desorption

temperatures were consistent with the observed graphitization temperature (ref. 4) and very little hydrogen-bonded

carbon existed on the reconstructed diamond surface prepared by the annealing in vacuum up to about 1000°C (ref.

5). The assumption (2) is supported by the experimental results that carbon film was formed at relatively low oxygen

gas pressures and at relatively low temperatures of 650°C-750°C (ref. 3), where the standard free energy for the

formation of H20 was lower than that for the formation of CO (ref. 6). The assumption (3) is supported by the

experimental results (ref. 8) that water vapor reduced the oxygen etching rate of diamond by an order of magnitude

below 750°C. Therefore the assumption (4) and (5) seem to be the main problems.

On the assumption (4), Tanaka, Ikawa and Tanaka reported (ref. 11) that the graphite film, formed on a diamond

surface by heating up to 900°C in vacuum, was very thick when it was heated with iron powder and thin when it was

heated with nickel powder. Though the diffusion process of metal in the graphite film is unknown, we can say

diffusion process should be playing an important role, as y-iron has much higher solubility limit of carbon atoms, 5.6

at. %, than Ni, 0.9 at. % at 900°C.

On the assumption (5), partial oxidation of diamond at the diamond-graphite interface can be considered to

have a function reducing the activation energy for phase transformation. If it is a real fact, it's first step would be the

oxidation of carbon forming diamond structure by the oxygen diffused through the graphite film, namely, formation

of _C-O-C- bonding. The next step would be the formation of layer structure with the progress of oxidation, like a

formation of graphite oxide (ref. 12). Then some graphite oxide layers would become free from the diamond

structure. Though there is no evidence verifying the role of oxidation on the diamond-graphite phase transformation,

we believe this is why a lot of CO2 and CO were evolved by the heating of diamond powder (ref. 4).

REFERENCES

1. Ikawa, N. and Tanaka, T.: Wear of Diamond Abrasive Grain in Grinding Pure Iron. J. Jap. Soc. of Prec.

Eng., vol. 37, no. 11, Nov. 1971, 56.

2. Tanaka, T. and Ikawa, N.: Thermal Aspect of Wear of Diamond Abrasive Grain in Grinding Mild Steel.

Bull. Jap. Soc. of Prec. Eng., vol. 7, no. 4, 1973, 97.

3. Phaal, C.: Surface Studies of Diamond. Ind. Dia. Rev., vol. 25, no. 300, 1965, 486.

4. Matsumoto, S., Sato, Y. and Setake, N.: Effect of the Preceding Heat Treatment on Hydrogen

Chemisorption of Diamond Powders. Carbon, vol. 19, 1981,232.

5. Pate, B. B.: The Diamond Surface: Atomic and Electronic Structure. Sufthce Science, vol. 165, 1986, 83.

6. Kubaschewski, O. and Alcock, C. B.: Metallurgical Thermochemistry. Pergamon Press, London, 1979,

268 - 356.

7. Fukuoka, T.: Private Communication, Taiho Kogyo Co., Ltd.

8. Phaal, C.: Surface Studies of Diamond II. Ind. Dia. Rev., vol. 25, no. 300, 1965, 591.

9. Smithells, C. J.: Metal Reference Book. Butterworths, London, 1976.

10. Miyoshi, K. and Buckly D. H.: Adhesion and Friction of Single-Crystal Diamond in Contact with Transition

Metals. Applications of Surface Science, vol. 6, 1980, 161.

11. Tanaka, Y., Ikawa, N. and Tanaka, T.: Thermal Attrition of Diamond. J. Jap. Soc. of Prec. Eng., vol. 36, no.

8, Aug. 1970, 60.

12. Hofman, U., Frenzel, A. and Csalfin, E.:Die Konstitution der Graphits_iure und ihre Reaktionen. Liebigs

Ann. Chem. Vol. 510, 1934, 1.

NASA/CP--2001-210948 220 August 6, 2001
4

DIAMOND SYNTHESIS, MODELING, AND

CHARACTERIZATION

Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

HIGH-TEMPERATURE ANNEALING OF COLOUR CENTRES IN DIAMOND.

Alan T Collins

Wheatstone Physics Laboratory, King's College London, Strand, London WC2R 2LS, UK.

E-mail: _A_ia__n:_C_oJli_n_s__k_c_l:a_,c_._u_'_ Telephone: +44 (0)20 7848 2584 Fax: +44 (0)20 7848 2420

ABSTRACT

Nitrogen-vacancy centres have been created by radiation damage, and annealing at 800 °C, in relatively

dislocation-free type Ia diamonds and type Ib diamonds. Some specimens have been progressively annealed at

temperatures up to 1600 °C in an inert gas or in vacuo; other similar specimens have been annealed at the highpressure, high-temperature (HPHT) conditions used to enhance the colour of natural brown diamonds. In the

diamonds containing (N-V-N) ° centres, a substantial concentration of (N-V-N) centres was formed after a 1 hour

anneal at 1500 °C. The annealing data suggest that this charge compensation may occur when a nitrogen-related

centre known as H lb dissociates.

In those diamonds that were heated from 20 °C to 2300 °C over a period of 140 seconds, at a pressure of

5 GPa, the (N-V-N) ° and (N-V) centres completely, or almost completely, annealed out, and no (N-V-N)

centres could be detected. In brown type Ia diamonds, annealed in this way, those optical centres are present in

medium to high concentrations. The present results therefore indicate that the production of these centres in HPHTannealed type Ia diamonds is a dynamic process with the generation and annealing in competition.

Keywords: point defects, absorption and luminescence spectroscopy.

INTRODUCTION

In this investigation, type Ia and type Ib diamonds have been subjected to radiation damage and annealed in the

temperature region 1000 to 2300 °C. The work has been carried out to reach a better understanding of some of the

phenomena observed when natural brown gem diamonds are annealed at high pressure and high temperature (HPHT)

to improve their colour.

Recent research (refs. 1,2) has shown that when natural type Ia brown diamonds are annealed under HPHT

conditions, the colour changes from brown to yellow or yellow-green. The annealing temperatures employed are of

the order 2000 °C, and rather higher temperatures are required to produce the green component of the colour.

Spectroscopic analysis of the diamonds after annealing shows that the yellow colour is produced by absorption in the

H3 band, and the green colour is a result of combined absorption in the H3 and H2 bands. Both of these absorption

bands can be produced in diamond by radiation damage and annealing, and are a result of vacancies being trapped at

the A aggregate of nitrogen (refs. 3,4). The A aggregate is a nearest-neighbour pair of substitutional nitrogen atoms

(ref. 5) and, on trapping a vacancy, a local rearrangement occurs to produce the N-V-N structure. In type Ia

diamond (in which virtually all the nitrogen is present as aggregates) this defect is present in the neutral charge state,

(N-V-N) °, which is the H3 centre. However, if a sufficient concentration of single substitutional nitrogen is present,

some of the defects will be present in the negative charge state, (N-V-N) , namely H2 centres. (Single nitrogen in

diamond acts as a deep donor, from which a charge transfer occurs to nearby defects.)

An efficient way to produce H2 centres is to start with an irradiated type Ib diamond which has been annealed to 800

°C. The dominant defect in such a specimen will be (N-V) centres, produced by the trapping of vacancies by the

single nitrogen (ref. 6). On annealing at higher temperatures, typically 1500 - 1700 °C, the (N-V) centres migrate

through the diamond and, by means of a vacancy-enhanced aggregation process, both A aggregates and N-V-N

centres are formed (ref. 7). Some of these will be in the neutral charge state (H3) and some will be in the negative

charge state (H2), depending on the proximity of single-nitrogen donors.

NASA/CP--2001-210948 223 August 6, 2001
H2centresarealsoproducedwhennaturalbrowntypeIadiamondsareannealedunderHPHTconditions

(ref.1).Infigure1weseethatsubstantial absorptionhasbeenproducedintheH2andH3bands,andabsorptionat

1.945eVdueto(N-V) centresisjustdetectable. Thedefectresponsible fortheoriginalbrowncolourisnotknown,

butmostspecimens showevidenceofplasticdeformation. It isargued(ref.1)thatatthetemperature usedfor

annealing, somehealingoccursoftherupturedbondsassociated withtheplasticdeformation, soreleasingvacancies

whichbecometrappedattheAaggregates toformN-V-Ncentres. Thishealingprocessalsoresultsinareduction

ofthebrowncoloration. Furthermore, atthesehightemperatures, asmallfractionoftheAaggregates dissociateso

formsinglesubstitutionalitrogen(ref.8).SomeofthesedonateanelectrontotheN-V-Ncentres, generatingH2

centres, andafewtrapvacanciestoproduce(N-V) centres. DissociationoftheAaggregates ismoreprobableat

hightemperatures, andthisiswhythelowerannealingtemperatures producepredominantly (N-V-N)° centres,

whereasannealingathighertemperatures alsoproduces(N-V-N) and(N-V)centres(ref.1).

_=

8

_o

,<

20

15

lO

H2

5

0

H3

Photon energy (eV)

Figure 1. Absorption spectrum of a natural brown diamond after HPHT processing.

Many natural type Ia diamonds contain some nitrogen in the form of B aggregates which comprise four nitrogen

atoms symmetrically surrounding a vacancy (ref. 9). A vacancy trapped at a B aggregate produces an H4 centre

which absorbs in a similar spectral region to an H3 centre. However, at high temperatures the H4 centres dissociate,

producing some H3 centres in the process (ref. 10). Therefore, even though some brown diamonds contain the

majority of the nitrogen as B aggregates, no H4 absorption is detected in the specimens after HPHT annealing.

EXPERIMENTAL

For this investigation a selection of type Ia natural diamonds and type Ib synthetic diamonds was used that had

been previously irradiated with either neutrons or 2-MeV electrons, and then annealed at approximately 800 °C.

Details of the specimens are given in table 1. One set of diamonds was further annealed at 1000, 1200, 1350, 1500

and 1600 °C for nominally 1 hour at each temperature. An inert atmosphere, or vacuum for the 1600 °C anneal, was

used to minimise the catalytic graphitisation of the specimens caused by the presence of oxygen. For the anneals at

1000 and 1200 °C the diamonds were pushed to the centre of the furnace once it had come up to temperature. For the

higher temperatures the specimens were placed at the centre of the furnace at 1200 °C, and the temperature was then

ramped up to the set point once the flow of inert gas, or a good vacuum, had been established.

The second set of diamonds was annealed under HPHT conditions at the NovaDiamond facility in Provo, Utah.

Here the temperature of the diamonds was increased to approximately 2300 °C over a period of 140 seconds, and

then the temperature decreased rapidly as the heating power was switched oft'. During the heating cycle the diamonds

were subjected to a pressure of approximately 5 GPa.

NASA/CP--2001-210948 224 August 6, 2001
Table1.Detailsofdiamonds, irradiationsandtheannealings.

Approx. nitrogen Irradiation Anneal Diamondtype concentration (ppm)

Ia A>>B 600 neutron 1000to1600°C

Ia A>>B 600 electron 1000to1600°C

2xIb 70,140 electron 1000to1600°C

Ia A>>B 600 neutron 2300°C

Ia A>>B 600 electron 2300°C

IaB>>A verylow electron 2300°C

Ia A>>B 450 neutron(highdose) 2300°C

IaB>A 700 neutron(highdose) 2300°C

2xIb 130,180 electron 2300°C

Aftereachanneal, absorptionspectraofallthespecimens weremeasuredonthreeseparatespectrometerso

covercontinuously thespectral regionfrom400nmto25gm(25000to400cml). Formeasurements inthevisible

andnear-infraredgionsthediamondswerecooledto77K.Thisensuresthathezero-phonon linesresultingfrom

electronictransitionsintheseregionsaresharpandwellresolved. Spectraintherange4000to400cm1arehardly

affectedbycooling, andwererecordedatroomtemperature.

RESULTS

AnnealingtypeIaAdiamondto1600°C

Figure2showstheabsorptionspectrumoftheneutron-irradiated typeIaAdiamondafterannealingat1600°C.

Asubstantial H2absorptionhasbeencreated, wherenoneexistedforannealingtemperatures blow1200°C.

(D

¢ 2

8

8 1

..Q

,<

H3

H2

1.5

0 .......

1.0 2.0 2.5 3.0

Photon energy (eV)

Figure 2. Neutron-irradiated type IaA diamond annealed at 1600 °C.

NASA/CP--2001-210948 225 August 6, 2001
OnlyasmallreductionhasoccurredintheintensityoftheH3band,butacontinuumabsorptionuderlyingthe

H3band,andwhichincreasedprogressively towardshigherenergies, hasbeensignificantlyreduced. Thecolourof

thediamond, originallyorange/yellowisnowyellow/green. ItisdifficulttoquantifythechangeintheH3absorption,

becausethetransmission atthepeakofthezero-phonon linewasoriginallytoosmalltomeasurereliably.Forthe

electron-irradiated diamondtheintensitydroppedtoapproximately 90%oftheoriginal value.(Thatcompareswith

adroptoapproximately 30%observedarlierafterannealingfor5hoursat1500°C(ref.10).)

Anabsorptionlieat595nm(2.086eV)canfirstbedetectedinelectron-irradiated typeIadiamondafter

annealingat275°C(ref.10).Thereisnoevidenceforvacancymigrationatthistemperature, andsothecentreat

whichthisabsorptionarisesiassumedtoinvolveaninterstitialtom.Using12-houranneals, theintensityoftheline

continuestogrowuntil800°C,thenisdestroyedbytwofurtherheattreatments at900and1000°C(ref.10).Asthe

2.086eVlineisdestroyed, twonewlinesareproducedinthenearinfraredregion(ref.11);theseareknownasHlb

at2024nm(0.612eV)andHlc at1934nm(0.641eV).Theavailableevidencesuggests(ref.11)thathecentres

givingrisetotheseabsorptionlinesareformedwhenpart,orall,ofthedefectwhichcomprises the2.086eVcentre

istrappedattheAaggregate (toformHlb)ortheBaggregate (toformHlc).

Figure3showsannealingdataforthe2.086eV,HlbandH2linesobservedfortheneutron-irradiated diamond

usedinthepresent investigation. Wesee,asreportedpreviously(ref.11)thatheHlb linegrowsasthe2.086eV

lineisdestroyed, althoughhereafinaltemperature of 1350°Cfor1hourisrequired, comparedwith12hoursat

1000°Cfoundintheearlierwork.Thesignificantewfindingtoemergefromthesemeasurements is thatthe

destructionftheHlbline,onannealingfrom1350to1600°C,isaccompanied byagrowthoftheH2absorption.

>,

6_

E

o

Z

1.0

0.5

, . . . ,

2.086 eV

\

\

\

• . , . . . ,

Hll/\f

\ / ,

800 1000 1200 1400 1600

Annealing temperature (°C)

Figure 3. Annealing data for neutron-irradiated type IaA diamond.

HPHT annealing of type IaA diamond to 2300 °C

Two type IaA diamonds, one electron-irradiated and the other neutron-irradiated, together with two very heavily

neutron-irradiated diamonds, were annealed using the same conditions as those used to change the colour of natural

brown diamonds. The heavily neutron-irradiated diamonds were initially opaque to visible and infrared irradiation,

even after annealing at 800 °C. After the HPHT annealing all specimens had a yellow/orange colour due to the

presence of isolated substitutional nitrogen, but in none of them could any H2 or H3 absorption be detected. Typical

spectra are shown in figure 4. For the heavily neutron-irradiated diamonds, the diamonds types, and approximate

nitrogen concentrations, were determined after the annealing, and are shown in table 1.

HPHT annealing of a type IaB diamond to 2300 °C

After radiation damage and annealing at 800 °C, the type IaB diamond exhibited a weak H4 absorption. A

further anneal at 2300 °C completely destroyed the H4 absorption, but it was still possible to detect an extremely

weak H3 absorption. This is consistent with earlier work (ref. 10) which showed that annealing for 5 hours at

1500 °C substantially reduced the H4 absorption and increased the H3 absorption. The interpretation of this

NASA/CP--2001-210948 226 August 6, 2001
phenomenon isthatheH4centre(4nitrogenatomsplus2vacancies) dissociatesoformanH3centre(2nitrogen

atomsplusavacancy) andotherminorproducts(ref.10).Thediamondusedinthepresent investigationstheonly

onethatstillexhibitedH3absorptionaftertheannealt2300°C.Thespecimenisalsountypical inthatit hasavery

lownitrogenconcentration. It is desirabletorepeatthismeasurement withtypeIaBdiamondscontainingafew

hundredppmofnitrogen.

AnnealingtypeIbdiamondto1600°C

Figure5showstheabsorptionspectrumoftheelectron-irradiated typeIbdiamondafterannealingat1600°C.

Examination ofthedefect-induced one-phonon spectrashowsthat,afterannealingat1350°C,someaggregation of

singlenitrogentoAaggregates canjustbedetected, andthatafterannealingat1600°Casubstantial fractionofthe

nitrogenisintheformofAaggregates. Thespectruminfigure5isthenaswewouldexpect fromtheworkofMitaet

al.(ref.4), withsignificant concentrations of H2andH3centresbeginningto form.Thevacancy-enhanced

aggregation processcanbewrittenas(ref.7)

N-V + N _ N-V-N _ N-N +V

N +V---_N-V

andeachvacancyisbelievedtogoroundthisloopapproximately 100times(ref.7).Somevacancies are,however,

lostintheprocess, andtheintensityofthe(N-V) absorption(zero-phonon lineat1.945eV)infigure5hasfallento

approximately 10%ofthestartingvalue.

30

= 20

g

g

..Q

<

0

Hi]_fo re anneal

2.086 eV )¥ After ._

2.0 2.5 3.0

Photon energy (eV)

Figure 4. Neutron-irradiated type IaA diamond before and after 2300 °C anneal.

6 .....

4

(9

8

•_ 2

g

o

..Q

,(

0

1.0

H2

,_J._.______

1.5 2.0

1.945 eV

H3

2.5 3.0

Photon energy (eV)

Figure 5. Electron-irradiated type Ib diamond annealed at 1600 °C.

NASA/CP--2001-210948 227 August 6, 2001
IrradiatedandannealedtypeIbdiamondsexhibitanabsorptionpeakat1450cm1,knownasHla,associated

withalocalisedvibrational mode. Measurements odiamondsopedwithapproximately equalconcentrations of 14N

and 15N have established that the centre consists of a nitrogen interstitial (ref. 12), and studies of this centre in

diamonds grown from mixtures of 12C and 13C have established that the nitrogen occupies a position midway

between two nearest-neighbour carbon atoms, displaced off the axis between them (ref. 13). Woods (ref. 14) showed

that the intensity of the Hla peak is reduced to half of its maximum value after 4 hours' annealing at 1400 °C. Figure

6 shows annealing data for the type Ib diamond used in the present work, comparing the intensities of the Hla, Hlb

and H2 lines. We note that the growth of H2 is uncorrelated with the decay of either Hla or Hlb.

>,

6_

f,

E

O

Z

1.0

0.5

, . . . , . . . , . . . , . . . ,

800 1000 1200 1400 1600

Annealing temperature (°C)

Figure 6. Annealing data for electron-irradiated type Ib diamond.

HPHT annealing of type Ib diamond to 2300 °C

Two electron-irradiated type Ib diamonds, having an intense pink colour due to absorption in the 1.945 eV band

associated with the (N-V) centre, were annealed using the same conditions as those used to change the colour of

natural brown diamonds. After the HPHT annealing no 1.945 eV absorption could be detected in either specimen;

spectra for one diamond are shown in figure 7. The H2 and H3 absorption systems, produced in similar diamonds

after annealing to 1600 °C (figure 5) have also completely annealed out.

40

30

E

¢ 20

oo

__

g

..o

<

0 . , i

1.8 2.0

• .945 eV"

Before HPHT anneal

After_

, i , i , i ,

2.2 2.4 2.6

Photon energy (eV)

Figure 7. Electron-irradiated type Ib diamond before and after 2300 "C anneal.

NASA/CP--2001-210948 228 August 6, 2001
DISCUSSION

Annealing phenomena in diamond are greatly affected by the temperature used, and the time spent at that

temperature. There is also evidence from this work, and earlier work, that annealing rates are affected by the

concentrations of nitrogen in the diamonds. Furthermore, it would not be surprising if electron-irradiated diamonds

behaved differently to neutron-irradiated specimens. It is therefore difficult to intercompare data already in the

literature, or to compare the present data with earlier work, without apparent discrepancies appearing. Nevertheless

important insights have emerged from this preliminary study.

The results clearly show that H2 absorption can be produced in type Ia and type Ib diamond by irradiation

damage and annealing to approximately 1600 °C (figures 2 and 5). The mechanisms are believed to be different for

these two types of diamond.

For type Ib diamond the annealing produces an appreciable aggregation of nitrogen to form A-aggregates

through a vacancy-enhanced process for which the formation of N-V-N centres is an intermediate step (ref. 7).

Because the diamond still contains a substantial concentration of single-nitrogen donors, then, as demonstrated in

reference 4, some of the N-V-N centres will be in the negative charge state (H2 centres).

For type IaA diamond, irradiation and annealing to 800 °C produces predominantly 2.086 eV and H3 centres.

Annealing to 1350 °C in the present investigation completely destroys the 2.086 eV absorption (figure 3) and the

intensity of the Hlb absorption reaches its maximum intensity at this same temperature. This finding for the neutronirradiated diamond confirms the complementary annealing behaviour observed previously (ref. 11) for the 2.086 eV

and H lb lines, although the behaviour was less convincing for the electron-irradiated diamond used in the present

work. There is currently no detailed atomic model available for the Hlb centre, but the available evidence (ref. 11)

indicates that it comprises all or part of the 2.086 eV centre trapped at an A aggregate.

Figure 3 shows that, for type IaA diamond, as the Hlb line anneals out, there is a corresponding growth in the

intensity of the H2 band. We can speculate that dissociation of the Hlb defect generates isolated substitutional

nitrogen atoms which donate their negative charge to nearby H3 centres. However, it is puzzling that there is no

obvious reduction in the intensity of the H3 absorption as the H2 absorption grows. An alternative explanation is that

the dissociation of an H lb centre directly generates an H2 centre. If that is the case, figure 6 clearly shows that this is

not the mechanism in type Ib diamond, for which a substantial increase in the H2 absorption occurs while the

intensity of the H lb absorption hardly changes.

We have seen for the neutron-irradiated type Ia diamond (figure 3) that a 1 hour anneal at 1500 °C reduces the

intensity of the Hlb line by almost 90 %. If we assume that the centre is dissociating, the decay of the intensity will

follow first-order kinetics, and we can write that the number of centres N(t) at time t is related to the number N(0) at

the start of the anneal by N(t) = N(0) exp (-t / _), where • is a lifetime. The probability P of dissociation is

P = v exp (-E / kBT) where v is the attempt-to-escape frequency, E is a thermal activation energy, T is the absolute

temperature and kB is Boltzmann's constant. Putting 1/• = P and setting v = 3.9 x 1013 Hz, the vibrational

frequency of the diamond lattice, allows us to estimate E = 5.9 eV.

This simplistic approach does not explain satisfactorily the shape of the annealing curve for the Hla centre in

figure 6. This more nearly follows second-order kinetics, for which, using the same symbols as previously, [ 1/N(t) -

1/(N(0) ] -- const, x Pt. In this case an estimate for E can be obtained by using data for two different annealing

temperatures.

There are insufficient data in this preliminary investigation to determine reliably the order of the kinetics or the

activation energies. What is clear, however, is that if the estimated energies are used to determine the concentrations

of the optical centres under study, following an anneal in which the temperature rises to 2300 °C in 140 seconds, in

almost all cases the concentrations will be negligible whether the simulation is performed using first-order or secondorder kinetics. This is consistent with the experimental observation that the H2, H3 and 1.945 eV absorption bands

could not be detected in the irradiated samples following such an anneal.

NASA/CP--2001-210948 229 August 6, 2001
Thedisappearance oftheseabsorptionbandsinirradiateddiamondscontrastsmarkedlywiththeproductionof

thesamebandsinnatural browndiamondsannealedunderthesameconditions(figure1).Herewemusttherefore

concludethatwehaveadynamicequilibriuminwhichvacanciesarebeingcontinuallyreleasedandtrappedatthe

nitrogenastheannealingproceedswhile,atthesametime,theopticalcentresbeingcreatedarealsobeingdestroyed

throughdissociation. Theendresultthereforedependscriticallyonthetimeandtemperature pofile,aswellasthe

propertiesofthestartingmaterial.

Inanearlierinvestigation(ref.1)it wasdemonstrated thatif atypeIadiamond, freefromplasticdeformation, is

subjectedtoanHPHTanneal, noH3orH2absorptionisproduced. However, theannealeddiamondhasayellow

colourbecauseof thedissociationf someaggregated nitrogentoproducesinglesubstitutionalitrogenwhich

absorbsintheblueregionofthevisiblespectrum. Theresidual yellowcolourinthetypeIairradiateddiamondsued

inthepresent investigationsalsocausedbythiscontinuumabsorption(figure4).InthetypeIbdiamondssubjected

to theHPHTanneal, approximately 25%of thenitrogenremainsin singlesubstitutional form,producingan

identical continuumabsorption(figure7).

In thisreportwehaveconcentrated onthemajorabsorptionfeaturespresent intheirradiatedandannealed

diamonds. Wehavealsoobservedotherminorcomponents in thenearinfraredandredspectral regions.For

example, inthetypeIaAdiamonds, thespectrashowedtheHlf andHlglines(ref.15)andmorethan10previously

unreportedlinesbetween5000and11000cm-1,andat644and649nm(1.925and1.910eV),thatexhibitvarious

annealingstagesinthetemperature rangeinvestigated. Theseareassumednottohaveadominant influenceonthe

phenomena observed.

ACKNOWLEDGMENTS

Dataforfigures4 and6wereobtainedbytheauthoratNovaDiamond, Provo,Utah.Theauthorisgrateful to

DavidHallandJoeFoxfortheirinvitationtovisitNovaDiamond, andformakingavailablefacilitiestocarryoutthe

research. Dataforfigures2,3,5and6wereobtainedbyAlexConnorduringhisfinal-year pojectatKing'sCollege

London.

REFERENCES

1. Collins, A. T., Kanda, H. and Kitawaki, H.: Colour changes produced in natural brown diamonds by

high-pressure, high-temperature treatment, Diamond and Relat. Mater., 9, 2000, 113-122.

2. Reini_,I.M.,etal.:Identification of HPHT-_eated yellow to green diamonds, Gems & Gemolog> 36,

2000,128-137.

3. Davies, G.: The effect of nitrogen impurity on the annealing of radiation damage in diamond,

J. Phys. C: Solid State Phys., 5, 1972, 2534-42.

4. Mita, Y.,etal.: Phomchromismof H2 and H3 cen_esin synthetic typeIb diamonds, J. Phys.: Condens.

Ma_eg 2, 1990,8567-8574.

5. Davies, G.: The A nitrogen aggregate in diamond - its symmetry and possible structure, J. Phys. C:

Solid State Phys., 9, 1976, L537-L542.

6. Davies, G. and HamegM. E: Opticalstudiesof We 1.945 eV vibronic bandin diamond, Proc. R. Soc.

(London) A ,348,1976,285-98.

7. Collins, A.T.: Vacancy enhanced aggregation ofni_ogenin diamond, J. Phys. C:Solid State Phys.,13,

1980,2641-2650.

NASA/CP--2001-210948 230 August6,2001
8.Brozel,M.R.,EvansT.andStephenson, R.E:Partial dissociationofnitrogenaggregates indiamond

byhigh-temperature high-pressure treatments, Proc.R.Soc.Lond.A,361,1978,109-127.

9.Loubser, J.H.N.andvanWyk,J.A.:ModelsforH3andH4centresbasedonESRmeasurements,

DiamondConference, Reading, UK,1981.(Unpublished abstracts.)

10.Collins,A.T.:Hightemperature annealingof electron-irradiated typeI diamond, Inst.Phys.Conf.

Series, 46,1979,327-333.

11.Woods,G.S.andCollins,A. T.:Newdevelopments in spectroscopic methodsfor detecting

artificiallycoloureddiamonds, J.Gemm., 20,1986,75-82.

12.Woods, G.S.andCollins,A.T.:.The1450cm-1infraredabsorptioninannealedlectron-irradiated

typeI diamond, J.PhysC:SolidStatePhys.,15,1982,L949-52.

13.Kiflawi,I.,etal.:Nitrogeniterstitialsindiamond, Phys.Rev.B,54,1996,16719-26.

14.Woods, G.S.:Infraredannealingstudiesoftheannealingofirradiateddiamond, Phil.Mag.B,50,

1984,673-88.

15.Kiflawi,I., et al.:Newinfraredabsorptioncentresin electronirradiatedandannealedtypeIa

diamonds, DiamondandRelatedMater.,8,1999,1576-80.

NASA/CP--2001-210948 231 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

DEFECT AGGREGATION AND DISSOCIATION IN DIAMONDS DURING

ANNEALING UNDER HPHT CONDITIONS

Filip De Weerdt

Hoge Raad voor Diamant (Diamond High Council), Hoveniersstraat 22, 2018 Antwerp, Belgium; e-mail:

research @hrd.be

ABSTRACT

Under influence of high pressure (+ 7 GPa) and high temperature (+ 2000°C) point defects in diamond can

aggregate and dissociate. When performing this type of annealing on brown coloured type Ia diamonds the colour

changes from brown to intense greenish yellow. Sample characterisation before and after HPHT annealing with UVVIS-NIR and FTIR spectroscopy indicate a dissociation of a nitrogen containing defect and the so-called IR active

"Amber" centre. Other defect centres like the H3, H2, [N-V]- and [N-V] ° defects can be formed during the

annealing sequence.

Keywords: HPHT annealing, point defect, LW-VIS-NIR spectroscopy, FTIR spectroscopy.

INTRODUCTION

Defect aggregation of defects in diamond can be achieved by applying high pressure and temperature to the

diamond so that impurities can overcome the energy barrier required for the diffusion of the defect. The new defects

created by the annealing strongly depend on the diamond type and defect content.

During irradiation with high energy particles, vacancies are created (GR1 centre with a ZPL line at 741 nm). In

type Ib diamonds with mainly C defects (single substitutional nitrogen) a large number of those vacancies capture an

electron donated by the C centre [1,2]. The GR1 centre is then converted in ND1 centre with a ZPL line at 393 nm.

During subsequent annealing at about 600 to 700°C, vacancies become mobile and can be captured by C centres

themselves, with a resulting N-V centre. That centre can also capture an electron resulting in a negatively charged NV centre with a ZPL line at 637 nm. The neutral charge state, however, is believed to be associated with a ZPL line

at 575 nm [2,3].

The annealing sequence is different for type Ia diamonds. Those diamonds contain nitrogen in an aggregated

state: A and B defects (resp. 2 nearest neigbouring nitrogen atoms and 4 nitrogen atoms surrounding a vacancy)

which convert in H3 and H4 centres respectively when capturing a vacancy. The neutral charge state of the N-V-N

centre has its ZPL line at 503 nm and the negative charge state of the N-V-N centre has its ZPL at 986 nm [3,4].

Massive irradiation damage and annealing at 300°C of type I diamonds results in formation of a centre with a

ZPL at 594 nm [5-7]. It is believed to appear after a change in the Fermi-level due to the capturing of electrons

originating from the C centres who act as deep donors in diamond by electron traps in irradiated diamond. The

nature of the traps is not yet fully understood.

During HPHT annealing of natural diamonds a number of those described centres can be annealed in. Other

centres can be annealed out: dissociation of platelets and A centres [8,9,10]. The components in which the platelets

dissociates are not known.

EXPERIMENTAL SETUP

UV-VIS-NIR measurements were performed using a diode-array spectrophotometer equipped with a grating

with 300 g/mm and a quartz optical fibre in a way that commercially polished diamonds could be measured. The

transmission spectrum of the diamond from 400 to 820 nm is measured that way. By rotating the grating, the

transmission spectrum can be measured up to 1020 nm. Because illumination of the sample is performed before the

grating, the fluorescence and absorption of the sample are measured simultaneously.

FTIR measurements are performed with a commercial Bio-Rad FTS-40 spectrophotometer equipped with a

DRIFT accesory. This allows us to measure FTIR spectra of polished diamonds.

Fluorescence imaging was done with a De Beers DiamondView. This instrument uses a short wave UV lamp

(maximum centered around 255 nm) to excite luminescence in a skindeep portion of the diamond. The image is

captured by a digital camera.

NASA/CP--2001-210948 232 August 6, 2001
Withall thosemeasurement tchniquesthepolisheddiamondismeasuredin onetimewhichrendersthe

absorptionduetoimpuritiesindiamondimpossibletocorrelatewiththeconcentration ofimpuritiesbecausethepath

lengthofthetransmittedlightbeamisimpossibletodetermine.

Photoluminescence spectraweremeasuredwithaRenishawmicro-Raman1000spectrophotometer usinga

HeCadlaser(325nm).

RESULTS AND DISCUSSION

Samples

HPHT annealed diamonds were on research loan from the NovaDiamond company. Those samples were only

studied after annealing. These diamonds are of type Ia with a brown colour before annealing. After annealing, they

became yellow with a strong green tone due to the strong green H3 fluorescence. Annealing conditions reported to us

by NovaDiamond and others were around 2000 °C at a pressure of about 6 Gpa, for approximately 30 minutes

[10,11,12].

Another brown type Ia diamond was cut in half and one part was kept as a reference. The other part was

subjected to annealing at 1900 _+70 °C and 7 GPa for 10 hours in a BARS press [13]. Both diamonds were polished

in plates with approximately the same width. The annealed diamond sample will be hereafter referred to as HRD1,

while the original sample will be referred to as HRD2.

All samples showed strong birefringence in transmitted light illumination when the diamond is placed between

two polarisation filters. That indicates that the diamond was subjected to strong plastic deformation. The brown

colour is attributed to the presence of this plastic deformation [10,14]. This supposition is supported by the fact that

brown coloration in type Ia diamonds correlates with strong plastically deformed diamond (111) layers. In pink type

Ia diamonds a similar feature is seen: sheets of pink colour concentration correlates with strong plastically deformed

regions.

Measurements

The diamond samples deliverd to us by NovaDiamond exhibited large cracks and were not cooled to liquid

nitrogen temperature because of the risk of enlarging the cracks. UV-VIS-NIR measurements clearly showed the

presence of strong green H3 fluorescence. Some diamonds exhibited very weak [N-V]- absorption at 637 nm. In the

NIR spectrum, all samples a very strong H2 ZPL at 986 nm was observed together with a very strong sideband. In

the FTIR measurements a very weak maximum was observed at 1344 cm -1, indicating the presence of C centres. The

so-called "Amber-centre", often observed in brown type Ia diamonds, was not detected. This Amber-centre in FTIR

is a broad absorption band centered around 4170 cm -1.

The HRD1 sample did not display the Amber-centre in the FTIR spectrum, while the HRD2 sample does display

the defect in the FTIR spectrum. The UV-VIS-NIR spectrum at liquid nitrogen temperature of the HRD1 sample

clearly shows an increase of H3 absorbance and fluorescence (see figure 1). H2 was not observed in the HRD2

sample, but it is clearly measurable in the NIR spectra of the HRD1 sample.

©

¢o

c

"_ 0.10

E

6_

C

F--

400

0.05

.

HRD1

5_o 5_o 7_o

Wavelength (nm)

0.6

0.5

0.4

0.3

0.2

0.1

Figure 2. UV-VIS spectra of the HRD1 and HRD2 sample. For reasons of clarity the spectrum is only

displayed in the region between 400 and 720 nm. The transmittance scale on the left is for the spectrum of the

HRD1 sample (black graph) while the right scale is valid for the HRD2 sample (red graph). A strong increase

of H3 fluorescence is clearly visible.

NASA/CP--2001-210948 233 August 6, 2001
Observations madewiththeDiamondViewdeviceshowthatthegreenH3fluorescence is notuniformly

distributedinthediamond, butrelatedtothegrowthhistoryofthediamond(seefigure2).Sampleinhomogenity isa

majorprobleminallsamplesunderstudy,andit isasourceoferrorsinourmeasurements aslocalabsorptioncanbe

largeror smallerthanmeasuredwithourspectroscopic techniques. Asaconsequence, thFermi-level is not

homogeneous indiamond, anda"local"Fermi-level varyingthroughthediamondanddependingonlocaldefect

concentrations shouldbeused.ThefactthatalowerH3luminescence isobservedinsomesectorsofthediamond

canbeduetoalowerH3concentration orduetoarelativelyhigherconcentration ofresidual Adefects, quenching

theH3luminescence.

PLspectraweretakenapproximately atthecentreoftheHRD1andtheHRD2sample.Fromthespectra, it is

clearthathereisasubstantial reductionofthe491nmline,associated withdecoratedsliplines.Thismeasurement

isconsistent wihthedataofref.10.

-.-;.-- -.;:

• o._°

Figure 2. DiamondView image of a HPHT annealed diamond with green H3 fluorescence. It can be observed

that some of the luminescent centres decorate slip lines, while others are dearly confined to sectors related to

the growth history of diamond.

,m

ee40000t I

30000

a an

20000

10000 .

I 491 nm_J

/

\

\

1

20000

400 500 600 700

wavelength (nm)

30000

25000

15000

10000

5000

0

Figure 3. PL spectra of the HRD1 and the HRD 2 sample at 80 K. There is a clear reduction of the 491 nm

line. Oscillations in the spectrum in the region between 600 nm and 650 nm are real, and also observable in

absorption measurements [15]. The H3 fluorescence was very strong in the HRD1 sample and the exposure

time had to be reduced by a factor of 2. The laser beam had to be attenuated to a factor of 10 relative to the

laser beam in tensity used for measuring the HRD2 sample.

Discussion

From the UV-VIS-NIR absorption spectra, it can clearly be observed that there is an increase of H3 defects.

The source of the vacancies is not fully understood but thermally activated, moving or interacting dislocations might

be a source of vacancies [10,16,17]. However, it is calculated that the activation energy for dislocation motion is

about 3,3 eV, but in the presence of other defects like hydrogen, the activation energy might be sufficiantly lowered

[ 18] in order to be thermally activated under the applied annealing conditions.

NASA/CP--2001-210948 234 August 6, 2001
Naturaldiamondsamplesdodisplayalowconcentration ofGR1defectsbecauseofaradiationdamageina

skindeeplayerof thediamond, butitsconcentration is toolowin orderto beresponsiblefortheincreasedH3

absorption/luminescence.

Thefactthathe[N-V]-andtheH2defectsappearafterHPHTannealingofthediamondpointsatthefactthat

donorcentreshavebeenformedinthediamonds, relocatingthe"local"FermilevelaroundtheH3andthe[N-V]°

defects.PossibleelectrondonorsareCcentres, however themagnitude oftheabsorptionduetotheH2defectinthe

diamondsHPHTannealedbyNovaDiamond isveryhighcomparedtotheH2intheHRD1sampleandonlyalow

concentration of Ccentresidetectedinthediamonds. Othersourcesofelectrondonorsshouldbeconsidered.

HPHTannealingconditionsemployedbyNovaDiamond areveryclosetothethermodynamic diamond/graphite

equilibrumline,andthismightfavourtheformationofotherelectrondonors.

CONCLUSION

It isseenfromtheobservations thatH3andH2iscreatedduringannealingconditions, butthesourceofthe

vacanciesisnotyetfullyunderstood. Alsothesourceoftheelectrondonors, givingrisetothepresenceofH2and

[N-V]-centresisnotfullyclearedout.ResultsoftheHPHTannealingdependsonsampleandannealingconditions.

TheAmber-centre intheFTIRspectrumisclearlyannealedoutbytheappliedHPHTparameters. Itcomponents in

whichit dissociates isalsonotfullyunderstood.

ACKNOWLEDGEMENT

I wouldlikethankNovaDiamond Inc.forprovidingHPHTannealedsamplesandalsotheresearchers fromthe

InstituteforMineralogyandPetrography, RAS,Novosibirskfortheannealingofoursamples.

REFERENCES

1. G.Davies, Chargestatesofthevacancyindiamond, Nature, vol.269,6october1977,p.498-499

2. A.Mainwood, Nitrogenandnitrogen-vacancy complexesandtheirformationidiamond, Phys.Rev.B,

vol.49,no.12,15March1994,p.7934-7940

3. K.Iakoubovskii, G.J.Adriaenssens, M.Nesladek, Photochromismofvacancy-related centresindiamond, J.

Phys.:Condens. Matter,vol.12,no.2,17January2000,p.189-199

4. S.C.Lawson, G.Davies, A.T.Collins,A.Mainwood, The"H2"opticaltransitionidiamond: theeffectsof

uniaxial stressperturbations, temperature andisotopicsubstitution, J.Phys.:Condens. Matter,vol.4,no.13,

30March1992,p.3439-3452

5. G.S.Woods, A.T.Collins,Newdevelopments ispectroscopic methodsfordetectingartificiallycoloured

diamonds, J.Gemm., vol.20,no.2,1986,p.75-82

6. G.Davies, H.M.Nazare, Uniaxial splittingsofEtoEtransitionsattrigonal centresincubiccrystals: the

594nmbandindiamond, J.Phys.C:SolidStatePhys., vol.13,no.22,10August1980,p.4127-4136

7. Y.Mira,Changeofabsorptionspectraintype-Ibdiamondwithheavyneutronirradiation, Phys.Rev.B,vol.

53,no.17,1May1996,p.11360-11364

8. M.R.Brozel,T.Evans,R.EStephenson, Partialdissociationof nitrogenaggregates indiamondbyhigh

temperature-high pressuretreatments, Proc.R.Soc.Lond.A,vol.361,1978,p.109-127

9. I. Kiflawi,J.Bruley,Thenitrogenaggregation sequence andtheformationofvoiditesindiamond, Diamond

Relat.Mater., vol.9,no.1,1Januari 2000,p.87-93

10.A.T.Collins,H.Kanda, H.Kitawaki, Colourchangesproducedinnatural browntypeIadiamonds byhighpressure, hightemperature treatment, DiamondRelat.Mater.,vol.9,no.2,February/March 2000,p. 123-

128

11.T.Tempelsman, NovaDiamond introducesnewenhancement, Rappaport, vol.23,no.1,7January2000,

p.1

12.R.Weldon, S.King,NewHPHTdiamondenhancement revealed, Professional Jeweller, 2december 2000

13.Yu.N.Pal'yanov, A.EKhokhryakov, Yu.M.Borzdov, A.G.Sokol,V.A.Gusev, G.M.RylovandN.V.

Sobolev,Growthstructureandrealstructureof syntheticdiamondcrystals,RussianGeologyand

Geophysics, vol.38,1997,p.920-945

NASA/CP--2001-210948 235 August6,2001
14.T.R.Anthony, Restoration of theas-growncolourofgem-quality diamond, DiamondRelat.Mater.,tobe

published

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16.H.S.Leipner,C.G.Htibner,T.E.M.Staab,M. Haugk,A. Sieck,R.Krause-Rehberg, T.Frauenheim,

Vacancyclustersin partiallydeformedsemiconductors, J.Phys.:Condens. Matter,vol.12,no.49,11

december 2000,p.10071-10078

17. D.Fisher, privatecommunication

18.M.I.Heggie,S.Jenkins, C.REwels,RJemmer, R.Jones, RR.Briddon,Theoryofdislocations idiamond

andsiliconandtheirinteractions withhydrogen, J.Phys.:Condens. Matter,vol.12,no.49,11december

2000,p.10313-10318

NASA/CP--2001-210948 236 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FULLERENES AS A CO-CATALYST FOR HPHT SYNTHESIS OF DIAMOND

A. Ya. Vul', S. V. Kidalov,

Ioffe Physico-Technical Institute, 26 Polytechnicheskaya st., St.Petersburg 19402 l, Russia

V. M. Davidenko, V. A. Yashin

Ilyich Abrasive Company, 23 Krasnogvardeiskii, StPetersburg 197342, Russia

S. S. Ordanyan

StPetersburg Technological University, 26 Moskovsky pr., StPetersburg, 198013 Russia

ABSTRACT

The paper reports on a study of the effect of fullerene co-catalyst for high pressure - high temperature (HPHT)

synthesis of diamond from graphite. It has been shown that the fullerene co-catalyst results in the increase of the

yield of the graphite-diamond transformation in 1.8 times.

Keywords: diamond, fullerene, HPHT synthesis, catalyst

INTRODUCTION

One year later of the Huffinan Krfitschmer arc method discovery for fullerene production (ref. 1) the paper of

R.S. Ruoff and A.L. Ruoff with the intriguing title, "Is C60 stiffer than diamond?", has been published (ref. 2). It

was a starting point for research of fullerenes as a material for synthesis of superhard carbon. During the previous

years it has been confirmed that some phases of polymerized fullerenes obtained at about 1000 °C and under

pressures of about 13 GPa indeed have hardness properties that are close to those of the diamond. It has been shown

that 3D C60 polymers have not only high hardness, comparable with diamond, but also unusual plasticity. In some

cases hardness even higher than diamond has been reported (ref. 3). It has also been to attempt to use fullerenes as a

starting carbon material replacing graphite for synthesis of diamond by industrial HPHT method. One might expect

that the partial sp 3 bonding configuration characteristic of C60 would work in favor of diamond synthesis and will

allow reduction of pressure/temperature conditions.

To our knowledge the first patent related to use fullerenes in HPHT diamond synthesis was applied in 1993

(ref. 4) wherein fullerenes was suggested as a starting material. As shown in reference 5, the C60 can be used as the

carbon source to carry out the metal-catalyzed synthesis of diamond under hydrostatic pressure of 6.7 GPa and at

temperatures between 1200 and 1850 °C, no advantage of the replacement of graphite by fullerenes was found.

Furthermore, the yield of diamonds from fullerenes was equal to that from graphite.

We have been using different approach (refs. 6 to 7). Fullerenes were used as a co-catalyst and small amount of

fullerenes was added to graphite with metal catalyst.

EXPERIMENTAL TECHNIQUES

The experiments were carried out with the use of a press with axial force 10 MN in a industrial high-pressure

chamber of a 4.4 cm 3 volume. Graphite in a mixture with nickel-manganese catalyst was used as a starting material.

The heating of the high pressure chamber was done by a.c at 50 Hz in the electric power stabilization regime to the

accuracy with not less than 1%. Industrial chemical extraction of diamond crystal was used. The diamond crystals

are characterized by the size and strength.

NASA/CP--2001-210948 237 August 6, 2001
RESULTS AND DISCUSSION

The catalytic effect of fullerenes has been observed. The effect of fullerenes was not depended on the

method, which was used for their addition. Addition of fullerene soot, made by Kr_itschmer-Huffman arc

method, fullerene mixture (C60 and C70), or the pure fullerene C60 leaded to the same positive effect. The effect

has been repeatedly confirmed at using a current industrial technology of the HPHT diamond synthesis from

graphite under pressure 4.5 - 5.5 GPa and temperature 1200 °C. The effect of the fullerene co-catalyst depends

on a value of pressure. The yield of diamond from graphite is 35% under a pressure of 4.5 GPa and ca 75%

under a pressure of 5.5 GPa. The last value is in 1.8 times more than that without the fullerene co-catalyst.

It would be mentioned, that the mechanism of metal-catalyst influence at HPHT diamond synthesis is debatable

yet (refs. 8 to 9). Many scientists have supported the solvent-catalyst approach, which accepts, that the metal acts as

solvent for graphite and that the solvent has to provide additional properties for diamond synthesis (ref. 8). Under

this approach is accepted that when graphite dissolves in solvent, the bonds of graphite are completely ruptured,

dissolution continues until the molten metal is saturated relative graphite and the saturated solution is supersaturated

relative to diamond. As result, diamond is growing from supersaturated solution in the thermodynamically stable

region of diamond. Withal the model accepts that the dissolved carbon must be positively charged, and the

dissolved carbon forms intermediate compounds (metal carbides) that subsequently decompose to form diamond.

The P-T region of diamond formation at HPHT synthesis is bounded by two distinct lines, namely a graphitediamond equilibrium line and a carbon-solvent eutectic line. Therefore, the P-T parameters vary depending on the

catalyst metal and diamond crystals are formed on the boundaries of the catalyst. The minimum temperature and

pressure are obtained by using the invar system (Fe Ni) to be 4.8 GPa, and 1160 °C (ref. 10).

In this connection, it is important to mention that use of amorphous carbon as a source of carbon instead of

graphite at HPHT synthesis is stopped of diamond synthesis (ref. 5). It means that graphene sheets are needed for the

diamond synthesis and the idea of the solvent-approach that the bonds of graphite are completely ruptured when

graphite has dissolved in the solvent is not correct. As shown in reference 11 the addition of 3 wt % of soot

produced by acetylene burner or hydrocarbons as like as paraffin or naphthalene to catalyst-metal interrupts synthesis

of diamond from graphite.

It can be supposed that C60 dissolves in metal-catalyst at lower temperature than graphite, thus actually

decreasing the reaction temperature. Moreover those C60 molecules, probable their fragments curved graphitic plates

are acting as precursors for diamond growth. The idea that fullerene carbon can significantly improve the nucleation of

diamond at the HPHT synthesis was also mentioned in (ref. 12).

The research was supported by the Russian State Programme "Fullerenes and Atomic Clusters.

REFERENCES

1. Kr_otschmer W., et al.: Solid C60: A New Form of Carbon. Nature, vol. 347, 1990, 354 - 358.

2. Ruoff R. S., Ruoff, A. L.: Is C60 stiffer than diamond Nature, vol. 350, 1991, 663 - 664.

3. Sundqvist B. Fullerenes under high pressures. Advances in Physics, 1999,vol. 48, no.1. 1-134.

4. Nunez-Requiero M.: Utilization de fullerenes comme produits de depart dans la fabrication de diamant. PCT

WO 93/02212, July 10, 1992, National Priority date (France) July 15, 1991.

5. Bocquillon, G., et al.: C 60 fullerene as carbon source for diamond synthesis. J.Phys.Chem., vol. 97, 1993,

12924-12927.

6. Vul' A.et al.: Method for production artificial diamonds. PCT application. Priority date Aug. 22, 1997,

RU1997000114492, Patent RU N2131763.

7. Baidakova, M.V., et al.: Diamond nanocrystals made from fullerene soot. Extended Abstracts and Program. 23 ra

Biennial Conference on Carbon. 18-23 July, 1997. Penn State University. The American Carbon Society, 1997,

354 355.

8. Setaka, N.: Development of Diamond Science and Technology in Japan. In: Synthetic Diamond: Emerging

CVD Science and Technology. Eds. K.E.Spear and J.P.Dismukes. John Wiley&Sons, Inc. NY, 1994, 57 - 90

9. Kanda H., Akaishi M., Yamaoka S.: New catalysts for diamond growth under high pressure and high

temperature. Appl.Phys.Lett., vol. 65., 1994, 784 - 786.

10. Fugunaga, O et al.: Pressure and temperature control in flat-bait type high pressure apparamres for reproducible

diamond synthesis, Diam. Relat. Mat., vol.8, 1999, 2036 2042.

11. Polyakov V.; et al.: Influence of hydrogen on polymorphous transformation of carbon in solution of Me-C under

high pressure. Doklady of Academy of Sciences of USSR, vol.307, 1989, no. 3,602 - 605 (in Russian).

12. Wei B., et al: The transformation of fullerenes into diamond under different processing conditions.

Journ.Materuial Processing Technology, vol. 63, 1997, 573-578.

NASA/CP--2001-210948 238 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Formation of synthetic single and polycrystalline diamonds

Vladimir Prokofievich Poliakov

Federal University of Parana, Department of Mechanics,

C.P. 19011, CEP 81531-990, Curitiba, Parana, Brazil.,

e-mail vladimir@demec.ufpr.br

ABSTRACT

There arc general rcgularities in the formation of single crystal (SD) and synthesized polycrystaUine diamond

(SPD). It is found that the diamond nuclei are formed under high pressure (P) and temperature by the fluctuational

or martensitic transition of the graphite crystallites that are contained in the initial carbon material in the presence of

metallic melt. If the initial carbon material does not contain graphite crystallites of sufficient size (rGcr,, > rDcr,, ) or if

they cmmot pass into the metallic melt then the formation of diamond nuclei is hindered. The basic difference

between SPD formation and SD one is that the initial porous polycrystalline graphite under high P and T far from

the line of equilibrium G _ D is impregnated with metallic melt with high rate (few seconds). The impregnation of

graphite bulk with metallic melt at the total high P in the reaction cell is accompanied by very fast solid phase

transformation of graphite lattice to diamond one (martensite transition). This results in a diamond polycrystalline

composite with inclusions of metallic phase on the grain boundaries.

The mechanism of metallic melt with dispersed refractory particles transfer to the bulk of SPD was realized for

the increase in mechanical strenglh of SPD through the addition of particles A1203 and TiN in the Ni - 15% Mo

alloy.

Key words: Single crystal diamond; Synthesized polycrystalline diamond; Impregnation; Thermal stability;

Mechanical strength.

INTRODUCTION

A major effect of initial carbonaceous materials structure on the formation of SD in the carbon-metal system at

high pressure and temperature for the first time was shown in (ref, 1, 2). It was showed experimentally that the yield

and size of SD depend on the capacity of the initial carbon material for homogeneous graphitization, i.e., for

formation of a graphite structure under heat treatment. According to the colloidal theory of diamond nucleation

(ref. 3), diamond nuclei (sp3 bonds) are formed in the metallic melt (metal or alloy) during the fluctuational or

martensitic transition of the graphite crystallites (sp: bonds) that are present in the initial carbon materials. If the

initial carbon material does not contain graphite crystallites of sufficient size (rGcri, > rz_r, )or if they cannot pass into

the metal melt then the formation of diamond nuclei is hindered. Since highly dispersed carbon material, e.g., nongraphitized carbon black, dissolves much more rapidly in metallic melt than does polycrystalline graphite (ref. 4),

then introducing carbon black prior to the metal or alloy, one can control the dissolution of polycrystalline graphite

and, hence, the passage of graphite crystallites into the melt. The above-mentioned mechanism was confirmed

experimentally. Indeed, the diamond was not synthesized at all at a certain amount of carbon black dissolved prior to

metallic melt (ref. 5).

It is well known, that the area of pressure-temperature carbon diagram for SD growth is near the equilibrium line

for the graphite-diamond transition (P >_4.5 GPa, T> 1600 K) figl (ref. 6) and the grox_lh of the large SD (> 2 mm)

by diffusion process takes more than a few hours.

On the contrary, the SPD is formed at P and T far from equilibrium line (ref. 7), and the process of SPD

formation with diameter is equal to 5 mm is completed no more than for 10 seconds (fig. 1).

FORMATION PROCESS OF SYNTHESIZED POLYCRYSTALLINE DIAMONDS

The mechanism of SPD formation had been discussed in the literature (ref. 9, 10)), but the attention was focused

on the discussion of the polymorph transition graphite-diamond and it was not allowed for transport of metallic melt

into the bulk of porous polycrystalline graphite. It was marked (ref. 11) that initial graphite prior to impregnated by

various organic liquids did not undergo the transformation to diamond, hnpregnation phenomena of graphite with

Cu and also Ni-Cr alloy under conditions, when the graphite-diamond transformation did not occur, (P > 3 GPa, T =

1500 K) was found by SEM/EDS method in (ref. 12). Impregnation process of the graphite with metallic melt in this

case is determined by both the value of high pressure (liP) in reaction cell (r.c.) and the value of pressure (P) into

NASA/CP--2001-210948 239 August 6, 2001
thepores(por.)ofgraphiteandcanbedescribedbyNovier-Stokes equation. Undertheactual conditionsofSPD

synthesis(P=7-8GPa,T _>2000 K) the value of AP (P in HP r.c. - P into the por.) close to the total high pressure

in the HP cell (the vapor phase into pores of graphite is highly compressible), while the value of capillary pressure is

no more than 0,2 GPa, when the graphite pores have a radius of file order of 10 lun (the valor of capillary pressure

was calculated on the bases of Novier-Stokes equation) (ref. 12).

10,0

,_, 8.0

_, 6.0

4.0

2,0

A

Diamond :::::::::::::::::::::::::::::::

Graphite

I I | I

1400 1600 1800 2000 2200

T(K)

Figure 1."A" - Region of the commercial high pressure synthesis of diamond from graphite.

SD - single crystal diamond (ref. 6); SPD - synthesized polycrystalline diamond (ref. 7, 8).

It is clear, therefore, that the value of zip is the dominant factor for the impregnation process, which occurs with

high rate. Thus, one can make the conclusion, that the impregnation process is accompanied by very fast solid phase

bulk transformation of graphite to diamond (sp 2 _ sp bonds), in the absence of diffusion in solid (martensite

transition) (ref. 13) and in the absence of the graphite visible dissolution into the metallic melt. This results in a

diamond polycrystalline composite, SPD, with inclusions of metallic phase (Ni+15% Mo) on the grain boundaries.

The mechanism proposed for SPD formation (impregnation process of graphite with metallic welt + martensite

lransition of graphite bulk to diamond Gs => Ds) allows also to explain the inheritance by SPD (ref. 14) of initial

graphites and other carbon materials features structures (texture, micro- and macrodefects and etc.).

Mechanism of macrostmcture formation in synthesized polycrystalline diamonds

In the work fief. 8) it was studied strucawal changes that occur in carbon material (carbon black obtained by

pyrolysis of methane) during heat treatment in the presence of Ni-Mn alloy in high-pressure chambers (HPC) of the

"lens" and "toroid" type (ref. 6).

The total lfigh pressure and the existence of shear stresses because of the non-hydrostatic nature of the pressure

in the HPC accelerate graphitization (ref. 15). Metallographic and X-ray studies of samples obtained at 3,7 -

4,3 GPa and 1750 - 2070 K revealed that they had a texture: the c axis of the crystallites of the carbon material

formed was parallel to vertical axis of the HPC. At the pressure rises the nature of the structure of the recrystallized

sample changes. At pressures and temperatures higher than 4,3 GPa, 1750 K and 5,5 GPa, 1750 K, respectively, the

grain boundaries are arranged in such a way as to form spherical (sphere-like) structures. Samples of spectral purity

graphite were also treated thermobarically at 5,5 and 7,7 GPa, 1990 K, in HPC.

At pressure 5,5 GPa the graphite undergoes texturing, similar to that observed in carbon black. At pressure

7,0 GPa the degree of texturing decreases essentially, indicating some change in the pressure field in HPC.

The pressure field in HPC during the formation of diamond polycrystal also depends on the impregnation of

graphite bulk with metallic melt at the total high pressure fief. 8). Under the conditions of the polymorphic graphitediamond transformation in HPC the initial graphite undergoes intensive recrystallization which ends in the

formation of a spherical texture of the resulting diamond polycrystal of ballas type. The fact that this texture can

form at 4,3-7,0 GPa is probably due to the existence of pressure gradients and volume effect of the graphitediamond transformation, as and in the formation of a sphere-like graphite texture during the recrystallization of

carbon black. At pressure above 7,0 GPa the sphere-like structure of diamond polycrystals is not formed because of

the role of the pressure gradientes decreases in comparison with the role of the total pressure (P > 7 GPa). The

metallic melt, which is adhesion-active with respect to graphite, can fill graplfite pores with diameter of 1 nm or less,

NASA/CP--2001-210948 240 August 6, 2001
causingfracturingof thegraphite(Rehbinder effect)and,therefore, fragmentation onthefinalstructureof the

sinthesized polycrystalline diamond.

THE INCREASE OF SPD PHYSICO-CHEMICAL AND MECHANICAL PROPERTIES

Synthesized polycrystalline diamonds (SPD) obtained at high pressures (P) and temperatures (T) have high level

of physico-mechanical properties and they are cheaper than the other policrystalline diamonds, which are produced

by methods of sintering. This superhard material is utilized an a large scale in Russia, China and other countries in

the production of dies, cylindrical heads, cutting, sawing, grinding and polishing tools, etc.

One must note, that commercial SPD has two shortcomings, namely, low thermal stability (beginning of

oxidation and graphitization at temperatures over 850 K) and visible decrease of mechanical strength after heating.

Heating the SPD to temperature above 1020 K results in structural changes which cause a worsening of its

mechanical properties. These structural changes are connected with the processes of surface graphitization and

oxidation of SPD. That does not allow to use intact SPD and also the diamond powder of grain fraction (250-

850 lun), obtained as o result of SPD trituration, on heavy work tools (as those used for processing marble and

granite, borehole drilling, etc.) or io replace scarce natural diamonds on tools.

The mechanism of metallic melt with dispersed refractory particle transfer to the bulk of SPD was realized for

the increase in mechanical strength of SPD through the addition of particles A1203 and TiN in the Ni-15% Mo alloy

(ref. 16, 17). Experimental data of SPD strength dependence on the content of dispersed dopes in the initial alloy are

explained, apparently, by the influence of these particles on the strength of the SPD non-diamond phase.

The latest experhnental results (ref. 19, 20) have shown the possibility for the essential increase in thermal

stability (resistance to oxidation of the SPD) and preservation in mechanical strength of the SPD up to higher

temperature through the addition of refractory particles (TiB2, BN _u_)in an initial Ni + 15% Mo alloy (fig. 2, 3).

12oo 4 _ '-'0 _o o

98O

o o

'J o

o v

1,4- BNcub

2,3- TiB 2

2 4 6 8 10 20

TiB2, BNcu b concentration (%,mass.)

25 30

Figure 2. Dependence of maximum temperature (T_) of diamond polycrystals drop

of strength on TiB2 (][IN cub) content in an initial alloy (1,2 - air, 10 min; 3,4 - vacuum, 20 min).

Figure 3. The influence of TiBz content in initial alloy on the minimal temperature

of oxidation (To) and polyerystal loss of mass (Am/m).

Before to 10% leads to the increasing of resistance to oxidation which is characterized by the increasing of To

(To - minimal temperature of oxidation) from 910 up to 1040 K and by decreasing of Am/m from 22% to 11%. The

further rise of TiB2 content does not change the characteristics of SPD oxidation (fig. 3). The same character of

NASA/CP--2001-210948 241 August 6, 2001
oxidation takes place when the initial alloy with TiB_ also contains cubic boron nitride (BN_b). The results obtained

were explained by the formation of boron oxide BzO3 as protective layer from oxidation and graphitization of the

SPD.

The other reason in the increase on Tm is the decrease of residual stresses level in SPD after healing and therefore

the diminution of probability of heat treatment cracks formation in diamond grains as BN _b has thermal expansion

coefficient close Io the one of diamond Under condition of diamond phase prevention of oxidation and grafitization

(TiB2) and of the decrease of residual stresses (BN _b) it is possible the further increase Tm of the SPD. In fact, the

use of initial alloys containing simultaneously TiBz and BN _b allows to synthesize the SPD keeping their strength

properties after heating up to 1290 K (rigA).

F, (N)

_ 1 Ni15%Mo 100- ]E 2-Nil 5%Mo+BNcu b

3-Ni15%Mo+BNcub(<1tt.m)÷Ti132

90 _ m)÷TiB2

80-

60- I I

900 1000 1100 1200 1300

Temperature of treatment, (K)

Figure 4. The influence of heat treatment (air 10 min.) and initial alloy composition

on SPD strength (40/45 mesh, 400/315 ltm grit size).

CONCLUSIONS

The above short review accumulates only part of the basic approaches to the selective improvement of SPD

physical chemical and mechanical properties of synthesized polyerystalline diamonds. The results of our study of

the processes diamond polycrystal formation makes it possible to determine practical ways of controlling these

processes. In particular, it seems promising to use initial materials with prescribed properties, e.g., metals or alloys

with different viscosities and graphites with a prescribed porous structure and degree of graphitization. This

probably results in a further improvement of SPD properties and extends the sphere of applicability of these

superhard materials in different areas of science and technology.

Acknowledgments

J would like to accord a thanks to my colleagues Dr. Victor V. Lapin, Dr. Andrei A. Ermolaev, Dr. Alexander I.

Laptev, Dr. Nicholas T. Loladze, Dr. Sergei A. Terent'ev, and others for fruitful mutual collaboration at the

Department of High-Temperature Materials of Moscow State Institute for Steel and Alloys.

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vol. 7, no. 1, 1975, pp. 41-47.

11. Valov, I.I.; Kalashnikov, Y.A.; Gumenin, A.V., et al.: Vliianie Nekotoryx Veschestv na Process

Perexoda Grafit- Almaz v Prisutstvii Metalla. Almazy i Sverhtverdye Materialy,

no. 6, 1975, pp.l-2.

12. Eliutin, V.P.; Poliakov, V.P.; Shalimov, M.D.: Issledovanie Processa Obrazovaniia Polikristallicheskix

Almazov. Doklady Academii Nauk SSSR,

vol. 275, no. 1, 1984, pp. 135-139.

13. Palatnik, I.S.; Gladkih, L.T.; Arinkin, A.V.: O Mexanizme Obrazovaniia Polikristallov Sinteticheskix

Almazov. Sverhtverdye Materialy,

no. 1, 1980, pp.7-13.

14. Vereschagin, L.F.; Apollonov, B.N.; Borovikov

issledovame polikristallicheskix almazov. Kristallografia,

vol. 19, 1974, pp.655-659.

N.F., et al.: Elcklronno-mikroskopicheskoe

15. Noda, T.: The influence of shear stresses in the reaction cell on graphitization of diamond. J. Carbon,

no. 6, 1968, pp. 125 128.

16. Poliakov, V.P.; Ermolaev, A.A.; Laptev, A.I.: Rossiyskaia Nauchno-Tehnicheskaia Conferencia,

MGATU, Moskva, Tesizy Dokladov, Novye Materialy i Technologii, 1995,pp. 168-171.

17. Poliakov, V.P.; Skury, A.L.; Ermolaev, A.A, et al.: The influence of refractory particles (A1203, TiN)

on the mechanical strength of carbonado type synthesized polycrystalline diamonds. Diamond and Related

Matcrials,

no. 7, 1998, pp. 422-426.

18. Poliakov, V.P.; Potemkin, A.A.; Ermolaev, A.A.: The Influence of Catalyst Composition on the

Oxidation Process of Carbonado Type Synthesized Polycrystalhne Diamonds. Advanced Performance

Materials,

NASA/CP--2001-210948 243 August 6, 2001
no.7,1998,pp. 422-426.

18.Poliakov, V.P.;Potemkin, A.A.;Ermolaev, A.A.:TheInfluenceof CatalystCompositionon the

Oxidation Process of Carbonado Type Synthesized Polycrystalline Diamonds. Advanced Performance

Materials,

no. 4, 1997, pp. 297-304.

19. Poliakov, V.P.; Ermolaev, A.A.; Laptev, A.I. el al.: The Influence of Boron Compounds Particles (TiB2,

BN _b) on file Thermal Stability and the Mechanical Strength of Synthesized "Carbonado After Heating.

Manuscript has been accepted for publication in Diamond and Related Materials, (07.03.2001).

NASA/CP--2001-210948 244 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

THE MULTI-ANVIL SPHERICAL TYPE DEVICE OF HIGH

PRESSURE WITH CUBIC COMPRESSION CHAMBER

Guerold Bobrovnitchii

Laboratory of Advanced Materials of State University of the North Fluminense,

Av. Alberto Lamego, 2000, CCT, Horto, Campos dos Goytacazes, RJ, Brazil,

CEP: 2815-620, Tel/Fax +55-24-7263733, e-mail: gaero!d(iz)a_nflbI

Eduard Perslkov

Laboratory of High Pressure of Institute of Experimental Mineralogy of AS of Russia.

Abstract

The problem of obtainment of diamond and cubic boron nitride in conditions of high pressure and temperature

has been successfully resolved after creation of multi-anvil device of high pressure. Compression chamber of this

device allows ensuring hydrostatic condition of pressure. We elaborated multi-anvil spherical unit of new design this

device makes it possible to produce: diamond powder up to 600 gk; monocrystals up to 1 carat; policrystalls up to 10

mm in diameter; composition in the form of superhard diamond plates up to 20 mm in diameter. Besides, there were

elaborated conditions to produce powder compositions for cutters, based on cubic boron nitride, for example

_Hexanite-R". The main difference of the offered device is: all the inner anvils are confined in spatial frame that

reciprocally orient anvils and at the same time, provides their fast change together with container. The device

includes the force block, a table, a transformer, a cooling system, a hydraulic drive, and a control panel. Power block

consists of two semi bodies with semispherical spaces, two semispherical elastic diaphragms, six or eight outside

anvils, six inside anvils, control device. Two semi rings keep semi bodies. The technical data of device: treatment

pressure - 4,5 ... 7,5 GPa; treatment temperature up to 1800°C; volume of cubic chamber - 30 ... 50 cm3; dimensions

- 2060 x 1970 x1820 mm; weight - 4000 kg.

Key words: high pressure, superhard materials, multianvil device.

INTRODUCTION

The first synthesis of artificial diamonds was executed in 1953 in the famous multi-anvil the device of von

Platen (ref.1, 2). However industrial synthesis is carried out on devises of type _Belt" (ref.3). Nevertheless, multianvil device continued the development (ref.4 to 25). In the beginning, they were applied only for scientific

purposes. These devices cautiously take root into last years in an industry (ref. 24, 25). Among devices of this design

the significant interest is represented by devices with hydrostatic drive of anvils or split sphere (ref.2, 6 to 11, 14 to

16, 21,23 to 25). All known designs of spherical devices have the features. However in all variants the chamber of

compression has in cubic or octahedral form. In some is volume of these chambers their almost hydrostatic

conditions of compression. The devices of this type have additional advantages: (i) opportunity of change of volume

of the chamber of compression; (ii) high synchronization of a movement of anvils to the center of the chamber of

compression; (iii) good cooling of anvils; (iiii) opportunity of simultaneous measurement of several parameters and

realization of long processes. All this gives appreciable advantages at growth crystals of diamond by method of a

temperature gradient (ref. 24,25). The difficulties of application of spherical device in industrial conditions are

connected to duration of installation of container and anvils and application of workmanship.

In the present work it was developed and the design of the device of type split sphere with cubic chamber of

compression was investigated. Such design enables to mechanize installation of anvils, container and deformable

gaskets (ref. 25).

NASA/CP--2001-210948 245 August 6, 2001
EXPERIMENTAL PART

Development of the construction of the device.

The idea of the new construction consists in branching internal anvils from outside and association of them in a

single unit together with the container. It was offered to insert all replaceable components into the uniform block,

using a dismountable crate from high-strength plastic. This block is submitted on a figure 1.

9 13 12

,,

Figure 1. An arrangement of components in the replaceable bloc.

The block consists of a skeleton of a crate 1, its top part 2 with elements of fastening 3, cubic container 4 with a

reactionary cell, six internal anvils 5 and gaskets 6,7 between the anvils. The reactionary cell has ceramic tablets 8,

conductors 9 and 10 and reactionary mix 11.

The installation of the block is carried out as follows. The gaskets 6 and 7 are pasted on all inclined planes

bottom and top anvils 5, and also on two planes two horizontal anvils 5. Bottom anvil 5 is established in the bottom

square aperture of a skeleton of a crate 1. Then four are mounted horizontal anvils 5. In space between anvils 5 the

container 4 with components 8 - 11 is established. Last mount top anvil 5 and top part 2 crates 1. After a fastening

the block is ready to installation in the device. For measurement of temperature the crate has the appropriate

apertures for installation of thermocouples. The container 4 is executed from a ceramic material on the basis of

oxides of metals. The anvils 5 were made from a hard alloy (WC 94 % + Co6 %).

During development the new technical decisions ensuring the best conditions for work of the device were found.

Among them it is possible to allocate the following:

Material and design of a membrane transmitting hydrostatic pressure on outsides anvils;

Form of shields, closing fissure between outside anvils;

Method of fastening of shields;

Materials both designs of the deformable container and gaskets;

Way entrance of cooling and electric power of heating;

Design of the device of disassembly of force parts of the device.

The device consists of the force block 1, table 2, transformer of heating 3, electronic system 4, hydraulic drive 5,

case of management 6, system of cooling 7 (figure 2 and 3).

NASA/CP--2001-210948 246 August 6, 2001
' o °: , ,_.......... o

• , : ', ,,-: :-.,, .: :

I' :, (t , _-' ::, ,

• i a , °. %.. .m , _., . , _ : .. , .... 4

': /i _" _ .....' ",:-:-:........

, I:/.,', .,,..._.:

.... .,. _,.. .:: :_ }: ..... _._. : :¢_':: : :

/. o oo.oo° ° ° ° ° °.._i _........ g ..° ° ° o • .... i p ,, ,

_............. ....... +.: •_............ ..............-_.-_-.-.4 _, l) ................ ' _ _ _ _Z_'....... _ ' :' :' 1820

1325 _................. 4 : :........................¢ _ ' _---..--, .-, :.'-. ,. :: ::

o...-- ----o- -'_. ,_'. -,_o_-_-_• .... ._,, _ ._" % "o %, **,_,'' , ," .

• _. _-_-_----'---_,ti_.', ::" :':' :':' f.",:'

500 • : :.

,.....i/: :::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::.....i

2

3 1

I /

/ f

I

t

|

|

i]k

, ' 1300

', 2060

i

i

1970

Figure 2. A general view of the developed device.

NASA/CP--2001-210948 247 August 6, 200 l
Figure 3. The device in an initial situation.

The force block includes two semi-cases with semi-spherical cavities, two semi spherical elastic membranes, six

outside and six internal anvils, units of heating and cooling, accident protection devices. During realization of

process the half-case are connected in a single unit to the help of two semi-rings from high-strengtla steel. The force

block is established on a table, where tlae drives for moving half rings and opening - closing of the top half-case are

placed.

The transformer of heating can be established about a table or inside it. The secondary winding of the

transformer is connected to electxoinputs of the force block.

The hydraulic drive includes the pump of low pressure 6.3 MPa and adjustable pump on pressure 250 MPa by

tlae appropriate control and distributive equipment. As a working liquid the mineral oil is applied.

The panel of management contains control and measuring devices, the programmed controller and other

necessary electrical devices. For realization of synthesis in accordance with mathematical models and of scientific

researches it can be used.

The device works in hand-operated (from separate buttons) and automatic modes. The following operations are

consecutively carried out: loading internals anvils with the container in the force block, closing of the top semibodies, moving of semi-rings on the force block, fast rise of pressure up to 6,3 MPa, increase of pressure up to

working meaning, inclusion of heating and cooling, maintenance of parameters of synthesis during the given time,

inclusion of heating and cooling, decrease of pressure up to atmospheric, cultivation of half rings, rise of the top halfcase, unloading internals anvils with the deformed container and product of synthesis from the force block, clean-up

of the device.

NASA/CP--2001-210948 248 August 6, 2001
Specificationsftlaedevice.

Thedeviceisintendedformanufacture of superhardmaterialsusedformanufacturing ofthevariouscutting

tools.

Thespecifications ofthedeviceallowtoreceivediamonds, high-strength aMmoreresistancetemperature, by

thesizeupto600gm,monocrystals ofweightupto 1carat,polycrystals byadiameter upto 10mm,andalso

composites byadiameter upto20mm.Besides, inthedevicetheconditionsforreceptionofplatesfromcBNcanbe

reproducedNowadaysdeviceistestedforworkwithtwotechnologies- receptionof diamondpowdersand

monocrystals.

Characteristics ofdevice:

1. Maximal pressureinhydrosystem2. Pressureinthechamber ofcompression -

3. Temperature inchamber ofcompression -

4. Volumeofthechamber ofcompression -

5. Productivity uponsyntlaesis ofpowdersofdiamond6. Productivity ofmonocrystals -

7. Timeofprocessing bypressure

minimal -

maximal -

8. Preparatory time9. Dimensions10.Weight250MPa;

4,5...7,5GPa;

upto1800C;

30...50cm3;

250...300x 103car/year;

100... 150pcs/ year;

10 min;

72 hours.

1...10 min;

2060xl 970xl 820 mm;

4000 kg;

RESEARCH OF THE CRIATED DEVICE

At the first stage we studied influence of a material and initial sizes of the deformable container on size of

generated pressure in the chamber of compression.

The pressure in the chamber of compression was determined through a knowntechnique (ref. 26) with use of Bi,

T1, P1Se, Ba.

For a timing of an optimum ratio between the initial size of a side A of the container and size an of a side of

internals anvils a series of deformable containers with the relations A/a =1,1... 1,45 was made. As the material of the

container it was applied a compacted mix of powders ZrO2 91-96% and CaO 9-4%. The deformable gaskets were

made of tlae identical sizes from pirofillite.

Distribution of deformation in volume of the compressed container and the synchronization of movement of

anvils was studied with the help of a copper cube, cut half-aM-half. On a plane of section the grid with the size of a

cell 1+o,o2x 1+o,o2mm was rendered by a typographical metlaod. The angular deformation of a cell no more than 5 %

was accepted for a designation of volume witla almost hydrostatic conditions of compression.

The deviations in size of gaskets formed between anvils showed a degree of synchronization of movement of

anvils to the center of compressed volume. We investigated also influence of deviations of tlae sizes of the assembled

block internal anvils on symmetry of distribution of deformations in compressed volume.

The materials, number and configuration gaskets between anvils varied. Both simple, and compound gaskets

were applied. The estimation of their influence was made on efficiency and on resistance of anvils. For efficiency the

minimal meaning of the primary pressure necessary for achievement of given pressure in the chamber of

compression was accepted.

With the help of the thermocouples established at the center of the compressed container and on periphery of

anvil, it was studied the efficiency of cooling.

On the device it was realized synthesis of powders of diamond in system Mn-Ni by a method of spontaneous

crystallization and growth of monocrystals by a metlaod of a temperature gradient within 48 hours.

RESULTS AND DISCUSSION

From the diagrams submitted on figure 4, follows that there is some certain relation A/a, at which the pressure Pl

is minimal for achievement of the given pressure in the chamber of compression. For example, for transition Bi I-II

(2,54 GPa) the relation A/a is equal 1,22, for Ba (5,5 GPa) it reaches magnitude 1,35. Thus, the dependence of the

optimum relation A/a for interval of pressure Pz of 2,5 - 7,5 GPa can be submitted by the following expression:

A/a =1,18+1,42.10 3"pl ;

NASA/CP--2001-210948 249 August 6, 2001
13..

v

13..

160.00--

120.00

80.00

40.00

Ba

0.00 ' I ' I ' I

1.00 1.20 1.40 1.60

A/a

Figure 4. Dependence of pressure Pl in hydrosystem of the device in function of the ratio A/a.

The received dependence of pressure Pz in the chamber of compression from pressure p: in hydrosystem

represent almost direct lines each equipped block and show efficiency of use of the device. The results of researches

will be well coordinated with the data of work (ref. 27).

The measurement of sizes of deformations of cells on the average section of the container has shown, that the

hydrostatic conditions form at 25-28 of % of volume of the chamber of compression. By indirect parameters, this

zone becomes wider at increase of temperature besides; the degree of deformation is sharply reduced to the center.

The maximal deviations in thickness of deformed gaskets at pressure p: equal tol00 MPa did not exceed 0,07

mm, deviations in deformed cube - 0,13 mm. Was found out eccentricity of the application of loading from outside

horizontal anvils to internal ones. She reached 0,23 mm. It specifies presence of influence of weight horizontal

outside anvils. The increase of thickness of rubber washers between bottom and horizontal outside anvils will allow

reducing this influence.

Research of efficiency of cooling of anvils have shown, tlmt at maximal temperature

(1400°C) at the center of compression volume the conditions of thermal balance comes in 2,5 min after inclusion

of heating. Thus temperature of internal anvils on an outside surface did not exceed 90 - 95 o C. Increase of the

charge of a cooling liquid with 1,0 1/s up to 1,5 1/s has reduced temperature only on 5 - 8o C. Gradient of

temperature inside the chamber of compression reached 18-20 gr/mm. It is a little bit above titan meanings received

in work (ref. 27).

At syntllesis in system Mn-Ni-C the crystals of diamond were received yellow - green color without appreciable

impurity. Majority of crystals was in the form of octahedral. These diamonds settled down in a zone of hydrostatic

compression. Their sizes ranged from 400 to 650 mk. Outside of this zone the crystals in the form of cube and

cubeoctahedral with the sizes 40-250 mk in size were synthesized. The degree of transformation of graphite in

diamond reached 28-31 of %, that the meanings received in the work (ref.27) are a little bit higher. The received

monocrystals had yellow-brown coloring and form as octahedral, truncated on the part of the basis on one third. The

sizes of monocrystals ranged from 4 up 6,5 mm. Unfortunately, one third of monocrystals had metal inclusions.

NASA/CP--2001-210948 250 August6,2001
CONCLUSIONS

Theconstruction ofthesplit-spheretypehighpressureandtemperature hasbeendevelopedandresearched,

whichcanbeappliedforindustrial synthesis ofsuperhard materials.

REFERENCES

1. LianderH.Artificialdiamonds. ASEA.Joumal.V.28, No5-6,1955,p.97-98.

2. B. yonPlaten.Chapter6.Amultiplepiston,highpressuretchniques. In:Modemveryhigh

pressuretechniques. EdbyR.H.Wentorf. Butterworflas, London, 1962,p.366.

3. HallT.H.Somehighpressure, hightemperature apparatusdesignconsideration: Equipment for

useat100.000atmospheres and3000oC.TheRev.ofSci.Instrum.V.29 (4),1958,p267-285.

4. BallhausenC.Vorrichtungzumerzeugenhoherdrucke. Ind.Diam.Abst.Bd.23,269, 1958,s.103-

195.

5. HallH.T.HighPressure-Temperature Apparatus. Metallurgyat HighPressureandHigh

Temperature. V.22,No4,N.Y.1964,p.45-49.

6. ZeiteinA.Highpressures equipment. Sci.Amer.V.212,1965,pp.38-46

7. KawaiN. A statichighpressureapparatuswithtaperingmulti-pistonsforminga sphere.

Proc.Jap.Acad.V. 42N°4,1966,pp.385-388.

8. KawaiN.Equipment forgeneratingpressureupto800kbar.In:Accuratecharacterization ofthe

highpressureenvironment. Ed.C.C.Lloyd,NBS.Proc.of a Symposiumheldat theNBS.

Caaithersburg, MD(October 14-18,1968). Washington. D.S.20234.

9. Highpressuremethodsinsolidstateresearch. Ed.byC.C.Bradley. London. Butterworflas. 1969,p.

291.

10.KawaiN.Highpressuregnerating methodsandapparatus. Patent N°1200934, Japan, B01j 3/00,

05.08.1970.

11.KawaiN.,EndoS.Thegenerationfultrahighhydrostaticpressurebyasplitsphereapparatus.

TheRev.ofSci.Instrum. V.41,N°8,1970,p.1178-1181.

12.Wakatsuki M.,IchinoseK.,AokiT.Characteristics oflink-typecubicanvil,highpressure- high

temperature apparatus. Ibid.V.11,No4,1972,pp.578-580.

13.KumasawaM.Theorygenerationfveryhighstaticpressurebyanexternal force.J.HT-HP,No

5,1973,p.212-220.

14.Bobrovnitchii G.S. and.MaximovIu.L.Multianvildevicesuchasspherefor creationof

superhighpressure. DeviceandEngineering ofExperiments. No2,1974,p.220-222(inRussian).

15.MalinovskyI. Ya.,RahnE.N.Dependence of "splitsphere"outerstageffortananvilnumber.

In:Experimental investigation onmineralogy. Ed.byIGM.Novosibirsk, 1975,p.135-143(inRussian)

16.Wanagel J.andRuoffA.L.Asplitsphere60.000TonPress.ComingGlassWorks.Coming, N.

Y.,1976,p.840-845.

17.BundyF.P.Designingtaperedanvilapparatusforachievinghigherpressures. TheRev.ofSci.

Instmm.V.6, 1977,p.317-324.

NASA/CP--2001-210948 251 August6,2001
18.SpainJ.L.Ultrahighpressureapparatusandtechnology. In:Equipment, design, materialsand

properties. Highpressuretechnology. V.1.Ed.byJ.L.Spain,J.Paauwe. NY:Dekker.1977,p.395-

423.

19.MalinovskyI. Ya.etal.Theanalysisofthebasicparameters andchoiceofanoptimumdesignof

multi-anvilsblockofatypesplitsphere. WorksofanInternational Seminar _Superhardmaterials".

Kiev.V.1,1981,p.45-46

20.Wakatsuki M.,IchinoseK. Multianvildevices. In:HighPressureReseatedin Geophysics.

Eds.S.Akmoto, M.H.Manghani (Tokyo:CenterofAcademicPublications). 1982,p.13-26.

21.AkifumiOnodera. Octahedral-anvil high-pressure devices. J.HT-HP.V.19,1987,p.579-609.

22.BundyF.P.Ultrahighpressures apparatus. PhysicsReports. V167(3),1988,p.133-176.

23.MalinovskyI.Yu.,ShurinYa.I.,RahnE.N.Analysisof effectiveness andchoiceof optimal

constructionf two-stagemultianvilunitof splitsphereapparatus. Proceedingof XI AIRAPT

International Conference. HighPressureScienceandTechnology. V.4.Kiev,1989,p.83-85.

24.MalinovskyI.Yu.,ShurinYa.Ietal.A newtypeofthe_Splitsphere"apparatus(BARS).In:

Phasetransformations athighpressureandtemperatures applications ofgeophysical andpetrologic

problem. Japan. Misasa. 1989,p.12-17.

25.PersikovE.S.,Bobrovnitchii G.S.KimD-Tch.,Deviceforgeneratingofhighpressure. Russian

Patent N°2082492, B01j3/06,1997.

26.Bobrovnitchii G.S.andKlebanovIu.D.Measurement of electrical propertiesof conductors at

highpressurebothtemperatures. DeviceandEngineering ofExperiments. N°4,1972,21-222.

27.FedorovI. I.,SoninV.M.etal.ToaquestiononsynthesisofdiamondatpresenceoftheMn-Nicatalyst.In: Experimental researchesin connectionwitha problemtopmanta.Ed.by IGM,

Novosibirsk, 1982,p.90-103.

NASA/CP--2001-210948 252 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

The Optimized Cell Design for High Pressure Synthesis of Diamond

James C. Sung

KINIK Company, Taipei, Taiwan & Taipei University of Technology, Taipei, Taiwan, R.O.C.

Abstract

More than 1O0 tons of diamond grits are consumed each year in sawing constructional

materials (e.g., concrete, granite, marble). These diamond grits are typically synthesized under high

pressure and high temperature using graphite as the source material and a molten metal (e.g., Fe, Co,

Ni or its alloy) as the catalyst. Graphite and catalyst may be either stacked in the reaction cell in

alternative layers or they can be mixed as powders. The reaction cell may be small (< 50 cubic

centimeters) with large gradients of pressure and temperature, as in the case of using a cubic press, or

they can be large (up to 500 cubic centimeters) with relatively uniform distribution of pressure and

temperature, as in the ease of using a large (1000 tons) belt press.

The current arts are such that if the reaction cell is small, and the charge contains layered

graphite and Af graptite andeatalys t, then t is about 1 carat per cubic centimeter of the compressed

cell, with about 1/3 of the grits that are desirable in size and quality. If the reaction cell is large, the

diamond yield is about 2 carats per cubic centimeter with about half as desirable grits. If powdered

graphite and catalyst are thoroughly mixed, and seeded crystals are dispersed uniformly in them, the

diamond yield in the later case could be boosted to about 3 carats per cubic centimeter with about 3/4

as high grade products.

By mixing powdered graphite and catalyst and pressed the mixture in thin layers, catalyst

coated diamond seeds may be planted uniformly into these disks. With the careful selection of

source materials, and the full utilization of reaction volume, the production yield may attain more than

4 carats per cubic centimeter with more than 3/4 as high grade saw diamonds. For a ten thousand

ton press with optimized design of high pressure assemblage, the reaction cell can attain a volume of

about 500 cubic centimeter. Hence, it is possible to yield 2 kilograms of high grade saw diamond in

a single run.

NASA/CP--2001-210948 253 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Liquid Phase Synthesis of Diamond in Hydrogen Atmosphere

Hsiao-Kuo Chung 1, Ming-Show Wong 3, Yong Liu 3, James C. Sung TM

2.

3.

1. KINIK Company, Taiwan, R.O.C.

Tung-Hwa University, Taiwan, R.O.C.

Chinese Academia Sinica, Taiwan, R.O.C.

4. Taipei University of Technology

Abstract

Graphite can be converted to diamond under high pressure by the catalytic influence of certain molten

metal. The transition is reversed at low pressure. However, if hydrogen atoms is available in the atmosphere,

graphite may be gasified to form methane, and the latter deposited on the molten metal as diamond.

NASA/CP--2001-210948 254 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

THE INVESTIGATION OF FORMATION AND AGGREGATION OF UDD

(ULTRADISPERSED) DIAMONDS ON SPECTRA OF THEIR IR (INFRARED)

ABSORPTION

A.Koretz, E.Mironov,V.Malyi

Krasnoyarsk Slate Technical University, Krasnoyarsk, Russia, prcom@kgtu.runnet.ru

E.Petrov

Federal scientific research center "Altay", Byisk, Russia

ABSTRACT

The ultradispersed diamond (UDD) produced by detonation method from mixture of trinitrotohiene and

hexagen (TH) consists of system of particles with average grain dimensions 3-grim and specific surface 250m2/g

and more. Most investigators consider that the considerable ratio of surface energy to volumetric energy for these

particles must make distortions of energetic bands and crystaUine structure and consequently can cause for many

phenomena typical only for nanometric objects. The high thermodynamic stability of UDD mentioned in most

experimental works, allows us to suppose that there is a mechanism of the UDD thermodynamic stability. We have

suggested that the high stability is connected with the decreasing of the surface energy portion in the total energy of

this system at the expense of physical aggregation of the UDD grains.

The experimental measurement of the UDD aggregates dimensions was performed on the basis of light

scattering. The identification of molecular groups "cover" was performed on the basis of infrared spectroscopy.

The object of this investigation was UDD, obtained by different groups of researchers: UDD(Kmsnoyarsk),

UDD(Chelyabinsk) and UDD(Novosibirsk) were obtained from TH mixture in mass ratio 60/40.

UDD(Krasnoyarsk) were refined from detonation products by means of boron anhydride. UDD(Chelyabinsk) and

UDD(Novosibirsk) were refined by means of acid treatment. The synthesis of UDD(Novosibirsk) was performed in

1989. Sample HS(Byisk) was obtained from hexagen with an addition of soot. There were investigated size and

properties dependence on the UDD-aggregates from physical effect (the thermal effect, the reactor's gammaradiation) in this work.

Table 1. The average integral dimension of the UDD-aggregates

sample dimension nm sample dimension ran

UDD(Krasnoyarsk) l01 UDD(Novosibirsk) 107

UDD(Chelyabinsk) 146 HS(Byisk) 375

According to the performed investigation the most informative range of IR-spectra of UDD-particles is 1000-

1300cm _. There is competitiveness of two bands: 1280cm "1- the band connected to the nitrogen defect of A-type in

the crystalline lattice of nanodiamonds characteristic for UDD(Krasnoyarsk), 100-1150cm "1- the band of valent CO vibration, characteristic for UDD(Chelyabinsk).

There was performed thermal procedure of samples UDD(Chelyabinsk), UDD(Krasnoyarsk) in the temperature

range 373_K, 673°K, 973°K. At heating to 673°K for these investigated samples the decreasing of the sample

dimension was observed.

Table 2. The temperature influence on the UDD-aggregates dimensions

sample The aggregate dimension at temperature

373C'K 673OK 973OK

UDD Krasnoyarsk 101.Snm 73.0nm 118.6nm

Some two irradiated samples were investigated - detonation soot, containing diamonds, and

UDD(Krasnoyarsk). IR-spectra of the grattfite-diamond soot from the explosive chamber did not change after -

irradiation. The second sample was irradiated by fast neutrons (E> 1MeV) with F=(I,30 :tO,19) 1017neutr./sm 2, and it

did not show significant changes in the structure of the IR-spectrum either. The UDD-sample was exposed to 7-

irradiation with value 0.5 + 0,17Mrad and after irradiation There were observed some absorption changes

connected with vibrations of alkanes hydrocarbon groups (1439, 2859, 2930 H 2970CM'1).

Therefore after the result it was decided to use only y-irradiation with increase of doze up to 183 +_24 Mrad

(with energy up to 2,4MeV). After the sample HS (Byisk) exposure it was clearly observed that the band had shifted

from 2945cm'tto 2920cm -1. The shift was explained by the radiation effect on vibrations of alkanes hydrocarbon

groups. For all these samples the average integral dimensions of aggregates decrease about 10-20%, in spite of the

considerable radiation energy.

NASA/CP--2001-210948 255 August 6, 2001
Keywords: Ultradispersed diamonds; Infraredabsorption; A-defect; Aggregation

KoretzAnatoly, professor

StateTechnical University

Kirensky street 26

Krasnoyarsk

660074 Russian Federation

Tel: (391-2) 49-72-66

Fax: (3912) 43-06-92

E-mail: prcom@kgtu.nmnet.ru

NASA/CP--2001-210948 256 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CHEMICAL STATES CARBON ATOMS ON UPPER LAYERS OF ULTRA

DISPERSION DIAMOND (UDD)

AND NATURAL DIAMOND (ND) AFTER H-TREATMENT

A.P. Dementjev, K.I. Maslakov, O.O. Zabusov

RRC "Kurchatov Institute", Kurchatov sq., Moscow 123182, Russia.

E-mail: A.Dem@g23.reicom.ru

ABSTRACT

A chemical state of carbon atoms on ultra dispersion diamond (UDD) and natural diamond (ND) surfaces after

in-situ H-tJceatment has been studied by X-ray excited Auger electron spectroscopy (XAES) combined with C 1s and

valence band XPS methods. The CKVV Auger spectra before and after H-treatment have revealed the unchanged

chemical state of carbon atoms on upper layers of UDD that was identified as UDD-like. The UDD-like chemical

state has been classified on the base of CKVV Auger spectra as sp2 state principally differing from that of HOPG

and fullerene. The transformation of the chemical state of carbon atoms on ND surface during H-treatment was as

follows:

sp2/sp3 ---> sp2(fullerene-like) --->sp2(UDD-like).

Key words: AES, XPS, Surfaces, Diamond, Ultra dispersion diamond.

INTRODUCTION

According to Johan C. Angus (ref. 1) _A better understanding of the nature of diamond nucleation is

required... Understanding of the nucleation process could help lead to the growth of a large diamond crystals from

the vapor. Of particular importance are diamond precursors..."

It is well known, that UDD can be used as a diamond growth precursor at CVD process.

Apparently, identification of the chemical state of carbon atoms on a UDD surface will give valuable

information on nucleation and growth of diamond from gas phase and will allow to find conditions of creating of

this state on any surfaces.

As was shown (ref. 2) the Auger spectra of carbon atoms with sp 2-, sp 3- and mixture of sp2/sp3-bonds have a

distinguishing feature due to the 6- and _z-levels positions in the valence bands. The information depth of Auger

emission is equal to about 2-3 monolayers.

The chemical state of UDD and ND specimen surface prepared ex-situ was identified as a new unknown

chemical state with sp2-bonds differing in principle from that in graphite (ref. 2). These data were established by

means of ex-situ experiments though in-situ experiments could provide more reliable results. Therefore the goal of

this work is to study in-situ the UDD and ND surfaces after H-treatment by XAES combined with Electron Energy

Loss at Cls and valence band XPS methods. Information depth of the methods are about 2, 7 and 10 monolayers

respectively.

EXPERIMENTAL

XPS and XAES data were obtained by a MK II VG Scientific spectrometer. The experimental conditions were

the same as previously describe in detail (ref. 2).

UDD powders (3-5 nm) were received from the Kiev Institute of Superhard Materials and from Krasnoyarsk.

The UDD powder were deposited on a Si plate (10xl 5 mm) from an alcohol suspension in an ultrasonic bath. Si2p

XPS peak was used to measure surface coverage by analogy with (ref. 3). The coverage appeared to be equal to

100%. The crystal plane (110) of a natural diamond of size 6*5 mm 2 and thickness 1.5 mm was mechanically

polished by means of a cast iron wheel. The roughness was equal to about 1 nm. An ultrasonic rinse in acetone and

then in alcohol was performed prior to the mounting of the sample in the vacuum chamber.

The surfaces of UDD and ND were analyzed before and after H-treatment by hot-filament (2000 K) in-situ in

the preparation chamber of the spectrometer. The pressure of H2 during H-treat_nent was equal to 10 .6 mbar. The W

hot-filament as spiral was located one centimeter above the sample. The substrate temperature was equal about

600 K. The analysis of the gas in the preparation chamber under H-treatment was carried out by Supavac

masspectrometer with mass range detection 1 - 100. W filament temperature was measured by a pyrometer.

NASA/CP--2001-210948 257 August 6, 2001
RESULTS

Theanalysisofthegasinthepreparationchamber duringUDDandNDH-treatment hasbeencarriedoutby

Supavacmasspectrometer. TheesultsofUDDH-treatment aregivenintable1.Twolowerowsrepresent the

pressureofOandN2withhotfilamentturnedoft'.Onlythemassesrepresented intwoupperandtwolowerowsof

thetablehavebeendetectedduringH-treatment ofND.It isnotsurprisingbecausetheareaofUDDsurfaceis

incomparably largerthanthatof ND.Thehydrocarbons iTabletestifiesthatduringH-treatment takeplace

chemical etchingof diamondsurface. Theformationof thehydrocarbons in gasphaseis notpossibleinour

experimental conditionsbecausequitelowpressure. Thefactofthechemical etchingofNDsurfacebyH-atomsin

thecommercial MW-PECVDsystemASTeX-PDS19 wasshown(ref.2).

Table 1. The results of the residual gas analysis in vacuum chamber during the H- treatment of UDD.

Mass Possible species Partial pressure,

mbar

1 H 5.5.10 .7

2 H2 2.5.10 .6

44 CO2 4.0.10 -l°

42 C3H6, C2H20 2.0.10 -1°

39 C3H3 2.0.10 -1°

16 O, CH4 2.0.10 -1°

26 C2H2 1.0.10 -1°

28 CO, C2H4 1.2.10 .9

29 C2H5 2.0.10 -1°

N2 1.0.10 -11

O 1.0.10 -11

Figure 1 represents the CKVV UDD spectra after the H-treatment in comparison with HOPG and fullerene. The

UDD spectra before and after H-treatment practically coincide with each other and have the shape inherent to the

sp2-bonds but principally differing from that of HOPG and fullerene.

As an Auger spectrum shape depends on chemical bonds (ref. 2) we can suppose that these spectra will allow to

identify an unknown chemical state on UDD surfaces. This chemical state could be designated as UDD-like.

1.0

C.K..VV.;: M .......... ,,.,," """ "

.° - ." ,.." _. .

>:, ...... Fullerene - \\

0.4 ...... Graphite ".\_

c ....\._

{ 0"2 "" "_

0"0

I i I i I

240 260 280

Kinetic energy, eV

Figure 1. The CKVV Auger spectra of UDD, fullerine and graphite.

NASA/CP--2001-210948 258 August 6, 2001
Figure2showsthemodificationfCKVVAugerspectraofNDaftertherowofsuccessive H-treatments and

interruptions forcarryingoutXAESandXPSanalysis:

1-(asreceived), 2-(7minofH-treat.), 3-(6minofH-treat.), 4-(6minofH-tzeat.), 5-(15minofH-treat.), 6-(15

minonlywithhot-filament without H2)

The spectrum 1 can be considered as the mixture of the sp 2 (30%)- and sp 3 (70%)-spectra.

The 2, 3 and 4 spectra coincide with that of fullerine. These comparisons lead us to suggest that the state has the

fullerene-like valence band structure after the 2nd and the 3rd H-treatment. The spectra take the UDD-like shape

after the 4th, 5th and 6th treatments. So, take place a transformation the chemical state carbon atoms on ND surface

sp2/sp 3 _ sp2(F-like)---_ sp2(UDD-like),

where F-like is Fullerene-like. The Auger spectra of the in-situ ND H-tzeat_nent coincide with that of the exsitu ND H-treated (ref. 2). This is the evidence of the surface state inertness of ND after H-treat_nent as well as UDD

with respect to the environment.

1.0

0.8

0.6

._ 0.4

_c 0.2

0.0

_ _ CKW

_ 56

, I , I , I , I , I

240 250 260 270 280

Kinetic energy, eV

Figure 2. The modification of CKVV Auger spectra of ND during H-treatment.

Electron energy loss spectra at Cls have been used for identification of sp3-bonds (ref. 4) in carbon. Fig. 3

shows modification of EEL at Cls XPS spectra after the row (as in Fig. 1) of successive H-treatments of the ND

surface. The intensities of all Cls XPS peaks were normalized. Changes in ELL spectra intensities after H-tzeatment

testify to H influence on the chemical state of carbon atoms in the depth up to 7 monolayers. These spectra represent

sp 3 bonds in contrast to Auger spectra.

Fig. 4 shows the valence band XPS spectra of ND. Some insignificant changes in the spectrum shapes can be

seen and this is the evidence of H-treat_nent influence on the deeper layers. Our spectra are in good agreement with

the data presented in (ref. 5 to 7).and correspond to sp 3 - bonds. An insignificant change in valence band spectra

after H-treatment was pointed out in(ref. 7) and our results confirm this fact.

NASA/CP--2001-210948 259 August 6, 2001
0 10 20 30 40 50

Relative Binding energy, eV

Figure 3. The modification of Electron Energy Loss at Cls XPS spectra of ND during H-treatment.

o9

.m

¢-

.m

O9

¢-

c

2.0

1.5

1.0

0.5

0.0

0 5 10 15 20 25 30

Binding energy, eV

Figure 4. The modification of valence band XPS spectra during H-treatment.

DISCUSSION

Thus, the results of the Auger, ELL at Cls and valence band XPS spectra having different information depth

can be summarized in the conclusion that spLbonds (UDD-like) forms after HqJceatment of ND only in 2-3 upper

monolayers.

NASA/CP--2001-210948 260 August 6, 2001
The following qualitative scheme of the chemical reactions can be proposed as a generalization of the Htreatment results:

C(initial) + H (H2) _ CnHm

CnHm _ C(F-like) + [H2 + CxHy]desorption

C(F-like) + H (H2) _ C(UDD-like) + [H2 + CxHy]desorption

C(UDD-like) + H (H2) _ CkH1

CkH1 _ C(UDD-like) + [H2 + CxHy]desorption

(1)

(2)

(3)

(4)

(5)

where C(initial) is the chemical state on UDD and ND surface before the treatment, CnHm and CkH1-

intermediate chemical compounds on ND and UDD surfaces during H-eatching, [H2 + CxHy] - hydrogen and

hydrocarbon desorbed from the surface.

CnHm forms on the surface as the result of the reaction of H and H2 with ND during 6 min (reaction 1) and

C(fullerene-like) forms after desorbtion of CxHy (reaction 2).

The next stage of H-treatment is the reaction of H and H2 with C(fullerene-like) and C(UDD-like) forms as the

result (reaction 3).

C(UDD-like) is the stationary state which don't change under H-treatment. Only the surface etching results

from this treatment. The data on surface roughness change from 1 up to 200 nm can be an evidence of such an

etching (ref. 2) as well as the detecting of hydrocarbons under UDD H-treatxnent.

C(UDD-like) state presents on the initial surface of UDD therefore only the reactions 4 and 5 take place under

H-treatment. This state remains the same after an annealing of UDD at air during 7 min at 900 K. This is the

evidence of unique chemical stability of the C(UDD-like) state.

The concentration of hydrocarbon in the most widespread CVD processes is only a few percents of hydrogen

content. Therefore we suppose that the 1-5 reactions take place during diamond growth and C(UDD-like) state is

stationary. It is well known that UDD enriches nucleation and growth of diamonds. This is why the C(UDD-like)

state could be a precursor and perhaps the growth does not start until its formation on a surface.

The fact that the sp2-state of carbon atoms is the precursor looks to be surprising.

We should understand as the result of which reactions does the substrate with C(UDD-like) state transform to

sp 3 during diamond growth.

One should identify which of two possible mechanisms do take place - Langmuir-Hinshelwood or Eley-Rideal:

C(UDD-like) state + (CHn)substrate _ diamond

or

C(UDD-like) state + (CHn)gas _ diamond

XAES is the most powerful method to solve the problem of identification of critical steps responsible for

diamond growth.

CONCLUSIONS

One of the main results of this work is the demonstration of possibility of the real C(UDD-like) state

identification by XAES and XPS methods.

The H-treatxnent of UDD does not result in changes of chemical states of carbon atoms on surfaces as well as in

near-surface regions (in the limits of information depth - 2+10 monolayers).

The chemical state on the UDD surface of about 2 monolayers always remains sp2 but it differs in principle

from that of HOPG and fullerene.

The H-treatment of ND changes dramatically chemical states of carbon atoms only in upper 2-3 monolayers.

The transformation of the chemical state carbon atoms on ND surface from mixture during H-treatment is the

following: sp2/sp 3 _ sp2(F-like) _ sp2(UDD-like).

REFERENCES

1. Johan C. Angus, Thin Solid Films 216 (1992) 126.

2. Dementjev A, New Diamond and Frontier Carbon Technology,. 11(2001) 35.

3. Eliyahu E., Buehler J., Ben-Chorin M., et. al. Diam. Relat. Mater. 8 (1999) 146.

4. McFeely F.R., Kowalczyk S.P., Ley L. et. al.: Phys. Rev., B9 (1974) 5268.

5. Kulik J., Lifshitz Y., Lempert G. D., et. al. J. Appl. Phys. 76 (1994) 5063.

6. G. Francz, P. Oelhafen, Diamond Relat. Mater. 4 (1995) 539.

7. P. Reinke and P. Oelhafen, Phys. Rev. B60 (1999) 15772.

NASA/CP--2001-210948 261 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

DL4MOND (ItETEROEPITAXIAL)NUCLEATION BY ION BEAM IMPACT:

A REALITY

S. T. Lee, 1, Belle, and 12. S. Lee

Center of Super-Diamond and Advanced Films (COSDAF) &

Depmtmeot ofPfD'sics atul Matenats Sa'iertce

Ci_' Umv_rsi_ of Hang Koug

Hon_Kong SAR, China

_ABSTRACT

The cotttmtled ioo _om modred was m_ccessfulty m:Wiuyed tbr lira ficst finu: for the systematic

_u¢leatlon of diamond ovm'a wide nmge of expenmeotal coadJtions Silicon sabamtzs (l_kl at _700°C)

were I_ombarded by _80-200 eV ions from a source fed by a mixture of hy_ocarbo_hydrogeo/argon gases

a_ a low (10 "4 "Ion) pre_s_tre. High-resolution TF_I analyses ¢o_-firraexl the format/on of diamoml

¢.,Tstallites in diff_ent oacleatioa sites (nn tke silicon s',d_s"aa',eLa_!j,acm_tt_ SiC cvyslalli_es, embedded m a

noa-diamoM carbon matrix). Heteroepita.,dal diamoncl crystallites with no miserientatioa to the sillcnn

_ubstmtes we,x fon_ed on sil.icm_ sleps. The diamonrl crS,stallites _xisted m the./brm ogbatll known carbon

conggnlratIons, i.e., cable mad he_gonai diamond. Most strikingly, among the dlamon_I nuctei we find the

firs_ ¢onglusive eviderce re1 tJte fa_oatioo _f a new diamond p_tN_arph, designated 9R, with a _eflie|_'

of 9 layers a:_d a rhombN:edral Nructor¢. The iestdls greatly iloprove our carton" uzdermm_tfing of the

cltanmnd mtcleattoo proce_, opelta new, controtlable way ofgrouing heterogpita.,daI diamond films, and

demonstrate the capaflili_y of leabeam_ m oaaosuuctmmg carbon materials.

luvile_l talk: &DCFFC'F 200t, A_gmst 6-t0, 200!, .auburn Universi 5, Hotel and Conference Center,

AobunL AlaL'ama USA

Correspondh_g mrttmr S. T Lee

Ceatea of Super- Diamond and Advanced Fgms

(COSDAb3 and

Department of Physics and Materials Science

Cby Unb.'ersity of Hang Kong

83 Tat Chee A_w.,

Kowloon, Hang Kong,

SAR China

Fax; (852) 2784 4696

"t_:i:(852) 2788 9606

_-mail: apazmal=__.;¢ityu =d-_tm

NASA/CP--2001-210948 262 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

HRTEM study of the early stages of diamond nucleation on Si(ll 1) dimples

J.C. ARNAULT, S. PECORARO, J. WERCKMANN and F. LE NORMAND

Groupe Surfaces-Interfaces, Institut de Physique et Chimie de Strasbourg, IPCMS-GSI,

UMR 7504, B_t 69, 23, rue du Loess, 67037 STRASBOURG France

N. MOTTA and R. POLINI

Physics department, Untversita Tor Vergata; Via della Ricerca Scientifica 00133 ROME, ITALY

Highly oriented diamond films are required for diamond applications in kiP / HT devices [1]. In this way, a

better knowledge of the elemental mechanisms occurring during the nucleation and growth is essential to control the

early stages of diamond growth. Indeed, both the geometry of the suitable nucleation sites and the chemical nature of

the diamond precursor still remain unclear. Nucleation studies by HRTEM, STM and SEM FEG performed in Si(l 11)

dimples exhibiting both multi-vicinal and thin fiat areas will be presented:

1) Well defined Si(111) multi-vicinal surfaces were induced by an in-situ UHV thermal treatment at 1500K on

dimpled samples which underwent any destructive pre-treatment. The terrace formation was checked by STM UHV

observations [2]. Then, HFCVD diamond nucleation and growth were studied on this well characterised distribution of

defects by HRTEM and SEM FEG. We observed a significant enhancement of the nucleation density on such stepped

surfaces (up to three orders of magnitude) compared to virgin silicon monocrystals. More, the diamond nuclei are

mainly located in the vicinity of these steps.

2) Thin Si(111) areas were obtained via a thinning down of the dimple centre. Diamond nucleation was followed by

HRTEM on thinned areas after successive HFCVD deposits. After 7 min of deposition, 13.SIC nanoerystals were

observed with a pseudoepitaxial SiC{220}//Si{220} relationship [3]. Then, after a second deposit, nanodiffraction

experiments on carbide nanocrystals reveal the presence of a new phase with parameters close from those of diamond.

These results suggest a nucleation mechanism occurring from 13-SIC nanocrystals in fair agreement with a recent study

[4].

References

[1] <<Diamond films : recent developments >>D.M. Gruen and I. Buckley-Golder, MRS bulletin (Sept 98).

[2] ] J.C.Arnault, S. Pecoraro, J. Werckmarm and F. Le Normand Diam. Rel. Mat. (in press, 2000)

[3] F. Le Normarid, J.C.Arnault, S. Pecoraro and J. Werckmarm, Appl. Surf. Sci. (in press, 2000)

[4] D. Wittorf, W. Jager, C. Dieker, A. Floter and H. Guttler, Diam.Rel.Mat. 9 (2000) 1696.

Keywords Diamond Nucleation Silicon Step Transmission Electron Microscopy

Corresponding author

Jean-Charles ARNAULT

IPCMS / GSI

23, rue du Loess

67037 STRASBOURG Cedex

France

Tel: (33) 3 88 10 71 02

Fax: (33) 3 88 10 72 48

e-maih Jean-Charles.Amault@i pcms.u-strasbg.fr

NASA/CP--2001-210948 263 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ULTRADISPERSE DIAMOND STRUCTURE, PHASE TRANSFORMATION AND

APPLICATION FOR CVD DIAMOND FILMS GROWTH

A. E. Alexenskii, M. V. Baidakova, N. A. Feoktistov, V. G. Golubev, S. Grudinkin,

A. T. Dideikin, V. G. Melehin, V. Yu. Osipov, V. I. Siklitski, A. Ya. Vul'

Ioffe Physico-Technical Institute, 26 Polytechnicheskaya, StPetersburg 194021, Russia

A. M. Panich, I. Shames

Department of Physics, Ben-Gurion University of the Negev, P.O.Box 653,

84 105 Be'er-Sheva, Israel

W. Kempinski

Institute of Molecular Physics, Polish Academy of Sciences, Smoluchowskiego 17, 60-179 Poznan, Odolandw,

Poland

A. Kriiger, E. Osawa, M. Ozawa

Toyohashi University of Technology, Toyohashi, Japan

ABSTRACT

The paper reports on study on the structure and phase transformation of a new form of carbon cluster

material - ultradisperse diamond (UDD) with mean size of the cluster of 50 angstrom. High concentration of

paramagnetic centers (102o spin/g) has been detected; it is supposed these centers are related to defects of the

sp3-hybridized carbon of the UDD core. A quantum size effect of the g-factor has been detected in UDD by EPR

method. Possibility of using the UDD clusters as precursors for CVD diamond film growth is discussed.

Keywords: carbon, cluster, diamond, phase transition.

INTRODUCTION

Ultradisperse diamond (UDD) is one of the few carbon cluster substances that can be produced in a large

amount by detonation method (refs.1 to 2). Studies of the properties of UDD started in 1988, with major attention

focused on the technology of preparation of this material and its mechanical properties. One of the unusual features

of UDD is the small size distribution, with most of the particles having dimensions of 40 - 50 A. The possibility of

transforming onion-like carbon to diamond by electron beam and low-heat treatments and an assumption that

diamond rather than graphite is a thermodynamically stable form of carbon in small cluster suggests that UDD can

be used as a model material to investigate the phase transformation in carbon clusters.

We have recently shown that a UDD cluster has a complex structure consisting of a diamond core of about 45 A

in size and a shell made up of sp 2 - coordinated carbon atoms. We used X-ray diffraction and small angle X-ray

scattering to answer the question how the shell structure and thickness vary with the explosion conditions and

chemical purification of UDD clusters from the soot produced by the detonation process (refs. 3 to 4).

A specific feature of clusters as nano-sized objects is the crucial effect of the surface on their structural and

electronic properties, since the number of surface atoms becomes commensurate with that of bulk atoms.

In this paper we submit results of comprehensive studies of the UDD by EPR, NMR, TEM and also, Raman,

SEM and AFM characterization of diamond films grown by HF CVD method on silicon substrate with UDD

precursors.

SAMPLES AND EXPERIMENTAL TECHNIQUES

The samples used were obtained from carbon-containing explosives (TNT and hexogen) by the detonation

method, with the detonation carbon cooled in different ways more specifically by gas (CO2) (the so-cooled "dry"

NASA/CP--2001-210948 264 August 6, 2001
techniqueandbywater(the"wet"technique) (ref.3).Thus,thestartingmaterialsdifferedinthecoolingrateofthe

detonationcarbonitheclusterstructure.

Thecleaning(purificationfUDDfromsp2hybridizedcarbonatoms)wasachievedbytreatment iwater

solutionofnitricacidatatemperature of200 250°C.Followingrepeatedrinsinginwater,theUDDwasdriedin

airat200°Ctobecameapowderwithgrains1 10_tminsize.

Thefabricationofdiamondfilmsisofgreatinterest becauseoftheirpotential usefordifferentapplications.

Themostcommonmethodsof surfacepretreatment to enhancediamondnucleationandgrowthincludethe

activationofthesubstratesurfacebymechanical abrasionandultrasoundagitationwhichusediamondpowders. In

thisstudy,diamondfilmsweredepositedbyHotFilamentChemical VaporDeposition(HFCVD)onsilicon,

sapphire, andquartzsubstrates seededwithUDD.

A UDDcolloidwassuspended inisopropyl acohol. Thesubstrates werecoatedwithathinlayeroftheUDD

colloid.ThetypicalHFCVDconditionsincludedthefilamenttemperature of2200°C,thesubstratetemperature of

650 850°C,thegaspressureof8-50Torr,theflowrateof100-500sccm,andthemethaneconcentration of0.5-

2.5%.

RESULTS AND DISCUSSIONS

It is shown that, irrespective of the type of synthesis ("dry" and "wet") the relative content of the diamond (sp3)

phase increases after annealing in hydrogen within the anneal temperature range of 450 to 750 °C, but the increase

being more pronounced in the samples prepared by "dry" synthesis. The analysis has shown that the effect is due to

the transformation of sp2 shell into sp3 diamond phase.

The EPR signal from detonation carbon is shown to consist of two lines: (1) a broad structureless line due to

the presence of metallic impurities which can be removed by chemical purification and (2) a narrow line attributable

to paramagnetic centers produced by defects on the core of a UDD cluster. It is shown that the EPR technique can be

effectively used to evaluate the degree of UDD purification. High concentration of paramagnetic centers (1019 spin/g)

has been detected; it is supposed that these centers are related to defects of the sp3-hybridized carbon of the UDD

core.

We have found a change in the g-factor of the electron on the paramagnetic center, Ag - (5.2 ± 0.3) 10 4,

associated with the size-effect in a diamond cluster of 40 A in diameter. The anomalous reduction in the spin-lattice

relaxation time of 13C in UDD relative to natural diamond can be accounted for by the interaction between the

electrons localized on paramagnetic centers and nuclear spins.

We have studied the dependence of the technological conditions on the diamond film properties. The growth

rate of CVD diamond film on silicon substrate covered by UDD precursors amounts to 5 gm/hour. The crystallite

quality of deposited CVD diamond films was monitored by Raman measurements. XRD was used to determine the

structure and grain size of the films. The surface morphology and film thickness is characterized by using a scanning

electron microscope. It is shown that UDD seeding is a good nondestructive method for both opaque and optically

transparent substrates.

The Ioffe Institute authors acknowledge support under the Russian State Program "Fulllerenes and Atomic

Clusters".

REFERENCES

1. Lyamkin, A. I., et al.: Diamond production from explosives. Sov. Phys. Dokl. vol. 33, 1988, 705 - 708.

2. Greiner, N. R., et al. Diamond in detonation soot. Nature, vol. 333, 1988, 440 - 442.

3. Alexenskii, A.E et.al.: The structure of diamond nanoclusters. Phys.Solid.State., vol..41, no. 4, 1999, 668 -

671.

4. Baidakova, M.V. et al.: Ultradisperse diamond nanoclusters. Fractal structure and diamond-graphite phase

transition. Chaos, vol. 10, no 12, 1999, 2153 - 2163.

NASA/CP--2001-210948 265 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

The properties of nucleation and growth for CVD diamond on

mirror-polished Si pre-seeded with nanosized diamond powder

Zengsun Jin, Changzhi. Gu, Yue Sun and Ye Sun,

State Key Laboratory for Superhard Materials, Jilin University,

130023 Changchun, China

Fax: 0086-431-8923907 Email: zsjin@mail.jlu.edu.cn

A pretreatment method by seeded nanosize diamond powder was used to enhance

diamond nucleation on mirror-polished Si substrates. The seed process were

performed using the liquid including 5 nm diamond powder by a brushing technique

for obtaining uniform seeded nuclei in large area and low interface state density

between the diamond and Si substrate. High temperature annealing process in

hydrogen plasma were performed on seeded substrates in microwave plasma

enhanced chemical vapor deposition (MPECVD) chamber to improving the quality of

film and increasing the adsorption between the diamond and substrate. Diamond films

were deposited on seeded substrates by MPECVD method. We studied the nucleation

density, grain size, growth rate and film morphology by means of scanning electron

microscopy (SEM). The effects of seeded conditions and annealing parameters on

diamond nuclei and films were detailed. The Raman spectroscopy (RS) was applied to

analysis the phase purity of diamond films. The results show that high quality

diamond films with large area and low interface state density can be deposited on

4-inch mirror-polished Si substrates.

Keywords: CVD diamond, diamond powder, seeded nuclei, annealing, growth

NASA/CP--2001-210948 266 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

EFFECT OF CYCLIC TIME INTERVAL OF CH4 FLOW RATE ON THE DIAMOND

NUCLEATION DENSITY ON GLASS SUBSTRATE

*S.-H. Kiln., E.-J. Bae

Department of New Materials Chemistry, Silla University, Pusan 617-736, Korea

T.-G. Kiln

Department of Materials Engineering, Miryang University, Miryang 627-702, Korea

Y.-H., Kiln

Center for Terahertz Photonics, Department of Electrical Engineering, POSTECH, Kyungbuk 790-784, Korea

T. Hosomi, T. Maki, and T. Kobayashi

Division of Advanced Electronics & Optical Science, Faculty of Engineering Science, Osaka University, Osaka 560-

8531, Japan

ABSTRACT

We deposited diamond films on the 1.0 x 1.0 cm 2 patterned glass substrate in a vertical-type microwave plasma

enhanced chemical vapor deposition (MPECVD) system. The pattern could be achieved by unidirectional scratching

on the surface of glass substrate using ca. 1 gm size diamond powders. The low temperature (ca. 450°C) diamond

deposition in this work could be achieved by keeping the substrate distance (ca. 3 cm) from the plasma under the

relative low microwave power (400 W) and total pressure (1.73 kPa) condition. Using the cyclic process, we can

certainly enhance the diamond nucleation density on glass substrate. Surface morphologies and diamond qualities of

the films have been investigated as a function of the cyclic time interval of the CH4 flow rate. The diamond

nucleation densities on the pretreated glass substrate would be much more enhanced under the condition of the less

CH4 flow off time interval in the cyclic modulation. The enhancement of the diamond quality for the grain by the

cyclic process was noticeable under the more CH4 flow off time interval in the cyclic modulation. Based on these

results, we discuss the trend and the cause for the enhancement of the diamond nucleation density on glass substrate

according to the cyclic time interval of CH 4 flow rate.

Keywords: cyclic process, diamond nucleation density, cyclic time interval, low temperature deposition

* o

Corresponding author E-mail: sl-dcim(o2silla.ac.kr Phone: +82 51 309 5619 Fax: +82 51 309 5176

INTRODUCTION

To practically apply synthetic diamond films to field emission displays (FEDs), the enhancement of the

efficiency of field emission characteristics, such as low operating voltage, high emission current, emission

uniformity, and the achievement of low temperature deposition of diamond film, namely, the diamond film

deposition on indium tin oxide (ITO) glass substrate is preferential (refs. 1 to 3). From the manufacturing point of

view, the patterning of the diamond films to be suitable for the diamond field emitter is also indispensable. However,

unfortunately, the patterning of the diamond films has been regarded as the most difficult barrier because of the

hardness and chemical inertness of diamond (ref. 4). Recently, we were able to achieve the patterning of the diamond

film via the enhancement of the diamond nucleation density on the pretreated glass substrate by the cyclic process

(ref. 5). Furthermore, the enhancement of the diamond nucleation densities has a favorable influence upon the

increase of the emission current and the emission uniformity.

Compared with the ex-situ process, the in-situ process has the advantage for the further enhancement of the

diamond nucleation density, because one can apply the in-situ technique on the ex-situ treated substrate,

consequently obtaining more enhanced nucleation density. The cyclic process is the in-situ process that has the

advantage for the further enhancement of the diamond nucleation densities, consequently increasing the efficiency of

diamond field emission characteristics. For the cyclic process, we incorporated the cyclic modulation of CH4 flow

rate during the initial deposition stage. The cyclic modulation was carried out through the on/off control of CH4 flow

rate. Namely, it was started from H2 + CH4 plasma (CH4 flow on), we call it the growing time, and ended in H2

plasma (CH4 flow of[), we call it the etching time (ref. 6).

NASA/CP--2001-210948 267 August 6, 2001
Thisworkpresentsthemethodtoobtainawell-patterned diamondfilmhavinghighnucleationdensitiesunder

therelativelowtemperature (about450°C)depositioncondition. In detail,wedepositedthediamondfilmby

changingthegrowing/etching tmeintervalinthecyclicprocess. Weinvestigatedthevariationof thediamond

nucleationdensitiesincludingthegrainsizesandthediamondqualitiesasafunctionofthegrowing/etching tme

interval.Finally,wecouldobtainthewell-patterned diamondfilmhavingthehighest nucleationdensitiesonthe

pretreatedglasssubstrateundertheconditionoftherelativelowratiooftheetchingtimeinterval inthiswork.Based

ontheseresults,wediscussthecausefortheenhancement oftheselectivedepositionof diamondfilmonglass

substrateunderthelowratioof etchingtimeintervalconditionandtheincreaseinthediamondqualitieswith

increasingtheetchingtimeratiointhecyclicprocess.

EXPERIMENTAL

Wedepositeddiamondfilmsonthe1.0x 1.0mm2patternedglasssubstrateinavertical-typeMPECVDsystem.

Thepatterncouldbeachievedbyunidirectional scratchingontheselectedareaofglasssubstrateusingca. 1 btm size

diamond powders. Actually, unidirectional scratching was achieved by manual for 3 times with patch incorporating

diamond powders and acetone. Before patterning, we exposed the selected area for the scratch via covering the

residual glass substrate surface using tape. In this way, we made both the pretreated and the untreated areas on glass

substrate, consequently obtaining the patterned glass substrate.

The low temperature (ca. 450 °C) diamond deposition in this work could be achieved by keeping the substrate

distance (ca. 3 cm) from the plasma under the relative low microwave power (400 W) and total pressure (1.73 kPa)

condition (ref. 7).

Before the deposition reaction, we cleaned the substrate with H2 plasma for a few minutes. CH4 and H2 were

used as source gases. H2 flow rate was fixed as 200 sccm (standard cm 3 per minute). CH4 flow rates were varied as 4

and 6 sccm according to the reaction processes.

For the cyclic process, we incorporated the cyclic modulation of CH4 flow rate during the initial deposition

stage. The cyclic modulation was carried out through the on/off control of CH 4 flow rate. It was started from H2 +

CH4 plasma (CH4 flow on) and ended in H2 plasma (CH4 flow off). Namely, it was proceeded as H2 + CH4 --->H2 --->

H2 + CH4---> H2--+ .... then depositing diamond film. The total on/off CH4 flow modulation time was fixed as 10 min.

After the cyclic modulation, we deposited the diamond film for 2 h and 50 min. So, the total reaction time in this

case was 3 hours. For the normal process, we deposited the diamond films for 3 h without the incorporation of the

cyclic modulation of CH4 flow rate during the initial deposition stage.

We investigated the detailed surface states, and the grain morphologies using scanning electron microscopy

(SEM). The qualities of diamond grains on the patterned glass substrate were investigated by a micro-Raman

spectrometer (Renishaw 2000) with ca. 1 btm spot size using an Ar laser source.

RESULTS AND DISCUSSION

To elucidate the effect of the growing/etching time interval in the cyclic process on the enhancement of the

diamond nucleation densities and the selective deposition of diamond film on the patterned glass substrate, we

deposited the diamond film as a function of the growing/etching time interval under the different CH4 flow rate

conditions. We first investigated the surface images on the pretreated glass substrate under 4 sccm CH 4 flow rate

(Fig. 1) and 6 sccm CH4 flow rate (Fig. 2) conditions. SEM images of the surface were shown for the normal process

(Figs. la and 2a) and for the cyclic process having the growing/etching time = 180/30 sec (Figs. lb and 2b) and

30/180 sec (Figs. lc and 2c), respectively. At 4 sccm CH4 flow rate, the number densities of nuclei on the pretreated

glass substrate were counted as ca. 4.0 x 107(normal process), ca. 1.2 x 108 (180/30 sec cyclic process) and ca. 1.0 x

108 (30/180 sec cyclic process) (nuclei/cm2). At 6 sccm CH 4 flow rate, the number densities of nuclei on the

pretreated glass substrate were counted as ca. 5.0 x 107(normal process), ca. 1.4 x 108 (180/30 sec cyclic process)

and ca. 1.1 x 108 (30/180 sec cyclic process) (nuclei/cm2). From these results, we can undoubtedly confirm the

enhancement of the diamond nucleation densities by the cyclic process, as previous report (ref. 5). In addition, by

comparing the results between 180/30 sec and 30/180 sec cyclic process, we can suggest that the diamond nucleation

densities would be much more enhanced under 180/30 sec, namely, the less etching time interval condition in the

cyclic process (compare Figs. lb with c and Figs. 2b with c, respectively).

NASA/CP--2001-210948 268 August 6, 2001
'_e • " ' ,*_ ", _ ,,_ _','c_g_5_g'd,;',_:':

_N_;_:_.:'. _V.._N,_:,f4x_._,tgNlf4 : _".'.Z.___'.(.'_.: e__'::_"+:::+:':+:::+

Figure 1. SEM images of as-deposited flints, under 4 sccm CH4 flow rate condition, on the pretreated

surface, in case of (a) the normal process, (b) 180/30 sec cyclic process,

and (c) 30/180 sec cyclic process, respectively.

,::_:,_:,,::x:_:: i:- -',_ :_. INN'S.,'_'2_,,__

::::::::::::::::::::::::::::::: iii .i!., i ::i

Figure 2. SEM images of as-deposited films, under 6 sccm CH4 flow rate condition, on the pretreated

surface, in case of (a) the normal process, (b) 180/30 sec cyclic process,

and (c) 30/180 sec cyclic process, respectively.

To closely measure the nucleation densities and the grain sizes, we investigated the surface states of the

pretreated surfaces by high-magnified images of SEM (see Figs. 3and 4). At 4 sccm CH4 flow rate, the grain sizes of

these nuclei were measured as ca. 0.45 (normal process), ca. 0.54 (180/30 sec cyclic process), and ca. 0.48 gm

(30/180 sec cyclic process) in diameter. At 6 sccm CH4 flow rate, the grain sizes of these nuclei were measured as

ca. 0.50 (normal process), ca. 0.62 (180/30 sc cyclic process), and ca. 0.55 gm (30/180 sec cyclic process) in

diameter. These results also indicate a little increase in the diamond grain size by the cyclic process (ref. 5),

especially under the condition of the less etching time interval in the cyclic process.

. ..

i i : il ii

............... . .,. ,,,,,, ° o, ° ,,. %, - °°° ,°°°° ° ° ° ° * • * •. • - • •, °,, ,,, ° , • • • • •,,,,,,. ..............

Figure 3. High-magnified SEM images of as-deposited films, under 4 sccm CH4 flow rate condition,

on the pretreated surface, in case of (a) the normal process, (b) 180/30 sec cyclic process,

and (c) 30/180 sec cyclic process, respectively.

NASA/CP--2001-210948 269 August 6, 2001
...................................................... . .... . • ° °°°ooooo°°°°..°.°....°o°°°ooo°°

Figure 4. High-magnified SEM images of as-deposited films, under 6 seem CH4 flow rate condition,

on the pretreated surface, in case of (a) the normal process, (b) 180130 see cyclic process,

and (e) 301180 see cyclic process, respectively.

To manifest the enhancement of the selectivity of diamond film deposition by the cyclic process, we investigated

the surface images of as-deposited diamond films after 6 h deposition reaction, especially around boundary areas

between the pretreated and the untreated areas on glass substrate. Figs. 5 and 6 show SEM images around boundary

areas of as-deposited diamond films in the normal process (Figs. 5a and 6a) and in the cyclic process having the

growing/etching time = 180/30 sec (Figs. 5b and 6b) and 30/180 sec (Fig. 5c and 6c), respectively. As shown in these

figures, on the untreated area of glass substrate, we could not observe any distinct difference of the diamond

nucleation densities between the cyclic process and the normal process. On the pretreated area of glass substrate,

however, we could already observe the enhancement of the diamond nucleation densities (ref. 5). Indeed, we can

obtain the highest diamond nucleation densities on the pretreated area, with little change of the diamond nucleation

densities on the untreated area, by using the cyclic process having the less etching time interval. Therefore, we

suggest that the cyclic process having less etching time interval may be the optimal condition for the enhancement of

the selectivity of diamond film deposition on glass substrate.

, Z.... ........ e ........... ,I "_ e _ ..... ,.

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::: ::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::: :::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

Figure 5. SEM images of as-deposited films, under 4 seem CH4 flow rate condition, around the boundary

areas between the pretreated surface and the untreated surface, in ease of (a) the normal process,

(b) 180130 see cyclic process, and (e) 301180 see cyclic process, respectively.

Figure 6. SEM images of as-deposited films, under 6 seem CH flow rate condition, around the boundary

areas between the pretreated surface and the untreated surface, in ease of (a) the normal process,

(b) 180130 see cyclic process, and (e) 301180 see cyclic process, respectively.

NASA/CP--2001-210948 270 August 6, 2001
Thecausefortheenhancement of theselectivityof diamondfilmdepositiononthepretreatedglass

substratebythecyclicprocessmaybeattributedtothecyclicmodulationftheH2/CH4concentration ratiointhe

sourcegas.Aspreviousreports(ref.7),thepretreatedareahasroughsurfaces. Theroughsurfacecanenhancethe

suitablediamondnucleationsitesbythecyclicprocess. Inthecaseoftheuntreatedarea,however, therelatively

smoothsurfacestatesseemtobeinsufficient forreadilydiamondnucleationbythecyclicprocess. Consequently, in

caseofglasssubstrate,hecyclicprocesscanenhancetheselectivedepositionof diamondfilmviatherelative

increaseinthenucleationdensitiesonthepretreatedareathanontheuntreatedarea.

Inaddition, theoptimal conditionfortheenhancement oftheselectivityofdiamondfilmdepositionglass

substratemayberelatedtotheinjectionamount ofcarbonsourcespeciesduringthecyclicprocess. Thelessetching

timeinterval(180/30sec)conditioncanintroducethemoreamount ofthecarbonspecies, comparedwiththemore

etchingtimeintervalcondition(30/180sec)inthecyclicprocess. Therefore, 180/30seccyclicprocessmayimpose

morediamondnucleationdensitiesthan30/180seccyclicprocess, duetotheabundant amount ofthecarbonspecies.

Consequently, 180/30seccyclicprocesswouldbetheoptimalconditionfortheenhancement of diamondfilm

depositionglasssubstrate.

Toinvestigatetheenhancement of thediamondqualityby theapplicationof thecyclicprocess, we

investigated thegrainsonthepretreatedglassubstrateusingmicro-Raman spectroscopy asshowninFigs.7a- c.

Therelativeintensityratio(IJIa)ofdiamond(at1332cm-1)toamorphous carbon(around1500cm-1)increasedin

thecyclicprocess(compareFigs7awithb andc),indicatingtheincreasein therelativefractionof diamond

component ithegrain.It revealsthereadilyformationofdiamond, insteadof amorphous carbon, bythecyclic

process.

• • ............................/............/ .............../....../.................

........................................................................................

. ......../..//.................................../................/......,../ ......

•.........................................................................................

ii!i!iiiii!i!iiiiiiiii!!}:iiii!iiiiiiiii!iiii!!i_! ::!!i!!iiiiii!!iiiiiiii!ii!_iiiiiii

::::::::::::::::::::::::::::::::::::::::::::::::::::

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

..,._-, taj.y.:.:.:.: :.:.,...,_.:.:.:.:.:.:.i-1._:::::.:.:.: :.:.:.:.:.:.:.:.:.:.:.i.:.:.:.:.:.:.:

,--m _:::!:!::::::::::/:::::::::: :::::::::::::::::::::::::::::::::::::::::::::::::::::

__ ::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

iii!i!iiiiiii_'iiiiiiiiiiiiiiiiiii!iiiiiiiiiiiiiiiiii_:i:_iiiiiiiiiiiiiiiiii!iiiiii

iiiii!i:;!!_iiiiiiiii{i:_ iiiiiiiiiiiiiiiii!_!_T!-iii:iiiii:i;_i+_#2

12o0 1300 1400 1500 1600 1700

Raman shift (cm -_)

Figure 7. Raman spectra of as-deposited films, under 4 sccm CH4 flow rate condition, in case of (a) the

normal process, (b) 180/30 sec cyclic process, and (c) 30/180 sec cyclic process, respectively.

In case of the cyclic process, we also observed the increase of Id/Ia with increasing the etching time interval

in the cyclic process (compare Figs. 7b with c). Obviously, it reveals the enhancement of the diamond quality for the

grain by the cyclic process having more etching time interval. The relative high ratio of atomic hydrogen during the

30/180 sec cyclic process seems to be the cause for the diamond quality enhancement of the grain, because atomic

hydrogen can more readily etch away amorphous carbon than the diamond component in the grain. Consequently,

the more etching time interval in the cyclic process can produce the higher relative fraction of diamond component in

the grain.

NASA/CP--2001-210948 271 August 6, 2001
CONCLUSIONS

Thediamondnucleationdensitiesonthepretreatedglassubstratewouldbemuchmoreenhancedunderthe

conditionofthelessetchingtimeinterval inthecyclicprocess. Consequently, thecyclicprocesshavinglessetching

timeinterval maybetheoptimal conditionfortheenhancement oftheselectivityofdiamondfilmdepositionglass

substrateviatherelativeincreaseinthenucleationdensitiesonthepretreatedarea,comparedwiththeuntreatedarea.

Thecauseforthismaybeattributedtotheincreaseintheamount ofcarbonspeciesduringthecyclicprocesswith

decreasing theetchingtimeinterval.

Theenhancement of thediamondqualityforthegrainbythecyclicprocesswasnoticeableunderthemore

etchingtimeinterval inthecyclicprocess. Therelativehighratioofatomichydrogenduringthehighetchingtime

interval cyclicprocessseemstoberesponsible fortheenhancement ofthediamondquality.

ACKNOWLEDGEMENT

ThisworkwassupportedbygrantNo.(2000-1-12100-002-1) fromtheBasicResearchProgramoftheKorea

Science&Engineering Foundation

REFERENCES

1. T. D. Moustakas : "Growth of diamond by CVD methods and effects of process parameters," in Synthetic

Diamond: Emerging CVD Science and Technology, eds. K. E. Spear and J. R Dismukes, John Wiley & Sons,

1994, pp. 145-192.

2. Eddy C. R., et al.: Diamond thin film growth on silicon at temperatures between 500 and 600 degree C using an

electron cyclotron resonance microwave plasma source, Surface & Coatings Technology, Vol. 48, 1991, pp. 69-

79.

3. Li I.,; Aslam D. M.; and Kwon S. J.: Field emission from undoped polycrystalline diamond deposited by MPCVD

at 520-660 C, Technical Digest of IVMC'97, Edirak, 1997, pp. 509-512.

4. Tolt Z. L.; Fink R. L.; and Yaniv Z.: Electron emission from patterned diamond film, Technical Digest of

IVMC'97, Edirak, 1997, pp. 171-175.

5. Kim S.-H.; Kim T.-G.; and Kim Y.-H.: Effect of cyclic process on the selective deposition of diamond film on

glass substrate, Proceedings of the Sixth International Symposium of Diamond Materials. The electrochemical

society, 1999, pp 110-119.

6. Kim S.-H., et. al.: Effect of cyclic process on the { 100 }-oriented texture growth of diamond film, Appl. Phys.

Lett., Vol. 69 No. 15, 1996, pp. 2184 -2186.

7. Kim S.-H., et. al.: Effect of the substrate state on the formation of diamond film in a low temperature microwaveplasma-enhanced chemical vapor deposition system, J. Vac. Sci. Technol. A, Vol. 13, No. 3, 1993, pp. 1619-1623.

NASA/CP--2001-210948 272 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Diamond nucleated by arc discharge and grown by MWPCVD

Changzhi. Gu and Zengsun Jin

The State Key Laboratory for Superhard Materials, Jilin University,

130023 Changchun, China

Fax: 0086-431-8923907 email: czgu@public.cc.jl.cn

An arc discharge method was applied to form diamond nuclei on mirror-polished

Si substrate, the properties of nucleation at different distances between substrate and

electrode and growth by microwave plasma chemical vapor deposition (MWPCVD)

were studied and compared with those on scratched Si substrate with diamond powder

and mirror polished Si substrate nucleated by bias enhanced nucleation (BEN) method.

Scanning electron microscopy (SEM) and Raman Spectroscopy (RS) were using to

analyze the film morphology and quality. The results indicate that high nucleation

density and high quality diamond films can be formed on arc discharge treated mirrorpolished Si substrate under selected deposition parameters.

Keywords: diamond, CVD, arc discharge, nucleation, growth

NASA/CP--2001-210948 273 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

HOMOEPITAXIAL DIAMOND GROWTH AT HIGH AND LOW

TEMPERATURES: DESIGNER DIAMONDS TO LARGE CRYSTALS

Yogesh K. Vnhra

Department of Physics

University of Alabama at Birminghmn (U,,M3)

Binrtingham, AL 35294-1 t70, USA

E-mail: ykv0hra,_c_uab.edu

ABSTRACT

This abstract will cover recent advances in d_e homoepitaxial dimnond growth by microwave

plasn_l chemical vapor deposition. We will explore _vo growflt regimes, one at fl_e temperatures in fl_e

range of 1200 to 1300 °C and another one at temperatures near 800 °C. tn a collaborative projec! with

Lawrence Livermore National Laboratory, high quality single crystalline dimnond films have been grown

on diantond mwils as snbsl_ales with encapsulated metal nticreciretdts to fabricate "Designer Dimnond

Anvils" for applications in high-pressure science and technology.

Designer Dimntmd in reflected light

showing exposed microprobes on d_e surface

lk

l 2ram

Designer diamond in Ozmsmitted lighl

showing ci_tt bmied microprobes

3.6 rant x 3.8 refit x 2amn

homoepitaxiaIty grown diamond

cu'stal bt a 6kW microuave plasma

c\rD syslem on a type Ib yellow

substrate

Single c_Tstal diamond can be grown honloepita,:ially at high temperatures of 1300 °C wilh g_'m'dh ':ate

approaching 30-40 microns/hour with methane/hydrogen/nitrogen chemistry on a type Ib substrates. Large

dmmond cry'slals up to 4 nun em_ be grown but the edge of cWst',ds show multiple t_im_htg and lhmt the

ultimate size of these diantonds. Alternative, low temperature chemistu- of methane,qtyd rogergox'ygea tins

been probed to grow large crystals starting with type iia diamond snbstrates. The eustallino quNity of the

homoepitaxial dimnond eO'stals was evaluated by the miero-tknnmt spectroscopy and single c_Tslat x-ray

diffraction using the rocking curve measurements of (400) diffraction peak from dimnond. Future research

d_rections in large area hontoepitaxi'ed diamond _ov_th wilt be d_scussed.

We acknowledge support from the National Science Foundation (NSF) _mder Gn-mt ,No. 9704428

and support by the B-Division al the Lawrence Livennore National Laborator_ (LLNL) under the auspices

of the U.S Department of Energy, under contract No. W-7405-ENG-48. Author would also like to

acknowledge contributions from Ns coworkers Dr. Shane A. Cattedge and Dr Ch_h-Sltiue Van from UAB

and Dr. Jag,m_Akel!a _md Dr. Sam Weir from the Lawrence Livermore National Laboratory..

Ke>avords: homocpitaxy, designer diamonds, defects, twimting

NASA/CP--2001-210948 274 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

BORON-DOPED HOMOEPITAXIAL DIAMOND (100) FILM INVESTIGATED BY

SCANNING TUNNELING MICROSCOPY

Blng Xlao, Welhal Fu, and Saeharla Albln

Microelectronics Laboratory

Department of Electrical and Computer Engineering

Old Dominion University, Norfolk, VA 23529, USA

Jason Moulton, and John Cooper

Department of Chemistry and Biochemistry

Old Dominion University, Norfolk, VA 23529, USA

ABSTRACT

Conducting epitaxial diamond films of high quality are essential for many diamond studies and diamond

electronic device fabrication. We have grown boron-doped epitaxial diamond films on type IIa natural diamond

(100) substrates by microwave plasma chemical vapor deposition. A gas mixture of H2/CH4 was used. Boron doping

was done by placing solid sources of pure boron in the microwave plasma. Homoepitaxial films with atomic

smoothness were achieved under the following growth conditions: substrate temperature 900 °C, gas pressure 40

Torr, and gas flow rates of H2/CH4 = 900/7.2 sccm. The growth rate was 0.87 gm/hr. Surfaces of the homoepitaxial

films were studied by scanning tunneling microscopy (STM). STM images show smooth and continuous surface with

ripple-like features on micrometer scale. On nanometer scale, alternating terraces of 2x 1 and lx2 dimerization were

clearly observed.

Keywords: diamond thin film, homoepitaxy, scanning tunneling microscopy (STM).

INTRODUCTION

Boron-doped homoepitaxial films are commonly used for the study of various properties of diamond thin films,

such as surface structures (refs. 1 to 3), electrical or electronic characteristics (refs. 4 to 6), and electrochemical

behavior (refs. 7 and 8). In addition, such films are used in many diamond electronic devices (refs. 9 and 10). For

epitaxial growth of diamond, microwave plasma chemical vapor deposition (MPCVD) has been widely used and

studied (refs. 1 to 6, 9). A lot of work has been done on surface morphology and structures of polycrystalline (refs.

11 and 12) or single-crystalline (refs. 1 to 3, 13) CVD diamond thin films, which show strong dependence on growth

conditions. In these studies, images from optical microscopy and scanning electron microscopy (SEM) are usually

used to show surface morphology on micron or sub-micron scale, while scanning tunneling microscope (STM) and

atomic force microscope (AFM) are used to observe the surface from micron level to atomic level.

In this study, STM was utilized to investigate the surface morphology of boron-doped homoepitaxial diamond

(100) films on both micron and nanometer scale.

EXPERIMENTS

Type IIa (100)-oriented natural diamond substrates (3x3x0.25 mm 3) were used in this study. Diamond substrates

were degreased in acetone and ethanol, then cleaned in a 1:1 solution of HNO3 and HF, and in a 1:3 solution of

HNO3 and HC1. They were finally rinsed with deionized water and mounted in a shallow dip on a Si wafer which

was used as the sample holder.

Diamond epitaxial growth was carried out in the 6-inch cylindrical chamber of a MPCVD system (ASTeX).

Hydrogen and methane were used as reactant gases. The gas flow rates of H2 and CH4 were 900 and 7.2 sccm,

*Author to whom all correspondence should be addressed; email: salbin@odu.edu, fax: (757)683-3220, phone:

(757)683-4967.

NASA/CP--2001-210948 275 August 6, 2001
respectively. ThediamondsubstratemountedontheSisampleholderwasplacedonagraphiteheaterintheCVD

chamber. Theheatertemperature wascontrollableduringthewholeCVDprocess. Thegaspressurewasmaintained

at40Torrduringthediamondgrowth.

TheCVDchamber wasfirstpumpeddowntoabasepressureblow10 -4 Torr. Then diamond substrate was

heated up in vacuum. When graphite heater reached 900 °C, hydrogen was introduced into the chamber and

hydrogen plasma was ignited by a 2.45 GHz microwave input with a power of 1000 W. After the sample was treated

in the hydrogen plasma for 5 min, CH4 was introduced into the chamber to begin diamond growth. The growth time

was 1 hr. Boron doping was done by placing four small pure boron pieces around the diamond substrate. Microwave

power and CH4 flow were turned off to stop diamond growth, and the sample was cooled down to 600 °C in H2. At

600 °C, hydrogen plasma was started again to treat the diamond surface for another 5 min. Finally, the diamond

sample was cooled down to room temperature in H2 gas ambient.

(a)

(b) (c)

Figure 1. Height mode STM images of diamond (100) surfaces: (a) surface before CVD growth; (b)

and (c) surface of the boron-doped homoepitaxial film.

NASA/CP--2001-210948 276 August 6, 2001
(a) (b)

Figure 2. STM images of the homoepitaxial diamond (100) film showing 2×1 surface reconstruction. (a)

A large-area height mode image shows alternating terraces of 2×1 and 1×2 dimerization. (b) Current

mode image shows that individual dimers are resolved.

Diamond samples were investigated in air by a commercial STM system (NanoScope III). Mechanically cut PtIr or ac-etched tungsten tips were used to probe the diamond surface. For scanning range on micrometer scale, height

mode operation was applied while images with atomic resolution were obtained in both height mode and current

mode.

RESULTS AND DISCUSSIONS

The homoepitaxial growth rate of 0.87 gm/hr was calculated from the sample weight increase after CVD

deposition. After the sample was boiled in H2SO4 or HNO3, the surface of the epitaxial film still had a relatively high

conductivity of approximately -10-1 _.cm. This result demonstrate that the electrical conduction was created by

boron doping, not by the surface hydrogenation since the high-conductivity layer near the surface, induced by

hydrogen incorporated in the surface region, can be removed by oxidation of the surface using strong acid solutions

such as H2SO4 and HNO3 (ref. 6).

Figure l(b) and l(c) are STM images showing the surface of the deposited diamond (100) film on micron and

sub-micron scale. For comparison, Figure 1(a) shows the STM image of the polished diamond surface prior to

epitaxial growth. The diamond substrate was treated in hydrogen plasma for 5 min to obtain surface conduction

needed for STM imaging. As shown in Figure l(a), the parallel-groove pattern is considered to be formed during the

surface polish process (ref. 14), and the surface roughness is 0.583 nmm_ and 4.509 nmpp. After homoepitaxial

growth, the parallel-groove pattern is still visible but smaller groove depths make a smoother surface with a

roughness of 0.384 nmm_s and 2.781 nmpp (Figure l(b)). In Figure l(c), the long-range height variation is also due to

the parallel-groove pattern and the surface appears consisting of nano-scale features which are believed to be the

terraces formed during the diamond deposition (ref. 15). The surface roughness in this 500x500 nm 2 area is 0.096

nmn_ and 0.893 nmpp. It is expected that the surface height variation could be further reduced with a longer

annealing time in hydrogen plasma and a thicker film grown for longer time.

The atomic-resolution STM pictures are shown in Figure 2. Terraces with alternating 2xl and lx2 dimerization

are clearly displayed and individual dimers are also resolved. This kind of surface structure, i.e. double-domain

surface, is considered to be created by two-dimensional nucleation (ref. 15). The average terrace width is -4 nm

while adjacent terraces are separated by a single-layer step.

SUMMARY

Boron-doped homoepitaxial diamond (100) films have been grown on type IIa natural diamond substrates. The

surface morphology was observed using STM. Compared with the diamond surface before CVD growth, the

NASA/CP--2001-210948 277 August 6, 2001
depositedsurfaceissmoother withlessurfaceheightvariation. Atomic-resolution STMimageswereobtained

revealingthesurfacerconstruction of2xldimerization.

ACKNOWLEDGMENT

ThisworkissupportedbyagrantfromNational ScienceFoundation.

REFERENCES

1. Kuang, Y., Wang, Y., and Lee, N., et al.: Surface Structure of Homoepitaxial Diamond (001) Films, A

Scanning Tunneling Microscopy Study. Appl. Phys. Lett., vol. 67, no. 25, Dec 1995, 3721.

2. Kawarada, H., Sasaki, H., and Sato, A.: Scanning-Tunneling-Microscope Observation of The Homo

epitaxial Diamond (001) 2xl Reconstruction Observed Under Atmospheric Pressure. Phys. Rev. B, vol. 52,

no. 15, Oct 1995, 11351.

3. Takami, T., Kusunoki, I., and Nishitani-Gamo, M., et al.: Homoepitaxial Diamond (001) Thin Film Studied

by Reflection High-Energy Electron Diffraction, Contact Atomic Force Microscopy, and Scanning

Tunneling Microscopy. J. Vac. Sci. Technol. B, vol. 18, no. 3, May/Jun 2000, 1198.

4. Shiomi, H., Nakahata, H., and Imai, T., et al.: Electrical Characteristics of Metal Contacts to Boron-Doped

Diamond Epitaxial Film. Jpn. J. Appl. Phys., vol. 28, no. 5, May 1989, 758.

5. Shiomi, H., Nishibayashi, Y., and Fujimori, N.: Characterization of Boron-Doped Diamond Epitaxial Films.

Jpn. J. Appl. Phys., vol. 30, no. 7, Jul 1991, 1363.

6. Hayashi, K., Yamanaka, S., and Okushi, H., et al.: Study of The Effect of Hydrogen on Transport Properties

in Chemical Vapor Deposited Diamond Films by Hall Measurements. Appl. Phys. Lett., vol. 68, no. 3, Jan

1996, 376.

7. Yanagisawa, M., Jiang, L., and Tryk, D.A., et al.: Surface Morphology and Electrochemical Properties of

Highly Boron-Doped Homoepitaxial Diamond Films. Diamond Relat. Mater., vol. 8, no. 11, Nov 1999,

2059.

8. Swain, G.M., Anderson, A.B., and Angus, J.C.: Applications of Diamond Thin Films in Electrochemistry.

MRS Bulletin, vol. 23, no. 9, Sept 1998, 56.

9. Shiomi, H., Nishibayashi, Y., and Fujimori, N.: High-Voltage Schottky Diodes on Boron-Doped Diamond

Epitaxial Films. Jpn. J. Appl. Phys. Part 2, vol. 29, no. 12, Dec 1990, L 2163.

10. Ariki, T., Shikama, S., and Suzuki, S., et al.: Efficient Field Effect in Heavily Doped Thin-Film Diamond

Metal-Insulator-Semiconductor Diode Employing BaTiO3 Insulator Film. Jpn. J. Appl. Phys. Part 2, vol. 33,

no. 6B, Jun 1994, L 888.

11. Zhou, D., Gruen, D.M., and Qin, L.C., et al.: Control of Diamond Film Microstructure by Ar Additions to

CHJH2 Microwave Plasmas. J. Appl. Phys., vol. 84, no. 4, Aug 1998, 1981.

12. Fitzgerald, A.G., Fan, Y., and John, R, et al.: Characterization of the Surface Morphology and Electronic

Properties of Microwave Enhanced Chemical Vapor Deposited Diamond Films. J. Vac. Sci. Technol. B, vol.

18, no. 6, Nov/Dec 2000, 2714.

13. Shiomi, H., Tanabe, K., and Nishibayashi, Y., et al.: Epitaxial Growth of High Quality Diamond Film by

The Microwave Plasma-Assisted Chemical-Vapor-Deposition Method. Jpn. J. Appl. Phys., vol. 29, no. 1,

Jan 1990, 34.

14. Spear, K.E., and Dismukes, J.: Synthetic Diamond: Emerging CVD Science and Technology. John Wiley &

Sons, New York, NY, USA, 1994, 322.

15. Tsuno, T., Tomikawa, T., and Shikata, S.: Diamond (001) Single-Domain 2xl Surface Grown by Chemical

Vapor Deposition. Appl. Phys. Lett., vol. 64, no. 5, Jan 1994, 572.

NASA/CP--2001-210948 278 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CATHODOLUMINESCENCE TOPOGRAPHY OF HOMOEPITAXIAL DIAMOND MADE BY

MICROWAVE PLASMA CVD

Hisao Kanda, Satoshi Koizumi, Kenji Watanabe

National Institute for Materials Science( NIMS ), 1-1 Namiki, Tsukuba, Ibaraki 305-0044, Japan

KANDA.Hisao @nims._o.jp, TEL&FAX: +81-298-51-4005

Tokuyuki Teraji

Department of Electric Engineering, Osaka University, 2-1 Yamada-oka, Suita, Osaka 565-0871, Japan

ABSTRACT

Diamond film has been made with adding PH3 in order to prepare n-type semiconducting diamond.

The diamond film was deposited on {001 } high pressure synthetic diamond by a microwave plasma CVD

method, and was characterized with cathodoluminescence spectroscopy and topography to obtain information

of defects. A 239 nm peak, an exciton bound to P, and a broad band with a maximum around 270 nm were

found, indicating that phosphorus is doped in the film grown on { 100 } substrate as well as {111 }. In addition,

it was also found that sharp peaks at 415,484, 500, 514 and 532 nm, which are common in CVD diamond, are

located in small diamond particles deposited with orientations independent of diamond substrate.

Keywords: CVD diamond, P doping, Cathodoluminescence

INTRODUCTION

Physical properties of diamond are sensitively influenced by impurities or defects present in diamond, and it

is important to understand structure and character of the defects, and then to establish methods to control

them.

Luminescence spectroscopy is an easy and sensitive method to detect defects in diamond. However,

spectroscopic data do not directly provide information of the defects such as defect structure and

concentrations. In order to increase power of the method, it may be a way to accumulate data which give

correlation between spectra and conditions under which luminescence bands appear or disappear.

There are a number of reports on the luminescence spectroscopy(ref. 1), but most of them are on high

pressure and natural diamond. As for CVD diamond, relatively few reports have been published. In CVD

diamonds, luminescence bands unique in CVD diamond have been observed as well as those common in

natural and HPHT diamonds. In this report, we describe luminescence bands observed in CVD diamond

doped with phosphorus during growth, and distribution of the bands is also described.

EXPERIMENTAL

We made diamond film using a microwave assisted plasma CVD system in which a metallic container was

used to avoid contamination of silicon impurity. High pressure Ib diamond (001) surface with 2 mm

square( Sumitomo Electric Industries Co.) was used as a substrate. Conditions of the synthesis are: 0.1-0.5%

of methane content relative to hydrogen, 80 - 200 Torr of total gas pressure, 0.2 - 1% of PH3 concentrations,

860 - 920 °C of substrate temperature, 2 - 20 hr of growth time.

Surface of the grown diamond film was observed using an Olympus optical microscope and a JEOL JSM5410 scanning electron microscope.

Cathodoluminescence( CL ) spectra and images were taken using a Topcon SX-40A scanning electron

microscope fitted with a SPEX 1/3 m spectrometer. CL light monochromitized by the spectrometer was

detected by a photomultiplier to display the signal on CRT to show CL image. Princeton Instruments CCD

NASA/CP--2001-210948 279 August 6, 2001
wasusedtorecordCLspectra. InordertodetecttheCLsignalwithhighintensity, slitwasopenedtolmm

wide,leadingtorelativelywiderluminescence peaks.

Sampleswerecoatedwithgold,buthecoatwassothinthatluminescence passedwithoutconsiderable

absorption. Thesamplewascooledto-160°Cwithliquidnitrogen.

RESULTS AND DISCUSSION

Surface morphology

Fig.3(a) shows an example of surface morphology of the grown diamond. Growth steps and fine particles

are seen, as described previously on (001) homoepitaxial diamond(ref. 2). Growth steps are high enough to be

clearly seen even by SEM. They are undulated, and sometimes form growth hillocks with square shape.

It is noted that fine particles are formed on the surface. They are diamond with 10 to 20 pm in diameter and

exhibit cubo-octahedral morphology. Twin planes are commonly seen in the particles. Their orientation is

independent of that of substrate. (111) surface of the particle is sometimes seen parallel to (001) of substrate.

Even if (001) of the particle is parallel to (001) of the substrate, <110> edges of the particle were not

coincident with those of the substrate.

Cathodoluminescence spectra

Several luminescence peaks were observed from the grown surface of diamond. They all have been

documented(ref. 1).

Peaks related to phosphorus impurity(ref. 3) are shown in fig. 1, which demonstrates two spectra taken from

different samples. A peak, (b), between two free exciton peaks, (a) and (c), is the P-related exciton peak,

which is located at 239 nm. A broad band, (d), with a maximum around 270 nm is also related to phosphorus.

Similar peaks have been described in boron doped diamond(ref. 4). However, peak position of the exciton

bound to phosphorus is 2 nm larger than that of the exciton bound to boron. Position of the broad band is also

different. Peak position of the P-related one is 10 nm larger than that of B-related one. Presence of the Prelated peaks confirms that phosphorus are incorporated into growth layer deposited on { 100 } substrate as

well. Previous reports on P-doped diamond were based on investigation of diamond film made on only { 111 }

surface(ref. 3). The broad band, (d), is present in one of the two spectra shown in fig. 1, suggesting that the

phosphorus related peaks, (b) and (d), are not directly correlated.

Q_

0 .......................................................................... ,

(a) (d)

20 (b) /p .... _,w.._

/

15 , _

220 240 260 280 300

Wave length( nm )

Fig. 1 Two examples of cathodoluminescence spectra of phosphorus doped {001 } CVD homoepitaxial

diamond in a UV region. Two peaks, (a) and (c), are free exciton peaks, and a peak (b) is an exciton peak

bound to phosphorus. A broad band, (d), is also phosphorus related.

NASA/CP--2001-210948 280 August 6, 2001
Otherpeakswhicharenotrelatedtophosphorus werealsoobserved. Twoexamplesofspectraareshowni

fig.2.

Spectrum(1)inFig.2exhibitsfivesharpeaksat415,484,500,514and532nm.Theyareunique, but

commonlyobservedinCVDdiamond(ref.5to13).However, theircharacter anddefectstructurearepoorly

understood. Relativeintensityofthefivepeakswasvariableinspectraaccordingtomeasurements for6films,

indicatingthatseveral typesofdefectsareresponsible forthepeaks. Inthisstudy,thediamondfilmwasso

thinthatluminescence fromthesubstratewasobservable. TheH3bandand2.56eVbandproducedfromthe

substrate, i.e.highpressuresyntheticdiamond, weresometimes superimposed inspectrameasured, butit was

possibletotakespectraavoidingtheluminescence fromthesubstrate, assectordependence ofluminescence

observedinthesubstratewasconsidered. Thus,it wasconfirmedthathevariablerelativeintensitymeasured

isnotaffectedbytheluminescence ofthesubstrate.

Fig.2 Cathodoluminescence spectraofphosphorus doped{001} CVDhomoepitaxial dmondinavisible

region.Spectra(1)and(2)weretakenfromdiamondparticlesandepitaxial layer,respectively. (a)- (e)are

peaksfounduniquelyfromCVDdiamond, whereas(f)and(g)areseenfromnaturalndHPHTdiamondas

well.

Itisreportedthatintegral intensityofthe484nmisproportional tothatofthe514nmpeak(ref. 11).Most

ofmeasurement inthisstudywasingoodagreement withthereport,buttherewasacasethathe484nmpeak

wasstrongwhilethe514nmwasnegligible. It isexpectedthathis484nmpeakhasdifferent originfromthat

whichisproportional tothe514nmpeak.

Thepeakat532nmmustbeascribedtoatleasttwodifferent defects, ashasbeenreported(ref. 6).There

wasaspectruminwhichapeakat532nmhasashapeconsistingoftwopeakswithdifferent width.Asharp

532nmpeakhappened tobeobservedaccompanied bythe575nmband,(g).Thisfeaturehasbeendescribed

byVavilovetal(ref.5).

Abroadband,(f),withamaximumaround430nmseeninbothspectra, (1)and(2),isthe"bandA",which

arecommonlyobservedveninhighpressurediamondandnatural diamonds. Thepeak,(g),isazerophonon

lineofthesocalled575nmbandascribedtoanitrogen-vacancy complex.

Theluminescence peaksdescribedabovewereinhomogeneous inadiamondfilm,andluminescence images

areshownithefollowingsection.

Cathodoluminescence images

Monochromatic CLimagesweretakenatthreedifferent wavelengths of430,500and575nm.Anexample

oftheimagesisshowninfig.3.Theimagesat430nmand575nmindicatedistributionfthebandAandthe

NASA/CP--2001-210948 281 August6,2001
575nmband.Theimageat500nmindicatesdistributionofthefourpeaksbetween480and532nm.Wehad

expectedthatdistributionofeachpeakofthefourisfoundfromimagesif theyaretakenat484,500,514and

532nm.However, itwasnotpossibletotakethemathighresolutionsothateachpeakisresolved, becausethe

slitatthespectrometer hadtobeopenedsothatCLsignalisstrongenoughtogiveimages. Wewereafraid

thatimagesat500nmisinfluencedbythatofthebandA.However, it hasbeenconfirmedthatheimageat

500nmisnotofthebandA.

Asshownifig.3,thereisacontrast inCLintensitybetweensmalldiamondparticlesandepitaxial growth

layer.CLintensityat430nmishighattheparticlesincontrast thatit ishighontheepitaxial layerat575nm.

TheimagesindicatehatbandAisstrongaroundtheparticlesashasbeenpublishedbefore. Itisreasonable

thathecrystal ishighlydefectivetoproducethebandAluminescence, becausecrystal latticeiscombined

withdifferent orientationattheinterfacebetweentheparticlesandtheepitaxial layer.

ItwasfoundforthefirsttimethatheCLpeaksaround500nmmainlyarestrongintheparticlesasseen

fromfig.3(c).PositionsofthebrightspotwellfitwiththediamondparticlesshownitheSEM

image(fig.3(a), andalsowiththebrightspotofthebandA.Itmaybepossiblethathebrightspotsat500nm

areduetocontributionftailofthebandA.However, itisreadilyconfirmedthathebrightspotsintheimage

at500nmarenotofthebandA.Comparingfig.3(b)and(c),it isfoundthathereissomedifference. Regions

indicatedbyarrowsarebrightat430nm,butnotat500nm,wherediamondparticlesarenotpresent. The

observation revealsthatheCLpeaksaround500nmarelocatedatdiamondparticlesandthathebandAis

observedvenatregionswhereadiamondparticleisnotpresent.

Fig.3 Cathodoluminescence imagesofadiamondfilmgrownona(001)HPHTdiamondtakenatdifferent

wavelength. (a)SEMimage, CLimagestakenat(b)430nm,(c)500nmand(d)575nm.Twospotsindicated

byarrowsarebrightat430nm,butdarkat500nm,confirmingthatbrightspotsin(c)arenotduetobandA.

Atrianglebrightregionseenin(c)showsagrowthsectorofthesubstrate.

NASA/CP--2001-210948 282 August6,2001
Ithasbeenreportedthathepeakat500nm,whichisoneofthepeaksobservedinthediamondparticles, is

strongin{111} growthsectorofadiamondparticlegrownbytheCVDmethod(ref. 12).Thepresent

experimental resultthatCLimageat500nmisbrightatthediamondparticlesmaybeconsistent wihthis

previousreport,becausethediamondparticlesexhibit{111} growthsector, whereastheepitaxial layerhas

only{001}surfaces. Luminescence spectraweretakenfromseveral diamondparticles. Accordingtothe

measurements, therewastendencythathefivepeaksat415,484, 500,514and532nmarestronger from

particlesonwhich{111}surfacesaredominant thanfrom{100} dominant particles. However, moredetailed

observationsrequiredtoconfirmthesectordependence ofthefivepeaks.

CLimagesoftheP-relatedpeakswerenotakensuccessfully, becausetheluminescence intensitywastoo

weaktoberevealedontheCRTdisplay.

SUMMARY

Phosphorus doped diamond film was deposited by the microwave assisted CVD method on a {001 }

diamond surface, and was characterized with cathodoluminescence spectroscopy and topography. An exciton

peak and a broad band, both of which are P-related, were detected, indicating that phosphorus is incorporated

even in the {001 } epitaxial layer. Luminescence peaks at 415,484, 500, 514 and 532 nm, which are uniquely

in CVD diamond, are preferentially observed from particles which nucleated on the substrate independently.

ACKNOWLEDGEMENT

This work was partly supported by the FCT project, which is cosigned to JFCC by NEDO.

REFERENCES

1. Zaitsev, A.M.: Handbook of Industrial Diamonds and Diamond Films, edited by M.A.Prelas,

G.Popovici, L.K.Bigelow, Marcel Dekker, New York, 1997, pp.227-376

2. Takeuchi,D et al: Homoepitaxial diamond films grown by step-flow mode in various misorientation

angles of diamond substrates, Diamond and Relat. Mater. vol.9, no.3, 2000, 231-235

3. Sternschulte,H. et al: Optical evidence for 630-meV phosphorus donor in synthetic diamond,

Phys.Rev. B, vol.59, no.20, 1999, 12924-927

4. Lawson,S.C. et al: Cathodoluminescence from high-pressure synthetic and chemical-vapor-deposited

diamond, , J.Appl.Phys. vol.77, no.4, 1995, 1729-1734

5. Vavilov, V.S. et al:, Investigation of the cathodoluminescence of epitaxial diamond films, Sov. Phys.

Semicond. vol. 14, no.9, (1980) 1078-1079

6. Collins,A.T.; Kamo,M.; Sato,Y.: Optical centres related to nitrogen, vacancies and interstitials in

polycrystalline diamond films grown by plasma-assisted chemical vapour deposition,

J.Phys.D:Appl.Pys. vol.22, no.9 1989, 1402-1405

7. Collins,A.T.; Kamo,M.; Sato,Y.: Intrinsic and extrinsic cathodoluminescence from single-crystal

diamonds grown from by chemical vapour deposition, J.Phys.: Condens. Matter. vol.1, no.25, 1989

4029-4033

8. Ruan,J; Choyke,W.J.: Cathodoluminescence and annealing study of plasma-deposited polycrystalline

diamond films, J.Appl. Phys. vol.69, no.9, 1991, 6632-6636

9. Ruan,J.; Kobashi,K; Choyke,W.J.: On the "band-A" emission and boron related luminescence in

diamond, Appl.Phys.Lett. vol.60, no.25, 1992, 3138-3140

10. Collins,A.T.: The characterization of point defects in diamond by luminescence spectroscopy,

Diamond and Relat. Mater. vol. 1, no.5-6, 1992, 457-469

11. Khong,Y.L.; Collins,A.T.: Temperature dependence of cathodoluminescence from CVD diamond,

Diamond Relat. Mater., vol.2, no. 1, 1993, 1-5

12. Robins,L.H.; Farabaugh,E.N.; Feldman,A.: Cathodoluminescence spectroscopy of free and bound

excitons in chemical-vapor-deposited diamond, Phys.Rev.B, vol.48, no. 19, 1993, 14167-181

13. Khong,Y.L.; Collins,A.T.; Allers,L.: Luminescence decay time studies and time-resolved

cathodoluminescence spectroscopy of CVD diamond, Diamond Relat. Mater., vol.3, no.7 1994, 1023-

1027

NASA/CP--2001-210948 283 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ORIGIN, EVOLUTION AND SOME PROBLEMS OF DIAMOND CVD

B.V.Spitsyn and A.E.Alexenko

Institute of Physical Chemistry RAS, Moscow, Russia

On totality of properties diamond may qualified as supermaterial. However it potential in

science and industry disclosed properly only with appearance of modern CVD methods of diamond and

diamond films (DF)synthesis. Issued on 1981 publication in the subject of Russian researchers

apparently triggered relative work of Japanese, and then and other our colleagues on all five continents.

Development of various methods in the diamond CVD, and impressive academic and practical results

are obliged to efforts of thousands researchers in hundreds laboratories. The authors of this short

review and their colleague in our Lab continue a modest contribution to some of the most challenging

directions. We prolonged in-situ doping of epitaxial DF by phosphorus and sulfur. Hall method

confirmed formation of diamond with an electron type conductivity. The combination of different

methods of vapor phase activation allows us essentially, in 2-3 times, to increase growth rate of

polycrystalline DF, with retaining of their quality. Obtained for the first time epitaxial A1N films on -

(111) diamond, is, supposedly the way to advanced heterostructures for solid state electronics.

Appreciating an impressive success of the diamond community, we shall note some unsolved

following problems: the mechanism of nucleation and growth CVD-diamond, synthesis continuous

ultra-thin DF, low-temperature DF deposition, DF heteroepitaxy, and controlled DF with n-type

semiconductivity.

The successful advance in the solution of the above-stated problems requires interdisciplinarary

approaches and international scientific cooperation. Outstanding opportunities for this provides us

forthcoming ADC/tZCT'01.

NASA/CP--2001-210948 284 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

COMPUTER SIMULATION OF LARGE AREA DIAMOND DEPOSITION BY DC ARC JET

WITH ARC ROOTS ROTATING AND OPERATING AT GAS RECYCLING MODE

F.X. L u, T.B. Huang, W.Z. Tang, J.H. Song, Y.M. Tong

University of Science and Technology Beijing, Beijing 100083, P.R. China

ABSTRACT

A computer model have been set up for simulation of the flow and temperature field, and the radial distribution

of atomic hydrogen and active carbonaceous species over a large area substrate surface for a new type dc arc plasma

torch with rotating arc roots and operating at gas recycling mode. The model was based on the structural feature of

the particular type of plasma torch, and on the principles of fluid dynamics, plasma physics, and plasma chemistry.

Rotating arc was treated as a localized arc column with a radius close to that of the torch orifice. Whilst the nonequilibrium arc plasma was treated as in a stead state of thermodynamically localized equilibrium in the scale of the

digitized cell size for numerical calculations. A gas recycling radio of 90% was assumed. In numerical calculation of

plasma chemistry, the Thermal-Calc program and a powerful thermodynamic database were employed. Calculated

results were used for prediction of the growth rate and the quality as well as their radial uniformity of diamond film

deposition over large area substrate.

Keywords- computer simulation, dc arc plasma torch, rotating arc roots, gas recycling, large area diamond

deposition

1. INTRODUCTION

Of the very many deposition methods, high power DC arc plasma jet is generally considered as the most

promising technique [1]. However, the conventional design of plasma torches based on the technologies of plasma

thermal spray or high power industrial arc-heaters suffer the disadvantages of high gas consumption, low heat

efficiency, the complicity and high investment of the equipment, and the most of all, the difficulties for large area

uniformity of diamond deposition. In our previous publications [2-4]we have disclosed a new type of magnetic and

fluid dynamic controlled large orifice long discharge tunnel plasma torch which guarantees the large area uniformity

by rotating arc root. In the present paper, a computer model have been set up for simulation of the flow and

temperature field, and the radial distribution of atomic hydrogen and active carbonaceous species over a large area

substrate surface for this particular type of dc arc plasma torch. Theoretical predictions on growth rate and quality

and the uniformity of diamond film deposition over large area substrate are presented.

2. THE COMPUTER MODEL

2.1 The large orifice arc plasma torch with rotating arc roots operating at gas recycling mode

Fig.1 (a) shows the photograph of the 100kW arc jet plasma torch with rotating arc roots operating at gas

recycling mode by which a growth rate as high as 40-50gm/h over a substrate area of _ll0mm maximum for

uniform diamond deposition can be realized [2-4] Fig.1 (b) and (c) show the arc plasma and the principle of arc

rotating for large area uniformity diamond deposition. As explained in our previous publications [2-4], carbon source

gas (e.g. CH4) can not be introduced through the cathode, and must be feed at a suitable position downstream.

However, the inlet cold carbon source stream will not automatically mix with the hot arc column, but instead cause it

to shrink. As shown schematically in fig.1 (c), the rotating arc acts just like a firewall near the exit nozzle of the

NASA/CP--2001-210948 285 August 6, 2001
anode, whichwillrapidlyheatupthecoldgassheathandcauseit toexpandandbeexpelledthroughthenozzle.

Besidestherotatingarcalsobehavesaastirrer,whichfurtherguarantees theuniformityoftheplasmajet.

(:old carbon

\

Arc

Anode

The

-Plasma jet

(a) (b) (c)

Figure 1 (a) Large orifice arc jet plasma torch with rotating arc root operating at gas recycling mode, (b) photograph

of arc plasma (a dark filter was used), the substrate area was dPll0mm, (c) schematic showing the principle of arc

root rotating

2.2 The computer model

In order to set up the computer model, the discharge tunnel of the plasma torch shown in fig. 1 is simplified as

composed of two columnar spaces (see fig.2). The cathode is located on the top of the discharge tunnel, whist the

_r c2th,,a_ substrate is placed a certain distance below the anode exit nozzle. Only newly feed

¢" _ i_i24ctf 4 Ar is allowed to pass though the cathode, whilst the newly feed Hz+CH4 are mixed

with the recycled gases before entering the tunnel. The principles of gas recycling

can be found elsewhere [2-4], and is similar to that used by Partlow et al [5'6].

1 Principles of the theories of fluid dynamics, plasma physics, and plasma chemistry

i have been applied. The detailed descriptions can be found elsewhere [71. In order to

Anode i

l calculate the heat generated due to arc discharge, it is assumed that the discharge is

i

1 constrained inside the arc column, and is uniform within its whole volume. The arc

i

i

11................ column is sustained between the tip of the tungsten cathode and the copper anode. By

: i t '_ proper manipulation of the magnetic field and the fluid dynamics, it is possible to locate

/./ /; "',. the arc roots exactly at the position near the exit nozzle of the anode, and rotate rapidly

I Me sub,_trate [ (see fig.1 (c)). The diameter of the arc column can be given as :

] ] Qi = ki.Q0.1/(r2+ra 2) (1)

Figure 2 Simplified Where Qi denotes the heat generated within the ith element of the arc column, Q0 is the

geometry of the total heat could be generated by the input power, ki is a coefficient, whilst r and ra are the

discharge tunnel radius of the anode nozzle and the arc column respectively. At the position of the

"firewall" shown in fig.l, ra is assumed to be equal to r.

Plasma chemistry was calculated by assuming that the stead state localized equilibrium in the scale of the

digitized cell size for numerical calculations could be easily established, although the plasma was thermodynamically

non-equilibrium in nature. The Thermo-Calc program together with a powerful database of thermodynamics was

used for calculation of the plasma chemistry.

Boundary conditions were chosen to match the actual experimental conditions as much as possible. They are as

follows: At all the water-cooled surfaces of the discharge tunnel: temperature was 323K, radial and axial velocity

were zero. At the Mo substrate surface: temperature was 1123K, radial and axial velocity were zero. Initial axial

velocity of the total inlet gases can be calculated from the total gas mass flow rate and the cross section near the

entrance of the arc discharge tunnel. The initial radial velocity was zero. The initial temperature of the inlet gas was

taken as 323K. The diameter of the nozzle exit and the water-cooled Mo substrate were qb60mm. Input power for this

NASA/CP--2001-210948 286 August 6, 2001
sizeoftorchexitandsubstratewaslessthan20kW.Aratioofgasrecyclingof90%wasassumedforallcalculations,

unlessotherwisetated. Boundarylayereffectswerenotakenintoaccount forthereasonsofsimplicity.

3. RESULTS AND DISCUSSION

3.1 Flow and temperature field

Fig.3 shows the calculated flow and temperature field for the particular type of plasma torch shown in fig.1 with

the following operating parameters: Argon to hydrogen ratio: 0.8:1.0; Inlet gas mass flow rate: Ar: 4.4 SLM, H2:5.5

SLM, CH4: 0.1SLM; Total gas mass flow rate: 100 SLM (inlet gases + recycled gases); Distance between the anode

nozzle and the substrate surface: 2 cm; Chamber pressure: 3kPa; Input power: 18kW.

g

/t

/

_ o'88_

(a) flow field (b) temperature field

Figure 3 Calculated flow and temperature field for plasma torch operating at 18kW and 3kPa. The unit for distance is

in cm, and for temperature is in Kelvin, the shaded area is the Mo substrate, the anode nozzle exit is located at

y-2cm, the cathode tip is placed at the position of x-0.00, y-16.00cm.

Mo substrate is placed at the position of y - 0.00, which is 2 cm away from the anode nozzle (y - 2.00). The

cathode is located at the top (y - 16.00). It can be seen from fig.3(a) that there is no turbulent flow and the flow over

the substrate surface is smooth. Small vortexes shown in fig.3(a) is due to over simplification, which do not happen

in the real case. However, very big temperature gradient can be observed near the cathode region, whilst the

temperature distribution is quite uniform over the _60mm Mo substrate, which may due to the rotating arc. The exit

2.00 _ I I I I I I I I I I I

100

0 50

0.00_

3.00 2.50 200 1.50 100 050 0.00 050 100 1.50 2.00 2.50 300 -3.(ZO-2.'50-2.'(8-1.'50-1.'00-0.'50 (R_)O 050 1.00 1.50 2_70 250 3.0[

(a) with arc roots rotating (b) without arc rotating

Figure 4 Temperature field over the Mo substrate surface

NASA/CP--2001-210948 287 August 6, 2001
temperatureof 3500-5000Kisingoodagreement asestimatedfromopticalemissiondata.Fig.4(a)further

showstheuniformityofthetemperaturefieldoverthesubstratesurfaceinthecaseof arcrootsrotating,whilst

thetemperaturefield withoutarcrootsrotationis obviouslynon-uniform(seefig.4 (b).Thisclearly

demonstratesheadvantageof arcrootsrotating,whichis of vitalimportanceforlargeareahighquality

diamondfilmdeposition. Fromfig.4(a),it canbeseenthathereisathinlayeroverthesubstratesurfaceacross

whichthereisalargebutradiallyuniformtemperaturegradient. Thethicknessofthislayerislessthan2ram,

whichisreferredas"temperature boundarylayer",andmayberelatedtotheboundarylayerinfluiddynamics.

3.2 Plasma chemistry

Fig.5 shows the calculated gas phase composition over the Mo substrate surface. The most abundant chemical

species are H, C, C2H2, C2H, C2, CH, whilst the concentration for CH4 and CH3 are negligibly small, and therefor not

shown. It can be seen from fig.5(a) that the radial distribution of those most abundant chemical species over the

substrate surface are quite uniform. For comparison, fig.5(b) shows the calculated radial distribution for the case

without arc roots rotating, where it can be seen that the distribution is rather non-uniform. Fig.5(c) and (d) shows

¢60mm diamond films on Mo substrate deposited with and without arc roots rotating at identical conditions as for

the numerical calculations, which is in direct coincidence with the radial distribution of chemical species (fig.5(a)

and (b)) and the distribution of gas temperatures over the substrate surface (fig.4). Thus we have demonstrated again

the advantage of arc roots rotating. Calculated results shown in fig.5 have also been partly verified by optical

emission spectroscopy, where only C2 and CH radicals were shown [sl

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Figure 5 Radial distribution of calculated gas phase chemical species over the substrate surface and the consequent

uniformity in diamond film deposition. Concentrations for CH4 and CH3 are negligibly small, therefore are not

shown. (a) and (c): with arc roots rotating; (b) and (d): without arc roots rotating

3.3 Prediction of _rowth rate and film quality and their uniformity for large area diamond film deposition

The absence of CH3 and CH4 radicals from our calculated results means that the CH3 mechanisms for diamond

growth cannot be applied. Controversial points of views on which radical should be responsible for CVD diamond

NASA/CP--2001-210948 288 August 6, 2001
growthcouldbefoundintheliterature. Initialarguments werefocusedonCH3andC2H2, andmainlyforHFCVD

andMWCVD, andexperimental evidencefavorableforbothCH3andC2H2werereported[9,i01. LohandCappeli [HI,

andGoodwin[:21suggestedhatCH3wasresponsible fordiamondgrowthindcarcplasmajet.However, ithadbeen

reportedthat,forhighspeedhightemperaturesystem, CandC2mightbecomethemainconstituents inthegas

phase, andsowouldresultindiamondgrowth,and[CH2]mayalsobeaneffectiveprecursor, whichcouldaccount

forasmuchas17%of observeddiamondgrowth[:243]. MorerecentlyGruenetal.reportedthatnanocrystalline

[14] diamond films could be grown from pure Argon and CU4, where C2 was the effective precursor

Based on our calculations it is reasonable to argue that C and C2 may be the main precursors for diamond growth.

However, C2H2, CzH, and CH2 may also be of importance.

Goodwin et al. [:sl had related the growth rate and the quality of CVD diamond films with the gas phase

concentration of atomic hydrogen and methyl radicals as:

G=9xl011f * [CH3IH] (2)

5×1o

c (3)

[def] _ [H]2

where G is the growth rate, f * is a constant, [dej] represents the density of defects, and therefore can be regarded as a

measure of the quality of diamond films. Apparently atomic hydrogen is the key factor for diamond growth. In our

numerical calculations it was shown the concentration of atomic hydrogen was pretty high (see fig.5(a)), therefore

high growth rate and high quality diamond deposition could be expected. Since [CH3] is negligibly small, so

equation (2) can be modified to fit the C2 mechanism as:

Z[CI[ H] (4)

G =k:" k2 + _q]

where kl and k2 are constants, which can be obtained using experimentally measured G and Z[C]i. Thus equation (4)

and (3) can be used for prediction of the growth rate and the quality as well as their radial uniformity for large area

diamond film deposition by DC arc plasma jet.

We have done a number of "virtue experiments" on the effects of process parameters on large area diamond

deposition. Very interesting results have been obtained and compared to experimental observations. For example,

fig.6 shows the effect of the distance from substrate surface to the exit of the anode nozzle. Where it can be seen that

the growth and quality as well as their radial uniformity all decreases with increasing substrate to anode nozzle

distance. Obviously this is in contradictory to that reported in literature for arcjet diamond film deposition [:61

However, the predicted results coincide quite well with the experimental observations. The reason for this

discrepancy is the design feature of arc root rotating which was also taken into account in our computer model.

20-

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Radial chstance (_) Radial distance (cra)

(a) growth rate (b) film quality

Figure 6 Effect of the distance of substrate to anode nozzle on growth rate and quality and their radial uniformity

over a qb60mm deposition area

NASA/CP--2001-210948 289 August 6, 2001
Becausetherotatingarcislocatedneartheanodenozzle,andthearcjetis"soft"ascomparedtomostoftheother

arcjets, whicharesometimes supersonic, thereforethetemperature fieldtendtodecreaseandtobemorenon-uniform

withincreasingdistancefromtheanodenozzle.

4. SUMMARY

A computer model has been established for a special type of dc arc plasma torch with magnetic field and fluid

dynamics control and arc root rotating operating at gas recycling mode. Flow and temperature field as well as plasma

chemistry has been simulated. The advantages of arc root rotating have been demonstrated. Growth rate and quality

as well as their uniformity for large area diamond deposition can be predicted. It has been shown that it is possible to

use the model for simulation of the effects for a series of important process parameters on large area diamond

deposition.

ACKNOWLEDGEMENT

The authors wish to express their thanks to the NAMCC (National Advanced Materials Committee of China) for

financial support under the contract of 863-715-Z38-03.

REFERENCES:

1. Busch, J. V. and Dismukes, J. P., Diamond and Related Materials, 3 (1994) 295

2. Lu, F.X., Zhong, G.F., Li, G.H. et al., Proc. Inter. Diamond Symposium Seoul, (1996) 115

3. Lu, F.X., Zhong, G.F., Sun, J.G. et al., Diamond and related Materials, 7/6 (1998) 737

4. Lu, F.X., Tang, W. Z., et al., Diamond and Related Materials, 9 (9 10), (2000) 1655

5. Partlow, W.D., Schreurs, J., Young, R.M., Martorell, I., Dighe, S.V., Swartzbeck, G., and Burton, J., Application

of Diamond Film and Related Materials, Feldman A., et al. eds., NIST Special Publications 885, USA, 1995

(519)

6. Martorell, I.A., Partlow, W.D., Young, R.M., Schreurs, J.J., Saunders, H.E., Diamond and Related Materials 8

(1999) 2926

7. Huang, T.B., Ph.D thesis, University of Science and Technology Beijing, March, 2001

8. Zhong, G.F., Ph.D thesis, University of Science and Technology Beijing, March, 1998

9. Harris, S.J. and Martin, L.R., J. Mater. Res., 5, 1990, 2313

10. Johnson, C.E, Weiner, W.A. and Cerio, F.M., J. Mater. Res., 7, 1992, 1427

11. Loh, M.H. and Cappelli, M.A, Appl. Phys.Lett., 70, 1997, 1052

12. Goodwin, D.G., Appl. Phys.Lett., 59, 1991, 277

13. Yu, B.W. and Girshick, S.L., J.Appl. Phys., 75, 1994, 3914

14. D. A. Homer, L. A. Curtiss, and D. M. Gruen, Chemical Physics Letters 233,243 (1995)

15. D.G.Goodwin, J.Appl.Phys. 74, 6888 (1993)

16. Cappelli, M.A, in Hand Book of Industrial Diamonds and Diamond Films, Chapter 23, Prelas, M.A.,

Popovici, G., and Biglow, L.K. eds., Marcel Dekker Inc., 1997, 865

NASA/CP--2001-210948 290 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Gas-phase concentrations and Temperature measurements

of C2 in a diamond depositing DC-arcjet CVD system

J.A. Smith', K.N. Rosser, P.W. May, H. Yagi t, A.J. Orr-Ewing and M.N.R. Ashfold

School of Chemistry, University of Bristol, Bristol, BS8 1TS, U.K.

lSchool of Engineering, University of Ehime, Matsuyama, Japan

Tel: (117)-9288312; Fax: (117)-9251295; e-mail: mike.ashfold@bris.ac.uk

Keywords: DC-arcjet CVD, ground state C2 concentration and temperature measurements, Cavity

Ring-down spectroscopy

Abstract

DC-arcjets operating with a hydrocarbon / H2 / Ar gas mixture enable chemical vapour deposition (CVD) of

high quality diamond films at growth rates unobtainable using the more traditional hot filament or

microwave reactors)

Previous gas-phase studies on microwave plasma enhanced diamond deposition have concluded that the

dicarbon (C2) species concentration may be inversely proportional to the diamond film quality) In

contrast, optical emission studies of a DC-arcjet diamond CVD reactor operating on Ar / H2 / CH4 gas

mixtures have found that high C2 concentration correlates with high quality diamond growth. 3'4

Our current studies axe focussing on first concentration measurements of ground state C2 radicals and their

rotational temperature, within a DC-arcjet plume using the ultra-sensitive laser based absorption technique

of Cavity Ring-down spectroscopy (CRDS). This type of spectroscopy has been used previously in gasphase studies of CH3 and CH radicals in Hot-filament 5 and Oxy-acetylene reactors 6 respectively, but has

not been applied to spatial profiling species concentrations and temperatures within a plasma jet.

References :

1. M.A. Cappelli and T.G. Owano, in Low-Pressure Synthetic Diamond, (13. Dischler, C.Wild,

(Eds.)),Springer, (1998), pp. 59-84.

2. H. Yamaguchi, M. Ishii, K. Uematsu and S. Morimoto, Jpn. J. AppL Phys. 35 (1996) 2306.

3. W. Zhu, A. Inspektor, A.R. BadzJan, T. McKenna and R. Messier, J. Appl. Phys., 68 (1990) 1489.

4. J.A. Smith, K.N. Rosser,H. Yagi, M.I. Wallace, P.W. May and M.N.R. Astffold, Diam. Relat. Mater,

10 (2001), in press, as part of Diamond 2000, Oporto, Portugal, Sept 2000.

5. E.H. Wahl, T.G. Owano, C.H. Kruger, Y. Ma, P. Zalicki and R.N. Zare, DiamondRelat. Mater. 6

(1997) 476.

6. R.L. Stolk and J.J. ter Meulen, DiamondRelat. Mater. 8 (1999) 1251.

NASA/CP--2001-210948 291 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

NUMERICAL SIMULATION AND REALISATION

OF NOVEL MICROWAVE PLASMA REACTORS FOR DIAMOND CVD

C. Wild, E. Pleuler, W. Miiller-Sebert, P. Koidi

Fmunhofer-Institut far Angewandte Festk_rperphysik, Tullastrasse 72, D-79108 Freiburg, Germany

E-mail: _v_i.l._l ._._.._g_._, Homepage: www.cvd-diamond.com

ABSTRACT

The developmem of efficient microwave plasma reactors for diamond CVD is an important issue for the

production of large area, high quality diamond disks. In this context numerical simulations using FEM algorithms

that allow a self-consistent determination of the electric field and plasma density distribution in microwave plasma

reactors have found considerable interest. The talk will describe the numerical concepts applied and the validation of

the simulation results.

Two types of microwave plasma reactors have been realized. One exhibits an ellipsoidal cavity with two focal

points. By coupling microwave energy into one focal point, high electric fields can be generated in the second one.

This in turn can be exploited for the excitation of intense plasmas. This ellipsoidal reactor is now used as a

"workhorse" for the routine production of diamond disks.

As an alternative a second reactor concept has been studied. In this reactor the microw_.ve enters the plasma

chamber through a ring-shaped circumferential window. The microwave is guided to the window by radially

expanding a coaxial waveguide. Simulation results, measurements performed with a miniature microwave sensor

and preliminary performance tests will be presented for this CAP-reactor (_ircumferential antenna l/lasma reactor).

Furthermore, the properties of CVD diamond disks and their applications in optics (radiation windows and

lenses), thermal management (heat spreaders) and mechanics (wear resistant components) will be described.

Keywords: microwave plasma reactor, numerical simulation, diamond disks

NASA/CP--2001-210948 292 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ELABORATION AND MECHANICAL PROPERTIES OF DIAMOND COATINGS

OBTAINED BY FLAME PROCESS

D. Paulmier, M. Schmitt

LPMM-ERMES, CNRS-UMR 7554

2, avenue de la for_t de Haye 54516 Vandoeuvre Cedex France

Daniel.Paulmier@ensem.inpl-nancy.fr

Phone : 00-33-(0)3-83-59-56-16 Fax : 00-33-(0)3-83-59-56-17

M. Mermoux

LEPMI-ENSEEG Domaine universitaire, BP 75 38402 Saint Martin d'H6res Cedex France

ABSTRACT

The flame process offers two advantages : it is, on the one hand, relatively easy to realise, and on the other

hand, gives a crystals growth speed more than hundred times faster than the CVD processes one.

This method presents however problems which are still not yet solved. It leads to "local" coatings (lower than

1 cm 2) : the covering of large surfaces consequently needs a torch motion across the sample surface, that can induce

a crystals graphitization. The use of an argon shield or a combustion chamber can avoid that.

The nucleation density, the crystals size and orientation and the growth direction are hardly controlled as they

greatly depend on many factors (gases speed, C2H2/O2 ratio, combustion instability); these parameters variations

lead to flame modifications inducing irregularities in the coatings.

In spite of these drawbacks, the flame process remains an attractive method for covering cutting tools. The

tribological behaviour of these coatings is different from the one of those obtained by CVD as the dangling bonds

are here mostly saturated by oxygen.

This paper presents a review dealing with the technological and scientific knowledge as well as the evolution of

this process during the ten past years.

Keywords : diamond coatings ; flame process ; mechanical properties ; cutting tools

INTRODUCTION

Among the different elaboration methods of diamond coatings, subject in vogue during the past ten years, the

flame process is very attractive as it, in theory, does not need expensive apparatus. But it is observed that since its

development [1], even if the elaborating devices were improved, particularly with the combustion chamber

principle, the obtained results do not keep, nowadays, their promise. Many studies were however realised and allow

a better understanding of the process itself as well as the deposition of coatings showing high optical and mechanical

qualities, but only on relatively small surfaces (in comparison with the low pressure CVD process).

We analyse hereafter the various parameters influencing the nature of the coatings and the process evolution

since its creation.

THE FLAME PROCESS

The CVD processes present a major drawback as they require a vacuum chamber, which is a relatively

expensive equipment, as well as components capable of resisting high temperatures.

To remedy these financial and material problems, Hirose and Kondoh in 1988 devised the flame process, using

the fact that a flame is in reality a plasma in air and can thus play the role of a reaction chamber giving thus the heat

necessary to the creation of radicals involved in the formation and growth of diamonds [1,2].

Experimental device

This process, based on the dissociation of the carbon originating from the acetylene, during its combustion in

oxygen, is comparatively easy and not onerous as it needs :

NASA/CP--2001-210948 293 August 6, 2001
-anoxy-acetylene torchequippedwithamovingsystem,

-twoflowmeters tocontrol theoxygenandcombustible flows,

-acoppersubstrateholderwithacoolingdevice,

-apyrometer checkingthesubstratesurfacetmperature (figurela).

t } +'+"" pyrometer ...... 1_'..,. _...I

dif_asionflame ] I ,'v'_qeatvlana £_nth_r

"' A & M" .... -' ..............

--l___Z¢__;-____'_'i' substrate '_"%:ff " "'_"I" ;"" '

............ t" ";:_ ..... dfffasi,'n _lame ;?) V

substrate holder with a cooling system .:

a b

Figure 1. Flame process (a) experimental device ; (b) flame structure

This device, outwardly not elaborate, requires nevertheless the mastery of many parameters to succeed in the

deposition of diamond coatings satisfying the imposed conditions of purity, crystallinity and morphology.

Principle of deposition

It is directly linked to the flame configuration ; three zones can be defined (figure lb) :

o:o zone 1 : the inner cone

This bright and sharp area is the hottest of the flame : temperatures around 3160°C were measured there [3] ; its high

temperature near its base ionises the hydrocarbon by thermal plasma, generating thus chemical vapour species useful

to the diamonds deposition. Oxygen and hydrocarbon burn there in stoechiometric conditions :

C2H2 + 02 -_ 2 CO + H2 + 107,2 kcal/mol (1)

o:o zone 2 : the acetylene feather

This area is formed when using an excess of acetylene and is made of carbon monoxide, hydrogen and radicals

originating from the hydrocarbon dissociation : C °, C2°, CH °, C2H °.

The substrate is set here so that the previously generated atomic carbon induces the formation of carbonaceous

structures like diamonds.

°:° zone 3 : the diffitsion flame

Hydrogen and carbon monoxide which were not entirely burnt during the primary combustion in zone 1, react with

the atmosphere (oxidation process) leading consequently to the formation of carbon dioxide and water vapour :

2 CO + H2 + 3/2 O2 _ 2 CO2 + H20 + 203,7 kcal/mol (2)

A peripheral zone (shield) is sometimes added to the three precedents when argon is used with the oxygenhydrocarbon mixing : this gas which has no effect on the previous reactions is intended to reduce the oxidation

phenomena occurring on the jet periphery by focusing the flame.

The deposition method of diamond coatings, beyond its low cost, differs from the CVD process essentially on

four points :

4, all the reactions occur in ambient air : the oxygen of the surrounding atmosphere could diffuse through the flame

periphery to the inner cone ; but it seems [4] that the attack abilities of the graphite by atomic oxygen (and OH

radicals) are two to three times higher than those of atomic hydrogen : a high concentration of atomic oxygen (and

OH radicals) at the flame periphery would contribute to the development of diamonds showing high quality [5].

o:o the dangling bonds of the diamond crystals are not necessarily saturated by hydrogen [6] as it is the case for the

coatings obtained by CVD.

o:o the crystals growth speed is high, 100 to 200 gm/h, reaching sometimes 600 to 900 gm/h but including in that

case graphitic impurities.

NASA/CP--2001-210948 294 August 6, 2001
o:othethermal energyneedful tothediamondssynthesisdirectlycomesfromthecombustion energyoftheprecursor

gas.

Theflameprocessshowsnevertheless somedrawbacks, moreparticularlyif thesurfacewhichhadtobecoated

islargeandthesubstratetotallyheldintheflame:

o:ointhisconfiguration, thesubstratesidesareoverheated ;whirlpoolsintheatmosphere arethengeneratedthere,

inducingdeteriorationsfthecoatingsqualityandhomogeneity.

o:oturbulences canappearintheflamewhenthegasescrossthesubstrate, thatleadstothepresenceof the

surrounding atmosphere onthesubstrate.

Infact,oneofthemainproblemoftheflameprocessistheobtainingofaverystableflame.It iswellknown

thatacousticwaves, inducinglocalandtemporal variationsofthepressure, takeplaceonthenozzlesides;these

haveeffectsontheflamecircumference, modifyingthusthekineticsandcombustion. Theseinstabilitiesdisturbthe

diamondcrystalsgrowthandleadtotheformationofsuccessive stpsonthecrystallinefacets;theycanhowever be

reducedbymachiningozzleshavingaconvergent-divergent sightcombinedwithanefficient electrolyticpolishing.

WhenthelowpressureCVDprocesses assistedbyplasmaandothermethodsleadtocoatingswhomcrystals

havearegularsize,thoseobtainedbyflamepresent variouscrystalssizeasthetemperature, thechemical species

distributioninthegaseousphase, thegasspeedandtheincident anglevarywiththedistancetotheflameaxis;these

coatingsareconsequentlyonhomogeneous andtheirmechanical propertieschangeable. Inordertoattenuatethese

phenomena, andtocoatlargesurfaces, alinearscanning[7]orarotatingmovement ofthesubstrate[8]areused.

Nevertheless, manystudieshavealreadybeencarriedoutconcerningtheprocessimprovingandthedeposition

conditionsoptimisation.

THE FLAME PROCESS IMPROVING ; INFLUENCE OF VARIOUS PARAMETERS

Flame process with combustion chamber [9,10]

In this device devised by Murakawa, the substrate and the nozzles are located in a chamber (figure 2) ; the

pressure inside this combustion chamber is set constant at a value slightly higher than the ambient one.

::;:..+ secondary

:::::::+,,I_--- flame

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view window

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combustton ::::::::::::::::::::::: +'_:....::::::::::.

chamber _:i:i:i:i.':_-:-:_:-:-'_:.:.:.:.:.:.:.:.:.

_:-:+:+:+:-::_:+.'._..'+:1:.-.-.-.-.-.-.-.-.:l substrate holder with

[:.i:.i:.i:.iiiiiiiiiiiiiiii: i_i:iiiiiiiiiiiii :__i-._.-. .}... cooling system

substrate

Figure 2. Flame process with combustion chamber

During the deposition with the classical flame process, both the carbon monoxide and hydrogen which are not

burnt during the primary combustion (1) are used in a secondary combustion ; but, when using the combustion

chamber, the oxygen needed for the secondary combustion is not introduced, and, as the initial oxygen has already

been used for the first combustion, the atmospheric oxygen is indispensable for continuing the reactions.

Consequently, in the case of flame process with chamber, the secondary combustion occurs out of the reactor, when

NASA/CP--2001-210948 295 August 6, 2001
thegasesproducedduringthefirstcombustionwouldhavebeenevacuatedbyaflarestackforthispurpose. This

depositionprocessinducesanidealchemical reactioninanatmosphere madeofcarbonmonoxideandhydrogen.

Thefirstcoatingsobtainedbythisprocessshowedatoohighapparent roughness whichpreventsthemtobe

usedinanytribologicalpplications ; thisproblemwassolvedbyworkingoutatwotimedeposition. Beforethe

beginningofthedepositionitself,thesubstratesurfaceissubmittedtoa pre-treatment : based on the fact that the

diamonds nucleation is not easy on a rough surface (the preferential nucleation sites are rather the convex slopes of

the asperities), diamond powder (5 nm in diameter) is distributed on the substrate previously polished by electrolytic

process ; the small surface cavities, thus partially filled up by these diamonds accumulations, become favourable

sites to the nucleation.

The process modification consists in the addition of a step during the crystals growth, step in which the

C2H2/O2 ratio is modified : it was shown that, for the classical flame process, an acetylene excess or a too low

substrate temperature induces a decrease of the non-diamond species attacks, that leads to the formation of a thin

DLC layer on the top of the pure diamond coatings. The growth by steps uses this thin DLC layer as a new

nucleation site for diamonds : it stops the formation of big crystals, diminishing consequently the final roughness of

the diamond coatings. The process begins first by a 25 min deposition with a C2H2/O2 ratio of 1.13 ; this ratio is then

increased to 1.43 during 30 sec before coming back to its initial value till the end of the test (40 min). The so

obtained coatings are made of small crystals, their roughness is similar to the one of diamond films deposited by hot

filament CVD, and the growth speed is still 30 gm/h.

In order to improve these coatings, Murakawa wanted to increase their thickness ; the last experimental device

allows the deposition of coatings having 300 gm thickness but in the form of non uniform layers. The setup was

then modified : the substrate can be rotated, and the nozzles, tilted, can be moved. The coatings uniformity was

consequently improved, but for this thickness (300gm) the diamonds quality was affected (whereas it is excellent for

30 to 50 gm coatings) showing thus that the modifications of the flame process with combustion chamber had to be

pursued.

Torch modifications

As the torch is the basis of the flame process, every modification of it will inevitably have consequences on the

coatings sight :

o_oan increase of the torch diameter going with an enlargement of the gases speed leads to a turbulent flame ; but, if

the proportions between the sizes of the inner cone and the acetylene feather remain constant, diamonds of very high

quality are obtained, even if the growth speed is lower [11,12].

o,*oa tilted torch induces variations of the coatings morphology [13] : with an incidence angle (angle between the

substrate surface and the torch axis) of 90 °, coatings are made of {111} facets, whereas at 60 °, {100} facets are

rather prevailing ; if this angle is decreased to 45 °, dense and homogeneous coatings are constituted of a mixing of

{ 111 } and { 100} facets. High purity polyhedral crystals are then obtained when the incidence angle is 30 °.

However, the tilted torch can create problems [5] : the so deposited coatings show delaminations and breaks,

which are essentially due to the bad adhesion induced by the decrease of the gaseous flows speed perpendicularly to

the substrate (the smaller the momentum of deposition particles running into the substrate, the lower the probability

of these particles diffusing into the substrate) ; the so observed growth speeds are moreover lower than those

obtained with the classical flame process.

The use of multinozzles [14] leads to uniform coatings, having a good quality, on 12 mm diameter surfaces ;

nevertheless, the presence of small non-diamond species indicates that this device had to be improved. Recently,

Hollman obtained diamond coatings of high quality with a 9 nozzles apparatus [15].

The substrate position in the acetylene feather

Diamonds can theoretically be obtained in all the acetylene feather ; however, if the distance (Ps) between the

extremity of the inner cone and the substrate increases, the coatings uniformity as well as the diamonds quality vary

[8]. On the contrary, if this distance is decreased, the growth speed is enlarged and the diamond quality improved ;

the best distance seems to be as close to the inner cone as possible [16].

The substrate preparation

Most of the diamond coatings are used to coat cutting tools made of tungsten carbide containing cobalt ; if the

Co concentration of the substrate surface is higher than a few percents, amorphous carbon appears at the diamondsubstrate interface during the deposition : this affects not only the substrate adhesion, but also the diamond growth

and cristallinity [17]. It is then essential to impoverish the surface in Co by using selective attacks with acid

NASA/CP--2001-210948 296 August 6, 2001
solutionsorbyadirectattackduringthediamonddeposition[18,19].Othersurfacetreatments foralloyedcarbide

(6%Co)weretested, andtheirinfluences onthecoatingscompared[20]:

o:.Oxidisingflamepre-treatment (C2H2/O2 < 1):thatinducesanoxidisationofthetungstencarbideleadingtoa

decreaseofthediamondparticlesdensityandconsequently toabetteradhesionaswellasarelativelysmooth

surfacestateattheendofthedeposition; however, if theCoisnotpreviouslypulledout,thecontactwiththe

oxidisingflameallowstheformationofCo2PandCoOwhicharefastenedtothetungstencarbidegrains,creating

thusaprotectivelayerpreventingthenthefavourableactionoftheflame.

o:.theultrasonicscratchingwithdiamondandironparticlesleadstoanexcessivesurfaceroughness, inducinga

relativelyowdiamond-substrate adhesion.

o:.theseedinggivesthehighest diamondsensity, thatconsequently improvestheadhesionadreducesthecoatings

roughness.

o:.thebestresults(lowestroughness, highestadhesion) areobtainedwhenthesubstrateisfirstsubmittedtoan

oxidisingflame,thenattackedbyacidsolution, andfinallyseeded.

Othertypesofsubstrates,peciallysilicium,werepre-treatedtoenlargethenucleationdensitybyoxidisingflame

[21]orbyapplyingabias[22].

Biasvoltageapplicationbetweenthenozzleandthesubstrate

Anegativebiasappliedtothesubstrateduringthedepositionleadstotheappearance ofanelectricfieldinthe

plasma, thathasaninfluenceonthecrystalsgrowth:largemonocrystalline diamondscanthusbeobtained[23].

Thewaythecoatingsgrowis alsomodified[24]: acolumnargrowthisreplacedbyasmallgrainssize

development whenapplyingabiasvoltage, then,whenathresholdisreached, nanocrystalline diamondsareformed.

Themechanisms explainingthesevariationsarenotyetwellknown:

o:.ontheonehand,thebiasapplicationmodifiesthedistributionofthereactivespeciesdistributionintheflame,and

moreparticularlyonthesubstratesurface, asindicatedbythebrightercolouroftheflameinthepresenceofabias,

o:.ontheotherhand,it isalsopossiblethathebiasinducesaionsbombardment ofthesubstratesurface, thatcan

introducedefaultswhichcanplaytheroleofnewnucleationsites;thiscouldexplaintheobservedsecondary

nucleation.

Oneobviousadvantage ofthebiasapplicationistheformationofsmoothdiamondcoatingsinonestep(nopre

orposttreatment) withalwaysahighgrowthrate(50gm/h)similartotheoneoftheclassical flameprocess.

Thedifferent kindofhydrocarbons

Acetyleneisthemostwidespread hydrocarbons usedintheflameprocess;themainreasonofthischoiceliesin

thefactthatontheonehandtheflamespeedofthiscombustibleisaround150cm/sinambient air(whereasit is

only75cm/sforethyleneand40cm/sformethane[25]),andontheotherhandit isneedful thisvalueishightoget

favourablegrowthconditionsearthesubstrate. Acetyleneisnevertheless arelativelyexpensivegas:thatiswhy

ethyleneandmethaneweretestedandcomparedtoacetylene[26]:

o:.forthesethreehydrocarbons, theobtaineddiamondsshowagoodquality,andhaveareasonablegrowthrate

(lgm/h),

O the range of the oxygen/hydrocarbon ratio values allowing the deposition of diamonds is larger for ethylene than

for acetylene,

O the first tests realised with methane seems to indicate it could give results similar to those of acetylene and

ethylene.

Other carbon precursors were also tested :

O MAPP, commercial mixing of methyl-acetylene and propadiene in LPG gas (propylene) at 70-180 torr, allows the

deposition of continuous diamond coatings [27],

O propylene alone, at 180 torr, leads to well crystallised diamonds [28].

Morphology modifications

The morphology is not the same in all the points of a diamond coating [14] :

o:o 0-3.5 mm from the centre, the crystals show {111 } facets,

o:. on and after 4 mm (transition zone), this is rather {100} facets on the top of { 111 } pyramids,

o:. at the periphery (on and after 5.5 mm), there are essentially {100} surfaces.

These observations, realised on diamond coatings obtained by flame process with multinozzles, are exactly

opposed to what these same authors have noticed in the case of a single nozzle [29]. This kind of opposition is

relatively widespread in the literature : for Kobashi [30], an enlargement of the methane concentration in hydrogen

modifies the crystals orientation from {111 } to {100}, and if the concentration continues to increase, it induces the

NASA/CP--2001-210948 297 August 6, 2001
formationofDLC; Spitsyn[31]obtained{111} orienteddiamondcoatingswithhighmethaneconcentration, and

{100}coatingsintheoppositecase.

Thediamonddepositionbyflameprocessoccursin conditionswhicharefarfromthethermodynamic

equilibrium, thisimpliesthatthesmallest modificationintheobtainingmethodhasa significant effectonthe

coating, explainingthusthedisparityofthepresented results.

Thesubstratetemperature andtheO2/C2H2 ratio

The substrate temperature (Ts) plays an important part in the growth rate, diamond nucleation and crystals

quality, but also in the adherence to the coating ; according to the research workers, the ideal range for the substrate

temperature varies between 500 and 1300°C [32] or 700 and 900°C [8] (depending on the experimental conditions).

The ratio R - O2/C2H2 had to be in a relatively restricted range (0.85-0.95) [15] to obtain diamonds having good

quality, without contamination by non diamond phases ; it also influences the crystals orientation [16,33].

Polycrystalline coatings with a minimum of amorphous carbon and graphite can be obtained with the following

values of R and Ts : R> 0.95, Ts between 950 and 1050°C [34].

The 02/C2H2 ratio generally remains constant during the all deposition ; however, a variation of this ratio as the

test goes on has repercussions on the coatings characteristics. Ravi [35,36] has shown that a decrease of this ratio

induces the previous formation of a DLC layer which has a marked effect on the nucleation density of the diamond

crystals growing on it, as well as on their morphology. In the same way, by diminishing this ratio from 0.9 to 0.7

after 5 rain deposition, and during 30 sec before coming back to 0.9, the nucleation density and the adherence

obviously increase [37] ; the sensitivity to the substrate roughness is nevertheless decreased compared with the one

observed during a deposition when 02/C2H2 remains constant, and the obtained coatings are rather mixtures in

which the sp 3 amount is lower than the one observed in the classical process.

Murakawa realised in 1997 diamond coatings of high quality with a two steps deposition using a flame process

with combustion chamber [9].

Large size coatings

Many authors succeeded in the synthesis of large size crystals [38,39] ; the coatings deposition on large

substrate is fundamental for numerous tribological applications.

Murakawa [40] obtained two types of coatings (ball-like and well-crystallised) of 10x50 mm 2 by moving the

torch across the substrate. This device was then improved [41] : the nozzle is still moving, but a flame cover is

located around the nozzle, at the substrate surface level, in order to prevent the air introduction in the deposition

zone and consequently to avoid the secondary combustion; this stops then the oxide film formation. The

nondiamond phases are moreover removed by recovering during the successive crossings of the torch ; the so

obtained coatings reach size of 50x50 mm 2.

Recently [8], an experimental device allowing the obtaining of large coatings on various substrate was realised:

the torch is tilted and the flame surrounded by an argon shield, the substrate is rotated and cooled by a heat

exchanger. Those coatings effectively have a big size (1 cm 2 on a Si3N4 triangle), but their growth rate is lower than

the one observed for the classical apparatus ; because of its rotation, the substrate is out of the flame deposition zone

during the two third of the test, that induces "growth/attack" cycles on the diamonds when these latest are in contact

with the oxygen of the ambient air.

Coating-substrate adhesion

Various processes were used to improve the diamond coating adhesion on a substrate, but the most widespread

method is the scratch test [42,43] ; the substrate preparation (Co impoverishment, ionic bombardment, interlayer,

bias voltage application) is also of great importance to get a good adhesion and consequently a longer lifetime of the

coating ; the substrate temperature during the deposition plays an important part too : the formation of a diamond

coating on a nitride chromium substrate at around 900°C leads to a continuous film, but this temperature had to be

decreased to 250-400°C (to avoid high residual stress) to improve the adhesion [44]. In this case, scotch tape tests

realised on these coatings show that the highest adhesion is obtained at the centre of the coating.

Modelling of the flame structure

A numerical approach of the flame structure was realised by Riou Okkerse-Ruitenberg who developed a

coupled kinetic and reactor model. This allows the simulation of interaction mechanisms between the gaseous phase

and the surface as well as the description of the crystal growth on (100) diamond surface [45,46]. This model

nevertheless does not take into account the sonic instabilities due to the nozzle and which are detrimental to a

controlled behaviour of the flame.

NASA/CP--2001-210948 298 August 6, 2001
MECHANICAL PROPERTIES OF DIAMOND COATINGS OBTAINED BY FLAME PROCESS AND

APPLICATIONS

The lifetime of a diamond coating depends, most of the time, on its adhesion to the substrate. The main

mechanical application field of the coatings obtained by flame process is the cutting tools one ; those coatings are

sometimes used to coat small size dies intended for deep drawing.

In the case of the cutting tools, the cut essentially leads to the application of compressive normal stresses, but

the shaving deformation principally induces tangential stresses. When the coatings well support compressive

stresses, their breaking or unsticking are observed as those applied tangential stresses reach a critical value ; this

latest decreases as residual tangential stresses, having a thermal origin, appear between the coating and the substrate.

The tangential forces applied by the shaving to the coating are high, more especially as the coating roughness is

increased by a bigger mechanical hooking. It is a good thing to realise coatings having very small crystals size

and/or being polished by diamond powder to decrease the shaving friction on the tool as well as the tangential

stresses between the coating and the substrate.

Concerning the deep drawing, coatings are essentially submitted to normal stresses and consequently have long

lifetime [9]. It is however needful that the relative motions of both the die and the metal which has to be drawn, are

small as the fresh surfaces on the piece created during the plastic deformation, are very reactive, leading thus to an

important transfer of metal to the coating, especially when the roughness is high.

DIAMOND COATINGS STUDY

The flame process was studied and developed in our Laboratory ; the tribological behaviour of the so deposited

diamond coatings was investigated, but their characteristics had to be known in order to well interpret the obtained

results. One of the used coatings during the tests was {100} oriented ; the SEM image indicates that the coating is

homogeneous and the crystals joined (figure 3a). Its high crystallinity and purity (lack of graphite) are shown by the

EDS analysis (important carbon peak) and Raman spectroscopy (figure 3b).

,e_ 9000

___8000

•_'_ 7000

"_ 6000

5000

¢M 4000

3000

diamond peak

I , I , I

1200 1400 1600

Wavenumber (cm -1)

a b

Figure 3. (a) SEM image and (b) Raman spectroscopy of a {100} oriented diamond coating

The scanning Raman spectroscopy on large surfaces (600 x 100 gm 2) was also used to characterise the coating

(figures 4,5).

NASA/CP--2001-210948 299 August 6, 2001
lm

iiiii!!_!_i._iii_i_!!_!i!i!_iii!ii_!_!_iiiii_i_ii;i_i_i_iiiiiiiiiiiiiiiiiiiiiiiiiiii_

300!i!! i i iiiiiii ii 100

200 400 600

(a)

gm

i_%iiiiii!_iii/iii_iiiii_iiii!iii!!i!!i!i!!!iiiiiiiiiiii!iiiiiiiil I

250 :::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::1.0

_iiil;_iiiiii!iNi!_iiiiiiii!i!i!!ii?ii!i!!iiiiiiiiiiiiiiiiiiio._ _00_iNiiii_i!!ii!iiNiiiiii!i_ii_iiiiiiiiii!iiiii_iilN!iii_i_iiiiiiiii:

o._

, i , i , i _m

200 400 600

(b)

Figure 4. Surface analysis by Scanning Raman spectroscopy of {100} oriented diamond coating

(a) diamond peak intensity ; (b) I.... ph ...... bon / Idi .... d

• o,:.:.:: .:.:.'.:o• ,.:,:.

;:.::

::::::::::::::::::::::::::::::::::::::::::::::: ..._::::.• _._. "'"'"""''""''''_"'""""' ;a

iii!::i!ii:!:ii!!!!!!!!!!!_!!!:!:i:i._:i:_:i:i:!'! 0 :::.:,

:::::::::::::::::::::::::::::::::::: :'.:_ s ,

Figure 5. Crystals analysis by Scanning Raman spectroscopy of {100} oriented diamond coating

The surface analysis indicates that the coating is relatively homogeneous (figure 4), and the crystals analysis

(figure 5) shows that the centre of the facets contains practically no amorphous carbon and seems to be submitted to

compressive stresses more important than in the periphery.

NASA/CP--2001-210948 300 August 6, 2001
The diamond tribological behaviour is very complex ; to simplify, two parts are involved in it : on the one hand,

the diamond coating itself, characterised by, among other things, its crystals size and orientation ; and on the other

hand, the transfer which can present different sights (shaving or agglomerated particles) and chemical compositions.

All these parameters interfere in the friction process, leading to many-sided and imbricated phenomena. That incited

us to study some of them, and more particularly the influence of the crystals orientation. The main results obtained

with {111 }, {110} and { 100} oriented coatings sliding against High Speed Steel counterface in vacuum, oxygen and

water vapour under two different applied normal loads [47,48,49] are summarised as follows :

o:° SEM observations and EDS analyses showed that the crystals orientation, on the one hand, modifies the nature of

the transfer as it can present the sight of a dense film or little balls which play then the role of a lubricant in the

contact, and on the other hand, influences the way it is fastened on the coating. It also appeared that, except iron, the

main constituent of the counterface, chromium and cobalt are transferred from the disc to the pin in a particular way,

according to the crystals orientation.

o:° the nature of the indentation under oxygen also depends on the diamonds orientation, for a given applied normal

load or for a constant disc hardness.

o:° the way the diamond coatings are modified during the sliding under water vapour is greatly influenced by the

crystals orientation ; two tendencies can be observed : for the {111 } coating, it is rather a diamond deterioration

inducing the formation of disordered amorphous carbon. The same phenomenon occurs for the {100} coating, but in

an attenuated way ; there is moreover asp 3 to sp 2 transfbrmation leading to graphite at low load. Analogous results

are obtained under oxygen, but the observed phenomena occur faster because of the activation due to water vapour.

With the object of coating the cutting tools, it seems that {100} oriented coating without water vapour could be

selected.

CONCLUSION

The flame process analyse, in the aim of obtaining diamond coatings of high quality, shows that the main

problem is to rigorously control and remain constant all the parameters directing the coatings elaboration. The slight

flame instabilities, due to the interactions between the gas flow and the nozzle sides, represent a major difficulty

which has to be solved ; the non-uniform interaction of the flame with the substrate surface is, in the same way, an

important phenomenon preventing the obtention of diamond crystals having constant size and orientation.

The size of the coated surfaces remains relatively moderate, limiting thus, at present, their use in industrial

applications essentially to cutting tools.

The process evolution, even if it induces an increasing complexity of the devices, portends however interesting

developments in a next future.

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NASA/CP--2001-210948 303 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Optical Quality Diamond Wafer Synthesis by Multi-Cathode DC PACVD

Kwang Yong Eun*, Jae-Kap Lee and Young-Joon BaLk,

Thin Film Technology Research Center, Korea Institute of Science and Technology, Korea

Wafer scale diamond films with optical transparency and high thermal conductivity have been synthesized by

the multi(7)-cathode direct current plasma assisted chemical vapor deposition (DC PACVD) method. Diamond

wafers have been grown on the metal substrates with 76 nun diameter under the deposition pressure of 100 Tort

and the input power of 15 kW, respectively. Synthesized free-standing diamond films have been polished and

their Ra values range several tens nm. Optical(visible and IR range) and thermal properties of diamond films

have been observed. Depending on the methane concentrations, there were large differences in the measured

values. A transparent diamond wafer deposited at an optimum condition showed high transmission of 70 % at

the 10.6 pm wavelength and high thermal conductivity of 21 W/cmK at room temperature. Variations of

transmission and thermal conductivity within a wafer were + 10%. The tangent loss and Roman spectroscopy of

the diamond films have been measured and the included impurity levels have been determined analysis

depending on the wafer location and the film thickness by RBS measurement.

Key words: CVD Diamond, de plasma, optical transmission, thermal conductivity

Prime Novelty: Properties of diamond films deposited by the multi-cathode DC PACVD method

Corresponding author : Kwang Yong Eun

Address : RO. Box 131, Cheongryang, Seoul, 130-650, Korea

Tel : 82-2-958-5491, Fax : 82-2-958-5489, e-mail : kyeun@kist.re.kr

NASA/CP--2001-210948 304 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Metastable Growth of Diamond in Molten Alloy

Yen-Te Chen, Kai-Hon Hsu, Hsiao-Kuo Chung, James C. Sung

Kinik Company, Taipei, Taiwan, R.O.C., and

National Taipei University of Technology, Taipei, Taiwan

Abstract

Diamond is typically grown in molten metal of Group VIIIB (Iron

Group) at high pressure where diamond is stable (Sung 1997). The carbon

source is graphite and the molten metal serves as a catalyst-solvent (Sung &

Tai 1997a). Diamond can also be deposited in partial vacuum where

graphite is stable by thermal decomposition of carbonaceous gas (e.g.,

methane). Such a CVD process often requires the presence of hydrogen

atoms as catalyst. At low pressure, diamond can be easily converted into

graphite when brought in contact with molten Group VIIIB transitional

metals. However, Roy (1992) found that if hydrogen atoms are present, the

reverse reaction could take place, i.e., diamond could be formed from

graphite inside the molten metal even though graphite is the stable phase.

Since then, numerous experiments have been conducted in synthesizing

metastable diamond with molten metal under the atmosphere of hydrogen

atoms. However, it has been disputed if the diamond so formed is derived

from carbonaceous gas as in the case of CVD methods (Chung & Sung

2001), or diamond is indeed nucleated inside molten metal that is overly

saturated with carbon (Mallika et al, 1999). It was proposed that diamond

could be formed from overly saturated carbon atoms provided these carbon

atoms are surrounded by hydrogen atoms inside the pseudo lattice of molten

metal (Sung & Tai 1995). In this case, the melting point of the metal could

be suppressed well below the eutectic point of metal-carbon due to the large

incorporation of hydrogen.

In order to further elucidate the mechanisms of the diamond formation

in liquid phase, various binary metal systems (Cu-Mn, Cu-Ni, Cu-Pd, Ni-Pd,

Mn-Pd, Ni-Mn) were prepared from their component powders. These

powders were mixed with graphite powder and used as a precursor in a hot

filament CVD system. The atmosphere contained either pure hydrogen or it

also incorporated about 1% methane. The metal precursors were melted at

NASA/CP--2001-210948 305 August 6, 2001
100-200 °C below eutectic points of metal-carbon and held in place for a

period of about 3 hours.

Diamond crystals of a few microns across were found either exposed on

the surface of frozen molten metal, or they were wrapped inside a thin metal

skin, or both. It would appear that diamond was formed via competing

mechanisms of both CVD route and liquid phase synthesis. The latter

mechanism is limited by the diffusion of hydrogen atoms to permeate the

molten alloy. The alloys used in this study fall in two camps. One group

contains hydrogen getters, such as Pd, and to some lesser degree, Ni. The

other group contains carbide or hydride former, such as Mn. If the diffusion

of hydrogen atoms is fast, as in the case of using Pd alloys, diamond could

form inside molten alloy before CVD deposition of diamond took place on

the surface of molten alloy. On the other hand, as in the case of using Mn

alloys, only CVD diamond may be formed.

Key Words: Diamond Synthesis, CVD Diamond, Liquid Phase Synthesis,

Palladium

Introduction:

The feat of diamond synthesis marked a major milestone for the human

advancement of technology. In early 1950s two independent routes were

pursued for the synthesis of diamond. On one hand, graphite was squeezed

under ever increasing pressure and temperature (e.g., by General Electric

scientists) with the hope that it could be transformed into the stable form of

diamond. On the other front, methane was decomposed at high temperature

in the presence of diamond seeds (e.g., by Union Carbide scientists) with the

intent to form metastable diamond on the surface. Both routes have led to

the commercial production of diamond practiced today. The former has

become the dominant way of making diamond grits; and the latter, diamond

films.

The success of making diamond commercially depends on the finding of

a right catalyst for the synthesis. In the case of high-pressure synthesis, the

catalyst is a molten metal that also acts as a carbon solvent (e.g., iron, cobalt,

nickel, or their alloys). The molten metal can dissolve and disintegrate

graphite into solute atoms or suspended clusters. The catalytic action can

then convert these atoms and clusters into diamond (Sung & Tai 1997a). If

these catalysts are not used, graphite could also be converted directly into

diamond (Sung & Tai 1997b), however, the kinetics of such conversion is

NASA/CP--2001-210948 306 August 6, 2001
extremely slow so the pressure becomes too high for any practical

application (Sung et al 1996).

On the other hand, the metastable growth of diamond can be accelerated

by using hydrogen atoms as catalyst (Sung & Tai 1995). The hydrogen

atoms can stabilize the sp 3 bonds of carbon atom and convert sp 2 bonds, if

formed, to sp 3.

If graphite is brought in contact with molten metal of iron group at low

pressure (below one atmosphere), diamond cannot be formed. On the

contrary, diamond will be rapidly graphitized. This is why diamond

superabrasive, although extremely effective in cutting almost every material,

is powerless to machine steel at high temperature. However, if hydrogen

atoms are present in the molten iron, not only diamond is no longer back

converted to graphite, but also the reverse process may take place. In other

words, graphite may convert to diamond under the influence of hydrogen

atoms even at ambient pressure. It would appear that hydrogen atoms can

catalyze diamond formation in liquid phase just as in gas phase.

There are two schools of thought as why diamond can be formed molten

metal at low pressure. Most believe that diamond, although in contact with

metal, is actually deposited from gas phase. Hence, molten metal only

serves the role as a substrate. Indeed, experiments demonstrated that if the

gas did not contain methane, diamond did not seem to form (Chung & Sung,

2001). Thus, diamond appears to derived from decomposed methane rather

than from dissolved graphite.

However, there are evidences that hydrogen atoms do permeate through

molten metal. Many metals (e.g., Cu, Ag) that do not dissolve carbon or

hydrogen can incorporate atoms of both kinds in substantial quantities (e.g.,

more than one tenth of the metal atoms). As a result, the melting point is

greatly (e.g., more than 100 °C) suppressed below the ordinary eutectic

point. There are also indications that diamond could be formed inside

molten metal. If this is indeed the case, hydrogen atoms may play a vital

role. It was proposed that small atoms of carbon and hydrogen form

networks inside the voids of large atoms of molten metal. Such networks

may stabilize carbon atoms in sp 3 bonding, and hence diamond formation

(Sung & Tai 1995).

It is the objective of this research to compare the rates of diamond

formation via these two competing mechanisms. Moreover, the

effectiveness of synthesizing diamond in various binary systems is assessed.

NASA/CP--2001-210948 307 August 6, 2001
Experimental

This research was conducted in an atmosphere controlled by hot

filament chemical vapor deposition (HFCSD). A sketch of the reactor is

shown in Figure 1. The hot filament was used to decompose gas as well as

heating the sample.

VIEWING

PORT

CAPACITANCE

GAGUE --_

CE

To

ROTARY PUMP

J

FILAMENT II

TO POWER SUPPLIERE

GAS INLET

1

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H2

MASS FLOW

CONTROLLER

CH4

Fig. 1. The schematic representation of the HFCVD reactor.

The samples were prepared according to various combinations of binary

metals as shown in Table 1.

Table 1" Binary Composition of Samples

System Ratio

Cu-Mn 60:40, 70:30, 80:20

Cu-Ni 10:90, 50:50, 90:10

Cu-Pd 10:90, 50:50, 90:10

Mn-Pd 18:82, 32:68, 42:58

Mn-Ni 68:32, 58:42, 48:52

Ni-Pd 30:70, 40:60, 50:50

NASA/CP--2001-210948 308 August 6, 2001
Microns sized powders of pure metal and graphite were intimately

mixed by hand. Each five gram mixture contains 2 wt% of graphite. 0.5-

gram samples were pressed to form pallets of about 10 mm in size and 1 mm

in thickness. The pallets were then placed in the holder made of graphite

that was latter set on the substrate stage. The reactor was evacuated to a

pressure of about 20 torts. Hydrogen gas with or without 1% methane was

added to the chamber at a flow rate of 300 sccm. The filament was then

heated to about 2200 °C. Each experiment lasted about 3 hours before the

power was cut off.

Results and Discussions

Some of the above experiments yielded diamond of various amounts. It

would appear that the type of the alloy (not the composition) is the dominant

factor in determining the outcome of diamond output. Figure 2 shows the

representative results from these experiments.

60Cu-40Mn 80Cu-20Mn+G

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NASA/CP--2001-210948 309 August 6, 2001
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NASA/CP--2001-210948 310 August 6, 2001
o Iii: ...'_!iii.i_:i.>_":'--'_I''" _ _ ""'i'i'i.":i'!:!':::"".. _ "''':

WITH CH4 WITHOUT CH4

Fig. 2: SEM micrographs of diamond formed in various binary alloys.

Based on the above SEM micrographs, it would appear that diamond

crystals had nucleated inside the molten alloys. Many crystals wrapped

inside are several microns in size and they tend to show euhedral crystal

shape. The appearance shows that diamond was formed in liquid phase.

Moreover, if methane was added to hydrogen atmosphere, smaller diamond

crystals (many are submicron in size) could form. But these secondary

crystals were exposed outside the molten metal. Such an occurrence

indicates that they were precipitated out of gas, similar to the CVD growth

of diamond. The exposed diamond was also the dominant occurrence for

liquid phase synthesized diamond reported in literature. Thus, this research

proves that diamond can be nucleated and grown inside molten metal at

partial vacuum.

The amount of diamond formed also suggests that the ability for the

molten alloy to nucleate and grow diamond increases in the order of Mn, Ni,

Cu, and Pd. Both Mn and Ni are powerful catalysts for synthesizing

diamond under high pressure (Sung & Tai. 1997a). On the other hand, Cu is

well known for its inability to react with carbon. It does not dissolve carbon,

nor catalyze it to form diamond. It seems that the ranking of catalytic power

for molten metals is just in a reverse order between diamond synthesized

under high pressure and at low pressure.

NASA/CP--2001-210948 311 August 6, 2001
Moreover, although Pd is also a catalyst for high-pressure synthesis, it is

a less powerful one. However, it becomes the most effective catalyst for

low-pressure diamond synthesis.

The above seemingly puzzling observations are all anticipated by the

crystal chemical model proposed by James C. Sung. According to this

model, the catalytic power of a molten metal is determined by its atomic size

and the electronic configuration (Sung & Tai 1997a). The atomic size must

be such that it is as close to the distance (2.46 A) between every other

carbon atom on a graphite basal plane. The electronic configuration must be

such that it contains as close to half filled d-orbitals as possible (Fig. 3).

(a)

b

i

NASA/CP--2001-210948 312 August 6, 2001
(b)

I3

0.

R ! R

C

A

A

i

_...(

----i

-t

Fig. 3: Direct puckering of graphite into diamond. An effective catalyst has the right

atomic size (large circles) to match every other carbon atom (small circle) on the basal

plane (0001) of graphite (top diagram).Its atoms can also strongly interact with the

matching carbon atoms. The interaction will pull half of the carbon atoms (lightly dotted)

away from the basal plane of graphite and toward the metal atoms. The result is chain

movement that may pucker graphite structure into diamond's (lower diagram).The power

of a catalyst is manifested in the thickness of a graphite flake that it can pucker into a

diamond nucleus. The middle diagram illustrates the catalytic conversion (all black atoms

moving upward and white ones, downward) of rhombohedral graphite (3R or ABC type)

into a cubic diamond. The lower diagram show a similar catalytic conversion (all white

atoms moving upward and black ones, downward) of a hexagonal graphite (2H or ABA

type or it derivative AAA type) into Ionsdaleite (hexagonal diamond).

NASA/CP--2001-210948 313 August 6, 2001
Based on this model, it was predicted that the most powerful catalysts

for high-pressure synthesis of diamond are in the decreasing order of Co, Fe,

Mn, Ni, Cr. This prediction matches well with empirical experience as these

five transition metals are indeed the most used catalysts for industrial

synthesis of diamond (about 350 ton a year) under pressure.

According to this crystal chemical model, Cu cannot catalyze the

diamond formation as its atoms lack any vacancy in d-orbitals, the essential

requirement for attracting the valence electron of carbon atoms. This is

exactly observed in the industry that Cu has no ability to make any diamond

under high pressure.

The crystal chemical model is also applicable for low-pressure synthesis

of diamond under the atmosphere of hydrogen. In this case, the atomic size

is applicable as the container for carbon and hydrogen solutes. However, the

electronic configuration is less important because the dissolved hydrogen

atoms will provide the electron for attracting carbon atoms (Figure 4).

O C _.IOD: 2,1 air OCl_lJ'_;iI0]. 511(:

O M Ot,',:, i3._ :, ,:':ralu';;Tal ai:_

Fig. 4: Dissolved C-H4 cluster in a liquid metal. Each unit cell contains four metal

atoms, four octahedral sites that share an edge, and eight tetrahedral sites that share a

comer. Each octahedral site is surrounded by eight tetrahedral sites, and each tetrahedral

sites is surrounded by four tetrahedral sites, all share a face. Out of the four octahedral

sites of each unit cell, one is occupied by the C atom. Out of the eight tetrahedral sites

that surround this C-occupied octahedral site, four are occupied by hydrogen atoms that

every other site. As a result, the stoichiometry of the composite is M4CH4. In such an

arrangement, each C atom is surrounded by four hydrogen atoms that are separated by

metal atoms. In other words ,the structure may be view as containing interstitially

dissolved C-H4 cluster in void of the metal frame work.

NASA/CP--2001-210948 314 August 6, 2001
Based on this model, it was predicted that the catalytic power of the four

metals we experimented would decrease in the order of Cu, Ni, Mn, Pd

(Sung & Tai 1995). However, the above reference states "The ranking of

metal solvents as discussed assume that hydrogen molecules in the

atmosphere can be readily dissociated to form hydrogen atoms on the

surface of the molten metal. But, if the dissociation of hydrogen molecules

becomes the rate-limiting step for the dissociation of hydrogen atoms in the

metal, then the ranking of metal solvents may be different. For example, it

is known that Pd is particularly effective in dissociating hydrogen molecules

and incorporating the dissociated hydrogen atoms into its voids. In 1995,

the author (Sung 1995) proposed coating a proper metal solvent with Pd for

the growth of metastalbe diamond in liquid phase. It was hoped that the

dissociated hydrogen atoms could diffuse through Pd to the metal solvent to

help catalyze the diamond formation. This idea was proven effective by the

experiment performed by C. C. Hong in his 1996 master thesis at Chin Hwa

University of Taiwan, R.O.C. entitled: Liquid phase synthesis of diamond in

microwave plasma."

Thus, it would appear that it was Pd, the powerful hydrogen getter, that

made diamond much faster than any other metal studied in this research, and

hence the catalytic power for the diamond synthesis at low pressure under

hydrogen atmosphere decreases in the predicted order of Pd, Cu, Ni, Mn, the

exact ranking sequence observed in this research.

Conclusion

Diamond formed by molten alloy at low pressure under hydrogen

atmosphere can proceed in two competing routes: via thermal decomposition

of hydrocarbons gas (e.g., methane), and via dissolution of carbon solids

(e.g., graphite). In the latter case, the rate-limiting step is the diffusion of

hydrogen atoms to fill the tetrahedral voids in the molten metal. If the latter

becomes a bottleneck, naked diamond crystals will form on the surface of

metal droplets. On the other hand, if hydrogen atoms can permeate the melt

rapidly, unexposed diamond crystals will form inside the molten metal.

The two mechanisms are competing in the experiments of the current

research. For Pd and Cu alloys, the permeating of hydrogen atoms are fast,

hence diamond crystals formed inside the molten metal are much larger than

those exposed outside. For Mn and Ni alloys, their catalytic powers are

lower than Pd and Cu, hence, submerged diamond crystals are scarce or

NASA/CP--2001-210948 315 August 6, 2001
entirely lacking, leaving only sporadic exposed diamond crystals deposited

by CVD.

The above observations agrees well with the crystal chemical model

proposed by James Sung (Sung & Tai 1995) that the catalytic power of

metals used for this research will decrease in the order of Pd, Cu, Ni, Mn.

References

Roy, R. 1992, Press conference for diamond synthesis by the low-pressure

solid-state source, Penn State University, U. S. A.

Chung, H. K., and James C. Sung, 2001, Low Pressure Synthesis of

Diamond in Molten Metal, Diamond & Related Materials, in press.

Mallika, K., R. C. DeVries, & R. Komanduri, 1999, Thin Solid Films, 339,

19-33.

Sung, James C., & M. F. Tai, 1995, High Temperatures-High Pressures, 27,

611-628.

Sung, James C., M. F. Tai, S. S. Cheng, Q. S. Huang, Y. D. Yao, 1996, High

Temperatures-High Pressures, 28,499-521.

Sung, James C., M. F. Tai, 1996, High Tem

Sung, James C., 1997, High Temperatures-High Pressures, 29, 253-293.

Sung, James C., & M. F. Tai, 1997a, Int. J. Refractory Metals & Hard

Materials, 15,237-356.

Sung, James C., & M. F. Tai, 1997b, High Temperatures-High Pressures, 29,

631-648.

NASA/CP--2001-210948 316 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FABRICATION OF DIAMOND MICROPATTERNS BY ECR OXYGEN PLASMA

WITH A METAL NAPTHENATE MASK

Shuji Kiyohara and Katsumi Mori

Department of Electronics and Computer Science, Science University of Tokyo in Yamaguchi,

1-1-1 Daigaku-dori, Onoda, Yamaguchi 756-0884, Japan, e-mail address: kiyohara@ed.yama.sut.ac.jp,

fax number: +81-(0)836-88-3844, phone number: +81-(0)836-3500 ext.653

ABSTRACT

The fabrication of diamond micropatterns by ECR oxygen plasma with a metal naphthenate mask has been

investigated using electron beam lithography technology. The use of metal naphthenates as mask materials with

resistance to oxygen plasma in order to form an oxide film on the surface was proposed. The exposure characteristics

of metal naphthenates and the etching characteristics of chemical vapor deposited (CVD) diamond and metal

naphthenate films processed with electron cyclotron resonance (ECR) oxygen plasma were investigated.

Furthermore, the crystal structure of diamond micropatterns fabricated by this process was evaluated using Raman

spectroscopy. It was found that the etching resistance of Yttrium (Y), barium (Ba) and copper (Cu) naphthenates

respectively is low to oxygen plasma. Therefore, the resistant mask material which forms a compound oxide was

prepared by mixing the metal naphthenates with a Y:Ba:Cu molar ratio of 1:2:3. It was found that the metal

naphthenates exhibited negative exposure characteristics upon electron beam irradiation. The sensitivity and the

gamma value were 2.4 x 10 -3 C/cm 2 and 1.6, respectively. The maximum etching selectivity (CVD

diamond/YBa2Cu307 naphthenate) of 10 was obtained under the ECR oxygen plasma etching conditions of a

microwave power of 300 W and oxygen gas flow rate of 3 sccm. The diamond micropatterns were fabricated under

the optimum electron dose of 6.6x 10 -3 C/cm 2 and plasma etching conditions (microwave power: 300 W, gas flow

rate: 3 sccm) of high etching selectivity. The CVD diamond micropatterns of 1 _m-line, 0.5 _m-line and 0.5 _mcomb, which have a height of 1.2 _m for the etching time of 1 hour were fabricated. The crystal structure of the

diamonds after etching and the micropatterns fabricated by this process remained constant; Raman spectra indicated

only the presence of a diamond peak at 1333 cm -1.

Keywords: diamond micropattern, metal naphthenate, ECR oxygen plasma, electron beam lithography

INTRODUCTION

The many unique properties of diamond, such as a wide band gap of 5.5 eV, high hardness, high thermal

conductivity, low thermal expansion, negative electron affinity (NEA) on a hydrogenated diamond surface and

chemical inertness (resistivity against acids, heat and radiation), make it a promising material for mechanical,

electronic and optical applications. The micropatterning technique of diamonds is essential to the fabrication of

many electronic and optical devices, and micro-gear and motor made of diamond using micromachines. Therefore,

electron cyclotron resonance (ECR) oxygen plasma etching of chemical-vapor-deposited (CVD) diamond films has

been investigated. We found that ultra fine processing of CVD diamond films is possible with ECR oxygen plasma

(ref. 1). In order to apply this etching technique for diamond-based devices, the micropatterning of CVD diamond

films using electron beam lithography technology was investigated. In this case, the most important point in the

formation of micropatterns is the mask resist materials. In general, resist materials [polymethyl methacrylate

(PMMA)] used in the fabrication of semiconductor devices are decomposed to carbon monoxide (CO) and hydrogen

monoxide (HO) on exposure to oxygen plasma, indicating that they are composed of carbon, hydrogen and oxygen.

Consequently, the use of metal naphthenates as mask resist materials with resistance to oxygen plasma in order to

form an oxide film on the surface was proposed (ref. 2). In this paper, we report the exposure characteristics of metal

naphthenates and the etching characteristics of CVD diamond and metal naphthenate films processed with ECR

oxygen plasma. Furthermore, the crystal structure of CVD diamond micropatterns fabricated by this process was

evaluated using Raman spectroscopy.

NASA/CP--2001-210948 317 August 6, 2001
EXPERIMENTAL APPARATUS AND PROCEDURE

ECR oxygen plasma etching of CVD diamond and metal naphthenate films

Polished polycrystalline diamond films 12 _m thick (arithmetical mean deviation of profile Ra: 10 nm) and metal

naphthenate films approximately 1 _m thick on silicon wafers (10x 10x0.5 mm 3) were used as samples. The

polycrystalline diamond films were synthesized by the hot filament CVD method on silicon-crystal sub strates (10 x 10

x5 mm 3) using a gas mixture of methane and hydrogen gases. Typical deposition parameters used to grow CVD

diamond films were as follows; filament temperature: 2150 °C, substrate temperature: 850 °C, filament-substrate

distance: 10 mm, CH4 flow rate: 5 sccm and H2 flow rate: 1000 sccm.

Metal naphthenates consist of cyclopentanes or cyclohexanes, methylene chains -(CH2)n-, -COO- and metals,

which have mainly following structure: [(cyclopentane)-(CH2)n-COO]-m - M m+, where M is metal atom (ref. 3).

Yttrium (Y), barium (Ba) and copper (Cu) naphthenates (Nihon Kagaku Sangyo Co., Ltd.) were used as mask

materials. They were sticky liquids at room temperature and stable in air. The silicon wafers were spin-coated with

the metal naphthenates at 2500 rpm for 10 s and then dried in air at 110 °C for 15 min to form metal naphthenate

films. The resulting film thickness of the metal naphthenates was approximately 1 _m. These samples were

processed using a plasma (Advanced Film Technology Inc. AFTEX PS-501) etching apparatus with an ECR-type

oxygen source, as shown schematically in figure 1.

Bias voltage_hermocouple

electrod Ill e_ctrode

- -- _ t -- -- _--_ Magnet(875G) _Gas

, , 1_7_ Quartz window

l il Molybdenum_- -- J_ ITM

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Working i sDia%l% nd _-_ I.-_--_-_,-_-_--_.1 I _ I Recta II ngular

7vaciuu I • owe .uoo,.

pump

Figure 1. Schematic diagram of the plasma etching apparatus with an ECR-type oxygen source.

As shown in figure 1, the samples were fixed on a molybdenum holder positioned approximately 5 cm from the

plasma chamber outlet. The molybdenum holder was equipped with a thermocouple and a bias voltage electrode.

Etching steps were made in CVD diamond and metal naphthenate film surfaces by striping a silver paste mask and an

aluminum plate, respectively. Then, step heights were measured using a diamond stylus surface profilometer (Tencor

Instruments ALPHA STEP 200).

Fabrication of diamond micropatterns

Figure 2 shows the micropatterning process of CVD diamond films. The diamond/Si substrates were first spincoated with the metal naphthenates at 2500 rpm for 10 s and then dried in air at 110 °C for 15 min. An electron beam

drawing system was manufactured in-house by modifying a conventional scanning electron microscope (SEM)

(TOPCON DS-130S) (ref. 4). An electron beam was scanned on the metal naphthenate films according to patterns

entered in the computer. Metal naphthenate micropatterns were developed by the removing the unirradiated area with

toluene. The samples were processed with an ECR oxygen plasma under etching conditions of high etching

selectivity. Finally, the mask micropatterns of the metal naphthenates remaining on the diamond films were removed

with phosphoric acid at 80 °C. In contrast to the micropatterning process using lift-off method (ref. 5), this one has

certain advantages including few residue and simple process because it is not necessary that the film (aluminum) is

deposited over the resist and the substrate.

NASA/CP--2001-210948 318 August 6, 2001
:._,V_D:diamon:d.:

,Ik

/re

,................

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,!k

Spin-coating with

Metal naphthenates

2500rpm, los

110°C, 15min, in air

Electron beam

drawing 3

6.6x10 C/cm 2

liiiii'iiiii'ii iiiiil

Development

with toluene

,,I,,

o÷,oL o*, oT

t , i i . ,

ECR oxygen

_.._._:i:i_ plasma etching

.................. 300W, 3sccm, lh 4,

Mask removal with phosphoric

acid at 80°(;

Figure 2. Micropatterning process of CVD diamond films.

EXPERIMENTAL RESULTS AND DISCUSSION

Exposure characteristics of metal naphthenates

In order to obtain the optimum electron beam drawing conditions, the dependence of the remaining film

thickness of metal naphthenates on the electron dose after development, as shown in figure 3, was investigated. The

remaining thickness is normalized with the thickness of the spin-coated film being equal to 1. As shown in figure 3,

the metal naphthenates exhibited negative exposure characteristics. The sensitivity and the gamma values of metal

naphthenates were 2.4x 10 -3 C/cm 2 and 1.6, respectively. The sensitivity is ten times lower than that of PMMA resist

and higher than that of electron-beam-induced CVD (ref. 6). The optimum drawing conditions to form micropatterns,

based on these results, were a spot size and probe current of the electron beam of 10 nm and 150 pA, respectively.

The electron dose was typically set to be 6.6x 10 -3 C/cm 2.

¢.D

O0

(D

z-Z

N

E

o

z

1.o ......................i...................... ' ..............

00.50................... i..................... l......................

10-4 10 3 10 2 10 1

Dose (C/cm 2)

Figure 3. Dependence of the remaining film thickness of metal naphthenates

on the electron dose after development.

NASA/CP--2001-210948 319 August 6, 2001
ECR oxygen plasma etching characteristics of CVD diamond and metal naphthenate films

In order to fabricate diamond micropatterns under etching conditions that give high etching selectivity, the

dependence of the plasma etching rates of CVD diamond and metal naphthenate films on the oxygen gas flow rate

was investigated, as shown in figure 4. At an oxygen gas flow rate up to a limit of 10 sccm, the etching rate of CVD

diamond films first increased with increasing gas flow rate and reached maximum rate at 3 sccm, then decreased

gradually with increasing gas flow rate. In the low-gas-flow-rate and high-gas-flow-rate regions, the etching rate

decreases because of the lack of oxygen gas reaching the film surface and the reduction in the amount of reactive

species reaching the film surface, respectively (ref. 7). The etching rate of Y, Ba and Cu naphthenates films first

increased with increasing gas flow rate and reached maximum rate at 3, 3 and 7 sccm respectively, then decreased

gradually with increasing gas flow rate. It was found that the etching resistance of Y, Ba and Cu naphthenates

respectively is low to oxygen plasma. Therefore, the resistant mask material which forms a compound oxide was

prepared by mixing the metal naphthenates with a Y:Ba:Cu molar ratio of 1:2:3. The etching rate of YBa2Cu307

naphthenate films first increased with increasing gas flow rate and reached maximum rate at 7 sccm respectively,

then decreased gradually with increasing gas flow rate. We believe that the YBa2Cu307 naphthenate as mask material

has resistance to oxygen plasma in order to form oxide film on the surface. Figure 5 shows SEM micrographs of the

mask micropatterns of Y, Ba, Cu and YBa2Cu307 naphthenates before and after ECR oxygen plasma etching

(microwave power: 300W, oxygen gas flow rate: 3sccm, etching time: lh). From figures 5 (a) and (b), it can be seen

that Y and Ba naphthenates were not able to fabricate fine mask patterns, and then they broke after ECR oxygen

plasma. From figure 5 (c), it can be seen that Cu naphthenate was able to fabricate fine mask patterns, but that has

low etching resistance. Therefore, YBa2Cu307 naphthenate used as the mask material because it was able to fabricate

fine mask patterns as shown in figure 5 (d), and then has high etching resistance as shown in figure 4.

The dependence of the etching selectivity (CVD diamond film/YBa2Cu307 naphthenate film) on the oxygen gas

flow rate is shown in figure 6. A maximum etching selectivity of 10 was obtained under the plasma etching

conditions of a microwave power of 300 W and oxygen gas flow rate of 3 sccm.

e.-

E

v

,u

..C:

o

ILl

CVD diamond film • 0

Metal naphthenate film 'A

II

t

,q,

,,0"- ',

,e

/

Y

Ba

Cu

Y [:?,a2Cu307

,_:7 ,,--..... _,.______--:_ lb ...__..._._t

5 10

Oxygen gas flow rate (sccm)

Figure 4. Dependence of the plasma etching rates of

CVD diamond and metal naphthenate films on the

oxygen gas flow rate. (microwave power: 300W)

1

CVD Diamond film/

YBa2Cu307 naphthenate film

>,

>

•_ 10

__e

.E

cO

-_ 5

iii

; I I I I i I I I I ] 5 10

Oxygen gas flow rate (sccm)

Figure 6. Dependence of etching selectivity (CVD

diamond film/YBa2Cu307 naphthenate film) on the

oxygen gas flow rate. (microwave power: 300)

Fabrication of diamond micropatterns

The micropatterns of CVD diamond films were fabricated under the optimum electron dose of 6.6x10 -3 C/cm 2 and

etching conditions (microwave power: 300W, oxygen gas flow rate: 3 sccm) of high etching selectivity. The SEM

micrographs of the resulting CVD diamond micropatterns of (a) 1 _m-line, (b) 0.5 _m-line and (c) 0.5 _m-comb,

which have a height of 1.2 _m for the etching time of 1 hour, are shown in figure 7.

NASA/CP--2001-210948 320 August 6, 2001
Before oxygen pla_m_ etching After oxygen p|a_ma etching

(a) Yttrium

• o.,.,.,.,.,.,...... ..............................

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::i_i_i_i_i_i_!:!_i_i_i_i:!_i_i_i_i_i_!_i_i_i_i_i_!_i_i_i_i_i_!_i_!_i_i_

:i:i:i:!_:i:i:!:i-')::i:i:!:i_._i:i:i:i:'$i:i:i:i:i'i:i:i:i:i:i$1:i:i:

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

(b) Bari um

:""" """_;ii"••:ziii"•"i!_i_iiiiiiii!_;!iiiiiii!_;_i!iiiii_i!_;i

:" y • .;;;;., ==================================

(C) CopPer

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

_?kZ_iiii!!?kZ_iiii!!?i_.:iiii!!i_.'iiii!!i_'iiii!!iNii._:iii -_

SNii'i'_ii'i'!:i$'."i'i:!:i_'i'i'!:._'i'i'!::_Z_'i"i._

:::::::::::::::::::::::::::::::::::::::::::::::::::::

(d) YBa_Ous07 __'_

Figure 5. SEM micrographs of the mask micropatterns of metal naphthenates before and after ECR oxygen

plasma etching (microwave power: 300W, oxygen gas flow rate: 3sccm, etching time: lh).

_;_:i:i:i:ie!:i_::i:i:i:i:i:_:i:i:i:i_'N:___ iiiiiiiiiiii!!ii::".:iiiiiiiiiiiiii._i._._::::!_ii:" .:iiiiiiiiiii::""""

..::":.iiiiiiiiiii_iiiiiiiiiii!!:i!;:iiiil:ii!_:.::.iiii!i'i_::ii i': ..::f-" ':':':':"':':" ":':':':':':" :'':_:'" ":':':':" "'

..:'.::i:i:i:i:::_'Y.:i:i:i:i:i:::!::..::i:i:i:i::'::'".:i::" ::'..::" ..!".."

,,,,,,,,,,,, o...., _:::: .._,-'::_,-',_'-,_-';-'. iii::..:. ,::::." :ii /.::."

Figure 7. SEM micrographs of the resulting CVD diamond micropatterns of (a) lttm-line, (b) 0.Sttm-line and

(c) 0.Spin-comb (microwave power: 300W, oxygen gas flow rate: 3sccm, etching time: lh).

NASA/CP--2001-210948 321 August6,2001
Ramanspectroscopy wasperformedusinganargonlaser(514.5nm)primarybeamtoexaminethecrystal structure

of theCVDdiamondmicropatterns fabricatedbythisprocess, aswellasthefilmsbeforeandafterECRoxygen

plasmaetching. Ramanshiftsweremeasuredoverarangeof500-2000cm-1atroomtemperature. Figure8showsthe

RamanspectraoftheCVDdiamondfilms(a)beforeand(b)afteretching(microwavepower:300W,oxygengas

flowrate:3sccm,substratetemperature: 100°C,etchingtime:1h),andtheresultingmicropatterns (c)1_m-line,

(d)0.5_m-lineand(e)0.5_m-comb. Fromfigure8,it canbeseenthatheRamanspectraindicateonlyadiamond

(sp3bonding)peakat1333cm-1.ThereforethecrystalstructureoftheCVDdiamondfilmisnotchangedafter

etchingandtheresultingmicropatterns.

CVD diamond film

Microwave power: 300W

Gas flow rate: 3sccm 520(Si)

_ 1333(diamond) j

(C -- "

t_

c (d_

-- (e_

2000 1500 1000 500

Raman shift (cm -1)

Figure 8. Raman spectra of the CVD diamond films (a) before and (b) after etching (microwave power: 300W,

oxygen gas flow rate: 3 sccm, substrate temperature: 100 °C, etching time: 1 h), and the resulting

micropatterns of (c) 1 pro-line, (d) 0.5 pro-line and (e) 0.5 pro-comb.

CONCLUSIONS

The fabrication of diamond micropatterns by ECR oxygen plasma with a metal naphthenate mask was

investigated. It was found that the etching resistance of elemental Y, Ba and Cu naphthenates respectively is low to

oxygen plasma. The use of YBa2Cu307 naphthenates prepared by mixing with a Y:Ba:Cu molar ratio of 1:2:3 as

mask resist materials with resistance to oxygen plasma in order to form an oxide film on the surface was proposed.

The CVD diamond micropatterns of 1 _m-line, 0.5 _m-line and 0.5 _m-comb, which have a height of 1.2 _m for the

etching time of 1 hour were fabricated. This technique is expected to enable applications for diamond-based devices.

REFERENCES

1. S. Kiyohara, Y. Yagi, K. Mori: Plasma etching of CVD diamond films using an ECR-type oxygen source.

Nanotechnology, vol. 10, no. 4, 1999, 385.

2. S. Kiyohara, K. Ayano, K. Mori: Micropatterning of Chemical-Vapor-Deposited Diamond Films in Electron

Beam Lithography. Jpn. J. Appl. Phys., vol. 39, no. 7B, 2000, 4532.

3. A. Kakimi, et al.: Fabrication of Ferroelectric B i4Ti3012 Thin Films by Dipping Pyrolysis of Metal

Naphthenates and Micropatterns by an Electron Beam. Jpn. J. Appl. Phys., vol. 33, no. 9B, 1994, 5301.

4. S. Okamura, et al.: Micropatterning of Ferroelectric Bi4Ti3012 Using Electron-Beam-Induced Reaction of

Metal Octylate Films. Jpn. J. Appl. Phys., vol. 35, no. 12B, 1996, 6579.

5. S. Shikata et al.: Microfabrication technique for diamond devices. 2 nd International Conference on the

Applications of Diamond films and Related Materials, MYU, 1993, 377.

6. S. Matsui, K. Mori: New selective deposition technology by electron beam induced surface reaction. J. Vac.

Sci. Technol., vol. B4, no. 1, 1986, 299.

7. B.N. Chapman, V. J. Minkiewicz: Flow rate effects in plasma etching. J. Vac. Sci. Technol., vol. 15, no. 2,

1978, 329.

NASA/CP--2001-210948 322 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

MORPHOLOGICAL CONTROL OF DIAMOND FILMS IN A 60 KW MICROWAVE

PLASMA CVD REACTOR

Y. Yokota _, Y. Ando, T. Tachibana", A. Watanabe, Y. Nishibayashi,

and K. Kobashi

Frontier Carbon Technology Project / Japan Fine Ceramics Center

6F, Center for Advanced Research Projects, Osaka University

2-1, Yamada-oka, Suita, Osaka 565-0871, Japan

T. Hirao and K. Oura

Faculty of Engineering, Osaka University

2-1, Yamada-oka, Suita, Osaka 565-0871, Japan

ABSTRACT

Growth orientation of diamond crystals and the a parameters for different growth conditions in a 60 kW 915

MHz CVD system have been investigated in order to control morphology of diamond fihns of 152 mm in diameter.

The contour of the c_parameter for the 60 kW system on the graphical plane of the substrate temperature and the CH4

concentration shifts to the higher substrate temperature or the lower CH4 concentration as compared with the results

of the conventional CVD systems. It was also found that the parameter range of diamond growth by the 60 kW

reactor is broader than that of the conventional CVD systems. Furthermore, <100> oriented diamond grains with

pyramid shape were successfully grown almost uniformly on the entire surface ofa 152 mm Si substrate by choosing

an appropriate process condition, which was identified in this study.

Keywords: diamond growth, a parameter, morphology, 60 kW, scale up

INTRODUCTION

Diamond possesses unrivalled semiconductor properties such as wide bandgap, high carrier transport speed,

high breakdown feld, and high thermal conductivity, coupled with high optical transparency and chemical stability.

It is projected that this material will be utilized for various components including robust electronic devices and

packaging for the coming generation (ref. 1). Since chemical vapor deposition (CVD) processes and apparatus of

diamond films were established approximately two decades ago by a Japanese group (ref. 2), many studies have been

carried out to understand the material properties and prove its potential for practical applications. Unfortunately,

most of the R&D efforts were limited in laboratory scale despite the fact that they are primarily motivated by

industrial applications, and less emphasis was placed on scale up technologies. Many issues, including scale up, cost

reduction, reproducibility, etc., have been left unresolved in the individual eflbrts.

Irrespective of specific applications of CVD diamond films, large size, production type CVD apparatus is

necessary to upgrade the present CVD technology. Several efforts have already been made to develop diamond CVD

apparatus and processes for high deposition rate and large area coverage. For instance, 76 mm silicon (Si) wafers

were unitbrmly covered by polycrystalline diamond films though its growth rate was as small as 0.5 _tm/h (ref. 3).

On the other hand, at an increased rate of-5 _tm/h, high quality diamond films of 0.1-1.2 mm thickness were grown

on a wafer of 57 rmn in diameter using a CVD reactor equipped with a 5 kW microwave source (ref. 4). Schelz and

co-workers compared 915 MHz microwave with conventionally used 2.45 GHz, in an attempt to deposit diamond

films on large size substrates. They concluded that diamond films with similar quality can be synthesized as long as

the power density of the plasma is similar (ref. 5). Regarding the large scale apparatus, Sevillano and Williams

successfully constructed a large microwave CVD reactor using 915 MHz microwave to enlarge coating area (ref. 6).

* CoFesponding author: YoshNko Yokota, E-mail:yo-yokota@rd.kcrl.kobelco.co.jp, Fax:+81-6-68794147, Phone:

+81-6-68794146.

aPresentaddress:Electronics Research Labor_ory, Kobe Steel, Ltd.,l-5-5, Yak_sukadai, Nishi-ku, Kobe, Hyogo

651-2271,Japan.

NASA/CP--2001-210948 323 August6,2001
Koidl and co-workers, sharing similar awareness of the present issue, developed a reactor of unique design and

demonstrated diamond films on Si wafers of 100 mm in diameter (ref. 7).

In the present study, the 915 MHz microwave plasma CVD was employed to develop basic CVD process

technologies for high purity and high quality diamond films including nucleation and growth orientation control,

elimination of impurity and residual stress, and cost analyses and minimization. Numerous data accumulated so far

by using lower power (-1.5 kW) microwave plasma CVD was a good guide to explore process conditions in the

large reactor, while it was also of our interest to investigate whether larger scale plasma could produce high quality

diamond fill at high rate.

Some preliminary investigations have been already reported (ref. 8). The present paper focuses on the growth

orientation of diamond crystals and the _ parameter under various growth conditions using the 60 kW 915 MHz

CVD system.

EXPERIMENTAL

An overmoded microwave plasma CVD reactor equipped with a 60 kW power supply (ASTeX/Seki Technotron)

was used to explore process conditions for diamond growth. The apparatus and its operation procedure were

described in ref. 8 in detail.

The substrates were silicon wafers of 25, 51, and 152 mm in diameter. The 25 and 51 mm substrates were

singlecrystalline, while the 152 mm substrates were polycrystalline, chosen by considering their availability and for

comparison. The source gas was 1200-5000 sccm of H2, 4-200 sccm of CH4. The gas pressure Pr was 11.3-17.7 kPa

(85-133 Tort). The microwave input power Pw was 15-60 kW. The substrate temperature T_, a dependent parameter,

was ranged from 950-1370 K (680-1100 °C) as measured by a two-wavelength optical pyrometer.

The morphology was observed using a scanning electron microscope. The a parameter was calculated by

comparing the facet shapes of an isolated diamond particle before and after the growth. The growth orientation was

determined by means of X-ray diffraction. The quality of diamond films was evaluated with Raman spectroscopy

using Ar laser operated at 514.5 nm.

RESULTS AND DISCUSSION

Figure la-c show dependence of the growth rate Rg and the Raman signal width WR on the gas pressure Pr (la),

the microwave input power Pw (lb), and the substrate temperature 7_,(lc), respectively. The flow rates of hydrogen

and methane were fixed as 2000 sccm and 100 sccm, respectively. As seen in fig. la and lb, when neither Pw nor Pr

is fixed both R_,and WR shows no dependence. While, it is obvious in fig. Ic that both R_, and WR increase as T,

increases. Substrate temperature seems to be the most influential parameter. As a matter of fiact, T_is a function of

Pw and Pr. Substrate temperature thus reflects the energy density of the plasma.

When Pw and Pr is 60 kW and 16.7 kPa (125 Tort), corresponding T_ is 1320 K (1050 °C), Rg and WI_were 7

gm/h and 11 cm4, respectively. In this case, the conversion efficiency from CH4 to diamond on a 152 mm substrate

exceeds 10 %, calculated by the mass increase of the substrate. It is two orders of magnitude higher than that

80 90

10

E 8

:1

_6

e)

n_ 4

(.9

0

10

Pr (Torr)

100 110 120 130

0

[]

L(oc)

600 700 800 900 1000

10

12

10 v

9 j=

7

E

5

18

_" (b)

E 8

r,- 4

xz

o 2

0

0

12

11ff--"

Eo

10= 9

7g

E

6_

5

80

E 8

_6

e) N

IZ 4

£ 2

(.9

(c)

o [] ,,Ill o " | o """

[].... [] , , o , P , ,

12 14 16 20 40 60 900 1000 1100 1200 1300

Gas Pressure Pr (kPa) Microwave Power Pw (kW) Substrate Temperature _ (K)

1100

12

10

8_

7g

E

Figure 1. Dependence of growth rate Rg (square) and Raman signal width WR (full width at half

maximum; circle) of diamond on the gas pressure Pr (a), the microwave input power Pw (b), and the

substrate temperature T_ (c).

NASA/CP--2001-210948 324 August 6, 2001
g1400

1300

_ 1200

o.. 1100

E

I-- 1000

_ 900

800

O9

70O

0.0 1.0 2.0 3.0 4.0 5.0

CH4 Concentration Cm (%)

b

1200

1100

1000

9OO

8O0

7OO

6OO

5OO

4OO

6.0

0

o

v

a

5/Jm

Figure 2. Gray areas show <Ill> and <100> oriented domains determined by X-ray diffraction. Dotted

lines show ct parameter 1.5 and 3. The typical film morphologies observed by SEM are also shown. The

rhombus points indicate the experimental points.

obtained by an ASTeX 1.5 kW system.

it might appear that given the methane concentration C,, and T,, the domains for different film morphologies

would be similar to those of small (i.e., 1.5 kW) reactors (fig. 2). It was however found that the domains of <100>

and <111> oriented growth in the Cm-T,.plane are totally reversed between the 60 kW and small reactors (ref. 2). In

small reactors, <111> oriented domain occupies higher temperature regions than the <100> domain. It is not certain

at the present stage how the crossover of the domain takes place as the gas pressure is increased from about 4.0 kPa

(30 Tort) for small reactors to 13.3 kPa (100 Torr) for the 60 kW reactor.

it is also of interest that the films possess unique morphologies. For the <100> oriented film (fig. 2a), the

corners surrounded by { 111} faces, like a pyramid, are sticking out of the film surface. On the other hand, for the

<111> oriented film (fig. 2b and 2c), the majority of the diamond faces is {111} that seem to be mutually twinned.

The c_parameters were evaluated from the change in the isolated crystal shape with time, and the approximate

curves for c_ 1.5 and 3 are shown in fig. 2. Compared with the results ofref. 1, it is seen that the curves are shifted

toward the higher temperature side. Further, there was indication that when c_> 3, <111> oriented diamond films

were grown, which is in contradiction to the conventional result that the orientation is <100>.

Based on the results, growth orientation and film morphology were manipulated. For example, when the

condition range was CHa/(CH4+H2) 2 %, 60 kW microwave input, 14.0-17.7 kPa (105-133 Tort), and 1170-1280

K (900-1010 °C), the grown diamond grains have a pyramid shape surface and a <100> orientation almost uniformly

on the entire surface of 152 mm substrates (fig. 3). The surface appears like as a black velvet by the naked eye;

lusterless to the normal direction to the substrate, but glossy to the oblique angle. That is originated to that the facets

have almost the same angle -54.7 ° to the substrate plane. The used 152 mm substrates were polycrystalline so that

the orientation behavior is independent of substrate.

NASA/CP--2001-210948 325 August 6, 2001
b

lO/lm

Figure 3. Photograph of the <100> oriented diamond film on a 6 inch diameter substrate (a), and SEM

images at the center (b), the intermediate (c), and the edge.

CONCLUSION

In order to control of morphology of large area diamond films, diamond deposition using the 60 kW microwave

plasma CVD system and characterization with SEM, XRD, and Raman have been carried out. The 60 kW CVD

system shows the growth rate of 7 gm/h and the conversion efficiency of excess 10 % from CH4 to diamond. The

contour of the a parameter shifts toward the higher substrate temperature and the lower CH4 concentration as

compared with the results of the smaller CVD systems.

It was also found that the parameter range of diamond growth by the 60 kW reactor is broader than that of the

conventional CVD systems. Furthermore, <100> oriented diamond grains with pyramid shape were successfully

grown almost uniformly on the entire surface ofa 152 mm Si substrate by choosing an appropriate process condition.

REFERENCES

1. Pan, L. S.; and Kania, D. R.: Diamond: Electronic properties and applications, Kluwer Academic, Boston,

1995.

2. Kamo, M., et al.: United States Patent 4434188 (1984).

3. Naseem, H. A., et al.: W. D. Brown, Thin Solid Films 308-309 (1997) 141.

4. Ralchenko, V., et al.: Diamond Relat. Mater. 8 (1999) 189.

5. Schelz, S.; Campillo, C.; and Moisan, M.: Diamond Relat. Mater. 7 (1998) 1675.

6. Sevillano, E.; and Williams, B. E.: Diamond Films Tech. 8 (1998) 73.

7. Ftiner, M.; Wild, C.; and Koidl, R: Appl. Phys. Lett. 72 (1998) 1149.

8. Tachibana, T, et al.: Diamond Relat. Mater., to be published as a special issue ofProc. 7 th Intl. Conf. New

Diamond Sci. Tech., Hong Kong, July 2000.

NASA/CP--2001-210948 326 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

A HIGHLY EFFECTIVE SETUP FOR DIAMOND COATING DEPOSITION

v. K. Pashnev, O. A. Opalev, I. I. Vyrovets, V. E. Strel'nitskij, V. A. Belous

National Science Centre "Kharkov Institute of Physics & Technology", Akademicheskaya St. 1,

Kharkov, 61108, Ukraine,

Z. I. Kolupaeva,

Kharkov State Polytechnical University, Frunze St. 2, Kharkov, 61002, Ukraine,

Yu. F. Shmal'ko

Institute of Mechanical Engineering Problems, Dm. Pozharskogo St. 2/10,

Kharkov, 61046, Ukraine

ABSTRACT

Active research is in progress on polycrystalline diamond coating (DC) deposition on substrates from various

materials by the CVD method at NSC KIPT Kharkov, Ukraine. For activation of working gas stabilised by a magnetic field, a glow discharge has been used. The characteristics of this type of discharge correspond to the "normal"

glow discharge with a current density greater than 1 A/cm 2. The current channel rotates in a crossed magnetic field

that allows one to carry out the DC deposition over large areas. The experiments were performed on the laboratory

set-up at an input power of about 5 kW, the DC growth rate at this power achieved 3 gm/h on a diameter of about 5

cm. That makes the set-up comparable with the best world results for installations with the same power level. The

mixture of hydrogen and methane was used as a working gas. To remove impurities from hydrogen, a metal-hydride

hydrogen generator was used. According to the experimental dependence discovered by Angus, the set-up productivity increases with the electrical power put into the discharge, this being true for different types of DC installations. In this case, it is possible to predict that at a power of 50 kW the productivity of the set-up under study will be

about 1 g/h. This is competitive with the best existing MW and arc plasma torch set-ups with a power level of about

200...300 kW.

Keywords: diamond coatings, glow discharge, high-efficiency set-up, deposition parameters, coating properties.

INTRODUCTION

By now considerable progress has been made in the field of polycrystalline diamond coating (DC) synthesis by

the CVD method. The investigations have demonstrated a wide applicability of synthesized DC in various fields of

science and engineering. There is a general understanding of processes of DC deposition, but a wide commercial

introduction of this technology is impeded by its relatively high cost. To reduce the DC cost, a highly productive

equipment should be developed to provide DC deposition over large areas with a high energy effectiveness (i. e.,

with minimum power inputs for DC quantity production) and with a low cost of facilities.

The analysis carried out by J. Angus, CWRU (ref. 1) has shown that the productivity of equipment and its energy effectiveness rise with an increasing electrical power put into the discharge for any type of working gas activation. At the same time, the set-up productivity increases with the power increasing proportionally (to the power

1.5). This reduces the influence of the electrical equipment cost on the DC net cost as the production capacity grows.

For working gas activation in the DC synthesis, different types of gas discharge are generally used. The MW

discharge providing a high-quality DC on rather large areas at a relatively high DC growth rate has gained the widest acceptance. The main disadvantage of this method is a high cost of MW elements and a low resource of their

work. Another method of working gas activation is the arc discharge that allows a rather high power to be put into a

small volume, thereby providing a high efficiency of the process at a low cost of energy equipment. However, the

essential disadvantages of the arc discharge technique is the difficulty of DC deposition over large areas and the DC

contamination by a significant amount of metal impurities resulting from the cathode erosion.

There is also a " hot filament" method of DC deposition. The given method permits DC deposition over large

areas. Its main disadvantage is a low power put into the discharge unit volume. For this reason, the hot filament

method fails to provide high efficiency of DC deposition per energy unit applied.

At NSC KIPT Kharkov, Ukraine, the method of DC synthesis with the help of glow discharge stabilised by a

magnetic field was successfully put to an evaluation test (refs. 2 to 5, 10). It was shown that with this method one

NASA/CP--2001-210948 327 August 6, 2001
canattainthespecificpower,putintothedischarge, comparable withthearcdischargeataminimumamount of

impuritiesincomingfromthecathode. Besides, thisdischargemakespossibletheDCdepositionlargeareas.

Theobjectiveofthiscommunication hasbeentoacquaint theexpertsengagedintheDCsynthesiswithspecial

featuresandprospectsofapplicationftheglowdischargestabilisedbyamagneticfieldforDCsynthesisinproductionquantities.

EXPERIMENTALRESULTS

Theinvestigations werecarriedoutontheexperimental set-up,theschematicofwhichisgiveninFig.1.The

peculiarityof thisset-upisthepresenceofthering-shaped anode, insidewhichthesubstrateholdermoves. The

glowdischargeisignitedacrossthemagneticfeldbetweenthecathodeandtheanode. UndertheactionoftheLorenzforcethecurrent channel rotatesandactivatestheworkingasinagreatvolume, thisprovidestheDCsynthesisoveralargearea.Theworkingaspressurerangedbetween50and300Torr;hydrogenanditsmixturewith

methane(upto3 %)wereusedasaworkingas.Themagneticfeldintensitywasabout200mT;thesubstrate

holderdiameter was5cm,themaximal powerenteredintothedischargedidnotexceed7kW.

\

4 / 3 /

Figure. 1. The scheme of the experimental set-up: 1-cathode; 2-anode;

3-annular magnet; 4-substrate holder; 5- current channel.

The current-voltage characteristics of the glow discharge stabilised by the magnetic field are given in Fig. 2. It is

seen from the figure that at the given working gas pressure the voltage Vp between the anode and the cathode does

not depend on the current value I, i.e., by the existing terminology, the "normal" glow discharge is observed. The

distinctive feature of this discharge should be a constancy of current density j, so the increase in the current value

occurs due to the growth of the current channel area.

For power supply of the set-up we have used a three-phase rectifier without special filters to smooth the ripple.

Therefore, pulsation in the voltage and discharge current, intricate in shape (non-sinusoidal), with a basic frequency

ff

of about 300 Hz, was observed. The voltage pulsation is about __ _ 2%, and the current ripple achieves

U

_< 30%. As it has been shown in reference 3, in the /_ I __ << discharges all current fluctuations go with the

I U I

current density held the same due to the current channel variation.

The dependence of current and voltage values in the discharge on pressure is shown in Fig. 3. The power supply

allowed the current adjustment over wide limits. Fig. 3 shows the minimal and maximal current values together with

the corresponding voltage values. It can be seen that with an increase in the working gas pressure P the currents

linearly decrease, while the voltage linearly increases. If we introduce the following notation: j - average current

density in the current channel cross-section, S - current channel area, tJ - average conductivity, and E - electrical

NASA/CP--2001-210948 328 August 6, 2001
1100

1000

>

0

>

800

5

4

3

S

7_0 I I I

4 5 6 7

Current, A

2

1

Figure. 2. Volt-ampere characteristics of glow

discharge in hydrogen. 1-P=50 Torr; 2-P--60

Torr; 3-P--80 Torr; 4-P--100 Torr; 5-P=120 Torr.

I,A

7

U_, B

3

50

I in ax

U_ "%-,,,%,

! ! !

100 150 200 P, 'i'opp

1200

lOOO

800

600

Figure. 3. Minimal (I .,i.) and maximal (Imax) discharge current and voltage between the cathode

and the anode vs. pressure in the discharge

chamber (P).

I

ll__pG and I /-/--pry where 1 is the

field strength, then from I = jS and j = rYE we obtain: j =__ j- g S

length of the current channel. So different functional dependences I=I (p) and Up=U (p) are possible only in the case

if the current density remains the same with pressure variations in the range from 50 to 250 Torr.

Spectral and probe measurements (ref. 2) have allowed us to determine the size of the current channel and to estimate the current density in the hydrogen discharge to be about j_ 1.2 A/cm 2. Such a high current density makes the

power characteristics of the magnetic filed-stabilised glow discharge closer to the arc discharge characteristics.

It has also been indicated in ref. 2 that on the substrate, electrically insulated from the anode and the cathode,

during the discharge there arises a negative potential of up to 400 V relative to the anode. Further experiments (refs.

3, 4) have suggested the conclusion that apart from the basic discharge (between the cathode and the anode), the

cascade discharge is observed in the circuit: the cathode - substrate along the magnetic field and the substrate - anode across the magnetic field. In the hydrogen discharge, the share of current participating in the cascade discharge

does not exceed 20 %, but the addition of methane to the working gas can cause, in some cases, a rise of current in

the cascade discharge up to 60 %. As demonstrated in reference 3, the substrate potential shows a linear dependence

on pressure, similar to the voltage dependence of the main discharge.

Fig. 4 presents the power put into the discharge and the substrate temperature versus the working gas pressure

for the minimal and maximal currents (ref. 3). It is evident that the values of pressure for the maximum of power put

into the discharge and the substrate temperature are not always coincident. According to reference 2, this points to

the effect of ion bombardment on the substrate heating when the substrate is at negative potential. As it was indicated in reference 4, about 10 % of the power input goes for heating the substrate. A half out of this value is lost for

radiation, and the rest is taken by the cooling system.

The discharge voltage and the substrate potential have appeared appreciably dependent on the methane concentration. It was shown in reference 3 that this dependence has a linear character. This fact is advantageously used for

the glow-discharge diagnostics and the methane concentration measurements at transient processes occurring during

DC deposition.

Investigations into optical radiation of the magnetic field-stabilised glow discharge were carried out (ref. 5). It

was shown that the radiation spectrum in the range from 230 to 700 nm was essentially in agreement with the data

obtained by other authors who used various working gas- activation methods in the set-ups (refs. 6,7). The differences were only observed in the relative intensity of hydrogen Balmier lines of the series Ha, H_,Hv, Hs, He and the

molecular lines H2 (462.8; 581.3; 602.8 rim), and also in the line distribution in the molecular bands of C2 (swan

band) and CH radicals. It was also shown that the highest concentration of C2 and CH radicals could be attained near

the substrate with the help of special engineering solutions. The linear dependence of the radiation intensity of C2

NASA/CP--2001-210948 329 August6,2001
T_,oH.ca. W,KBT

4O

3O

2O

10 i i i 3

50 100 150 200 P, Topp

Figure. 4. Power (W) input into the discharge and

the substrate temperature (T) vs working gas

pressure (P). The curves are built for the maximal and minimal current values.

I, oH. ca.

D

o - Up /_

[] - C2 ,_

, - CH _/_" .--_-_'-

J

jj e j-'_--

_::51 ........

I ¢ I I I

0 1 2 MeTaH, %

Up, B

1150

1100

1050

Figure. 5. Emission line intensities of C2 (516.5

nm) and CH (431.4 nm) radicals and the discharge gap voltage Up vs methane concentration

in the discharge.

and CH radicals on the methane concentration in the hydrogen discharge (see Fig. 5) was observed, that could also

be used for the discharge diagnostics.

The optical radiation spectra of the glow discharge stabilised by the magnetic field were studied to check for the

presence of metal impurities. In particular, much attention was given to the presence of Mo impurity in the discharge, because the cathode and the substrate holder were made from Mo. Fortunately, the optical spectrum exhibited no Mo lines. At the same time, the installation of a copper specimen in the substrate holder gave rise to Cu I

lines (324.7 nm and 327.4 nm) with an excitation potential of about 3.8 eV.

The optical measurements have made it possible to determine the frequency of current channel rotation in the

glow discharge stabilised by the magnetic field (ref. 2). It has been found that the rotational frequency changes as a

function of discharge parameters in the range from 100 up to 200 Hz and has different functional dependences: f oc 7---_' U, .

for the hydrogen discharge; f _: "I (PH__DAF +PA_/_ for the hydrogen discharge with the

Ar addition; and f _ (P--_H_ DH_ _ 4[I - AIR ] for the discharges with the methane addition. Here PH2, PAr, PCH4 are the

partial pressures of hydrogen, argon and methane, respectively; and the factor A depends on the geometrical

parameters of the discharge. The experiments made with the use of single and double electrical probes have fully

confirmed the data of optical measurements.

The thermocouple measurements of the working gas temperature and the results obtained with the help of electrical probes have given the basic parameters of the magnetic field-stabilised glow discharge at pressure P _ 150 torr

(ref. 2). It is shown that the highest gas temperature in the current channel reaches 5000 °C, the average gas

temperature in the region of current channel motion is about 2000 °C, the electron density ne_ 1017m -3, time of

electron collisions t_3.10 -10 sec, and the parameter m% *e_l, where WBeis the electron cyclotron frequency.

An attempt was made to derive the analytical solution for the frequency of current channel rotation at the discharge, using a set of equations of gas dynamics and magnetic hydrodynamics. It is shown that in the case of an extremely simplified model it is possible to obtain the analytical solution to the given problem (ref. 4). It has appeared

that the speed of current channel rotation can be presented as the balance between the speed of rotation under the

action of Lorenz force, with being hindered by the neutral gas, and the recombination rate of plasma particles. The

process is noticeably influenced by radial plasma diffusion.

In the magnetic field-stabilised glow discharge, the DC growth rate was investigated as a function of various parameters (ref. 4). It is found that a rather flat maximum is observed on the DC deposition rate vs substrate temperaNASA/CP--2001-210948 330 August6,2001
turecurve.Theoptimumsubstratetemperature valueliesbetween980and1010°Candisdependent oncertaindischargeparameters, inparticular, onthepurityofhydrogenintroducedinthedischarge. Thehighest DCgrowthrate

achieved3 gm/honadiameter of5cm.TheDCgrowthratewasfoundtoincreasewithagrowingflowrateof

methaneatthegivenrateofhydrogenflow;however, nochangeintheDCgrowthratewasobservedwiththevariationsinthehydrogenflowratefrom230cm3/minupto450CM3/minataconstant methaneflowrate.Inotherwords,

theDCgrowthrateisdependent onthemethaneflowrateinthehydrogendischargeratherthanonthemethane

concentration. Themethaneconcentration essentiallyinfluencesthequalityof coatingdeposited; thebestresults

wereobtainedatamethaneconcentration ofabout1%.

Theexperiments haveshownthenecessityofusingahigh-purityhydrogenforDCsynthesisinthedischarge

(ref.5).Therefore, forpurificationofindustrial hydrogenthemetal-hydride generator ofhydrogen(MHGH),developedbytheInstituteofMechanical EngineeringProblemsofNASofUkraine,Kharkov, wasused(ref.8).Inthe

MHGHofthegivendesign, about2801ofpurehydrogencanbestored. ThepurityofhydrogenleavingtheMHGH

afterspecial purificationprocedures canattain99.99%.

TheuseofMHGH-purified hydrogenhasnoticeablyinfluencedtheprocessofDCsynthesis. Firstofall,theDC

growthratehasincreased(abouttwice,insomecases); thisbeing,inouropinion,duetobothahighpurityofhydrogensuppliedanditsthermoemission activationatdesorptionfromhydride-generating intermetalliccompounds

(refs.8,9).Then,withtheuseofMHGHhydrogen, theprobabilityfortheglowdischargetoturnintothearcdischargeessentially decreases, andthisindirectlyconfirmsthepresenceoftheearlierestablished (ref.9)conditionsof

dischargeautostabilization inpressure. Forthesamereasons, theamount ofthecathodematerial (Mo)dropletsin

thedepositedDCsampleconsiderably decreases.

StudiesintodiamondnucleationandtheDCstructuresynthesisedinthemagneticfield-stabilised glowdischargewerecarriedout(ref.10).Itisfoundthatatinitialstagesofdiamondcrystalformationtheso-calledeffectof

decoration, associated withthediamondcrystalnucleationatthegrain-boundary substratecarbides(alwayspresent

intherealsubstrate), manifestsitself.

Theintermediate carbidelayerandtheDCformeduringthesynthesiswereinvestigated withthemethodsof

X-raystructureanalysis(ref.10).Itisestablished thathecarbidelayerhastheMo2CorW2Cstructure, depending

onwhetherthesubstratematerial MoorWwasused.FortheMosubstrate, theintermediate carbidelayerthickness

makes7...8.5gmataDCthicknessofmorethan40gm.

Thelatticeconstantsforthediamondcoatingandthecarbidesubstratewerecalculated. Theresultsallowusto

suggest theabsenceofanappreciable amount ofhydrogeninthesynthesised DCanditspresenceinthecarbidesubstrate. Itisalsoshownthathediamondcoatingsareundertheactionofcompressive stressofabout1.5GPa,while

thecarbidesubstrateisundertheactionofstretchingstress. Thestresses intheDCdependonlyweaklyontheDC

thickness, andthisallowsthesynthesisofDCwithathicknessupto1mm.

Asit isseenfromtheabovexperimental dta,themagneticfeld-stabilised glowdischargeusedforworking

gasactivationenablesonetoattainasufficientlyhighDCgrowthrate.Thecomparison ofperformance characteristicsoftheabove-mentioned set-upwiththetechnical parameters ofotherset-upsforDCsynthesis, describedinreference1,allowsustoassert thatbytheproductivityofDCsynthesisandbytheefficiencyofenergyusagepercarat

ofthesynthesised DCthepresent methodofworkingasactivationisinonegroupwiththebestworldresultsobtainedwiththeuseoftheMW/arcmethodsofworkingasactivation, andthehotfilamentmethodatthesame

powerlevel(5kW)putintothedischarge.

Atpresent, 20kWand50kWset-upsarebeingdesignedatNSCKIPT,thatwillprovide, throughvariousmodificationsofthemagneticfeld-stabilised discharge, the0.25g/hand> 1g/hproductivities ofDCsynthesis, respectively.Thecostofthediamond-coated carbideinsertcanbereduceddownto$0.2perpiece.Thediameter ofthe

substrateholderforthe50kWset-upcanreachupto25cmandover,dependingontherangeofproductsobe

coatedbyDC.

CONCLUSIONS

Theinvestigations madeonthelaboratoryset-uphaveshownthatheglowdischargestabilisedbyamagnetic

fieldhasseveral essential advantages incomparison withother,widelyusedCVDmethodsoftheDCsynthesis.

Thecurrent channel rotationinthemagneticfeldmakesitpossibletousethegiventypeofworkingasactivationforDCdepositionverlargeareas. Thebasicharacteristics ofthedischargeinthepressurerangefrom50to

300tortandatcurrentsfrom3to8Acorrespond totheclassical _normab>glowdischargewithacurrentdensity

NASA/CP--2001-210948 331 August6,2001
higherthanI A/cm2.Thiscurrentdensityprovidesahighspecificpower(higherthan200W/cm3inthecurrent

channel) neededfortheDCdepositionoverlargeareas.Despitesuchahighspecificpowerinthedischarge, the

opticalspectrumexhibitsnoemissionlinesofmaterials, fromwhichthemostelectricallyoverloadedset-upelements,inparticular, thecathode, weremade.ThispointstothepossibilityofdepositingDCofhighpurity.

Thepresenceofthenegativepotential upto-400Vacrossthesubstrateholder(electricallyinsulatedfromthe

anodeandcathode) hasbeendemonstrated experimentally. Thepotentialrisesasaresultofthecascadeischarge

inthecathode- substrate- anodecircuitandinfluencesthesubstrateheating. Theinvestigations havenotrevealed

anynegativeinfluenceofthispotential ontheDCgrowthrateandquality.TheDCgrowthrateonadiameter of5

cmachieved3gm/h,thisbeingcomparable withthebestworlddatafortheset-upswithapowerinputof about5

kWintothedischarge.

Theabsenceofimpuritiesin hydrogenandtheenergeticstateofhydrogenatomsandmoleculesxertsanappreciablyfavourableeffectontheDCqualityandgrowthrate.Inourexperiments, ametalhydridegenerator was

usedforhydrogenpurificationtoprovidehydrogenof99.99%purityinamountsnecessary forDCsynthesis. It has

resultedinatwo-foldincreaseoftheset-upproductivity.

Diamondcoatings, 40gmandoverthick,onMoandWsubstrates wereproduced. Betweenthesubstrates and

theDCthickcarbideinterlayerswereformed. ThosewereMo2C,upto8.5gmthick,andW2C,>4gmthick,forthe

MoandWsubstrates, respectively. Thediamondcoatingisundertheactionofcompressive stressof 1.5GPa,and

thecarbidelayerisundertheactionof stretchingstresses. TheDCstressvaluedependsonlyweaklyontheDC

thickness, andthispermitsthesynthesisofdiamondcoatingsupto1mmandoverinthickness.

REFERENCES

1. J.E. Butler, H. Windischmann. Developments in CVD Diamond Synthesis During the Past Decade, MRS Bulletin, September 1998, p. 22.

2. V.K. Pashnev., O. A. Opalev, V. A. Belous, V. E. Strel'nitskij, Some features of direct current glow discharge

in a transversal magnetic field for diamond films deposition.//ISDF - 4, 1999. P. 18.

3. V.K. Pashnev, O. A. Opalev, I. I. Vyrovets, V. E. Strel'nitskij, I. K. Koval'chuk, V. A. Belous, Glow discharge

for diamond coating deposition over a large area. 12 m International Symposium "Thin Films in Electronics ",

April, 2001,. Kharkov, Ukraine (in Russian).

4. O.A. Opalev, V. K. Pashnev, I. K. Koval'chuk, V. E. Strel'nitskij, V. A. Belous. Diamond coating synthesis in

the glow discharge stabilised by a magnetic field. VANT, series " Physics of radiation damages and materials

science", 4 (78) 2000, P. 158-164. (in Russian).

5. V.K. Pashnev, O. A. Opalev, I. I. Vyrovets, V. I. Grytsyna, V. E. Strel'nitskij, V. A. Belous, V. G. Konovalov,

A. N. Shapoval, Yu. F. Shmal'ko, A. I. Ivanovskiy. Spectral research of the glow discharge stabilized by a magnetic field. 12 th International Symposium "Thin Films in Electronics", April, 2001,. Kharkov, Ukraine (in Russian).

6. W. Zhu, A. Unspektor, A. R. Badzian, T. Mckenna and R. Messiez. Effects of noble gases on diamond deposition from methane-hydrogen microwave plasmas. J. Appl. Phys. 68 (4), 15 August 1990.

7. R.R. Stalder and R. L. Sharpless. Plasma properties of a hydrocarbon arc jet used in the plasma deposition of

diamond thin films. J. Appl. Phys. 68 (12) 15 December 1990.

8. Yu. F.Shmal'ko, V.V.Solovey and M.V.Lototsky. Use of Metal Hydrides in Systems for Supplying Vacuum

Physical-Energy Installations. / In.: Hydrogen Energy Progress X. Proc. 10-th World Hydrogen Energy Conf.,

Cocoa Beach, Florida, USA, 20-24 June 1994.- Ed. by. D.L.Block and T.N.Veziroglu, Int. Association for Hydrogen Energy, 1994, vol.2, p. 1311-1319.

9. V.N.Borisko, Ye. V.Klochko, M.V.Lototsky, A.A.Petrushenya and Yu. F.Shmal'ko. Pressure auto-stabilised gas

supply systems for vacuum-plasma installations using reversible hydrogen getters.- VTT (Vacuum Engineering

and Technology), 2000, v. 10, No. 1, p. 15-19 (in Russian).

10. V. K. Pashnev, V. E. Strel'nitskij, O. A. Opalev, I. I. Vyrovets, V. I. Grytsyna, V.V. Bryk, K. A. Malyava, A.

A. Parkhomenko, I. M. Neklyudov, Z. I. Kolupaeva. Structural characteristics of diamond coatings obtained in

the glow discharge stabilized by a magnetic field. 12 th International Symposium "Thin Films in Electronics ",

April 2001,. Kharkov, Ukraine (in Russian).

NASA/CP--2001-210948 332 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

High growth rate and high quality CVD diamond growth

H.-G. Jentsch, S. M. Rosiwal, R. F. Singer

Friedrich-Alexander UniversiNt Erlangen Ntirnberg, Institut t'tir Werkstoffwissenschaften,

Lehrstuhl Far Werkstoftkunde und Technologie der Metalle

Corresponding author: Dipl.-Ing. Hans-Georg Jentsch

Institut Far Werkstoffwissenschaften, Lehrstuhl II

Friedrich-Alexander Universit_t Erlangen-Ntimberg

Martensstrage 5

D-91058 Erlangen, Germany

Tel.: + 49 - 9131 - 8527520 Fax: + 49 - 9131 - 8527515

email: jentsch@ww.uni-erlangen.de

H. Eibisch

Neue Materialien Fttrth GmbH

Abstract

CVD diamond is the most promising material for thermal management due to its very high thermal

conductivity. Nevertheless wide spread use of CVD diamond in heat spreading applications has not been

successful, due to the high price of the material, One major issue is the slow growth rate (typically 1 gm/h)

which yields the appropriate diamond quality for heat spreading applications in low pressure CVD

processes. For effective heat spreading of e.g. laser diode bars a necessary thickness of the diamond heat

spreader is calculated to be 100 gm. Increasing the growth rate without loosing diamond quality has not

been successful yet. But high thermal conductivity can only be achieved when superior phase purity and

crystallinity is provided in the diamond coating. We demonstrate that it is possible to grow CVD diamond

with 8 gm/h with high phase purity. We utilize the etching capabilities of reactive species in the micro

wave plasma witch contains carbon, hydrogen and oxygen. We found that the usage of high oxygen ratios

in the plasma leads to a better selectivity of the etching species in the plasma with respect to sp 3 and sp2

carbon.

In order to achieve high diamond quality it is most necessary to avoid graphitic inclusions. This is usually

achieved with g small growth rate and high atomic hydrogen concentrations in order to selectively etch sp2

hybridized carbon from the growing diamond surface. Enhancing the selective etching process would

therefore allow to increase the growth rate without diminishing the diamond quality. Kinetic calculations of

hydrogen, methane and oxygen mixtures show that the formation of carbon monoxide is slower than the

formation of hydroxyl radicals or water. Hydroxyl radicals are therefore likely to have an influence on the

growth of CVD-diamond. We measured the growth rate online and determined the activation energy of

diamond growth. This activation energy is raised from standard 1% methane 99% hydrogen mixtures with

50 kJ/mol to 85 kJ/mol when diamond is grown from gas mixtures with higher concentrations of methane

and oxygen. Detailed confocal micro raman analysis were performed to characterize the phase purity of the

deposited diamond films. Initially the phase purity drops with higher growth rate, as has been frequently

reported. A further increase of methane and oxygen leads to an increase in phase purity while the growth

rate is maintained at a high level.

Keywords: thermal management, roman spectroscopy, growth mechanism, oxygen

NASA/CP--2001-210948 333 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

The Rapid Growth of Thin Transparent Films of Diamond

Hsiao-Kuo Chung and James C. Sung

KINIK Company, 64 Chung San Road, Ying-Kuo Taipei Hsien 239 Taiwan.

E-mail : cgwrd@kinik.com

Abstract

Thin transparent films of diamond have only been grown by CVD methods at very slow rate. Using hot

filament CVD reactor of large area (30 cm by 40 cm), such films were deposited onto highly polished copper

substrate at a rate about l gin/hr. The cost of manufacturing thin transparent films of 20lam thick may be as

low as $ 1/cm 2. At such a low cost, thin transparent films should have diversified applications, such as X-ray

windows, high frequency (up to 10 GHz) SAW filters, and other advanced devices.

Key Words: CVD Diamond, Optical Window, SAW Filter.

NASA/CP--2001-210948 334 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Diamond particles synthesized on glassy and oxide substrates

Yoshiki TAKAGI & Fumitomo ONISHI

Teikyo University of Science & Technology

2525 Uenohara-cho, Kitatsuru-gun, Yamanashi-pref., JAPAN

.t_a_k___g i__ n__U.____.j_.

tel: +81-554-63-4411, fax: +81-554-63-4431

ABSTRACT

The conventional techniques for depositing diamond at low pressure utilize a flow of a hydrocarbon-hydrogen gas

mixture with complicated gas tubing for introducing the reaction gas into a reaction chamber and evacuating reactant

gas from it. For few years, we had been concentrated on developing completely closed diamond

synthesizing system aimed for high gravity or microgravity conditions. With our preliminary experiments, we

successfully confirmed that diamond was synthesized with this completely closed system, and that this system was

very suitable for mounting on large centrifuges and in spacecraft.

In our study, we used graphite rod for heater and carbon source, at the same time. Diamond particles were

synthesized on Si substrates with graphite heating on our completely closed reaction chamber ( method A ). We tried

other substrates, glassy materials ( silica glass and other glasses ) and ceramics ( aluminum oxide ).

We also tried to synthesize diamond with conventional hot filament CVD method with hydrogen and methane gas

mixture ( method C ) and graphite heating on hydrogen flow experiments ( method B ).Synthesized particles were

observed with SEM and confirmed as diamond with Raman spectroscopy.

With 60 minutes experimental time on Silica glass substrates, we obtained 1.5 to 3.0 micron sized diamond with

method B, 3.0 to 4.0 with method A and over 20.0 micron big particles with method C. On Aluminum oxide

( alumina ) substrates, 5.0 microns with method A and B, and over 30.0 microns with method C in 60 minutes. With

Raman spectroscopy, method B gave us the best diamond ( smallest FWHM valued Raman peaks attributed to

diamond particles ) compared with other two methods. We used no bias voltage, substrate pretreatment nor plasma.

Discussions for carbon source effects ( solid graphite and gaseous methane ) and flow effects ( with source gas

flow system and completely closed system ), and more details of experimental methods and results including SEM

photographs and Raman spectra will be shown in the conference.

Keywords: closed system, silica glass, aluminum oxide, graphite heating, diamond.

INTRODUCTION

With diamond synthesis, crystalline substrates, such as Si and Mo are usually used. In this study, we used oxide

amorphous for example, quartz glass, and crystalline oxide such as aluminum oxide for substrates. In last year, on

MRS Spring meeting, we reported the synthesized diamond on quartz glass[l]. In this study, we tried synthesis with

conventional flow system and with methane and hydrogen mixture gas.

NASA/CP--2001-210948 335 August 6, 2001
EXPERIMENTAL

In this study, we tried three deferent methods listed on Table 1. For each method, schematic figures are shown in Fig.

1.

Method A is so called completely closed system. After hydrogen was introduced on the chamber with suitable

pressure, all valves were closed, and graphite rod was heated up with electricity. To keep suitable temperature of

graphite rod and substrate, which was heated with strong radiation come from graphite rod, diamond particles were

synthesized on the substrates.

Method B is flow system with graphite rod, which installed on the center of the chamber. Reaction gas (hydrogen)

was introduced and reactant gas was evacuated to keep suitable pressure.

Table 1 Experimental method

A B C

gas flow closed flow flow

system system system

carbon graphite graphite OH4

source rod rod gas

Method C is common and conventional thermal filament method. For reaction gas, mixture of hydrogen and

methane was used. For filament, spiral shaped tungsten wire was used.

Experimental conditions were listed on Table 2. The synthesized particles were shown in Fig. 2, for Method A,

Fig.3 (method B) and Fig.4 (C), with SEM photographs and Raman spectra.

AC power atl_ly

reaction _-_

gas iifie?_

!1

reaclant

gas o111 et

alve

AC tx_er attty

gas

_o*_ II

T :II

gasortla

AC tz__r attly

ga_otllet

A B

Figure 1. Schematic figure of apparatus for method A, B and C.

C

NASA/CP--2001-210948 336 August 6, 2001
Experimental methods

subs'Eates

Initial pressure / Torr

Total pressure / Torr

Flow rate (H z:CH4) / SCCM

Pyrogen temperature/

Substrate temperature / *(3

Deposition lime/min.

Distance;substrate -

pyrogen / mm

Sample number

SiO2

glass

2000

Table 2 Ex

A

Alumina

50

x

)erimental conditions

B

SiO2 Alumina

glass

x

100

C

SiO2 Alumina

glass

X

100

x 50 : 0 99 : 1

I 2300 2100 2300

750 750

6O

3

3 4

60

700~750

60 60 I 60 120

=

3

1 2 5 6 I 7 8

RESULTS AND DISCUSSION

Figure 2 shows SEM photographs and Raman spectra for particles synthesized by method A on silica glass (left) and

on alumina (right). The particle sizes were about 4 microns on silica glass and about 6 microns on alumina. The

particles were confirmed as diamond with characteristic diamond peaks on Raman spectra.

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::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

........ ooo°°o°oooooo°oo°ooo°o°o°,

,°°%°°°°°°°°°°°°°°°°°°°°°°°°%°°°°°°%°°°°°°°°°°°°°°°°°°°°°°°°°°°,

,:°:°:°1.:°:°:°:o:°:°:o:°:°:°:°:°:o:°:°:°:°:°:o:o:°:°:°:°:°:°:°:o:,

,.°°°°*°°.%°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°°,

:::::::::::::::::::::5::::::::::::::::::::::::::::::::::::::::::::l_m

:.:.:.:.:.:.;.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:

xlO000

50

30

2O

10

0 I I I I

1800 1600 1400 1200 1000 800

on Alumina (Sample number 2)

_'_¢.-:.::::;..;::::::::::::::::::::::::::::::::::::::::::::::

1,111111".'1".111" "t o . ili.. t i%_ lffillil_ll :.'-:."".:.:.:.:.:.:.._.-i.'-" :,:. ",'+ •

.:..: .'.:.:.:.:.:.>..:.:;'..'::..:.:.:°:.:.:.:..,:.: .:.:.:.:.:.:::,

i:i:i:i:!:i.i:!_:!<:!:!:_:i:??!:i:!:i:_:i:i:i:;_i:_:i:_,:i

::::::::::::::::::::::::::::::::::::::::::::::::::::::::

i

5_tm

x5000

Raman s hift/cm -_

Figure 2 SEM photos and Raxnan spectra (method A)

Figure 3 shows SEM photos and Raman spectra for the particles synthesized with method B on silica glass (left) and

on alumina (right). The particle sizes were about 2 and 4 microns each. The fivefold shaped particles were observed

on alumina, which axe typical on CVD diamonds. Raman spectra showed diamond peaks on both samples.

NASA/CP--2001-210948 337 August 6, 2001
onSiO_ _lass (Sample number 3)

_i_,,_._}_._.i,_ _,;,;.;.;.;4,;,,.;..;,'-.'-.;".- " ,',',,,_,-', ,",'_,,,',',_

1

l_m

x20000

on Alumina (Sample number 4)

.................................

:.:.:.-.:.:.: .-x...:.:.:.:.:.:.:.-.:.:.:.:.:.:.:.:.:..--, x.:.:.. B...., .....................:_.g:::::..............×.....::.ll

:.:.:.:.:.:.:.:.:.:.:.,._..:.'.'.'.'...... ":::.x:.:."

•:.:.:.:.:.:.:.:.,:::..:.x." "::..".. ":.:.::x::

:::::::::::::::::::::::::::"" ..!:!:::...'8."

::::::::::::::::::::::::::::::::::::::::

!:_:i$i.'.:!'_'i'i:i'i:i'i'ig.'.. iS;i:!:

]

1400

5O

3O

2O

i0

i

1800 1600 1200 i000 800 1800 1600 1400 1200 i000 800

Raman s hift/cm -_

Figure 3 SEM photos and Raman spectra ( method B )

Figure 4 shows SEM photographs for particles synthesized by method C on silica glass (left top) and on alumina.

With method C, tungsten wire was used as a filament. In all cases except for sample 6, a filament was gradually

expanded and finally touched with the substrates.

+_.on SiO 2 glass_ 41

(Sample number 5) (Sample number 6)

iiiiiii!iiiiiiiiiilililil

i!iiiiiii!ii!ii!'.i'.':!'-i'.:.

i i

15/_m x1500 5/_m x5000

( 60 INn ) (60min)

on Alumina •

(Sample number 7) (Sample number 8)

:: !:i!iiiii:%iii:!:ii?ii:!:i:i:!:i!!i{:!:i<: !.:;;_i__:jO.: "._._'._.'_'<:

i i

15/.tm x1500 15/.tm x1500

(60iNn) (120iNn )

Figure 4 SEM photos with method C experiments

Figure 4 shows that all particles obtained by the experiments with filament touch to the substrates, gives very large

particles such as 1 micron.

Figure 5 shows Raman spectra for each sample. With all Rmnan spectra, diamond peaks were confirmed.

NASA/CP--2001-210948 338 August 6, 2001
(Sample number 5)

25

20

15

10

5

0

1800 1600 1400 1200 i000 800

(Sample number 6) (Sample number 7) (Sample number 8)

25 50

20 40

15 30

i0 20

5 5 i0

0 0 0

1800 1600 1400 1200 i000 800 1800 1600 1400 1200 i000 800 1800 1600 1400 1200 i000 800

Raman shift/cm q

Figure 5 Rmnan spectra for method C

Figure 6 shows deposited particles were grown very well and touch each other.

Figure 6 SEM photograph for sample 8

SUMMARY

We confirmed diamond on all methods A, B and C. With Raman spectra analysis, method B gives the best

crystalline diamonds.

With method C, in which methane was used as carbon source, synthesized diamond was the lmcgest ones among all

three methods.

With these results,

Graphite was very effective for synthesize good crystalline diamonds.

Methane was very useful for high crystal growth rate, and to synthesize lmcge diamond pmcticles.

REFERENCE

1. Rie Hayashi, Fumitomo Onishi, Yoshiki Takagi, Mat. Res. Soc. Syrup. Proc. Vol. 616

pp.195-197 (2000).

NASA/CP--2001-210948 339 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

High quality diamond growth on iron and steel substrates

Eri NAKAMURA, Manabu OHYAMA, and Kenji K. HIRAKURI

e-mail, fax number, and phone number of corresponding author

e-mail: manabu@ff.f.dendai.ac.jp

fax: +81-492-96-6413

phone: +81-492-96-2911 ext. 3134

Department of Electronic and Computer Engineering, Faculty of Science & Engineering,

Tokyo Denki University

Ishizaka, Hatoyama, Saitama 350-0394, Japan

Nobuki MUTSUKURA, and Yoshio MACHI

Department of Electronic Engineering, Faculty of Engineering, Tokyo Denki University

2-2 Chiyoda-ku, Kandanishiki-cho, Tokyo 101-8457, Japan

Gernot FRIEDBACHER

Institute of Analytical Chemistry, Vienna University of Technology

Getreidemarkt 9/151, A- 1060, Wien, Austria

Abstract

It is difficult to produce diamond grains on iron-based materials due to the catalytic effect and the rapid diffusion

coefficients of carbon species into the iron-based materials, such as iron and stainless steel (SUS) containing

chromium (Cr) 18 % and nickel (Ni) 8%. However diamond growth on iron-based materials is extremely important

for mechanical and electrical applications because the iron-based materials are widely used in domestic science.

In our previous study, the diamond nucleation and subsequent growth are precisely controlled by residence time

of source gas as an essential parameter. We have carried out the diamond growth on iron-based materials using the

hot-filament chemical vapor deposition (CVD) technique with varying the residence time. At lower residence time,

diamond grains with practical growth rate are grown on the iron-based materials. The growth rate of diamond grains

on SUS substrates was a factor of about 7 greater than that for a normal silicon substrate at optimum condition.

Furthermore, the growth rate exponentially increases with the growth time. Diamond film of about 0.1 mm

thickness was obtained at growth time of 10 hours.

Raman spectroscopy and scanning electron spectroscopy were utilized for estimation of the quality for grown

diamond grains. The FWHM of diamonds on iron-based materials is comparable to it of natural diamonds.

Moreover diamond grains were also observed not only at the surface but also at the bottom of the SUS substrates.

Key words : diamond growth & characterization, heated filament, chemical vapor deposition,

Raman spectra

1. Introduction

Substrate surface treatment such as mechanical polishing, plasma etching, ion implantation, and thin film

deposition, have brought various influences on the nucleation of diamond and subsequent growth 1'2). The conditions

of the substrate surface are very important in the growth of the diamond nuclei 3). Furthermore, the formation of SiC

layer on the Si substrate surface is observed at the initial stage of diamond nucleation. Iron thin films on Si substrate

enhanced diamond nucleation density by high diffusion co efficiency against carbon. It is very difficult to produce

the high quality diamond growth on iron and steel substrates 4).

In this study, diamond deposition on iron and steel substrates have been attempted using the residence time of

source gas as essential parameter of diamond growth. The significant enhancement in the growth rate of diamond and

the improvement of crystalline quality were obtained with remarkably difference in comparison with the diamond on

normal Si substrate.

NASA/CP--2001-210948 340 August 6, 2001
2. Experiments

The hot-filament chemical vapor deposition (CVD) was used for diamond growth. The double chamber of

stainless steel was used for diamond growth. The substrate and filament temperatures were controlled at 850 and

2100 °C, respectively. The CH4 concentration was kept at 0.5 % during diamond growth and total gas pressure was 5

kPa. The depositions were carried out on normal Si, pure iron (99.99 %), and 18-8 stainless steel (including 18 %

chrome and 8 % nickel) substrates.

The morphology of the deposits was studied by means of scanning electron microscopy (SEM) images obtained

with a JEOL JSM-5310LVB system. Raman spectroscopy (JASCO NRS-2100) with an Ar ion laser was also used to

identify the crystalline quality of the grown diamonds.

3. Results and discussion

Figure 1 shows the growth rate of grown diamond at different residence time. The growth rates of all samples

increase with decrease of the residence time. The maximum growth rate of stainless steel substrate was a factor of

about 7 greater than that of normal Si substrate.

8

j=__7 --e-si substrate

---41.-- SUS substrate

5 _k- Fe substrate

,= 3

0 I I

0.001 0.01 0.1 1

Residence time [sec]

Figure 1. Relationship between growth rate and residence time.

SEM micrographs of diamond grains grown on normal Si, iron, and stainless substrates under maximum growth

rate are shown in Fig.2. At stainless steel substrate, single crystals without any facets can be observed (Fig. 2 (b). In

contrast, twin structures were obtained on the normal Si substrate (Fig. 2(a)).

(a) Si substrate (b) Stainless steel substrate (c) Iron substrate

Figure 2. SEM images.

In order to compare the crystalline quality of grown diamonds at different substrate, micro Raman spectroscopy

was utilized. Figure 3 shows Raman spectra of diamond grains prepared on each substrate. The grains were identified

as diamond due to the sharp peak at 1332 cm -1. The full width half maximum (FWHM) of the peaks was used to

characterize the quality of diamond grains. Typical average FWHMs of the diamond peak are listed in Fig. 3 (as

referred natural diamond). The FWHM of diamond on stainless steel substrate is comparable to that of natural

NASA/CP--2001-210948 341 August 6, 2001
diamond. Bycontroltheresidencetimeoftotalgas,thediamondswithfinecrystallinequalitywereproduceon

stainless steel.

F.W.H.M. Si Substrate

3.2 crn'l_ SUS Substrate

7,2co __FeSubst te

3.8 cm'l I Natural Diamond

i i i i

1200 1300 1400 1500 1600 1700

Raman shift [cm -1]

Figure 3. Raman spectra.

4. Conclusions

Diamond growth on stainless steel and iron substrate with convenient growth rate and quality has been realized

using residence time as one of the growth condition. The growth rate wa found to increase by a factor of 10 when the

residence time was kept at 5 ms. The effect of the increased growth rate may be because of the increased amount of

atomic hydrogen and diamond precursors.

References

1) K. Kobayashi, T. Nakano, N. Muts_a, and Y. Machi, Vacuum, 44 1 (1993).

2) C. P. Chang, D. L. Flamm, D. E. Ibbotson, and J. A. Mucha, J. Appl. Phys., 63, 1744 (1988).

3) D. N. Belton, S. J. Harris, S. J. Schmieg, A. M. Weiner, and T. A. Pewy, Appl. Phys., 54, 416 (1989).

4) T. Sato, S. Narumi, S. Ito, and K. Akashi, Thin Solid Films, 316, 29 (1998).

NASA/CP--2001-210948 342 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

The Uniformity of 8 Inch CVD Diamond Films

Sheng-Yi Lee, Shtao-Kuo Chang, Jow-Lay Hung, & James C, Sung

KINIK Company, Taipei, Taiwan, R.O.C.

Abstract

Diamond films of 8 inch in diameters were deposited by CVD using hot filaments. These films were uniform in

thickness and other properties. Such diamond films are suitable for many industrial applications including cutting

tools, heat spreaders, SAW filters, and boron doped electrodes.

Key Words: CVD Diamond, Diamond Films, Hot Filament

NASA/CP--2001-210948 343 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Synthesis of diamond under high pressure using hot-filament CVD method

Eri NAKAMURA, Toshihiro KOBAYASHI, and Kenji K. HIRAKURI

e-mail, fax number, and phone number of corresponding author

e-mail: eri @ ff.f.dendai.ac.jp

fax: +81-492-96-6413

phone: +81-492-96-2911 ext. 3134

Department of Electronic and Computer Engineering, Faculty of Science & Engineering,

Tokyo Denki University

Ishizaka, Hatoyama, Saitama 350-0394, Japan

Nobuki MUTSUKURA, and Yoshio MACHI

Department of Electronic Engineering, Faculty of Engineering, Tokyo Denki University

2-2 Chiyoda-ku, Kandanishiki-cho, Tokyo 101-8457, Japan

Gernot FRIEDBACHER

Abstract

Institute of Analytical Chemistry, Vienna University of Technology

Getreidemarkt 9/151, A-1060, Wien, Austria

In previous work, we have shown that it is possible to achieve diamond growth by hot-filament chemical vapor

deposition (CVD) under atmospheric pressure and that the residence time is the most important parameter for

diamond deposition. Furthermore, the relationship between the quality of deposited diamond and the residence time

has been also investigated by Raman spectroscopy. The results have shown that growth rate and quality of diamond

are controlled by the residence time.

In this work we have tried to increase the growth rate by deposition under a pressure of 200 kPa. Moreover, the

structure of the substrate surface after deposition was characterized with Auger electron spectroscopy (AES).

At 200 kPa a significant increase of the growth rate compared to conventional pressure (5 kPa) and atmospheric

pressure (106 kPa) was observed, since the optimum residence time of the precursor is abruptly increasing with the

deposition pressure. In addition to that, Raman analysis revealed that the quality of diamonds grown at 200 kPa is

better than that of diamonds produced at low pressures. Furthermore, high quality diamond growth was achieved

with a long residence time of the input gas at high pressure. Under the applied deposition conditions a peak

originating from silicon carbide could be observed after a growth time of 15 minutes.

Key words : diamond growth, diamond characterization, heated filament, chemical vapor

deposition, grain size

NASA/CP--2001-210948 344 August 6, 2001
1 Introduction

Chemical vapordeposition(CVD)ofdiamondfilmshasreceivedconsiderable attentionduetotheattractivefilm

propertiessuchashighthermal conductivity, mechanical hrdness, widebandgap,andchemical inertness 1'2).If rapid

growthandimprovedqualityof CVDdiamondscanberealizedthentheyareexpectedtobeutilizedwidely

industriallyinthefieldofelectronics, mechanics, andoptics.

Uptonowmostresearchers commonlyusepressure, flowrate,gasratios,substratebias,powerandsubstrate

temperature tocharacterize thegrowthconditions. Astheseparameters aenottheonlycrucialdepositionparameters

comparison ofresultsachievedwithexperimental set-upsandcommercial systemsisnormallyveryrequired. Thus,

another standardize ableparameter isrequired.

Inourpreviouswork,theresidencetimewasproposedasawellcontrollableparameter foCVDdiamondgrowth

andtheresultsontwodifferentsystemswereingoodagreement wihtheproposedparameter 1).A significant

enhancement inthegrowthrateof CVDdiamondwasobtainedunderatmospheric pressure1).Furthermore, XRD

analysisrevealedthatthequalityof diamondsgrownatatmospheric pressurewasbetterthanthatof diamonds

producedatlowpressure. Inordertoincreasegrowthrate,wehaveattempteddiamondgrowthat200kPa.

1. Experimental procedures

A conventional hotfilamentCVDsystemhasbeenusedfordiamondgrowthunderstandardconditionsas

describedpreviously 2).Thesystemisequippedwithaccessories forhighpressureprotectionandthermal safety.The

vacuumchamber madeofstainlessteelcontainsamolybdenumsubstrateholderwithathermocouple andatungsten

filament.Thevacuumchamberwasevacuatedbyarotarypumpto abackgroundpressureof 0.1Paafterthe

substrates weremountedonthesubstrateholder.Thehydrogengaswasintroducedintothechamber throughamass

flowcontroller atatmospheric pressure. Thefilamentwasheatedupto2370Kbypassinganelectriccurrent through

it andthemethanegaswasmixedwithhydrogenafterastablefilamenttemperature wasreached. Thesubstrate

temperature waskeptat1120Kwhenthesubstrates werelocatedatadistanceof5mmfromthefilament.

Thediamondgrowthwascontrolledbytheresidencetimetr [sec] of the precursor which is given by

PxV

t r = k s --

Q

where P is the gas pressure [Pa], Q the flow rate [sccm], V the volume of the deposition chamber [cm3], and kr

the constant value of 10.51) . In order to investigate the influence of the methane concentration on the quality of the

grown diamond, it was varied from 0.2 %.

The surface morphology of the grown diamond was investigated using scanning electron microscopy (SEM :

JEOL JSM-5310LV, resolution = 4 nm). Micro Raman spectroscopy (JASCO : NSR-2100) was used to identify the

quality of the diamond films grown. The growth rate was estimated by the size of the diamond grains and the growth

time.

2. Results and discussion

Figure 1 shows the growth rate of diamond at different residence time. At low pressure conditions (5 kPa, and

50 kPa), the maximum growth rate was achieved at a residence time of 0.05 s. It is striking to note that under low

pressure the dependency follow similar curves, indicating that diamond growth cannot be characterized by pressure

and flow rate alone. Thus, the residence time is suggested to be used as a standardized parameter for CVD diamond

growth.

NASA/CP--2001-210948 345 August 6, 2001
5

.-e--5kPa

4 ----Am 50kPa

___ --_-- lOOkPa /l_Vl ""

3

2

i,

0 v0.01 0.1 1 10

Residence time [sec]

Figure 1. Growth rate of CVD diamond grains produced

different pressure as a function of the residence time.

at

With increasing pressure up to atmospheric pressure and above, the growth rate dramatically increases with

residence time. The maximum growth rate at 200 kPa was about 4 times higher than that at low pressure conditions.

The reason for that increase is discussed in a previous paper 2_. The effect of the increased growth rate is because of

the increased amount of atomic hydrogen and diamond precursors on the substrate surface resulting from a prolonged

residence time.

The maximum growth rate of diamonds was strongly depended upon the residence time. Then they were

picked up and traced the tendency. Figure 2 shows the maximum growth rate as a function of the residence time. It

can be seen that the maximum growth rate increases monotonously with the residence time up to 2.3 s. The possibility

of increase of the growth rate was found from this Figure.

5

=, 4

93

S

S

S

I I

1 2

Residence time [sec]

0 3

Figure 2. Maximum growth rate of CVD diamond

grains under different pressure condition.

In order to compare the quality of the diamonds grown at different conditions, micro Raman spectroscopy was

applied. Figure 2 shows Raman spectra of diamond grains produced at pressure condition between 5 kPa and 200 kPa.

NASA/CP--2001-210948 346 August 6, 2001
Ineachcasethegrainswereidentifiedasdiamondduetothesharpeakat1332cm -1. Significant peaks at 1350 cm -1

(D band) and 1530 cm -1 (G band) originating from disordered and glassy amorphous carbon around the grains can be

observed in the Raman spectra of the diamond at 5 kPa 3_. On the other hand, D band and D band did not appear at the

spectrum of diamond grown at 200 kPa.

This result is similar to the dependence observed for the growth rate as a function of methane concentration as

shown in Fig. 3. At low methane concentrations, comparative ratio between atomic hydrogen and methyl radical

increases with decreasing methane concentration. Thus, non-diamond species can be etched oft', which is also

supported by the SEM micrographs in Fig. 3.

' "-.-.,,__...

.= 1_ 200 kPa

i I i I

1200 1300 1400 1500 1600 1700

Raman shift [cm 1]

Fig. 3: Raman spectra of CVD diamond grains grown

at different pressure condition.

3. Conclusions

Using the hot-filament CVD technique, the influence of the gas pressure on diamond growth has been

investigated. A dramatically increase of diamond growth rate could be observed with gas pressure. The significant

enhancement of the growth rate at high pressure allows production of high quality diamond.

Acknowledgements

The authors would like to thank Mr. T. Toya for performing SEM analysis. This work has been partially

supported by the Center for Research of Tokyo Denki University.

References

1. T. Kobayashi, K. K. Hirakuri, N. Mutsukura, and Y. Machi, Diamond and Related Materials, 8 (1999)

1057-1060.

2. K. K. Hirakuri, M. Yoshii, G. Friedbacher, and M. Grasserbauer, Diamond and Related Materials,

6(1997) 1031-1035.

3. S.M. Leeds, T. J. Davis, E W. May, C. D. O. Pickard, and M. N. R. Ashfold, Diamond and Related

Materials, 7(1998)233-237.

NASA/CP--2001-210948 347 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

DIAMOND SYNTHESIS BY M1CROX_,AVE AND DC HYBRID PLASMA CVD

Keiji Fuji, Taku Sumitomo, 2Xkimitsu Hatta

DepartmEnt of Electronic and Photonic Systems Engineering, Kochi University o f Techn_lob-'3'

185 Miyanokuchi, Tosayamada, Kami-glm, Kochi 782-8502, Japan

e-mail:ahatta_elekochi-techac.jp FAX: +81-887-57-2120 TEL;+81-S87-57-2]13

ABSTRACT

For production of tlfin and uniform disk shaped high density plasma for diamond grm_lh at higher pressurE,

mlcrowave and DC hybrid d,schalge is proposed.

It has been succeeded that high quality diamond is synthesized by the microwave plasma chemical _,apor

deposition (CVD) method Crystalline diamond can be slowly grown form source gases of a several kPa in presser

fltmugh chemical reaction enhmlced by the plasma. One of the diffictdties [i_r industri',d applicaUon of CVD

diamond is high production cost due to file smalt grox_1h rate.

II is expected that increasing the reaction pressure will increase the growth rate. For production of an uniform

plas_r_a at higher pressure over the substrate size of a few inches, it is necessary to increase the microwave power

and to use an welt designed microwave cavity. In a cavity for the microwave discharge, the plasma of a large volume

is produced by such the huge microwave poweI, as shown in Fig (a). 1[ the plasma can be produced only around the

substrate, as shown in Fig (b), the microwave power can be reduced.

Here, WE propose a hybrid discharge by using microwave and DC powers to produce the plasma only the

substrate. This method _s expected that the Iota! power fol plasma produciion by using both rite m_crowave ancl DC

will be smaller than dmt for die production of large-x'olume plasma by microwave alone.

The experimental apparatus was a convenlional micIowave plasma CVD appmalus with two electrodes inserted

The experimental condition was at 13kPa of It: with 420W micro_vave and 70W DC, or wilh 500W microa'ave

alone.

When it was discharged with microwave and DC, a thin plasma disk abou_ hum thick 16mm ut :limneter was

produced on the surface of a substrate. When it was discharged with microwave only, plasma was coniracled at the

corner of mtbstrale.

tt has been reported that file negative voltage of the substame E1dmnced diamond _mcleation bat obslructed

c_'stal growth. At higher pressure, however, the energy will be lowered by collision. It is expected that the ion

bombardment at lowered energy will enhance the co'_al gro_vlh rather 1t_1 rite ob_acle.

[ ]

Sobs [rate

(a) (b)

Fig (a) Plasma ball and (b) plasma disk.

NASA/CP--2001-210948 348 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CHEMICAL PROCESSES IN ECR-MICROWAVE PLASMAS CAPABLE OF DIAMOND

FILM DEPOSITION

Rik Blumenthal

Department o f Chemistry, Auburn University, Auburn AL 36 8 4 9

ABSTRACT

The chemical compositions of 4% carbon in hydrogen/deuterium electron cyclotron resonance (ECR)

microwave plasmas based on ethane, ethylene, acetylene and methane have been investigated using Supersonic

Pulse, Plasma Sampling Mass Spectrometry. Kinetic modeling based on the simulation of the plasma as a steadystate source of hydrogen atoms using 54 reversible chemical reactions and their isotopic analogs has successfully

modeled the composition of all of the plasmas studied. A unique set of parameters was found that generated

simulated spectra that matched all of the experimentally measured spectra. The success of this modeling effort

provides strong evidence that the chemistry of ECR-microwave plasmas is dominated by neutral molecule reactions.

Keywords: Microwave, Diamond, Deposition, Hydrocarbon, Plasma.

INTRODUCTION

Diamond films and diamond-like coatings (DLCs) have found a number of practical applications, including

scratch resistant films on sunglasses, thermally conductive films for cooling of semiconductor devices, and as

mechanical strength enhancers on razor blades. From the time of the initial discovery of low-temperature diamond

deposition, the enormous potential benefits of diamond films has driven extensive research into both the practical

and fundamental aspects of the process. In a practical sense, significant progress has been made in understanding the

relationships between the deposition conditions and the morphology and physical properties of the film. An

amazingly wide array of techniques have been demonstrated to be capable of depositing diamond films, including

flame torches, hot filament reactors, microwave and high-density plasmas. In contrast to the physical properties of

the deposited film, which are characterized ex-situ or after deposition is completed, and despite significant efforts,

the chemical phenomena responsible for diamond film deposition have not been understood to an extent even closely

approaching the understanding of the morphological issues.

In this work, we report both the results of supersonic pulse, plasma sampling mass spectrometric measurements

of 4% carbon in hydrogen/deuterium electron cyclotron resonance (ECR) microwave plasmas based on four different

hydrocarbons, methane, acetylene, ethylene and ethane, along with the results of kinetic modeling of the plasmas

based on a simplified model of the plasma that support prior interpretations made solely on the basis of the mass

spectrometric results. Supersonic pulse, plasma sampling mass spectroscopy has been shown to be capable of

determining the chemical composition of a plasma, including the observation of reactive species and radicals that are

not observed using other methodologies, such as the high percent dissociation of C12 during the etching of silicon in a

high-density plasmas. Kinetic modeling of diamond film deposition is also well-established, and has been previously

applied to hot filament reactors, plasma torches and microwave plasmas.(refs. 1-3) Unfortunately, the results of

those kinetic modeling studies could only be tested by comparison with secondary observations, such as film

deposition rates, or with the of concentration of a single molecular species within the complex environment.

In prior kinetic modeling of plasma deposition, a majority of the calculational effort was directed to the solution

of transport properties of the gas. In contrast, we have decided to return to the 1-D method employed in the hotfilament reactor simulations and focus our calculational efforts on simulation of the chemical composition based on

the time evolution of the feed gas mixture under the influence a steady-state concentration of atomic hydrogen

(simulating the plasma) using 750 forward reactions (54 reversible reactions resulting in 375 reversible variations

when isotope effects are taken into account) taken from the combustion literature. The resulting kinetic modeling

suffers from two significant deficiencies. The first deficiency of our model is our indirect treatment of the electronmolecule reactions that generate the reactive species that are ultimately responsible for diamond growth. In our

model, the plasma is replaced by a steady-state concentration of atomic hydrogen. In diamond deposition plasmas

with >90% H2 in the feed gas, it does not seem unreasonable to assert that the majority of the electron impact

NASA/CP--2001-210948 349 August 6, 2001
chemistrywillresultinthegenerationofatomichydrogen. Hence, thisfundamental assumption ofourmodel is

similarlynotunreasonable. Theseconddeficiencyofourmodel isthecompletenglect ofwallreactions. No

experimental evidenceexiststosupport suchasignificant assumption, andourchemical intuitionwouldleadusto

believethatneglectingwallreactionsisapoorassumption. However, thequalityoftheresultsofthemodel,the

abilitytofitthetotalcompositions ofeightplasmas(oneforeachfeedgasinbothhydrogenanddeuterium) provides

strongevidencethatwallreactionsmusteitherplayamuchmoreminorrolethanintuitiondictates, orthatinthe

steady-state ofheplasma, oneormoreofthegasphasereactionsincludedinthemodelsimplyreplacetheactual

wallreactionsoftheplasma.

Overall, thechemical compositions of4%carbonihydrogen/deuterium ECR-microwave plasmasbasedon

methane, acetylene, ethyleneandethanehavebeenmeasuredandsimulatedusingasimplifiedmodel inwhichthe

plasmaisreplacedwithasteady-state concentration ofatomichydrogen. Thequalityoftheagreement beweenof

themodelandexperiment overanunprecedented rangeofconditionsandexperimental detailsprovidesstrong

evidencethathemajorreactionpathwayinECRmicrowaveplasmasisneutral reactions.

EXPERIMENTAL

All ofthediamonddepositionplasmasweregenerated usinganASTeXmodelAX4300compact ECRmicrowaveplasmasource.Typical operatingconditionswereatotalpressureof-0.5Pa,aforwardmicrowave

powerof200W.Theconfinement magnet wasoperatedat-16.5Aandashapingmagnet, located-40cm

downstreamoftheplasmasourcewasadjustedtoyieldcylindrical plasma.Feedgasesweremixedthroughingas

manifoldconstructed ofthreegasflowtubestoobtainthedesiredratioofthegases.Gasflowswereadjusted

manuallytoobtainthedesiredratiosof-10:1H2:CH4and-22:1ratiofortheotherthreehydrocarbons.

Supersonic pulse,plasmasamplingmassspectrometry

Supersonic pulse,plasmasamplingmasspectrometry hasbeendescribedindetailpreviously.(ref. 4) Briefly,

400uspulsesofargonarereleasedattheedgeoftheplasmaregion.Thegaspulsesexpandthroughteplasma

region,coolingandultimatelycondensingtoclusters. Themolecular andatomicspeciespresent intheplasma

regionactascondensation nucleifortheclusteringprocess, resultingintheplasma-born speciesbeingincorporated

intothecenteroflargeclustersofargon.Theclustersthenprotectthespeciesfromfurtherreactionuntiltheyreach

theExtrelQ50quadrupole massspectrometer. Themasspectrometer wasoperatedat30eVelectronimpact

ionizationenergyinsinglemassmode.Masspectrawereaccumulated byadvancingthemassby0.1AMUperAr

pulseatapulserepetitionrateof10Hz.All dataareaverages of100massscansandsevenpointsmoothed.

Kineticmodelingofplasmacomposition

Thefundamental propositionfthesimulationisthereplacement oftheplasmawith a steady state concentration

of hydrogen atoms. The justification for this simplistic model is found in the fundamental physics of ECR

microwave plasmas. In theses plasmas, the input microwave energy is coupled to the plasma by direct absorption by

electrons in cyclotron orbits. Chemical excitation of the gas is realized through the collision of energetic electrons

with the molecular species of the gas. Since the vast majority of the gas is hydrogen (initially 96 or 98% molecular

hydrogen), it is reasonable to assume the majority the chemical excitation is achieved through the fragmentation of

H2 to yield atomic hydrogen, hence the assumptions of the model used in the simulations is reasonable.

Kinetic modeling of the composition of the plasmas was based on re-calculating the composition at fixed time

steps based on changes in the composition calculated by the Arrhenius rate equations. The rate constants, barriers

and free energies used in the Arrhenius calculations were taken from the literature. A total of 375 reversible neutral

species reactions, including isotope effects, were derived from 54 principle reactions, shown in Table 1. Where

possible, rate constants were taken entirely from Harris et al. Additional reactions were taken from the GRI-Mech

database. All reactions parameters were re-tabulated using the smaller of the forward and reverse barriers calculated

from the published parameters using the free energies in the individual publications. For the purposes of this work,

the larger of the barriers was recalculated using free energies from the JANAFF tables. This unusual approach was

required to prevent slight differences in the free energies used in different sources from creating artificial negative

energy barriers ,which disrupted the mass balance of the calculation.

Each simulation was begun with the input of a total gas pressure, fraction of hydrocarbon and a temperature, and

was repeated for successively increasing hydrogen gas dissociation percentages of 10%. Within each calculation, the

gas composition was started at the specified feed gas composition with fixed hydrogen dissociation percentage. The

gas composition was allowed to evolve in time for a specified period until the calculation was terminated and a

NASA/CP--2001-210948 350 August 6, 2001
simulatedmassspectrumwascalculatedbasedonthefinalcomposition ofthegas.Theterminationtimewasvaried

aroundtheexperimentally determinedresidence timeofagaspulseintheplasmachamber, 200ms.Intheinitial

stagesofthesearchforasinglesetofinputparameters thatcouldsimulateallspectra, thetimestepforthe

Arrheniuscalculations was50gs,andduringthelatterstagesofrefinement, thetimestepwasreducedtolgs.Mass

spectrawereconstructed forcomparison texperimental data,usingtheexperimentally determined masspectrafor

thestablenon-deuterated hydrocarbons (measured intheabsenceoftheplasma)ndisotopicbranchingratiostaken

fromstandardreferencespectra.

RESULTSANDDISCUSSION

Theexperimentally measured massspectraareshowniFig.1,andthebest-fitresultsofthemodelingare

showniFig.2. Theinputparameters ofthekineticmodeling, seeabove, wereallvariedoverawiderangeandthe

resultswerevisuallycomparedtotheexperimental dataovertheC2HxDvmassrange,m/z22-36.Onlyonesetof

parameters wasfoundthatcouldeffectivelyreproducetheexperimental resultsforalleightplasmassimultaneously.

Theresulting"bestfit"wasfoundtobe40%dissociation ofhydrogengas,atemperature of490K,atotalpressureof

106Paanda200ms.Notvariedinthecalculationwasthepercent hydrocarbon, whichwasfixedat4%forCH4and

2%forC2H2, C2H4, C2H6.Thecalculationprovedtobeextremelysensitivetothetemperature, witha30Kincrease

resultinginacalculatedC2H2inD2patternwith the middle of the three peaks (m/z 27) dominant and a decrease of

30K resulting in a C2H2 in D2 pattern with only a single peak at m/z 28. Sensitivity to the percent dissociation of the

hydrogen was slightly less than the dependence of temperature. For instance, at 20% dissociation of H2, the C2H2 in

D2 pattern inverted (m/z 26 > m/z 27 > m/z 28), while at 60% dissociation of H2, the middle peak (m/z 27) of the

C2H2 in D2 pattern became drastically reduced as was the m/z 28 peak of the C2H4 in H2 pattern. Pressure also

played a strong role in the calculation. At a pressure of 10.6 Pa, the C2H2 in D2 pattern displayed only a single peak

at m/z 26, and no variation of other parameters could be found to correct it to the three peak pattern of the

experimental data. At increased pressure the m/z 27 peak of the C2H4 in D2 pattern was significantly reduced and

the m/z 26 peak was enhanced. It should be noted that what often appeared to be a positive trend in the variation of

pattern, as observed in compared to a single experiment, almost always resulted in an equally bad trend when

compared to another spectrum. The only best-fit parameter that did not closely agree with the experimental value

was the pressure, which was approximately two orders of magnitude higher than we measured. The only

explanations for discrepancy lie in the remote location of our ion gauge, which is both obstructed from the plasma

region and closer to the pump, and the low sensitivity of the gauge to H2, which results in a low pressure reading.

In conclusion, we have measured the chemical compositions of a number of ECR-microwave plasmas capable of

depositing diamond and compared the results to a simplified kinetic model in which the plasma is replaced with a

steady-state concentration of atomic hydrogen and only neutral molecule chemistry is considered. The results of the

calculations agree with prior models (ref. 5) that the primary reaction pathway is the stripping of hydrogen/deuterium

from the C2Hx to generate acetylene, the only hydrocarbon species that undergoes significant hydrogen addition

reactions. Coupling of the C2Hx reaction manifold to the C1Hx reaction manifold is through the reaction of the

ethane radical and a hydrogen atom to make two methyl radicals, while coupling of the C1Hx reaction manifold to

the C2Hx reaction manifold is through the three-body recombination of methyl radicals to form ethane. The quality

of the agreement between of the model and eight experimental mass spectra provides very strong evidence that the

maj or reaction pathways in ECR microwave plasmas are neutral-neutral reaction presented previously.

REFERENCES

1. Harris, S.J.; and Weiner, A.M.: Methyl radical and H-atom concentrations during diamond growth. J. Appl.

Phys. 67(10), 1990, 6520-6526.

2. Coltrin, M.E.; and Dandy, D.S.: Analysis of diamond growth in subatmospheric dc plasma-gun reactors. J.

Appl. Phys. 47(9), 1993, 5803-5820.

3. Hyman, E. et al. : One-point numerical modeling of microwave plasma chemical vapor deposition diamond

deposition reactors. J. Vac. Sci. Technol. A, 12(4), 1994, 1474-1479

4. Gaddy et al., : Supersonic pulse, plasma sampling mass spectrometry: theory and practice. Plasma Chem.

Plasma Proc., 19(4), 1999, 513-544.

5. Webb, S.F. et al. : Investigation of ECR-microwave plasmas using ethane as the hydrocarbon source gas. J.

Vac. Sci. Technol. A, 17(5), 1999, 2456-2462.

NASA/CP--2001-210948 351 August 6, 2001
Table 1. Reactions and Arrhenious Rate Parameters

Reaction A

C2H6 + H _ C2H5 + H2 5.4 x 102

k = A (T) B e-Ea/RT

B

3.5

E_ (J)

21,725

C2H5 + H _ C2H 4 ÷ H2 8 xl012 0

C2H 4 + H _ C2H 3 ÷ H2 1.5 xl0 TM 42,614

2 xl013

6 x 1013

C2H 3 + H _ C2H2 + H2

C2H 2 + H -+ C2H + H 2 99,015

CH 4 ÷ H _ CH 3 ÷ H 2 2.2 x 104 3 36,765

CH3 ÷ H _ CH2 ÷ H2 7.2 x 10 TM 63,086

CH2+H --* CH+H2 4 x 1013

C2H 6 + CH3-+C2H 5 + CH 4

5.5 x 10 -1 34,676

C2H 5 ÷ CH3--*C2H 4 ÷ CH 4 7.9 x 1011 0

C2H 4 ÷ CH3--*C2H 3 ÷ CH 4 2.27 x 105 2 38,511

C2H 3 ÷ CH3-+C2H 2 ÷ CH 4 7.9 x 1011 0

C2H 3 ÷ H +M----_C2H 4 +M 2.6 x 1017 403,584 + AGr_I

C2H2+ H +M---_C2H3+M 4.2 x 1041 -7.5 94,838 + AGr_I

C2H+H+M---_C2H 2 +M 4 x 1016 44,703 + AG_

4x 105

8 X 1026

1 x 1016

5.4 x 102

C2+H+M _C2H+M 2.4

-3

3.5

CH3+H+M _CH4+M

CH2+H+M _CH3+M

C2H 5 + H2-+C2H6+ H

4,178

370,321 + AG_,_

21,725 + AG_,_

C2H 4 +H2----_C2H 5 +H 8 xl012 Aarn _

C2H 3 +H2----_C2H 4 +H 1.5 xl0 TM 42,614 + AG_I

C2H 2 +H2-+C2H 3 +H 2 xl013 Aarn _

C2H+H2--+C2H2 +U 6 x 1013 99,015 + AGr_

CH 3 + H2---_CH4+ H 2.2 x 104 3 36,765 + AG_

CH2 + H2_CH3 + H 7.2 x 10 TM 63,086 + AG_

CH+H2 --* CH2+H 4 x 1013

C2H4+ H-+C2H5 5 x 1015 -2.4

AGr>al

153,328 + AG_,_

C2H2÷ H 2 ÷ M _ C2H 4 ÷ M 2.6 x 1017 240,646 + AG_,_

C2H 5 ÷ H _ CH 3 ÷ CH 3 2.8 x 1013 0

C2H 4 + H -+ CH 3 + CH 2 2 x 1013 AGrxa_

C2H 4 + H _ CH 4 + CH 6 x 1013

C2H 3 x 1013 3 + H _ CH 3 + CH

AGr>a_

AGr>al

C2H 6 -+ CH 3 ÷ CH 3 9.4 x 1045 -9 420,713

5 X 1018

3.4 X 1041

2.6 X 1017

C2H5-+C2H4 + H

C2H 6 + M _ CH 3 + CH 3 + M

C2H 4 +M-+C2H 2 + H 2 +M

C2H4+M-+C2H 3 + H +M 2.6 x 1017

-2.4

-7.03

153,328

11,543 + AGr_

240,646

403,584

C2H 3 +M ---_C2H 2 +H+M 4.2 x 1041 -7.5 94,838

C2H 2 +M---+C2H+H+M 4 x 1016 44,703

CzH +M--+C 2 + H +M 4 x 10 s 2.4 4,178 + AG_

CH4+M _CH3+H+M 8 X 1026 -3

CH3+M _CH2+H+M 1 x 1016

AGr>a_

370,322

CH 3 ÷ CH 3 ÷ M _ C2H 6 ÷ M 3.4 x 1041 -7.03 11,543

CH3+CH3-+C2H 6 9.4 x 1045

2.8 x 1013

-9

CH 3 + CH 3 -+ C2H 5 + H

420,713 + AG_,_

AGr>al

CH 3 + CH 3 --* CH 4 + CH 2 1 x 1013 AGrxa_

CH 3 + CH 2 _ C2H 4 + H 2 x 1013 0

CH + CH 4 -+ C2H 4 ÷ H 6 x 1013 0

CH + CH 3 _ C2H 3 + H 3 x 1013

C2H 5 + CH4-+C2H 6 + CH 3 5.5 x 10 -1 34,676 + AG_,_

C2H 4 + CH 4 --* C2H 5 + CH 3 7.9 x 1011 AGrxa_

2.27 x 10 s 2

7.9 x 1011

C2H 3 ÷ CH4--+C2H 4 ÷ CH 3

C2H 2 + CH4-+C2H 3 + CH 3

38,511 + AG_

AGr>al

CH 4 + CH 2 --* CH 3 + CH 3 1 x 1013 0

NASA/CP--2001-210948 352 August 6, 2001
1.o ] Acetylene in Hydrogen 1.o Acetyelene in Deuterium

_0.4_ _ _ 00.4 j O0 ''1' '1''1 '' I''l' '1 '' I' '1''1' '_''1 '_ 0.0 " ' I- , ........................ _,_,"T',

12 14 16 18 20 22 24 26 28 30 32 34 36

rn/z

1.0 1.0

1.0

Ethylene inHydrogen

0.8--

•_->"0.6 -

o_

t-"

-.-- 0.4-

0.2-

0.0

Ethane in Hydrogen

''l''l''l''l''l''l''l''l''l''l''l''

12 14 16 18 20 22 24 26 28 30 32 34 36

rn/z

12 14 16 18 20 22 24 26 28 30 32 34 36

rn/z

0.8

._0.6

I-"

_O 0.4

¢-

0.2

0.0

1.0

Ethylene in Deuterium

_, 2k ,_

''1 ''1''1''1''1''1''1''1''1''1''

12 14 16 18 20 22 24 26 28 30 32 34 36

m/z

0.8

Ethane in Deuterium

1.0 1.0

0.8 t Methane in Hydrogen ]0.8 ethane in Deuterium

06 -_ 0.6

0.4 _ 0.4

0.0 i i

o2

0.0 i"l"l"

12 14 16 18 20 22 24 26 28 30 32 34 36 12 14 16 18 20 22 24 26 28 30 32 34 36

m/z m/z

Figure 1. Supersonic pulse, plasma sampling mass spectra of 4% carbon in hydrogen and 4% carbon in

deuterium ECR-Microwave plasmas using four different hydrocarbons as feed gases.

NASA/CP--2001-210948 353 August 6, 2001
1.0

0.8

---_0.6

C

-.-- 0.4

C

0.2

0.0

Acetylene in Hydrogen

II__

I I-I I I I I I I I I I I I I I I I I I I I I I

12 14 16 18 20 22 24 26 28 30 32 34 36

m/z

0.8

Ethylene in Hydrogen

0.6

tt)

c 0.4

0.2

 I,I l_

IIIIIIIIIIIIIIIIIIIIIIII

12 14 16 18 20 22 24 26 28 30 32 34 36

m/z

0.3

>, 0.2 -

t/l

C

121

o.1-

Ethane in Hydrogen

.m.I.

illlllllllllllllllll

12 14 16 18 20 22 24 26 28 30 32 34 36

m/z

0.7

06 t

0.5

_0.4

Acetyelene in Deuterium

c 0.3t II

0.2

0.1

0 I I I I I--I I I I I I I I I'1 I I I--I--I--I I I I

12 14 16 18 20 22 24 26 28 30 32 34 36

m/z

0.5

0.4--

2o.3_

C

0.2-

C

0.1-

0.4

>,°3 t

0.2

Ethylene in Deuterium

.....Ill

IIIII IIII I II II- I I I IIII IIII

12 14 16 18 20 22 24 26 28 30 32 34 36

m/z

Ethane in Deuterium

0.

12 14 16 18 20 22 24 26 28 30 32 34 36

m/z

• _ Methane in Deuterium ]

I IIII I I

m c: 0

0

12 14 16 18 20 22 24 26 28 30 32 34 36 12 14 16 18 20 22 24 26 28 30 32 34 36

m/z m/z

Figure 2. Simulated mass spectra of 4% carbon in hydrogen/deuterium for four different hydrocarbon feed

gases assuming the conditions: 10.6Pa, 490K, 200ms residence time and 40% dissociation of the

hydrogen/deuterium.

NASA/CP--2001-210948 354 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

EXPERIMENTAL CONFIRMATION OF POSIBILITY OF SYNTHESIS OF

DIAMOND SINGLE CRYSTALS SIMILAR TO NATURAL

T.A.Nachalnay, V.G.Malogolovets, G.A.Podzyarey, S.A.Ivahknenko, O.A.Zanevskiy,

Institute for Superhard Materials of NAS of Ukraine, Kyiv, Ukraine

I.A. Novoselova

Institute of General & Inorganic Chemistry of NAS of Ukraine, Kyiv, Ukraine

ABSTRACT

Using EPR and IR spectroscopy we have studied crystals grown in nonconventional systems, prepared

by temperature-gradient growing on a seed crystal, and synthesized from hydrocarbons. We discuss special

features spectroscopic characteristics, which demonstrate the feasibility of synthesizing diamond crystals which

do not differ from natural diamonds in defect-impurity composition (and thus physico-mechanical properties).

Keywords: spectroscopy, diamond, temperature gradient, nitrogen, hydrocarbons

INTRODUCTION

Among aggregate nitrogen impurity centers typical of natural diamonds, the A center is the most

common one. Diamond crystals containing impurity nitrogen solely in the form of the A center axe rare in

occurrence. Although the A center is usually accompanied by other aggregate nitrogen impurity centers - B 1

and B2 (like the A center they axe non-paxaxnegnetic centers), it is dominant. The occurrence of natural

diamonds which axe placed into the category of nitrogen-containing crystals (type Ia), exceeds 95%.

It is appropriate to emphasize here that the concentration of aggregate impurity nitrogen in highquality natural diamonds (including gem-quality crystals) can be as high as 102o cm 4. Absolutely opposite to

natural type Ia diamonds axe impurity-free crystals classified as type IIa. Such crystals are quite rare (their

occurrence is less than 5%. These crystals have the highest thermal conductivity (ref. 1), which is the reason of

an increased interest they attract. Commercial synthesis of diamonds yields crystals, which axe classified as type

Ib according to spectroscopic characteristics. The impurity nitrogen in these crystals is solely in the form of

single isolated substitution atoms (the paxaxnagnetic nitrogen C center) whose coming into the lattice obeys the

random distribution law (as distinct from ingress of impurity nitrogen into natural diamonds). These crystals axe

somewhat inferior to natural diamond in some physico-mechanical and performance characteristics. Therefore,

the attempts of scientists to synthesize diamond close to type Ia and IIa natural diamonds in impurity-defect

composition seem to be justified. The impurity nitrogen, being an indicator of the diamond crystal growth

conditions, taxes an active part in the formation of strength, optical, thermal, and electrophysical properties of

these crystals. The difference in properties between synthetic and natural diamond crystals depends greatly on

the structural state of impurity nitrogen in them. It is known that during high pressure-high temperature action

on synthetic diamonds in their thermodynamic stability region the impurity paxaxnagnetic nitrogen transforms to

non-paxaxnagnetic A and Blnitrogen centers typical of natural diamonds (ref. 2).

An alternative way of synthesizing diamond crystals, which would be close to natural ones in

impurity-defect composition, is to provide conditions facilitating the formation of such crystals directly during

the synthesis. The prerequisites for successful implementation of this alternative include the possibility of

changing the impurity composition of the crystals by varying the synthesis conditions, the presence of variously

aggregated paraxnagnetic nitrogen complexes in statically synthesized diamond and yielding large diamond

crystals by the growth on a seed. Given below axe the results of investigation of synthetic diamond crystals and

diamond materials by these methods along some routes where, in our opinion, significant advances have been

made in synthesizing diamond crystals close to type Ia and IIa natural diamond in defect-impurity composition.

EXPERIMENT AND SAMPLES

All investigations were performed at room temperature. We obtained the EPR and FMR spectra using

an updated SE/X-2547 microwave spectrometer (Radipan), and the IR absorption spectra using a UR-20 prism

spectrophotometer. We have studied synthetic diamond monocrystals of about 2 mm in size, which were grown

on a seed by a temperature-gradient method in the Ni-Fe-C system; diamond crystals which were prepared in

the Mg-C system by gloving on a seed with some additives; diamonds synthesized from hydrocarbons in a

NASA/CP--2001-210948 355 August 6, 2001
standardhigh-pressure apparatus without addingsolvent-metaJs, attheInstituteofHigh-Pressure Physicsofthe

RussianAcademy ofSciences.

RESULTS AND DISCUSSION

Synthetic diamonds seed-grown by a temperature-gradient method.

A comparison of EPR spectra given in Figure 1 shows that the spectrum of the seed-grown crystal is

a superposition of a line triplet g = 2, 0025 (due to a single substitution nitrogen atoms) and a wide single line of

exchange-bound paJcamagnetic nitrogen pairs and complexes. The total concentration of paJca_nagnetic nitrogen

in these crystals is about 4.1019 cm -3. The fraction of paJcamagnetic nitrogen complexes is 70 to 80% of the total

content of paJca_nagnetic nitrogen (ref. 3). The EPR spectra measured at liquid-nitrogen temperature exhibit no

line with g = 2.032 due to nickel atomic impurity (the EPR spectra of natural diamond do not contain this line

either).

1 2

"b

g

.4

i /'-'_j ,

/[

\

2 / t : \<

A_____/'/ a\ tq

_f J '%_.

Figure 1. EPR spectra of a seed-grown crystal (1) and synthetic diamond containing nitrogen in the form

of single substitutional atoms (2).

Figure 2. IR absorption spectra of statically synthesized diamond (1), a crystal grown on a seed by the

temperature gradient method (2), and type Ia natural diamond (3).

A special feature in the single-phonon (impurity) region of the IR absorption spectrum (Figure 2) of

the seed-grown diamond crystal is increased absorption intensity at 1282 cm -1, which is not typical of spectra of

the type Ib synthetic diamonds. This can be attributed to neither non-para_nagnetic impurity D center (nitrogen

and boron atoms in neighbor substitution positions (ref. 4)) because there is no sharp band at Raman frequency

(1332 cm -1) nor absorption due to non-compensated boron acceptors (as in the case of semiconductor diamonds)

because otherwise there should be intensive 2465 and 2810 cm -1 lines due to purely electron transitions.

Therefore, the excessive absorption at 1282 cm -1 in the spectrum of the seed-grown crystal can be only due to

non-paJca_nagnetic nitrogen in the form of paired impurity A centers, which agrees qualitatively with the EPR

data on an increased content of paramagnetic nitrogen pairs and complexes in these diamond crystals.

The concentration of the C and A centers (Arc and Na, respectively) in the seed-grown crystals was calculated by

the equations which had been derived in the course of solving a spectrophotometric problem for a twocomponent system and were similar to equations (14) in (ref. 5). In doing so, the coefficient of proportionality

between the concentration and the absorption factor at 1135 cm -1 was taken as 1.6.10 TM cm -2 (ref. 6) in the case

of paJca_nagnetic nitrogen and 5.8"10 TM cm -2 (ref. 7) in the case of non-paJca_nagnetic nitrogen (at 1282 cm 1). The

calculations have demonstrated that the content of the A centers in the samples varies from 3.1019 cm -3 to 7.1019

cm 3. In view of the EPR data on para_nagnetic nitrogen, it gives the fraction of 40 to 60% of the total content

of impurity nitrogen in the crystals. The volume distribution of impurity nitrogen was studied by a method of

local spectrophotometric analysis on several samples. We compared the IR absorption spectra of peripheral and

central zones of the crystals. The area of these zones is equal to about two thirds of the total cross-sectional area

of the sample. It turned out that the content of impurity nitrogen in the central zones of the crystals is higher and

NASA/CP--2001-210948 356 August 6, 2001
theNo/NA ratio noticeably lower in comparison to the peripheral zones. When passing from the central to the

peripheral zones, Nc decreases, on the average, by 17%, while NA by about 36% (i.e. almost twice), which

agrees with the common concept of the structure of these centers. With a random distribution of impurity

nitrogen in diamond crystals synthesized by spontaneous crystallization, the fraction of non-para_nagnetic

nitrogen impurity centers is negligibly small (and is not revealed in the IR absorption spectra) (ref. 8) even in

the case of high concentrations of paxamagnetic nitrogen. Therefore, the formation of a large amount of the A

centers can be attributed to the special features of crystal growth under the condition of a temperature gradient

(low growth rates, long p,T-action, etc.). Analysis of the experimental findings shows that the A centers form

mainly through entering of atomic nitrogen into a lattice point which is adjacent with one already occupied by a

nitrogen atom. In this case, the entering of nitrogen atom into the crystal volume does no longer obey the

random distribution law: according to this law, with the total nitrogen concentration in diamond crystals being

about 5"1019 cm -3 the fraction of non-para_nagnetic centers must be less than 10% (ref. 8). The formation of the

A centers through transformation of the C centers seems less likely. First, using the experimentally found value

of nitrogen aggregation rate at 1700 K (ref. 9) and applying a kinematic equation proposed in (ref. 2) we can

find that under the p,T parameters of growing the diamond crystals under study it will take more than three

months to reach a 50% degree of the C--->A center transformation. Second, one of the grown crystals still had a

seed crystal whose impurity nitrogen was solely in the form of the C centers.

The results of investigations demonstrate that under specially preset and controlled growth conditions

one can obtain diamond crystals, which would approach the most abundant type of natural diamonds, i.e. type

Ia, in the state of impurity nitrogen. This proves the potential of synthesizing large crystals close to natural ones

in their physicomechanical properties, by a temperature-gradient method of growing on a seed. The present

findings suggest also that the formation of type Ia diamonds under natural geological conditions might occur in

a similar way, i.e. at a temperature gradient.

Special features of the EPR spectra of diamonds grown in non-conventional growth systems.

Earlier (ref. 10) we reported the results of spontaneous diamond synthesis in the Mg-C system with

various additives, which allowed us to obtain diamond crystals with controllable semiconductor properties. The

main disadvantage of this method is that the synthesized crystals axe very small (about 0.5 _tm). The samples

studied in the present work were prepared in the Mg-C system by growing on a seed rather than by the

spontaneous synthesis method. The EPR spectrum (Figure 3) of these diamond single crystals is a classic

nitrogen triplet (AHm - 0.3 Gs, g = 2.0025) with a low nitrogen concentration (Nc - 1016 spin/cm3). The sample

contains neither atomic Ni impurity nor ferromagnetic inclusions. In all parameters this EPR spectrum is similar

to that of natural diamond crystals with low nitrogen content.

4Gs

.... ft 1

Figure 3. Manifestation of the 13C (1) and lSN (2) isotopes in the EPR spectrum of the single crystal grown

on a seed in the Mg-C system.

Among distinct features of the grown sample under consideration are its good cut, large size (2-3

mm), an increased content of non-paxa_nagnetic nitrogen aggregate centers (according to the IR spectroscopy)

and full absence of paxamagnetic non-impurity structural defects, which may be revealed by the EPR method.

NASA/CP--2001-210948 357 August 6, 2001
Contrastingly, thespontaneously synthesized crystalswiththesamelowconcentrations oftheCcenterscontain

usuallyalargeamount ofstructural defectssuchasdislocations, interstitial atomsandvacancies, whichmanifest

themselves inthesuperposition ofasingledefect line(withg=2.0027andg=2.0029) inthespectrumcenter

andleadtoanincreased integral intensityofthecentral lineincomparison tothesatellites.

Onthissample(owingtoitsuniqueproperties) wewereabletoobserveanumber offineeffects,

namely: thesuper-ultrafine interaction ofelectrons ofsinglenitrogencenters(whichdonotinteraction witheach

other)the13C(I = 1/2)isotopeatomslocatedinthenearest neighborhoodf nitrogen(ref.11)(without

purposeful dopingthenatural occurrence ofthe13Cisotopeisabout1%).Neareachoftheofthethreelinesof

themainitrogentripletthereaxethreepairsofgood-resolution weaklines.Thecalculated distances fromthese

pairsoflinestothecenterofthemainlinesofthenitrogentriplet(2.2,4.6,and6.4Gs)andthedistances

betweenthelinesofeachpair(4.4,9.2,and12.8Gs)axeingoodagreement withtheresultsreportedin(ref.12).

Besides, theEPRspectrumofthissampleexhibitswoweaklinesassociated withthepresenceofthe15N

isotope with its own nuclear spin I = V2 (natural occurrence of 15N is 0.365). The validity of attributing these

lines to the 15N isotope is confirmed by the fact that the ultrafine interaction constants calculated from the

experimental spectrum (All = 57.2 Gs, A± = 41.0 Gs) coincide very closely with the results obtained in (refs. 12,

15

13) for natural diamonds and synthetic diamonds doped with the N isotope.

Thus, the crystal under study is an analog (in the impurity-defect composition) of natural diamonds

and can serve as an ideal object for investigation of ultrafine and super-ultrafine interactions with the 13C and

15N isotopes of limited occurrence. For this purpose there is no need of an expensive doping of synthetic

diamonds with the isotopes.

Diamonds synthesized from hydrocarbons.

We studied diamonds synthesized from hydrocarbons in a standard high-pressure apparatus without

adding solvent-metaJs, at the Institute of High-Pressure Physics of the Russian Academy of Sciences. Thermal

destruction of hydrocarbons was carried out at thermodynamic parameters corresponding to the diamond

stability region (ref. 14). The crystals synthesized under these conditions axe mainly in the form of octahedrons,

they are colorless and transparent.

The EPR spectra of these diamonds have the following main features:

1. The EPR spectrum contains the triplet of lines (typical of diamond powders) with g = 2.0025 due to

nitrogen atoms which substitute carbon in the diamond lattice. The nitrogen concentration is about

(1-2).1018 cm -1. We revealed no EPR signals from any paxamagnetic structural defects. Also, we did not

observe a line with g = 2.0324 (nickel atomic impurity).

2. No FMR spectrum was recorded. This suggests that these diamond crystals contain no ferromagnetic

inclusions.

3. By adding nitrogen-containing organic materials to initial hydrocarbons, one can easily vary the content of

impurity nitrogen in diamond from 10 TM to 10 20 cm-3. Diamond crystals with a nitrogen concentration _>1019

cm -3 are yellow in color.

4. Adding of boron-containing organic compounds results in a decreased concentration of paxamagnetic

nitrogen centers down to 1017 cm -3 and below. In this case, these diamond crystals become black and

increase their electric conductivity.

5. As distinct from the crystals synthesized from graphite in the presence of a Ni-Mn solvent, the nitrogen

centers revealed by the EPR method in diamond crystals synthesized from hydrocarbons exhibit a deviation

from random distribution.

6. In the case of diamonds synthesized from hydrocarbons, the width AHm (between the extremum points on

the derivative of the absorption curve) reaches 3.5-4 Gs (in the case of synthetic diamonds prepared using a

metal solvent, AHm for nonsaturated EPR lines does not exceed 2.5 Gs).

The revealed deviations from random distribution of the nitrogen impurity can be associated both

with initial caxbon-containing material and with the synthesis process proper. In hydrocarbons with amine

groups the nitrogen distributed is quite orderly rather than random. Conceivably, when entering into synthetic

diamond crystals as an impurity, nitrogen atoms retain a part, though quite insignificant, of the above-mentioned

orderly pattern.

The EPR lines of nitrogen with a width AHm = 3-4 Gs axe observed only in the natural diamond

crystals which possess a considerable concentration of paxaxnagnetic nitrogen. In such crystals nitrogen is

usually distributed no uniformly and its concentration grows in the direction from the center towards the edge of

the crystal (ref. 11), and the ratio between concentrations of single and complex nitrogen centers is indicative of

the clustering of nitrogen impurity. The presence of the EPR spectrum of nitrogen with lines of different width

suggests that the conditions of formation of natural diamond crystals during their growth undergo abrupt

changes. It is the yellow-green outer zone which the EPR spectrum of nitrogen with lines of width AHm = 3-4

Gs can be attributed to. It can be assumed that the conditions of diamond formation from hydrocarbons simulate

NASA/CP--2001-210948 358 August 6, 2001
toacertainextent somestagesofgrowthofnatural diamonds. Itislikelythatnatural diamonds toohavesome

orderlydistribution ofimpuritynitrogeninsideindividual zones.

CONCLUSION

Thus,in all thedirectionsstudiedweobtainedthedatawhichindicatethatthespectroscopic

characteristics ofthesyntheticdiamonds understudyareclosetothoseoftypeIaandIIanatural diamonds. The

findingsofthepresent workcanserveasscientificsubstantiation ofthefeasibilityofsynthesis ofdiamond

crystalswhichwouldnotdifferindefect-impurity composition (andthusinphysico-mechanical properties)

fromthebestsamples ofnatural diamonds.

AcknowledRments. Theauthorsofworkaregrateful toSTCU(project1622)formaintenance offinancial

support ofthecarriedoutresearches.

REFERENCES

1. Berman, R., Hudson,P. R.W., Martinez, M., Nitrogen in diamond: Evidence from thermal conductivity. J.

Phys. C: Solid State Phys., vol. 8, 1975, p. L430-434.

2. Chrenko,R. M., Tuft, R. E., Strong, N. M. Transformation of the state of nitrogen in diamond. Nature, vol.

270, no. 5633, 1977, p. 141-144.

3. Zaadtsky, I. M., Tikhonenko, K. A. Processing of EPR spectra consisting of wide and narrow absorption

lines. Ukrainskiy Fizicheskiy Zhurnal, vol. 14, 1969, p. 125-128.

4. Malogolovets, V. G. Impurity nonpaacamagnetic D center in synthetic diamond crystals. In: The Influence of

High Pressures on the Structure and Properties of Materials [in Russian], ISM AN UkrSSR, Kiev, 1983,

p. 26-29.

5. Bokiy, G. B. et al. Natural and Synthetic Diamonds [in Russian], Nauka, Moscow, 1986.

6. Nachalnaya, T. A., Podzyaacei, G. A., Prikhna A. I. et al., Spectroscopic investigations of polycrystals

sintered from diamond micron powders. Sverkhtverdye MateriaJy, no. 3, 981, p. 23-27.

7. Kaiser, W., Bond,W. L. Nitrogen a major impurity in common type I diamond. Phys. Rev., vol. 115, no. 4,

1959, p. 857-863,

8. Nachalnaya, T. A., . Shulman, L. A. On relation between the paacamagnetic nitrogen concentration and

synthetic diamond crystal habit according to EPR data. Sinteticheskiye Almazy, no 3, 1973, p. 11-13.

9. Evans, T., Qi, Z. The kinetics of the aggregation of nitrogen atoms in diamond. Proc. Roy. Soc. London. A,

vol. 381, 1982, p. 159-178,

10. Nachalnaya, T. A., . Podzyaacei, G. A., Shulman, L.A. et al. Special features of EPR spectra of

semiconductor diamonds grown in Zn-Mg-C system," in: II All-Union Meeting on Deep States in

Semiconductors (Tashkent, 1980). Part II, pp. 88-89 [in Russian], Izdatelstvo Min. Vysshego

Obrazovaniya SSSR, Moscow, 1980.

11. Ralchenko, V.G., Smolin, A. A., Konov, V. I. et al. Laacge-area diamond deposition by microwave plasma.

Diamond Relat. Mater. no. 6, 1997, p. 417-421, .

12. Loubser, J. H. N., Du Preez, L. New lines in the ESR spectrum substitutional nitrogen donors in diamond.

Brit. J. Appl. Phys., vol. 16, 1965. p. 4571-4582

13. Klyuev, Yu. A., Nepsha, V. I., Naletov, A. M. Electron paacamagnetic resonance spectra and IR absorption

spectra of diamonds doped with nitrogen isotopes. FTT, vol. 16, no. 11, 1974, p. 3259-3264

14. Yakovlev, E. N., Voronov, O. A. Diamonds from hydrocarbons. Almazy i Sverkhtverdye Materialy, no.7,

1982, p. 1-2.

NASA/CP--2001-210948 359 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

The FTIR Studies of (100) Oriented Diamond Films grown on Si

Substrate Using Temperature Gradient Across the Substrate

E.Titus 1, A.K.Sikder2,U.Paltnikarl,M.K.Singh 1, and D.S.Misra 1@

1Department of physics, Indian Institute of Technology, Bombay, INDIA

Mumbai-400076

2Center for Microelectronics research, University of South Florida, Tampa, FL 33620

Abstract

Diamond films with dominant (100) texture were grown with a temperature gradient

across the Si(100) substrates using hot filament chemical vapor deposition technique.

Deposition was carried out with 0.8% CH4 in balance hydrogen at average substrate

temperature of 880°C. The deposition pressure was varied between 26.61x 102 Pa to

159.96x 102 Pa. Films were characterized using X-ray diffraction (XRD), scanning

electron microscopy (SEM) and Fourier Transform Infrared Spectroscopy (FTIR). XRD

shows very strong (400) reflection in all the samples. SEM results shows smooth

diamond surface comprised of (100) platelets. As the (100) diamond plates were grown

on top of the (100) oriented silicon substrate the faces are more or less aligned parallel

with the substrate surface resulting in a relatively smoother diamond surfaces. FTIR

studies show novel features in the films. Quantitative analysis was carried out to measure

the H content in the films.

@author for correspondence.

E-mail: mrdesh@phy.iitb.ac.in

NASA/CP--2001-210948 360 August 6, 2001
Experimental

The deposition of the films was carried out in a HFCVD apparatus whose details are

discussed elsewhere (16). Mirror polished p-type Si (100) wafers were used as substrate.

The substrates were treated with 2gm diamond powder and cleaned ultrasonically in a

acetone bath. Prior to the deposition of the thin film heater the substrates were dipped in

HF followed by de-ionized water cleaning in order to remove the active oxide layer of

silicon. Few depositions were done on the substrates without treatment with diamond

powder for comparison. Chromium thin film heater was deposited in thermal evaporation

unit on the backside of the substrates for the generation of the gradient Prior to the

introduction of the reactant gases, the gradient was created across the sample by applying

dc power to the heater. The gradient was measured at both ends of the sample using

chromel-alumel thermocouples. The reactant gases were methane and hydrogen at the

flow rates of 1.6 and 200 sccm, respectively. The average substrate temperature was

measured at the center of the substrate and was held constant for all depositions at 880°C.

The substrates were treated with hydrogen plasma for half an hour prior to the deposition.

Self-standing diamond sheets were used for IR analysis by making a window of diameter

6-7 mm in the center by chemical etching.

Results and Discussions

Typical SEM micrographs of the samples deposited with temperature gradient across

substrate show that the films contain dominantly (100) oriented grains. Moreover the size

of the grains increases significantly with the increase in temperature gradient. At lower

gradients the crystal size was less than 1gm. As the gradient increases further the crystal

size exceeds lgm subsequently increasing to about 3.6 gm for a gradient of 150°C cm -1.

NASA/CP--2001-210948 361 August 6, 2001
Thedensityof thegrainshoweverdecreasesasafunctionof thegradient.Any further

increasein thegradientresultedin lossof theorientationandthedensityof thegrains.

Thesampledepositedat 159.96x102Pawithoutgradientshowsa surfacemorphology

typically observedin verypoor qualitydiamondfilms Thegrainsaretypically having

cauliflowerstructureevidentlycontaininglargeconcentrationof non-diamondcarbon

impurities.Appearanceof strong(400)peakin XRDpatternalsoindicatesthegrowthof

(100)orienteddiamondfilm.

In orderto understandthemechanismresponsiblefor theformationof (100)oriented

grainswe depositedthefilms onthesiliconsubstratesfor variousdepositiontimeswith

andwithout the gradient.Bareuntreatedsilicon (100) substrateswere subjectedto

Hydrogenplasmaetchingat 880°Cwith andwithoutthegradientwith anobjectiveto

revealthestructuraltransformationonthesurfaceof thesubstratethatmightleadto the

orientationof thegrains.It is notedthatthesiliconsurfacedevelopsmicro-poreswhen

subjectedtohydrogenplasmaetchingin presenceof thetemperaturegradient.Thesizeof

theporesonthesurfacerangesfromsubmicronto 1-2micron.In contrastnoporescanbe

seenonthesiliconsurfacesubjectedto Hydrogenplasmaetchingwithoutthegradient.

We stronglybelievethatthegrowthof (100)orienteddiamondgrainsis directlyrelated

to theformationof theporesonthesiliconsurface.Theetchingof the siliconsurface

resultsin the exposureof underlying(100)orientedsiliconsurfacethat mightactasa

nucleationcenter.

Recently,considerabledebatepertainingto thegrowthmechanismof (100)and(111)

orienteddiamondgrainsis goingon. Severalgrowth modelshavebeenpresented.

HREELSmeasurementswereconductedin situ on as-grownhomoepitaxialdiamond

NASA/CP--2001-210948 362 August6,2001
(111)surfaceandhighlyoriented(111)diamondfilmsgrownonSi (001) andthepeaks

correspondingto the stretchingandbendingC-Hvibrationwere foundto be present.

Basedontheseresultsit hasbeenconcludedthatthe(111)growthoccursin atwo-stage

process.Using HREELSstudiesthe diamond(100) surfacesgrownat 800°Cwere

investigated.TheHREELSspectrashowedthepresenceof onlythepeakscorresponding

to thevariousmodesofsp3CH2radicals.Thisstronglysupportsthatthegrowthof (100)

orienteddiamondtakesplacevia CH2species.IndeedweobservethattheFTIRspectra

in the(100)orienteddiamondfilmsconsistmostlyof sp 3 CH2 bands.

The FTIR spectra in the range 2800cm -1 and 3100cm -1 of (100) oriented and (111)

oriented diamond films. Both the films were deposited at a pressure 159.96x 102 Pa and

temperature 880°C. The spectrum of (100) oriented samples show dominant sharp peaks

at 2930 cm -1 and 2860 cm -1 corresponding to asymmetric and symmetric stretch bands of

sp 3 CH2 group. The spectrum also shows weak bands at 2960 cm -1 and 2880 cm -1 related

with sp 3 CH3 asymmetric and symmetric stretch bands. It is noteworthy that no peak

corresponding to sp 3 CH vibration is present. The spectrum of (111) oriented films on the

other hand contains the multiple bands corresponding to sp 3 CH3, sp 3 CH2 and sp 3 CH

groups which appears as a broad band. It should be mentioned however that the quality of

the films deposited without the gradient at 159.96x 102 Pa is very poor and that may also

affect the FTIR spectra. From these results we conclude that the growth of (100) oriented

diamond proceeds via dominantly CH2 species. This is also confirmed by HREELS

studies in which it is observed that (100) surface of diamond are terminated by CH2

radicals.

NASA/CP--2001-210948 363 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

Diamond Micro Crystals Grown by CVD Method

Abstract

Hsiao-Kno Chung and James C. Sung

KINIK Company, Taiwan, R_O.C.

Hot filament CVD method can be used to grow micron sized diamond crystals in the quantity of several

millions per square centimeter. These crystals are euhedral with fully developed facets. Such uniform shaped

crystals can be sized easily with tight distribution. When they are used to polish workpieee, they may not cause

conspicuous scratches. Moreover, the smooth surfaces of these whole crystals can minimize the amount of

adhered impurities, so the contamination of workpiece may be avoided. The micro diamond crystals so grown

are the ideal superabrasive for polishing high valued materials such as semiconductors and other electronic

devices.

Key Words: Chemical Synthesis, Hardness

NASA/CP--2001-210948 364 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

Diamond Synthesis in Perspective

James Sung

China Grinding Wheel Corporation, Taipei, Taiwan, ROC

Taipei University of Technology, Taipei, Taiwan, ROC

Abstract

Diamond may be converted directly from a non-diamond carbon source in microseconds by a displacive

movement of atoms without diffusion, or synthesized catalytically in minutesby reconstructive means with

the transport of atoms. The diamond formation pressure and temperature are dependent on the time scale

and the structure of the carbon source. The instant conversion may occur with zero-dimensional carbon

(point-like sp° atoms) by PVD, one-dimensional carbon (line-like sp I chains) by dynamite detonation, or

two-dimensional carbon (plane-like sp2 surface) by shock wave compaction. The timely synthesis may

proceed with two-dimensional (sp 2) graphite (e.g., by puckering recrystallized flakes in a molten catalyst,

or with three-dimensional (sp 3) diamond-like molecules such as solute atoms in a molten catalyst (e.g., by

temperature gradient method), or caged atoms in methane molecules (e.g., by CVD).

The approximate activation energies (E) for the major diamond synthesis methods may be estimated from

their reaction mutes. The threshold pressure and temperature for each of these synthesis routine may be

calculated based on the estimated activation energy. The results are summarized in the following table.

C Source State Method Route Time EteV) PG_CG_p__

Plasma sp e PVD Direct Instant 0.33 16 3500

Dynamite sp I Explosion Direct Instant 0.25 12 2700

Graphite sp 2 Shock Wave Direct Instant 0.17 8.1 1800

Graphite sp 2 Puckering Catalytic Timely 0.13 6.2 1250

Solute sp a Precipitation Catalytic Timely 0.13 6.2 1250

Methane sp a Adjoining Catalytic Timely 0.09 4.3 800

Keywords: Diamond Synthesis, Diamond Catalyst, PVD Diamond, CVD Diamond

NASA/CP--2001-210948 365 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

The Eastern Wind of Diamond Synthesis

James C. Sung I & John C. Chen z

1, KINIK Company, Taipei, Taiwan, R.O.C. and Taipei University of Technology, Taipei, Taiwan, R.O.C.

2. Fine Abrasives Company ofTaiwan, Taoyuan, Taiwan, R.O.C.

Abstract

More than 300 tons of diamond superabrasives are made each year by high pressure method. China

makes more than 1/3 of the world production and its share is rapidly increasing.

The commercialization of CVD diamond films are yet to succeed. Chinese producers are more

competitive in this area than high pressure synthesis. Hence, the world will also see Chines dominance in this

area in the future.

Facing with such a pressure of Eastern Wind of diamond synthesis, diamond producers in Japan and

Western world must concentrate in making high grade products that Chinese cannot produce.

NASA/CP--2001-210948 366 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

MODERN THERMODYNAMICS FOR ACTIVATED LOW PRESSURE CVD

DIAMOND SYNTHESIS

Ji-Tao Wang

Dept. of Electronic Eng., Fudan University, 220 Handan Rd., Shanghai 200433, China

Tel: +86 21 6564 2389, Fax: +86 21 6564 8783, Email: j tx_v_an.g@_'_u.d__ao.._.a..c.,c.n

ABSTRACT

In activated low pressure chemical vapor deposition (CVD) process the transformation from graphite to

diamond under low pressure is a nonspontaneous reaction, i.e. (dG_)z p > 0. There must be at least another

spontaneous reaction in the system. That is mainly the association of superequilibrium atomic hydrogen, i.e.

(dG2)zp< 0 and I(dGt)zpl<l(dGgr+l. Therefore, the whole reaction is still going toward the right side, i.e.

[(dGt)zp> 0, (dG2)zp< 0 & (dG)r.p_<0]. Here, (dGl)r.p, (dG:)zp and (dG)zp={(dGOzp+(dGa)zp } are Gibbs free energy

changes of reaction 1, reaction 2 and the whole reaction (or system) at constant temperature and pressure,

respectively. The extreme criterion of [(dG_)zp> 0, (dG2)zp< 0 & (dG)r.p = 0] is the basis of calculation for phase

diagrams. However, the calculated phase diagrams are nonequilibrium (stationary) phase diagrams. All these

calculated phase diagrams agree excellently with a lot of experimental data reported in the literature by several

research groups. These phase diagrams play an important role for guiding the industrial production of diamond

films. The situation of [(dG1)zp> 0, (dQ)zp< 0 & (dG)zeg_<0] is called reaction coupling or coupled reaction.

Coupled reactions are of great importance in biological processes and many fields of new science and technology,

including the frontier carbon technology.

After quantitative verification of reaction coupling in the activated low pressure CVD diamond process, the

equality of the 2 nd law is no longer a necessary and sufficient condition for a complex system being in equilibrium.

For a complex system including coupled reaction the equality part of the 2_ law, i.e. [(dG)zp = 0], can be subdivided

into an equilibrium part of [(dG1)zp = 0, (dG2)zp = 0 & (dG)zr = 0] and a nonequilibrium part of [(dGl)zr > 0, (dG:)zp<

0 & (dG)r,p = 0]. Therefore, the classification of thermodynamics has been drastically changed, as shown in Fig. 1.

There is a new field for the nonequilibrium (stationary) phase diagrams. The activated CVD diamond system may be

the simplest system including reaction coupling, and can be quantitatively calculated.

Modem thermodynamics

[(dG)r.p <-0]

/divided

Zero-dissipative thermodynamics Dissipative thermodynamics

[(dG)z p = 0]

........-."'_bdivide_"'---....._

Equilibrium Nonequilibrium

zero-dissipative - zero-dissipative

thermodynamics thermodynamics

[(dGl)r,p=O,(dG2)r,p=O [(dGt)r,p>O,( dG2)r,e<O

and (dG)r,p= 01 and (dG)r,p = 01

I representative ! event Classical The a thor's

.,._s(duG)r,p < 0]

bdivide"d""-----_

Linear dissipative Nonlinear dissipative

thermodynamics thermodynamics

[(dG)r.r<0 and [(dG)r.p<0 and

linear relation nonlinear relation

between J & X] between J & X]

Onsager s Prigogine's

equilibrium nonequilibrium (stationary) reciprocity dissipative

phase diagrams phase diagrams relation structure

Fig. 1 Systematization of modem thermodynamics for complex systems including reaction coupling

Note: X and J represent thermodynamic force and thermodynamic flux for irreversible process, respectively.

Keywords: Diamond, Thermodynamics, Phase diagrams, Reaction coupling, Chemical vapor deposition

NASA/CP--2001-210948 367 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Preparation of Diamond Surfaces

Dr. James E. Butler

Gas/Surface Dynamics Section

Naval Research Laboratory

Washington DC 20375

butler_ccf.nrl.navy.mil

202-767-1115

202-767-3321 fax

Abstract:

Fundamental to understanding of how diamond interfaces and interacts with other materials and

environments is the Diamond Surface, its structure and chemistry. However, production of single crystal

diamond surfaces with reproducible structure and chemistry is problematic due to the metastablity of

diamond and the damage caused by traditional surface preparation techniques such as polishing, abrasive or

laser cutting, ion sputtering and annealing. The quality of the diamond surface is important to fundamental

surface science studies, the use of the surface as a template for homoepitaxial growth, and as an interface

for adhesion and bonding to other materials. This work will describe our current understanding and

methods employed for preparing high quality diamond surfaces using polishing, chemical cleaning, and

plasma treatments.

NASA/CP--2001-210948 368 August 6, 2001
Proceedingsof the SixthApplied DiamondConference/Second

FrontierCarbonTechnologyJoint Conference(ADC/FCT2001)

Y. Tzeng, K.Miyoshi,M. Yoshikawa,M. Murakawa,Y. Koga,

K. Kobashi,G.A.J.Amaratunga(Editors)

INFLUENCE OF ELECTRON BEAM IRRADIATION ON

WETTABILITY OF DIAMOND SURFACE

Kazuya OGURI, Akira TONEGAWA

Department of Physics, Tokai University,

I l 17 Kitakaname, Ilixatsuka, Kanagawa, 259-1292, Japan

Nobuhiro IWATAKA and Yoshitake NISHI

Department of Materials Science, Tokai University,

1117 Kitakaname, Hiratsuka, Kanagawa, 259-1292, Japan

TEL:+81-463-58-1211 ext 4206, FAX:+81-463-58-1812

am026429@keyaki.cc.u-tokai.ac.jp

ABSTRACT

A diamond coating enhances the resistance to scratch, and application of a diamond eoating to the

hard window of infrared sensors is beneficial. A serious problem with the infrared windows of sensor

is misting of the diamond window. If the infrared ray is transmitted through the misted diamond

window, accurate results would not be obtained. Misting-free windows would reduce the energy loss of

analytical instruments such as fourier transform infrared microscope. The sheet electron beam

irradiation (SEBI) may induce clear vision without residual impurity atoms. The misting-free dentist

mirror has been developed by SEBI treatment [1]. Therefore, we studied the influence of electron

beam irradiation on wettability of the diamond surface.

Sheet electron beam irradiation (SEBI) was homogeneously applied with an electron curtain

processor (Type:CBI 175/15/180L, Energy Science Inc., Wobum, MA; Iwasaki Electric Group Company,

Tokyo ). The sheet beam was generated by a tungsten filament in a vacuum. Since the sheet

electron beam is the most important feature, homogeneous treatment of the electron beam was

performed under protective nitrogen gas with the pressure inside the apparatus kept at atmospheric

pressure. The concentration of impure oxygen gas was below 400 ppm.

The SEBI decreased the contact angles of water on diamond surface. Namely, the water drops on

a SEBI-treated surface have a lower contact angle and higher interfacial energy than those on a non

SEBI-treated surface.

Wettability depends on the identity of the atoms that are terminated on the diamond surface [2-4].

If there is'an increase in hydrogen atoms that are terminated on the diamond surface, the wettability

decreases. On the other hand, if there is an increase in oxygen atoms that are terminated on the

diamond surface, the wettability increases.

Based on the result of x-ray photoeleetron spectra, the magnitude of the O (ls) peak on the x-ray

photoelectron spectrum increased. Since the terminated atoms on the diamond surface changed from

hydrogen atoms to oxygen atoms, SEBI enhanced the wettability and decreased the contact angle.

Keywords: electron beam, diamond, wettability, surface, x-ray photoelectron spectra

[1 ] Y.Nishi, H.Izumi, J.Kawano, K.Oguri, Y.Kawaguchi, M.Ogata, A.Tonegawa, K.Takayama,

T.Kawai and M.Oehi, J. Mater. Sei., 32 (1997) 3637.

[2] R.S.Butter, D.R.Waterman, A.H.Lettinglon R.T.Ramos and EJ.Fordham, Thin Solid Films, 311,

107 (1997).

[3] T.Ando, N.Ishii, M.Kamoshige, Y.Sato, New Diamond, 10, 2 (1994).

[4] R.Sappok and H.P.Boehm, Carbon, 6, 573 (1968).

NASA/CP--2001-210948 369 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

TIME-RESOLVED ELECTRICAL RESISTANCE MEASUREMENTS OF

GRAPHITIZED LAYERS IN DIAMOND DURING PULSED LASER IRRADIATION

S.M. Pimenov, V.V. Kononenko, E.V. Zavedeev, V.I. Konov

General Physics Institute, Vavilov str. 38, 119991 Moscow, Russia (e-mail: pimenov@nsc.gpi.ru)

M. Kehrli, W. Liithy, V. Romano, H.P. Weber

Institute of Applied Physics, University of Bern, Sidlerstr. 5, CH-3012 Bern, Switzerland

R.A. Khmelnitskiy

P.N. Lebedev Physical Institute, Leninsky avenue 53, 117924 Moscow, Russia

ABSTRACT

During short-pulsed laser processing of diamond and diamond-graphite heterostructures, various fast phase

transitions (such as graphitization, melting, evaporation, and diamond regrowth) may occur on the surface and in the

bulk of diamond materials. Data on these laser-induced structural modifications result usually from post-processing

examination of the material properties, while valuable information on both the occurrence and kinetics of the laserinduced phase transitions can be obtained from real-time monitoring of the variations in the optical/electrical

properties of a laser-irradiated region during (and after) a laser pulse action.

In addition to the transient reflectivity technique (ref. 1), we report on time-resolved measurements of electrical

resistance of graphitized layers in diamond and its changes induced by nanosecond pulses of an ArF excimer laser

(wavelength)v=193 nm, pulse duration x=20 ns) and a Nd:YAG laser ()v=-l.06 btm, x=20 ns). Two types of structures

were tested (Fig. 1). The first structure is a laser-graphitized layer on the surface of a 350-btm-thick diamond plate: a

graphitized pathway was produced by laser ablation and electroplated with gold, and after removal of gold from the

central part of the pathway a line test structure with two gold electrodes and a graphitized region in between was

formed (Fig.l(a)). The thickness of the laser-graphitized surface layer was determined from the surface profile

measurements (of laser-etched regions) with an atomic force microscope. The second structure is a buried graphitic

layer formed by 350-keV He + ion implantation in a diamond type-IIa single crystal (Fig. l(b)). The upper boundary

of the buried layer is estimated to position at the depth of 680 nm, and its thickness - of about 110 nm.

The current transients were measured with a fast oscilloscope (LeCroy 9354AL, 500 MHz) across the 50-£2

load resistance. It was found that the transient conductance of the graphitized surface layers depended on the

wavelength and incident fluence. At low laser intensities (for both wavelengths) the current pulses were identical to

the laser pulse shape (photoconductive response) and they were significantly broadened with increasing fluence,

thus showing a temperature-dependent component in the signals. As an example, Figure 2 shows the current pulses

induced by Nd:YAG laser pulsed irradiation in the first test structure at fluences lower (a) and higher (b) than the

laser-induced plasma threshold. For the buried layers, similar (but a more complicated) behaviour of the transient

conductance was observed in dependence on the incident fluence, but, in addition, the current pulse parameters were

found to vary with the number of shots. The obtained results are analysed in connection with the laser-induced

melting of carbon and internal graphitization during laser ablation of diamond films and laser irradiation of the

buried graphitic layers in diamond.

The work was supported by the Swiss National Science Foundation (projects 7SUP J062184 and 7IP62635).

Keywords: diamond-graphite heterostructures, laser irradiation, phase transitions

NASA/CP--2001-210948 370 August 6, 2001
a)

Au film (1 #m thick) electroplated

,/selectively onto a laser_blated groove

Laser_blated (graphitized) groove of

----------___'__'__'__'__'__'..'__'__'__'__'__'__'__'__'__'__'__'__'__'__:._n l 5 mm length and 0.4 mm width

/

CVD diamond plate 1

Irradiation of laser_jraphitized diamond surface wil_

J,$_, ArF (X=193 rim) or Nd:YAG (2v=1.06 lain)laser pulses

r:7:7:7:7:7:7:7:71 r:T:T:T:T:T:T:T:q

......................

CVD diamond plate

b)

Irradiation ofa graphitized buried layer

J,; with Nd:YAG (X=1.06 pro)laser pulses

................... -_------F--- Ionqmplanted buried layer

Diamond single crystal

Figure 1. Schematic view of two line structures tested in the experiments.

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fluences lower (a) and higher (b) than the laser-induced plasma threshold.

REFERENCES

1. P. Tosin, W. Ltithy, and H.P. Weber, "Observation of liquid carbon on diamond films under pulsed IR and UV

laser irradiation", Phys. Rev. B 53 (1996) 543-546.

NASA/CP--2001-210948 371 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

MECHANISM OF ELECTRICAL CONDUCTIVITY APPEARED IN SURFACE

REGION OF UNDOPED DIAMOND FILM

Takeshi Hosomi .1'3), Kiyoshige Tsuji 1'3) and Takeshi Kobayashi _'2'3)

1)Division of Advanced Electronics & Optical Science, Graduate school of Engineering Science,

Osaka University, Machikaneyama 1-3, Toyonaka, Osaka 560-8531, Japan.

Phone: +81-6-6850-6311, Fax: +81-6-6850-6341

*e-mail: hosomi @ sup.ee.es.osaka-u.ac.jp

2_FCT Collaborating Member

3)CREST, Japan Science and Technology Corporation (JST), Japan.

ABSTRACT

A p-type surface conductive layer appeared by calcium fluoride (CaF2) deposition treatment on the oxygenated

undoped diamond. Though a p-type surface conduction layer of hydrogenated undoped diamond disappears by

oxygenation and it reappears by surface re-hydrogenation, we obtained the surface conductive layer without

hydrogen plasma treatment. For reappearance of the surface conductivity, CaF2 deposition was carried out at several

temperatures employing electron beam evaporator. And the appearance of the surface conductivity depended on the

CaF2 deposition temperature, e.g. as the temperature became high, the values of carrier density increased in this

experimental range. The surface conductivity of CaF2-treated diamond had similar electrical properties to that of

hydrogenated diamond surface. The effects of CaF2 treatment on oxygenated diamond surface could be supposed to

abstract oxygen from diamond surface and to fluorinate the diamond surface leading to the elimination of surface

states and the stabilization of surface structure. And the surface conductivity depended on surface ambient.

Keywords: diamond, chemical vapor deposition, surface conductivity, calcium fluoride, hydrogenation

INTRODUCTION

Owing to its good properties such as high thermal conductivity, wide band gap and large breakdown field, diamond

is expected to be a promising candidate for applications to semiconductor devices suitable for hard electronics. Due

to the facility of the manufacture, a p-type surface conductive layer (refs. 1, 2) induced by hydrogen adsorption on

the diamond thin films is often utilized as a channel of diamond metal-insulator-semiconductor (MIS) or metalsemiconductor (MES) devices (refs. 3, 4).

This surface conductive layer has been investigated at the various points of view. Hole density is about 1013 cm -2 in

homoepitaxial undoped diamond films, Hall mobility is about 30 - 40 cm2/V,s (ref.5) and several atmospheres

influence these electrical properties (refs. 6, 7). The surface conductivity of hydrogenated diamond disappears by

surface oxygenation and the surface conductivity reappears by exposing oxygenated diamond surface to hydrogen

plasma. Though hydrogenated diamond surface has less surface states, oxygen adsorbed on diamond surface cause a

large number of surface states and it degrades electrical properties of diamond surface (refs. 3, 8). Therefore, oxides

are not suitable for insulator on diamond and non-oxide material is hoped to get a good insulator/diamond interface.

In our previous study, it has examined that calcium fluoride (CaF2) is a good candidate for insulator of MIS structure

(ref. 9). CaF2 as an insulator is able to suppress the amount of surface states density because it is a chemical

compound made up of calcium and fluorine without oxygen.

In this study, we show that surface conductivity also appears by the treatment of the oxygenated diamond thin

films with CaF2. It is a characteristic of CaF2 that surface conduction reappears by high temperature deposition of

CaF2 on oxygenated diamond surface without re-hydrogenation. It is thinkable that this surface conduction is caused

by the fluorine adsorption on the diamond thin films. This surface conductive layer is also applicable to the diamond

MIS or MES devices as the channel.

NASA/CP--2001-210948 372 August 6, 2001
EXPERIMENTAL

UndopeddiamondthinfilmswerepreparedonsyntheticIb (100)diamondsubstrates bymicrowaveplasma

assistedchemical vapordepositionmethodwithH2+ CH4gases. As-growndiamondfilmswereexposedtooxygen

plasmatoobtaininsulateddiamondsurfaceeliminatingthehydrogenated surfacelayer,andthenCaF2wasdeposited

ontheoxygenated diamondsurfaceatseveral temperatures byelectron-beam(EB)evaporation apparatus. Duringthe

EBevaporation, partof CaF2is dissociated foundoutbyemissionspectroscopy of evaporatingCaF2.Afterthe

removal of CaF2layeronthediamondbyhydrochloricacid,electrical measurements werecarriedoutwithAu

electrodes. Surfacestructurewasalsoevaluatedmployinglowenergyelectrondiffraction(LEED)method.

RESULTS AND DISCUSSION

Figure 1 shows the sheet conductivity of the CaF2-treated diamond

surface conductive layers as a function of CaF2 deposition

temperature. The data reveal marked temperature dependence and E

0

the presence of the onset temperature of around 400 °C, below which co

current does not flow. Following might serve as a clue to understand ,_=

the mechanism for reappearing of the surface conduction. One is the >,

chemical reaction between evaporated CaF2 and oxygenated .->

diamond surface. Ando et al. reported that fluorination of *5

-o

oxygenated diamond surface requires high temperature and the onset _-

o

temperature is about 400 °C when CF4 plasma is employed (ref. 10). O

This temperature is close to the onset temperature of reappearance of _

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surface conductivity by CaF2 treatment. The other is the _c

60

graphitization of diamond surface during CaF2 deposition under high

temperature. Therefore, the crystallinity of the CaF2-treated diamond

surface was evaluated by LEED method as shown in Fig. 2. These

LEED patterns exhibit clear diffraction spot. Therefore, diamond

surface did not covered by graphitic/amorphous layer. Electrical

measurements also suggest non-graphitization of the surface, as

described later.

20

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0

300 400 500 600

CaF2 Deposition Temperature (°C)

Figure 1. Sheet conductivities of CaF2-

treated diamond surfaces as a function

of CaF2 deposition temperature

(a) (b)

Figure 2. LEED patterns of CaF2 treated diamond surface. Accelerating voltages are (a)

130 V and (b) 180 V.

The surface conductivity was investigated by the Hall effect measurement. The conduction type of the surface

conductive layer was p-type. Hall mobility and sheet carrier densities of the CaF2-treated diamond surface

conductive layers are shown in Fig. 3 as a function of CaF2 deposition temperature. Sheet carrier density reveals

larger temperature dependence than that of Hall mobility. The high carrier density at high temperature should reflect

the amount of reaction between CaF2 and diamond surface region. The effects of CaF2 deposited on oxygenated

diamond surface could be supposed to abstract oxygen from diamond surface and to fluorinate the diamond surface.

NASA/CP--2001-210948 373 August6,2001
Oxygenated diamondsurfacehasahigh-density surfacestatespinningthesurfacefermilevel.Therefore, oxygen

shouldbeabstractedfromdiamondsurface,.g.replacement ofoxygenwithfluorineandoxygenabstractionby

calcium, forreappearance ofthesurfaceconductivity. Andit isprobablethatfluorinationofdiamondsurfaceoccurs

byfluorineradicalreachedtothediamondsurfaceathightemperature becausepartofCaF2evaporatinginEB

evaporator isdissociated byEBirradiation. Hallmobilityalsorevealstemperature dependence. Asthetemperature

becomeshigh,thevaluesofHallmobilityincreasebutit decreased suddenlyatabout550°C.Thereasondoesnot

declaredbutdiamondsurfacemaybecomeroughduringthereactionbetweenCaF2anddiamondsurface,.g.

diamondwasetchedoffbycalciumandfluorinexcitedbyEBirradiation. Atanyrate,surfaceconductivelayer

appearedonoxygenated diamondbyCaF2depositioneasily.Andanimportant factortoobtainsurfaceconductivity

willbeaneliminationofsurfacestates, achievedbyremoval ofoxygen.

40

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Figure 3. Dependence of (a) Hall mobility and (b) sheet carrier density of the CaF2-treated diamond

surface conductive layers on CaF2 deposition temperature.

The influence of atmosphere on surface conductivity was

also measured. Figure 4 shows the conductance of the CaF2-

treated diamond under low-pressure atmosphere (-10 Pa),

which depends on time and temperature. When air was

exhausted from the closed container, the conductance began to

decrease and the value seems to saturate. And then the

conductance decreased by a little annealing. At the room

temperature, when air was again introduced into the container,

the conductance began to increase and was saturated to the

previous value. This tendency was also observed on the

hydrogenated diamond surface. Ri et al. also reported that

resistance of surface conductive layer of hydrogenated

diamond decrease in low-pressure atmosphere and it recover

the previous value in the air. The surface conductivity of the

CaF2-treated diamond is sensitive to atmosphere in the same

way of that of the hydrogenated diamond. The Hall effect

measurements shown above performed after samples were

placed in the air more than 30 hours.

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Figure 4. Dependence of conductance of

the CaF2-treated diamond surface on

time and temperature.

CONCLUSIONS

Surface p-type conductive layer of CaF2-treated undoped diamond was investigated. After surface conductive

layer disappeared by oxygenation, it reappeared when CaF2 was deposited on the oxygenated diamond surface by EB

NASA/CP--2001-210948 374 August 6, 2001
evaporationathightemperature. Theconductivityof CaF2-treated diamondsurfaceshowedmarkedeposition

temperature dependence andtheonsettemperature wasatabout400°C.Asthetemperature becamehigher,the

carrierdensitybecamelargerandthisresultshouldreflecttheamount ofreactionbetweenCaF2anddiamondsurface

region.TheeffectsofCaF2depositedonoxygenated diamondsurfacecouldbesupposedtoabstract oxygenfrom

diamondsurfaceandtofluorinatehediamondsurface. ThesurfaceconductivelayerofCaFz-treated diamondhas

similarelectrical propertiestothatofhydrogenated diamondsurface. Thesurfaceconductivities dependonsurface

ambient andthenappearance ofsurfaceconductionseemstobeassisted.

REFERENCES

1 M.I. Landstrass and K. V. Lavi: Hydrogen passivation of electrically active defects in diamond, Appl. Phys.

Lett., vol. 55, no. 10, 1989, pp. 975-977.

2 T. Maki, S. Shikama, et al.: Hydrogenating effect of single-crystal diamond surface, Jpn. J. Appl. Phys., vol.

31, no. 10A, 1992, pp. L1446-L1449.

3 Y. Yun, T. Maki, et al.: Electrical properties of A1/CaF2/i-diamond metal-insulator-semiconductor fieldeffect-transistor fabricated by ultra high vacuum process, Jpn. J. Appl. Phys., vol. 37, no. 11A, 1998, pp.

L1293-L1296.

4 K. Tsugawa, K. Kitatani, et al.: High-performance diamond surface-channel field-effect transistors and their

operation mechanism, Diamond & Relat. Mater., vol. 8, 1999, pp. 927-933.

5 K. Hayashi, S. Yamanaka, et al.: Investigation of the effect of hydrogen on electrical and optical properties

in chemical vapor deposited on homoepitaxial diamond films, J. Appl. Phys., vol. 81, no. 2, 1997, pp. 744-753.

6 Ri S. G., T. Mizumasa, et al.: Formation mechanism of p-type surface conductive layer on deposited

diamond films, Jap. J. Appl. Phys., vol. 34, no. 10, 1995, pp. 5550-5555.

7 F. Maier, M. Riedel, et al.: Origin of Surface Conductivity in Diamond, Phys. Rev. Lett., vol. 16, no. 16,

2000, pp. 3472-3475.

8 Y. Otsuka, S. Suzuki, et al.: Fermi level pinning in Metal-Insulator-Diamond structure, Jpn. J. Appl. Phys.,

vol. 34, no. 5A, 1995, pp. L551-L554.

9 Y. Yun, T. Maki, T. Kobayashi: Surface state density distribution of semiconducting diamond films

measured from the A1/CaF2/i-diamond metal-insulator-semiconductor diodes and transistors, J. Appl. Phys., vol.

82, no. 7, 1997, pp. 3422-3429.

10 T. Ando, J. Tanaka, et al.: Diffuse reflectance Fourier-transform infrared study of the plasma-fluorination of

diamond surfaces using a microwave discharge in CF4, J. Chem. Soc. Faraday Trans., vol. 89, no. 16, 1993, pp.

3105-3109.

NASA/CP--2001-210948 375 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

APPLICATION OF RAMAN SPECTROSCOPY TO DETERMINE STRESS IN

POLYCRYSTALLINE DIAMOND TOOLS AS A FUNCTION OF TOOL GEOMETRY

AND TEMPERATURE

J.D. Comins, R.M. Erasmus and V. Mofokeng

School of Physics, University of the Witwatersrand, Johannesburg, Private Bag 3, WITS, 2050, South Africa

Z. Martin

Department of Materials, University of Oxford, Parks Road, Oxford, OX 1 3PH, United Kingdom

ABSTRACT

Polycrystalline diamond (PCD) tools commonly consist of a PCD layer sintered onto a cobalt-tungsten carbide

(Co-WC) substrate. These tools are used in diverse applications such as drilling for oil exploration, machining of

hard metal alloys and the cutting and finishing of wood chipboard. A knowledge of the stresses present in the PCD

layer is thus of great value to both tool manufacturer and end user, as both the magnitude and distribution of these

stresses affect tool behavior. In this paper, the stresses in PCD drill-bits are investigated by means of micro-Raman

spectroscopy. The position of the diamond Raman line reveals both the nature of the stress present (compressive or

tensile) and the magnitude of the stress.

It was found that the surface preparation technique, i.e. rough lapping, fine lapping or polishing, influenced the

average stress present in the PCD layer. In all three cases, the PCD surface layer was in compression, with the largest

stresses encountered in the roughly lapped sample (1.4 GPa) and the stress values decreasing with fine lapping

(0.8 GPa) and polishing (0.1 GPa). Spots with small tensile stresses were found in some polished samples.

Measurements of stress as a function of temperature for roughly lapped samples indicated a linear trend of

decreasing stress values with increasing temperature, although the stress remained compressive. Measurement of

stress at room temperature (after the heating cycle) indicated some evidence for a degree of annealing of the surface

compressive stresses.

Keywords: polycrystalline diamond (PCD), tool-bits, Raman spectroscopy, stress, temperature

INTRODUCTION

The applications of diamond, apart from it role as gemstone, are quite diverse (ref. 1). Currently it is mainly

exploited for its ultra-hardness property in cutting, grinding and polishing. The very property that contributes to its

desirability as an industrial ceramic also leads to it being extremely brittle. The solution to this problem is to make

use of sintered polycrystalline diamond (PCD), frequently with cobalt (Co) as a sintering agent. This combines the

hardness properties with increased toughness, which translates to better wear and fracture resistance under extreme

loading conditions (ref. 2).

The application environment of PCD tools is often quite hostile, and can include circumstances such as high

temperature and abrasive liquids. Certain demands are placed on the tool by the end user in terms of cost and life

expectancy. The stress state of the surface of a PCD tool is important in its functionality. A compressive stress state

opposes the formation and propagation of median cracks under impact loads, as well as the propagation of microcracks already present in the PCD. A tensile stress state, on the other hand, is conducive to crack formation and

propagation. It is thus clear that the stress state of the surface of a PCD tool is of interest to both manufacturer and

end user.

Raman spectroscopy has been used quite widely in the measurement of stress in thin diamond films (see e.g.

refs. 3 to 7), but has had less application to PCD tools (ref. 8). Reference 9 compares X-ray diffraction and microRaman spectroscopy as techniques to measure residual stress in diamond coatings. In the case of thin diamond films,

stresses at the diamond/substrate interface and within the diamond films have been analyzed. Residual compressive

stresses up to 7.1 GPa have been measured in thin films without delamination occurring (refs. 4 and 5), and have

been attributed to the difference in the thermal expansion coefficients between the diamond film and the substrate, as

well as the high density of microcrystalline defects in the films (refs. 3 and 4). Different facets of isolated crystals of

NASA/CP--2001-210948 376 August 6, 2001
CVDdiamondhavebeenshowntohavedifferent typesofresidual stressdepending onthedominant growthsector

(ref.7).Inreference6acorrelationwasfoundbetweenthedegreeofsp 2 bonding present in diamond films and the

magnitude of residual compressive stress. The work on PCD tools (ref. 8) showed that micro-Raman spectroscopy

can be applied to the measurement of residual stress in PCD cutting tools with reasonable accuracy. It also illustrated

that the stress distribution in the PCD layer is not necessarily hydrostatic, and a bi-axial stress model was found to be

more appropriate.

THEORY

In Raman spectroscopy, the sample under study is illuminated with a laser beam with photons at a specific

frequency. A small proportion of the photons interact with the lattice phonons in the sample material and are

scattered inelastically. Inelastic scattering from optic modes is known as Raman scattering and the scattered photons

have a frequency shifted from the incident photons by the frequency of the mode which is characteristic of the

sample material. Diamond has a single, triply degenerate Raman peak at 1332.5 Rcm -1 (ref. 10) (units are in

wavenumbers relative to the incident laser beam wavenumber). The position of the Raman peak thus conveys

information on the nature and structure of the material under study.

In an ideal crystal, the sample would be stress free and the Raman peak would be measured at its theoretical

position. As soon as the crystal is subjected to stresses, be it externally imposed or due to internal imperfections, the

frequency of the lattice phonons is subtly affected and thus also the position of the Raman peak. In this way, careful

measurements of the shift of the Raman peak relative to its unstressed position can yield information on the stresses

present in the sample. Compressive stress results in a shift of the peak to higher frequencies, and tensile stress to

lower frequencies. In order to convert a degree of shift to a stress quantity, use is made of the coefficient of stressinduced shift as follows:

v s. -v r = get (1)

where vr is the position of the stress-free Raman line, vs is the measured position of the Raman line, er is the

stress present (in GPa) and g is the coefficient of stress-induced shift.

The value of g has been measured under hydrostatic conditions in diamond anvil cells by several authors, and

the average value of these measurements is 2.88 + 0.17 cm-1/GPa (ref. 8). For a PCD layer on a Co-WC substrate, it

is, however, more reasonable to assume a bi-axial stress model (ref. 4). With this assumption, and following the

derivation of reference 9, the value of g for the bi-axial model is 1.92 + 0.11 cm-1/GPa, in agreement with reference

6. This value of g was used for the calculations in this paper.

EXPERIMENTAL DETAILS

The samples used in this work were cylindrical drill-bits, consisting of a Co-WC substrate with a PCD layer

sintered onto the top. The average grain size was 10 btm. The cylinders were approximately 17 mm high, 19 mm in

diameter, with a PCD layer of 0.6 to 1 mm thick, depending on the processing. Three different batches of samples

were examined, each batch with a different surface preparation. The first batch of samples had roughly lapped

surfaces, the second batch of samples had finely lapped surfaces and the third batch of samples had surfaces polished

to a mirror finish.

Room temperature (303K) Raman spectra were acquired with a Jobin-Yvon T64000 Raman spectrometer

operated in single spectrograph mode with an 1800 grooves/mm grating and a CCD detector. The 514.5 nm line of

an argon ion laser was used as the excitation wavelength. Spectra were acquired with a confocal micro-Raman

attachment, where the laser beam was focused to a 50 gm diameter spot. A motorized X-Y microscope stage was

also available and used to obtain high spatial resolution measurements. Certain samples were measured in the macro

chamber of the spectrometer, where the diameter of the spot size was - 300 gm. The PCD layer was quite opaque,

thus limiting the penetration depth of the laser to several tens of microns.

High temperature measurements were made in situ with a custom built furnace, fitted into the macro chamber of

the spectrometer. The furnace was mounted on an X-Y-Z adjustable stage with micrometer adjustments. Prior to

heating the furnace was evacuated to a high vacuum, and backfilled with low pressure argon gas to promote a

uniform temperature of the sample. High temperature measurements were only performed on the samples with

roughly lapped surfaces, as these samples gave the best signal-to-noise ratios in the spectra.

NASA/CP--2001-210948 377 August 6, 2001
Fortheroomtemperature measurements utilizingthemotorizedX-Ystage, datapointswereacquiredas

follows:twotraverses weremeasured acrossthediameter ofeachsample, withthepointsineachtraversespaced

1000gmapart,andthetraversesseparated by500gm.Anotherpairoftraverseswerealsoacquiredatrightanglesto

thefirstset.AteachpointaRamanspectrumwasacquired, thusyieldingamapoftheRamanspectrumintheshape

ofacrossacrossthesurfaceofthesample. Forthehightemperature measurements performedinthefurnace, onlya

singlelineofpointsspacedat2000gmintervalswasmeasured.

Thepeakpositionwasdetermined byfittingaVoigtfunctiontothepeakdata.TheVoigtfunctionisa

convolutionfaLorentziandaGaussianlineshape. Certainsamples, especially thosewithapolishedsurface,

werefoundtohaveveryweakdiamondpeaksatcertainpoints.Inthesecases, orwherethepeakwasnotdetectable,

datapointsaremissingonthecorresponding graphs.

AstheexactpositionoftheRamanpeakisimportant indeterminingthenatureandmagnitudeofstresses

measured, thespectrograph wasverycarefullycalibratedbeforeachsetofacquisitions. Forthiswork,the

5460.74AlineofaHgdischargelampwasusedasthecalibrationstandard. Asareferencefortheunstressed

positionofthediamondRamanline,thepositionofthediamondRamanlineasmeasuredonatypeIIadiamondwas

used(ref.3).

ThediamondRamanlineshiftsduringchangesoftemperature (ref.12).Thisshiftwascompensated for

accordingtothisreferencebforecalculatingtheresidual stresses inthePCDduringtheheatingexperiments.

RESULTS

TheEffectofSurfaceProcessing

Theeffectofsurfaceprocessing onthesurfacecompressive stressiillustratedinfigure1.Thetraversesare

shownforonesampleforeachofthedifferent surfacefinishes. Thesearerepresentative ofourfindings. Although

largevariationsinstressoccurfrompointopoint,thereisadefinitetrendshowingaloweringoftheaverage

compressive stresswithsuccessive surfaceprocessing stages. Thesurfacecompressive stressaveragedoverthetwo

traversesfortheroughlylappedsamplesis1.4GPa,forthefinelylappedsamplesit is0.8GPaandforthepolished

samplesit is0.1GPa.

Asisillustratedbyfigure1,therearenoobviousstressdistributions alongthediameter ofthesamplesandthe

variationsinthemagnitudeofthestressarelargeoverelativelyshortdistances. Thestressmeasurements were

repeatedfortheroughlyappedsampleswithalargerdiameter laserspot,i.e.100gminsteadof50gminanattempt

to"average" overalargernumber ofgrainswithasinglemeasurement. Theseresultswereverysimilartothose

presentedinfigure1andsuggest thathereisarandomdistributionoffluctuatingstressvaluesacrossthesurfaceof

thesample.

TheEffectofTemperature

Theeffectoftemperature onthesurfacecompressive stressiillustratedinfigures2and3.Infigure2,the

averagesurfacecompressive stresses fortworoughlylappedsamples(AandD)arereported. Itisobservedthathe

averagesurfacecompressive stressdecreases linearlywithincreasingtemperature. Inthelimitednumber of

experiments conductedit appearsthatoncoolingtoroomtemperature, thesurfacecompressive stressrecoverstoa

valuethatisslightlylessthantheroomtemperature startingvalue.Theerrorbarsonthemeasured valuesrepresent

thestandarddeviationofthestressvalues. Itcanbeseenthathestandarddeviationdecreases withincreasing

temperature, especiallyforsampleD,andagainrecoverstopreviousvaluesoncooling.

Inordertoinvestigatetheapparent lineartrendinstressrelaxationuponheating, andtheinfluenceofhigher

temperatures, adifferent drill-bitwasmeasured to893K.Theseresultsareillustratedinfigure3.Itcanbeseenthat

theinitialroomtemperature compressive stressvalueishigherthantheaveragevalueextractedfromthedataof

figure1fortheroughlylappedsample, butcomparable tothehighervaluesmeasured infigure1.Onheating, the

compressive stressvaluesdecreaseinasimilarwaytothatobservedinfigure2.Thestandarddeviationofthe

measurements decreases withincreasingtemperature toalargerdegreethanfortheresultsreportedinfigure2.

NASA/CP--2001-210948 378 August6,2001
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I I I I I I I I I I

0 2000 4000 6000 8000 10000 12000 14000 16000 18000 20000

Position (lum)

Figure 1. Compressive stress as a function of surface finish across the diameter of three different samples.

Circles (o) correspond to a roughly lapped sample, diamonds (1) to a finely lapped sample and the crosses (+)

correspond to a polished surface.

1.4

1.2

1.0 12.

(5

"-" 0.8

• 0.6

,,,-,

O0

• 0.4

_0 0.2

,0

E 0.0

o

(D -0.2

-0.4

• Sample A

• Sample D

JIL

/

7

I , I , I , I , I

303 503 703 503 303

Temperature (K)

Figure 2: Change of average surface compressive stress with temperature for two roughly lapped samples.

The error bars correspond to the standard deviation of the variations across the surface.

NASA/CP--2001-210948 379 August 6, 2001
n

C3

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if)

if)

0)

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if)

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o

O

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0.8

0.6

0.4

0.2

0.0

• Sample 1 I

,IL

.LL

,LL

I , I , I , I

297 493 693 893

Temperature (K)

Figure 3: Change of average surface compressive stress with temperature for a different type of PCD drill-bit.

The error bars correspond to the standard deviation of the variations across the surface.

DISCUSSION

If it is assumed that the PCD layer on the Co-WC substrate can be treated as a thin film, it is reasonable to

expect a compressive stress in the diamond layer and a tensile stress in the substrate induced by the differential

thermal contraction between the PCD and substrate during cooling from the sintering temperature (ref. 11). If it is

further assumed that the stresses are transmitted uniformly through the diamond layer, then the induced compressive

stress can be calculated to be between 1.5 GPa and 2.7 GPa. These values are comparable to the actual values

measured in this work.

When investigating the effect of temperature on compressive stress, it is clear that a reversible mechanism must

be considered to account for the largest observed changes, with an irreversible mechanism possibly contributing to a

lesser degree. A plausible reversible mechanism is the differential thermal expansion between the PCD layer and the

substrate. Upon heating the differential expansion decreases, leading to a relaxation in the measured stress, with an

increase in stress upon returning to room temperature. The apparent linear relationship between temperature and

compressive stress also supports the suggestion that differential thermal expansion is the dominant mechanism at

work during heating and cooling. The temperature data thus also suggests that the assumption that the PCD layer acts

as a thin film is valid.

There are several possible mechanisms that will account for the observed decrease in surface compressive stress

after annealing, including relaxation due to annealing in the substrate and annealing of the cobalt within the PCD

layer. The data here are not sufficient to make a conclusion as to which of these mechanisms may be operative.

The lapping process is known to introduce a compressive surface layer, primarily due to the plastic deformation

of the surface that occurs when chips are formed (see e.g. ref. 13). The observed surface compressive stress for the

lapped samples is in agreement with this.

The polishing process can be considered to be a combination of chipping and plastic flow taking place on a very

small scale. Frictional processes play a large role in the material removal, with the result that local hot spots are

formed on the diamond surface. These hot spots allow for the plastic flow of diamond and subsequent decrease in

localized residual stresses. If the annealing process is sufficiently localized, it can result in spots of tensile stress

developing.

NASA/CP--2001-210948 380 August 6, 2001
InconclusionitcanbestatedthataPCDtoolisacomplexmaterial, involvingseveral variablesinits

manufacture anduse.It issuggestedhatusebemadeofFiniteElement Modelingtoinvestigatestressdistributions

fromacomplementary,heoretical pointofview.

ACKNOWLEGDEMENTS

TheauthorswishtothanktheNational Research Foundation ofSouthAfricaforfinancialssistance.

REFERENCES

1. Field, J.E. (ed.): The Properties of Natural and Synthetic Diamond. Academic Press, London, 1992, 607-

636.

2. Wentorf, R.H.; DeVries, R.C.; and Bundy, ER: Sintered Superhard Materials. Science, vol. 208, 1980, 873-

880.

3. Ager III, J.W.; Veirs, D.K.; and Rosenblatt, G.M.: Spatially Resolved Raman Studies of Diamond Films

Grown by Chemical Vapor Deposition. Phys. Rev. B, vol. 43, no. 8, 1991, 6491-6499.

4. Ager III, J.W.; and Drory, M.D.: Quantitative Measurement of Residual Biaxial Stress by Raman

Spectroscopy in Diamond Grown on a Ti Alloy by Chemical Vapor Deposition. Phys. Rev. B, vol. 48, no. 4,

1993, 2601-2607.

5. Catledge, S.A.; and Vohra, Y.K.: High Density Plasma Processing of Diamond Films on Titanium: Residual

Stress and Adhesion Measurements. J. Appl. Phys., vol. 78, no. 12, 1995, 7053-7058.

6. Bergman, L.; and Nemanich, R.J.: Raman and Photoluminescence Analysis of Stress State and Impurity

Distribution in Diamond Thin Films. J. Appl. Phys., vol. 78, no. 11, 1995, 6709-6719.

7. Nugent, K.W.; and Prawer, S.: Confocal Raman Strain Mapping of Isolated Single CVD Diamond Crystals.

Diam. Related Mater., vol. 7, 1998, 215-221.

8. Catledge, S.A., et al.: Micro-Raman Stress Investigations and X-Ray Diffraction Analysis of Polycrystalline

Diamond (PCD) tools. Diam. Related Mater., vol. 5, 1996, 1159-1165.

9. Mohrbacher, H., et al.: Comparative Measurement of Residual Stress in Diamond Coatings by Lowincident-beam-angle-diffraction and Micro-Raman Spectroscopy. J. Mat. Res., vol. 11, no. 7, 1996, 1776-

1782.

10. Field, J.E. (ed.): The Properties of Natural and Synthetic Diamond. Academic Press, London, 1992, 428.

11. Lin, T.R, et.al.: Residual Stress in Polycrystalline Diamond Compacts. J. Am. Ceram. Soc., vol. 77, no. 6,

1994, 1562-1568.

12. Herschen, H.; and Cappelli, M.A.: First-order Raman Spectrum of Diamond at High Temperatures. Phys.

Rev. B, vol. 43, no. 14, 1991, 11740-11744.

13. Richerson, D.W.: Modern Ceramic Engineering. Marcel Dekker, Inc., New York, 1982, 260-273.

NASA/CP--2001-210948 381 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

EVALUATION OF CRYSTAL DEFECTS IN (100) AND (111)

SINGLE-CRYSTAL DIAMONDS BY CATHODOLUMINESCENCE

SPECTROSCOPY

Sadao Takeuchi and Masao Murakawa

Department of Mechanical Engineering

Nippon Institute of Technology

4-1 Gakuendal, Miyashiro, Saitama 345-8501 Japan

take@nit.ac.jp Tel: +81-480-33-7624, Fax: +81-480-33-7566

ABSTRACT

The practical application of diamond films synthesized by the chemical vapor deposition (CVD) method to the

field of semiconductors is highly expected. However, even a slight crystal defect or impurities of the material may

lead to a decrease in the carrier mobility of semiconductors. Under such circumstances, the establishment of a

technique for analyzing crystal defects in diamond has been desired. In this study, we use type Ib (100) and

(111) diamonds synthesized by the high-pressure and high-temperature method; i.e., the diamond includes nitrogen

impurities. A crystal defect is introduced into the diamond mechanically using a diamond indenter, followed by the

observation of the CL spectra. The following results are obtained. 1) A broad band A peak that corresponds to

nitrogen-related defects is observed at the wavelength of 400 nrn. 2) A broad band A peak that corresponds only to

crystal defects is observed at the wavelength of 430 nm.

Keywords: diamond, band A, crystal defects, nitrogen concentration

INTRODUCTION

Since the production of high-quality diamond by the ehemical vapor deposition (CVD) method has become

possible, its application in the field of semiconductors is expected. However, even a slight crystal defect may lead

to a decrease in the performance of the semiconductors. Therefore, the establishment of a technique for analyzing

crystal defects in diamond has been desired.

In eathodoluminescence (CL) spectroscopy, emission caused by electron-beam excitation is observed,

through which it is possible to identify crystal defects and nitrogen impurities in diamond (ref. 1). It was reported

that the band A peak centered around the wavelength of 430 nm (ref. 2 to 4), which corresponds to the imperfectness

of the crystal, is broad and that the band A peak has strong correlation with mechanical crystal defects (ref. 5).

However, the relationships among the band A peak, the crystal orientation and the nitrogen concentration have not

been clarified thus far. In this study, type Ib diamond that contained uniformly dispersed nitrogen was used, and

crystal defects were mechanically introduced onto crystal faces having different crystal orientations and nitrogen

concentrations. Subsequently, changes in the position and intensity of the band A peak were observed.

EXPERIMENTAL METHODS

A commercially available type Ib tingle-crystal diamond plate (2 x 1.5 x 0.5 mm) synthesized by the highpressure and high-temperature method, was prepared. The crystal orientations to be observed were (100) and (111)

facets. Minute cracks which do not cause breakage were introduced to the samples by indentation using a singlecrystal diamond indenter (Rockwell C-Scale indenter, the tip diameter of which was 0.2 mm) for 30 min to apply a

NASA/CP--2001-210948 382 August 6, 2001
100Nload. Forobservation ofthe CL spectra, a scanning electron microscope in which a liquid-nitrogen cooling

stage used to cool the samples and the CL observation system were incorporated, was used under the following

conditions: acceleration voltage, 20 kV; sample current, 20 rtA and sample temperature, 80 K.

RESULTS AND DISCUSSION

Figures l(a) and l(b) show scanning electron microscopy (SEM) results for (100)diamond and (111)

diamond, both comaining minute cracks. As shown in the figure, square and hexagonal indentation marks were

obtained for (100) diamond and (111) diamond, respectively. It has been reported that when diamond is fractured

by a spherical indenter, square and hexagonal indentation marks are formed on (100) and (111) facets, respectively

(ref. 6). Therefore, from the shape of the indentation marks obtained in this ,study, the observed diamond surfaces

were identified to be (100) and (111 ) facets

m

a)

:iiiiii!iiiii!iiiiiiii:_iiiiiii_i:i::'::':':'_:'i::'::i:i:iiiiiiii

::::::i:::_:_

:::::::_:_

b)

Figure 1 SEM images of the a) (100) diamond and b) (111) diamond at a surface

portion after indentation.

Figure 2 shows the CL images of the above samples. An emission distribution of 503 nm wavelength,

which corresponds to the H3(nitrogen-vacancy-nitrogen) center structure of nitrogen(tel. 7), is observed. In the

case of (100) diamond, a uniform emission distribution is obtained. On the other hand, in the case of (111)

diamond, a strong beltlike emission is observed, and the indentation mark is found to overlap on the boundary area

of the bdtlike emission. This finding demonstrates thal similar crystal defects are introduced into areas having

different nitrogen concenlrations on (111) diamond;

:::::::::::::::::::::::::::"'_ : ::."

:_::.-:::::: ::::::::::::::::::::::

:" ":_2:::_: " ",".'"_"

a)

i:i:i:!:!;_j::

b)

Figure 2 CLimages (wavelength: 503 nm) of the a) (100) diamond and b) (111) diamond

after indentation.

NASA/CP--2001-210948 383 August 6, 2001
Figure3 shows the emission distribution of the band A peak (430 nm) corresponding to the imperfectness of

the crystals. For both (100) and (I!l) diamonds, strong emissions are observed in spots in the vicinity of the

indentation marks. In addition, in the case of (111) diamond (Fig. 3(b)), band A emission corresponding to the

belt-shaped area with high nitrogen concentration is observed. In other words, the number of crystal defects in the

area where nitrogen is segregated increases.

Figure 4 shows the CL spectra of areas (A)-(F) indicated in Fig. 3. A strong peak of 503 nm wavelength,

which corresponds to the H3 center structure of nitrogen, is observed in the spectra of (A), (B) and ((2) for areas

outside the indentation marks, regardless of crystal orientation, In the spectrum of (C) for the area of (11l)

diamond with high nitrogen concentration, a weak band A peak centered at 400 nm is observed. On the other hand,

band A emission in spots (for areas (D), (E) and (F)) corresponding to pure crystal defects in the vicinity of the

indentation marks has a broad peak centered around 430 nm, regardless of crystal orientation and nitrogen

concentration.

m

Figure 3

":::....'-5_

!8?i:i::""

::;::::

a) b)

CL images (wavelength: 430 nm) of the a) (100) diamond and b) (111) diamond

after indentation.

Figure 4

1,.--.¢

L)

Band A (40Ohm) i_ (C)

•-- mr-- )

tit Band A (430nm)

(D)

I

II

_ _.___ (E)

.. I I I

200 300 400 500 600 700 800

Wavelength (nm)

Measurement results of CL spectrum around indentafiCm mark in Fig. 3 a) and b).

NASA/CP--2001-210948 384 August 6, 2001
CONCLUSIONS

Changes inbandA emissionspectraoftypelb diamond synthesizedby thehigh-pressureand high-temperature

method were examinedby introducingmechanicalcrystaldefects.The resultsindicatedthatthebandA emission

spectracorrespondingto mechanicalcrystaldefectsexhibiteda broadpeak at430 nm forboth (100)and (III)

diamonds. Furthermore,the band A peak thatcorrespondsto nitrogenimpuritiesshiftedto a 30 nm shorter

wavelengthand exhibiteda broadpeakat400 rim.

REFERENCES

1. A.T. Collins: The characterization of point defects in diamond by luminescence spectroscopy, Diamond Rel.

Mater. eel. 1, 1992, 457.

2. I. Kiflawi and A. R. Lang: Linearly polarized luminescence from linear defects in natural and synthetic

diamond, Philos. Mag. eel. 30, 1974, 2.

3. A.T. Collins, M. Kamo and Y. Sate: Intrinsic and extrinsic cathodoluminescence from single-crystal diamonds

grown by chemical vapour deposition, J. Phys. Condens Matter ,eel. 1,1989, 4029.

4. H. Kawarada, ctal.: Dominant free-exaction recombination radiation in chemical vapor deposition diamonds,

Appl. Phys. Lett. eel. 64,no.4, 24, January 1994, 451.

5. S. Takeuehi and M. Murakawa, Thin Solid Films: Synthesis and evaluation of high-quality homoepitaxial

diamond made by the combustion flame method, eel. 311-378, 2000, 290.

6. D.A. Nelson and Jr., A. L. Ruoff: The compressive strength of perfect diamond, J. Appl. Phys., vol. 50, no. 4,

1979, 2763.

7. S.C. Lawson and H. Kanda: Cathodoluminescence from high-pressure synthetic and chemical-vapor-deposited

diamond, J. Appl. Phys. vol. 77, no. 4, 15 February 1995, 1729.

NASA/CP--2001-210948 385 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

VACANCY IN BULK DIAMOND AND DIAMOND FILM:

A MOLECULAR DYNAMICS STUDY

X. J. Hu Y.B. Dai H.S.Shen Z.M.Zhang Y.Z.Wan

State Key Lab of MMCM's, Shanghai Jiaotong Univ.

Shanghai, China, 200030

E-mail:xche@mail.sjtu.edu.cn

X.C.He

ABSTRACT Therelaxation process of adjacent atoms around a vacancy in bulk diamond as well as the movement of a

vacancy in diamond film were investigated by molecular dynamics simulation. The atomic interaction is described by

Tersoff's many-body empirical potential. The results show '_'_' the average relaxation magnitude of the nearest neighbor

atoms of the vacancy is outwardby 0.02rim, which is in a _ greementwith the results obtained by ab-initio method and

duster theory. The relaxation orientation of the nearest atoms around the single vacancy deviates from <111> _ion by

0.5°~3.0 ° before introducing the vacancy and by 0.05o-0.3 ° after introducing the vacancy into the supercelL In addition, if

one vacancy is introduced into the second layer near (001) surface of diamond film at 1800K, the vacancy diffusing from the

second layer to the first layer is observed.

Keywords: diamond, vacancy, relaxation, molecular dynamics

INTRODUCTION

Vacancy is an important kind of defect in diamond, Most properties of diamond like optical and electronic

characterization are closely related to vacancy. However, the nature of vacancies in diamond is complicated

and is not well understood since it is difficult to observe vacancies directly. From the theoretical point of view,

molecular dynamics (MD) simulation is proved to be a powerful technique that provides an opportunity to

observe and measure dynamical phenomena in diamond

While vacancies are introduced into diamond, the localized lattice relaxation inevitably happens. This

relaxation influences the perfection of diamond lattice, which plays a very important role in improving the

conductivity of diamond. Many researchers had investigated this aspect with various methods, but the results

were controversial. Mainwoods (re£ 1) and Watkins (ref,2) concluded that the nearest neighbor atoms around

the single vacancy moved with inward relaxation by ~10%. Beruer and Briddon (re£3), using ab-initio

method and cluster theory, predicted that the nearest neighbor atoms moved with outward relaxation of similar

magnitude. Joubert (ref.4) undertook Car-Parrinello molecular dynamics to calculate the movement of the

nearest and second nearest neighbor atoms, and his result reached a good agreement with the cluster model

calculation. These results were obtained by ab-initio method or duster theory, which need to expend much

calculation time and in which the calculated system is small (refs. 1 to 4). While less accurate than ab-initio

methods, molecular dynamics simulation with empirical potential is invaluable for the treatment of complex

and large system. In this paper, we will use molecular dynamics with Tersoff's many-body potential (ref.5),

which has been particularly successfully in modeling covalent system, to study the relaxation process and

orientation of every adjacent atoms around single vacancy in bulk diamond.

In addition, surface defects play a very important role in the nucleation and growth of diamond so an

understanding of defects at atomic dimensions is highly desired today. Bar-Yam and Moustakis (re£6)

calculated a high concentration of vacancies growing into a diamond surface and concluded that the vacancies

would stabilize the growth of diamond compared to graphite, at temperature and pressure in the graphitestable part of the phase diagram. Therefore, the behavior of the vacancy near the surface of diamond films is

NASA/CP--2001-210948 386 August 6, 2001
veryimportant forunderstanding theCVDmechanismof diamond film. Halicioglu (ref.7) calculated the

formation energy for a single vacancy in the first and second layer of diamond and concluded that the

formation energy of single vacancy in the second layer is higher than that in the first layer. It implied that the

vacancy in the second layer probably move to the first layer. In addition, Mainwood (ref. 1) got the conclusion

by CL experiment that the vacancies were located within the top few Ixm of the samples after the growth

process ended. However, they didn't give the movement of the vacancy at atomic dimensions. Another target

of this paper is to simulate the movement of one vacancy in second layer near (001) surface of diamond film

at high temperature.

CALCULATION METHODS

We used Tersoff's (ref5) empirical many-body potential to describe the interatomic interaction and Verlet

algorithm (ref.8) to solve Newton's motion equations. The detailed molecular dynamics method and the

parameters of Tersoff's potential about carbon can be found in reference 5. The simulation was undertaken in

the canonical ensemble.

For bulk diamond, the temperature was 30OK. The total run time was 30ps with the integration time step

of lfs After the perfect lattice relaxed 5ps, one atom was removed from the center of diamond supercell and

resulted in a single vacancy. In the calculation of the relaxation of atoms around a vacancy, some atoms in the

supercell are always fixed (ref.4, 9). Moreover, because of the limited atoms used in the molecular dynamics,

the effects of surface on the neighbor atoms around the vacancy must be taken into account. To avoid this

spurious effect, two ways can be used: cyclical boundary condition or fixing walls whose thickness equals the

cut-off radius S. The latter method can save calculation time and is more suitable for determining the changes

in the positions of the atoms due to the relaxation (ref.10). In this paper, we used 5ax5axSa cubic (a is

diamond lattice constant 0.357nm) with 1000 atoms, and fixed the atoms in outmost three layers, whose

thickness is 0.267nm and larger than the cut-offradius 0.21nm

For diamond films, the supercell consisted of 20 layers and each layer had 25 atoms. The initial positions

of carbon atoms were on the sites at OK. The atoms in bottom three layers were fixed as the static substrate of

diamond film. Two-dimension periodic boundary condition was exposed on X-axis and Y-axis of upper

seventeen layers, and Z-axis was free. The temperature was 1800K. The total run time was 25ps with a time

step as long as 1Is. After the perfect lattice relaxed 3ps, one atom was removed from the second layer near

(001) surface of diamond film. To record the moving trajectory of a vacancy was more difficult than that of an

atom in diamond. While a vacancy in the bulk or on the surface can move only via its nearest neighbor atoms

by exchanging their position (ref 11), we recorded the trace of the nearest neighbor atoms to speculate on the

moving of the vacancy. In addition, the initial position of the vacancy was considered as a criterion and the

distance between the criterion and the nearest neighbor atoms of the vacancy at every time step can be

calculated. Then, the difference between the distance and the bond length of diamond (0.154nm), called

relaxation distance "d", can be gotten. (In bulk diamond, relaxation distance has the same meaning.) If"d" of

one of the nearest neighbor atoms around the vacancy is smaller than a critical value, we consider that this

atom have been occupied the vacancy. Thus, the vacancy has been moved to the position of the atom.

RESULTS AND DISCUSSION

Single vacancy in bulk diamond

NASA/CP--2001-210948 387 August 6, 2001
Theaveragerelaxationdistanceof the nearest and second nearest neighbor atoms around the single

vacancy in diamond versus time evolution is shown in Fig. 1. The nearest neighbor atoms relaxed outward by

0.02nra, which is in a good agreement with the results obtained by ab-initio method and cluster theory (refs 3

to 4). The average relaxation magnitude of the second nearest neighbor atoms is O.0007nm outward, which is

much smaller than that &the nearest neighbor atoms. In Fig.2, it is displayed that the relaxation of the nearest

and second nearest neighbor atoms and the energy of the system within 40fs before and after introducing the

vacancy. We can see that the energy rapidly increases at the moment of introducing the vacancy. For diamond

has very. strong bond, removal &one atom from the perfect lattice results in four dangling bonds and destroys

the balance of the system. Therefore, the atoms around the vacancy must move to relax the system to a lower

energy state. From Fig.2, it is seen that four nearest neighbor atoms move outward and most second nearest

neighbor atoms move inward the moment of introducing the vacancy. Due to the strong localization of carbon

valence electron wave functions and the fact that the nearest carbon atoms of the vacancy are in secondnearest-neighbor distance, the C-dangling bonds do almost not overlag Those facts result in a substantial

outward relaxation of the nearest atoms around the vacancy in diamond (ref. 12). After lOfs of the appearance

of the vacancy, the second nearest neighbor atoms relaxed inward to their maximum magnitude, the bond

length (L) of the nearest neighbor atoms and their corresponding nearest neighbor atoms is 0.149nm, and the

energy reached its extreme small value. However, the nearest neighbor atoms continue move outward. When

the distance L is larger than OA49nm, the attractive force is gotten between the nearest and second nearest

neighbor atoms. So, the second nearest neighbor atoms move outward and the energy increased. At 3017fs,

the nearest and second nearest neighbor atoms relaxed to their maximum magnitude 0.030nm-0.035nm and

O.003nm-OOOgnm respectively. The maximum relaxation magnitude of the second nearest atoms after

introducing the vacancy is O.O04nm larger than that before introducing the vacancy. Due to the full relaxation,

the energy of the system decreased by 1.8eV. Summarily, according to the relaxation process of the nearest

and second nearest neighbor atoms, we can know that the nearest neighbor atoms have not a tendency to move

inward The second nearest neighbor atoms move inward at the moment of the introducing of the vacancy.

Later, they vibrate outward or inward, and the average magnitude of relaxation is outward by OO007nm,

which is much smaller than that of the nearest neighbor atoms.

The relaxation orientation of the nearest neighbor atoms of the single vacancy was investigated In

diamond, the directions of the four nearest neighbor atoms and the vacancy is along four body diagonals [ 111],

[ 1 1 1], [ 1 1 1 ] and [ 1 1 1 ], For the four nearest neighbor atoms are equivalent, we use <t 11> to represent

four body-diagonals direction. Before introducing the vacancy, the relaxation orientation of those four atoms

deviated from <111> direction by 0.5°~3.0 °, and by 005°--0.3 ° after introducing the vacancy into diamond.

The average angle deviated from <111> is 0.14 °, Thus, the nearest neighbor atoms are restricted to move

along <t 11> direction, as shown in Fig.3_ Due to the extreme bond strength of diamond, significant distortion

of the local tetrahedral atomic geometry would cost too much elastic energy. Consequently, small atomic

displacements occur and a significant Jahn-Teller splitting could not be achieved. Therefore, the Ta symmetry

is nearly conserved, which is in agreement with the results of Zywietz (ref. 12) and is different from that of

cluster study (re£3) predicting a Jahn-Teller distortion from Td to D2dsymmetry. Obviously, the length L will

change after introducing the vacancy. Our result shows that the average length L is 0.150nm_ shortened by

O.O04nm, as shown in Fig3.

The equation of vacancy formation energy in bulk diamond from reference 13 is as follows:

N-1 EN

e_, = EN_ _ N (1)

where e_ is vacancy formation energy, Ey total energy of perfect diamond lattice, EN._total energy of diamond

NASA/CP--2001-210948 388 August 6, 2001
lattice with a vacancy, N the number of atoms in the calculation. According to this equation, we got the

vacancy formation energy 647eV. Zywietz (ref 12) used the same equation, based on the density function abinitio method, and got the vacancy formation energy 6.97eV. Mehaandru (ref.14) removed the bulk

coordination centrally located carbon atom from the C3sH_ and let it bind to the 3-fold site on the (111)

surface of a C22H2_cluster model. The difference between the energy of removal the bulk coordinated carbon

atom to infinity (15.8eV) and the bond strength to a 3-fold surface site (9.8eV) is the vacancy formation

energy 60eV. Bemhole (ref.15) also obtained the vacancy formation energy 7.2eV, based on local density

function pseudopotential. Empirical calculation of Swalin (ref.9) produced a value of 4.2eV for the vacancy

formation energy on the assumption that two bonds form to the surface. That calculation gave small vacancy

formation energy because the calculated stabilization energy for relaxation of atoms around the vacancy is too

high. In this work, if we take the stabilization energy (l.8eV) into account, the vacancy formation energy is

4.67eV, which is in agreement with the result of Swalin.

One vacancy in the second layer of diamond film near (001) surface

The relaxation distance of the nearest four atoms of the vacancy with the evolution of time is displayed in

Fig.4. From Fig4, we can know that at about 15ps the relaxation distance between ava and the vacancy nearly

equals -0.154nm, which shows that a_ has absolutely occupied the site of the vacancy. In other words, the

vacancy diffuses from the second layer to the first layer, which demonstrates the results gotten by Mainwood

(ref.1) that the vacancies were located within top few gm of the films at the end of diamond film growth. We

study the movement of those four nearest neighbor atoms of the vacancy and find that a_ and a_ firstly move

upward and their z-coordinate value increase 0.03nm and 0.02nm respectively. After that, they move

downward. Especially, from 5ps to abom 12.5ps, the distance between avn and the vacancy is about 0.05nm,

while that of a_ is about 0.12nm Obviously, a_ has a bigger chance to move into the site of the vacancy In

fact, a_ and a_ both can move to the site of the vacancy, but which one can locate at the site at last depends

on initial condition of the introducing of the vacancy. At about 12.5ps, a_ begins to move to the place of the

vacancy while a_ moves to its original place. On the other hand, at the beginning of the vacancy's

introducing, avBand a_4 vibrate away from the vacancy by 0.02nm, which is in a good agreement with that of

the nearest neighbor atoms of the vacancy in bulk diamond. When a_ occupies the site of the vacancy, a_

and a_4 move to their original sites, as shown in Fig4. Therefore, when a_ has absolutely occupies the place

of the vacancy, other three nearest neighbor atoms have moved to their original sites. It is obvious that the

diffusion of vacancy from the second layer to the surface of the film can decrease the distortion of diamond

lattice produced by the introduction of the vacancy. Also, in this procedure, the total energy decreases about

3.4eV. This result validates that of Halicioglu (ref.7), who suggested that the diffusing of single vacancy from

the second layer to the surface was energetically favorable.

We investigated the bonding process ofa_ during it moved to the site of the vacancy, as shown in Fig.5. L_,

L2, L3and L4 represent the distance between a_ and a_f4,a_ and a_, a,_ and a_, ava and a_ respectively. The

curves are the change of Lx (x=1,2,3,4) with the evolution of simulation time. Clearly, the lowest energy

diffusion path will be the one that involves the stretching or breaking of the smallest number of bonds (ref.3).

From Fig.5, we can know that during the diffusion of the vacancy, only one bond (L4) was broken and three

bonds were gotten. Therefore, when a_ has absolutely been on the site of the vacancy (from 15ps to 25ps),

a,_ has bonded with three nearest neighbor atoms of the vacancy; so three dangling bonds, among four

dangling bonds produced by the vacancy, are eliminated. Therefore, the lattice distortion produced by the

vacancy is greatly reduced. It implies that our simulation of the vacancy diffusion is energetically favorable.

NASA/CP--2001-210948 389 August 6, 2001
CONCLUSIONS

We calculated the relaxation of adjacent atoms of single vacancy in bulk diamond and the movement of

vacancy in diamond film by molecular dynamics simulation. For vacancy in bulk diamond, the results show

that the average relaxation magnitude of the nearest neighbor atoms of the singte vacancy is outward by

0.02nm, which is in agreement with the results obtained by ab-initio method and cluster theory. The second

nearest neighbor atoms around the single vacancy move inward the moment the vacancy is introduced. Later,

they vibrate outward or inward, and the average relaxation magnitude of them is 0.0007nm. In addition, the

relaxation orientation of the nearest atoms around the single vacancy deviated from <111> direction by

0.5°-3,0 ° before introducing the vacancy and by 0.05°~0.3 ° after introducing the vacancy. For one vacancy in

the second layer of diamond film, the vacancy diffusing from the second layer to first layer is observed at

1800K One of vacancy's nearest neighbor atoms in first layer moves to the site of the vacancy, and other

nearest neighbor atoms move to their original sites after the vacancy diffuses to the surface. It implies that the

difl3asion of the vacancy can decrease the lattice distortion caused by the introduction of the vacancy.

0030

0 025

0.020

ToTs

.o0010

0005

0 0_

second nearest neig_or _om_

/

Io t (_) 3o 4o

Fig.l.The average relaxation distance of tlmenearest

and second nearest neighhor atoms around a single

vacancy in a diamond snpercell

O 1476"_ft,

4g

Im

Fig. 3 the schematicgraph of relaxation orientation

and bond length of the nearest atoms around a

vacancy.

the square represents a vacancy, circle 1 ' and 1

represent the site of the nearest atom before and

after the introduction of the vacancy respectively,

and circle 2, 3, 4 represent the sites of the nearest

atoms of circle 1 after introducing the vacancy into

the supercelL Bond lengths are in nm

' • _ . • v , t _

_=oo2_ * t..:_ '.';I ;

ooi_ .;;i"_ 1 w

-o-oq-_-_-_--i:_--' _.& '_.__]:'o__4

Fig. 2 The relaxation distance of four nearest and

four second nearest neighbor atoms around a single

vacancy in a diamond supercell and the related

energy

goo

g

°°°'

,8%E

s _Ot (ps) _ 20 2_

Fig.4 the relationship between simulation time and

the relaxation distance of the nearest neighbor atoms

of a single vacancy in second layer of diamond film.

an, am, a_ and a_s_ represent four nearest neighbor

atoms of the vacancy in film; a_ n and av_ are in first

layer and a_ and a_r_in third layer of the film.

NASA/CP--2001-210948 390 August 6, 2001
0,20

0.20

E

0.24

c 0.22

"o 0.20

010

-_. 01e

0.14

2.

3.

4.

6.

7.

8.

9.

4 g --..,_.n .,_,-41

I , I , I , I , | ,,,

5 10 t5 20 _5

t (ps)

Fig.5 The interatomic distance between am and

corresponding atoms versus the evolution of

time, the meaning of L1, Lz, L_ and L4 is in text.

References:

1. Mainwood A.: Point defects in nature and synthetic diamond: What they can tell us about CVD diamond.

Phys.stst.sol.(a), 172, 1999, 25~35

Watkins G.D.; Messmer R.P.: Many-electrun effects for deep levels in solids: the lattice vacancy in

Diamond. Phys. Rev. Lett., 32, 1974, 1244

Breuer S.L; Briddon P.R.: Ab initio investigation of the native defects in diamond and self-diffusion. Phys.

Rev. B, 51, 11, 1995, 6984-6994

Joubert D.R; Li L.; Lowther J.E.: Supercell calculation of relaxation at the vacancy in diamond. Solid State

Communications, 100, 8, 1996, 561~563

5. Tersoff J. Modeling solid-state chemistry: interatomic potential for multicomponen! systems. Phys. Rew B,

39, 8, t989, 5566~5568

Bar-Yam Y.; Moustakas TD.; Defect-induced stabilization of diamond films. Nature,342, 14, 1989,

786-787

I-Ialicioglu T.: Calculations for defects formed on diamond surfaces. Thin Solid Films, 228, 1993,293_296

Allen M.P.; Tildesley D.J.: Computer simulation of liquids. Oxford, Clarendon press, 1987,79~81

Anderson A. B.; Grantscharova E J.: Molecular-orbital theory of monatomic and diatomic

substitutional defects as shallow n-type dopants in diamond. Phys. Rev. B, 54, 20,1996,14341 ~14348

10. Dominguez J.; Andribet E.P.; Carmen Perez-Martin A.M., Molecular dynamics study of the relaxation

processes induecd by defects in metals. Surface and Coatings Technology, 83, 1996, 55-59

11. Duan X. M.; Sun D.Y.; Gong X. G: Hypermolecular dynamics simulations ofmonovacancy diffusion.

Computational Material Science, 20, 2001,151-156

12. Zywietz A.; Furthmuller J.; Bechstedt F.: Neutral vacancies in group-IV semiconductors. Phys. Stat. Sol.(b),

210, 1998, 13

13. Puska M.J.: Point defects in silicon, fu'st-principle calculations. Computational Material Science,

17, 2000,365-373

14. Mehandru SP.; Anderson A.B.; Angus J.C: Hydrogen binding and diffusion in diamond. J. Mater. Res.,

7, 1992,689-695

15. Bernhole J.; Antonelli A.; Del Sole T.M.; Y.Bar-Yam, et.al.: Mechanism of self-diffusion in diamond. Phys.

Rev. Lctt., 61, 1988, 2689~2692

NASA/CP--2001-210948 391 August 6, 2001

5

APPLICATIONS OF DLC, NANO-DIAMOND, AND NOVEL

CARBONS

Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

THIN FILM DIAMOND-LIKE CARBON DIELECTRICS

Richard L.C. Wu, H. Kosai, and J. Wrbanek

K Systems Corporation, 1522 Marsetta Drive, Beavercreek, Ohio 45432, USA

Tel. (937) 429-5151, Fax (937) 429-1122, e-mail: rlwal@aol cm-.t

M. Freeman

DARPA TACTICAL Technology

3701 North Fairfax Drive

Arlington, VA, USA

S. Fries-Carr and J. Weimer

Air force Research Laboratory

Wright-Patterson AFB, OH 45433, USA

P.B. Kosel

Department of Electrical and Computer Engineering

University of Cincinnati, Cincinnati, OH 45221

ABSTRACT

Diamond-like carbon (DLC) films possesses the unique properties of high dielectric strength, high

resistivity, low loss, high decomposition temperature, chemical inertness, radiation hardness and good thermal

conductivity. It has been demonstrated that very thin DLC film can be deposited directly onto smooth aluminum

surfaces with good adhesion. These DLC films are amorphous and highly flexible, making them suitable for the

production of wound capacitors.

A manufacturing technique for producing high energy density DLC capacitors has been designed,

developed and demonstrated. A key feature of this technique is that oppositely directly beams of hydrocarbon ions

are produced by two 6 cm x 22 cm linear RF (13.56 MHz) inductively coupled ion guns. This allows DLC films to

be directly deposited on opposite sides of thin capacitor grade aluminum foil in a continuous mode. Pairs of the

DLC-coated aluminum foils are then rolled in a winding system to produce high energy density capacitors.

Key Word: Diamond-like carbon, ion beam technique, and high energy density capacitor

Capacitors are a pervasive technology in every commercial and military application. Millions are used in

the power systems of transportation vehicles such as automobiles, aircraft, and spacecraft. Capacitors often fail due

to increasing environmental temperature and reliability limitations arising from low voltage breakdown strength,

high dissipation factors and possible long-term drift of the dielectric constant. Increased performance and small

physical size have been the main focus of present research into capacitors with DLC dielectric.

DLC films are amorphous and flexible, and can be deposited directly onto aluminum foils, which provide

physical support to the DLC dielectric and serve as the capacitor electrode material. This allows the metal foil to be

thicker than the dielectric film thereby providing capacitors with higher energy density storage and lower power

losses than polymer capacitors of comparable size and construction. Also, since DLC films are thermally stable

above 200 0C, high temperature operation should be possible with these devices.

A schematic of the manufacturing system for producing high energy density DLC capacitors is shown in

Fig 1. The system consists of two 370-liter stainless steel chamber modules that are bolted together. Each module is

fitted with an ion gun for coating one side of an aluminum foil, which is pulled between the modules. Two 6 x 22 cm

linear RF (13.56 MHz) inductively excited ion sources generate various hydrocarbon ions with high ion beam currents.

The use of a computer controlled web-handling system permits efficient production of uniform DLC films on both sides

of capacitor grade aluminum foil in a continuous mode.

NASA/CP--2001-210948 395 August 6, 2001
Ion

8oun_

Figure 1. Dual Ion Beam Deposition System

During deposition, a quadrupole mass spectrometer (Stanford Research System Model CIS 200) is used

with and without the ionizer to monitor the ion and neutral compositions. The ion source can operate either in RF or

current regulated mode. The typical background pressure is 1 x 10 .5 Pa, and during deposition it can rise to the order

of 10 -1 Pa. Typical parameters which are monitored during deposition of DLC films are RF power (275-425 W), ion

energy (75 - 900eV), ion current (100 mA), gas mixtures (H2/CH4 - 2,3,4), and substrate temperature (19-26 o C).

The composition of the DLC films was determined by Rutherford Back Scattering (RBS) and hydrogen

forward scattering (HFS) techniques. The electrical properties of the DLC films were determined by performing

three different types of measurements: (a) capacitance, (b) leakage current, and (c) voltage breakdown. Given the

area of the capacitors and film thickness, the dielectric constants, the resistivities and breakdown strength of the

films were obtained. The dissipation factor (Df) is the fraction of the energy lost to the dielectric when the capacitor

is charged. The Df is dependent on the resistivity and dielectric constant, and is a good measure of the quality of the

material being considered for a dielectric. It is usually measured at a frequency of 1 kHz.

A method was developed for the fabrication of parallel-plate test capacitors on glass and aluminum foil

substrates using two levels of metal electrodes, which were deposited through shadow masks. Initial wound DLC

capacitors were constructed manually to check capacitor performance.

It has been demonstrated that DLC electrical properties are profoundly dependent on the process

parameters. In general, as the ion beam energy is increased, the dielectric constant and dissipation factor are both

increased while the resistivity and breakdown strength of the DLC are decreased. The capacitance values of DLC

capacitors were found to be constant over a wide range of frequency from 10 Hz to 1 MHz with very low dissipation

factors, high DC resistivity and high breakdown strength. The capacitance values were found to be stable up to

250 °C.

The currently developed high energy density DLC capacitors offer a 50 % decrease in size, weight and

volume, and greater than a 50% increase in operating temperature over similar value capacitors built with existing

technology. These DLC capacitors are an enabling technology for the AF More Electric Aircraft, directed-energy

weapons, high power microwave and pulsed power applications, and for commercial applications in utilities, well

drilling, power supplies, aircraft, communication satellites, trains, automobiles and medical instruments.

NASA/CP--2001-210948 396 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FROM DLC TO DLC-SiO_ HYBRID LOW-k DIELECTRICS FOR ULSI

INTERCONNECTS

ALFRED GRILL

IBM - T.J.WATSON RESEARCH CENTER, YORKTOWN HEIGHTS, NY, USA

Low dielectric constant (k) insulators, with k significantly lower than that of presently used

SiO2, are needed for reducing the interconnect capacitances of the continuously shrinking ULSI

chips. The reduced capacitance will improve the switching performances of the ULSI devices and

reduce their power consumption.

Diamondlike carbon (DLC), which has found a variety of applications based on its attractive

mechanical, tribological, and optical properties, is also a dielectric material whose electrical

resistivity can reach values of 1016 £_2 .cm. However, the amorphous carbon materials with

diamondlike properties are characterized by dielectric constants (k) similar, or even higher than

that of SiO2 (k-4). Hydrogenated diamondlike carbon (DLC) films, typically prepared as a wear

and corrosion resistant coating, can be modified by the adjustment of the PECVD deposition

conditions to obtain materials with dielectric constants ranging from 3.3 to 2.7. Incorporation of F

in the DLC further reduces the dielectric constants to values as low as 2.4 1,2

While these are attractive values, the integration of the materials in the ULSI chips imposes

a significant number of requirements that are not easily achieved by the new dielectrics. In order

to improve the integrability and reliability of the low-k materials, a hybrid composition of DLC

and SiO2 has been developed. Carbon doped oxides, or SiCOH films, comprising Si, C, O and H

and deposited by PECVD, have achieved dielectric constant values as lower than 2.83.

One of the principal criteria for the choice of a new interconnect dielectric is its extendibility

to later chip generations, requiring further reduced k values. The reduction of the dielectric

constant of the SiCOH films could be achieved by increasing or introducing porosity in the films.

Such enhanced porosity could be produced by depositing multiphase PECVD films containing at

least one thermally unstable CH phase in addition to the SiCOH phase and annealing the films to

remove the labile phase from the material 4. With proper choice of the organic precursors and

adjustment of the deposition conditions the dielectric constant of the SiCOH films can be reduced

to values below 2.1 s.

All the materials mentioned above have been deposited by RF PECVD in a parallel plate

reactor using a 13.56 MHz rf power supply on Si substrates placed on the powered electrode, thus

being at a negative self bias. The hydrogenated DLC films were deposited from a pure

hydrocarbon, while the FDLC films have been deposited from fluorinated hydrocarbon

(fluorocarbon), or mixtures of the fluorocarbon with hydrogen. The SiCOH films have been

prepared from a helium diluted, ring structured precursor containing all components of the films.

Dual phase films have been prepared by adding an organic precursor the SiCOH precursor and

adjusting the deposition parameters. The films were annealed in helium for 4 hours at 400 °C to

NASA/CP--2001-210948 397 August 6, 2001
removetheunstablefractionsfromthefilms.Thedetailsof thedepositionof eachtypeof

materialcanbefoundelsewhere3'4'6.

Thedielectricmaterialshavebeencharacterizedbydifferenttechniquesintheas-deposited

stateandafterannealingandtheircompositionandphysicalpropertiesarereportedin the

previouslymentionedreferences. Thedielectricconstantsweremeasuredat 1 MHz,onMIS

A1/low-k/Si capacitorsfabricatedon0.002_-cm,n-typeSisubstrates. Breakdownandleakage

currentsweredeterminedfromI-Vcurvesmeasuredonthesamestructures.

WhileDLCfilmsofpracticalinterest,withk<3,arenotthermallystableintheas-deposited

state,theycanbestabilizedbyastabilizationtreatment 1. Thedielectriconstantmayincrease

slightlyafterthestabilizationanneal butcanbemaintainedbelow3.Theleakagecurrenttrougha

stabilizedfilmhavingadielectriconstant k 2.7was10.7A/cm2at0.5MV/cm,sufficientlylow

tomakethematerialusableastheinterconnect dielectric.

Theintegrationof DLCwithcoppermetallizationhasbeendemonstratedfortwolevelsof

Cu wires embedded in three levels of DLC dielectric 7. While significant efforts have still to be

invested in optimizing the integration processes of DLC in the interconnect structure of an ULSI

circuit and characterizing its reliability, the results obtained so far indicate that DLC has the

potential of a low-k interconnect dielectric with Cu metallization in the BEOL of ULSI chips.

The FDLC films appeared thermally at 400 °C, based on mass loss and thickness changes,

however, at 400 °C the fluorine in the films is sufficiently mobile to interacts with the layers in

contact with FDLC, weakening the interface between the layers and causing delaminations 2.

FDLC materials do not appear to have, therefore, the potential of becoming a low-k interconnect

dielectric.

The annealed SiCOH films are thermally stable and their dielectric constants can be reduced

to ultralow values of 2.15. The materials are characterized by leakage currents of about 1.10 .9

A/cm 2 at 1 MV/cm and breakdown fields in the range of 5 - 6 MV/cm. These values indicate that

the films are suitable candidates for the interconnect dielectric of the ULSI circuits. As it is

expected that the PECVD deposited films will have superior mechanical properties to spin-on

films of similar dielectric constants, the PECVD low-k films have the potential to become a

strong competitor to the spin-on low-k films as the interconnect dielectric.

REFERENCES

1. A. Grill, C. Jahnes, V. Patel, J. Ott, J. Hummel, R. Mih, and J. Liu, Mat.Res. Soc.Symp.

Proc., 511 (1998) 223.

2. A. Grill, V. Patel, and C. Jahnes, J. Electrochem. Soc., 145 (1998) 1649.

3. A. Grill, L. Perraud, V. Patel, C. Jahnes, and S. Cohen, Mat.Res.Soc.Symp.Proc., 565 (1999)

107.

4. A. Grill and V. Patel, Mat. Res. Soc. Symp.Proc., 612 (2000) D2.9.1

5. A. Grill and V.Patel, submitted Surf.Coat.Technol.

6. A. Grill, V.Patel, S.A.Cohen, D.C.Edelstein, J.R.Paraszczak, and C.Jahnes, in R.Haveman,

J.Schmitz, H.Komiyama, and K.Tsubouchi (ed.) Advanced Metallization and Interconnect

Systems for ULSI Applications in 1996, Materials Research Society, Pittsburgh, PA, 1997, p.417.

7. A. Grill, Diamond and Related Materials, 10 (2001) 234-239.

NASA/CP--2001-210948 398 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ENHANCED FIELD EMISSION FROM NANOSTRUCTURED CARBON FILMS

Igor Pavlovsky, ILL. Fink, LE Thuesen, Zvi Yaniv

Field Emission Picture Element Technology, Inc.

Robert Espinosa

Microwave Power Technology

ABSTRACT

Nanostructured carbon thin films with enhanced field emission characteristics have been grown on a

variety of substrate materials. The crystalline structure, sp2/sp 3 ratio, and morphology of carbon films

depended on the method and parameters of the deposition. Good adhesion properties for the films on glass,

ceramics, silicon and metals have been achieved. Average distance between the emitters is on the order of

their height, significantly improving the field enhancement conditions over the emission area of the films.

Typical threshold electric field is 1.5 to 2 V, emission current densities reach 10 mA/cm 2 at a field of 3

V/_tm, and over 1 A/cm 2 at a field of 7 V/lam (see Fig. 1). Such low fields can be attributed to high aspect

ratios of the emitters, or low electron affinity of emission areas, depending on the type of carbon material.

The long-term stability graph for the cathodes tested in DC mode in a sealed device is shown in Fig. 2. An

emission current degradation of less than 10% after 2000 hrs operation at constant voltage was observed.

1.4

1.2 ,Z'D

---D-- Pl Ised Mo • //Er

1.0 _,

---X--D, afterP IseTest /

0.8 E

0.6

0.4

0.2

2 3 4 5 6 7 8

Electric Field. V/lain

Fig. 1. I-V curves for a Si-backed cathode

60

50_

40

30

20

10

0 ' " '

0

_>-o-o- -o-.-o..o__>.o--.o-._.o

800 1000 1800 2000

Time, hours

Fig.2. Lifetime test of a cathode in a sealed device

We demonstrated the selective growth of carbon films by patterning a layer of catalytic material or by a

surface passivation technique. Low-cost carbon deposition processes developed at FEPET can be scaled up

to larger substrates (currently the substrate size is ca. 6"). The patterned emitters can find applications in

gated electron sources, field emission displays, and microwave electronics.

Keywords: field emission, carbon cold cathodes, electron guns.

Igor Pavlovsky

FEPET, Inc.

3006 Longhorn Blvd. #107,

Austin, TX 78758 USA

Fax: (512) 339-5021

Tel: (512) 339-5020x131

E-mail: ipavlovs@carbontech.net

NASA/CP--2001-210948 399 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CONDUCTING ION TRACKS IN DIAMOND-LIKE CARBON FILMS

J. Krauser, A. Weidinger, M. Waiblinger, V. Hoffmann

Hahn-Meitner-Institut, Glienicker Str. 100, 14109 Berlin, Germany

Tel.: ++49(0)30 80622673, Fax: ++49(0)30 80623199, email: krauser@hnu.de

C. Trautmann

Gesellschaft ftir Schwerionenforschung mbH, Planckstr. 1, 64291 Darmstadt Germany

B. Schultrich

Fraunhofer-Institut IWS, Winterbergstrasse 28, 01277 Dresden, Germany

H. Hofs_iss, C. Ronning

Universit_it G_ttingen, Bunsenstr_e 13, 37073 G_ttingen, Germany

ABSTRACT

Electrically conducting thin filaments were produced in diamond-like carbon (DLC) films by heavy ion

irradiation. For this purpose, 1 GeV uranium ions were chosen since they provide the largest electronic stopping

power (dE/dx) possible and therefore lead to the highest temperature in the tracks. Due to the high temperature a

transformation of the insulating, diamond-like form of carbon (sp3-bonding) into the conducting, graphitic

configuration (sp2-bonding) occurs. The separation of the tracks from one another is determined by the ion beam

fluence while their length is determined by the thickness of the film. Atomic force microscopy (AFM) was used to

measure the topography and current mapping of the irradiated films. Hillocks of approximately 4 nm height and

conducting channels with a current enhancement of 3 to 4 orders of magnitude were found at the ion impact sites.

The resistivity of the ion tracks is in the range of 40 to 250 _cm.

Keywords: Diamond-like carbon, ion tracks, field emission.

INTRODUCTION

Amorphous diamond-like carbon (DLC) films have been studied extensively with respect to electron field

emission and application as cold cathodes in field emission displays (FEDs) (ref. 1 to 4). It was found that pure

diamond is less suitable for this purpose than defect-rich materials; the latter have lower threshold voltages for

electron emission (ref. 5). For field enhancement due to a large aspect ratio (length to diameter), long and thin

conducting filaments embedded in an insolating film and reaching from the substrate to the surface are expected to

be optimum. However, no satisfying realization has been found yet. Carbon nanotubes, with their extremely large

aspect ratio, have been applied for field emission showing very low threshold fields and high emission currents

(ref. 6 to 7). Nevertheless, long-term stability of the emission sites is limited due to continuous burning-off (ref. 8).

In this paper we present studies on conducting ion tracks in DLC films which might be useful for field emission

applications.

SAMPLE PREPARATION

The DLC films were produced by either ion beam techniques (Univ. G6ttingen) or by plasma deposition with

magnetic filtering (filtered arc method, FhG Dresden) on highly doped silicon substrates. In both cases, the C ions

were implanted into the growing film with an energy in the order of 100 eV. This is the optimum ion energy to

create diamond-like carbon by "subplantation". Such films are amorphous, contain 70 - 80 % sp 3 bonds and have a

high resistivity. In the present case the films are 50 nm (FhG Dresden) and 100 nm (Univ. G6ttingen) thick.

NASA/CP--2001-210948 400 August 6, 2001
ION TRACK FORMATION

The ion irradiation of the DLC films was performed at the heavy ion accelerator facility UNILAC of GSI

Darmstadt with uranium projectiles moderated down from 2.7 GeV by an aluminum foil to an energy of ~ 1 GeV

or 4.2 MeV/amu (amu = atomic mass unit). The irradiation dose for the data presented here was 1X101° ions/cm 2.

Figure 1 illustrates the evolution stages of the track formation process (ref. 9). The high energetic U ions hit the

sample which is at room temperature under a normal incidence (fig. la). They pass the DLC film and are stopped

deep inside the Si substrate (~ 55 _m). The first step in the ion track formation is the energy transfer from the high

energetic ion to the electrons present along its path through the DLC layer (electronic stopping). The energy

deposition is uniform and amounts to approximately 28 keV/nm. It occurs within a diameter of a few nanometers

forcing the atoms to move and to rearrange again (fig. lb). Due to the high energy deposition the temperature

along the ion trajectories increases up to several thousand degrees. As a result the material melts and is

transformed from insulating diamond-like (sp3-bonding) to conducting graphite-like carbon (sp2-bonding) leading

to thin electrically conducting channels embedded in an insulating matrix (fig. lc). The hillocks are formed since

material flows out of the region of the ion track during the hot stage of the track formation process. The outflow is

driven by the density reduction during the transformation from the diamond-like (P ~ 3 g/cm 3) to the graphitic

(P ~ 2.3 g/cm 3) phase.

b

high energetic ion

e.g. uranium, 1 GeV

I

:¢:g::g_:a:l:_:lt:k:_:5:::e::._

• . ; '_'a¢/

.--°%. _ • ,. °.o. _ _ •

, .:v'_... ° ,2_'_." _ .,...,* _,,_,._,-. ,

DLC-film (e.g. 50 nm)

T _ 5000 K

, :,:_.. :..: :.: i..a_,o_, _._,. hillock 3 - 5 nm

x ._-la,._,.-. _. _ .:.-. ,. conducting ion track

Figure 1: Evolution of conducting ion track formation (ref. 9).

Experimentally it is found (ref. 10 to 12) that continuous tracks are only formed if the energy deposition along

the ion path exceeds a critical value, which depends on the target material. Since the energy loss per path length or

stopping power dE/dx increases with the ion mass and has its maximum around 1 to 5 MeV/amu, only high

energetic heavy ions form continuous ion tracks. This is demonstrated in fig. 2a, where a TRIM (ref. 13)

calculation of dE/dx for different heavy ions is shown as a function of the beam energy. It can be seen that the

stopping power dE/dx falls off rather steeply below 1 MeV/amu. Thus low-energy heavy ions are not suited for the

formation of continuous ion tracks.

NASA/CPlC, I2001-210948 401 August 6, 2001
Figure2bshowstheevolutionofthedamagemorphology independence ofthestoppingpowerasit is

proposedbyToulemonde etal.(ref.12).Actuallythisworkisdoneforoxidematerialslikee.g.Y3FesO12, butwe

assumethatheseresultscanalsobeadoptedtotheDLCfilmsunderinvestigationhere.Thisisalsosupportedby

theworkofPawlak(ref.14),wheretheiontrackformationiamorphous carbonisstudied. Atlowvaluesof

dE/dx,smallspherical defectsappear whichareseparated fromeachother(regimeII).WithincreasingdE/dxthe

damagedregionsstarttooverlapformingsomewhat extendedcylindrical defects(regimeIII).Butthesecylinders

stilldonotformacontinuous track.ForstillhighervaluesofdE/dx,thesecylindrical defectsgrowfurther(regime

IV)andfinallybuildalongcylinderwherethedamageishomogeneous (regimeV).Formostmaterials

investigated byToulemonde etal.continuous trackformationstartsatdE/dxvaluesgreaterthan20keV/nm.

Therefore, fromfig.2aitisclearthatheavyionbombardment oftheDLCfilmwithAuorUataspecificenergy

between2and5MeV/amuisthebestchoicetocreatecontinuous iontracksinthismaterial. In(ref.15)

conductingiontracksinDLCwerefoundinspiteofirradiationwithrelativelyightXeions(1.1MeV/amu), but

theconductivity ofthetrackswasmuchpoorerthaninthepresent caseinagreement wihtheargumentation

showninfig.2.

35

_3o

25

K 20

_ 15

©

10

3

o

2 5

0

0

• i i

Carbon

I _ I _ I _ I

2 4 6 8

specific energy (MeV/amu)

u

Xe

Ar

i

10

©

a b

IV

___ III

_) _) _ II

damage morphology

Figure 2: TRIM (ref 13) calculation of the electronic stopping power dE/dx versus specific beam energy for

different ions in a carbon target with p = 2.9 g/cm 3 (a) . The corresponding damage morphology is illustrated in

(b): The classification in regimes from H to V is based on the work of Toulemonde et al. (ref 12) and links the

defect creation process to the deposited energy in the film (see texO.

AFM-MEASUREMENTS

With regard to resolution and the possibility to measure the topography and the conductivity of the sample

simultaneously, atomic force microscopy with a conducting tip is the best choice for visualization of the ion tracks.

In the following we present AFM measurements recorded in contact mode with an air AFM using a highly doped

(0.01 - 0.025 f_cm) silicon cantilever with a tip coated with polycrystalline conducting diamond. The current

mapping was done with a fixed voltage of 2.5 V between the tip and the substrate in case of the 50 nm thick films

and 5 V for the 100 nm thick films.

Topography

Figure 3 shows the two- and three-dimensional image of the surface topography (1 X 1 _m 2) of a 50 nm thick

DLC film irradiated with 1X101° ions/cm 2 uranium ions at ~ 1 GeV. Within experimental uncertainties the number

of bright spots in the left picture corresponds to the applied ion fluence (100 ions per _m2). Hillocks with a few nm

NASA/CPlq2001-210948 402 August 6, 2001
inheightwhicharetheresultofthetrackformingprocessdescribedaboveareseenattheionimpactsites.The

apparent width(FWHM)ofthehillocksis17nmasshowninFigure4b.Essentiallythisreflectsnothetrue

dimensionfthematerial throw-off butreflectstheconvolutionftheAFMtipandthehillock.

i:i:5_i:i:i

• t,

1 _m

Figure 3: Topography of a so nm thick DLC film (I X I tim 2) irradiated with uranium ions of I GeV and a dose

of IXIO 1° ions/cm 2 in two-dimensional and three-dimensional view. The hillocks have a height of 2 to 5 nm and

are due to the outflow of material from the ion tracks during the hot stage of the track formation process.

The topography image in fig. 4a is taken from the same film as in fig.3 just recorded with a higher

measurement resolution.

iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii

__._._._._._._._._._._._._._._._._._._._._._._._._._._._._._._._._._._._._._._._._

iiiiiiiiiiiiiiiiiiiiiiiiiiiiiii

.°°..,***°..°..°°..°***°***.°*.****.....°.°°.°...*......°*..... ...°°°°,°,-,,,,

...................... • ............. ,,,,,

..................................... • ,,,

°, .......................................

______%_____________________*......._.__________________________________________.

,,,,,,., .................................

• ........................................

................................... • .,°°.

.................................. • ,,,,,,

___;_;_;__

T

b_

0.00 53.01

X[nm]

a b

Figure 4: AFM image of the topography (a) of a 50 nm thick DLC film irradiated with uranium, 1XIO 1°

ions/cm 2 at 1 GeE A topographical profile (b) is made through one of the impact sites (black line in (a))

showing a width of17 nm (FWHM) and a maximum height of approximately 3.3 nm.

NASA/CPlq2001-210948 403 August 6, 2001
Current

ThecurrentfromthetiptothesubstratethroughtheDLCfilmwasmeasuredinthecontact AFMmode

together withthetopography. Figure5showsclearlythatforeachillockseeninthetopography (fig.5a)abright

spotappearsinthecurrentmapping(fig.5b).Thewidthofthecurrentpeakis16nm,i.e.inthesamerangeasthe

widthfoundforthehillocks.However, notethatheapparent widthagainisaconvolutionfthewidthofthetrack

andtheAFMtipcurvature. Therealdiameter oftheiontracksisestimatedtobeapproximately 10nm(ref.12).

a

iiiiiii:i:iiiiiiiii:;iiii:::iiiiiii::iiiiiiiiiiii:!iiiiiiiiiiiii

::::::::::::::::::::::::::::::::::::::::::::::::::::::::

:_;:'_/gi:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i&',i_

b

7

0.00 64.15

X[nm]

c

Figure 5: AFM images of the topography (a) and the current mapping (b) of a 50 nm thick DLC film irradiated

with uranium, 1XIO 1° ions/cm 2 at 1 GeV. A current profile (c) is made through one of the hillocks (black line in

(b)) showing a width of16 nm (FWHM) and a maximum current ofl.56 nA flowing from the AFM tip to the

substrate (tip-substrate voltage = 2. 5 V).

Figure 6 shows a three-dimensional view of the current (fig. 6a) and a current versus voltage curve (fig. 6b) for

a single ion track of a 100 nm thick DLC film irradiated with uranium, 1)<101° ions/cm 2 at 1 GeV.

OO,'fONf

c ".:.:.:..:.:

!! i":!i!i!i! ili!!iiii::ii:.

10 _

100

10 _

10 _

10 _

0.01

b

i

DLC-fihn

irrad, with uranium (1 GeV)

off track

o ocooo

Figure 6: Three-dimensional current image (a) (tip-substrate voltage = 5 V) of a 100 nm thick DLC film

irradiated with uranium, 1XIO _ ions/cm 2 at 1 Ge V,,and currentvoltage curve (b) for a single track (AFM tip on

top of the track). For comparison the currentvoltage curve in the off track position is also shown.

Again each of the current spikes corresponds to an ion track. If the AFM tip does not contact a track the

measured current is practically zero, indicating the overall noise level of the unirradiated parts of the sample. For a

spot on the ion track the current through the track is a function of the applied voltage. In fig. 6b these two cases are

NASA/CPlC,12001-210948 404 August 6, 2001
termedas"offtrack"and"ontrack",respectively. Thedifferenceinthecurrent betweenthistwotippositions

exceeds3ordersofmagnitude. Assuminganiontrackcrosssectionof100nm2,themeasuredcurrentat5Vgives

acurrentdensityof8×103A]cm2andthuswith100nmasthelengthofthetracktheresistivityP at this selected

point is 62.5 _cm. The different P values found for this kind of samples is in the range of 40 _cm to 250 _cm.

CONCLUSION

With heavy ion irradiation it is possible to form electrically conducting ion tracks in an insulating diamondlike matrix. This was demonstrated in this work for 50 nm and 100 nm thick DLC films which where bombarded

with 1 GeV uranium ions. AFM measurements with a conducting tip show that there is a clear correlation between

the topographic surface image and the current mapping confirming that the current only flows at the impact

positions of the ions. From theoretical considerations and from experiments in other systems it can be assumed that

the tracks go straight through the film and have a diameter of several nanometers. The length of the tracks is

determined by the film thickness and can be extended up to several micrometers. The energy of the ions can be

chosen such that the electronic energy loss per unit length along their path through the film remains approximately

constant even for micrometer-thick films. The average separation of the ion tracks from one another is determined

by the ion fluence and can be adjusted in such a way that the emission sites in field emission applications have the

required density. The field enhancement factor for a conductor with 10 nm in diameter and a length of 100 nm

embedded in a dielectric medium is still fairly low. For a significant improvement of the field emission properties,

thicker DLC films, which then yield a larger aspect ratio of the tracks, should be used. We are presently extending

our work in this direction.

REFERENCES

1. Bonnof, A.M., John, R, Kawarada, H., Lettington, A.H., Sitar, Z., Zachai, R. (Eds.): Diamond Relat.

Mater 8, 1999, 1

2. Papers from the 11 th Int. Vacuum Microelectronics Conference, J. Vac. Sci. Technol. B 17, 1999, 506-701

3. Gr6ning, O.: Field emission from DLC films. Appl. Surf. Sci. 111, 1997, 135-139

4. Amaratunga, G.A.J., Silva, S.R.R: Nitrogen containing hydrogenated amorphous carbon for thin-filmed

field emission cathodes. Appl. Phys. Lett. 68, 1996, 2529-2531

5. Zhu, W., et al.: Defect-enhanced electron field emission from chemical vapor deposited diamond. J. Appl.

Phys. 78, 1995, 2707-2711

6. De Heer, W.A., Chatelain, A., Ugarte, D.: A carbon nanotube field-emission electron source. Science 270,

1995, 1179-1180

7. Rinzler, A.G., et al.: Unraveling Nanotubes: Field Emission from an Atomic Wire. Science 269, 1995,

1550

8. Monteiro, O.R., et al.: Microstructure and electronic emission properties of films prepared from singlewall and multi-wall nanotubes containing powders. Appl. Phys. A 71, 2000, 121-124

9. Schiewitz, G.: private communication

10. Trombello, T.A.: Distribution of damage along an MeV ion track. Nucl. Instr. Methods B 83, 1993, 508-

512

11. Klaumtinzer, S., Gutzmann, A.: Effects of High-Energy Heavy Ions on Amorphous Materials. Nukleonika

39, 1994, 125-140

12. Toulemonde, M., Bouffard, S., Studer, F.: Swift heavy ions in insulating and conducting oxides: tracks

and physical properties. Nucl. Instr. Meth. In Phys. Res. B 91, 1994, 108-123

13. Ziegler, J.F., Biersack, J.B., Littmark, U.: The stopping and range of ions in matter, Pergamon, New York,

1985

14. Pawlak, F.: Irradiation du carbone deu6r6: un mat6riau de transition. PhD Thesis, Univ. de Caen, 1997

15. Waiblinger, M., et al.: Electrically conducting ion tracks in diamond-like carbon films for field emission.

Appl. Phys. A 69, 1999, 239-240

NASA/CPlC,12001-210948 405 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FIELD ELECTRON EMISSION FROM NANOSTRUCTURE

DIAMOND ON POROUS SILICON

W. L. Wang, K. J. Liao, C.G.Hu, J.Xu

Department of Applied Physics, Chongqing university, Chonqging 400044, R R. China

ABSTRACT

The field emission properties of nanostructured diamond on porous Si substrate were investigated. The

porous silicon was prepared by chemical etching. The porosity was estimated to be 75% by weight method. The

thickness of the porous layer was about 10_tm. The films were grown by hot fila_nent chemical vapor deposition.

The experimental results showed that the surface structure and post-treatment have a considerable influence on

the field emission behavior of nano-diamond films. The threshold electric field was found to be 1.5MV/m after

annealing at 500°C for 6h, but 2.8MV/m at an annealing temperature of 350°C for 3h. This may be ascribed to the

elimination of some traps due to the annealing treatment. The results obtained also showed that the threshold

voltage was decreased with increasing the thickness of the porous layer. This is due to increasing surface

roughness.

Keywords: Field emission; Diamond films; Porous Silicon; CVD; Nanostructured diamond.

INTRODUCTION

With the development of flat panel displays, it becomes more and more important to approach new types of

cathode materials with a high emission current under a low eclectic field. In the last few years, chemical vapor

deposited (CVD) diamond was considered as a possible cold cathode materials due to its chemical intentness, high

thermal conductivity, negative electron affinity, and compatibility to conventional microfabrication techniques. The

experimental studies have shown that it has an emission current of 10mA/cm 2 under a low electric field of 3MV/m,

1-3 orders of magnitude lower than that of ordinary materials. Generally, electron emission from CVD diamond

films was thought to originate from defects or surface states below the conduction band or from the conduction band

itself (1 to 4)

In this paper, field emission properties of CVD nanostructured diamond grown by hot filament chemical vapor

deposition were investigated. The porous silicon was used as a substrate material. Experimental studies showed that

the surface structure has a considerable influence on the electron emission of diamond films.

EXPERIMENTAL DETAILS

The porous silicon was prepared by

chemical etching. The mirror-polished n-type

silicon wafers with resistivity of 0. l_)-cm were

firstly etched by NaOH solution in order to

remove damaged layer on silicon surface, then

cleaned by dilute HC1 and deionized water.

There were two kinds of chemical etching

solution i, e : HF: HNO3:H20=1:3:4 and HF:

HNO3:H20=l:2.5:3. Etching time was about

30min. The porosity was estimated to be

75% by weight method. The thickness of the

porous layer was about 10pm.

Diamond thin films were prepared by a

conventional hot filament CVD apparatus (5).

The porous Si substrate was ultrasonically

cleaned in acetone for 15min and subsequently

in a methanol bath for 5 min. After a short

rinse in the methanol, the substrate was placed

on a graphite holder. The filament temperature

Figure 1. SEM image of nanocrystalline diamond films

on porous Si and flat Si (a) fiat Si; (b) porous Si.

NASA/CP--2001-210948 406 August6,2001
duringgrowthwas2100°C, andsubstrate temperature was720°C. Themethaneconcentration inhydrogenwas3.5%

atthepressureof4xl03Paandflowof200sccm. Thefilmthicknesswasabout200nm. Fig.la-bshowstheSEM

imageofdiamondfilmsonflatandporoussilicon. Thefilmsontheporoussiliconconsisted offinegrainswithout

crystal facets.

Thefieldemissionmeasurements werecarriedoutinanion-pumped chamber atanominal pressureof10aTorr.

Tocollectelectrons, apositivelybiasedITOcoatedglasselectrode wasmounted100gmabovethediamondfilms.A

biasof0-1100Vwasappliedbetweenthesampleandtheanodewithanareaof3mm2.TheI-Vmeasurements were

conducted withakeithley237-source masurement unit.AgandAuwasusedasohmiccontactsofanodeand

cathode, andaresistorf5M_wasconnectedothecathodetoprotect themeasurement equipment.

RESULTS AND DISCUSSION

The field emission current-voltage (I-V) plot of diamond films on the porous silicon is shown in Figure 2. The

I-V characteristics of a diamond coated porous silicon were measured at different annealing temperatures. The

emission current stability was considerably improved after the annealing treatment. The threshold voltage decreased

with the increasing annealing temperatures. The threshold voltage was about 150V when the annealing temperature

was 500°C for 6h, but 280V after annealing at 350°C for 3h. The annealing treatment was performed in a vacuum of

10-* Torr.

200

< !60

,_ i20

80, <.2

"_ ,it

.m

0

0 t00 200 7_00 40,0 509 600

,A>pl_ed Voltage ( V )

Figure 2. I-V curve of field emission for

nano-diomond films on porous Si at different

annealing temperature: (a) g00°C for 6h

(b) 410°C for 3h (c) 3g0°C for 3h

200

c t60

.o

t20

b,

co 80

<3

40

.,._#

Figure 3.

f /

• -_- _ J i L

150 300 45i? 600 750

Apl_lied Voltage ( V )

Field emission I-V plot; (a)

nano-diamond coating on porous Si; (b)

diamond films on fiat Si substrate

Figure 3b shows the field emission I-V curve of polycrystalline diamond films on the mirror-polished silicon

substrate. Fig.3a shows the electron emission of diamond coated porous silicon without annealing treatment. It is

quite obvious that there is a high threshold voltage for polycrystalline diamond films as compared with the films on

porous silicon, which is about 450V. This is because the two materials have totally different morphologies as

evidenced by the SEM images in Fig. 1.

Figure 4 shows the changes of threshold voltage with the thickness of porous Si layer. Experimental results

showed that the threshold voltage was decreased with increasing the thickness of the porous layer. The threshold

voltage of the field emission decreased from 332 V to 283 V when the thickness of the porous layer increased from

5bu-n to 30btm. Virtually, the roughness of Si substrate surface was increased with increasing porous layer.

Micro-Raman spectra revealed that there was a small peak of nanocrystalline diamond centered at 1200 cm -1

and two bands centered at 1390 cm -1 and 1590 cm -1 of amorphous carbon phases (6) as shown in Fig 5b. The peak of

nano-diamond was enhanced, and the bands of disordered carbon phase became narrow with increasing treating

temperature as shown in Fig.5c-e. This implies that some structure of the films is changed with the annealing

treatment. These changes mainly concentrated on the interface between the porous silicon and nano-diamond. On the

other hand, porous materials easily absorbed some gas, which have effused during the annealing, Fig.5a shows

NASA/CP--2001-210948 407 August 6, 2001
RamanspectrumofdiamondfilmsonflatSi.Thesamples fromthoseshowniFig.5c-ewereusedfortheemission

measurements showniFig.2a-c, andsamples inFig.5a-bwerethesameasthatinFig.3a-b. Itisapparent thathere

isaquitedifferent qualityandstructureinthetwokindsofthefilms,so,thefilmsdeposited onporousSibehave

differentlyfromthoseontheflatSiinfieldemission

35O

325

31111

275

t

250

[)

4-

+

-4-

I0 20 30

Thickness of Poroils Layer (IX in)

)

I I

I

"\ \ \1

01)0 1201 1400 1600 18110

ll.aman Shift (cra")

Figure 4. Changes of threshold voltage with

the thickness of porous layers.

Figure 5. Micro-Raman spectra of diamond films.

(a) Diamond films on flat Si; (b) Nano-diamond films

on porous Si without annealing; (c) after annealing at

500 ° for 6h; (d) 410°C for 3h (e) 350°C for 3h

Figure 6 shows the effect of annealing temperature on

measurement are the same as that in Fig.5. The stress in the

films was estimated by Raxnan spectra and the measured

curvature of the substrate. From Fig6, the stress was

decreased with increasing annealing temperature. This

suggested that the impurities or vacancies were decreased

with increasing annealing temperature (7).

As mentioned above the experimental results, the

mechanism of field emission from nanocrystalline diamond

films on porous silicon can be proposed. These are two kinds

of cases to be effect on the properties of the field emission.

Firstly; there is an extraordinary rough surface on porous

silicon. The emitters may consist of Si tips with a thin

nano-diamond coating. Thus the emission mechanism is

geometric electric-field enhancement at a high aspect ratio

conductive cone. The emission current density, J, is related to

the local electric field E at the emitting surface by the Fowler

Nordhein equation,

internal stress in the films. All the samples used in the

8

Z

4

0 100 200

• | I . i , i

300 400 50(}

Annealing Temperature ( oC )

j = aE2e b/e (1) Figure 6. Internal stress in the nano-diamond

films with annealing temperature

where a and b are constants. However, this emission model can explain a part of the experimental results.

On the other hand, the thin nanocrystalline diamond coatings on porous silicon contained a great number of

defects and impurities as shown in Fig 5-6. This suggested that there is a high concentration of donors and shallow

traps in the thin films. The model assumed a rather high carrier concentration in the thin nano-diamond films to

establish a thin depletion layer. The electrons from high doped porous silicon were tunneling into a conduction band

of thin diamond layer through the thin depletion layer at the Si-diamond interface, and transported electrons to the

NASA/CP--2001-210948 408 August 6, 2001
surface, wheretheyareaccelerated toenergiessufficient tobeescaped intovacuumundernegativeelectionaffinity

conditions.

FromFig.2.thereaxetworegionsinI-Vcurve.Thecurrent belowthesaturationwascorresponded tothe

space-charge-limited currnt densitylaw.Thecurrent saturation regionwascausedbyemissionfelectrons itoa

vacuumfromthesurfacestates, whichdepended onthesurfacepotential height. Athighervoltage, rapidincreasein

current isduetothegeometric electric-fieldnhancement. Theannealing treatment mayeliminatesomeofthetraps

inthefilms.Thus,thethresholdvoltageandinternal stressweredecreased withincreasing annealingtemperature

( seeFig5-6). Another possibilitystheproductionfgraphitization orconductive particlesonthesurfacedueto

hightemperature treatment. Tofurtherunderstand themechanismoffieldemission, it isnecessary furtheranalyze

theemissioncharacteristics accordingtothemodelandthesolvingofPoisson's equationforthebalancedpositive

chargeinthethinfilmsindetail.

CONCLUSION

Thefieldemissionproperties of poroussiliconemitterwiththinnano-daimond layerswereinvestigated.

Experimental resultsshowedthatthethresholdvoltageofelectronemissionwasdecreased withincreasing annealing

temperature andthickness ofporouslayer.Thismaybeascribedtotheeliminationofsometrapsinthefilmsand

increasing surfaceroughness. Themodel ofelectronemissionthroughtethindepletionlayerwasusedtoexplain

theobserved electronemissionatlowelectricfields.Further workneedstobedoneforthemodel fromexperimental

studiestotheoretical calculations.

ACKNOWLEDGMENT

The work was financially supported by the National Natural Science Foundation of China under Grant No.

19904016.

REFERENCES

1 Ding M.Q, Gruen D.M, Krauss, et al: Studies of field emission from bias-grown diamond films.

J.Vac.Sci.Technol. B 17(2),Mar/Apr. 1999, 705-709.

2 Xu N.S, Tzeng Y and Lathaxn R.V, Similarities in the "cold' electron emission characteristics of diamond coated

molybdenum electrodes and polished bulk graphite surface. J.Phys.D: Appl. Phys. 26, 1993, 1776-1780.

3 Zhu W, Kochansky G.R Jim S: Low-field electron emission from undoped nanostructured diamond. Science,

282, Nov. 1998, 1471-1473.

4 Frolov V.D, Karabutov A.V, Pimenov S.M, et al. Electron properties of the emission sites of low-field emitting.

Diamond Relat.Mater. 9, 2000, 1196-1200.

5 Wang W,L, Polo M.C, Sanchez G, et al. Internal stress and strain in heavily boron-doped diamond films grown

by microwave plasma and hot filament chemical vapor deposition. J Appl. Phys. 80(3),Aug. 1996,

1846-1850.

6 Liao K.J, Wang W.L, Feng B. Nucleation and initial growth of diamond by biased hot filament chemical vapor

deposition. Acta Phys.Sinica. 47(3),Max. 1998, 514-519.

7 Wang W.L, Liao K.J, Gao J.Y, et al. Internal stress analysis in diamond films formed by d.c. plasma chemical

vapor deposition. Thin Solid Films. 215, 1992, 174-178.

NASA/CP--2001-210948 409 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ELECTRON FIELD EMISSION PROPERTIES OF NANO-DIAMONDS

SYNTHESIZED BY CHEMICAL VAPOR DEPOSITION PROCESS

I-Nan Lin a, Gia-Ming Lin, and Tong Hsu

Department of Material Science and Engineering, "Material Science Center, National Tsing-Hua University,

Hsin-Chu(300), Taiwan 300, ROC

ABSTRACT

Field emission devices (FED) with high emission current density have been attained in metal tip I and

silicon tip arrays 2. They have great potential for applications as electron emitters in flat panel displays and

have attracted thorough investigations. N-type diamond films had been reported possessing better field

emission properties than the p-type diamond films 3, and nitrogen-doped films perform better than the

phosphorous-doped. However, how the incorporation of nitrogen species into the diamond films modified

the characteristics and field emission properties of the films is not yet completely understood. In this paper,

the modification on nucleation behavior, morphology and the field emission properties of the diamond

films due to incorporation of nitrogen and the application of bias voltage was systematically investigated.

The possible mechanism is discussed.

Diamond films were grown on silicon substrate by a microwave plasma enhanced chemical vapor

deposition (MPECVD) method 4, using a ASTeX 5400 system. The diamonds were directly nucleated on

mirror smooth silicon or gold/silicon surface using a -170 V bias voltage for 9-28 min and then deposited

without bias voltage for lh. In addition to the CH4 and H2 gases used, nitrogen-species were incorporated

into the diamond films by using urea vapor, which was maintained at 23°C. The total pressure and

microwave power were controlled at 70 torr and 2500 Watts, and the substrate temperature was maintained

at around 900°C.

Nano-diamonds, doped with nitrogen, were successfully synthesized by using urea/methanol saturated

solution as nitrogen source and in-situ application of negative bias voltage in growth period. SEM and

Raman spectroscopic examinations reveal that increasing urea/methanol ratio in gas mixture markedly

influences the field emission properties of diamond films, but insignificantly alters their structure and

morphology. The diamond films thus obtained possess good electron field emission properties, (E0= 2.35

V/gm, Je=30.2gA/cm -2 , and effective work function 0e=0.028 eV). Converting the grains of diamond

films from submicrons into nano-size (-50 nm) via the application of bias voltage in-situ further improves

these field emission properties. The turn-on field decreases to E0'=1.88 V/gm, the field emission capacity

increases to Je'=l18.2 gA/cm 2 (at 6.5 V/gm applied field) and the effective work function decreases to

0e=0.017 eV.

1)C. A. Spindt, I. Brodie, L. Humphrey, and E. R. Westerberg, J.Appl. Phys. 47, 5248(1976).

2) G. G. R Van Gorkom and A. M. E. Hoeberechts, J. Vac. Sci. Technol. B 4, 108(1986).

3)M. W. Geis, J. C. Twichell, N. N. Krohn, and T. M. Lyszczarz, Appl. Phys. Lett. 68, 2294 (1996).

4)j. S. Lee, K. S. Liu andI. N. Lin, Appl. Phys. Lett. 67, 1555(1995).

Keywords: Nano-diamonds, Field Emission, Chemical Vapor Deposition

Corresponding author: I-Nan Lin

101 Kuang-Fu Rd., Sec. 2, Materials Science Center, National Tsing-Hua University, Hsin-Chu, Taiwan

300, ROC

J:din (_Inx.ntbu.edu.tw

Fax: +886-3-5716977

Tel: +886-3-5742574

NASA/CP--2001-210948 410 August 6, 2001
m,,,, ......... o,,.....

::::::::::::::::::::::::::::::::::::::

¢+:.X¢,_+:,:.:o?:+_X_.;;::_:::5-:,::_.:-:,:.:.;-:

Fig. 1 SEM micrographs of nitrogen-doped diamond films, nucleated with 0 sccm urea and grown with 0-8

sccm urea, which are designated as DN0-DNs, respectively.

4O

35

3O

E

o 25

2o

C21 15

-N lO

O

0,

0

• NoJ _ 0 Eo 104V/pm.¢_ 0052ev"

• N=J 15#A/cm2 Eo 402V/pm.¢. 0025ev"

• N4J 150#A/cm2 Eo_96V/pm.¢" 0021ev"

• NeJ _ 300#A/cm2 Eo_35V/pm.¢" 0028ev"

• NeJ _ 120_A/cm2 Eo_317V/_m.¢" 0(X34ev"

2

• ooO

e , _- °

_0_ _o oO

_ , _ , _ , _ , _ , _ _o

• I/E (_n_) V_'n. _000 O0

1 2 3 4 5 6 7

Electric Field (V&m)

Fig. 2 The electron field emission of nitrogen-doped diamond films, nucleated with 0 sccm urea and grown

with 0-8 sccm urea, which are designated as N0-Ns, respectively.

NASA/CP--2001-210948 411 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

SUB 10 NM TttICK A.MOIlPHO[;S CARBON" OVERCOXF I-OR "[Hf,_" FILI_,I

MAGNE'I'IC _.IEDIA

i,fl_ayuld YA,f_I_IOTO am! I_iitoyuk: blYOIlO

Fojilsu Lal;orateries L|d. Atsagi, 243-0197 Japan

Takashi TO "(OGL'CIII

Yamagata Fuj;,tsu Ltd. Higashine. 999-37,'),1 Japan

ABSTI_ACT

The area! recording density of h:_d disk drives has rapidly increased by over 100% p_r 3'ear in ,he t990s and is

expect ti_ :nc_eose _r,/_er in d_c 2ls_. cemtu_. Such a hlgb ra,.e has required i]I_pro,,'cmCP.l {_ device cantpomm su,;h

a_ a Gb,.nt Magnetoresisl:_e (GMI4.) bo.,,q :,Jtd a low ,IOIS¢ m':Oim_t, m¢cb:,Iror, Jcs god s_gr,:,l pcocessm2. ;]: :uldit:o:,.

tCncreduction of magnetic spacing has become an imperative for main,ainin£ such a _a,e...The magnebc sP.,-_.,o_°'_° is

defined l:,y the sam of flying be:g!tl and thickoess of lhi_ ovv.n:_;a_.:o die Jtead and medm:o, q'he/e_brc. I),i-re bas been

a big moli:.atiea to reduce _]te thlck,'.ess of the ov,etv,?at x_!61e t_::::Fhtg its :d;_: dulabf'.il,, ar.d sbcag c:_n:Jsion

resistance. An overcoat _s expected to be on :he order of nanomc:ers in an advance recording syst_:n.

Nil:ogenated or hyd:ogenated ao'orNtm_s :arbo_ Ii;m:. by DC l;tagqetr_;n sp:,lter;n S has boo:_ rudely ,cscd re:

,'ul overcoal ill 1_i_: [hill fi!aI lhcdhbll Vccadse riley llaJ_ p.ovid:.d su!,erm," p.:rfo_ mr, nee Jr( td'.,t_;'?jy dad ,o corrosb:n

res[stm_cc However. as :he overcoat [bJckness approsche_ to }es_ than l'd rap_ ,]_,:sputtered fihas may not mamta:n

tl_epo:tbro:aoce. 2b_mgb plasma chemi,;at vapor dcpoJ:tk ]: (;,C'.'D) o: ice beam de[,_siti,3o <;BD) appears ,._t,'_ :,:l

aiten:ative lc, lhc sput,.erhtg b_cause it provides films ,,,.ith high hurdItcss :-a:d _c,,?,I co',::lage ,t is still a coac_n:

whether or not Ibe deposition can continue to prey:de such good propcr,Jes when flirt, tb:ckness rcdu,;cs to a fe:_

Par_9)I_etel S.

FiRered CaIhodlc Arc tFCA) deposition can produce ven, hard amorphous carbon even at a fc_,, 5m_.ometer._.

compared wRh ocher __itoJuadss}.Icb a_ 1._Csp;_:ltCfil_g, [,C\,q_'_ :_d ;_B[). ';'he amorphous c,;rboa fi;m5 co:r,.;_;se,l of a

tet,agored sp_ bot,i;ng _tta; t"(,':o5 a d,..o.o,,J ..... s,_roc,o,c, , _ m:,] a i_:2_,xal , sFa b::_:d:og ,'b:,: [brms a gcai,h,_ ,%, ,....OC,O,c.

The film prepped by FCA deposition has an sp" content .Ua.lcl '_ -r i.,.mk",tho_e by the o,her methods. _,x,,¢can ebtain

snu;oth _,.'ndb:gbl_.. d¢l, se ft;lr,3.._S a re_S_iil, lhe FCA film _!t,:ws the so_;ng¢_t ,;,ear resis_,,n,:e wlu;,lg _;Jms prepared

by other m¢,.hods, as shew]t in Figure t. The F(.A film can be a slrong candidate for c,:t v.llra-thi]t o,.ercoa_ of

xb:ance modm

T.'ds _vt._& xv:,s sat:,D.'6_:! lr, p:,_t hy A_VI i:_ the Mi'FI'S R,_.D p:o;9am.

Figar,_ ! AFM in_.g_s 6fnanowear ca, lores tbr.,r_d a_ I0 _,?; aI_.d12 scans

Kcy_rords: overcoal, lhin fihr, magnetic media, amo_hous carbon. FCA. trlbelegy

Taka?_ki YAMAMOTO

f',_j;lsu 1.al:oratcr;:_s I.!:i.

I O-i ,Moriaosa'.o-_vakam,ya

A,.sugi 2S3-0197 Japan

Tel; -81 (46; 249-66°,5, Fax: +81 (46; 2504!2Sg

NASA/CP--2001-210948 412 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

IMPROVEMENTS OF TRIBOLOGICAL PROPERTIES OF METAL EVAPORATED

TAPES FOR DIGITAL RECORDING BY DLC PROTECTIVE COATING

Hiroyuki Osaki

Sony Corporation, Sendai Technology Center, 3-4-1 Sakuragi, Tagajo, Miyagi 985-0842, Japan

E-mail:ttir,_>yuki.Osaki @jp.sony.com, TEL:81-22-367-2452, FAX:81-22-367-2589

ABSTRACT

With the increase in the recording density of hard disk systems, the demand on the increase in the data recording

capacity of removal cassettes of tape drive systems as back-up systems for hard disk systems which could cause

"head crash" trouble is increasing. For that purpose, the volumetric recording density of tape drive systems, which is

one of the advantages of them, should be increased by increasing the areal recording density and decreasing the

thickness of the tape.

The mechanisms of tribological problems caused by the increased areal recording density and the methods to

improve them are investigated. It was found that the increase in the friction coefficient of the tape by rubbing causes

tribological problems, and that the combination of surface asperities, a wear resistant coating such as a DLC coating

and a lubricant is very effective to improve them. Even after a hundred million head traces, the wear of tape surface

could not be observed. By this combination, the metal evaporated tapes have obtained superior durabilities, as well

as superior recording characteristics, as the tapes for high capacity digital recording systems.

Keywords: tape drive systems, metal evaporated tape, surface asperities, wear resistant coating, DLC

INTRODUCTION

With the increase in the recording density of hard disk systems, the demand on the increase in the data recording

capacity of removal cassettes of tape drive systems as back-up systems for hard disk systems which could cause

"head crash" trouble is increasing. One of the advantages of the tape drive systems is high volumetric recording

density, which is obtained by high areal recording density and thin tapes. The trend of the volumetric recording

densities of helical scan recording systems, which are suitable for higher density recording, is shown in Figure 1. The

definition of the volumetric recording density is the product of the areal recording density (= the inverse of a half of

a minimum recording wave length times a track pitch) and the inverse of a tape thickness. The volumetric recording

density has been increased ten times in ten years. Each systems is going to next generation by increasing volumetric

recording density.

The areal recording density can be increased by introducing high performance tapes, such as metal evaporated

tapes with superior magnetic characteristics and smooth surface(ref. 1). Smooth surface reduces the spacing loss

between magnetic heads and a tape to increase the reproduced output signal, but produces higher friction coefficient,

which could cause tape damage by rotary heads and unstable tape runnability in the tape drives. The decreased tape

thickness will also make runnability unstable, especially when the friction coefficient is increased. The durability and

runnability of the metal evaporated tapes had been improved by forming surface asperities on it to reduce the friction

by decreasing real contact area, sacrificing a little of recording characteristics by spacing loss. However, the height

of surface asperities of metal evaporated tapes should be reduced to achieve higher areal recording density for

increasing demand. The decrease in the height of surface asperities results in the short life time until they are worn

out. Therefore, the wear resistant coating is required to keep the initial height of lowered surface asperities even

after many rubbing cycles. The thickness of wear resistant coating should be less than ten nanometers to reduce the

total spacing loss produced by the height of surface asperities and the thickness of the wear resistant coating. The

damage mechanisms and the improvement methods of the metal evaporated tapes are investigated and reported in

this paper.

NASA/CP--2001-210948 413 August 6, 2001
.+,

o

28

'.... +:+2--__-__--_-+;+-_--_ .... _......... /--]

74 i.... 1 "+ :,v++,,oo.P>_"

Recording

x _-5 i+,.,+_ Dv.._ty :

,__ :0' !......... _......... !....... : :.:.. ; , r

- -....... _ .'.,..i ......... _....

X I , '. ' ++" : : !

+ . : , +.., : +

V_s(+;+ .'., ! + +

' _C "° ' : ;

+ . ,'; :

'+Un,_t;++.'" + : vr_+e_P.+m_h

I ._ i : Te :T,"-aCkpib'Js

I °%

0_ ! i ... _.... :.......................... "+ •Tape _lock_'_

v_r

Figure 1.Trend of volumetric recording density of helical scan systems.

SPECIMENS AND APPARATUSES

STRUCTURE OF METAL EVAPORATED TAPES

The typical structure of metal evaporated tapes is shown in Figure 2. The magnetic layer is formed by vacuum

deposition of a Co0.8+l.0Ni0.2+0 alloy up to about 0.2 micrometers thickness on a polymer base film (4-10 microns

thickness). The materials for base film is polyethylene terephthalate (PET), polyethylene naphthalate (PEN), or

aramid. The back side of the base film is coated with particulate back coating layer. Some of the metal evaporated

tapes have DLC coating on the magnetic layer. The DLC coatings are formed by sputtering or CVD. The DLC

coating formed by sputtering is used in this paper. The top surface is covered with a lubricant. The lubricant used in

this paper is a modified perfluoropolyether (PFPE).

LUBRICANT

BACK COATING

( DLC COATING )

MAGNETIC LAYER

( SURFACE ASPERITIES )

Figure 2.Typical structure of metal evaporated tapes.

NASA/CP--2001-210948 414 August 6, 2001
STRUCTUREOFTAPEDRIVES

ThetypicalstructureoftapedrivesisshowninFigure3.Atapeisforwardedalongtapeguidesandastationary

drum.Rotaryheads, whicharemountedonandprotrudedfromarotarydrum,scanandrubthetapesurfacetorecord

(write)orplayback(read)asignal.In thecaseofvideosystems, themostseveremodefortapesisthestillmode

(pausemode,stand-bymode), inwhichmodetape-forwarding isstoppedandtherotaryheadsrubthesametrackon

thetapemanytimes.Forexample, theheadsrubthesametrack150timesasecondperheadforDV(digitalvideo)

systems. Thoughstillmodeisnotavailableinthecaseof datastoragesystems, it isoneofteststoevaluatehe

durabilitiesoftapes.

MAGNETIC TAPE ROTARY DRUM

_ I VIDEO HEAD

r ........... 1

, ', _O

STATIONARY DRUM ROLLER GUIDE

Figure 3.Typical structure of tape drives.

FRICTION TESTER BETWEEN TAPE AND HEAD

Establishing the friction coefficient between a tape and a rotary head is essential to analyze the wear mechanisms

of tapes (refs. 2 to 6). The friction tester is shown in Figure 4. It consists of a vertically fixed drum assembly, which

consists of two stationary drums and one rotary disk on which a rotary head is mounted. The head contacts a tape

surface from the inside of the drums.

Since only the rotary head mounted on the rotary disk rubs the tape surface through the slit between two

stationary drums, the friction force between the tape and the rotating head can be obtained by measuring the torque

by the torque meter (inverse magnetostrictive effect type) installed between the rotary disk and the motor.

A microscope with a stroboscope synchronized with the head rotation views the rubbing surface of the rotating

head. The rubbing surface of the head can be observed continuously through a TV monitor. The contact load

between the tape and the head can be controlled by the amount of head protrusion and the tape tension. The tension

is provided by the dead weight.

STATIONARY DRUMS

WEIGHT /

STROBOSCOPE

MAGNETIC TAPE

/

ROTARY HEAD

OTOR

QUE METER

(by INVERSE

MAGNETOSTRICTIVE EFFECT)

MICROSCOPE ,\

TV CAMERA

Figure 4.Schematic drawing of friction tester between tape and rotary head.

NASA/CP--2001-210948 415 August 6, 2001
RESULTS AND DISCUSSIONS

DAMAGE MECHANISM OF METAL EVAPORATED TAPES

The variations of the friction force between a video head and sample specimens of metal evaporated tapes with the

increasing number of head traces are shown in Figure5(ref. 4). The friction force of the sample as deposited, which is

indicated as "no asperity" sample in Figure5, increases quickly, and the magnetic layer is damaged at the critical

value of the friction force, which is indicated by a dotted line. The symbol x indicates the fracture of the magnetic

layer and a sudden decrease in reproduced output signal. The typical aspect of damaged tape surface is shown in

Figure 6(ref. 2). Many arc-shaped Helzern cracks and the debonding of the magnetic layer can be observed. It shows

the magnetic layer was peeled off by high friction force applied by the video head. This critical value of the friction

force is attributed to the bond strength between the magnetic layer and the base film(refs. 7 and 8).

e-

,4

LV

Ill

tO

nO

I.I.

Z

0

ItO

n"

I.I.

5

HEAD PROTRUSION : 30/1 m, 35/1 m I/

TAPE TENSION : 15 g

NO ASPERITY

35pm, 30pm

/P'/_- WITH ASPERITIES, 35/_m

I I I I I I r • • i ,

0 50 100

STILL TIME ( min )

I .... I .... I • • I I I I I l I i_ _ " i

0 5 10 15 20 40

NUMBER OF HEAD TRACES

i

I I I

25O

"'111

45

Figure 5.Variations of friction force between rotary head and metal evaporated tapes with and without

surface asperities.

DIREC_ON OF

HEAD TRACE

TAPE SURFACE

CRACKS

iiii',iii ,iiii',iiiiiiii!i!',!i!iiii:ii': i ;':iiiiii ii! !

!!i!i!iiii!iiii!!!!!i)i?22:;iii [_12i ;21 !iiii[!iiiiii!ii;i!

%*,%..%%*j.','. / ..... _ ._+ ,...,,...%,.**%.°%%,j,,.,,

DEBONDING

Figure 6.Micrograph of typical aspect of damaged tape surface.

NASA/CP--2001-210948 416 August 6, 2001
IMPROVEMENT BYSURFACEASPERITIES

In ordertoimprovethedurabilityinthestill modeof metalevaporated tapes,lowfrictionforceshouldbe

maintainedtoprevent themagneticlayerfrombeingpeeledoft'.Theincreaseinthefrictionforceseemstobecaused

bytheincreaseintherealcontact areawhichisattributedtosmoothingofthetapesurfacebywear(ref. 5),aswellas

bytheremoval ofalubricant andadsorbedcontamination. Thereductionoftherealcontactareawasattemptedby

introducingartificialsurfaceasperities(ref. 4),asshowninFigure7andFigure8(a).Theschematic drawingofthe

crossectionofthemetalevaporatedapewithsurfaceasperitiesisshowninFigure7.Themicrographofthetape

surfaceisshowniFigure8(a).Bytheseartificialsurfaceasperities, lowfrictionforcecanbemaintainedforalong

time,asshowniFigure5.Bythismethod, thefirstcommercial metalevaporated tapeforhigh-band8mmvideo

system(Hi8)obtainedenoughdurability.

Themetalevaporatedapesshouldbeimprovedintwopointsforthenewapplications ofhigherrecordingdensity

drives.Thefirstpointishigherreproducedoutputsignal.It shouldbeachievedbytheimprovement of magnetic

characteristics andthereductionof spacinglossbetweenatapeandahead.Therefore, theheightof surface

asperitiesshouldbedecreased.

Thesecondpointishigherdurability. Astherotatingspeedof therotaryheadsisincreasedtoincreasethe

transferrateofthedata,thedurabilityinthestillmodeshouldbeincreased. Iftheconventional metalevaporated

tapeisusedinsuchanadvancedsystem, thesurfaceasperities ofthemetalevaporatedapewillbewornoutbefore

practical lifetime.Thewornsurfaceoftheconventional metalevaporatedapewithsurfaceasperitiesisshowni

Figure8(b)(ref.4). Theartificialsurfaceasperitiesloosetheroletoreducethefrictionforcebywear(refs.4and9).

Therefore, theheightofsurfaceasperitiesshouldbeincreased. Thefirstpointandthesecondoneareinthecontrary

directions.

SURFACE ASPERITY

LUBRICANT

MAGNETC LAYER

BASE FILM

BACK COATING

Figure 7.Schematic drawing of cross-section of metal evaporated tape with surface asperities.

:::::::::::::::::::::::::::::::::::::::::::::::::::: '_.v,°,',',',',v,°,°,',',°,',', ,°,:,_,','°',',V°_,_,', °'_',', ,v,._,

,..,,°. °...°...°. _:.,°.,:(,.......,-..._.., -,,..,.1...,-_ .1.:, ;1,_ HEAD TRACE ::.::N:_,_::,_ °:_:_;,;:,,::_:N:,:_,X:::p.:_:_:::_

,:::::-_:_i;)i:,:.:::._:._::x_.:.x;:.:+:::,':<.:.::.'::-':;:.'..:-::;

"_.IK.'.'o5:_,.'o:._:_°" e*:_-;ol'°,','.°°.%"°°'°°°°_-,,:,_o',,*e.,.'.'o'_

/ :!" .._-:,._.- ....,,. ..,...: .-,..p; ..., •

!_i::.:::._."." '.".'. •._-":.:_-" _ '

ROLLER OF

WEAR DEBRIS ,lpm ,

(a) (b)

Figure 8.SEM micrographs of metal evaporated tapes with surface asperities

(a) before head traces

(b) after head traces.

NASA/CP--2001-210948 417 August 6, 2001
IMPROVEMENT BYDLCCOATING

Thesolutionforthecontrarydirectionsmentionedaboveiswearesistant coating,suchasaDiamondLike

Carbon(DLC)coating(refs. 10to 14),to keeptheinitialheightof loweredsurfaceasperities. Therequired

characteristic ofthecoatingistheresistanceagainst themildwearofsurfaceasperities, nothardness, becausethe

tapesurfaceisnotdamagedbyabrasivewear.Theheightof artificialsurfaceasperitieswasdesignedtomakethe

totalspacinglosscausedbythethicknessoftheDLCcoatingandtheheightofsurfaceasperitieslowerthanthetape

withouttheDLCcoating. Thevariationsofthefrictionforcebetweenthevideoheadandthemetalevaporatedapes

withandwithouttheDLCcoatinginthestillmodeisshowninFigure9(a).Bothtapeshavealubricant ontheirtop

surfaces. Thevariationsofstandardized rproducedoutputsignalofthesetwotapesinthestillmodeisshownin

Figure9(b).

WithouttheDLCcoating, thefrictionforceincreasesquickly,andtheoutputsignaldecreased quickly.Onethe

otherhand,withtheDLCcoating, lowfrictionforceismaintainedoverahundredmillionheadtraces. Andoutput

signal doesnotchange. Themicrographs ofthemetalevaporatedapewiththeDLCcoatingbeforeheadtracesand

afterahundredmillionheadtracesareshowninFigure10(a)andFigure10(b),respectively. Nowearofsurface

asperities canbeobserved.

¢-

e_

t0

v

UJ

O

nO

u_

Z

O

I--

O

tr

u_

10

0

i

0

Relative Speed : 10rn/sec

Tape Tension : 7.0g

Number of Heads : 2

Tapes : with lubricant

"'_without ;....... iiiiiiiiiiiii DLC Coating ............................

1 2 3 4 100

NUMBER OF HEAD TRACES ( × 106 )

t i i i _ t

1 2 3 4 93

STILL TIME (h)

(a)

4

I--

o. 0

I--

O

a -4

LM

(3

a -8

O

nQ.

UJ

nwith DLC Coating

_)iithout DLC Coating

Relative Speed : 10m/eec

_ Tape Tension : 7.0g

I Number of Heads : 2

LTapes_ - with lubricant

1 2 3 4 100

NUMBER OF HEAD TRACES ( × 106)

I I I I I _ I

0 1 2 3 4 93

STILL TIME (h)

(b)

Figure 9.Improvement of durability in still mode by DLC coating

(a) Variations of friction force for metal evaporated tapes with and without DLC coating

(b) variations of reproduced output signal for metal evaporated tapes with and without DLC coating.

NASA/CP--2001-210948 418 August 6, 2001
I 0.5/_ml

(a) before head traces

. ..

I 0.5 p m_

(b) after head traces

Figure IO.SEM micrographs of tape surface with DLC coating

(a) before head traces

(b) after head traces•

LUBRICATING PROPERTY AND DAMAGE OF DLC COATING

It was found that the durability of metal evaporated tapes in the still mode can be improved very much by the

adoption of a DLC coating. To evaluate the lubricating property of the DLC coating itself, the variations of the

friction force between a rotary head and the DLC-coated metal evaporated tapes with and without the lubricant are

compared, as shown in Figure 11. Without the lubricant, the friction force between the rotary head and the DLCcoated tape increases very quickly and has two parts of peaks of the friction force. It was found that the surface

asperities were not worn on the way to the first peak of the friction force. Therefore it seems that the quick increase

in the friction force of the DLC-coated tape without the lubricant up to the first peak is caused by the removal of

adsorbed contamination which has lubricating property, not by the increased real contact area caused by wear. At the

first peak of the friction force, the DLC layer is peeled by the increased friction force but cracks have not been

observed yet. By the rolling of the wear debris from the peeled DLC coating, as shown in Figure 12, the friction

force is decreased for a moment and increased again up to the second part of the peaks (friction force -- 10 in Figure

11) after the rollers are exhausted from the rubbing surface. These rollers are made of inorganic carbon, which

should be DLC. At the second part of the peaks, arc-shaped cracks through the magnetic layer are produced by the

increased friction force just behind the damaged area of the DLC coating, of which the friction force is increased by

the wear of surface asperities, as shown in Figure 13. After this process, the magnetic layer is peeled oft, as

mentioned before.

On the other hand, with the lubricant, low friction force is maintained. As the DLC coating at least used in this

paper has no lubricating property, DLC surface should be lubricated to improve the durability of the metal

evaporated tapes.

NASA/CP--2001-210948 419 August 6, 2001
A

= 15

,4

'_ 10

W

o 5

Z

O

I--

0

Relative speed :lOm/s

Tape tension :7.0g

ME tape with DLC coating

thout lubricant

with lubricant

i i i i I i i i i I i i i i I i i i i I i i i i

5 10 15 20 25

NUMBER OF HEAD TRACES ( x 10 4 passes)

Figure ll.Effect of lubricant on friction force between DLC coated tape and rotary head.

DIRECTION OF

HEAD TRACE

ROLLER FORMED BY ROLLING

OF WEAR DEBRIS OF DLC

i_ ._._.!:_:_._-...'_

! • ::

, 1/_m,

Figure 12.Roller formed by rolling of wear debris from damaged surface of DLC-coated metal evaporated

tape without lubricant.

NASA/CP--2001-210948 420 August 6, 2001
DIRECTION OF

HEAD TRACE

CRACK DAMAGED AREA

!!i_i_i_i__F!!!_

iii!iiiiiiiiiiiiiii!i!iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii!ii

g/e m

i r

Figure 13.Damaged surface of DLC-coated metal evaporated tape without lubricant.

CONCLUSIONS

The mechanisms of tribological problems caused by the increased areal recording density and the methods to

improve them are investigated. It was found that the increase in the friction of the tape by rubbing causes tribological

problems. Surface asperities decrease the friction force by decreasing real contact area, and the combination of wear

resistant coatings such as the DLC coating and a lubricant keep the initial height of surface asperities to keep the

friction force low. As the DLC coating dose not show lubricating property, the lubricant is required for the DLC

coating to show enough wear resistance. Therefore, the combination of surface asperities, wear resistant coating such

as the DLC coating and the lubricant is very effective to improve the durability of the metal evaporated tapes. Even

after a hundred million head traces, the wear of tape surface could not be observed. By this combination, the metal

evaporated tapes have obtained superior durabilities, as well as superior recording characteristics, as the tapes for

high capacity digital recording systems.

For more advanced systems with much higher recording density in future, the developments of new wear resistant

coatings, which is effective with the thickness below 5 nanometers is required to reduce the spacing loss.

REFERENCES

1. H. Naruse, K. Sato, H. Osaki, K. Chiba, T. Sasaki and H. Yoshimura, Advanced Metal Evaporated Tape

for Consumer Digital VCR's, IEEE Trans. Cons. Elec., vol.42, no.3, 1996, 851-859.

2. H.Osaki, K.Fukushi and K.Ozawa, Wear Mechanisms of Metal-Evaporated Magnetic Tapes in Helical

Scan Video Tape Recorders, IEEE Trans. Magn., vol.MAG-26, no.6, 1990, 3180-3185.

3. H.Osaki, E.Oyanagi, T.Kanou and J.Kurihara, Wear Mechanisms of Particulate Magnetic Tapes in

Helical Scan Video Tape Recorders, IEEE Trans. Magn.,vol. MAG-28, no. 1, 1992, 76-83.

4. H.Osaki, Role of Surface Asperities on Durability of Metal-Evaporated Magnetic Tapes, IEEE Trans.

Magn., vol.MAG-29, no.l, 1993, 11-20.

5. H.Osaki, H.Uchiyama and N.Honda, Wear Mechanism of Co-Cr Tape for Parpendicular Magnetic

Recording, IEEE Trans. Magn., vol.MAG-29, no. 1, 1993, 41-58.

6. H.Osaki, J.Kurihara and T.Kanou, Mechanisms of Head-Clogging by Particulate Magnetic Tapes in

Helical Scan Video Tape Recorders, IEEE Trnas. Magn., vol.MAG-30, no.4, 1994, 1491-1498.

7. D.R.Wheeler and H.Osaki, Metallization in Polymers, American Chemical Society Books No.440 Chapter

36,, American Chemical Society, 1990, 500-512.

NASA/CP--2001-210948 421 August 6, 2001
8. Q.Guo, H.Osaki, L.M.Keer andD.R.Wheeler, Measurement oftheintrinsicbondstrengthofbrittlethin

filmsonflexiblesubstrate, J.Appl.Phys. 68(4),15, 1990,1649-1654.

9. RBerar, K.Kato,H.Osaki, T.Terazono andRKapsa, WearMechanisms attheMetalEvaporated Magnetic

Tape-Head Interface, Trans.ASME,J.ofTribology, vol.121,1999,408-415.

10. H.Kurokawa, T.MitsuyandT.Yonezawa, Tech.Mtg.onMagn.Soc.ofJapan(inJapanese), vol.46,

1986,67.

11. K.Shinohara, H.Yoshida, T.KuniedaandM.Murai,J.Mag.Soc.ofJapan(inJapanese), vol.18,1994,

299.

12. H.Osaki,TribologyofVideoTapes, Journal ofWear,vol.2000, 1996,244-251.

13. H.Osaki,Tribologyofmetalevaporatedapes-forimprovement ofrecordingdensity-, Tribology

International, vol.31,no.9,1998,531-539.

14. H.Osaki,Flexiblemedia- recent developments fromthetribologypointofview,TribologyInternational,

vol.33,2000,373-382.

NASA/CP--2001-210948 422 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

PROPERTIES OF FLEXIBLE DLC FILM DEPOSITED BY AMPLITUDEMODULATED RF P-CVD.

Takahiro NAKAHIGASHI

Nippon ITF, Inc., 575 Kuze Tonoshiro-cyo, Minami-ku, Kyoto 601-8205 Japan

E-mail address:JDM04606@,nifty, ne.jp,

TEL 81-75-931-6040, FAX 81-75-931-6166

ABSTRACT

In recent years, DLC (Diamond-like-carbon) film is strongly expected to develop applications in

tribological elements. Contributing to this expectation is the properties of DLC films such as high hardness,

low friction coefficient, wear resistance, and chemical stability. DLC films have been mainly applied to

substrates of hard metals and ceramics because of their peculiar high internal stress.

Oil and fat have so far been added to improve the surface lubricity of the polymer materials such as

rubber and resin. When oil and fat run dry, however, drawbacks such as gradual increase in the friction

coefficient arise. Components and products made of rubber also tend to cling to the other materials used

together therein. We wondered if we could eliminate such oil and fat additives that cause:such bad effects.

We have proceeded to coating to the polymer materials such as rubber and resin while making the best use

of the feature of DLC.

We thought that three issues have to be resolved. (1) Low heat resistance of the polymer materials such

as rubber and resin. (2) Pollution of the polymer material surface by oil, fat, resin, and oxidation prevention

agents, etc., and (3) Transformation of the polymer materials. To resolve these possible problems: (1) We

have developed a processing method by using the Amplitude-Modulated RF Plasma Chemical Vapor

Deposition method which enables coating at lower temperature (below 80 degree C) and does not allow the

processing temperature to rise any higher. (2) To prevent the pollution, we decided to clean the polymer

surface by plasma. (3) To prevent the transformation, the film should be flexible enough to absorb the

polymer material transformation. We have modified the DLC film structure to permit expansion and

contraction.

The flexible DLC film is one of the novel techniques to deposit the DLC film on polymer materials

such as rubber and resin, which does not cause any peeling off of the film even with the deformation of the

substrates. The flexible DLC technique is expected to expand itself and to create wider applications of DLC

films.

Keywords: DLC, Plasma, CVD, low friction, polymer material

End of paper.

NASA/CP--2001-210948 423 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

HIGH PERFORMANCE DLC COATINGS FOR SEVERE WEAR APPLICATIONS

Suri A. Sastri, Nilesh Gunda and Richard Cooke

Surmet Corporation

33 B Street, Burlington, MA. 01803

Santosh Jha and Vimal Desai

AMPAC, University of Central Florida

12424 Research Parkway, Orlando, Florida 32826

ABSTRACT

This paper deals with diamond like carbon coatings that show very promising results in terms of their

tribological properties and better adherence to and compatibility with a wide range of substrates. These films were

deposited using a plasma assisted vacuum vapor deposition technique. Advantages of this technique include, the

ability to precisely control the coating thickness, and the achievement of a micro conformable (to substrate surface),

uniform and smooth films on a variety of substrates. Coatings were deposited at relatively low temperatures

(<100°C) and thicknesses exceeding 25 microns were achieved.

Results of the various advanced materials characterization techniques such as Raman Spectroscopy, Pin on Disc

wear testing and dust erosion testing of the coating on a variety of substrates will be discussed. A novel engineered

multi-layered nanocomposite coating based on DLC which exhibited excellent resistance to impact erosion by highenergy abrasive particles will also be discussed.

The authors will also provide examples of a series of applications in the industry that are taking advantage of

the unusual properties inherent in these DLC coatings. Large parts up to 8' in length have been successfully coated.

Examples of critical commercial components in aerospace, biomedical and semiconductor wafer handling industries,

among others, will be enumerated.

Keywords: DLC, coating, characterization, wear, hard

INTRODUCTION

Diamond and diamond-like coatings have attracted rapidly increasing interest since 1980's and these coatings

are today probably the most intensively researched surface coatings. The commercial interest in these coatings is

considerable, not only because of their great potential in tribological applications but also because of their potential

as new and more efficient semiconductor and optical thin film materials. Diamond-like carbon (DLC) is the name

commonly accepted for hard carbon coatings which have similar mechanical, optical, electrical and chemical

properties to natural diamond, but which do not have a dominant crystalline lattice structure. This paper deals with

diamond like carbon coatings that show very promising results in terms of their tribological properties and better

adherence to and compatibility with a wide range of substrates.

Low STRESS ULTRA-C TM AND ULTRAC-HT TM DLC COATINGS

In the present study diamond hard carbon coatings, namely, UltraC Diamond TM and UltraC-HT TM deposited

using low-temperature plasma assisted vacuum vapor technique, were investigated. UltraC Diamond TM coating is an

amorphous diamond hard carbon coating while UltraC-HT TM (High toughness UltraC coating) is an ultrafine

lamellar composite coating based on UltraC Diamond TM and SixCy with superior toughness compared to the parent

coating. Various substrates were coated for characterization of the coating. Low stress coatings of thicknesses more

than 15gm for UltraC Diamond TM and 30gm for UltraC-HT TM were deposited in this study. This is a significant

improvement over conventional DLC technology in which the thickness is limited to a maximum of 5gm due to

high internal stress in the coating.

NASA/CP--2001-210948 424 August 6, 2001
EXTREMELY HIGHWEARESISTANCE

Pin-on-Diskweartestingcharacterization ofourcoatingsonTi6A14ValloydiskswascarriedoutatFalex

CorporationofSugarGrove,IL.(ref.1)2.5inchdiameter Ti6A14Valloydisksweremachinedandpolishedby

metallographic techniques toasubmicronfinish.Thesediskswerecoatedwithdifferent combinations ofcoating.

Weartestingwascarriedoutwiththetestingparameters showninTable1.FollowingisalistofTi6A14Valloydiscs

withdifferent combinationofthecoatingsthatweredepositedandsenttoFalexforweartesting:

1)2_mthickUltraCDiamondCoating

2)UltraC-HT(Laminatedcompositestructure)

3)BareTi6A14Valloy

Table1:Wear Test Parameters for Pin on Disc Testing

Ball Material 0.5"A1203 ball at lON load

0.5" Silicon Nitride ball for 15.68N

load

Test Speed (rpm) 71

Test Load 10N and 15.68N

Test Duration 1000m

Temperature Ambient

Test Radius 13.5mm

Figure 1 shows the wear profiles obtained on the samples. Table 2 shows the wear volume obtained on the

samples after the tests.

_lm

-8

-32

Bare

Ti6A14V alloy

10 N Load

A1203 ball

_lm

_lm

0.0

-1

UltraC TM

2 gm thick

15.68 N Load

SiBN4 ball

UltraC-HT TM

6 gm thick

15.68 N Load

SiBN4 ball

Fig. 1. Wear profiles obtained on samples after wear testing. The red lines indicate the wear track area.

NASA/CP--2001-210948 425 August 6, 2001
Table2:Wearvolumefordifferent samples

Sample Wear volume Wear volume

(mm a) at (mm a) at

10 N 15.68N

UltraC NMW 0.02058

Laminated NMW 0.01355

Bare alloy 6.8606

NMW = No measurable wear

Wear testing was also carried out on a SAE52100 bearing steel coupon polished to a mirror finish and coated

with a 10gm thick coating of UltraC Diamond TM coating. The test was carried out using a Si3N4 ball with a 10N

load at the same conditions as the previous wear tests. The test did not produce any measurable wear. Wear profile

did not produce any measurable wear. Fig. 2 shows wear profile obtained on the sample after the test.

iw • ,. -.i...., .

4 ...... + -- -- %-- -- --

• -" ',' .'.

00'"' ' "" '

(." • _ .%.',

, ::.i:.;.:i::.i-b:::

d_-_ _._ _ d_i_ _.z-- -

. ,. • t. ] ' t " " p.....

I(-"_??T¢?,-_.+-'_ -_ . • ..

• . .' . --._---__._._

UltraC

10gm thick

10N Load

Si3N4 ball

Fig. 2. Wear profile obtained on samples after wear tests showing no measurable wear produced on the

sample

Wear data shows that UltraC Diamond TM based coatings even with thicknesses exceeding 10gm are extremely

wear resistant.

EXTREMELY EROSION RESISTANT ULTRAC-HT TM COATING

Erosion testing was carried out on 2gm thick UltraC diamond TM coating and 25gm thick UltraC-HT TM coating

with twenty-five ultrafine lamellar layers of SixCy (2500 A) and UltraC (7500 A). Erosion tests were done in a wind

tunnel facility at the University of Cincinnati (UC) using Arizona dust with silica particle sizes between 10-100 gm

and 9.5gm alumina at particle velocities of 600'/s. The samples with just the UltraC diamond TM coating did not

survive the tests but the multilayered (nanolaminated) coating of UltraC-HT TM displayed very good erosion

resistance. Results of the testing done on UltraC-HT TM (nano-laminated coating structure) are shown in Table 3 and

Table 4.

Table 3. Erosion test results on nano-laminated coating structure using alumina particles (9.51xm) at 600ft./s

Angle of Impact (o)

90

Mass Loading (gms)

5

Erosion Rate (mg/g)

0.092

90 5 0.074

90 20 0.05

90 30 0.03

90 100 0.03

160 (TOTAL) .0552

30 10 0.270

30 50 0.60

30 20 1.11

30 20 1.4

I lOO(TOTAL) I 0.845

NASAJCP--2001-210948 426 August6,2001
Table4.Erosiontestresultsonnano-laminated coatingstructureusingsilicaparticles(100-200_tm) at600ft./s

AngleofImpact(o)

9O

MassLoading(g)

100

ErosionRate(mg/g)

1.76

90 100 1.10

200(TOTAL) 1.43

ResultsofthedusterosiontestingshowedthatheUltraC-HTTM coating displayed extremely high erosion

resistance.

RAMAN SPECTROSCOPIC ANALYSIS

The sample of UltraC Diamond TM was analyzed using Raman spectroscopy (Model: LabRam, JY/Horiba). The

632.8 nm radiation of an air-cooled 6 mW He-Ne laser was focused to a round spot of about 1 mm. The investigated

sample spots were optically checked with a microscope using a magnification of 1000X. There was no degradation

of the coated specimens visible under these conditions. The Raman spectra were excited at a temperature of 22+2°C

under laboratory conditions. The different locations were checked on each specimen to have consistency of data.

The Raman spectrum is shown in Figure 3.

30000 "

!!i ii

I Intensity 25o]o:: :: :: , : .i :.

15000 "",'i:

c<,.,,

10000

doo iAo 14'oo doo 1_oo

Raman shift (cm -1)

Figure 3: Raman spectra of UltraC Diamond TM coating on Si-wafer

As can be seen from the above figure that the most significant peak appears at about 1514.45 cm-1. This peak

represents the amorphous carbon phase (clusters of sp2-bonded carbon embedded in a sp3-bonded carbon matrix).

This peak is generally assigned to the distorted graphitic structure due to C-C stretching vibration mode. Another

peak at 1319.3 cm-1, though less significant, is attributed to the diamond-like sp3 bonded carbon structure in the

film.

The above figure represents the nature of a typical Raman spectrum of a DLC film with characteristic D- and Gpeak at 1319.3 and 1514.4cm-1, respectively. The result is in good agreement with those reported in various

literatures (refs. 2 to 3). The important point that should be mentioned here is as the film composition changes from

a predominantly diamond phase to graphitic carbon rich phase, the intensity of D-band peak decreases while that of

G-band peak increases continuously.

NASA/CP--2001-210948 427 August 6, 2001
APPLICATIONS

UltraC DiamondTM based coatings are being evaluated for many applications because of its ability to virtually

TM

stop the formation of macro and submicroscopic wear debris, UltraC Diamond is increasingly being specified for a

wide range of precision applications in environments encompassing semiconductor wafer processing to biomedical

engineering. Some examples are listed below:

Electrostatic Clamp Surface

Electrostatic clamp or E-clamp is used to hold the wafer during Ion Implantation and other IC fabrication steps.

UltraC Diamond TM is being used as the outer surface of the E-clamp to reduce backside particle generation. The

coating dramatically reduces the particle contamination by at least three orders of magnitude when compared to a

polished alumina surface. A photograph of the e-clamp surface with the UltraC Diamond TM coating is shown in

Figure 4.

__-_-_-°-___._-_____._._.___._._______._.___._.___.___-___-_.___-_._._.___.___________-_-°-_-_-_- ,°° ..............................................

iiiiiiiiiiiiiiii ........... ijiiiiii i

::::::::::::::::::::::::iiiiiiiiiiii i ",,':::::::::::::iiii i

:_:_:_:_:_:?_:_:?_:?1:1:??_:_:_:_:_:;_:_:_:_:_:_:?_:1:_:_...;..._:_:1:_:1:_:_:_:_:?_:_:_:_:!

Fig. 4. Photograph of an electrostatic clamp with UltraC Diamond TM hard carbon coating on the top

surface to prevent particle formation

Air Bearings

UltraC Diamond TM coating is applied on air bearings used in semiconductor positioning equipment. These

bearings are rotated at speeds of 5000 to 10000 rpm. UltraC Diamond TM applied on a hard anodized surface prevents

wear of the bearing surfaces when spinning with and without the presence of air. The near frictionless and wear

resistant property of the coating allows the bearings to be run without particle contamination even when the flow of

air is interrupted.

Fig 5. Photograph of air bearings coated with UltraC Diamond coating

NASA/CP--2001-210948 428 August 6, 2001
Cryocooler partsonHubbletelescope

Surmet CorporationandGoddard's HubbleProject, successfully putthediamond-hard coating, UltraC

DiamondTM, through a rigorous test program to verify it could operate in the extremes of space. The coating covers

tiny parts of an experimental cryogenic cooler, or "cryocooler," that will be attached onto Hubble's Near Infrared

Camera and Multi-Object Spectrometer (NICMOS) during a servicing mission in 2001. The cryocooler cools the

detectors in the instrument and maintains them at 70 degrees Kelvin.

During the cryocooler's normal operation, a tiny circulator shaft spins at a speed of 6000 revolutions per second.

To minimize friction and wear, the circulator shaft and bearings required a very thin, hard, slippery coating. UltraC

Diamond TM coating virtually eliminated wear and tear on moving parts fabricated from a Ti alloy.

Fig. 6. Tiny bearings and turbine shafts like these are part of a new cooling system for Hubble's

infrared NICMOS instrument. To minimize friction and wear, shafts and bearings are covered with

a slippery, diamond-hard coating approximately 1/100th the thickness of a human hair. These tiny

parts are pictured with a dime for size comparison.

Left Ventricular Assist Pumps

UltraC Diamond TM coating is applied on parts of a left ventricular assist heart pump made of Ti6A14V alloy to

prevent wear on the moving parts. The coating is bio-compatible and a thin fluoropolymer overcoat is applied on

top of the UltraC Diamond TM coating to improve the thrombogenity and hydrophobicity of the surface.

Fig. 7. Photograph of parts of a left ventricular assist heart pump coated with UltraC Diamond TM coating

CONCLUSIONS

The low stress DLC technology described in this paper provides several competitive advantages:

• Ability to deposit on many types of substrates; metals, ceramics and polymers

• Coatings with a thickness of up to 15gm for UltraC Diamond TM and 30gm for UltraC-IIT TM can be

deposited with extremely low residual stress

• UltraC Diamond TM provides extremely high wear resistance with a very low coefficient of friction

• Multilayered coating of UltraC-HT TM provides high fracture toughness, low internal stresses and the ability

to control the total stress in the coating

• The process is highly scaleable to larger substrates with varying geometries

NASA/CP--2001-210948 429 August 6, 2001
REFERENCES

1. Suri Sastri, N. Gunda, Rick Cooke: Compressor Impeller Erosion Resistant Surface Treatment, Final

Report, SBIR Phase I Contract No. N68335-00-C-0335

2. A.Dorner, B. Wielage, C. Schurer: Improvement of the corrosion resistance of C/Al-composites by

diamond-like carbon coatings, Thin Solid Films 355-356 (1999) 214-218.

3. J.B. Cui, M. Stammler, J. Reistein and L. Ley: Role of hydrogen on field emission from chemical vapor

deposited diamond and nanocrystalline diamond powder,"J. App. Physics, 88 (2000) 3667-3673.

NASA/CP--2001-210948 430 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

LOW-TEMPERATURE TECHNOLOGY OF OBTAINING THE

PROTECTING LAYER FROM DIAMOND-LIKE FILM

Kassabian K., Panosyan Zh., Stepanyan A., Torosyan G., Yengibaryan Y.

State Engineering University of Armenia

105 Teryan St., Yerevan 375009 Armenia; E-mail: joseph@seua.am

Fax: (3741) 151-068; phone: (3741) 272202

ABSTRACT

The technological regimes of obtaining the diamond-like C (DLC) films with all the necessary characteristics

have been elaborated. We have studied the spectrum dependence of the reflection, transparence and absorbance

indices on the basis of which there were determined the values of the reflective index and the width of the banned

zone, where n=2 and Eg=2.5 eV. The new photovoltaic cell (PV cell) with the enlightenment layer from DLC films

was obtained. At thickness of the film equal to 80 nm, the silicon PV cell has received the most enlightenment. The

reflective index of the enlightenment layer at wavelength of 650 nm is 0.02. The obtained DLC films displayed

firmness to the concentrated acids, and the films kept their optical properties after the long affection of the

ultraviolet radiation.

Keywords: Diamond-like, photovoltaic, protecting, low-temperature, technology.

INTRODUCTION

The possibility of films growth obtained from diamond on Si, GaAs substrates etc., allows to develop new

optoelectrical and microelectrical devices, which have such unique properties like stability to mechanical, thermal,

chemical and radioactive influences, high transparence in large range of spectrum. These properties make them

perspective material for protective, heat-eliminated, resistive, enlightenment (antireflective) surfaces (ref. 1).

Anyway, the growth of diamonds and SiC on GaAs or Si substrate is carried out by using comparatively high

temperatures, expensive technologies and equipment (ref. 1). In connection with it the problem of development

simple but inexpensive technology of diamond-like C film growth by ionized plasma assisted deposition from vapor

phase at comparatively low pressure and temperatures, near to room temperature, causes a big interest.

The aim of the given investigation is the development of low-temperature obtaining the polycrystalline

diamond-like C film on Si substrate.

THE TECHNOLOGY OF OBTAINING THE FILMS ON Si PV CELL

The diamond or diamond-like C films can be obtained in many ways, but none of the existing methods can be

applied for using them as a protective layer in production, as the temperature of the used substrates exceeds 1000 °C.

In this case the PV cell from crystal Si is used as a substrate, which cannot keep the properties of p-n junction, lying

near to the deposited surface.

Usually the protective and enlightenment layers are deposited on such a surface, where the distance between the

surface and the p-n junction is 0.5+1 mkm.

In this work, in order to get protective and enlightenment layers, DLC films were obtained on the crystal Si

substrate with the help of plasma assisted deposition of ionized streams of C + and H +, which were obtained by

dissolution of hydrocarbon compounds in the electric and magnetic fields. The advantage of this method is its

simplicity and possibility of using uncomplicated technological equipment, applying relatively low pressure and the

temperature of the substrate near to room temperature.

For this purpose there were worked out technological regimes for obtaining DLC films at temperature below

150 °C. The technology of growing is based on the method of DC plasma enhanced chemical vapor deposition

(CVD) with simultaneous affect of the magnetic field. The technological plant is made on the basis of a standard

5

vacuum device and specially designed source of ions. After achieving the preliminary vacuum (-10 mm Hg), there

starts supply of cyclohexane (C6H12) vapor into the region of crossed electrical and magnetic fields. Plasma process

becomes stable with the help of selecting the supply regimes of C6H12 vapor and magnitude of magnetic and

electrical fields. The quality of films is greatly influenced by the average kinetic energy of positive ions, which is

NASA/CP--2001-210948 431 August6,2001
correctedwiththehelpof theelectricaldeflectionbetweenthesubstrateandcathodeof ion supply.The

investigations haveshownthatCfilmswithhighdensityanddiamond-like properties, grownwiththehelpofthe

giventechnology, aregeneratedat30+50eVof averagekineticenergyof ions(ref.2).In orderto choosethe

technological regimesof DLCfilmsobtainingtherewerecarriedoutinvestigations of thedensity,growthrate,

thicknessandstructureoftheobtainedfilms.

In thefilms,whichweredepositedinthisway,differentcarbonstages(diamond, diamond-like, graphiteor

otherCHbonds)canbeformed.Theirelativeformation, whichleadstoobtainingdifferentpropertiesoffilms,

dependsontheconditionsofdeposition, i.e.onphysico-chemical processes, takingplaceintheplasma.

In theprocessof cyclohexane andtoluoldissolution, variousradicalsaregenerated. Theirtypesandrelative

quantitystrictlydependontheperformance parameters ofthesourceofions(suchascurrent, voltage, magnitudeof

magneticfield,etc)andontheconfigurationof electricandmagneticfields.Formationof aradicalalsocauses

changesinthestructureandinthepropertiesofthegrowingfilms.

Fortheseinvestigations therewereworkedoutchangesofvoltagemagnitudebtweenanodeandcathodeandof

plasmacurrent(accordingly V=1500+5000 VandIp=50+100 mA).

Infig.1thereisshowndependence ofthedensityandgrowthrateofDLCfilmsontheanodevoltageandthe

substratepotential. Thus,wecanconcludethatwhilethevoltageisincreasing, thegrowingdensityprovesthe

changeofthefilmsformationmechanismofthefilms.It ispossibletochangethedissolutionprocesses whilethe

energyofionsisincreasingintheionsystem, andalsotochangetheprocesses ofadsorptionaddesorptionfthe

articlesonthesubstrate.

It isknownthatheincreaseofthefilmdensityisconditionedbytheamountof thehydrogenatoms.The

investigations resultsshowedthathebiggest magnitudeofgrowthrateanddensityofDLCfilmscorrespond tothe

voltageof3.5eVbetweenanodeandcathodeandtothezeropotential ofthesubstrate(heearthingcase).

TheX-Rayphaseanalysesof thefilmsgrownbymeansof thislow-temperature technology, havingthe

thicknessof0.3mkm,showedthathestagewithsp3covalent connections isprevailing. Thefurtherincreaseofthe

thicknesscausesdecreaseof thediamondportion.Thus,theamorphous layerwithsp2chemical connections,

conditionedbycarbon, isformed.

Inthiswork,DLCfilmsweremainlygrownonthemonocrystalline siliconfilm.ThegrowthofDLCfilmson

thesurfaceofmonocrystalline siliconcanbedescribedasthefollowing,resultingfromtheX-Rayphaseanalyses.

OntheinitialstageSiCpolycrystalline filmwiththethicknessof-102A°isobtained(ref.2).Thediamond-like

stageisgrownonthisfilm.WhenSisubstrateisbeingbombardedbyC+ions(withhighenergy), theyenterSi

ratherdeeplyandcreatethepolycrystalline layerofSiC.Thefurthergrowingof CintheareaofSi-Cjunction

decreases thesignificanceofbothSiandSiO2film.

At first,thediamondhexagonal crystallitesareformed;theyhavethethicknessof-2300A°andcreatethe

nucleusonthehexagonal sidesofthefilmsdensity(fig.1).

v, nm/c P' g/°na

0,25 ............................ 2,5 ................................

1 // 0,2

0,15

0,1

0,05

2/

0 • • , _ • • , L • , • J 2,1

0 2 4 6

U, kV

a

2,4

2,3

2,2

,/

,// 2

illlllllll

1 2 3 4 5

_kV

Fig.1. Dependence of the rate (a) and density (b) of DLC films, grown on the accelerating

voltage U at dissolution of C6H12. 1-earthed substrate, 2-substrate has varying potential

NASA/CP--2001-210948 432 August 6, 2001
OntheX-Rayphasecharacteristics, whichcorrespond toSiCanddiamond-likeC(besidesX-Raydiffraction

maximums, formedfromtheinteratomic distance) therewereobtainedpeaks, corresponding to(111)directionofSi

substrateandsometimes toSiO2thinfilm,existingonSisurface.

THE OPTICAL PROPERTIES OF DLC FILMS

It is known that besides diamond films there are also some other carbon films, such as amorphous carbon (c_-C),

tetragonal amorphous carbon (to_-C), diamond-like C (DLC), hydrogen amorphous carbon (o_-C:H, which contains

approximately 40% of hydrogen), C60 and others, which have different properties (refs. 3 and 4) . For most applied

tasks each of the afore-mentioned films is of particular interest.

DLC films, grown on the substrate of Si and A1203 have homogeneous surfaces with the mirror (reflection)

property. For the purpose of application of protection and enlightenment on the surfaces of PV cells, it is very

important to study the optical characteristics of the films, obtained through the above mentioned technological

regimes.

The optical characteristics of diamond films, obtained on Si and A1203 substrate, were received on the basis of

measuring reflection specters and absorbtion.

The curves in fig.2 are the interference images, which were obtained from the reflection of DLC films, grown

on Si(111) surface. They are spectral dependences of the coefficient of the diamond films reflection, having two

different thickness. The curves show that the interference images of the films with different thickness, are different.

R 0,35

0,3

0,25

0,2

0,15

0,1

0,05

0

1

i i i i

200 300 400 500 600 700 800 900 1000

_,, nm

Fig.2. Spectral dependence of the reflection coefficient of DLC films of different

thickness, grown on the monocrystalline Si, 1- d--1120 nm, 2- d--150 nm:

Using the spectral positions of the extremums, there was found the thickness of the given film by means of the

following correlation between the positions of extremums (in this spectral range) and the film thickness.

4n(£).d = m_,m m =1,2,3 ..... (1)

where n(X) is the refractive index, d is the thickness of the film, m is the order of the interference extremum, Xm is

the wavelength, corresponding to the m-order of the extremum.

The thickness of the film is determined by 2 interferential maximums and minimums in the following formula:

d= &&-:

2n('gm-2-'gin) (2)

where n is the refractive index, _n and _n-2 are the wavelength of the extremums of m-order and m-2-order

accordingly. Thus, the thickness of the films, grown on Si substrates, is 1250+40 nm for one film and 150 nm for the

other one.

NASA/CP--2001-210948 433 August 6, 2001
Therewasmeasuredc_()_) spectral dependence of DLCfilms,grownonA1203substrates (fig.3).Fromthese

curvesit isseenthatheabsorbtionedgeofthefilmsisneartothewavelength of500nm,whichcorresponds tothe

lengthofthebannedzone,whichhasthevalueof2.5eV.

50000

45000

40000

35000

30000

25000

20000

15000

10000

5000

0 ''

25O

@@

@

*@

%

@

%.

......."-.:.,,,.

300 350 400 450 500 550 600

_,, ran

Fig.3. Spectral dependence of absorbtion coefficient of DLC film, grown on A1203

TESTING AND MEASURING OF THE PARAMETERS OF PV CELLS WITH DLC FILM

Dependence of the reflection coefficient on the light wavelength of the films, obtained by the above mentioned

technology (as the enlightenment layer) can be seen in fig.4. The thickness of the enlightenment layer of d--80 nm

was selected so that the least value of the reflection coefficient R for the first interference minimum corresponded to

the falling light wavelength of )_m=650 nm.

R 0,2

0,15

0,1

0,05

0

200 300 400 500 600 700 800 900 1000

_,_ nm

Fig.4. Spectral dependence of reflection coefficient of DLC enlightenment film, grown on Si.

This value of )_n corresponds to the maximum photosensitivity of the Si convertor. As it is seen from fig.4, the

reflection coefficient in the range )_n=650 nm is 2%, and in the ranges )_--450 nm and )_--1000 nm it does not

exceed 10%.

These low values R of silicon convertor in the whole spectral range show that the enlightenment layer has been

obtained. Comparing the spectral dependence of the received R for the obtained DLC film with the R spectral

dependences of the one-layer enlightenment layer (which corresponds to the above mentioned one-layer case), for

example from ZnS and others, grown on the surfaces of the silicon convertors, it is seen that DLC film has

advantage to the other enlightenment layers.

Investigations of the mechanical hardness of DLC films, grown on the surface of the monocrystalline Si(111),

showed that they had coefficient 7.5-8.0 by the scale of Moss, which proves their high mechanical hardness.

After the investigations, these films displayed chemical firmness. It was clear that they have high firmness to

the concentrated acids, such as: H2SO4, HC1, HNO 3 and HF.

PV cells of the new type were tested by ultraviolet radiation in order to find out the changes of the performance

parameters under these conditions.

Alongside with the new PV cell with the enlightenment layer from the DLC film, the films of the same type,

grown on the transparent substrates (for example, on A1203) were tested by ultraviolet radiation. PV cell and films,

NASA/CP--2001-210948 434 August 6, 2001
forthepurposeofcomparison, weretestedbymeansofradiationwithintensityof5000hours.Theresultof the

investigations showedthatransparence oftheenlightenment layerdidnotundergoanychanges, neitherdidthePV

cellcharacteristics (efficiency, open-circuit voltage, short-circuit current)(refs.5to7).

Thegivenspectral dependence ofabsorbtioncoefficient (fig.3)showedthathefilmsweretransparent inthe

visibleandnearinfraredranges. Allthecurrent-voltage characteristics ofthePVcellsweremeasured, whichhelped

todeterminetheefficiencyofthePVcell.

Thus,makingof thenewDLCoutlayerallowstoincreasegreatlymechanical, chemical andradioactive

firmness, performance reliabilityandenduranceoftheenlightenment layerand,consequently, ofthePVcell,and

promotesincreaseof thegeneratedlectrical power,whichleadsto decreaseof thecostof theelectricenergy,

generated bytheconvector.

Besidesalltheafore-mentioned, applicationof theprotectiveandenlighteningdiamond-likelayersallowsto

decreaseinmanytimestheweightofthePVcellduetonotusingthetraditional protectivelayers.Theadditional

thickprotectivelayersfromglassandquartzleadtothesignificant icreaseoftheweightofcells.Thiscircumstance

isofgreatsignificance, especiallywhenusingtheminspace.

REFERENCES

1. Editors: R.J. Shul, F. Ren, M. Murakami, W. Pletschen. Wide-Bandgap Electronic Devices,- MRS

Symposium Proceedings 2000, v. 622.

2. Panosyan Zh., Arakelova E.R., Voskanyan S.S., Voskanyan A.S., Yengibaryan Y.V., Stepanyan A.V,

Torosyan G.G., Shaboyan A.S. The obtaining and investigation of silicone carbide films and diamond-like

carbon on silicone surfaces// Proceeding of ISTC seminar "Conversion Potential of Armenia and ISTC

Programs", Yerevan 2000, part 2, p.207-210.

3. Clough F.J., Lleinsorge B., Milne W.I, Robertson J. Tetrahedral amorphous carbon thin film transistors//

Materials Research Society, Symposium Proceedings, v.423, p. 39-43, 1996.

4. Keblinski P., Wolf D., Cleri F., Phillpot S.R., Gleiter H. On the nature of grain boundaries in

nanocrystalline diamond//MRS Bulletin/September, p. 36-41, 1998.

5. Voskanyan S., Yengibaryan Y., Karapetyan K., Panossian Zh. Technology of manufacturing and properties

of diamond-like carbon films. In: Proceeding of the 2nd national conf. "Semiconductor Microelectronic",

Dilijan, 1999, p.194-197.

6. Voscanyan S., Yengibaryan Y., Karapetyan K., Panosyan Zh. Optical properties of diamond-like carbon

films.- Proceedings of National Academy of Sciences of Armenia, Physics, v.34, N6, 1999, p 358-361.

7. AM Patent N 851, H01 L31/02, Semiconductor PV cell / Panosyan Zh., et. all,/20.08.99. Officical Bull.

N 2, 2000.

NASA/CP--2001-210948 435 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

The Applications of Amorphous Diamond deposited by Cathodic Arc.

James C. Sung

KINIK Company, Taipei, Taiwan, & Taipei University of Technology, Taipei, Taiwan, R.O.C.

Abstract

Amorphous diamond combines the advantages of supreme properties of diamond and coating

capability of diamond-like carbon (DLC). It has been applied to a variety of industrial products with

superb performance in wear resistance, dry lubrication, and acid protection. Amorphous diamond earl

be deposited rapidly by a cathodic arc system in large dimensions. The wide spread applications of

amorphous diamond coatings can drastically improve the productivity of many industrial products.

Key Words: Thin Films, Coatings, Hardness, Abrasion, Corrosion.

NASA/CP--2001-210948 436 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

SUPERLOW FRICTION CARBON FILMS FROM HYDROGEN RICH

HYDROCARBON PLASMAS*

Ali Erdemir

Energy Technology Division

Argonne National Laboratory

Argonne, IL 60439

USA

ABSTRACT

Carbon-based coatings combine many attractive properties that make them good candidates for a wide range of

critical engineering applications. Systematic studies in our laboratory on carbon films over the last decade have

resulted in the development of a new class of amorphous diamondlike carbon films providing friction and wear

coefficients of 0.001 and 10 -10 mm3/N.m, respectively (when tested in dry nitrogen or inert-gas environments).

These films were synthesized in a plasma-enhanced chemical vapor deposition system that uses high proportions of

gaseous hydrogen in addition to hydrocarbon source gases, such as methane and acetylene. Comprehensive

tribological studies of the films revealed the existence of a close correlation between the chemistry of the

hydrocarbon source gases and the friction and wear coefficients of the DLC films. Those films grown in source gases

with higher hydrogen-to-carbon ratios had much lower friction and wear coefficients than did films derived from

source gases with lower hydrogen-to-carbon ratios. Fundamental and surface analytical studies have led us to

conclude that the presence of protonic, atomic, and molecular hydrogen either within the films or near the sliding

surfaces plays a critical role in the friction and wear properties of these films. Based on the findings of experimental

and surface analytical studies, a mechanistic model is proposed to explain the superlow friction and wear properties

of the films

Key words: PECVD, Diamond-like carbon, super-low friction and wear, lubrication mechanism

INTRODUCTION

The reduction and/or proper control of friction between rolling, rotating, or sliding mechanical parts and components

is extremely important for achieving higher efficiency and lower emissions in most moving mechanical assemblies.

For years, liquid and grease lubricants have been used in combating friction and wear, but the increasingly

demanding application conditions of these assemblies have necessitated the uses of solid lubricants alone or in

addition to liquid lubricants. In certain applications (i.e., aerospace, high-temperature, seals, MEMS), the use of

liquid or grease lubricants is neither practical nor desirable. For such extreme applications, solid lubricant coatings

provide the only choice [1]. Some of the well-known solid lubricants include teflon, MoS2, graphite, Pb, and

diamond and diamond-like carbon (DLC) materials. Depending on test environments and other conditions, these

solids can provide friction coefficients ranging from 0.01 up to 0.2. Quite recently, new forms of carbon films were

developed and demonstrated to provide friction coefficients in the range of 0.001 to 0.01 [2].

Fundamentally, friction results largely from chemical and physical interactions between sliding surfaces.

Deformation and third body interactions can also have a significant influence on friction. In general, physical

roughness on a sliding surface causes high friction. When there is too much chemical or adhesive interactions

between two sliding surfaces, high friction is also observed. Theoretically, one should expect superlow friction from

atomically smooth surfaces with very little or no adhesive interactions.

Among the many solid films, diamond and DLC provide a unique opportunity to achieve superlow friction on sliding

surfaces. For one thing, they represent some of the hardest and the most rigid materials (thus insuring very small

contact area), plus they can also be made very smooth [3]. The amount of chemical or adhesive interactions between

NASA/CP--2001-210948 437 August 6, 2001
slidingdiamondandDLCsurfacescanbemadeverysmall.If welookatthetypesofbondsthatexistbetweensuch

surfaces, wecan mainlysee/findcovalent, vanderWaals,_-_*interactions, capillaryforcesandelectrostatic

attractionand/orrepulsion. Ionic,metallic,ormagneticattractionsdonotexistincarboncoatings(unlesstheyare

dopedwithmetalsorcompounds). Amongthebondsmentionedabove,covalent bondingisthestrongest andif

established betweenslidingcarbonsurfacescanbethedominant causeoffriction.Suchbondingmayoccurbetween

slidingdiamondsurfacesathightemperatures and/orinhighvacuumandleadstoveryhighfriction[4-6].Vander

Waalsbondingistheweakest andmostdesiredbetweenslidingsurfaces. Thisisthekindofbondthatholdsthe

sheetsofMoS2andafewotherlayeredsolidstogether. Theyareknownfortheirexcellent lubricity.

Bycontrollingdepositionparameters andsourcegaschemistryduringdeposition, DLCfilmscanbemadephysically

verysmooth, mechanically veryhardandchemicallyveryinert.Asaresult,onecanachievethesmallest contact area

andtheleastamount ofasperityandadhesiveinteractions. Themainquestionis" howcanwepreparesuchfilms

withextremequalities?" Thispaperwilldescribeindetailthe,inception, realization, andtribological demonstration

of suchasuperlowfrictioncarbonfilm.It will beshownthatwhendepositedonextremelysmoothandrigid

substrates (suchassapphire), thesefilmscanprovidedfrictioncoefficients of0.001-0.005 indrynitrogen[2].Abrief

descriptionfthefrictionandwearmechanisms ofthesefilmswillalsobepresented. Special emphasis willbeplaced

ontheeffectsofsourcegaschemistryandenvironmental speciesonfrictionandwearsincetheyrelatestronglytofilm

microstructure andchemistry. Abriefreviewofpreviousresearchisalsopresented.

SYNTHESISOF DLC FILMS

Brief Review of Recent Developments

In recent years, we have witnessed an overwhelming interest in the production, characterization, and diverse utilization

of DLC films for various applications [3]. The attractiveness of these films stems mainly from the fact that they

combine a very low friction coefficient with high wear and corrosion resistance. Such a combination of unique

properties is not available in most other materials. Furthermore, they are very easy to synthesize and apply on all kinds

of engineering materials including, metals, ceramics, and plastics. Our fundamental understanding of the chemistry and

tribology of DLC and other hard coatings has increased tremendously in recent years and has been used to design,

optimize, and customize coatings that can meet the increasingly stringent applications needs of future tribological

systems. In addition to phase-pure DLC films, metal doped DLC films were also developed and used in a variety of

applications [7]. These nanocomposite, duplex, or multilayered DLC films exhibit better mechanical, thermal, and

chemical properties and are well-suited for certain industrial applications [8,9].

Structurally, DLC films are amorphous and composed of very-short-range-ordered diamond (characterized by sp3-type

tetrahedral bonds) and graphitic phases (characterized by sp2-type trigonal bonds) [10,11]. They can be regarded as

degenerate forms of bulk diamond and/or graphite. Mechanically, they are quite resilient and can have Vickers hardness

values as high as 9000 kg/mm 2 [12]. They are electronically insulating and optically transparent to visible and

ultraviolet light [13]. D1C films can be deposited at temperatures ranging from subzero to 300°C by a variety of

methods, including ion-beam deposition, DC and RF sputtering, arc-plasma, PECVD, and laser ablation. PECVD

method is well-suited for the deposition of DLCs from hydrocarbon source gases (such as acetylene (C2H2), ethane

(C2H_), ethylene (C2H4), and methane (CH4). It can be done at room temperature and with relative ease. The use of RF

or pulse DC bias on substrate is essential for preventing charge build-up, since the DLC films are electrically

insulating. Deposition temperature, gas composition and pressure as well as bias voltage used during deposition can

directly influence the film properties. In particular, bias voltage, deposition pressure, and substrate temperature may

have the greatest influences on the type of C-C bonding as well as the fraction of hydrogen bonded to C.

Synthesis of Super-low Friction DLC Films

For the production of superlow-friction DLC films that will be discussed below, PECVD provides the easiest and

perhaps the most desirable method. The use of RF and/or pulse DC bias prevents charge build-up during deposition.

Source gas and gas mixtures can be precisely adjusted and with the aid of a mass spectrometer, one can monitor the

range of gaseous species present within the deposition chamber. Superlow friction DLC films can be produced on

NASA/CP--2001-210948 438 August 6, 2001
ceramics(suchassapphire) andonthepolishedsurfacesof metalsandalloys(suchasAISIM50ballsandH13steel

flats)aswellashardplasticsbythePECVDmethod. Inourlaboratory, weuseanR&DsizePECVDsystemtoproduce

theDLCfilmsonavarietyofsubstrates. Aforthesourcegases, wecanusepuremethane(CH4), acetylene(C2H2),

ethane(C2H6), ethylene(C2H4), andCH4 + hydrogen (H:) [2,14-16]. As will be elaborated later, our systematic studies

have clearly established a close correlation between film properties and the hydrogen-to-carbon (H/C) ratio in source

gases [2,14]. This ratio is essentially zero for films sputtered from solid carbon targets, whereas 1 for C2H 2 source gas

and 4 for CH 4. For hydrogen-rich CH 4 plasmas, we can add more hydrogen gas to achieve H/C ratios of 10 or more.

The procedure for forming DLC films on steel or ceramic substrates by PECVD involve sputter-cleaning of the

substrates in an Ar plasma for 30 min by applying a 1200-1700V bias. The substrates are then coated with a 50-70

nm thick silicon bond layer by switching to a sputtering mode and sputtering silicon from a target. Silane (Sill4) gas

can also be used to form a bond layer on substrates. Finally, carbon-bearing source gases (CH4, C2H2, C2H6, C2H4,

and CH4 + H2) can be bled into the chamber and the deposition of DLC on the substrates may start. The gas pressure

can vary between 10 to 30 mtorr and the RF bias is at 400-600 V. Further details of the deposition process can be

found in Refs..2,14-16.

TRIBOLOGY OF DLC FILMS

Brief Overview of Recent Work

Tribological properties of DLC films have been investigated extensively by many scientists. In general, it was

demonstrated that when prepared properly, these films can substantially reduce friction and wear between sliding

tribological applications. The reported sliding friction coefficients of DLC films range from 0.001 to 0.5, depending on

test conditions, deposition methodology, and counterface materials [2,17-26]. They can also reduce friction under oil or

water lubricated conditions. The friction coefficients are generally lower under heavier loads and reach low steady-state

friction coefficients faster at higher sliding velocities than at lower velocities [25].

Depending on test conditions, tribochemical (such as tribo-oxidation[27-31]) and structural changes (such as

graphitization [22-26,32-35]) may occur at or near the sliding contact surfaces of DLC films and strongly affect their

friction and wear behavior. It is well known that with respect to graphite, both the diamond and DLC films are unstable.

When subjected to thermal and/or mechanical loadings, they may transform to the more stable graphitic form.

Substantial frictional heating can easily occur at real contact spots of DLC films and the transient temperatures of these

spots can reach very high values. Repeated occurrence of such heating/cooling at real contact spots under high loads

can convert unstable sp3-bonded carbon to the more stable spLbonded form. A combination of TEM, electron

diffraction, bright-field/dark-field imaging, Raman, and FTIR spectra were used by Erdemir et al. and Liu et al. [32-35]

and Ramirez and Sinclair [36] to show that sliding contact surfaces of counterface pins were indeed covered by a thin

transfer layer and that the wear of DLC was controlled by a micrographitization mechanisms.

Chemi-sorption and/or desorption of active gaseous species may also have strong influence. Among others, oxygen

and water molecules have the greatest effect on friction of both the hydrogenated and hydrogen-free DLC films. For

example, in dry air and inert gases, friction coefficients of 0.001-0.02 appear feasible [2,14,15,18,25-28]. However, in

humid air, the friction coefficients may increase to 0.3 while sliding against both metallic and ceramic materials [27].

Increasing humidity has a beneficial effect on the frictional behavior of hydrogen-free DLC films [37]. Furthermore,

films with high doses of Si and Ti seem to be less sensitive to humidity [38,39]. The presence of atomic hydrogen in

test chambers appears to have a beneficial effect on the friction behavior of both the hydrogenated and hydrogen-free

DLC films [40]. However, molecular hydrogen has lesser beneficial effect.

Lubrication Mechanism

In general, the extent of friction between two sliding surfaces is largely governed by the physical and chemical

interactions between them. Rougher surfaces can cause greater mechanical interlocking and hence high friction. For

example, reduction or elimination of physical roughness in diamond films resulted in ultra-low friction [41,42].

When the contribution of surface roughness to friction is reduced, the extent of adhesive interactions becomes the

dominant cause of friction between two sliding surfaces. The strength of adhesive bonding across the sliding

interfaces of DLC films may result from a large variety of short- and long-range interactions. Some of these are very

strong in nature (i.e., covalent.) while others are rather weak (such as van der Waals forces, electrostatic attractions,

NASA/CP--2001-210948 439 August 6, 2001
capillaryforces,etc.).Inshort,hemakingandbreakingofbondsacrosstheslidinginterfacelargelydetermines the

extentoffrictionbetweenslidingDLCsurfaces.

In DLCfilms,covalent or o-bonds cause the strongest interactions (if available) across the sliding interface and

hence give rise to highest friction in these carbon materials. For example, when diamond and hydrogen-free DLC

films are friction-tested in ultrahigh vacuum or at high temperatures, they may exhibit friction coefficients of >1,

presumably because the o-bonds of surface carbon atoms are exposed and are free to form strong covalent bonding

across the sliding interfaces [6]. If moisture or other active gaseous species such as, oxygen, hydrogen, fluorine, etc.,

are introduced into the vacuum chamber, the friction coefficient of diamond and of hydrogen-free DLC drops

precipitously to as low as 0.05, presumably due to the passivation of the o-bonds and thus the elimination of strong

covalent interactions across the sliding interface. This is the most widely accepted explanation for the low-friction

mechanism of diamonds [4,5,43-45].

The other important bonds for the friction and wear behavior of diamond and DLC films are the relatively weak van

der Waals attraction and _z-bonding, which are more relevant to graphite than to diamond or high-quality DLCs.

When graphitic DLCs are tested or the sliding surfaces of DLC are graphitized (under the influence of high thermal

or mechanical loadings at high sliding velocities), the extent of van der Vaals and _z-rc*-attractions may become very

important and dominate the frictional performance of these materials [22, 32-35]. Lastly, in humid air, several mono

layers of water molecules and/or other gaseous species may precipitate on the surfaces of sliding DLC films. These

adsorbed layers may form a meniscus in and around the contact spots and then give rise to capillary forces that cause

high friction. In vacuum or well-controlled dry and inert-gas environments, such forces are greatly reduced or

essentially absent, but other residual gaseous species such as oxygen, hydrogen, etc., may chemisorb on these

surfaces and dominate the friction and wear behavior of diamond and DLC films.

Tribology of Super-low Friction Carbon Films

Systematic studies in our laboratory on DLC coatings over the past 10 years have indicated that the chemistry of source

gases (from which the DLC films are derived) can have dramatic effects on the friction and wear performance of the

resultant films. For example, when tested in dry N 2, a sputtered and/or arc-PVD-deposited DLC film provided friction

coefficients of 0.6 to 0.7. DLC films derived from methane exhibited friction coefficients of 0.015-0.02 (which are 30

to 40 times lower than those exhibited by the hydrogen-free DLC films) [ 16]. Films derived from C2H2Provided friction

coefficients of 0.3 under the same conditions. Furthermore, addition of hydrogen (up to 90%) to CH 4 or C2H 2 source

gases made the resultant DLC super lubricious,. In fact, the films grown in 75%H 2 + 25% CH 4 gas mixture provided

friction coefficients as low as 0.003 in dry N 2 [2,15,16]. Figure 1 summarizes the frictional performance of various

DLCs grown in hydrogen free to highly-hydrogenated gas discharge plasmas. Some of the actual frictional traces of

these films are shown in Figure 2. As is obvious, these results reveal a close correlation between the friction coefficients

of DLC films and the H/C ratio of the hydrocarbon source gases. In general, the higher the H/C ratio of source gases,

the lower the friction coefficient (suggesting that hydrogen plays an important role in frictional behavior of DLC films).

In open air with 50% relative humidity, the friction coefficient of the best film increases to 0.026 as shown in Figure 3,

suggesting that the capillary forces may be in action. The rigidity of substrate material seems to play a major role in

friction as well. The lowest friction coefficient was achieved on sapphire or WC substrates. Figure 4 shows the friction

coefficient of same DLC film produced on sapphire and steel substrates. As far as the wear life of the DLC films are

concerned, they can last rather long. In one test, a 1 micrometer-thick DLC film was able to last over 17 millions cycles

(that translated into 1300 km sliding distance).

Mechanistically, hydrogen can chemically bond to and effectively passivate the free o-bonds of carbon atoms in

diamond and DLC films and make the surfaces of these materials chemically very inert [4,5,43-45]. This is the

widely accepted view used to explain the low friction nature of diamonds. Note that the C-H bonding is covalent and

extremely strong (stronger than single C-C bonds). When extra H2 is used during DLC deposition, several important

events occur and determine the structural chemistry of the DLC. First, it is reasonable to claim that the amount of

hydrogen in the bulk, as well as on the surface, of the DLC films will increase when extra hydrogen is used in the gas

discharge plasmas. Most of the hydrogen atoms are expected to be paired with the o-bonds of carbon atoms, but

some unbonded free hydrogen may also be present as interstitial species in the structure. Increased hydrogen density

in the bulk and on the surface should effectively diminish or even eliminate the possibility of any unoccupied obonds remaining and participating in adhesive interactions during sliding. Second, hydrogen is known to effectively

NASA/CP--2001-210948 440 August 6, 2001
etchoutorremovethesp2-bonded orgraphiticcarbonprecursor fromthefilmsurfaceandthusprevent theformation

ofplanargraphiticphasesand/orcross-linkingthatcangiveriseto_z-bonding [46].WhenDLCfilmsarepreparedin

ahighlyhydrogenated plasmaandunderenergetichydrogenionbombardment, strongC-HbondingratherthanC=C

doublebondingshouldbefavored. Asexplainedabove, theexistenceofresidual _z-bonding thatcanresultfromC=C

doublebondsinDLCcangiverisetohighfriction.Finally,someofthecarbonatoms(atleastthoseonthesurface)

weredihydrated, thatistosay,twohydrogenatomswerebondedtoeachcarbonatomonthesurface. Thiscanoccur

ontheunreconstructed (100)surfacesofdiamondstructures undersomespecial orsuper-criticalonditions[46]that

mayhavebeencreatedbytheenergetichydrogenbombardment inahighlyhydrogenated dischargeplasma. The

existenceofdihydratedcarbonatomsonthesurfaceswillincreasehydrogendensityofthesesurfacesandprovide

bettershieldingorpassivation andthussuperlowfriction.Suchafrictionmodel ispresentedinFigure5forpartially

di-hydratedslidingDLCsurfaces.

Asmentionedabove,otherforcesuchasvanderWaalsandcapillaryforces,aswellaselectrostatic attractioncan

alsobepresent andcauseadhesionadhencefrictionattheslidinginterfacesof DLCfilms.Becauseweranthe

frictiontestsinaclean,drynitrogenenvironment, theextentofcapillaryforcesduetomoistureprecipitationthe

slidingsurfacesshouldbeminimal oressentially absent. AsforthevanderWaalsforces,theywillbepresent athe

slidinginterfacesbuttheirelativecontributions tooverallfrictional forceshouldbeinsignificant becauseweused

veryhighcontact loadsinourtests.Asfortheelectrostatic attraction, sincetheDLCfilmsareingeneral dielectric,

theirslidingsurfacescancertainlyaccumulate saticelectrical charges. Thenthemainquestioniswhetherthese

chargeswillcauseattractionorrepulsion. Whenthefreeelectronsofhydrogenatomspairwiththedanglingo-bonds

of carbon atoms, the electrical charge density is permanently shifted to the other side of the nucleus of the hydrogen

atom and away from the surface. Such a shift in charge density allows the positively charged hydrogen proton in its

nucleus to be closer to the surface than the electron which is used up by the o-bond of the surface carbon atoms.

Therefore, it is reasonable to claim that the creation of such a dipole configuration at the sliding interface should give

rise to electrostatic repulsion rather than attraction between the hydrogen-terminated sliding surfaces of the DLC

films.

In short, hydrogen plays a key role in both the type and extent of chemical/physical interactions and hence the

friction of DLC films. Films grown in hydrogen-rich plasmas are more likely to contain more hydrogen on their

sliding surfaces, in their microstructures than will the films grown in pure CH 4 so that these films are chemically

more inert and their covalent o-bonds (which can cause high adhesion/friction) are absent.

The wear resistance of super-low friction DLC films were rather high and correlated well with their friction

performance. Specifically, films with higher friction coefficients suffer higher wear rates. The difference between

wear rates of balls coated with CH 4 + 75% H2-grown DLC and those coated with CH4-grown DLC films is more than

9 3

one order of magnitude (i.e., 9 x 10- mm/N.m for methane-grown film vs. 4.6 x 10 -10 mm3/N.m for 25% CH 4 -I- 75%

H2-grown film).

SUMMARY AND FUTURE DIRECTION

DLC films have much to offer for future tribological applications. Over the years, great strides have been made in the

deposition and characterization of these films. Their widespread uses by industry will increase when and if the

reliability and quality control issues are adequately addressed, and the quality control mechanisms are developed.

Unlike most other engineering materials, DLC enjoys a combination of super low friction and high wear resistance

under a wide range of sliding contact conditions. Recent trends suggest that with proper process control and structural

engineering (e.g., control of gas plasma chemistry, graded interface and multilayer coatings, and metal-ion doping, etc),

DLC films can be tailored to provide superlow friction, last very long, and function well under demanding tribological

conditions.

ACKNOWLEDGMENTS

This work is supported by the U.S. Department of Energy, under Contract W-31-109-Eng-38. The author thanks his

students and collaborators who participated in the preparation, testing, characterization of DLC-coated samples.

NASA/CP--2001-210948 441 August 6, 2001
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28.Olsen, J.E.,Fischer, T.E.,andGallois, B.:In situ analysis of the tribochemical reactions of diamond-like carbon

by internal reflection spectroscopy Wear, 200 (1996) 233.

29. Marchon, B., Heiman, H., and Khan, R.: IEEE Trans. Magn., 26 (1990) 168

30. Strom, B. D. et al.: J. Tribol., 133 (1991) 689.

31. Ramirez, A. G. et al.: Tribol. Trans., 39 (1996)

32. Erdemir, A., Bindal, C., Pagan J., and Wilbur, P. (1995), Characterization of Transfer Layers on Steel Surfaces

Sliding Against Diamondlike Carbon in Dry Nitrogen, Surf. Coat. Technol., 76-77, 559-563.

33. Liu, Y. Erdemir, A., and Meletis, E. I.: A Study of the Wear Mechanism of Diamond-Like Carbon Films, Surf.

Coat. Technol., 82 (1996)48.

34. Liu, Y. Erdemir, A., and Meletis, E. I.: Influence of Environmental Parameters on The Friction Behavior of DLC

Coatings, Surf. Coat. Technol., 94-95 (1997) 463.

35. Erdemir, A., Fenske, G. R., and Halter, M.: World Tribology Congress Extended Absracts, Mech. Eng. Publ. Ltd.,

London, UK, 1977, p. 488.

36. Ramirez, A. G. and Sinclair, R.: Wear Effects on Microstructural Features of Amorphous Carbon Thin Films, Surf.

Coat. Technol., 94-95 (1997) 549.

37. Holmberg, K. et al.: Diamond Films and Technol., 4 (1994) 113

38. Neerinck, D.: Diamond-like Nanocomposite Coatings (A-C:H/A-Si:O) for Tribological Applications, Diamond Rel.

Mater., 7 (1998) 468.

39. Gangopadhyay, A. K.: Amorphous hydrogenated carbon films for tribological applications .2. Films deposited on

aluminium alloys and steel Tribol. Int., 30 (1997) 19.

40. Paulmier, D. et al: Tribological Behavior of Diamondlike Coatings - Effect of Active Gases in Atomic and Molecular

States, Surf. Coat. Technol. 62 (1993) 570.

41. Hayward, I. P., Singer, I. L., and Seitzman, L. E.: Effect of roughness on the friction of diamond on CVD

diamond coatings, Wear, 157 (1992)215.

42. Erdemir, A. et al: Durability and Tribological Performance of Smooth Diamond Films Produced by Ar-C_0

Microwave Plasmas and by Laser Polishing, Surf. Coat. Technol., 94-96 (1997).

43. Pepper, S. V.: Effect of Electronic Structure of the Diamond Surface on the Strength of the Diamond-metal

Interfaces, J. Vac. Sci. Technol., 20 (1982) 643.

44. Bowden, F. P. and Hanwell, A. E.: The Friction of Clean Crystal Surfaces, Proc. Roy. Soc., London, A295 (1966)

233.

45. Dugger, M. Peebles, D. and Pope, L.: Counterface Material and Ambient Atmosphere: Role in the Tribological

Performance of Diamond Films,"in Surface Science Investigations in Tribology, Experimental Approaches, (Y.-

W. Chung, A. M. Homolo, and G. B. Street, Eds.), ACS Symposium Series 485, American Chemical Society,

Washington D.C., (1992), pp. 72-102.

46. Harris, S. J. and Goodwin, G. D.: Growth on the Reconstructed Diamond (100) J. Phys. Chem., 97(1993)23.

Corresponding Author:

Ali Erdemir

Energy Technology Division

Argonne National Laboratory

Argonne, IL 60439

USA

Phone: 630-252-6571

Fax: 630-252-4798

e-mail: erdemir@m_l.gov

NASA/CP--2001-210948 443 August 6, 2001
0.7

0.6

0.5

_q_ 0.4

O

_D 0.3

0

0.1

ii iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii!!!!!ii!iii:

:: _:::i:::::::::::i::::::::::::i::: :::i::::: ::i::: -_

_ _i_i_i_  ̧ _i_ _i_ _

!!!!!!!!!i!!! !!

R_/_ i..... ; .... :::i_-:: 2?::2 .... i---

_M ! ............ ! ............ !............. !............. ',............

Source Gas Chemistry

Figure 1. Friction coefficients of DLC films derived from various source gases (all tests run in dry nitrogen).

E

I]3

O

O

cO

"r"

LL

0.9

0.8

0.7

0.6

0.5

0.4

0.3

0.2

0.1

I --@-- lOO% CH4

|--1_t-- 100% C2H2

_- .... 1 .... ] ' 'I_H-FreeDLC

_-[ : : "| -_25% CH4+75%H2

/

r--_t_....... :-. -H/C=e- - -:....

i ........... ' T

1'.

2-

- ! , I1 , , I_WaP• _ -

_-_ ................. "_ .................. i I_ - --: ......

.'_.': ,i ,li,._ _ ,,-_, : ..... . I I&IG "

.':ll II 'g .... 1_. iI_ , i I/_tl:

......... ; -_-'-I_ -I I I_1 I _I ,,',_' _ ,'I • ' " I'_

-,t.,_,,_.... '.......... '._. ,.. ;, '.._.,,._,...'.-

-• - , , -• • I, • • I•

-• II II .I I Coa3

--i II * , I • ,64 • I i Ir , -

_i , , i ' I ii ' -

-_ ' ' H/C=IO ''%;1 ' :

.............. H7C--'4 ............. '• ........

-I _I )_I I _I __I _i II --

_.______-_.__--_._--_._-__--_____. _, :_i___.-2.2__'_g2

Ig Failure

ing Failure

0 500 1000 1500 2000

Time (s)

Figure 2. Actual friction traces of some of the DLC films presented in Figure 1. Note that higher the H/C ratio of gas

discharge plasma, the lower the friction coefficient. (Test Conditions: load, 10 N; speed, 0.3 m/s; temperature, 22°C;

environment, dry nitrogen).

NASA/CP--2001-210948 444 August 6, 2001
0.06

0.05

0.04

,m

F:

ill

,_, 0.03

0

,to

0.02

'_.

LI.

0.01

\ Perhaps due _o capillar_

.............. ',-50%R:H. _ ..... " ....

...... " ...... :...... :. ..... " ...... :..... I

I

J

:0% R.H.

0 500 1000 1500 2000 2500 3000

forces

Time (s)

Figure 3. Effect of humidity on frictional behavior of DLC film grown in 25% CH4 and 75% H2.

0.03

0.025

c-

.- 0.02

E

o 0.015

¢-

o

0.01

.m

Ii

0.005

'-r--i_l_r-- "--r--r_r--r-- _

I ...... DLC-coated Sapphire f

i i i i i

0 1000 2000 3000 4000 5000 6000

Time (s)

Figure 4. Friction coefficients of DLC-coated sapphire and M-50 balls against DLC-coated sapphire and H13 disks,

respectively, in dry nitrogen (Test Conditions: load, 10 N; speed, 0.3 m/s; temperature, 22°C).

NASA/CP--2001-210948 445 August6,2001
Load Load

-_'Shear

Interface

+

Partially di-hydrated

Sliding carbon surfaces

H-terminated

C atoms

Figure 5. A mechanistic model for super-low friction behavior of hydrogenated DLC surfaces with partial dihydration. Situation for individual carbon atoms terminated by hydrogen is also illustrated.

NASA/CP--2001-210948 446 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FRICTION AND WEAR OF NANO-SIZE DIAMOND FILMS

WITH AMORPHOUS PHASE

Koichiro Wazumi

Japan Fine Ceramics Center, Joint Research Consortium of Frontier Carbon Technology

1-1 Higashi, Tsukuba, Ibaraki 305-8565, JAPAN e-mail:kwaznmi@nimc.go.jp

Yoshinori Koga, Akihiro Tanaka

Research Center for Advanced Carbon Materials, AIST Tsukuba Central5

Higashi, Tsukuba, Ibaraki 305-8565, JAPAN

ABSTRACT

Tribology experiments were performed with nano-size diamond films to investigate the influence of the

amorphous phase among the diamond films for the reduction of the friction and wear of the films and balls. Diamond

films were prepared by microwave plasma CVD with the mixture of hydrogen and methane gas under the high

methane content ration of 5, 9, 17 and 23 at% on the Si (100) wafer. The films deposited under 5 and 9at% methane

content ratio included large size grains (200-500nm). XRD analysis showed the sharp diamond (111) peak. On the

other hand the films deposited under 17at% and 23at% methane content ratio showed the nano size diamond grains.

XRD analysis showed the broad diamond (111) peak were observed in both films. Especially the film deposited with

23at% methane content showed the amorphous peak under the diamond (111) crystal peak indicating the existence

of the amorphized grains. Ball on disk experiments were done using the steel ball in dry air condition. The films

prepared with the methane content of 5 and 17at% showed the high friction coefficient of 0.5 and 0.22. On the other

hand the films deposited under 23at% methane gas showed the low friction coefficient of 0.11 and the wear rate of

the film and ball are 1.15xl0-7mm3/Nm and 4.5xl0-7mm3/Nm respectively which are relatively low value. Low

friction coefficients and the low wear rate of the ball in this case are considered due to the existence of amorphous

phase in nano-size diamond film from XPS analysis.

Keywords: Nano-size diamond, Microwave Plasma CVD, Friction Coefficient, Wear Rate, Steel-ball.

INTRODUCTION

Low friction and high wear resistive carbon films are currently great interest for the wear protection. Recent studies

showed nanocrystalline diamond film is one of the prospective materials as a solid lubricating materials (ref. 1,

2,3,4). It exhibits the smooth surface, low friction coefficient due to the small grain size and high wear resistivity due

to its high hardness. However nanocrystalline diamond film still has a problem of ploughing characteristics of the

films against the counterparts. As a results it causes the increasing of the wear rate of the ball which is much larger

than that of a-C films. In our previous work we synthesized the nano-size diamond films in microwave plasma CVD

in which grain size were varied from 20nm to 50nm (ref.5). However the results of the friction test indicated that the

wear rate of the steel ball was 3xl0-_mm3/Nm which was relatively higher value. From this point of view, we

recently prepared the nano-size diamond films under the high methane content up to 23at% and enhance the

amorphous phase in diamond films. These experiments resulted in the improvement of the tribological properties,

especially the specific wear rate of the steel ball has been improved. In this paper we firstly discuss the structural

properties and hardness of these films and next we discuss the wear properties of the films and balls.

EXPERIMENTALS

Diamond films were prepared by antenna-induced microwave plasma CVD. In this CVD system, eight metal

antennae around the cylindrical quartz tube transmit the microwave power inside the chamber. Therefore electrical

fields from each antenna are overlapped, and it makes possible to create the uniform plasma at the center of the

NASA/CP--2001-210948 447 August 6, 2001
chamber, whichresultintheuniformthicknessdistributionofthefilmsoverthe4-inchsubstrate. Eacheightmetal

antennae areconnected withthemicro-wavedipoleantennaperpendicular tothewaveguide.

Asapre-treatment ofthesubstrate, Si(100)waferswerescratchedbythefinediamondpowders(5nmand1Hm)

andtheywerecleanedultrasonically inacetonebathbeforeintroducingintotheCVDsystem. Depositionofnanosizediamondfilmswereconductedundertheconditionofhydrogenflowrateof 100sccmandmethaneflowrateof

5-30sccmatthepressureof 16.5-24Torrwiththemicrowavepowerof2000Wandthesubstratetemperature was

650°C.Depositiontimewas5hours. Thicknessof depositedfilmswereabout500nm. Wecallthesamplesprepared

withthemethaneflowrateof5sccm(5at%),10sccm(9at%),20sccm(17at%)and30sccm(23at%)assampleC1,

C2,C3andC4respectively. Scanningelectronmicroscopy (SEM)wasusedtodeterminethesurfacemorphology.

Thecrystallinityofthediamondgrainwasmeasured byX-raydiffraction(XRD).RAMANspectroscopy wasusedto

characterize thbondingstateofcarbonatoms. Frictioncoefficient wasmeasured byCSEMball-on-disktesterat1N

loadwiththeslidingspeedof 30mm/sinthedryairconditionattheroomtemperature byusingthesteelballwhich

containedseveral eements(1.5at%ofCr,lat%ofC,0.5at%ofMnand0.25at%ofSi).Hertizancontact sresswas

estimatedas0.31GPa. AWearratewascalculatedfromthedepthanalysisoftheweartracksofthefilmsandballsby

surfaceprofilemeter. Microhardness ofthefilmsweremeasuredbyNanoindenter XP(MTS).X-rayphotoelectron

spectroscopy (XPS)wasusedtoanalyzethechemical shiftofcarbonbondingbeforeandafterfrictiontestatthe

surfaceoftheweartrack.

RESULTS AND DISCUSSIONS

Fig.1 shows the SEM micrographs at the surface of sample C1,C2,C3 and C4 respectively. A large grains (50-

400nm) are observed at the surface of the sample C1 and 100-200nm grains are observed in sample C2. On the other

hand nano-size grains (30-50nm) are seen at the surface of the sample C3 , C4, which were deposited relatively

under high methane condition. RAMAN spectra are shown in Fig.2. Each spectra shows the diamond peak

m

Sample C1

"_ ._

Sample C3

Fig. 1

_.,. _...-..-o%-o .-.....-._..oO_....o.....-..o..\-.._ _.%5°.....o.o.,

',:4-" ".',:._::,:,:- g.:.:..X',." w.:, ,,',:,', "*:.",'.;-:,:

°'°°. -'. i'o_*, e'.:._ _°._.°. °%¢o" .'.'°'-%'.'°%'. "o" '.'.

°_ ........ .,°.°p;,,- .... _ ....... .'_ _.,......._>>......,..:.:......... _,......._,.,,......,,..

• ._.._. °.°..°._.oO_.._ %°°.° °°.- *..._.%..°.;-_._ .._..-°. o-... r_,,, ,;,:,;.° ..... eo,,ovX,,., ,: ....... .,,°,

_'.°°'._.° ".'._.%°.'°" "o°3".'._°_.'_. . o" • "._ ° ° _._'.."" • °i'."2

. ...°;......,,°.,,,_4,.....,.°,,,°.,,,.: °'.. .',_.:.I,:i--.,._:-",:.,.,%_

tf'_.':-.:,5¢.°..¢e.-...'_.,-,.''.°.'.._,;*°',,o.,-.- _o.'.-_-°°4,.,:

.._' %%% • g* .°o_.%° ...°*...%o. °. %..-.°.-.o .o._-._:o..; _-°

Sample C2

S ample C4

SEM micrographs

NASA/CP--2001-210948 448 August 6, 2001
._o

0a

._o

r-

, , , , i , , , , i , , , , i , , , , i , , , ,

ili iii

500 1000 1500 2000 2500 3000

-1

Raman shift e_ergy(cm )

rFig.2 RAMAN spectra

U_

e-

_D

.... ,.... ,.... ,.... .... ,....

/

....

, , , , , :• 7

,_rnV_C_:- :-:-: j't : -I

- 4 .... ' ,%% '

sam _)le C2',

"(_'I_ :).3-at°1o- ..... - - - ++ - _., '- - -

F CH :17at_ , , , _..,J,_

/ 4 _,f@-._-;-6-_i.

FsampleC4, . , : [_ '.

I" CH4--'23at_ ' _' , .'t, ,(,F"_

• . -, , _ _.._. ,

38 39 40 41 42 43 44 45 46

angle(2 0)

Fig.3 XRD patterns

at the energy of 1333 cm -1 . The broad peaks were observed at 1150, 1330 and 1450 cm -1 RAMAN shift energy.

The peak intensity of the diamond line decreased with increasing the methane gas flow rate during the deposition.

This suggests that grain size changed into small with the increasing of the methane flow rate. XRD measurement

from the glancing angle was carried out in order to examine the crystallinity of the diamond films. Fig.3 shows the

diffraction peaks from the diamond (111) plane which are divided into the crystal diffraction peak and broader

amorphous peak (hatched area in the each spectra). These figures indicate that the diffraction peaks of crystal line are

sharp in sample C1 and C2 but little wider in sample C3 and C4. The grain sizes were estimated in Fig.4 by using

S cherrer formula from the width of (111) diffraction peak after removing the influence of the diverging X ray which

is inherent to the optical system in this equipment. In sample C1 in which small grains are layered under the large

grains, average grain size is calculated as 85nm. In sample C2, average grain size is calculated as 155nm. The grain

sizes of sample C3 and C4 are calculated as 47nm, 32nm respectively which are in good agreement with the results

of SEM observations of the nano-size diamond films in Fig. 1. Fig.4 also shows the increasing of the amount of the

amorphous phase in the diamond films with increase of the methane gas concentrations during the deposition.

Especially between sample C3 and C4, amorphous phases increased significantly. An amorphous phase in sample C4

are distributed in the wide diffraction range. Although SEM micrograph in Fig.2 shows that morphologies of

sample C3 and C4 are the same, the results of XRD measurement suggest that the crystallinity of these samples are

different.

15_

N

m 1_0

5O

200 .... , .... , .... , .... , ....

.... i .... i .... i .... i ....

5 10 15 2(_ 25

Methane ratio(at%)

i00

8O

A

6o

_go40

t_

-r2O

.... i .... i .... i .... i .... i .... i .... i ....

, , , , , ,

iiii 

,,,i .... i .... i .... i .... i .... i .... i,,,

5 9 13 17 20 23 26

CH flow ratio (at%)

4

Fig.4Calculated grain size Fig.5 Hardness of the nano-size diamond

NASA/CP--2001-210948 449 August 6, 2001
0.8

o

._o 0.6

8

c 0.4

o

I_

, ,,_ _'_,_ c/.

, , , i , , , i , , , i , , , i , , , i , , ,

2000 4000 6000 8000 10000 12000

Number of revolutions

Fig.6 Friction coefficient (steel ball)

0.8

o

o

,i

__ 0.6

o

o

0.4

o

L-T 0.2

i i i I i i i I i i i I i i i I i i i I i i i

2000 4000 6000 8000 I0000 12000

Number of revolutions

Fig.7 Friction coefficient(SiC ball)

The significant increase of the amorphous phase in sample C4 was also examined by the micro-hardness

measurement shown in Fig.5. The average hardness of the sample C1, C2, C3 and C4 were 86Gpa, 74Gpa, 79Gpa

and 55Gpa respectively. These data indicate that the hardness remains constant at lower methane content below 17

at%, however hardness value decreases gradually over the 17% methane ratio. It is thought that the degredation of

the crystallinity at these ranges is in good accordance with the results of XRD analysis. As a whole the

morphologies of the CVD-diamond films in this system have changed from the films with coarse grains to the nanosize diamond films when methane content has increased to 17at% under the deposition condition of 2kW MW power

and the pressure of 16.5-24 torr with the substrate temperature of 650°C. Moreover amorphous phases have

increased when methane content has increased over 17at%.

The behaviors of the friction coefficient of the sample C1, C3 and C4 for the steel ball are shown in Fig.6. The

film of sample C2 was peeled off soon after a friction test started. The results of the friction test indicate that sample

C1 containing the coarse diamond grains showed a high friction coefficient of 0.6 constantly from the beginning to

the end and sample C3, which contains the nano-size diamond grains, also showed high friction coefficient of 0.8-

0.9 at the beginning of the friction test though it reduced to the 0.22 at the end of the test. On the other hand sample

C4 containing the amorphized nano-size grains exhibited the sudden reduction of the friction coefficient at the

beginning of the test and it reached to lowest value 0.11 at the 12000 revolution times compared with other samples.

The wear rate of the films of sample C4 was 1.15xl0-7mm3/Nm. And it did not have so much difference with that of

the sample C3 (1.54x 10-7mm3/Nm). However it should be noted that the wear rate of the ball materials showed a big

difference between sample C3 and C4. Because the wear rate of the steel ball for sample C3 turned out to be 3.5xl0-

_mm3/Nm and that of the sample C4 was 4.5xl0-7mm3/Nm ,lower value. This low wear rate of the ball is thought to

be due to the existence of the amorphous phase among the film.

In order to examine the effect of the amorphous phase in nano-size diamond films for steel ball, XPS measurement

was carried out. Cls state of sample C3 and C4 were analyzed at the pure surface of the films and the surface upon

the wear tracks. The results showed Cls spectra at the pure surface of the sample C3 exhibited the peak at 285.5 eV

which was near the peak position of Cls spectra of sp 3 bonding in CVD-diamond (ref.6). A shoulder peak was also

observed at the lower energy region indicating the existence of some amount of sp 2 bonding in the films. However at

the surface of the wear track, Cls peak was shifted to 285.1eV suggesting the increase of the carbon sp 2 bonding

after the friction tests. It is thought that friction coefficient has been decreased because of the transformation of the

carbon bonding after a several sliding. On the other hand Cls spectra at the pure surface of the sample C4 exhibited

the peak at 285.35 eV which was lower than that of sample C3 suggesting that the sp 2 ratio was relatively higher at

the pure surface of the sample C4 rather than sample C3. Cls peak at the surface of the wear track was positioned at

285.3 eV which was not shifted after the sliding test. Therefore it is thought that an existing of the relative amount of

amorphized sp 2 phase in nano-size diamond films gives influence to decrease the friction coefficient and wear rate

for steel ball. Fig.7 shows the friction behaviors of the deposited films for SiC ball. The friction coefficient of the

NASA/CP--2001-210948 450 August 6, 2001
sample C4 decreased to 0.11 at the end of the friction test compared with that of sample C3 (0.16). And the wear

rates of the ball are 2.3xl0-6mm3/Nm and 9. lxl0-7mm3/Nm for sample C3 and sample C4 respectively. These results

suggest that amorphous phase among the nano-size diamond grains takes good effect also for reduction of the

friction and wear rate of the SiC ball. However the wear rate of the films for sample C3 and sample C4 showed 6.3

xl0-Smm3/Nm and 3.9xl0-7mm3/Nm, indicating that film wear increased due to the appearance of the amorphous

phase when sliding with SiC ball. The reason is thought that amorphous sp 2 phase has influence for the surface of

the SiC ball to increase the adhesiveness at the interface. As a result the surface of the film was scratched and wear

rate increased, although friction coefficient decreased because of the decreasing of the shear strength. These results

indicate that the effect of the amorphous phase in nano-size diamond films depend on the ball materials. However it

is clear that it promotes the tribological properties in case of the steel balls

CONCLUSIONS

Tribology experiments were performed with nano-size diamond films to investigate the influence of amorphous

phase among the diamond films. Diamond films were prepared by microwave plasma CVD with the mixture of

hydrogen and high contents of methane gas. The films deposited under 17at% and 23at% methane ratio show the

nano size diamond. The film of 23at% methane ratio shows the broader XRD peak which can be separated from the

diamond (111) crystal peak indicating the existence of the amorphized grains. Ball on disk experiments were done

using the steel and SiC ball in dry air condition. The film deposited under 23at% methane gas shows the low friction

coefficient of 0.11 and the low wear rate of the films and balls which are 1.15x 107mm3/Nm and 4.5x 107mm3/Nm

in case of the steel ball. Low friction coefficients and the low wear rate of the balls are considered due to the

existence of amorphous phases in nano-size diamond films that were confirmed by XPS analysis.

ACKNOWLEDGEMENTS

The authors would like to express deep appreciation to Dr. Yamamoto in Research Center for Advanced Carbon

Materials for his advice to the XPS analysis and Ms. Gotoh for her sincere help in the XPS measurements. This work

was supported by New Energy and Industrial Technology Development Organization (NEDO) as a part of the

Frontier Carbon Technology (FCT) project of Agency of Industrial Science and Technology (AIST) of Japanese

government.

REFERENCES

1. K.Miyoshi, Mater. Sci. Eng., A209,1996,38.

2. K.Miyoshi, Richard L.C. Wu and Alan Garscadden : Surf. Coat. Technol., 54-55,1992,428.

3. Erdemir, G.R.Fenske, A.R.Krauss, D.M.Gruen, T.McCauley and R.T.Csencsits : Surf. Coat. Technol.,

120-121,1999, 567.

4. RHollman, O.Wanstrand and S.Hogmark : Diamond Relat. Mater. : 7, 1998, 1471.

5. K. Wazumi, Y.Koga and A.Tanaka : Proc. International Tribology Conference Nagasaki 2000 (in press).

6. J.Shirafuji, Y.Sakamoto, A.Furukawa, H .Shigeta and T.Sugio : Diamond Relat. Mater. : 4, 1995, 984.

NASA/CP--2001-210948 451 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

LOW TEMPERATURE GROWTH OF NANOSTRUCTURED DIAMOND FILMS

ON METALS

Paul A. Baker, Shane A. Catledge, and Yogesh K. Vohra

Department of Physics

University of Alabama at Birmingham (UAB)

Birmingham, AL 35294-1170, USA

E-mail: yto,ohra(_uab.edu

ABSTRACT

The field of nanocrystalline diamond and tetrahedral amorphous carbon films has been the focus

of intense experimental activity in the last few years for applications in field emission display devices,

optical windows, and tribological coatings. The choice ofsubstrate used in most studies has typically been

silicon. For metals, however, the thermal expansion mismatch between the diamond film and substrate

gives rise to thermal stress that often results in delamination of the film. To avoid this problem in

conventional CVD deposition low substrate temperatures (< 700 C) have been used, often with the

incorporation of oxygen or carbon monoxide to the feedgas mixture. Conventionally grown CVD diamond

films are also rough and would require post-deposition polishing for most applications. Therefore, there is

an obvious need to develop techniques for deposition of well-adhered, smooth nano-structured diamond

films on metals for various tribological applications. In our work, nanostmctured diamond films are grown

on a titanium alloy substrate using a two-step deposition process. The first step is performed at elevated

temperature (820 °C) for 30 minutes using a H:/CI-IdN2 gas mixture in order to grow a thin (approx. 600

nm) nanostructured diamond layer and improve film adhesion. The remainder of the deposition involves

growth at low temperature (< 600 °C) in a H:/CI-L/O: gas mixture.

1800

.i,.,,a

iw,,i

O

I I I I I I I

1500

1200

900

600

300

0 I I I I I I I

0 25 50 75 100 125 150 175 200

Time (minutes)

Figure 1: Laser reflectance Intefferometry (LRI) pattern during growth of a nanostructured diamond film

on Ti-6A1-4V alloy. The first 30 minutes are at a high temperature of 820 °C and the rest of the film is

grown at a low temperature of 580 °C. The fringe pattern is observed till the very end due to extremely low

surface roughness of 40 nm.

The continuation of the smooth nanostructured diamond film growth during low temperature deposition is

confirmed by in-situ laser reflectance interferometry (Figure 1) and by post-deposition micro-Raman

spectroscopy and surface profilometry. Similar experiments performed without the starting nanostructured

diamond layer resulted in poorly adhered films with a more crystalline appearance and a higher surface

roughness. This low temperature deposition of nanostructured diamond films on metals offers advantages

in cases where high residual thermal stress leads to delamination at high temperatures.

We acknowledge support from the "Advanced Materials Cluster" of the NASA-EPSCoR Alabama

Program under Grant No. NCC5-165.

Keywords: nanostructured diamond, metals, interferometry, adhesion

NASA/CP--2001-210948 452 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

The Tribology of Amorphous Diamond

Shenq Y. Luo t, Jui-Kang Kuo 2,Brian Yeh 2, James C. Sung 2,Chuang-Wen Dai 3, and Tsung J. Tsai 1

1. Department of Mechanical Engineering, Huafan University, Taipei, Taiwan, R.O.C.

2. Kinik Company, Taipei, Taiwan, R.O.C.

3. Department of Mechanical Engineering, Tamkang University, Taipei, Taiwan, R.O.C.

ABSTRACT

_ _ is_ _ _[_ carbon(DLC)wi_ ix_pmiesv_ close_om_e diama_ _ _

_ is the low flier(realcoetfdent and high wear tesislar_. _ WC dis_ were coa_ byamtxttx_ diamt_ deposited

by caltrxtic arc. These disks were mbl_ cyclicallyby pim rnade of akmaknzn-silicon alloy(4032), ¢mtxm steel (52100), anda_henim

c_nies (AI_). It was found _at frica_ _ atter abeperiodof"mak-in," was significantlyret_ed wt'en _ wiah

fiat ofuncoa_ carbides. However, the fdcfional coefficient(02) with steel was more than twice that with A1 alloy(0.9) and altmfim

(0.8).I1aisahxxmal high frictia-al coeti_cientisrx_sa"olydue tmhe clx_cal in_ _we_n arface _ of_ _ _

The wear mt_hanisms include mechanical attrition, chemical adhesion, and fatigue. The wear loss was the highest for the

softer A! aBoy,and lowest for 52100 steel. However, the wear of amorphous diamond was the highest when it was rubbed

against Al-al_. It was postulated that diamond was worn down primar_y by the hardened layers of A! alloy and ox)gen

formed _ duringthe rubbing actio_

NASA/CP--2001-210948 453 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

PEEM AND NEXAFS ANALYSIS OF WEAR DEBRIS OF SI INCORPORATED

DIAMOND LIKE CARBON FILMS IN VARIOUS ENVIRONMENTS

Sejun Park, Kwang-Yong Eun, Kwang-Ryeol Lee

Thin Film Technology Research Center, Korea Institule of Science and Technology, RO. Box 131,

Cheongryang, Seoul, 130-650, KOREA

Andreas Scholl, Frithjof Nolting, H. Padmore

Lawrence Berkeley National Lab. University of Berkeley, Berkeley, CA 94720, USA

Si incorporation into diamond4ike carbon (DLC) film has been known to improve the mechanical

properties and tribological behavior of the films. While the friction behavior of pure DLC film is

significantly dependent on the test environment such as humidity or oxygen, it has been reported that the Si

addition can reduce the environmental dependence of the friction behavior. However, Si incorporated DLC

film also shows a different friction behaviors depending on the test environment, even if the friction

coefficient between the film and steel ball in humid air is much smaller than that of pure DLC film. Among

three test environments of dry air, humid air and vacuum, improved friction behavior by Si addition was the

most significant in dry air.

In order to understand the mechanism of the different friction behavior, we investigated the chemical

structure of the debris by PEEM and NEXAFS at Advanced Light Source of LBNL. Tribological test was

performed using environmental ball on disk type wear-rig AISI52100 steel ball was used for the wear test.

The films were deposited on Si wafer by r.f.-PACVD using mixture of benzene and silane as the precursor

gas. Si concentration of the film was varied from 0 to 9.5 at.% by adjusting the concentration of silane.

The film was deposited at the negative bias voltage of 400 V and the deposition pressure 1.33 Pa.

In humid air, it was observed that the debris were partly polymerized and finely dispersed when the Si

concentration was larger than 5 at.%, which resulted in low and stable friction behavior. In dry air, the

chemical bond structure of the debris was essentially the same as those in humid air. However, the debris of

smaller size has more spherical shape. These results imply that the rolling action of the debris seems to play

major role in dry air condition. However, in humid air, the debris seems to conglomerate with absorbed

water molecules. In vacuum, quite different chemical bond structure and morphology of the debris was

observed. The debris at low Si concentration showed typical NEXAFS spectrum of a polymer

(polymethylmethacrylate). The polymeric component in the debirs decreased as the Si concentration

increased. Based on these observations, the environmental effect of the friction behavior will be discussed

in terms of the tribo-chemical reaction between the films and the steel ball and the behavior of the debris

during the wear test.

NASA/CP--2001-210948 454 August 6, 2001
6

DEPOSITION AND CHARACTERIZATION OF DLC AND

NOVEL CARBONS

Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Molecular Dynamics Simulations of the Structure of Amorphous

Carbon Containing Nitrogen

D.R.McKenzie

Department of Applied and Plasma Physics

University of Sydney

A. Merchant and D.G. McCulloch

Department of Applied Physics R.M.I.T. University

Melbourne

We present the results of a full molecular dynamics

simulation of the structure of amorphous carbons containing

nitrogen. We used an ab initio approach based on the Car

Parrinello method in the generalised gradient approximation.

Three different concentrations of nitrogen were studied.

The local spin density approximation was used to study

the effect of ionising the nitrogen site in order to

understand the non volatile memory effect in some nitrogen

doped carbons.

Professor David McKenzie,

Head, Department of Applied Physics, A28

University of Sydney

NSW 2006, Australia

Tel: +61-2-9351 3577

Fax: +61-2-9351 7725

email: mckenzie@physics.usyd.edu.au

http://www.physics.usyd.edu.au/apphys/applied.html

NASA/CP--2001-210948 457 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

EFFECT OF NITROGEN IN AMORPHOUS CARBON FILMS DEPOSITED BY

ELECTRON CYCLOTRON RESONANCE PLASMA SPUTTERING METHOD

T. OHANA 1, T. NAKAMURA 1, A. TANAKA 1'2, Yo KOGA 1'2

1 Research Center for Advanced Carbon Materials, AIST TSUKUBA Central 5, 1-1-1 Higashi, Tsukuba, Ibaraki

305-8565, JAPAN

Fax: +81-298-61-4573, Tel: +81-298-61-4821, E-mail: t.ohm:a@a!st.go.jp

A. GOTO _

2 Joint Research Consortium of Frontier Carbon Technology, JFCC, c/o AIST TSUKUBA Central 5, 1-1-1 Higashi,

Tsukuba, Ibaraki 305-8565, JAPAN

The diamond-like carbon (DLC) films are of interest in connection with protective materials. With the

development of computer technology, the storage capacity of hard disks has been an explosive increase. In order to

accomplish a high-density memory, the distance between the magnetic media and the magnetic head must be

decreased. The protective overcoat of the magnetic head plays an important role under an extreme condition.

Previously, we reported the deposition of DLC and carbon nitride films using an electron cyclotron resonance

(ECR) plasma sputtering method (ref. 1). Because the ECR sputter technique can give the stable plasma at low

pressure and the high plasma density easily (ref. 2 and 3), the formation of the hard film with low friction coefficient

can be expected. We found that the DLC films deposited by ECR plasma in low-pressure atmosphere showed

significantly low friction properties. However, the deposited film has high internal stress and readily exfoliated. The

properties of films were improved by the mixing of nitrogen atoms in the films. Although the excess of nitrogen

atoms in the films showed the high friction coefficient and high wear rate, the films contained a small amount of

nitrogen showed the low friction and wear resistance. However, the friction coefficients of the films contained

nitrogen were slight higher than that of DLC film. In this paper, we report the contribution of bonds for tribological

properties of the films containing a few amount of nitrogen atom.

The films were deposited on Si (1 0 0) substrates by sputtering of graphite target in either argon and/or nitrogen

gas atmosphere. Microwave (2.45 GHz) was introduced into a plasma chamber and a negative bias was applied to

the target with DC power supply. The typical deposition conditions were as follows, pressure; 2.0 x 10 .2 to 4.0 x 10 .2

Pa, target bias; -350 V, microwave power; 400 W, substrate temperature; room temperature, deposition time; 30 min,

and substrate bias; self-bias. The substrate was washed by acetone, ethanol, and water in ultrasonic bath for 5 min,

and then the surface was pre-sputtered by argon plasma for 10 min before deposition. The characterization of the

deposited films was investigated by means of Raman spectra and XPS measurements, and the tribological properties

of the films were investigated by means of ball-on-disk method. The friction measurement was conducted under lowhumidity (15 - 25 %) against SiC ball. The sliding speed is 100 rpm, and diameter is 5 mm. The ball was contacted

with films at a load 0.1 N.

Figure 1 showed the Cls region of XPS spectra of the DLC films. These spectra were decomposed to three

constituents. The peaks of curves are 288.7, 286.2, and 284.8 eV for the film deposited at 2.0 x 10 .2 Pa. The binding

energy at ca. 289 eV would be ascribed to the C-O. The oxygen would be ascribed to water existed in the plasma

chamber and the graphite target as an impurity. The ratio of the peaks area of 286 and 285 eV are 0.35, 0.48, and

0.66 for the pressure at 2.0 x 10 -2, 3.0 x 10 .2 and 4.0 x 10 .2 Pa, respectively. It is indicated that the ratio of the peaks

area of 286 (sp 3) and 285 eV (sp 2) increased with the increasing of pressure. Moreover, these films showed lowfriction coefficient around 0.1 to 0.2. With increasing of the pressure, the tendency of increasing of friction

coefficient was also observed. It is considered that there are some relationship between the structural analysis and

tribological properties, though the friction data contain a few errors.

The CNx films deposited in an argon atmosphere with nitrogen gas showed the similar relationship between the

structural analysis and tribological properties. Figure 2 showed the Cls region of XPS spectra of the CNx films.

With increasing of nitrogen atomic ratio, the Cls XPS curve split to two parts. It would be considered that the

nitrogen atom enhanced the contribution of constitution of the peak at 286 eV. The friction coefficients of the CNx

films are 0.31, 0.38, and 0.45 for figure 2 (A), (B), and (C), respectively. The friction coefficients increase with the

increasing of nitrogen atomic ratio. Therefore, the excess nitrogen atom in the films will interfere with low friction

NASA/CP--2001-210948 458 August 6, 2001
coefficient. ItwouldbenotonlythereasonwhytheCNxfilmscontainedxcessofnitrogenshowedthehighfriction.

TheeffectofnitrogenatomsinDLCfilmsisstillnotclear.It wouldbedifficulttocomparewiththespectrumof

CNxfilmandthatofDLCfilm,becausethespectrumofCNxarecontributedfrommanyspecieslikeC-N(sp3), C-N

(sp2), C-C(sp2),andC-C(sp2)(ref.4).However, weconcludethathechemical structureofthefilminfluencestothe

surfaceproperties.

(A)

1,1

;'1

¢

I

295

(B)

,' l

,, l

./:,. ,

........... _"......,.:"_".".:,..!%.....

i i i I I I I

290 285

Binding Energy/eV

(C) ,4,

i

i ,,,I

I ''J

1,1

I_ ii

I

..... .\

_ .. . "-.-,..._.'%._

I I I I

280 295 290 285 280 295 290 285 280

Binding Energy/eV Binding Energy/eV

Figure 1. Cls region of XPS spectra of DLC films, (A) deposited at 2.0 x 10 .2 Pa, (B) 3.0 x 10 .2 Pa, and (C)

4.0 x 10 .2 Pa.

292 290 288 286 284 282

Binding Energy / eV

"t,, 1

/4 i' j,/ '\'_, 'i

292 290 288 286 284 282

Binding Energy / eV

/i !t

292 290 288 286 284 282

Binding Energy / eV

Figure 2. Cls region of XPS spectra of CNx films, (A) deposited in N2/Ar = 1/9, (B) N2/Ar = 1/1, and (C) in

nitrogen atmosphere.

REFERENCES

1. Ohana,T., et al.: The structure and tribological property of amorphous carbon and carbon nitride films

prepared by ECR plasma sputtering method. Diamond Relat. Mater., in press.

2. Ono, T., Takahashi, C., and Matsuo, S.: Electron Cyclotron Resonance Plasma Deposition Technique Using

Raw Material Supply by Sputtering. J. Appl. Phys., vol. 23, no. 8, Aug 1984, L534.

3. Matsuo, M. and Ono, K.: Ion energy analysis for sputtering-type electron-cyclotron-resonance microwave

plasma. J. Appl. Phys., vol. 64, no. 10, Nov 1988, 5179.

4. Cheng, Y. H., et al.: Deposition of carbon nitride films by filtered cathodic vacuum arc combined with radio

frequency ion beam. Diamond Relat. Mater., vol. 9, 2000, 2010.

NASA/CP--2001-210948 459 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

AN INVESTIGATION OF SI-DLC FILM STRUCTURE AND CHEMISTRY

RELATIONSHIPS TO DEPOSITION PROCESS CONDITIONS AND MECHANICAL

PROPERTIES

Ryan D. Evans , Jeffrey T. Glass, Philip W. Morrison, Jr.

Dept. of Chemical Engineering, Case Western Reserve University, 1090 0 Euclid Ave., Cleveland, Ohio 44106

Gary L. Doll

Advanced Materials R & D, The Timken Company, 1835 Dueber Ave., S. W., Canton, Ohio 44706

ABSTRACT

The objective of this research was to develop and characterize hydrogenated silicon-containing diamond-like

carbon (Si-DLC) tribological coatings. A series of Si-DLC films was deposited by reactive sputtering in an

unbalanced magnetron sputtering system. Depositions were conducted with different tetramethyl silane flowrates

(TMS), thus varying the amounts of silicon and hydrogen in the films. All films were deposited onto an interlayer

of sputtered titanium to improve substrate adhesion and enhance mechanical properties. The goal of this study was

to understand the correlation between TMS flowrate, the bonding environment in the Si-DLC phase, and mechanical

properties. Hardness and elastic modulus were measured via microindentation, and internal compressive stress was

determined from surface mapping microscope radius of curvature measurements. Film chemistry was investigated

with x-ray photoelectron spectroscopy (XPS), Raman spectroscopy, and attenuated total reflection Fourier transform

infrared spectroscopy (ATR-FTIR). XPS and ATR-FTIR indicated that silicon carbide bonds are present in all of

the films deposited using TMS. Raman spectroscopy indicated that silicon incorporation may have reduced the

graphite cluster size. Conventional transmission electron microscopy revealed that the Si-DLC phase is amorphous,

and TiC exists at the Si-DLC/Ti phase boundary. It appears that after the silicon to carbon ratio in the films exceeds

0.04, the film matrix transforms from diamond-like/graphitic to that of a carbon-rich hydrogenated amorphous

silicon carbide containing small graphitic clusters.

Keywords: DLC, silicon, PVD, XPS, Raman

INTRODUCTION

Diamond-like carbon (DLC) coatings are known for being chemically inert and having low friction coefficients

and high hardness (ref. 1). A flexible, low friction, wear resistant DLC coating is desirable for use in hightemperature, high-pressure tribological applications such as bearings. It has been reported (refs. 2,3) that

incorporating silicon into DLC lowers internal compressive stress, reduces humidity sensitivity on friction

coefficients, and improves the high temperature utility by reducing thermal degradation. In this study, a series of SiDLC films with a Ti interlayer was deposited and characterized. Mechanical property trends were observed, and the

film structure and chemistry results were applied to understand their basis and aid in future process development

efforts.

EXPERIMENTAL

Coating Deposition Process

The films for this study were deposited on steel coupons and silicon wafers using a closed-field unbalanced

magentron sputtering system. The system utilizes field coils to extend the effectiveness of the magnetron permanent

magnets and has a 3-axis planetary motion substrate-holding carousel. A reactive sputtering process with a mixed

argon and tetramethyl silane (TMS) feed gas was used for the depositions. TMS has been used elsewhere (refs. 3 to

5) as a feed gas to deposit both hard and polymer coatings. The volumetric flowrate ratio of TMS to argon was

adjusted to deposit films with varying silicon concentrations. Two titanium and two carbon targets were

simultaneously loaded in the chamber and sputtered in the appropriate layer sequence. Six coatings with varying

silicon contents are considered in this report. They were all deposited using the same carbon target magnetron

power, DC pulsed substrate bias level, argon flowrate, and field coil current. For all of these depostions, the system

pressure was in the range of 10-15 mtorr. The Si-DLC layer deposition time was 2 hours for all processes.

* Email: rde3@po.cwru.edu, Phone: 216-368-2648, Fax: 216-368-3016

NASA/CP--2001-210948 460 August 6, 2001
Mechanical Properties

A CSEMmicroindentation/microscratch testerwithaVickersindentertipwasusedtogenerateloaddisplacement curves.Hardness(H)andelasticindentationmodulus(FindE/(1-v2)) werecalculatedfromthese

curvesusingthemethodofOliverandPharr(ref.6).Apower-lawindentationareacorrelationtocompensate fortip

roundingwasdevelopedbasedonacalibrationstandard. Themaximumindentationdepthforthesetestswas

~350nm, roughly30%ofthefilmthickness. Thisisgreater thanthe10%nanoindentation heuristic, sotheresults

wereadjustedforsubstrateeffects(ref.6).Theadjustedvaluesareonlyapproximate, asanestimateofPoisson's

ratiowasutilized(v=0.3).Theratioofyieldstrengthswasestimatedbyperformingvariableoadingstudiesand

fittingthedatatoanexponential decay.Thetrendfromthesesubstrate-adjusted measurements wasconsistent with

themeasuredhardnesstrend,t_n d is reported without assuming a value ofv. The internal compressive stress of the

films was estimated by measuring the radius of curvature at several points on the coated Si wafer using a Phase Shift

Technology MicroXAM Surface Mapping Microscope. The radius of curvature data were inserted into the popular

Stoney equation to deduce the stress level in the film (ref. 7).

Structure/Chemistry Characterization

Cross-sectional film imaging was performed with a 200kV Philips CM20 Scanning Transmission Electron

Microscope in TEM mode. Sample preparation included ion polishing with a Gatan Precision Ion Polishing System

with low angle (4°), 4 keV argon ion beams. Film thicknesses were obtained using a Calowear tester, and were

verified with TEM. X-ray photoelectron spectroscopy (XPS) characterization was done with a PHI 5600 ESCA

system using an A1K_ 1486.6 eV monochromatic x-ray source, equipped with a PHI 04-303 sputter gun. Multiplex

scans were conducted until a peak/noise ratio of 1000 was obtained for the Cls peak. Physical Electronics Multipak

software was used for curve fitting. The integral method was used for baseline subtraction. Peaks were fit with 85-

100% Gaussian peaks having a FWHM 1-2 eV. The films behaved like conductors during XPS analysis, and the

binding energy scale was set using near surface Ti and O species as a reference. Raman spectroscopy work was

done with a Jobin Yvon Horiba LabRam system. A 20 mW HeNe laser with a wavelength of 632.82 nm coupled

with optical density filters and an 1800 g/mm holographic grating was used. The reported measurements were

obtained using a 2 mW laser. LabSpec software was used for polynomial baseline adjustment and Gaussian peak

fitting. Attenuated total reflectance Fourier transform infrared spectroscopy (ATR-FTIR) work was done with a

Bomem MB-series system and a zinc selenide prism with a 45 ° angle of incidence. The films were pressed against

the crystal to provide good optical contact.

RESULTS AND DISCUSSION

Thickness, Composition, Structure

Figure 1 shows film thickness and Si/C ratio as detected by XPS in each film plotted versus the TMS/argon

flow rate ratio, QrMs/Qar. The sputtered Ti interlayer in each film was ~0.3 _tm thick for every coating. The SiDLC phase deposition rate increased linearly with increasing TMS flow rate. XPS depth profiling and TEM

revealed that a ~200 A thick surface layer of TiC and TiO2 exists on the surface of the films. This is attributed to a

final target-cleaning step which occurred after the film deposition and before the sample was removed. To obtain

the Si-DLC compositions, a surface layer of approximately 500A was sputtered away until there was no detectable

titanium present.

0.25 2.0

_-t= 0.20 1.5 '_

•-= o.15

..,= ! 1.o

O.lO

tO 0.5 0.05

0.00 0.0

0.00 0.10 0.20 0.30 0.40 0.50

QTMs/Q_r

Figure 1: Film compositions and thicknesses vs. TMS/Ar flowrate ratio

NASA/CP--2001-210948 461 August6,2001
Crosssectional TEMmicrographs andselectedareadiffractionforthefilmwithSi/C 0.04areshowni

Figure2.FromtheTEMobservations, theSi-DLCphaseappearedtobeamorphous inallofthesilicon-containing

filmsinthisseries.Thetitaniuminterlayer ispolycrystalline, anditsringpatternisconsistent wihamixtureof

cubicTiCandhexagonal Ti.TheTilayerappearstoconsist ofcolumnar gains.TheSi-DLC/Ti interfaceissemidiffuseand~75nmthick.

500nm

iiiiiiiiiiiiiiiiiiiiiiiii 

'o°°°.°o°°°°°°*°*°°°°°°.'.'°'.'°'.'.'.'.'.'''.'.''. ,° .........................

,o°°° .......................

iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii

___:::::::::::

...... • • -. • • • • • • • • • • • • • • • • • --.

::::::::::::::::::::::::::::::::::::::::::

:::::::::::::::::::::::::::::::::::::::::::::::::::::::

;::'," :::::::::::::::::::::::::::::::::::::::::::::::

:::::::::::::::::::::::::::::::::::.::._

Figure 2: (a) Coating cross-section; (b) View of Si-DLC/Ti interface; (c) SAD of Ti interlayer

Mechanical Property Characterization

Mechanical properties influence the performance and life oftribological coatings. Figure 3 (a) shows the

hardness and compressive stress vs. composition. A peak appears to be located at a Si to C ratio of 0.04 (flowrate

ratio QTMs/QAr 0.1). A connection between hardness and residual stress is obvious for these films. Another study

(ref. 2) has reported a peak in hardness with respect to silicon content in Si-DLC deposited with silane and benzene

utilizing rf-PECVD. These authors postulated that silicon incorporation initially increases interlinks and hardness,

but after a saturation point of~5 at%, Si-Si and Si-H bonds form that actively reduce the hardness. In order to

determine the applicability of this explanation to the present research involving reactive sputter deposition, chemical

analysis was performed on these samples as discussed in the next section.

A high level of residual internal stress is undesirable for tribological coatings, as it can cause poor substrate

adhesion and can lead to delamination of the film from the substrate. Figure 3 (a) shows that the residual stress

initially increased as Si was addedto the films and then decreased as the concentration exceeded Si/C 0.04. This

may be caused by an initial increase in Csp3-Csp 3 and Csp3-Si bonding that tightens the network, followed by the

formation of C-H bonds at higher Si contents that relax the network. A previous report (ref. 3) observed stress

reduction in PECVD-deposited Si-DLC without a peak. In the present case, the peak occurred after the transition

from a PVD (Si/C 0, no TMS addition) to a reactive sputtering (Si/C 0.04, TMS added to the system) deposition

process. The indentation modulus data are shown in Figure 3 (b). The modulus decreased as silicon and hydrogen

content was increased due to TMS addition. A high elastic modulus is typically associated with a large number of

C-C sp 3 sites in DLC, similar to the diamond lattice (ref 1). Added silicon and hydrogen lower the modulus,

presumably lowering the overall Csp3-Csp 3 bonding character. A film with a low E/H ratio is more elastic than a

film with a higher ratio (ref. 8). The silicon containing films appear to be more elastic than the Si/C 0 film.

30-

25-

2o

15

i"-

lOi

-i- 5

5

4 18o3_ o. 160

_ 140

2_ w 120

1

100

0.05

0 0 80

0.00 0.10 0.15 0.20 0.00 0.05 0.10 0.15 0.20

Si/C At% Ratio in Film Si/C At% Ratio in Film

Figure 3: (a) Hardness and stress vs. composition; (b) Indentation modulus vs. composition

NASA/CP--2001-210948 462 August 6, 2001
FilmChemistryCharacterization

X-rayPhotoelectron Spectroscopy (XPS)

Anexamination ofXPSpeaklocationswasutilizedtoevaluatethetypeandrelativequantityofchemical bonds

inthefilms.TheClspeakforthesefilmswasnon-Gaussian withaFWHM>2eV,characteristic ofamorphous

carbon.TheClsfeaturecanbefitwithfourpeaksrepresenting fourcarbonbondingconfigurations (refs.5,9,10).

Since~500Aofthesurfaceforeachsamplewasremovedbyargonionsputteringpriortospectrumcollection, there

maybeamoderatesp3tosp2conversion andexaggeration ofthegraphiticpeak(ref.11).Forthisreason,

quantitativejudgmentsaboutthebonddensityorsp2/sp3ratioarenotmadeinthisstudy.However, arelative

quantitativeevaluationismadebasedonthetrendsamongthesamples. Table1liststhepeakpositionsandareasfor

Ols,Si2p, andClspeakcurvefits.Possiblebondassignments areconsistent withreferences 5,9,10,and12.

Table 1: XPS peak fitting results

Possible Bond

Peak Si/C 0.00 0.04 0.06 0.10 0.14 0.18 Assignment

Pos. 531.8 531.9 531.9 531.9 531.8 531.9 O1s 1 C=O

Area/%) 93.7% 100.0% 95.5% 100.0% 100.0% 100.0%

Pos. 100.6 100.5 100.5 100.5 100.6

1 Si-C

Si2p Area (%) 81.5% 91.4% 92.6% 93.6% 95.3%

2 Pos. 102.4 102.4 102.6 102.4 102.6 HxCyO_iw

Area (%) 18.5% 8.6% 7.4% 6.4% 4.7%

Pos. 281.8 282.0 282.1 282.4 282.6 282.7 1 Si-C

Area (%) 0.9% 1.2% 1.5% 2.9% 5.5% 5.2%

Pos. 284.1 284.1 284.1 284.1 284.0 284.1

2 C=C

Cls Area (%) 76.1% 79.5% 84.3% 84.9% 81.3% 84.8%

Pos. 285.9 285.7 285.8 285.9 285.5 285.8 3 C-C

Area (%) 17.1% 14.9% 11.1% 9.6% 10.8% 8.2%

Pos. 288.4 288.2 288.3 288.3 287.7 288.1

4 C-O, C=O

Area (%) 5.9% 4.4% 3.1% 2.6% 2.4% 1.8%

Oxygen was detected in every film at 2-4 at%. The Ols peak for every film was located at the same binding

energy of 531.86 + 0.05 eV. This band is attributed to C O bonding (ref. 13). For the films with Si/C 0 and 0.06,

a second peak at higher binding energies was found and attributed to HxCyOz species (ref. 12). There were no Ols

peaks for SiO2 at 533.0 eV or detectable Ti, indicating the films do not contain significant pinholes that would

expose Si or Ti surfaces.

The Si2p features were fit with two peaks, one at 100.5 eV and the other at 102.5 eV. The silicon bands were

attributed to the Si-DLC phase, as there was no band at 99 eV to indicate the presence of elemental silicon. As

expected, no silicon peak was observed for the sample grown without the addition of TMS. The Si2p peak at 100.5

eV could correspond to a variety of HxCySiz compounds, including SiCc_ (ref. 12). This is believed to be due to

regions of hydrogenated amorphous silicon carbide. The Si2p peak at 102.5 eV could represent a variety of

compounds with the general formula CxHyOzSiq in the energy range of 102.4-102.7 eV (ref. 12). The 102.5 eV peak

decreased and the 100.5 eV peak became stronger with increasing Si content.

The binding energy of the first Cls fitted peak was ~282 eV. As the Si content increased, the binding energy and

peak area also increased. This peak was attributed to Si-C, as the binding energy for SiC_ is 282.5-283.4 eV (ref.

12). The second Cls band was located at 284.1 eV for all of the samples. This binding energy is consistent with

graphite powder (ref. 14) and is attributed to graphite-like islands in the films caused by sputtering of the graphite

targets. The C C peak area increased rapidly from Si/C 0 to Si/C 0.06 and then leveled off. The third Cls band

at 285.8 eV was also constant in position among the samples, but the peak area steadily decreased with increasing

silicon content. This band was consistent with diamond (285.5 eV) (ref. 14), or sp 3 hybridized carbon bonding

(285.8 eV) (ref. 11). There appeared to be less Csp3-Csp 3 bonding with increasing silicon and hydrogen content. SiC bonds increased with increasing Si, so it couldbe concluded that, as the Si content increased, the majority of

tetrahedral bonding involved Si, C, and H rather than C-C. The last Cls band at approximately 288 eV is popularly

assigned to C-O or C O bonding (refs. 5, 9, 10). This carbon-oxygen bonding component appeared to decrease with

increasing silicon content. This observation is consistent with a decrease in CxHyOzSiw at 102.5 eV in the Si2p band

as Si in the film increases.

Raman Spectroscopy

The films contained a broad Raman band between 1100 and 1500 cm -1 region. This band is generally believed

to be composed of G and D bands centered around 1575 cm -1 and 1355 cm -1, respectively. The G band is due to the

NASA/CP--2001-210948 463 August 6, 2001
RamanactiveE2gmodeofgraphiteorC Csp2stretchvibrations. TheD("disordered") bandisattributedtoahigh

phonondensityofstatesindisorderedmicrocrystalline graphite(ref.15).Althoughtephysical validityoffitting

andinterpretingDLCspectrawithDandGbandsioftendebatedintheliterature(ref.1),relativetrendsare

generallyacceptedtoindicatehegraphiticvs.diamond-like natureofthefilms.Thus,inthisworkpeakfittingwas

usedtoindicatesuchtrends.Thebroadpeaksat1575crff1andG-andD-bandpositionsforthefilmsdecreased

linearlywithincreasingsiliconcontent.ThishasbeenobservedlsewhereforDLCserieswithincreasinghydrogen

andsiliconcontent (ref.2,15).Figure4showsthatheIJI_bandintensityratiodecreases andflattensoutwith

increasingsiliconcontent.ThefilmwithSi/C 0hadthelargestID/I_ratioandsmallest Glinewidth.Itcould

containlargergraphiticclustersthantheSicontainingfilmsthathavesmallerID/I_ratiosandGlinewidths,

accordingtotheanalysisofreference15.Likewise, reference15reportedthatheDbandarisesasaresultofCsp3

sitesdisturbinggraphiticclusters. TheDbandintensitymaybereducedintheSicontainingfilmsbecauseilicon

andhydrogen, ratherthanCsp3sites,disturbedthegraphiticclusters.

...o

13

m

o"

,m

co

_=

1.8

1.6

1.4

1.2

1.0

0.8

0.6

0.4

0.2

0.0

0.00 0.05 0.10 0.15 0.20

Si/C At% Ratio in Film

Figure 4: II>/IG vs. composition

Attenuated Total Refection Fourier Transform Infrared Spectroscopy (ATR-FTIR)

ATR-FTIR was used to detect specific bonding types in the films, especially those involving hydrogen. Figure

5 (a) shows what appear to be interference effects (*) in the low Si films (Si/C 0 & 0.04). With the exception of

this phenomena, there were no distinguishable absorption bands. The films with Si/C 0.06 0.18, did display

absorption features as shown in Figure 5 (b). These same absorption features were also observed in external

reflection FTIR measurements. Popular peak definitions are as follows: Si-C stretch at 780 cm-1; C C stretch at

1600 cm-1; Si-H at 2160 cm-1; and C-H stretch at 2800-3000 cm -1 (refs. 16, 17). The band at 1000 crff 1 could be a

C-H rock/wag (ref. 9), a Si-CH2 and Si-CH3 deformation (ref. 16), or a Si-CH2-Si rock/wag (ref. 17). The Si-C bond

peak intensity increased with increasing silicon content. The peak at 1000 cm -1 increased slightly with increasing

silicon content. If this peak resulted from a Si-(CH2)-Si configuration, this evidence supports the claim that the film

matrix became more like hydrogenated amorphous silicon carbide with increasing Si content. The Si-H and C-H

peaks were also more distinct with increasing silicon content.

:: Si/C = 0.04 ""'-v

.oi * ' 'i /'i,

.............. 77 ....... 7,7 ........ 77 ....... 7,7 ....

_2 rS'i : ....

, b- "" ->" , ""

. ..

b_ '. •

........................................... .'-....

Figure 5: (a) ATR-FTIR spectra for Si/C = 0, 0.04; (b) Expanded view of Si/C = 0.06, 0.10, 0.18

NASA/CP--2001-210948 464 August 6, 2001
CONCLUSIONS

A series of Si-DLC coatings with varying silicon contents and a Ti-interlayer were deposited. The mechanical

propertie s were measured and interpreted with respect to structure and chemistry results for the Si-DLC phase.

From TEM analysis, the Si-DLC phase did not appear to contain long range order. XPS indicated that Csp3-Csp 3

bonding was relatively more prevalent in the film with no Si. Raman results indicated that this film also contained

the largest graphitic clusters. The observed peaks in hardness and stress when Si/C 0.04 are likely due to a

maximum in the tetrahedral bonding in the matrix, including Csp3-Csp 3 and Csp3-Si bonds. This sample contained

the lowest amount of hydrogen as measured by ATR-FTIR, and apparently has the appropriate amount of Si to

interact with the sputtered graphitic material to produce the hardest and most tightly bound matrix. As the silicon to

carbon ratio in the films increased beyond 0.04, the film matrix apparently transformed from diamond-like/graphitic

to that of a carbon-rich hydrogenated amorphous silicon carbide containing small graphitic clusters. This was

indicated by a reduction in hardness, stress, modulus, Csp3-Csp 3 bonds by XPS, D- and G-band positions by Raman;

and an increase in Si-C, C-H, and Si-H bonds by XPS and ATR-FTIR. Ongoing research involves a high resolution

and analytical TEM evaluation of the local bonding environment in the Si-DLC phase and the Si-DLC/Ti interfacial

region.

ACKNOWLEDGMENTS

We would like to thank The Timken Company for providing financial support for this project. Likewise we

thank Dr. Wayne Jennings and John Sears from the Center for Surface Analysis of Materials at CWRU; Dr. David

Cogdell and Doug Clouse for stress measurements; Martin Galehouse for mechanical property assistance; Brian

Osborn and Alexi Lazarides for assistance in coating laboratory; and Prof. Edward Evans (Dept. of Chemical

Engineering, The University of Akron) for many useful discussions.

REFERENCES

1. Lifshitz, Y.: Diamond-like carbon-present status. Dia. and Rel. Mats., vol. 8, 1999, 1659-1676.

2. Lee, K.R., et al.: Structural dependence of mechanical properties of Si incorporated diamond-like carbon

films deposited by RF PACVD. Thin Solid Films, vol. 308-309, 1997, 263-267.

3. Yi, J.W.; Lee, Y.H.; Farouk, B.: Effects of gas composition and r.f. power on properties ofa-C:H/SiC:H

composite films grown by PECVD. Thin Solid Films, vol. 326, 1998, 154-159.

4. Van Ooij, W.J.; Surman, D.; Yasuda, H.K.: Plasma-polymerized coatings oftrimethylsilane deposited on

cold-rolled steel substrates, Part 2. Prog. in Organic Coatings, vol. 25, 1995, 319-337.

5. Zhao, J.F., et al.: The effects of Si incorporation on the microstucture and nanomechanical properties of

DLC thin films. J. Phys.: Condens. Matter, vol. 12, 2000, 9201-9213.

6. Pharr, G.M.; Oliver, W.C.: Measurement of Thin Film Mechanical Properties Using Nanoindentation.

MRS Bulletin, July 1992, 28-33.

7. Hoffman, R.W.: Mechanical properties of non-metallic thin films. Phys. of nonmetallic thin fills, (NATO

Advanced Study Institutes Series: Series B, Physics), Dupuy, C. H. S., Ed., Plenum Press, 1976, 273-353.

8. Michler, T., et al.: Mechanical properties of DLC films prepared by bipolar pulsed DC PACVD. Dia. and

Rel. Mats., vol.7, 1998, 1333-1337.

9. Choi, W.K., et al.: Infrared and x-ray photoelectron spectroscopy studies of as-prepared and furnace

annealed rf sputtered amorphous SiC films. Journ. of App. Phys., vol. 83, no. 9, 1998, 4968-4973.

10. Ramm, M., et al.: Studies of amorphous carbon using x-ray photoelectron spectroscopy, near-edge x-ray

absorption fine structure and Raman spectroscopy. Thin Solid Films, vol. 354, 1999, 106-110.

11. Jackson, S.T.; Nuzzo, R.G.: Determining hybridization differences for amorphous carbon from the XPS

C(ls) envelope. App. Surf. Sci., vol. 90, 1995, 195-203.

12. Wagner, C.D., ed.: NIST x-ray photoelectron spectroscopy database. Distributed by: Standard Reference

Data, National Institute of Standards and Technology, Gaithersburg, MD 20899, 1989.

13. Barr, T.L.; Yin, M.: Concerted x-ray photoelectron spectroscopy study of the character of select

carbonaceous materials. J. Vac. Sci. Tecllol. A, vol. 10, no. 4, 1992, 2788-2795.

14. Taki, Y., et al.: XPS structural characterization ofa-C:H. Thin Solid Films, vol. 316, 1998, 45-50.

15. Schwan, J., et al.: Raman spectroscopy on a-C films. Journ. of App. Phys., vol. 80, no. 1, 1996, 440-447.

16. Varma, A., et al.: Tribological behavior of Si-DLC coatings. Surf. Engin., vol. 15, no. 4, 1999, 301-306.

17. Kim, D.S.; Lee, Y.H.: Roon_temperature deposition ofa-SiC:H thin films by ion-assisted plasma-enhanced

CVD. Thin Solid Films, vol. 283, 1996, 109-118.

NASA/CP--2001-210948 465 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

"METALORGANIC CHEMICAL VAPOR DEPOSITION OF RUTHENIUM-DOPED

DIAMONDLIKE CARBON FILMS"

M.K. Sunkara, M. Ueno

Department of Chemical Engineering, University of Louisville, Louisville, KY 40292

G. Lian and E.C. Dickey

Department of Chemical and Materials Engineering, University of Kentucky, Lexington, KY 40506

ABSTRACT

We investigated metalorganic precursor deposition using a Microwave Electron Cyclotron Resonance (ECR)

plasma for depositing metal-doped diamondlike carbon films. Specifically, the deposition of ruthenium doped

diamondlike carbon films was investigated using the decomposition of a novel ruthenium precursor, Bis-

(ethylcyclopentadienyl)-ruthenium (Ru(C5H4C2H5)2). The ruthenium precursor was introduced close to the

substrate stage. The substrate was independently biased using an applied RF power. Films were characterized using

Fourier Transform Infrared Spectroscopy (FTIR), Transmission Electron Microscopy (TEM) and Four Point Probe.

The conductivity of the films deposited using ruthenium precursor showed strong dependency on the deposition

parameters such as pressure. Ruthenium doped sample showed the presence of diamond crystallites with an average

size of -3 nm while un-doped diamondlike carbon sample showed the presence of diamond crystallites with an

average size of 11 nm. TEM results showed that ruthenium was atomically dispersed within the amorphous carbon

network in the films.

Keywords: MOCVD, Ruthenium, Me DLC, Nanocrystallites, ECR, HRTEM

INTRODU" TION

Diamondlike carbon (DLC) films are attractive for several commercial applications because of the range

over which their mechanical, optical, thermal and electrical properties could be tuned. Many techniques including

Radio Frequency (RF) discharges, DC Magnetrons, Ion Beams, Mass Selected Beams, Cathodic Arc and ECR

Plasma discharges have been demonstrated for deposition of hard amorphous carbon films. Please see the review

articles in references 1-2 and references within. Processing method, gas phase chemistry and the ion impact energies

play an important role in determining the diamondlike nature (sp 3 content) of the resulting amorphous carbon films.

Mechanical properties such as hardness and wear resistance depend upon the structure and the resulting stresses

within the films during deposition. Recently, it was shown that the use and choice of interlayer material depending

upon the softness of underlying substrate material could help reduce the stresses within resulting the amorphous

carbon films (ref. 3).

The structure of amorphous carbon is more open (density ranges from 1.5 - 2.5 gm/c.c, than the structure

of crystalline diamond (density of 3.5 gm/c.c). Because of its open nature, these amorphous carbon structures can

accept a wider range of dopants when compared to crystalline diamond. This presents an opportunity for

experimenting with a variety of dopant elements for imparting/improving new properties to the deposited

diamondlike carbon films and in synthesizing an entirely new class of materials based on sp3-bonded carbon

networks.

Silicon doping was shown to reduce stresses within the amorphous carbon films (refs. 4 to 5). Amorphous

carbon films doped with boron show p-type electronic behavior (ref. 6). There were reports of iodine doping into

RF plasma-deposited DLC films in which iodine was introduced as vapor from heated solid iodide sources (ref. 7).

Fluorine incorporated amorphous carbon films showed low surface energy, low moisture uptake and higher

NASA/CP--2001-210948 466 August 6, 2001
corrosionprotectionability(ref.8).NitrogendopingintoDLCfilmshasattractedattentionduetothepossibilityof

producingsuperhardcarbonnitridephase, C3N4. Nitrogen doping of amorphous carbon films with high sp 3 content

and low hydrogen content has been shown to affect electronic and electron emission properties (refs. 9 to 13).

Metal doping into amorphous carbon films is interesting for the following reasons: fundamental curiosity

on alloying behavior of carbon with metal; ability to manipulate electrical property through percolation of metallic

or metallic carbide inclusions inside amorphous carbon films; and ability to reduce stresses and improve adhesion

for improving mechanical properties (ref. 14). Co-sputtering and co-evaporation of metal were the primary methods

used for metal incorporation to deposit metal containing amorphous carbon films (refs. 15-22). In the films using

co-sputtered or co-evaporated metal, the metal was incorporated into amorphous carbon films in the form of

nanoclusters. These nanoclusters were found to be either metallic or metallic carbide depending on whether the

metal is carbide forming or noble type (refs. 17-22). In this study, we examined whether we can utilize metalorganic

precursor decomposition for metal incorporation into amorphous carbon films and how it impacts the structure of

amorphous carbon network and metal. The electrochemical properties of metal containing amorphous carbon films

depend upon the structure of the films (ref. 23).

EXPERIMENTAL

The deposition experiments were conducted using Electron Cyclotron Resonance (ECR) CVD reactor as shown

in the schematic in Figure 1. Microwave ECR discharge was created in a quartz bell chamber using a cooled

electron-magnet (ASTeX CECRP4025), a 250 W microwave power supply (ASTeX AX2000) and a coaxial slug

tuner (ASTeX AX3040). A 3" copper substrate stage was placed in a double walled stainless steel chamber at

approximately 30 cm below the ECR plasma discharge. Substrate stage was equipped with RF bias and cooling. The

generated bias voltage of the substrate above the chamber with applied RF power was measured using a DC

voltmeter and a RF choke circuit. Backside thermocouple was used to monitor the substrate temperature. Ruthenium

precursor, Bis-(ethylcyclopentadienyl)-ruthenium (Ru(CsH4C2H5)2, was introduced using a flow of carrier gas

through a bubbler containing the liquid precursor at room temperature. This precursor was introduced in the reactor

close to the substrate. Methane was used for depositing un-doped diamond-like carbon films by introducing it near

plasma discharge. The pressure was varied from 0.015 torr to 0.035 torr with a flow rate of 5 sccm and a MW power

of 200W. With a forward RF power of approximately 50W, the substrate bias was measured to be at -70V. The

measured deposition rates of ruthenium doped films and un-doped films at 0.025 torr pressure were -55 and 25

Angstroms/min, respectively. The deposited films were characterized using Infrared (FTIR) spectroscopy for C-H

bonding and four-point probe for conductivity.

---------4_ MW Generator

Exhaust _ECR Source

i [ @Electro

_ Magnet Exhaust

CH4 °r H2---*---_ l-- I I

Purge Gas--_ X , ____v_P___ _E____ T

Chemicm_ @_ cret;_ler vumps

WaterBath _J

RF Generator and

t.______aMatching Network

Figure 1. Schematic of the MW ECR chemical vapor deposition reactor.

NASA/CP--2001-210948 467 August 6, 2001
PlanviewTEMspecimens werepreparedwithmechanical poishingfromSisubstrate, followedbyionmilled

withAr+ionsatliquidnitrogencooledstage.CrossectionTEMspecimens werepreparedasthefollowing:the

sampleswerecutintotwopieceswithasizeof2.5xlmmandthefilmsideswerebondedtogether, thenmechanical

polishedtoabout10-20gmandionmilledtoelectrontransparency ataliquidnitrogencooledstage.Conventional

transmissionelectronmicroscopy(CTEM)analysiswasperformedonJEOL2000FX. Highresolutionelectron

microscopy (HREM)wascarriedoutusingJEOLJEM2010Ffieldemissionelectronmicroscopy, operatingat200

kV,equippedwithanannular darkfielddetector, ascanningunit,apostcolumnimagefilter(GatanGIF200)andan

Oxfordenergydispersivex-ray(EDX)detector.

RESULTS AND DISCUSSION

The electrical resistivities of the films deposited using ruthenium precursor varied from 1.3 x 10-2ohm-cm to

22 ohm-cm as the system pressure increased from 0.015 tort to 0.035 tort. The deposition rates were same over the

above pressure range. The conductivities of the films deposited at lower pressures were higher as shown in Figure 2.

This is consistent with a reported study in which pure ruthenium films could be deposited using the same precursor

(ref. 24). The results in Figure 2 with conductivity show that the ruthenium incorporation into the film could be

tuned with deposition variable such as pressure. The FTIR spectra behavior is similar to that observed with films

deposited using methane as precursor, i.e., the peaks associated with sp 3 hybridized bonding were sharper at 0.035

tort. The main difference is that the films deposited using ruthenium precursor at 0.035 tort showed a peak around

3100 cm -1 probably with higher amounts of sp 2 hybridized bonding.

The presence and distribution of ruthenium in these films was confirmed using electron dispersive

spectroscopy and electron energy loss spectroscopy (ref. 25). Figure 2 shows lattice images of two different film

samples deposited using methane and ruthenium precursors at 0.025 tort pressure. Both films showed the presence of

nanocrystallites. The selected area diffraction patterns are shown in Figure 3. The lattice plane spacings are

determined from the SAD patterns and are summarized in Table 1. These results confirm the structures of these

nanocrystals in both samples to be of cubic diamond. The formation of diamond nanocrystallites in un-doped

diamondlike carbon films deposited using MW ECR plasmas with applied RF bias energies of 150 eV was reported

(refs. 26-28). The results in this paper indicate that diamond nanocrystallites could be formed within diamondlike

carbon networks with typical applied RF bias voltages of-70V.

<

".4,'-'.; ;:"T:,:. :.

10

?

8

1

\

,01

Pressure, mTorr

Wavenumbers, cm "1

(a) (b)

Figure 2: (a) FTIR spectra films deposited using ruthenium precursor at different pressures. (b) Resistivity of

films deposited using ruthenium precursor at different pressures.

NASA/CP--2001-210948 468 August 6, 2001
Toourknowledge, nostudieswerereportedearlierontheformationofdiamondcrystallitesinanytypeof

metalcontainingamorphous carbonfilms.In thisregard,thepresent observations ofdiamondnanocrystallites in

rutheniumdopedamorphous carbonfilmsarecompletelyunexpected. Inaddition, anoblemetallikerutheniumdid

notexhibitclusteringwithintheamorphous carbonnetworkasexpected. Theconductivity ofthesample(resistivity

of 1ohm-cm) examinedinthispaperwassimilartothefilmsdepositedusingco-sputtering ofmetalthatexhibited

metalliclusters. Anotherimportant observation is thatheaveragesizeof diamondcrystallitesinthesefilmsis

muchsmaller(- 3nm)comparedtothose(-11nm)foundinthefilmsobtainedwithmethaneprecursor. Thisresult

indicatesthatheco-chemical vpordepositionfrutheniumalongwithamorphous carbonihibiteddiamondcrystal

growthandenhancednucleationfnewdiamondcrystalswithintheamorphous carbonnetwork.AsshowniFigure

5,thesizedistributionofthediamondnanocrystallites inthesetwoamorphous carbonfilmsisnarrower comparedto

un-dopeddiamondlike carbonfilms.Theresultspresented inthispaperindicatehatonecanproducedifferent types

of metaldispersionswiththeamorphous carbonnetworksusingacompletechemical vapordepositiontechnique

comparedtoahybridtechniquecombiningPhysical VaporDeposition(PVD)suchassputteringandCVD.

" 33; _. =============================

_. , ff,..._..o. .................

__ _,'_. ')..i_.::i:i:::::i:i::::

.-

_i_I _'_,_,::. :°,_i:,4: ::::::::::::::::::::::::::::::::::::

•::N.% '.','.'e+',".'-'.'-'." '.'.-'.'.'.'.'- :::::::::

• oo_,' o_ "oO¢. '°% o%°_%°o°o', .-.' ... - . *.. • •

_2k:;_'Y_"_::f4 :':-_:<':':':':::':-:::::::::::::::::::::::::::

•:_,:,1,:_",;,',-_ ::::::::::::::::::::::::::::::::::::::

==============================================

_;:;?,°,,N' ::::::::::::::::::::::::::::::::::::::::::

• _ °,o,_ _o.,% o,,-.,o,.. *...-........ o

:,'.o_i'°°o,'i,°o'o;,2o_.:2_._.11_i_._i_12.1.1.211.

_,_,_::;_::_:,-;, _,.?.,:_,.,:-.:;::., ..:............,.............

(b)

Figure 3: High resolution transmission electron micrographs (HRTEM) of un-doped and ruthenium-doped

diamondlike carbon films. Images show lattice images of nanocrystals of diamond.

(a) (b)

Figure 4: Selected area diffraction patterns of un-doped and ruthenium doped samples. Un-doped film was

prepared using methane as gas phase precursor and ruthenium doped film was prepared using the ruthenium

precursor at pressure of 0.025 tort and a bias of -70V.

NASA/CP--2001-210948 469 August 6, 2001
TABLE1:Thediffractiondatafromtheselectedareadiffractionpatternsobtainedfromun-dopedand

rutheniumdopedfilms.

JCPDSdata

CubicDiamond(a=3.57A)

dhkl I (%1 d(A)

111 100 2.06

200 1.79

220 25 1.26

311 16 1.07

400 8 0.89

331 16 0.82

Graphite-2H(a=2.47,c=6.721

dhld I (%1 d(A)

002 100 3.38

100 2 2.14

101 6 2.04

102 <1 1.81

004 4 1.68

103 1 1.55

110 3 1.23

112 3 1.16

Experimental observation

Un-doped Ru-doped Intensity

DLC d(A)

2.4

DLC d(A)

2.38

order

5th

2.07 2.06 1st

1.79 1.79 3rd

1.27 1.27 2nd

1.09 1.08 4th

0.82 0.80 6th

14,

', f_-:-::_&;?;;gf-_i

12 i { ...... Pure DEC

e 10'_ "

il

Z 4i i

2

0

0 5 19 15 20

Grain size (rm)

Figure 5: The size distribution of diamond crystallites within ruthenium-doped and un-doped samples.

SUMMARY

A novel deposition scheme using metalorganic precursor decomposition is demonstrated for ruthenium doped

diamondlike carbon films. The results showed that the conductivity of the resulting films depended strongly upon the

operating pressure. The ruthenium incorporation in to amorphous carbon films increased the nanocrystallinity of

diamondlike carbon films by increasing the density and reducing the average size of diamond crystallites from 11 nm

to an average size of 3 nm. These results show a process dependence on the structure of the resulting metalNASA/CP--2001-210948 470 August 6, 2001
containingamorphous carbonfilms,i.e.,thestructureofmetalcontainingamorphous carbonfilmsdepositedusinga

completechemical vapordepositiontechniquecouldbedifferentthanthosedepositedusinghybridtechniques

involvingbothPVDandCVD.

ACKNOWLEDGEMENTS

Authorsgratefullyacknowledge financial support fromKentuckyNASAEPSCoRprogramandNSF-MRSECat

UniversityofKentucky.

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NASA/CP--2001-210948 471 August 6, 2001
20.Schiffmann, K.I.,Fryda,M.,Goerigik, G.,Lauer,R.,andBulak,RA.,"SizesandDistances ofMetalClustersin

Au-,Pt-,W-andFe-containing Diamond-likeCarbonHardCoatings: aComparative StudybySmallAngleXrayScattering, WideAngleX-rayDiffraction, Transmission ElectronMicroscopyandScanningTunnelling

Microscopy", Thin Solid Films 347, 60 (1999).

21. Fan, X., Dickey, E.C., Pennycook, S.J., and Sunkara, M.K., "Z-contrast Imaging and Electron Energy Loss

Spectroscopy Analysis of Chromium-Doped Diamond-Like Carbon Films", Appl. Phys. Lett. 75, 2740 (1999).

22. C. Donnet, "Recent Progress on the Tribology of Doped Diamondlike and Carbon Alloy Coatings: A Review",

Surf. Coat Technol. 100/101, 180 (1998).

23. Sunkara, M.K., Chandrasekaran, H., and Koduri, R, "Electrochemical Studies of Metal-Doped Diamondlike

Carbon Films", New Diam. and Frontier Carbon Technol. 9(6), 407 (2000).

24. Aoyama, T., and Eguchi, K., "Ruthenium Films Prepared by Liquid Source Chemical Vapor Deposition Using

Bis-(Cyclopentadienyl) Ruthenium", Japan J. Appl. Phys. 38, L 1134 (1999).

25. Lian, G., Dickey, E.C., Ueno, M., and Sunkara, M., "Diamond Nanocrystallites in Un-Doped and Ru-Doped

Diamond-like Carbon Films", Submitted (2001).

26. Silva, S.R.R, Xu, S., Tay, B.X., Tan, H.S., and Milne, W.I., "Nanocrystallites in Tetrahedral Amorphous Carbon

Films", Appl. Phys. Lett. 69, 491 (1996).

27. Zarrabian, M., Fourches-Coulon, N., Turban, G., Marhic, C., and Lancin, M., "Observation of Nanocrystalline

Diamond in Diamondlike Carbon Films Deposited at Room Temperature in Electron Cyclotron Resonance

Plasma", Appl. Phys. Lett. 70, 2535 (1997).

28. Heiman, A., Gouzman, A.I., Christiansen, S.H., Strunk, H.R, and Hoffman, A., "Nano-Diamond Films

Deposited by Direct Current Glow Discharge Assisted Chemical Vapor Deposition", Diam. and Relat. Mater. 9,

866 (2000).

NASA/CP--2001-210948 472 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

NICKEL-CONTAINING CARBON FILMS DEPOSITED USING THE SCREEN

GRID TECHNIQUE IN AN ELECTRON CYCLOTRON RESONANCE

CHEMICAL VAPOR DEPOSITION SYSTEM

a a ,a* a a • • b

Q. F. Huang, S. F. Yoon , Rush , Q. Zhang J. Ahn, E.J. Teo b, T. Oslpowlcz , and F. Watt b

a School of Electrical & Electronic Engineering, Nanyang Technological University

Nanyang Avenue, Singapore 639798, Rep. of Singapore

b Research Center for Nuclear Microscope, Physics Department, National University of Singapore

10 Kent Ridge Crescent, Singapore 119260, Rep. of Singapore

ABSTRACT

Metal-containing carbon (Me-C:H) films are presently attracting wide interest due to their potential applications

in mechanical, tribological, electronics and optical fields. In this study, two sets of nickel-containing carbon (NiC:H) films were deposited using an electron cyclotron resonance chemical vapor deposition (ECR-CVD) system in

conjunction with two biased screen grids. The Ni-C:H films were characterized using Raman scattering (RS) and

UV-Visible spectrophotometry. Their resistivity, optical bandgap and hardness were evaluated as a function of the

gas flow ratio (CHJAr). Rutherford back scattering (RBS) analysis showed that the atomic fraction of Ni

incorporated in the carbon film decreases drastically from 35% to 1.4%, following an increase in the CH4/Ar flow

ratio. The hardness also decreases and the film with a Ni atomic fraction of 12% has a hardness of around 16GPa.

The film resistivity was found to increase by 11 orders of magnitude following the decrease in the Ni atomic

fraction. Vacuum annealing at 200°C for 1 hour immediately after deposition results in an increase in resistivity.

The threshold field for field emission decreases with decreasing Ni fraction in the film and a value of around 4V/gm

can be obtained for the film with the lowest Ni atomic fraction of 1.4%. The results showed that our ECR-based

screen grid deposition technique for Me-C:H is highly effective and flexible, with good control over the amount of

metal incorporated.

Keywords: Diamond-like carbon, electron cyclotron resonance, resistivity, field emission

1 INTRODUCTION

Metal-containing carbon (Me-C:H) films are being widely studied and have been shown to exhibit small friction

values, good conductivity, low abrasive wear rates and enhanced adhesion to metal substrates compared to pure

DLC films (refs. 1 to 2). The metal in these films can be introduced either by post-growth implantation (ref.3), or by

incorporation during film growth using chemical vapor deposition (CVD) technique (refs.4 to 5) or a combination of

physical vapor deposition (PVD, sputtering of metals) and CVD (refs.6 to 7). The technique used in our study is

based on an electron cyclotron resonance chemical vapor deposition (ECR-CVD) system in conjunction with two

metal grids situated within the deposition chamber. The deposition process is a combination of PVD and CVD, and

can be considered as a plasma reactive sputtering in which the sputtered metal atoms by the energetic Ar + are

incorporated into the growing a-C:H films. Compare to other techniques (refs. 1 to 2), our deposition technique is

flexible since the grids, instead of one of the electrodes, are used as the source of the metal. This implies that it can

be easily extended for the deposition of other Me-C:H films. In addition, such a system allows an independent

control of the plasma ionization and the metal sputtering rates, through the microwave power and the grid bias

respectively. This is in contrast to the commonly used glow discharge sputtering technique where the two important

growth parameters are correlated through the rf power (refs. 1 to 2).

In our previous paper (ref.8), we have studied in detailed W- and Mo-C:H films deposited using the above

technique. In this article, results for Ni-C:H films deposited at different CHJAr ratios are presented. A set of

concurrently grown films subject to post-deposition annealing in vacuum was also studied to investigate their

thermal stability.

NASA/CP--2001-210948 473 August 6, 2001
2EXPERIMENT

Thedetailsofthesetupusedinthisstudycanbefoundelsewhere(ref.9).WhenagasmixtureofArandCH4is

introducedintotheexcitationchamberthroughthegasinlet,aplasmais formedueto theexcitationbythe

microwavepowerandthedivergent magneticfeld.AnelectricfieldcreatedbytheDCbiasappliedatthescreen

gridshelpstodirecttheplasmaintothedepositionchamber throughtheNi screengrids.Intheprocessofpassing

throughthescreengrids,sputteringofNi fromthegridsbyAr+takesplace.ThesputteredNi isthenincorporated

intothegrowingfilms,whicharedepositedonsubstrates of(100)-oriented singlecrystalsiliconandcorninglass.

Forallthedepositions ithisstudy,thetotalchamber pressurewasmaintainedat8mTorr,andamicrowavepower

ofa400Wwasapplied. ThetwoscreengridswereshortedtoaDCvoltageof-330V,andtherf powerappliedto

thesamplewasmaintainedat60W,resultinginincreasingDCbiasrangingfrom-100Vto-135VovertheCH4/Ar

ratioof0.4to1.25.Priortodeposition, thechamber wasevacuated tobelow5x 10-_Torrusingaturbomolecular

pump.Nodeliberatehatingwasappliedtothesamplesduringthedepositionprocess, andthesampletmperature

wasestimatedtobe50to60°Cduringrowth.Forsamplesthatunderwent annealing, theywereheatedto200°Cfor

anhourinavacuumenvironment of5x 10-_Torrimmediately afterdeposition.

3 RESULTS AND DISCUSSIONS

Figure 1 shows the Ni/C fraction deduced from RBS for the Ni-C:H films at different CH4/Ar ratios. It is found

that the Ni/C fraction decreases with increasing CH4/Ar ratio, attributed to a decrease in the sputtering rate of the Ni

grids at lower Ar partial pressures. The samples at CHJAr = 0.75, 1 and 1.25 reveal a two-layer structure with

different fractions of Ni incorporated in each layer. For the two-layer structure samples, layer 1 refers to the top

layer whereas layer 2 refers to the layer next to the substrate. The average Ni, shown in the inset, and Ni/C fractions

for these samples were computed taking into account the thickness of the individual layers. At high CHJAr ratio,

the sputtering rate at the grids will be lower, and consequently the grids are increasingly coated by carbon. This in

turn will further reduce the sputtering rate, leading to the formation of the layered structure observed, with the top

layers (layer 1) having a smaller fraction of Ni incorporated. Simultaneously, there is an increase in the growth rate

for the films due to the reduced number of energetic Ni +, which can have an etching effect on the films. On the other

hand, at lower CHJAr ratios, the sputtering rates are higher and the grids are not likely to be coated. This will lead

to a uniform fraction of Ni across the depth of the films, and also lower growth rates.

0.6-

0.4-

0.2-

0.08 -

0.04 -

0.00

014 016 018 110 112

CH4/Ar

Fig.1. The Ni/C fractions of the samples deduced from RBS as a function of the CHa/Ar ratio. Two layers

fitting with different fractions are required for samples deposited with CHa/Ar = 0.75, 1.00, and 1.25. The

average Ni, shown in the inset, and Ni/C fractions for these samples were computed taking into account the

thickness of the individual layers.

Fig.2 shows the Raman spectra of the Ni-C:H films at different CH4/Ar ratios, measured using a SPEX1400

Raman system with the 514.5 nm line of an Ar + laser as the excitation source. Two characteristic peaks of a-C:H, the

G and D peaks, can be clearly seen in the Raman spectra. The intensity of the two peaks increases with the CHJAr

ratio, revealing the characteristics of glassy carbon (refs.10). The D peak around 1350 cm -1 is much more

pronounced compared with that of Mo-C:H films (ref. 11). This can be attributed to the difference in the impinging

ion energy during film growth, which ranges from -100V to -135V for the Ni-C:H films, compared to the lower bias

of-90V for the Mo-C:H films. A higher DC bias at substrates will result in more energetic ions impinging onto the

NASA/CP--2001-210948 474 August 6, 2001
substrates, whichamorphizes thefilmsandpromotesthegrowthof sp2bondedcarbon.Indeed,Mo-C:Hfilms

depositedathigherDCbiasabove-105Valsoexhibitedpronounced Dpeaks(ref.12).

A

D Peak J _\ G Peak

5

_ 210/20

i i i i i i

1200 1300 1400 1500 1600 1700 1800

Raman Shift (cm 1)

Fig.2. The Raman spectra of the Ni-C:H films deposited at different CH4/Ar ratios.

Hardness measurement, using nanoindentation technique by Nanoindenter II from Nano Instruments, Inc., was

performed for the films deposited on silicon substrates. The results shown in Fig. 3 are the average values measured

at six different locations. It is found that the film hardness decreases from -16 GPa to -8 GPa following an increase

in the CH4/Ar ratio. Compared to metal-free a-C:H films grown under the same process conditions with hardness of

around 10 GPa, the results clearly suggest that Ni incorporation increases the film hardness. In our previous study of

Mo-C:H films, a two-phase-model consisting of carbide precipitates and a-C:H matrix phases was employed to

explain the decrease in their hardness with the CH4/Ar ratio (ref.9). At lower CH4/Ar ratio, the Mo fraction, and

hence the carbide fraction was higher and predominantly accounted for the higher hardness. At increasing CHJAr

ratio with lower Mo fractions incorporated, the hardness was mainly determined by the a-C:H matrix and was

consequently lower. The increase in hardness for Me-C:H at larger carbide fractions was also observed by Klages et

al. (ref.1) in their W-, Ta- and Ti-C:H films. In contrast, Ni-C:H films prepared by incorporating Ni through

evaporation into the a-C:H grown on the rf-powered electrode in a plasma-CVD process from benzene or butane

was found to exhibit lower hardness, attributed to the inability of Ni to form hard carbidic precipitates (ref. 1). For

our Ni-C:H films, in the absence of XRD and XPS results to support the presence of Ni carbides, it is unclear at

present whether the decrease in hardness is related to the decrease in the carbide fraction.

16-

14-

12.

10-

8.

6

i i i J J

0.4 0.6 0.8 1.0 1.2

CH4/Ar

Fig.3. The film hardness as a function of the CH4/Ar ratio.

Fig.4 plots the resistivity against CHJAr ratio for both the as-deposited and the annealed Ni-C:H films on the

glass substrates. It is found that the resistivity for both sets of films increases sharply by eleven orders of magnitude

with increasing CHJAr ratio from 0.4 to 0.75, beyond which saturations occur. The sharp increase in resistivity with

NASA/CP--2001-210948 475 August 6, 2001
theCHdAratiowasalsoobservedinourW-andMo-C:Hfilms(ref.8),andwassimilarlyreportedforRu-,Ta-,

Co-andAu-C:Hfilmsatlowmetalfractions(refs.1,2,13). Thisindicatesthatheincorporated metal,evenatlow

atomicfractions, playsaveryimportant roleindeterminingthefilmresistivity. AtlowerCHdArratiotherewillbe

moreNiatomsbeingincorporated inthefilms,leadingtothelowerresistivityobserved. It isnotedthathefilmwith

CHdAr=0.4hasaresistivityaslowasabout10 -4 f2cm, comparable to that of metals, whereas the film at 0.75

exhibits insulating behavior with a resistivity of nearly 108 f2cm. After annealing, a slight increase in the film

resistivity was observed, which can be attributed to a denser structure in the annealed films. It is thought that

graphitization (especially in the presence of Ni (ref. 14)) and cross-linkage are two competitive processes which coexist in the films during annealing. A temperature of 200°C is not high enough to result in graphitization (ref. 14),

but can lead to a slight cross-linkage in their structures. This can be confirmed from our Raman results, where the

relative intensity of the D peak to the G peak (ID/IG), which is correlated to the sp2/sp 3 ratio, decreases for the

annealed Ni-C:H films (ref.11). The denser structure of the annealed films can also be caused by a structural

rearrangement. It is known that energetic bombardment of ions (mainly the Ni ions in our case) on the growing film

can result in large compressive stress (ref. 15). Upon one hour of annealing at 200°C there will be relaxation in the

films, leading to a relief of their compression stress. Thus the annealed films will be energetically more stable and

structurally denser. It is worth noting that our Ni-C:H films were annealed immediately after deposition, and

consequently post-oxidation will be much less compared to the as-deposited films. Therefore, the free radicals and

ions on or very near to the surface can cross-link to form a tighter carbon network structure. This may be also

responsible for the observed increase in their resistivity.

108

106

104

102

100

!

102_

!

104 -

./ l:ng

• Amlealing

i i i i i

0.4 0.6 0.8 1.0 1.2

CH4/Ar

Fig.4. The film resistivity as a function of the CH4/Ar ratio.

The optical absorption in the Ni-C:H films are obtained using samples deposited on glass substrates, through

their transmission and reflection spectra measured using a Perkin-Elmer Lambda 16 dual beam spectrophotometer.

• _/_ --i-- No annealing

4-- Annealing

• i , i , i , i , i , i , i , i , i

0.4 0.5 0.6 0.7 0.8 0.9 1.0 1.1 1.2 1.3

CH4/Ar

Fig.5. The Tauc gap (Eg) as a function of the CH4/Ar ratio.

Fig.5 shows the Tauc gap (Eg) for both the as-deposited and the annealed films as a function of the CHdAr ratio.

Except for CHdAr -- 1.0 where the Eg exhibits a sharp drop, Eg is found to increase following an increase in the

NASA/CP--2001-210948 476 August6,2001
CHJArratio,consistent withtheresultsobtainedforW- andMo-C:Hfilms(ref.8).Thissuggeststhatoptical

absorptionincreases withtheNi fraction. ThesharpdropinEgatCHJAr--1.0isthoughttoresultfromabigerror

in thethickness. A slightincreaseinEguponannealingsuggestsa slightdecreasein thesp2/sp3fraction,and

supportstheabsenceofgraphitization, butaslightcross-linkage, thusadenserstructurefortheannealedfilms,

consistent withtheresistivityresults.

ThefieldemissionbehaviorftheNi-C:Hfilmswasalsoinvestigated inourstudy.Itisfoundthathethreshold

fieldforemissiondecreases withdecreasing Niincorporationnthefilm.Fig.6showsthefieldemissionbehaviorf

theNi-C:HfilmsatthelowestNi atomicfractionof 1.4%.A thresholdfieldof around4V/btmatanemission

currentof1nAcanbeseenfromFig.6(a).Thelineardependence betweenln(I/E2)and1/E, as shown in Fig.6 (b),

indicates that the field emission in our Ni-C:H films exhibits the Fowler-Nordheim (FN) emission mechanism

(ref.16). It is noted that a threshold field of around 6 V/btm was also obtained for Ni-containing ta-C film with Ni

fraction of 1.91%, deposited using filtered cathodic vacuum arc technique. A further study on the field emission

behavior of the Me-C:H films is in progress.

1E5

1E6

1E7

1E8

1E 9

1E 10

1E 11

,44ff'""" '

i J i i i J

0 5 10 15 20 25

E (V/_m)

18.

20.

%

22.

24

0.00

i i i i i

0.05 0.10 0.15 0.20 0.25 0.30

1/E

(a) (b)

Fig.6. The field emission current-field characteristics (a) and Fowler-Nordheim plot (b) of the Ni-C:H films

with a Ni atomic fraction of 1.4 %.

4 CONCLUSIONS

Two sets of Ni-C:H films were deposited using a technique based on an electron cyclotron resonance chemical

vapor deposition (ECR-CVD) system in conjunction with two metal grids situated within the deposition chamber.

The films deposited at higher CHJAr ratios show a two-layer structure, and the average Ni atomic fractions in the

films decrease drastically from 35% to 1.4% following an increase in the CHJAr flow ratio. The hardness also

decreases with increasing CHJAr flow ratio, and the film with a Ni atomic fraction of 12% has a hardness of around

16GPa. The resistivity increases by 11 orders of magnitude following the decrease in the Ni atomic fraction. The

films annealed at 200°C exhibit a slight increase both in their resistivity and optical gap. The threshold field for field

emission decreases with decreasing Ni incorporation in the film and a value of around 4V/gm can be obtained for

the film with the lowest Ni atomic fraction of 1.4%.

NASA/CP--2001-210948 477 August 6, 2001
References

1. C. P. Klages and R. Memming: Microstructure and physical properties of metal-containing hydrogenated carbon

films. Material Science Forum 52 & 53, (1989), 609.

2. H. Dimigen, H. Hubsch and R. Memming: Tribological and electrical properties of metal-containing

hydrogenated carbon films. Appl. Phys. Lett. 50, (1987), 1056.

3. O.Amir and R.Kalish: Doping of amorphous-hydrogenated carbon films by ion implantation. Diamond and Relat.

Mater. 1, (1992), 364.

4. W.Luithardt and C.Benndorf: Deposition of Fe-C:H coatings from a ferrocene precursor in a plasma-activated r.f.

process. Diamond and Relat. Mater. 4, (1995), 346.

5. E.Kny, L.L.Levenson and W.J.James: Formation and properties of metallic organotin films. Thin Solid Films, 85,

(1981), 23.

6. C.Benndorf, M.Grischke, H.Koeberle, et al: Identification of carbon and tantalum chemical states in metal-doped

a-C:H films. Surf. Coat. Technol. 36, (1988), 171.

7. W.J.Meng and B.A.Gillispie: Mechanical properties of Ti-containing and W-containing diamond-like carbon

coatings. J. Appl. Phys. 84, (1998), 8.

8. Rusli, S.F. Yoon, Q.F. Huang, et al.: Metal-containing amorphous carbon film development using electron

cyclotron resonance CVD. Diamond Relat. Mater. 10, (2001), 132.

9. Q.F. Huang, S.F. Yoon, Rusli, et al.: Molybdenum-containing carbon films deposited using the screen grid

technique in an electron cyclotron resonance chemical vapor deposition system. Diamond Relat. Mater. 9, (2000),

534.

10. F. Li and J.S. Lannin: Radial distribution function of amorphous carbon. Phys. Rev. Lett. 65, (1990),1905.

11. Q.F. Huang, S.F. Yoon, Rusli, et al.: Raman scattering from the metal-containing carbon (Me-C:H) films.

Submitted to J. Appl. Phys.

12. Rusli, S.F. Yoon, Q.F. Huang, et al.: Investigation of molybdenum-carbon (Mo-C:H) deposited using an electron

cyclotron resonance chemical vapor deposition system. J. Appl. Phys. 88, (2000), 3699.

13. H.Koeberle, M.Grischke, F.Thieme, et al: A microstructural investigation of Au-a-C-H films. Surf. Coat.

Technol. 39&40, (1989), 275.

14. M. Yudasaka, K. Tasaka, R. Kikuchi, et al.: Influence of chemical bond of carbon on Ni catalyzed

graphitization. J. Appl. Phys. 81, (1997), 7623.

15. H. Windischmann: An intrinsic stress scaling law for polycrystalline thin films prepared by ion beam sputtering.

J. Appl. Phys. 62, (1987), 1800.

16. Dieter K. Schroder: Semiconductor Materials and Device Characterization. John Wiley & Sons, Inc. New York.

1998.

NASA/CP--2001-210948 478 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

GROWTH OF DLC AND BCN FILMS BY MAGNETRON SPUTTERING

USING DIAMOND AND C-BN TARGET

N. Ohtake, H. Momose and X. D. Yang

Department of Mechanical Sciences and Engineering, Graduate school of Science and Engineering

Tokyo Institute of Technology

2-12-1, O-okayama, Meguro-ku, Tokyo 152-8552, Japan.

ABSTRACT

This paper describes characteristics of DLC and BCN films grown from diamond and diamond/c-BN powder

mixture targets. Firstly, DLC films were prepared on silicon and silica glass substrates by DC-magnetron sputtering

using a sintered diamond (polycrystalline diamond, PCD) target. The PCD target was sputtered by argon and

hydrogen gases. The hardness of the DLC film increased with increasing hydrogen flow rate, corresponding to a

maximum at 10 cm3/min, followed by a drop above 15 cm3/min. The hardness of the DLC film grown from the PCD

target was higher than that for a film grown from a graphite target, indicating that the target material has a significant

effect on the quality of DLC films in sputtering deposition. The effect was very clear in the optical properties; it was

found that the optical transmittance of the DLC film grown from the PCD target with an Ar ion laser beam

(wavelength = 488 nm, 514.5 nm) was about four-fold larger than that grown from a graphite target. These results

suggest that the DLC film grown from a PCD target has the potential to be applied to novel optical coatings.

Secondly, BCN films were prepared from diamond and c-BN powder mixture target. The diamond/c-BN ratio of the

target was from 1 to 4, and the substJcate temperature was varied from 323-773K. The carbon content in the BCN

film decreased with increasing the substrate temperature. The BCN film deposited at 773 K showed high wear

resistance as well as high thermal resistance at 773 K in the air.

Key words: DLC, BCN, target material, magnetron sputtering, optical transmittance, hardness, hydrogen

addition, thermal stability

INTRODUCTION

The outstanding mechanical properties of diamond-like carbon (DLC) film suggest numerous potential

applications, including very-low-wear coatings, low friction coatings and chemical-resistant coatings 1)-4). The DLC

films have been prepared by several techniques such as ion beam deposition, DC 5) and RF 6) plasma CVD, DC and

RF magnetron sputtering 7)-12), ion plating 4),13), ECR deposition 14) and laser ablation 15) . Mechanical properties of

the DLC films prepared by these deposition techniques are high enough for the films to be applied to industrial uses;

however, the optical properties of the films are not good enough enable their use for the optical coatings due to the

absorption of visible light by the sp 2 carbon bonds remaining in the DLC films. In the case of the deposition of DLC

films by DC/RF sputtering of a graphite target, it was clearly observed that sp 2 carbon clusters were sputtered from

the target surface and formed a %loud" between the target and an anode. These sp 2 carbon clusters may play a

significant role in forming sp 2 carbon bonds on the substrate. In the deposition of DLC films by DC and RF

sputtering methods, the addition of metal precursors 7)9)11)and gas precursors s)10)n) has been carried out to improve

the mechanical and the optical properties of DLC films; however, the effect of target material on the quality of DLC

films has not been discussed.

The purpose of this paper is to reveal the effects of the target material on the optical quality of the DLC films

using a sintered diamond (PCD) target instead of the graphite target generally used for the deposition of DLC films.

The mechanical and optical properties of DLC films prepared using PCD targets are discussed in comparison with

the DLC films prepared using a graphite target. Effects of hydrogen flow rate on the hardness of the films are also

discussed. Moreover, to extend the temperature range of DLC coatings application, deposition of BCN films are

discussed using diamond and c-BN power mixture target. The thermal stability and the mechanical property of the

film are discussed in comparison with the DLC film.

NASA/CP--2001-210948 479 August 6, 2001
EXPERIMENTAL

Figure 1 shows a schematic of a DC-magnetron sputtering apparatus used for the preparation of DLC films. The

apparatus consists of a magnetron cathode, a substrate holder and a vacuum chamber. The cooled-down magnetron

cathode can supply sufficient carbon species, which are sputtered from the target at relatively low pressure. By

applying RF (13.56 MHz) power, the substrate holder is negatively charged and a negative DC voltage can be

applied. The pressure during deposition was set to 0.4-0.67 Pa. As the energy of the ion at the time it reaches the

substrate seems to be an important factor that affects film quality, the substrate biasing voltage Vac by self-biasing

was varied from 0 to -150 V.

The sintered diamond (PCD) target was 60 mm in diameter and the thickness of the sintered diamond layer was

about 0.5 mm. Since the sintered diamond included cobalt as a binder, the target was dipped in an acid bath for 24 h

to remove cobalt from the sintered diamond layer. The Si (100) and SiO2 substrates were about 20x20 mm 2 . The

substzate was rinsed in acetone and mounted on the substrate holder with a Ag paste. After sample introduction, the

chamber was pumped to approximately 1.3x10 -3 Pa, the chamber was backfilled to 2.7 Pa with Ar, and the plasma

was initiated at the substrate holder to remove contaminated layer at the surface of the substrate. After the substrate

pretreatment, DC glow plasma was generated at the sputtering cathode to grow films on the substrate. Deposition

parameters are shown in Table 1. The standard condition of hydrogen flow was 2.5 cm3/min. The substrate

temperature was measured using a thermocouple and kept at 323-373 K.

In the case of BCN film deposition, the powder target consisting of diamond and c-BN was used instead of

PCD target. The diamond/c-BN ratio of the target was varied from 1 to 4. The diamond and c-BN power were

mixed and pressed to approximately 2 mm in thick. Each powder consists of the large size grains, 100-120 gm in

diameter, and small size grains, 4-6 gm. Deposition parameters of BCN films are shown in Table 2. The standard

condition of nitrogen flow was 10 cm3/min. The substrate temperature was measured using a thermocouple and kept

at 323 K, 373 K, 473 K, 573 K, 673 K, 773K.

The samples were analyzed ex-situ using SEM, X-ray microanalysis (EDAX), AFM and Raman spectroscopy.

The transmittance of the film for visible light was measured using an Ar laser (wavelength=488 nm, 514.5 nm ). The

hardness of the film was measured by a nano-indentor. The thermal stability of BCN film was evaluated in the range

from 293 to 973 K in the air.

_,_ts_;_[3__o,,v_,= l

13 56MHz ::>-

Fig.l Schematic illustration ofDC magnetron

sputtering apparatus.

Table 1 Deposition conditions of DLC film

Target material Sintered diamond,

Graphite

DC-sputtering voltage -600 V

current 0.125 A

RF biasing frequency 13.56 MHz

RF biasing voltage (Vac) 0 - -150 V

Ar flow rate 10 cm_/min

H: flow rate 0 - 20 cm_/min

Pressure 0.67 Pa

Deposition time 1 - 3 h

Substrate Si(100), SiO_

Table 2 Deposition conditions of BCN film

RF sputtering power 200 W

RF sputtering frequency 13.56 MHz

DC biasing voltage 0 - - 150V

N: flow rate 10 cm_/min

H: flow rate 0 - 30 cm_/min

Pressure 0.4 Pa

Deposition time 1 h

Substrate Si(100)

Substrate temperature 323 - 773 K

NASA/CP--2001-210948 480 August 6, 2001
RESULTS AND DISCUSSION

DLC films grown from PCD target

Figure 2 shows the effect of hydrogen flow rate on the hardness of DLC films grown from PCD target. The

pressure was 0.67 Pa, and Vac was -100V for each film deposition. The hardness increases with increasing hydrogen

flow rate, corresponding to a maximum at 10 cm3/min after which it drops above 15 cm3/min. The film deposited at

20 cm3/min looked like a polymer-like carbon film, and it shows a very low hardness value. This result indicates that

adequate hydrogen addition improves the hardness of the DLC film in sputtering deposition using a PCD target.

According to this result, all the experiments described blow were performed at hydrogen flow rate of 10 cm3/min.

Figure 3 shows the SEM and AFM images of

DLC films grown from a PCD target (a) and grown 20

from a graphite target (b). The both surfaces are

smooth and uniform and approximately 200 nm thick.

The surface roughness of the films grown from PCD _" 15

and graphite targets were 6.3nm Rz and 8.1 nm Rz,

respectively. These results indicate that the influence "-_

of the target material on the morphology of the DLC __ 10

film is very small. In addition, no peaks from Co were _

detected by EDAX analysis of these films. The :_

influence of Vac on the deposition rate of DLC film is 5t

shown in Fig.4 for both target materials. The rate was I

estimated from the thickness of the film measured by

AFM. It is remarkable that the deposition rate of the

film grown from a PCD target is higher than that from 0

a graphite target at each Vac under the same deposition

conditions. The PCD target was porous because the

cobalt binder was removed; hence the apparent surface

area of the PCD target was larger than the graphite

target. This difference of apparent surface area in both

SEM

images

, I , I , I ,

5 10 15

H2 flow rate (cm3/min)

Fig.2 Effect of H2 flow rate on hardness of

DLC film.

Indenting force : 15 mN

Target: Sintered diamond

si (100 

Steel

q.--

2O

AFM

images

(a) Sintered diamond target (b) Graphite target

Fig. 3 SEM and AFM images of DLC films grown from sintered diamond target

(a) and graphite target (b).

NASA/CP--2001-210948 481 August 6, 2001
500

_'400

300 e,

©

:_ 200

_D

lOO

o

o

.... I .... I

!

Diamond target J

-'_,,,_'Graphite target JI

p=0.67 P,a I '

_-- Ar 10 crff/min I

% -_-_ H210 cm3/m in J -

I

N x

-50 -100 -150

Biasing voltage _o (V)

Fig.4 Influence of Vdcon the deposition rate of

DLC film grown from sintered diamond

target and graphite target.

targets could explain die difference in the growth rate.

Figure 5 shows the effect of biasing voltage on the hardness of

the DLC films that were grown from both targets in the range from 0

V to -150 V at a pressure of 0.67 Pa. The measurements were

performed more than five times at each condition. The solid lines

indicate the mean value of the measured hardness. The hardness of

DLC films prepared from a PCD target is higher than the film grown

from a graphite target at each Vclc, and increases as the negative

biasing voltage increases in the range from 0 V to -100 V, whereas

the hardness decreases with the biasing voltage in the case of the

graphite target in the range from -70 V to -100 V. These results

indicate that the target material has a significant influence on the

hardness of the DLC film.

It is very interesting that the hardness of the film grown from a

PCD target is more than twofold higher than the film grown from a

graphite target at a biasing voltage of 0 V. The difference in the

clusters generated from both targets also explains the great difference

in hardness in both films at 0 V. In some application of DLC wear

resistant coatings, an ion bombardment introduces the defects and

eliminates an electrical and a magnetic properties from the substrate.

Hence this sputtering technique will be one of the candidate of

protective coatings to be applied to such a bias-sensitive materials.

These results indicate that the carbon clusters play an important role

in forming the DLC films in sputtering deposition.

Raman spectra of the DLC films grown at the substrate biasing

voltage of-70 V are shown in Fig. 6. Each spectrum is consisting of

a G peak and a D peak, indicating that these films are DLC films. The

25

20

15,I

10,1

I

-@" Diamond _rgetI

Graphite target [

p=0.67_Pa I

Ar lOd"_m in I

laz10crt/m in ]

.i"

0 , , I ....

0 -50

Biasing voltage Vdc (v)

Fig.5 Effect of substrate biasing voltage

Vdc on hardness of DLC film.

-lOO

i

' I ' I ' I ' I ' I

' I ' I ' I ' I ' I '

(b) Graphite target

1700 1600 1500 1400 1300 1200 1100

Raman shift (cm-1)

Fig.6 Effect of target material on Raman

spectra of DLC films.

Substrate biasing voltage Vdd -70 V

G/D ratio of the DLC films grown from PCD and graphite are 1.73 and 1.53, respectively. The G/D ratio of the DLC

film grown from both targets at Vac=-100 V were 1.64 and 1.62, indicating that the difference of G/D ratio of both

films is small, whereas the mechanical property is significantly different in both films as shown in Fig.5. This result

indicates that the G/D ratio does not always represent the mechanical properties of the DLC film.

Figure 7 shows the transmittance of the DLC films deposited on quartz glass substrates (Coaning 7059) at Vac =-

100V and a hydrogen flow rate of 10 cm3/min. The thickness of the substrate is 1.0 mm. An Ar laser (main

wavelength = 488 nm and 514.5 nm ) was used in measuring the transmittance of the film. It is clear that the

transmittance of the DLC film prepared from the PCD target was about four-fold higher than that prepared from the

graphite target in spite of the large film thickness. It was also found that the transmittance of the DLC film at a

NASA/CP--2001-210948 482 August 6, 2001
hydrogen flow rate of 2.5 cm3/min was about two-fold lower

fiaan the film grown at 10 cm3/min. This result and the

hardness of the film (see Fig.2) indicate that optimal

hydrogen is essential to fabricate high quality DLC films in

normal magnetron sputtering deposition, which mainly deals

with relatively low energy ion species and carbon clusters.

These results led to the conclusion that high-quality DLC

films can be prepared at relatively low biasing voltages using

a PCD target instead of a graphite target in DC-magnetron

sputtering deposition. It will be suggested fiaat the DLC fills

grown from PCD targets have the potential to be applied to

novel optical coatings.

Preparation of BCN films

BCN films were prepared from a diamond and c-BN

powder mixture target The diamond/c-BN ratio of the target

was from 1 to 4 Firstly, the influence of the substrate

temperature on the stability of the film was checked The

diamond/c-BN ratio was 10 and a hydrogen flow rate was

constant at 10cm3/min It was found that the BCN films

•.... • • ..... • • ..... • ......... • • • • .................../ • .............../...

i i iii i: i:ili!:!iii!:ii :ii i::i !ii::li ! ili!: ili

(a) DLC grown from sintered diamond target

Thickness of DLC film : 210 nm

Transmittance:21.2 %

(b) DLC grown from graphite target

Thikness of DLC film : 150 nm

Transmittance : 4.8 %

Fig.7 Overview of DLC films on SiO2 substrate.

prepared at 323 K, 373 K, 473 K, 573 K, 673 K were not stable in the atmosphere; they peeled off from fiae substxate

in two days. The Elastic Recoil Detection Analysis (ERDA) revealed that the BCN fill deposited at the substrate

temperature of 573 K consists of B(0.057), C(0.228), N(0.171), O(0.114) and H(0.43). Since the content of

hydrogen in the film is as high as 43%, it is possible that the large amount of hydrogen as well as oxygen makes the

film unstable in fiae humid atmosphere. This phenomenon seems to be interesting in terms of hydrogen occlusion

application of the BCN film. In addition, the carbon content in the BCN film decreased with increasing the substrate

temperature.

The BCN film prepared at 773 K was stable in the air for more than 6 months, so the effect of the flow rate of

hydrogen on the hardness of BCN fill was investigated at 773 K. The substrate biasing voltage was kept constant at

-50 V. The results are summarized in Fig. 8. The hardness increases with increasing hydrogen flow rate,

corresponding to a maximum at 20 cm3/min after which it drops above 30 cm3/min. The tendency in this series of

experiments is similar to the tendency in DLC film as shown in Fig. 2. It is possible that the atomic hydrogen plays

an important role to terminate fiae dangling bonds in fiae BCN film as it is well known in DLC film deposition.

The thermal stability of BCN film was evaluated by a series of annealing experiments. Each BCN film was

heated up to 573 K, 673 K, 773 K and 973 K in the furnace at the atmospheric pressure, then kept at each set

temperature for 1 h, then the hardness was measured by

nano-indentor after the sample was cooled down. Figure 9 20 i

shows the effects of the annealing temperature on the

dynamic hardness of BCN film. The hardness is almost

constant in the range from 293 K to 773 K. Figure 10

represents the mass change of BCN films and DLC films

after the annealing test at 773 K. The mass of the BCN films

increase after the annealing, whereas the DLC films are

disappeared. The increase of mass observed in BCN film

will be explained by the oxidation of BCN films. It is

possible that the existence of hydrogen may have an effect

on the oxidation, because the mass increase is obvious in the

BCN film deposited at high hydrogen flow rate. It is widely

accepted that DLC films are stable to about 650 K, thermal

evolution of hydrogen from the films originate the

graphitization between 573 and 723 K 16)-17)

These results clearly indicate that the BCN film has

high thermal stability as comparison with the hydrogenated

DLC fill

i

I

a

I

0

i i

:= - si

1[.... .... v- • I N2flow rate-- 10 cmO/min I : _,

/ Biasingvoltage =-50 V 1 : 1

/ Substrate temprature= 773 K I :

' a , i ,

I . I

0 10 20 30

3

H2 flow rate (cm/min)

Fig.8 Relation between hardness of BCN

film and hydrogen flow rate.

DLC

NASA/CP--2001-210948 483 August 6, 2001
20

71c

0 ' I , I , I , I , I , I

100 200 300 400 500 600 700

Annealing temperature (oC)

Fig.9 Variation of hardness of BCN film with

annealing temperature.

0.2

0.1

%o 0

-0.1

-0.2

-0.3

-0.4

Fig.10 Mass change of BCN and DLC films by

annealing at 773 K in atmospheric pressure.

CONCLUDING REMARKS

Deposition of DLC films and BCN films by DC-magnetron sputtering using a sintered diamond target and a

diamond/c-BN mixture target were carried out. The hardness of the DLC film grown from PCD target was higher

than that for a film grown from a graphite target, indicating that the target material has a significant effect on the

quality of DLC films in sputtering deposition. The optical transmittance of the DLC fill grown from a PCD target

was about four-fold greater than that of a film grown from a graphite target. These results indicate that high-quality

DLC films can be prepared at relatively low biasing voltages using PCD targets instead of graphite targets in DCmagnetron sputtering deposition. The BCN film, which is prepared at 773 K in N2 and H2, showed high thermal

stability at 773 K as comparison with the hydrogenated DLC film. The optimal hydrogen addition improves the

hardness of not only DLC film but also BCN film in sputtering deposition.

ACKNOWLEDGMENTS

The authors express their gratitude to Dr. K. Suzuki and Mr. S. Watanabe of Tomei Diamond Co., Ltd, for

providing sintered the diamond targets. This work was partially supported by a Grant-in-Aid for Science Research

(No. 11750093, 12555196) from the Ministry of Education, Culture, Sports, Science and Technology.

REFERENCES

1) S. Aisenberg and R. Chabot: J. AppL Phys., 42(1971) 2953.

2) S. Miyake and R. Kaneko: Thin Solid Films, 212(1992) 256.

3) K. Miyoshi: Surf. Coat. Technol. 44(1990) 799.

4) S. Miyake, S. Watanabe, H. Miyazawa and M. Murakawa: DiamondFilms andTechnol., 3,3(1994) 167.

5) M. Ham and K. A. Lou: J. Vac. Sci. Technol., A8 (1989) 2143.

6) A. Grill, B. Meyerson, V. Patel, J. A. Reimer and M. A. Petrich: J. Appl. Phys., 61(1987) 2874.

7) C. Rebhoz, H. Ziegele, A. A. Voevodin, J. M. Schneider and A. Matthews: Surface Engineering, 13(1997) 375.

8)A. Erdemir, G R. Fenske, J. Terry and R Wilbur: Surface and Coatings Technology, 94-95(1997) 525.

9) A. A. Voevodin, M. A. Capano, S.J.P. Laube, M. S. Donley and J. S. Zabinski: Thin SolidFilms,298(1997) 107.

10) J. M. Sivertsen, G Wang, G-L. Chen and J. H. Judy: IEEE Transactions on Magnetics, 33(1997) 926.

11 )A. A. Voevodin, J. M. Schneider, C. Rebholtz and A. Matthews: Tribology International, 29(1996) 559.

12)A. A. Voevodin, R Stevenson, C. Rebholtz, J. M. Schneider, and A. Matthews:Vacuum 46(1995)723.

13) B.Rother, J.Siegel, K.Breuer, I.Muehling, S.Deutschmann, J.Vetter, G.Trommer and B.Rau,: J. Mater Res., 6(1991) 101.

14) F. S. Pool and Y. H. Shing: J. Appl. Phys., 68(1990) 62.

15) C. B. Collins et al.: DiamondFilms and Technology, 2,1(1992) 25.

16) Wu, Weng-Jin. Hon, Min-Hsiung. Materials Research Society Symposium Proceedings. 555(1999) 339.

17) S S Jr. Camargo, R A. Santos, A L Neto, R. Carius, F. Finger: Thin SolidFilms. 332(1998) 130.

NASA/CP--2001-210948 484 August 6, 2001
Proceedings oftheSixthApplied Diamond Conference/Second Frontier Carbon Technology Joint Conference (ADC/FCT 2001) Y.Tzeng, K.Miyoshi, M.Yoshikawa, M.Murakawa, Y.Koga, K.Kobashi, G.A.J. Amaratunga (Editors)

Formation ofa-C thin films by plasma-based ion implantation

Toshiya WATANABE

Japan Fine Ceramics Center, Joint Research Consortium of Frontier Carbon Technology,

c/o Research Center for Advanced Carbon Materials, AIST Tsukuba Central 5

1-1 Higashi, Tsulcuba, Ibaraki 305-8565, JAPAN

TEL : +81-298-61-9982 FAX: +81-298-61-4796

E-mail address : to i_. w_7..i..mc....g.o..j9.

Kazuhiro YAMAMOTO, Yoshinori KOGA and Akihiro TANAKA

Research Center for Advanced Carbon Materials, AIST Tsukuba Central 5

1-1 Higashi, Tsukuba, lbarald 305-8565, JAPAN

ABSTRACT

Ion-beam-assisted deposition is an efficient technique to synthesize metastable hard materials, such as cubic

boron nitride and diamond-like carbon (DLC), under lower-temperature and low-pressure conditions. However,

it is difficult to apply this technique to surface modification of three-dimensional or complex shaped objects

because of the line-of-sight process of ion beams. A new surface modification technique called "'Plasma Based

Ion Implantation (PBII)", is currently being studied. The PBII is a surface modification technique that applies

negative high voltage pulses to the object immersed in a plasma, and ions in a plasma are able to irradiate in any

direction to the object. In this study, DLC films were synthesized by using PBII technique, and the influence of

the raw gases and the duty ratio of the pulse bias to the target for the property of the DLC films were

investigated.

Ct-I4 and C2Hz gases were used as raw materials. The duty ratio of the pulse bias was changed form 0% to

50%, and the pulse voltage was -2kV. DLC films prepared by applying the pulse bias with a low duty ratio of

1% using either CH4 or C:H_ gas had very fiat surface and showed good tribological property. The C-H

absorption band measured by FT-IR was not observed within the detection limit in the DLC films formed by

using Cl-h gas as raw material, .although that in the DLC films formed by using C2H2 gas was observed. On the

other hand, hydrogen content remained in the film formed by using CH4 gas was more than that remained in the

film formed by using C:H2 gas by the measurement of ERDA. It is assumed that hydrogen atoms remained in

the film formed by using CH4 gas were not bonded with carbon atoms and they existed in the interstitial site in

the film. The tribologieal property of the DLC films formed by using CI-I4 gas was better than that of the DLC

films formed by using C2t-I2 gas. It is considered that these differences of the properties of the DLC films

formed by using either CI-h or C2H2 gas as the raw material occurred by the difference of the ratio of the film

thickness deposited when no pulse bias had been applied and the amount of irradiated ions in a pulse bias.

Keywords : plasma based ion implantation, ECR plasma with a mirror field, DLC film, tribotogieal property

NASA/CP--2001-210948 485 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

HIGHLY ADHESIVE DIAMOND-LIKE CARBON FILM BY ION MIXING METHODS

Kenji Fuchigami, Katsumi Suzuki, Masaaki Horiuchi, Kazuo Uematsu, Tadashi Sasa

Research laboratory, Ishikawajima-Harima Heavy Industries Co., Ltd.

e-mail:kerji_fuchiga mi (_ihi .co.jl"

TEL +813-3534-3338, FAX +81-3-3534-3333

Koichiro Wazumi

FCT Research Department, Japan Fine Ceramics Center

ABSTRACT

Highly adhesive diamond-like carbon (DLC) films on Inconel 718 were formed by applying ion beam mixing

method to plasma enhanced CVD process. In this study, it was investigated that effects of ion beam mixing process

on adhesion of the DLC films deposited by CVD. Adhesion strength of the DLC films were evaluated by 25N loaded

ball-on-disk friction test. Results of the friction test indicate that the DLC films formed with ion beam mixing have

high adhesion than that formed without ion mixing. It was indicated that the implanted ions and carbon atoms of the

DLC films were penetrated to the substrate and the concentration of carbon atoms was close to 40 at% by analyses of

Auger electron spectroscopy (AES). X-ray photoelectron spectroscopy (XPS) spectra of the penetrating carbon

atoms indicated that the carbon atoms formed carbide with metal atoms near an interface of the film and substrate.

These results suggest that the carbide layer near the interface formed by ion beam mixing plays an important role of

forming highly adhesive DLC films.

Keywords: DLC, ion beam mixing, adhesion, tribology

INTRODUCTION

Diamond-like carbon (DLC) have useful properties for tribological uses such as high hardness, low friction

coefficient and high wear resistance. However, poor adhesion strength of DLC films to metal materials such as steel,

alloy steel, super alloy and so on prevents to use DLC films for wide tribological uses. There have been many

attempts to improve the adhesion of DLC films by various methods, such as intermediate layer forming, multilayer

forming, metal doping, graded composition film forming, ion beam mixing and other methods. The ion beam mixing

is the process that can improve adhesion of thin films (ref. 1). However, few studies about forming adhesive DLC

films using the ion beam mixing process have been reported (ref. 2). The almost studies about the ion beam mixing

have been carried out by the process combined ion implantation system with PVD system (ref. 2). From an industrial

point of view, however, CVD process is more useful than PVD process on a basis of deposition rate of DLC film. In

this study, the ion beam mixing method was applied to plasma enhanced CVD process in order to obtain highly

adhesive DLC films with metal substrates.

EXPERIMENTAL

Fig. 1 shows a schematic diagram of the apparatus used in this study, which consists of ECR (electron cyclotron

resonance) plasma enhanced CVD system combined with high energetic ion implantation system. This apparatus

enables to irradiate mass separated gaseous ions, which are accelerated by up to 100kV, to a substrate settled in CVD

chamber. Current densities of the ion beam were measured by Faraday cups being placed at the position of the

substrate. Furthermore detailed descriptions of the apparatus were described elsewhere (ref. 3).

DLC films were prepared by ECR plasma enhanced CVD using ethylene (C2H4) gas. Inconel 718, which is

Ni-based super alloy, was used as a substrate of DLC coating. The base pressure of the apparatus was lower than

2×10 -4 Pa and the gas pressure range of CVD was about 0.1 Pa. The specimens were prepared with followings

processes. Deposition of DLC film of 10 nm thickness on the substrate and implantation ions accelerated by 50kV to

the substrate was performed alternately three times, which is called "ion mixing process" below. After that, DLC film

NASA/CP--2001-210948 486 August 6, 2001
Faraday-cups

ECR-CVD

chamber

Len_]]_

Differencial

pumping chamber

Micro wave supply

f Ion

Magnet for mas_

sepa rati°2,,,,j

Fig. 1 Schematic diagramof the apparatus.

was deposited on the substrate by CVD. The total thickness of the DLC film reached about 0.7 gm. C + ions were

used and the total doses of ions were changed in the range of about 1015 ions/cm 2 to 1017 ions/cm 2.

Adhesion strength of the films was evaluated by ball-on-disk friction test. A 9.5mm-dia. SiC ball loaded with

25N was used in the friction test, and the maximum Hertz's contact pressure was estimated at about 1.6 GPa.

Chemical compositions and bonding states of specimens were analyzed by Auger electron spectroscopy (AES) and

X-ray photoelectron spectroscopy (XPS), respectively. Thicknesses of affected areas of the specimens by ion mixing

process were measured by transmission electron microscopy (TEM).

RESULTS AND DISCUSSION

3.1 Evaluations of adhesion of films by ball-on-disk friction test

Heavily loaded ball-on-disk friction tests were used as an evaluating method about the adhesion of the DLC

films. The friction tests were carried out until a measured friction coefficient increased sharply, which was caused by

exposing substrate surface because of wearing out or peeling off the DLC film. Thus, comparing the total revolutions

of friction enables to evaluate the adhesion of the films.

Fig. 2 shows typical data of friction coefficient of two kinds of DLC films. One is the DLC film which was

formed by CVD with C + ion beam mixing with ion doses of about 4x1016 ions/cm 2 (called "ion-mixed DLC" below).

0.5

._ 0.4

0.3

©

= 0.2

._ 0.1

i DLC (without ion mixing)

• __ __ -_- -.-. -_-_ _ -_ - -___-_ . .........

.... i .... | .... | .... i .... i .... | .... i .... i .... i, ,

0 5000 10000 15000 20000 25000 30000 35000 40000 45000

revolutions of friction

Fig. 2 Friction coefficients of the DLC films.

NASA/CP--2001-210948 487 August 6, 2001
TheotheristheDLCfilmthatwasformedwithouttheionbeammixing(called"non-mixedDLC"below).Itis

showniFig.2thathetotalrevolutionoffrictionoftheion-mixedDLCislargerthanthatofnon-mixedDLCfilm.

Thedifferenceoftheresultwasexplainedbyresultsofmicroscope observationsftheslidingareasasfollowings:

thenon-mixedDLCfilmwaspeeledoffatearlierstageofthefrictiontest,althoughtheion-mixedDLCfilmendured

over40000revolutionsoffriction(ref.4).Thus,itisindicatedthatheion-mixedDLCfilmshavehighadhesion

thanthenon-mixedDLCfilm.

3.2Chemical analysesoffilms

Inordertoinvestigatetheatomicconcentration andchemical bondingstatesofcarbonatomsneartheinterface

betweenthefilmandthesubstrate, theDLCfilmsofabout30nmthicknessformedbyionmixingprocesswere

analyzedbyAESandXPS.Fig.3showsdepthprofilesofatomicconcentration oftwokindsofspecimens,

non-mixedandion-mixed(about4x1016ions/cm2dose)DLCfilm.Inthenon-mixedDLCfilm,carbonwasfoundto

somedepthpenetratetothesubstratelessthan30at%asshowniFig.3(a).Ontheotherhand,intheion-mixed

DLC,theclearlyexistenceofmixedlayerthatcontainscloseto40at%concentration ofcarbonatomswasobserved

"-" 100 "-" 100

= 80 80

© ©

60 _ 60

Ni a=

40 _ 40

¢O ¢O

= Cr = © ©

_ 20 Fe __ 20 0 ....... 0

0 10 20 30 40

m

C

t

/__" Ni

Cr

'LJI_ a|||a0

0 10 20 30 40

sputter time (min) sputter time (min)

(a) (b)

Fig. 3 AES deth profiles of the DLC films formed (a) without ion mixing and

(b) with ion mixing with dose of about 4 × 1016 ions/cm 2.

290 288 286 284 282 280 290 288 286 284 282 280

Binding energy (eV) Binding energy (eV)

(a) (b)

Fig. 4 Cls XPS spectra of the DLC films formed (a) without ion mixing

and (b) with ion mixing with dose of abaout 4 × 1016 ions/cm 2.

NASA/CP--2001-210948 488 August 6, 2001
DLC film

_'_'.*'.:_:i_ :::::::_::::__.'.'._.t._!

_:::::.: ,q:::::::::::::$_::::.'::..::::::

_!.:.".'i :_:!:!:i:i:i:i:_:i _.:?:i:i:_i:i:i:i:i:

, tg oeote e t'tt_ q ttetett.tt°e°ttt_ gttte'e't°e'e°6ee'eetet '

_..*_e __O :.:o_¢._ ,%1_0_%%_.0_%, ,%o:.:o_%:o:.:o:°:o:°.%,

,'." : .'.'......, ..........., ,....:.:...:............,

;.'.":-:':'..."I":_.:, :.:..";':..':':':' :.:':.'.:':':.:.:.'.:':.:

,:.:.:.:.:.:._ ,:.:.:.:.:.:.:.:.: ,:.:.:.:.:.:.:.:.:.:.:.:, :."...:::::::_ :::::::::::::::::::::::::::::::::::::::::::::

(a)

DLC film

.:::.....::.,.:_..." .... +: "'.;::::::

• m%el

i:: ..::i$_:i:i:i::

_:_.. .'.:::::.:.'.::

(b)

Fig. 5 TEM photographs of the DLC films formed (a) without ion mixing

and (b) with ion mixing with dose of about 4 x 1016 ions/cm 2.

near the interface as shown in Fig. 3 (b). Fig. 4 shows the Cls XPS spectra of the same specimens shown in Fig. 3.

The spectra of each specimen were obtained from the DLC film and near the interface between film and substrate,

where carbon concentration was about 20-30 at%. It is obvious that the peak energy of the Cls spectrum of the

ion-mixed DLC film is different between the DLC film and near the interface, the value of that is 284.5 eV and 283

eV respectively. Generally, the energy of C 1 s spectrum of metal carbides shows in the ranges of about 381 eV to 383

eV. Thus, it is considered that there is a certain metal carbide layer near the interface of ion-mixed DLC film.

Fig.5 shows the TEM photographs near the interface of the DLC films. It is considered that the dark area near

the interface between the DLC film and the substrate shown in Fig. 5 is the carbon penetration area indicated by AES

analyses as shown in Fig. 3, since the brightness of the area is different from that of the substrate in the photographs.

It is obvious that the thickness of the carbon penetration area of the ion-mixed DLC film is about 80nm, although

that of non-mixed DLC films is less about 30nm. The amount of the carbon atoms penetrating to the substrate by ion

mixing process is estimated to be much higher than the value of 4x 1016 ion s/cm2 which is the amount of implanted

100

80

©

"N 60

¢..) 40

©

20

0

100

80

©

"N 60

Ni _=

_/ "_ _ 40

....... Cr =¢

f .:,:."_ Fe ._

L_,_JL_L ' " I ' ' ' | ' _ 0

0 10 20 30 40

C Ni

Cr

..J .........

0 10 20 30 40

sputter time (min) sputter time (min)

(a) (b)

Fig. 6 AES deth profiles of the DLC films formed by ion mixing with

various ion doses of about (a) 2 x 1015 ions/em 2 and (b) 8 × 1016 ions/cm 2.

NASA/CP--2001-210948 489 August 6, 2001
carbon ions to the substrate. Therefore, the most of the carbon atoms penetrating to the substrate were supplied from

the DLC film by the ion mixing process. The mechanism of the penetration is considered as recoil mixing, cascade

mixing, radiation enhanced mixing and other processes reported in the preceding works of the ion beam mixing (ref.

5).

3.3 Effect of ion doses on adhesion of films

Ion-mixed DLC films formed with various C + ion doses in the ranges of about 1015 ions/cm 2 to 1017 ions/cm 2.

The DLC films were evaluated its adhesion by the friction test and were investigated the atomic concentration by

AES. Fig. 6 shows the typical AES depth profiles of the specimens. In the friction tests, the film shown in Fig. 6 (a)

peeled off in the early stage of the friction, although the film shown in Fig. 6 (b) was indicated to have high adhesion.

The results of the friction tests of the DLC films formed with various ion doses showed that the carbon concentration

near the interface of the high adhesive films is close to 40 at%. Thus, it is considered that the large amount of carbon

such as close to 40 at% is needed to form highly adhesive DLC films by ion mixing process.

CONCLUSIONS

Highly adhesive DLC films were formed by CVD process combined with ion beam mixing process. The DLC

films formed with ion mixing were able to endure in the heavily loaded friction test until the films were completely

worn out, although the films formed without ion mixing were peeled off the substrates at earlier stage of the friction

test. These results of the chemical analyses indicated that ion mixing process formed metal carbide layer between the

film and the substrate. It is also indicated that the highly adhesive DLC films contained large amount of carbon such

as close to 40 at% near the interface. Thus, it is considered that the carbide layer with high concentration of carbon

plays an important role of forming highly adhesion DLC film.

ACKNOWLEDGMENTS

This work was supported by New Energy and Industrial Technology Development Organization (NEDO) as a

part of the Frontier Carbon Technology (FCT) project.

REFERENCES

1. R.A. Kant, B.D. Sartwell, I.L. Singer and R.G. Vardiman: Nucl. Instr. and Meth. B7/8, (1985) 915.

2. S. Miyake, S. Watanabe, H. Miyazawa and M..Murakawa: Diamond Films Technol. 3,(1994) 205.

3. H.Nakai, et al.:H.Kuwahara, J.Shinohara, T.Kawaratani, T.Sassa and Y.Ikegami: Nucl. Instr. and Meth.

Bl12, (1996) 280.

4. K. Fuchigami, M. Horiuchi, K. Suzuki, K. Wazumi, K. Uematsu and T. Sasa: Proc. of ITC 2000

(publishing)

5. Z.L. Wang: Nucl. Instr. and Meth. B2, (1984) 784.

NASA/CP--2001-210948 490 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Preparation of diamondlike carbon films by high-intensity pulsed-ion-beam

deposition

X. X. Mei T.C.Ma

(State Key Laboratory for Materials Modification by Laser, Ion and Electron Beams,

Dalian University of Technology, Dalian 116024, China)

Diamondlike carbon (DLC) films were prepared onto silieon(Si) substrate by high-intensity

pulsed-ion-beam ablation of graphite targets. A 300 KeV, 250 A/cm2,100 ns beam, consisting

primarily of hydrogen and carbon ions, was focused onto a graphite target at fluence of 5-10

J/crn2. Amorphous carbon films were deposited at up to 10 nm per pulse, corresponding to an

instantaneous deposition rate greater than 1 mm/s. Raman spectra indicates that diamondlike

carbon is present in most of the films. Scanning electron microscopy revealed satisfactory film

equality. The characteristic Raman bands of these DLC films were investigated by X-ray

photoelectron spectroscopy (XPS) and X-ray excited Auger electron spectrcopy (XAES). The

analysis of the C l s peak is shown to be a very simple and direct method to evaluate the sp3

content in unhydrogenated DLC thin films.

Key words: DLC, HIPIBs, film

Principal Author (to receive correspondence):

FULL NAME:Xian-Xiu Mei

FULL ADDRESS: Department of Physics Dalian University of Technology

Dalian 116024_ China

Telephone Number: +86-411-4708380

Telefax Number: +86-411-4708389

e-mail Address: meili@dlut.edu.cn

NASA/CP--2001-210948 491 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

DEPOSITION OF DIAMOND-IAKE CARBON FILM IN A CLOSED-SPACE

CH4 RF PLASMA

N. Mutsukura and Y. Handa

Faculty of Engineering, Tokyo Denki University, 2-2 Kanda-Nishiki-cho, Chiyoda-ku,

Tokyo 101, Japan

(Corresponding author: N Mutsukura, E-mail: mutsu@d.dendai ac.jp, Fax: +81-3-5280-3565,

Phone: +81-3-5280-3421)

ABSTRACT

Deposition of diamond-like carbon (DLC) _lm and mass spectrometry measurements were carried out in a

closed-space CH4 rf(13.56 MHz) plasma (without both gas injection and vacuum pumping during the process).

At pressuresless than 0.6 Ton-, the thickness ofthe DLC film deposited was increased with an elapsed deposition

time, and reached a maximum value, but after this the film thickness was then turned to decrease because ofmaybe

ion-induced sputter etching. The maximum film thickness appeared at larger elapsed time forhigher deposition

pressure. The mass intensities ofhydrocarbonions indicated anomalous behaviors at earlydepositiontime, but

those of higher hydrocarbon ions were then clearly increased at where the film thickness was decreased. These

results suggested that when the the ratio of precursor CH3 radical density to total hydrocarbon ions density

(_cm/q_i_) in the CH4 plasma, became less than a omain critical value, the deposited rim was then re-etched

predominantly by the secondary higher hydrocarbon ions. At 1.0 Torr where a polymer-like soft carbon film was

deposited, such re-etching of deposited film was not observed.

Keywords: DLC film, Plasma CVD, Plasma diagnostics

INTRODUCTION

In general, it has been speculated that in a CH4 plasma one of main precursors is a CH3 radical, and reactions

of the CH3 radicals with hydrocarbon ions having relatively large kinetic energies on the growing flm smfa_

bring the DLC film formation (rely. 1 to 2). In our previous artides, it has also been speculated that higher

hydrocarbon ions generated by secondary reactions contribute more to the DLC film formation than do the primary

CH3 ÷ and CH4÷ions (re£3). Moreover, the deposition mechanisms for the DLC film have been discussed using

the results obtainedin raregas contained CH4 plasmas, wherethe densities of'both CH3 radical and hydrocarbon

ions were able to be changed by chemical reactions of rare gas excited atoms and ions with CH4 molecules (refs.

3 to 5). However, some questions on the contribution of ionic species during the DLC film deposition remain

even now. In this work, the ratio of CI-I3 radical density to total hydrocarbon ions density (_cm/_i®) in the CH4

plasma was able to be changed using a closed-space CH4 plasma, which depended on an elapsed deposition time,

and it seemed to be decreased with the elapsed time. Using the results obtained in this work, a mechanism on

the DLC film deposition was discussed in view of the action of hydrocarbon ions.

EXPERIMENTAL

DLC flm deposition and mass spectrometry measurements w¢_ carried out in a conventional plasma

chamber made era stainless steel (SUS) cylinder with aparallel plate configntation, previously descrived in detail

elsewhere (re[ 6). The parallel electrodes consist of SUS disks 153 mm in diameter with a spacing of 60 ram.

A power supply is a 13.56 MHz ffoscillator, andthe ffpower input was nurmally kept at 100 W. The deposition

chamber was evacuated to be less than 104 Tort, and then source CH4 gas was injected just prior to the fire

deposition. During the film deposition, all valves connected to a vacuum pump and a gas inlet were completely

dosed. A silicon wafer as a substrate was placed on a lower (ffpowered) electrode. The gas pressure in the

chamber was monitored with a capacitance manometer. The film deposition was carried out at gas pressure range

NASA/CP--2001-210948 492 August 6, 2001
of 0.07-1.0 Torr, for less than 15 min in order to avoid a eonsiderablecomtamination of the source gas due to a

leak in the closed-space deposition chamber. The leak pressure was measured to be about 0.3 mTorr/min at

these deposition pressures. For the mass spectrometry measuremems, aquadrupole mass filterwas aRchedto an

upper anode (grounded) electrode, and ions reaching the electrode were extracted through an ori_ce (200 mm in

diameter) formed at the cemet ofthe electrode.

RESULTS AND DISCUSSION

The tilm thicknesses of carbon Sims deposited at pressures of 0.07 - 1.0 Tort and rfpower of 100 W,

depending on an elapsed deposition time, are shown in Fig. 1. The negative film thickness in the figure means

0.6

0.5

0.4

0.3

0.2

0.1

k,1

0.0

-o.1

.0.2

-0.3

-0.4

I I I I

100 W 1 Torr

600 mTorr

_J

- mTo l-

-

4- I I I -I0 5 10 15 20

DEPOSITION TIME [min]

Figure 1. Film thickness of carbon film deposited at various

gas pressures, depending on an elapsed deposition time.

0.4

0.3

0.2

0.1

0.0

_- I I I I I -f

600 mTorr 100 W

- .._'_---._ _--- _ 15 0 W -

,/f/-

t , I .x.c/ow, 2s

0 2 4 6 8 10 12

DEPOSITION TIME [rain]

Figure 2. Film thickness of carbon film deposited at 0.6 Ton',

depending on an elapsed deposition time.

NASA/CP--2001-210948 493 August 6, 2001
the etching of the film, which was measured using a substrate with the DLC film deposited on it in a CH4 plasma

at same pressure with a normal gas flow, prior to the deposition processin the closed-space plasma. Each film

thickness increases with an increasing elepsed time, andindicates a maximum value at early deposition time, but

it then decreased monotonically, except the case in 1.0 Ton. : the maximum film thickness appeared at larger

elapsed time for higher deposition pressure; about 2 min at 0.07 Torr and about 8 min at 0.6 Tort. Although

the film thickness in 0.6 Tort indicated a faint decrease afi¢_ about 8 min, it beganto decrease considerably from

earlier dapsed time than 8 min when rfpower input was increased more than 100 W, as shown in Fig2. In

general, the kinetic energies ofions bombarding the growing film surface are known to increase with an increasing

ffpower input. Atthese deposition pressuresless than 0.6 Tort, theDLC films wcredeposited. At 1.0 Ton-,

on the other hand, the film thickness always increased with an ino'casing deposition time up to 15 rain (it was

confirmed to increaseuntil 20 rain), where the films deposited werepolymer-like soft films. From these results,

it can be speculated that the deercasein the film thickness is coming fi'omthe re-etching ofthe deposited film by

the sputtering action of ionic species having relatively large kinetic energies.

In order to make sure of above speculation, mass spectrometry measurements were then carried out, and the

temporal behaviors on several kinds of hydrocarbon ions were mainly measured. The results at both 0.07 and

0.3 Ton" are shown in Figs. 3 and 4. A temporal variation in pressure inside the deposition chamb_" is also

indicated at upper part in the figure. Mass signals of predominant hydrocarbon ions indicate anomalous

behaviors in both pressures. At first, those decrease rapidly, and then increase and reach a maximum value.

After this, those decrease rapidly again, and then increase slightly (only C2H5 ÷ ion density increases considerably).

The maximum value forboth CH4 +and CH5 +ions appears at somewhat larger elapsed time as compared with other

CH3*, C2H5 ÷ and C3I-I5 ÷ ions at both pressures. It is well known that secondary C2H5 + and CI-I5÷ ions are

"_ 120 I I i I

L..

 100 ...r

8o

_, 60 I I I I

10 r i i i i c2i-_ +-

1 70 mTorr

6

4 CHs+

,< 2 _. .................i 1o7'_ --::::::::::::::'" ............................ CH3+

OH- I I I I

0 2 4 6 8 10

DEPOSITION TIME [re_in]

Figure 3. Mass signals of several kinds of hydrocarbon ions

depending on an elapsed deposition time at 0.07 Tort.

A pressure change inside the deposition chamber is also given.

NASA/CP--2001-210948 494 August 6, 2001
-, 500

t..

450

'_' 400

,<

i i i i

350

300

250 I I I I -4

1.2_- I i I i -_

l• 300mTorr /

1.0_ _._.% 100W -_

i_, ° • ÷i

0.4 - o 4pNat._,_. _ CH$+ q

0.2 _,, o_ +

_-_'-22_ "" .... -_'_C3Hs÷ /

0.0- I I I I -fl

0 2 4 6 8 10

DEPOSITION TIME [rain]

Figure 4. Mass signals of several kinds of hydrocarbon ions

depending on an elapsed deposition time at 0.3 Tort'.

A pressure change inside the deposition chamber is also given.

produced by reactions of principal CI-I3+ and CH4* ions with parent CH4 molecules, respectively.

Ct'I3 ++ CH4 ---> C2H.5+ + H2 (1)

CH4 ++ CH4 ---> CI'I5+ + CH3 (2)

The C3I-B÷ ion will be predominantly produced through the reaction as follows (refs. 7 to 8).

C2H5++ C2H4 --> C3H5+ + CH4 (3)

From above reactions, it can be understood that CI-I5+ion signal follows the behavior ofCH4 +ion, and both C 2H3+

and C3I-I5+ ions go together with CH3 + ion.

The pressure in the chamber greatly increases at early elapsed time, because of the decomposition of source

CH4 molecules and the formation ofother molecules such as H2. It then reaches the maximum at around 2 and

4 min for 0.07 and 0.3 Ton', respectively. After these elapsed time, the pressure indicates a faint decrease, and

the densities of hydrocarbon ions begin to increase, as mentioned above, at where the _lm thickness decreases.

Simultaneously, the ratio ofprecursor CH3 radic_ density (q)_) to total hydrocarbonions density (¢'i,) in the

plasma was also considered to be decreased. This is supported by the following speculation. That is, C2H5 +

ion density (and accordingly total hydroca_on ion density) increased considerably, whereas C3H5 +ion density

indicated almost constant value or a faint increase aRer 2 and 4 min at 0.07 and 0.3 Ton., respectively, as shown

in Figs. 3 and 4. The C3H5 _ ion can be mainly generated through a channel ofreaction (3), and C2H4 molecule

NASA/CP--2001-210948 495 August 6, 2001
will be predominantly generated by the electron impact dissociation of C2H6 molecule [9] which can be mainly

produced by CH3-CH3 recombination reaction (ref 10). If parent CH4 molecules are mainly consumed to

produce C2I-B÷ion through the reaction (1), CI-IJradical density (and accordingly C2H4 molecule density) will not

increase in the plasma. As the result, the density ofC3I-E +ion does not increase considerably as compared with

that ofC21-E + ion, and also .the film _rmation will be greatly suppressed. Considering above results and

speculations, it can be suggested that when the ratio _cm/_io_ decreasesless than a certain critical value, the film

deposited will be then re-etched by the sputtering with the hydrocarbon ions, especially predominant higher

hydrocarbon ions such as C2I-b+, In otherword, there is a competition between the film formation and the film

sputtering in the plasma, and the @cm/_io, value may determine which of these two processes will occur

predominantly.

Up to the present, several surface reaction models and mechanisms far the deposition of hydrogenated

amorphous carbon film havebeen reported (refs. 11 to 14), where the incorporation ofprecursors into the growing

flm was considered to be intannediated by a physisod_ed layer. In the CH4 plasma, the incorporation of

adsorbed CH3 radicals induced by primary CH3+and CH4÷ions, and the direct incorporation ofthe CH3÷and CH4 +

ions arcincluded in the retrace reaction models (refs. 12 to 13). Although the chemical re-etching ofthe growing

film with atomic hydrogen from the plasma was Also taken into account (ref 12), it will be very slow at room

temp_a_ure (ref 15). In our p_wious paper, the contribution of secondary higher hydrocarbon ions was

considered to be predominant in the ion-induced incorporation ofphysisorbed precursors (ref 3), and also the

simultaneous ion-induced re-etching was suggested in the DLC flm deposition (ref 14). In this work, the

possibility of re-etching the surface of the growing film by the hydrocarbon ions was dearly pointed out, and it

should additionally be taken into account in the surface reaction model.

SUMMARY

Deposition of diamond-like carbon (DLC) film and mass spectrometry measurements were carried out in a

closed-space CH4 rf (13.56 MHz) plasma, in order to make sure of DLC film deposition mechanism. At

pressures less than 0.6 Ton+, the DLC film thickness was increased with an elapsed deposition time, and reached

a maximum value. But, it then turned to decrease because of ion-induced sputter etching. The elapsed

deposition time at where the maximum film thickness appeared, increased from about 2 rain for 0.07 Torruntil

about 8 min far 0.6 Tort. The densities of predominant ionic species in the plasma were simultaneously

measured by means of mass spectrometry. The mass intensities of hydrocarbon ions indicated anomalous

behaviors at early deposition time, but those of higher hydrocarbon ions such as C2H_J÷ ion wexe then deady

increased, where the film thickness was decreased. These results suggested that when the q_cm/@io_ value

became less than a certain critical value during the film deposition process with increasing the elapsed time, the

deposited film was then re-etched probably by the predominant higher hydrocarbon ions. These phenomena

were obse_ved in gas pressures less than 0.6 Torr, where the DLC tlms were deposited. Atl.0Torrwherea

polymer-like soft carbon film was deposited, such a process change fi'om the film deposition into the re-etching of

the film was not observed. From the results obtained in this work, the contribution ofionic speciesin the CH4

plasma was speculated that the hydrocarbon ions having relatively large kinetic energies are possible to act as both

the deposition presursor and the etchant tbr the depositing DLC film, and when the relative hydrocafl_on ions

density in the plasma exceeds a certain critical value, those act mainly as the etchant. In normal CH4 plasma

(with a source gas flow in the deposition chamber), the DLC film f_rmation may always overcome the

simultaneous film etching, and accordingly the DLC film will be continued to deposit.

REFERENCES

1. Zou, J.W., et al.: The Deposition and Study ofI-Iard CarbonFilms. J. Appl. Phys., vol. 65, no. 10, May

1989, 3914-3918.

2. Mutsukura, N, Inoue, S; Machi, Y.: Deposition Mechanism ofHydrogenated Hard-CarbonFilms in aCH4

RF Discharge Plasma. 3. Appl. Phys., vol. 72, no. 1, luly 1992, 43-53.

NASA/CP--2001-210948 496 August6,2001
3. Mutsukura, N.; Miyatani, K.: Deposition of Diamond-Like Carbon Film in CH4-He R.F. Plasma.

Diamond Relat. Mater., vol. 4, 1995, 342-345.

4. Mutsukura, N.; Yoshida, K.: Deposition of DLC Films in CH4/Arand CH4/XeR.F. Plasmas. Diamond

Relat. Mater., vol. 5, 1996, 919-922.

5. Mutsukura, N.; Yoshida, K.: Deposition of DLC Films in CH4/Ne and CH4/KrR.F. Plasmas. Diamond

Relat. Mater., vol. 6, 1997, 547-550.

6. Mutsukura, Net al.: Diagnostics and Control of Radio-Frequency Chow Discharge. J. Vac. Sci. Technol.,

vol. A12, no. 6, Nov/Dec 1994, 3126-3130.

7. Burr, J.A., et al.: Some Ion-Molecule Reactions ofI-13+ and The Proton Affinity ofH2*. J. Chem. Phys.,

vol. 52, no. 12, June 1970, 6062-6075.

8. Hiraoka, K.; Aoyama, K.; Morise, K.: AStudyofReaetionMechanisms of Methane in aRadio-Frequeney

Glow Discharge Plasma Using Radical and Ion Scavengers. Can. J. Chem., vol. 63, 1985, 2899-2905.

9. Tahara, H., et al.: Diagnostic Experiment and Kinetic Model Analysis of Microwave CH4/H2 Plasmas for

Deposition of Diamondlike Carbon Films. Jpn. J. Appl. Phys., vol. 34, no. 4A, April 1995, 1972-1979.

10. Toyoda, H.; Kojima, H.; Sugai, H.: Mass Spectroscopic Investigation ofthe CH3 Radicals in a Methane

RF Discharge. Appl. Phys. Lett., vol. 54, no. 16, April 1989, 1507-1509.

11. Kersten, H; Kroesen, G.M.W.: On the Temperature Dependence of the Deposition Rate of Amorphous,

Hydrogenated Carbon Films. J. Vae. Sci. Teclmol., vol. AS, no. 1, Nan/Feb 1990, 38-42.

12. MOiler, W.: Plasma and Surface Modeling of the Deposition of Hydrogenated Carbon Films from LowPressure Methane Plasmas. Appl. Phys., vol. A56, 1993, 527-546.

13. Gogolides, E.: A Synthetic Approachto RF Plasma Modeling Verified by Experiments: Demonstration of

a Predictive and Complete Plasma Simulator. Jpn. J. Appl. Phys., vol. 36, no. 4B, April 1997, 2435-

2442.

14. Mutsukura, N.: Deposition of Hydrogenated Carbon Film in a Magnetically Confined CH4 RF Discharge.

Vacuum, vol. 56, 2000, 129-132.

15. MOiler, W. et al.: Mechanisms of the Deposition of Hydrogenated Carbon Films. 1994 Dry Process

Symposium, ed. by Nishizawa, J., 1994, 117-124.

NASA/CP--2001-210948 497 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

A METHOD FOR DEPOSITION OF AMORPHOUS CARBON

Jawad Haidar

CSIRO Telecommunications and Industrial Physics, P O Box 218, Lindfield NSW 2070, Australia.

Fax: +61-2-9413 7200, email: ,awad.haidar@llp.csiro.au

ABSTRACT

A method for high rate deposition of amorphous diamond is described. An arc is operated in vacuum between a

graphite anode and an auxiliary cathode. Constriction of the size of the arc attachment to the anode is achieved by

the graphite material properties, cooling and also magnetic confnement. The present method has been used to

deposit films of hard carbon on substrates of glass, silicon and steel at a rate in excess of 1 gm per minute at more

then 40 cm from the arc. Raman spectra of the deposited materials show a single wide peak, consistent with

amorphous diamond. The high density of energy input from the arc to the anode tip in the arc-anode attachment

region together with the effects of ohmic heating in the anode by the arc current cause the carbon anode to evaporate,

resulting in a high pressure zone in the plasma region immediately surrounding the anode tip. The combined effects

of the high evaporation rate from the anode and the plasma thermodynamics produce a plasma jet from the region

surrounding the anode towards the outer regions of the arc chamber, at very high velocities. Numerical simulation of

the effects described above suggest that the plasma velocities can reach values of more than 30 km/s, corresponding

to energies for the plasma species of more than 50 eV.

Keywords: amorphous diamond, plasma, films, deposition, high rate.

INTRODUCTION

Amorphous carbon films (a-C) can be deposited using various processes, including filtered cathodic vacuum arc

process and the pulsed laser arc process. A comprehensive review of the various available technologies for

deposition of amorphous carbon films can be found in ref. 1. For the filtered cathodic vacuum arc process, materials

emitted from the cathode spots and ionized in the plasma are driven through a magnetic filter onto the substrate to

prevent macrodroplets emitted from the cathode spots from reaching the substrate. Cathodic vacuum arc systems

produce high quality films but their characteristic deposition rate is very low. For pulsed laser arcs, an arc of less of

duration less than 1 ms is operated at high frequency between a graphite cathode and an auxiliary anode. The arc is

triggered using a high power laser pulse. Heat from the laser pulse to the cathode surface results in the formation of a

localized carbon plasma region in front of the cathode surface, allowing the arc to start and then the discharge is

maintained for a period of less than one millisecond. For this technique, the substrate is usually placed along the line

of sight of the cathode spots, leading to higher deposition rate of up to more than 10 gm/hour at a distance target to

substrate of around 150 mm (ref. 1). However, the quality of the film is lower than for cathodic vacuum arcs.

We have developed a new process for deposition of amorphous carbon films, using an arc in vacuum between a

rod anode of graphite and a cathode disc also of graphite. Carbon vacuum arcs with a rod disc configuration have

been used for decades for production of various carbon structures (ref. 2). Recently, this configuration has been used

for production of fullerenes and nanotubes at moderate pressures below 1 arm using short arcs of less than a few mm

NASA/CP--2001-210948 498 August 6, 2001
(ref.3).Foranarcinvacuumbetweenarodanodeofgraphiteandadisccathodealsoofgraphite, thearcattachment

totheanodespreadsusuallyoveralargeareaoftheanodesurface, andmaterialsevaporated fromtheanodehave

lowenergies, makingtheprocessinadequate fordepositionfsp3richfilms.It isusuallyacceptedthatdepositionf

a-Cfilmswithhighsp3contentrequirestheenergyofthedepositedspeciestobearound30eV(ref.5). Wehave

extendedthisrod-discconfiguration forproductionofhardcarbonfilmswithahighsp3content ofmorethan50%.

Thefollowingsectionsprovideadescriptionftheexperimental arrangement andpresent examplesofdeposited

filmstogether withvariousanalyses.

Anode feed system

Anode

_i_ Vacuum

,_ _ Magnetic coil

kt

Cathode

Cathode shield

Figure 1: A schematic diagram of the arc and the electrodes

EXPERIMENTAL SET-UP

Figure 1 is a schematic representation of the arc and the electrodes arrangement. An arc is struck in a closed

vessel at a pressure less than 10 -4 torr between a water-cooled anode rod of graphite and a water-cooled cathode disk

also of graphite. The graphite anode rod is surrounded by a nozzle, preventing the arc from spreading over the

cylindrical top section of the rod. A mechanical system moves the anode rod downwards so as to compensate for

losses due to erosion from the tip. A cylindrical shield surrounds the cathode to prevent particles emitted from the

cathodic spots from reaching the substrate. The arc is started by moving the tip of the anode rod into contact with the

cathode, triggering an arc between the two electrodes. Subsequently, the arc length is maintained at 10-20 cm by

moving the anode rod downwards towards the cathode surface to compensate for losses at the anode tip due to

evaporation. As mentioned before, under normal operating conditions the arc would spread to cover a large area of

the anode surface, resulting in a low evaporation rate from the anode and also a low energy for the evaporated

species. For some graphite grades of relatively high electrical conductivity such as RW IV Ringsdorft, we find that

the size of the arc attachment to the anode rod is relatively small, allowing for deposition of hard carbon films having

a hardness of more than 15 GPa. However, operating under such conditions is somewhat unreliable as reproducibility

depends on various parameters, including the cathode materials and the initial triggering of the arc.

NASA/CP--2001-210948 499 August 6, 2001
Forthemethodpresentedhere,weincreaseboththeevaporation ratefromtheanodeandtheenergyof the

evaporated speciesbyconstrictingthesizeofthearcattachment totheanode. Thishastheeffectsofincreasingthe

pressureinthecarbonplasmaintheregionimmediately surrounding theanodetip,resultinginahighionization

degreeforthecarbonvaporfromtheanodeandalsoinahighvelocityjetdirectedtowardstheouterregionsofthe

arc.Toconstrict thearc-anode attachment regionweusevariousmeans, suchaslimitingthesizeofthesurfacearea

of theanoderodthatisexposedtothearc,andmoreadequately bycontrollingtheplasmawithamagneticcoil

placedaroundtheanoderod.Themagneticfeldduetothecoilcausesthearctoconcentrate ontheaxisofthecoil

thatcoincideswiththeaxisoftheanoderod.Fortheexamplesanddiscussions intherestofthispaperweuseacoil

tocontrolthearc-anodeattachment, producingamagneticfieldof upto 100Gaussonitsaxis.Thesubstrateis

placedalongthelineofsightoftheanodetipasshowniFigure1.Foraconductingsubstrate, abiasvoltagemaybe

appliedtothesubstratesoastoaffecttheenergyofchargedspeciesfromthecarbonplasmadirectedtowardsthe

substratesurface. Wefoundthathardcarbonfilmsindicatinghighsp3content canbedepositedonsubstrates of

silicon,metalsandalsodielectricswithoutapplyinganybiasvoltage.

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Figure 2: A magnified view of a film surface; area 80_m x 100 gin.

Figure 2 shows a magnified view of a film deposited on a Si substrate, corresponding to a surface area of 80 gm

x 100 gm. The dimension of the particulate inclusions seen in Figure 2 is less than one micron. These particulates

may be due to sources such as macroparticles from the cathode and/or dust from the internal walls of the arc vessel.

Also, it may be that the macroparticles are due to disintegration of parts of the anode tip. However, this seems to be

unlikely as physical effects governing heat transfer and evaporation from the anode tip do not favor explosive

emission of particulates from its surface. Emission of particulates is well known to occur at the cathode spots due to

effects associated with current constriction at the exit of the cathode and cooling of the cathode surface due to

electronic emission. For the cathode, the resulting combined effects may lead to a maximum temperature in the

cathode bulk right below the surface, providing conditions favorable for ejection of particulates from the surface. For

the anode, the surface in the arc attachment region is heated by the plasma and therefore the temperature is expected

to be highest at the surface, with no particle emission.

On a microscopic scale, Atomic Force Microscope micrographs for the film in Figure 2 show that the

microscopic irregularities in the film surface are small at less than a few tens of nanometers. The corresponding film

hardness, measured using nano-indentation is more than 30 GPa.

NASA/CP--2001-210948 500 August 6, 2001
35OO

3000

2500

2000

1500

1000

5OO

0

1000 1200 1400 1600 1800 2000

CM-1

I ....... GRAPHITE DLC

Figure 3: A comparison of Raman spectra of a deposited a-C film (otherwise referred to as diamond like

carbon DLC) and the graphite anode material.

Figure 3 compares Raman spectra for the a-C film in Figure 2 and for the graphite material of the anode rod. It

is seen that for the film, there is a broad peak centering around 1488 cm -1, typical for amorphous diamond (ref. 4).

As seen in Figure 3, graphite has two distinct peaks corresponding to the G and D bands. Furthermore, the sp3

content of the film has been determined using Electron Energy Loss measurements (EELS) for a film deposited on a

KCL crystal. The measurements indicate an sp3 content of around 50%. As the effects of the plasma sheath at the

interface between the non-conducting KC1 substrate and the plasma limit transfer of ions to the film, it would be

expected that films deposited on conducting substrates would have higher sp3 content.

THEORETICAL CONSIDERATIONS

As mentioned before, it is generally accepted that deposition of a-C films with high sp3 contents requires

energetic species with around 30 eV. For cathodic deposition processes, the energetic carbon species emanate from

the cathode and there has been an extensive discussion in the literature about the nature and mechanism of cathode

phenomena. References to the literature on these processes can be found in ref. 4 and also in the general literature on

vacuum arcs. For conventional anodic evaporation, it is commonly considered that it is somewhat similar to

conventional evaporation processes and as such the evaporated species would have low energies. As mentioned

before, for the process described here, films with 50% sp3 content have been deposited, suggesting that the carbon

species emanating from the anode tip may have energies equivalent to those in cathodic vacuum arcs.

The physical characteristics of the process described in this paper have been investigated through numerical

simulation of the arc and electrode configuration shown in Figure 1. A detailed description of the model as applied

to atmospheric pressure arcs is presented in ref.5. Here, we only outline the main features of the model and we

present brief details about features specific to the vacuum arc configuration used here. The aim of the present

calculations is only to predict the main characteristics and behavior of the carbon vapor emitted from the anode. The

model is based on the dynamic conservation equations of mass, energy, radial momentum, axial momentum and

current, and also the Maxwell equation and Ohm's law. The calculation domain is divided into an anode region, an

arc region and a cathode region. These regions are described relative to a cylindrical coordinate system assuming

symmetry around the arc axis. The plasma is assumed to be in local thermodynamic equilibrium and appropriate

boundary conditions are used at the edges of the calculation domain taken to be a zone around the arc and the

electrodes inside the arc chamber. Also, for the present calculations, we take the shape of the anode and also the

evaporation rate from experimental measurements as input parameters for the model. Other inputs required by the

code include the arc current, the electrode geometrical configuration and material properties of the electrodes as a

NASA/CP--2001-210948 501 August 6, 2001
functionoftemperature. Thematerial functionsofthecarbonplasmaatpressures between104atmand1atm,at

temperatures upto25000Karealsoprovidedasinputsforthecode.Theoutputofthemodel includes2dimensional

distributions oftemperature, velocities, pressureandelectricpotential. Forthepresent calculations, wetakethebase

pressureinthechamber tobe104arm.Theaimofthepresent calculations isonlytoinvestigatethebehaviorfthe

plasmaneartheanodetip,andalthoughtepressureusedfordepositionislessthanforthecalculation, theresults

wouldstillbevalidasthecalculations arelimitedtoasmallregionaroundthearcandtheelectrodes.

2,50 _ i

Anode

F.-.;>:"J.)/i} ",,+-'---.

-1.2_--i/il)}/ ",, _,,

I / / I , _\

b ."/tl / \ ",

J

<.,,o°,

-I2,5o

o,.oo _.25 2-_o 3.7_ 5.oo

R_cllml IImmlll©:n(cm)

Figure 4:Temperature contours for a 175 A arc. The outermost contour is 3000 K and contour intervals are

2000 K.

It is evident that the assumption of equilibrium used in the plasma may not be valid at the low pressures that

exist in the outer region of the arc chamber, where mean free path for the plasma species may be comparable to the

physical dimensions of the chamber. However, the results are valid near the anode tip where we find that the pressure

can increase by several orders of magnitude to near atmospheric, making the plasma in the anode tip region similar

to thermal plasmas.

In Figure 4, we present calculated temperature contours in the arc and the electrodes for a current of 175 A. As

mentioned before the base pressure in the vessel is taken to be 104 arm. The maximum plasma temperature predicted

here is more than 20000 K in the plasma region near the anode surface, suggesting that the plasma in this region is

fully ionized. Also, for this region near the anode surface, the pressure increases to more than 0.5 arm, justifying the

use of the equilibrium assumption there as mentioned before. The presence of the high pressure zone near the anode

is due to the compounded effects of large evaporation from the anode tip, ohmic heating in the plasma by the arc

current and the effects of the magnetic pinch force due to the magnetic field generated by the arc current. Because of

the resulting very steep pressure gradients between the region near the anode tip and the outer regions of the arc, the

plasma velocities calculated in the region near the anode tip exceeds 50 km/s in the vertical direction and more than

30 km/s in the radial direction. As the outer region of the arc and the chamber are at a very low pressure, it would be

expected that the velocities of the plasma species would be maintained as they travel towards the internal walls of the

chamber. The velocities calculated above correspond to translational kinetic energies of more than 25 eV, which may

explain the high sp3 content mentioned above.

NASA/CP--2001-210948 502 August 6, 2001
If theplasmaeffectswereeliminated, wecalculatevelocitiesof around2 km/sfor thecarbonvapor,

corresponding toenergyofaround0.25eV.Forthiscasewithouttheeffectsoftheplasma, thephysical conditions

wouldbeequivalent tohoseforconventional evaporation fromacarbonelectrodeatatemperature ofabout3500K.

Thismayalsobeequivalent toevaporation fromtheelectrodewithawidearc-anode attachment region.Thisenergy

of0.25eVisconsistent wihexperimental observations showingthatconventional carbonevaporationsinadequate

fordepositionfamorphous carbonfilms.

Forthepresent calculations, thecathodeisconsideredasanauxiliaryelectrodeandevaporation fromitssurface

isnottakenintoaccount. However, theresultsdiscussedabovefortheanodemaystillbeapplicableforthecathode.

Forthecathodeinavacuumarc,thearcisnaturallyconstrictedduetotherequirements ofdeliveringthearccurrent

fromitssurface. Theextremeconstrictionof thearcatthecathodespotsleadstoformationof localizedhigh

pressurezonesinfontofthecathodespotsduetothecombinedffectsoftheplasmaandthehighevaporation rate

fromthespots.Wesuggest thatheresultingsteeppressuregradientsinthearccausetheformationofaplasmajetat

highvelocities, consistent withexperimental observations. Thus,ionsareproducedthermallyinthehighpressure

zoneinfrontofthespotsandthenaccelerated duetopressuregradients, withnoneedforapotential humpinfrontof

thecathodesurfaceasusuallyadvanced toexplaintheformationofmultiplychargedionsintheplasma(ref.4).

SUMMARY

A new process for deposition of amorphous carbon films is described. The process may be used for deposition

of films over very large area with a deposition rate exceeding 50 gm/hour. Raman analysis and EELS measurements

suggest that the films are of amorphous diamond with an sp3 content exceeding 40%. Modeling of the process

suggests that the energies of the species evaporated from the anode tip have energies in excess of 25 eV, enabling

successful deposition of high quality amorphous carbon films.

REFERENCES

1. Wei, Q; and Narayan, J.: Superhard diamondlike carbon: preparation, theory and properties. Int. Mat. Rev.,

vol. 45, no. 4, 2000, 133.

2. Roger, B.: Growth,structure and properties of graphite whiskers., J. Appl. Phys., vol. 31, no. 2, 1960, 283.

3. Iijima, S.: Helical microtubules of graphitic carbon., vol. 354, November 1991, 56.

4. McKenzie, D.R.: Tetrahedral bonding in amorphous carbon., Rep. Prog. Phys., vol. 59, 1996, 1611.

5. Haidar, J.: A theoretical model for gas metal arc welding and gas tungsten arc welding. I., J. Appl. Phys.,

vol. 84, 1998, 3518.

NASA/CP--2001-210948 503 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

The Influence of Substrate Bias and Argon Flow on The Properties of

Amorphous Diamond Deposited by Cathodic Arc

Wen Ting Yeh 1,Sea-Fue Wang', James Sung I

1. Kinik Company, and National Taipei University, Taipei, Taiwan, R.O.C.

2. National Taipei University, Taipei, Taiwan, R.O.C.

Abstract

Amorphous diamond can be conveniently deposited by cathodic are. The arc contains

vaporized carbon ions that are driven toward the substrate by a negative bias. Amorphous diamond

is made of tightly packed carbon atoms jointed together by distorted graphitic (sp 2) bonds, or

diamond-like (sp 3) bonds. The sp3/sp 2 bond ratio is dependent on the kinetic energy of carbon ions.

This kinetic energy can be increased by the bias of substrate. However, with increasing of argon

gas pressure, the mean free path of carbon ions reduces, so is their kinetic energy. Hence it is

expected that sp3/sp 2 bond ratio to increase with the increasing bias, but it may decrease with

increasing rate of argon flow.

Amorphous diamond films were deposited onto cemented tungsten carbide (K10) substrates for 1

hour at two bias levels (-15 and -20 V) with two rates of argon (16 and 125 SCCM). The thickness

of these films was estimated by examining $EM micrographs of cross sections. The measured

values ranged from 0.65 to 1.25 microns. It was found that the thickness of amorphous diamond

tended to surge with the small increase of bias, possibly due to the increase of the flux speed of

carbon ions. However, when the rate of argon flow increased, the deposition rate also gained

slightly.

The hardness of amorphous diamond was determined by Vickers scale to range from 2484 to

3457 Kg/mm 2. It did not vary much with the small change of bias. However, with the increasing

rate of argon flow, the carbon ions lost some kinetics energy, hence the hardness reduced. The

adherence of amorphous diamond was tested by observing compressionaI spauling of films under

the load (150 Kg) of a Rockwell indenter. It would appear that the adherence was greatly

3 2

strengthened by increasing the rate of argon flow. This increase is likely due to the rise of sp/sp

bond ratio, and hence the reduction of compressional stress that is the driving force for delaminating

the amorphous diamond film from the substrate.

NASA/CP--2001-210948 504 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

STRUCTURE AND MECHANICAL PROPERTIES OF CARBON FILMS PREPARED

BY MASS-SEPARATED ION BEAM DEPOSITION

Kazuhiro Yamamoto*, Yoshinori Koga, and Shuzo Fujiwara

National Institute of Advanced Industrial Science and Technology (AIST),

1-1-1 Higashi, Tsukuba, Ibaraki 305-8565, Japan

Toshiya Watanabe, Koichiro Wazumi, and Fumio Kokai

Joint Research Consortium of Frontier Carbon Technology JFCC,

c/o National Institute of Advanced Industrial Science and Technology (AIST),

1-1 Higashi, Tsukuba, Ibaraki 305-8565, Japan

*corresponding author, e-mail; k-yamamoto @aist, g_',jp

Phone; +81-298-61-4503

Fax; +81-298-61-4474

ABSTRACT

An mass-separated ion beam deposition (MSIBD) system was developed with a high ion current density under

ultra-high vacuum conditions. In previous studies, carbon ion current was low, and the vacuum pressure during

deposition was insufficient. There seems to be an influence of residual hydrogen gas, and the damage to the film by

the high-speed neutral species produced by the charge exchanging with the residual gas molecules during the ion

transportation. Typical carbon ion beam density of this MSIBD system is 0.2 mA/cm 2 at the energy of 100eV under

a vacuum pressure of 6x10 -7 Pa. Carbon films were prepared and the structure and sp 3 bond fraction of films

were examined by transmission electron microscope(TEM) and electron energy loss spectroscopy(EELS). The

structure of all films are amorphous. The dependence of the sp 3 bond fraction on the carbon ion energy was

examined. The maximum sp 3 bond fraction of 84% is obtained at an ion energy between 75 and 100 eV.

Nanoindentation test of prepared films was performed and the hardness of the films with the highest sp 3 bond

fraction is 80 GPa.

Keywords: ta-C, DLC, sp 3 bond, MSIBD, Hardness

1. INTRODUCTION

Hard amorphous carbon films have recently attracted much attention. These materials have a disordered

network of sp 2 and sp 3 bonded carbon atoms, and the fraction of sp 3 bonding in the hardest film reaches 80%. (ref.

1) The physical and chemical properties of these materials are similar to those of diamond, therefore, they are called

diamondlike carbon (DLC) films. The characteristics of these films include extreme hardness, high electrical

resistivity, optical transparency and high thermal conductivity. Such properties are very attractive for a variety of

applications. These diamondlike carbon films are obtained by cathodic arc, (ref. 2) ion beam sputtering, (ref. 3)

direct carbon ion beam deposition (refs. 1, 4 to 5) and pulsed laser deposition. (refs. 6 to 8) It is reported that the

energetic carbon species which attach to the substrate are effective due to subplantation. (refs. 9 to 10) The incident

ions are quenched in the subsurface or displaced into the subsurface layer, and cause the formation of sp 3 bonding.

The threshold ion energy of this effect is less than 100 eV. (refs. 10 to 11).

Mass-separated ion beam deposition (MSIBD) is the most suitable method for the examination of DLC

formation since it can provide a monoenergetic ion species. Some studies of carbon films by MSIBD have been

reported previously ; (refs. 4, 12 to 14) however, the carbon ion current has been low or the vacuum pressure during

deposition has not been good. With low vacuum pressure during deposition, there seems to be an influence from the

residual hydrogen gas and damage to the film of the high-speed neutral species produced by the charge exchange

NASA/CP--2001-210948 505 August 6, 2001
duringtheiontransportation. Ontheotherhand,residual hydrogengasadsorptionrthesitewithlessnucleation

underlowioncurrent densityseemstohaveaninfluence.

Inthisstudy,anMSIBDsystemwithahighioncurrentdensityof0.20mA/cm2underanultra-highvacuum

of6x10-7Pawasdeveloped. Thecarbonfilmswerepreparedatvariousionenergies, andthesp3bondfractionof

thefilmswasexaminedby electronenergylossspectroscopy. Hardnessof filmswasalsodeterminedby

nanoindentation test.

2.EXPERIMENTAL

CarbonfilmswerepreparedonSi(100)wafersbyanMSIBDof 12C+ions.TheMSIBDsystemisshown

schematically inFig.1. Oursystemconsistedof anionsource, amasseparationmagnet, atransport tubewith

somemagneticlenses,adeflectionmagnet, decelerationelectrodes, anda depositionchamberwitha sample

exchangetransporter. Theplasma-filament ionsource, whichisoperatedatahigherdischargecurrentwithalonger

lifespanthantheconventional Freemanionsource, wasusedastheionsource. (ref.15)CO2wasusedasasource

gasofcarbonions,andArwasusedasafilamentplasmagas.Thecarbonionswereextractedfromtheionsource

atapproximately 34kVandmass-separated byasector-type mass-analyzer. Thecarbonionbeamwasthenfocused

byatripletquadrupole magneticlens.Thesteeringmagnet lensinfrontofthequadrupole lenswasusedtoposition

beamsinabeamtransport tube.Theionbeamwasbentbythedeflector magnet, andthefastneutrals, whichwere

producedbyinelasticcollisionsofthefastpositivecarbonionswithresidual gasmolecules, weretrapped.After

defection, theionbeamwasdecelerated bytripletelectrodes anddepositedonasubstrate.All magnetsandabeam

transport tubewerefloatedatanacceleration voltageof-34kV,andtheionsourcewasbiasedatapositivevoltage

withrespect togroundlevel.Thefinalionenergywasdetermined bythesourcevoltageplustheplasmapotential of

theionsourcewithrespect tothegroundedsubstrate. Thetotalioncurrentatasubstratereached1.0mA,andthe

carbonionbeamdensitywas0.20mA/cm2attheenergyof100eV.Thesystemequippedthefivecryopumps witha

pumpingspeedof 4000l/s,whichwassufficientto evacuatehehydrogengas. Thebasepressureof the

mass-separation tube,thebeamtransport tube,andthedepositionchamber were3.0x10-4, 1.0xl0-6,and1.0xl0-8

Pa,respectively.

V7

• Deceleration [_. Substrate

Magnetic electrodes /'_

quadrupole "_ __

lens Insulator I __,%_" [] a ; ,

1"2"21 _ k_,,,_1_j,," II-M. J Transfer

Mass [_P] ..... _'_'_ I_! chamber

maalysing _ _ _ _ y_._. Growth _'__.

magn_ ;-_mber _P

IcP_l" Icfl qioo_,.,. Ic_ Magnetic

_ _ ...... .e ' ' deflector

,_ magnetic CP ;Cryopump

Plasma- TMP ; Turbo molecular pump

filament ion

source

Figure 1. Schematic diagram of the apparatus for mass-separated ion beam deposition.

A Si substrate, which was cleaned with a solution of HF and HNO 3, was introduced into the UHV deposition

chamber. The substrate was not heated during deposition. The base pressure of the deposition chamber was 1.0x

10-8pa. The vacuum pressure was 6x10 -7 Pa during deposition.

The micro-structural characterization in plane view specimens and EELS analyses were carried out by

transmission electron microscopy (TEM) (Leo, model EM 922), operated at 200 kV. The microscope was equipped

with an OMEGA-type energy filter. The specimens for TEM observations were prepared by conventional

mechanical grinding and dimpling and thinned by ion milling with Ar + at an energy of 5 kV. The incident angle of

ions was 4 ° to prevent damage to the specimen by ion irradiation.

NASA/CP--2001-210948 506 August 6, 2001
Hardnessofpreparedfilmswasdeterminedbyshallowdepthindentationtester(Nanoinstruments, model

Nanoindenter XP).Thehardness idetermined bythecontinuous stiffnessmeasurement (CMS)technique. (ref.16)

TheBerkovitchdiamondtip wascalibratedonfusedsilica. Thetribological experiment wasperformedby

ball-on-disktypefrictiontester,whichwasconductedwitha1Nnormal loadindryairatroomtemperature. The

matingmaterial wasSiCballwiththediameter of6mmandslidingspeedwas20mm/sec.

3.RESULTSANDDISCUSSION

3.1.Structureandsp3bondfraction

CarbonfilmsweredepositedbyC+ionswithenergiesinarangeof30to 1000eV.Sisubstrates werenot

heated.Theioncurrentdensitywasapproximately 0.20mA/cm2,andthevacuumpressureduringdepositionwas

under9.0xl0-7Pa.The structureoffilmswasexaminedbyelectrondiffraction.Theelectrondiffractgraphs of

allfilmsshowahalopattern, thereforethestructureisamorphous. Theionenergydependence ofthecarbonK-edge

EELSspectraisshowninFig.2. There* peak at an energy of 284 eV can be clearly observed in the film prepared

at the ion energy of 1000 eV. The peak intensity of re* decreased as the ion energy decreased to 75 eV; it then

increased again to an ion energy of 30 eV. This dependence indicates that the sp 3 bond fraction is changed by ion

energy. The near-edge structure at the energy ranging from 320 to 340 eV is not observed for all films. The

structure of all films then is amorphous. This result agrees with the electron diffractgraph of films, which shows a

halo pattern.

S:

I

...j

280

i , i i i i i

300 320 340

Electron loss energy (eV)

Figure 2. Electron energy loss

spectra in the carbon K-edge

region of prepared fdms as a

function of carbon ion energy.

10(

8C

=

.o

6C

"g 4C

©

2(

...... i .......................

10 100 1000 10000

Ion energy ( eV )

Figure 3. Dependence of sp3 bond

fraction on carbon ion energy.

100

_-_ 80

60

_= 40

_z

20

o

10 100 1000

Energy (eV)

Figure 4. Dependence of hardness

on carbon ion energy.

The fraction of sp 3 bonds was calculated from the EELS spectra using the method of Berger et al. (ref. 17)

Their method was based on a comparison of the integral under the re* peak normalized by the integrated counts over

a defined energy width in the carbon K-edge spectrum to that of standard materials consisting of 100% sp 2 bonds.

2

Although graphite consists of only sp bonds, the shape of the K-edge spectrum is sensitive to the orientation of the

incident electron beam with respect to the c-axis of graphite. (ref. 18) In this study, a highly oriented pyrolytic

NASA/CP--2001-210948 507 August 6, 2001
3

graphite(HOPG)wasusedasthestandardmaterial, andtheincident electronbeamwasparallel toitsc-axis.Thesp

bondfractiondependence withionenergyisshowninFig.3.Thefractionofsp3bondsofthefilmpreparedatthe

ionenergyof 1000eVis41%andincreasesatheionenergyisdecreased. Themaximumvalueofthesp3bond3

fractionis84%attheionenergyintervalof75to 100eV;it thendecreases atlowerenergy.Thefractionofsp

bondsdecreases to70%attheionenergyof30eV.Toformthesp3bonds,theenergeticC+ionspeciesmustreach

thesubstrate. Thefilmwithansp3bondfractionofmorethan80%wasobtainedatanionenergybetween50and

125eV.

3.2.Mechanical properties

Hardness ofcarbonfilmsdepositedbyC+ionswithenergiesinarangeof30to1000eVwasmeasuredwith

nanoindentation tester.Theionenergydependence ofthehardnessisshowninFig.4. Thehardnessofthefilm

preparedattheionenergyof1000eVis46GPa,andincreases astheionenergyisdecreased. Themaximumvalue

ofthehardnessis80GPaattheionenergyinterval of 100eV,andthenit decreases atlowerenergy.Thehardness

decreases to35GPaattheionenergyof30eV. Thisisverysimilartothedependence of thesp3bondfractionon

theionbeamenergy.Thefrictioncoefficientsof thepreparedcarbonfilmsweremeasuredbyball-ondisktype

frictiontester.Theobtainedfrictioncoefficient are0.10,0.08,and0.11forthecarbonfilmspreparedatanion

energyof 30,100,and500eV,respectively. Thecarbonfilmpreparedat 100eVshowsthemostmechanical

properties, suchashardnessof80Gpaandthefrictioncoefficient of0.08.

4.CONCLUSION

AnMSIBDsystemwithahighioncurrent 2 densityunderanultra-highvacuumconditionwasdevelol]ed.

Thetypicalcarbonbeamdensitywas0.2mA/cmatanenergyof 100eVunderavacuumpressurebelow10--Pa.

CarbonfilmswerepreparedbyMSIBD,andthefilmstructureandthesp3bondfractionwereexaminedbyTEMand

EELS.Thestructureofallfilmswasamorphous. Thesp3bondfractionofthecarbonfilm,whichwerepreparedat

theionenergiesrangingfrom30to 1000eV,wasexaminedbyEELS.Thefilmswithahighsp3bondfraction

greaterthan80%wereobtainedationenergiesbetween50and125eV. Themaximumsp3bondfractionof84%is

obtainedatanionenergybetween75and100eV.Thefilmhardnessshowsthemaximumvalueof80Gpa,andthe

frictioncoefficient is0.08,forthefilmpreparedatanionenergyof100eV.

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5)R.G.Pregliasco, G.Zampieri, H.Huck,E.B.Halac,M.A.R.deBenyacar andR.Righini,Appl.Surf.Sci.,103

(1996)261.

6)F.Xiong,Y.Y.Wang, V.Leppert andR.RH.Chang, J.Mater.Res.,8(1993)2265.

7)D.L.Pappas, K.L.Saenger, J.Bruley,W.Krakow, J.J.Cuomo,T.GuandR.W.Collins,J.Appl.Phys.,71

(1992)5675.

8)EDavanloo, E.M.Juengerman, D.R.Jander, T.J.LeeandC.B.Collins,J.Mater.Res.,5(1990)2398.

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10)I.Koponen, M.HakovirtaandR.Lappalainen, J.Appl.Phys.,78(1995)5837.

11)J.Robertson, DiamondRelat.Mater., 2(1993)984.

12)Y.Lifshitz,G.D.Lempert, E.Grossman, I. Avigal,C.Uzan-Saguy, R.Kalish,J.Kulik,D.Marton,andJ.W.

Rabalais, DiamondRelat.Mater., 4(1995)318.

13)M.Hakovirta, J.Salo,R.Lappalainen, adA.Anttila,Phys.Lett.A,205(1995)287.

14)C.Ronning, E.Dreher, J.-U.Thiele,ROelhafen, andH.Hofs_is, DiamondRelat.Mater., 6(1997)830.

15)E.Yabe,Rev.Sci.Instrum., 58(1987)1.

16)W.C.OliverandG.M.Pharr, J.Mater.Res.,7(1992)1564.

17)S.D.Berger, D.R.Mckenzie, andRJ.Martin,Philos.Mag.Lett.,57(1988)285.

18)R.D.Leapman, RL.Fejes,andJ.Silcox,Phys.Rev.B,28(1983)2361.

NASA/CP--2001-210948 508 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ON THE EFFICIENCY OF PLASMA FILTERS FOR THE CATHODIC ARC

SYNTHESIS OF DLC FILMS

I.I. Aksenov, D.Yu. Zaleskij and V.E. Strel'nitskij

National Science Center "Kharkov Institute of Physics and Technology", Akademicheskaya St. 1, Kharkov 61108,

Ukraine

ABSTRACT

This work demonstrates the possibility and advisability of improving and optimizing the filtering properties of

plasma filters, their prediction for the systems under design through computer simulation of macroparticle motion in

the plasma duct with due account for multiple rebounds of particles from the walls. The results of computations performed for some most known magnetic filter versions are presented. The data give evidence that the commonly accepted indispensable criterion - the absence of the direct lines-of-sight between the cathode and the substrate - is not

sufficient for providing the needed degree of macroparticle removal from the plasma. When using graphite as a cathode. The proposed method of estimating the efficiency of macroparticle flow suppression in combination with the

computation of magnetic field distribution in the transport channel of the filter permits optimizing the filter by two

criteria: (i) maximum removal of macroparticles from the output plasma flow; (ii) minimum losses of ion component

of the plasma to be filtered. Curvilinear and rectilinear filters were studied and tentatively optimized using the

method of computation.

Keywords: diamond-like carbon films, vacuum arc, plasma filter, transport efficiency, macroparticles.

INTRODUCTION

The systematic studies into the synthesis of diamond-like carbon (DLC) by the vacuum-arc method were initiated at the Kharkov Institute of Physics and Technology (KIPT) by Strel'nitskij and co-workers in the early seventies. The first publications reporting the results of those studies (ref. 1 to 3) have appeared following the communication by Aisenberg and Shabot (ref. 4) and the registration of the patent by Golyanov and Demidov (ref. 5), who synthesized the DLC films by deposition from flows of ions and neutral atoms, respectively. The very first experiments

have demonstrated a multifold advantage of the vacuum-arc method in the growth rate of the DLC condensate and its

adhesion to the metal substrate. However, the films produced had a great many rough defects due to graphite

macroparticles (MP) coming to the substrate from the arc-discharge cathode spot. The defect density was so high that

this cast some doubt upon the feasibility and utility of the new method in practice. Yet, before long, the way to overcome the difficulties associated with MP's was found. The cardinal solution of the issue has become possible with

the advent of magnetic filters (or "separators"), i.e., devices for removing macroparticles of cathode material from

the cathodic-arc erosion plasma. The work at finding the means to suppress MP's in the vacuum-arc plasma were

started by Aksenov's team (KIPT) in 1974, and as early as in 1976 the application was registered for the invention of

a plasma filter with a curvilinear plasma guide bent as a quarter of torus, as well as with S- and _-shaped plasma

guides (ref. 6, 7). Having applied that filter in their investigations, the Strel'nitskij+Aksenov team obtained the results that removed the basic restriction on the use of vacuum-arc discharge to form high-quality DLC films (ref. 8). It

was just those results that gave impetus to the deployment of wide-scale studies in the field of vacuum-arc synthesis

of hydrogen-free DLC coatings in many countries of the world (ref. 9).

In general terms, the mechanism of magnetic filter cleaning of erosion plasma can be described as follows. Between the substrate and the active cathode surface emitting the plasma with MP's there is a certain barrier installed to

exclude the direct line-of-sight between the cathode and the substrate. A screen or the walls of the bent tube - plasma

duct - can serve as the barrier mentioned. In their motion in straight lines the MP's are confronted by this barrier and

fail to arrive at the substrate, while the ion component of the plasma flow, owing to the magnetic field of particular

configuration, goes around the barrier and reaches the substrate. The mechanism of ion transport in curvilinear magnetic-electric fields of plasma filters has been the subject of numerous investigations, the results of which are systemized in the surveys (ref. 9, 10). As is commonly conceived, the transport of the plasma ion component in curvilinear

filtering systems under electron magnetization conditions occurs along the magnetic field lines crossing the active

surface of the cathode. To ensure the advantageous passage of plasma through the filter, these lines should nowhere

cross the plasma duct surface. The efficiency of transportation is determined by the plasma losses as the flow travels

NASA/CP--2001-210948 509 August 6, 2001
alongthetransport channel.Theratioofthetotalionflowatthechannelxitothedischargecurrent(Ii/Id)iscommonlyassumedtobetheefficiencycriterionofplasmapassagethroughthefilter.Themacroparticles, havingavery

largemass(ascomparedtoions)andaweakcharge, practicallyshownoresponsetothefield.Whenenteringthe

curvilinear filter,theseparticlesareinevitablyconfrontedbyitswalls.

Sincenotallmacroparticles fullylosetheirkineticenergyevenatseveral successive collisionswiththewalls,a

noticeableamount ofthemreachestheplasmaductexit,andhence, arrivesatthesubstrate. Then,theefficiencyof

cleaningtheplasmaisthehigher,thelongeristheplasmaguide,thenarrower it isandthegreateristheangleofits

totalbend(ref.10).However, inthiscase,thelossesoftheuseful(ion)component oftheflowtransported necessarily increase; theproductivityofthesystemdrops,whilethecomplexityofitsmanufacture anditscostincrease. The

attemptstoincreasetheefficiencyofsystemsandtosimplifytheirdesignbyminimizingthelengthoftheplasma

guideductandtheangleofitsbendatitsrelativelywidecrossection(ref.11),orbyusingthesimplest magnetic

systemintheformoftwoseriesrectilinear solenoidsatanangleof45°(ref.12)havefailedtogivethedesiredresult:

essential simplifications ithedesignofthesystemarenotfollowedbyanadequateincreaseinitsefficiency. The

simplificationfthemagneticsysteminthecasesconsidereddeteriorates itstransportingproperties. Butthemost

seriousnegativeconsequence ofthose"simplifications" isthedegradation ofthefilteringpropertiesofthesystem,

becausethementionedwaysof"simplification" evidentlyenterintocontradiction withtherequirements, thefulfillmentofwhichisnecessary foranefficientcleaningofplasma.

Itshouldbealsonotedthatallknowncurvilinear filtersarecharacterized byaverynon-uniformdensitydistributioninthediameter oftheoutletplasmaflow.Theeffectivecross-section areaoftheflowforthemajorityofthese

filtersisonlyafewsquarecentimeters oseveral tenscm2,atbest.Inturn,thiscausesassociated inconveniences

whenusingthedevicesinpractice. Thetreatment ofitemshavingthedimensions thatexceedtheplasmaflowcross

section, thedepositionofcoatingswithuniformdistributionofproperties(e.g.,thickness) areimpossibletorealize

withoutextracomplicationofthesystemsbyintroducingscanningunits(ref.13).

Thispaperdescribestheresultsof comparisoninvestigations of magneticplasmafiltersusedinvacuum-arc

processes ofynthesizing hydrogen-free DLCfilms.

REMOVALOFMACROPARTICLES FROMTHEPLASMA

Theexistingmethodsofdeterminingthedegreeoffilteringthevacuum-arc plasmaarebasedoncalculatingthe

densityofdefectsproducedbytheMP'sinthecondensate (withandwithoutstickingtothesubstrate) (ref.9).The

qualityofplasmafiltrationcanbeestimatedcrudelybymeasuringthesurfaceroughness (Ra)ofthecondensate producedonapolishedsubstrate. Asforthepredictionofthedegreeofplasmafiltrationtobeprovidedbythefilterunderdesign, herethedesigner hasperhapsonlyonecriterioninadditiontoitsintuition,namely, thesystembeingdevelopedmustensure"theabsenceofanydirectline-of-sight betweenthesubstrateandthecathode". Thisstrongly

complicateshechallengeoffilteroptimization, particularlyinthosecaseswherethesystemdesignedisexpectedto

beusedforcleaningtheplasmaproducedbyvacuumarcwithagraphitecathodethatgenerates hardrecoilingMP's.

Theproblemis simplifiedbyresortingto computationof MPmotion.Thetaskwassolvedin thetwodimensional approximation foraxisymmetric andplane-symmetric systems. It wasassumedthatheMP'swerehard

andspherical inshape, theinnersurfacesof theplasmaguideandtheinterceptingribsweresmooth, theparticle

bouncingbackfromthewallswaspartiallyelastic,thefunctionofangularMP'sflowdensitydistributionhasthe

formN(oO=const. Here c_ is the angle between the normal to the cathode active surface and the direction of MP

emission. The computation results are the graphic representation of trajectories of both single particles that have left

the cathode spot at angle c_ and the set of particles leaving the cathode for any sector of the two-dimensional halfspace over the cathode at a given interval Ac_. The digital data of the computation allow one to estimate the ratio of

the outlet MP flow Nex to the flow Nev,_generated by the cathode spot. The ratio

NJ N_n_ characterizes the degree of MP removal from the plasma.

To simplify the computations, it was assumed that all the particles that had left the cathode had the same velocity

Vo. The MP's emerging from graphite were put to have the maximum velocity of 300 rrdsec (ref. 13). As a result of

finite number of collisions n, V0 the reduces to Vn = lrrdsec. At this velocity, the particle either adheres to the plasma

guide walls or departs from the substrate under the action of gravitational force and/or electrostatic repulsion from

the negatively charged substrate. The motion of the particle was computed either before it crossed the plane of the

exit aperture (or substrate) or up to its nth collision with the plasma guide. In the last case, the particle is assumed to

be absorbed by the filter. The n value is given by the relationship k _ = VJV o, where k is the recovery coefficient. For

the case under consideration we have k _ 0.5 and therefore rt = frt(V,//_,, )/frtl¢ _- 9.

NASA/CP--2001-210948 510 August 6, 2001
Ofcourse, thecomputations withthementionedassumptions donotprovidethequantitativeestimationofthe

trueNJNew_ value. Yet, since the errors introduced by these assumptions have a regular statistic character, the

method is rather useful for the comparison estimation of filtering properties of the systems with the plasma guiding

duct of different geometries. The results of computations performed for some most known magnetic filter versions

are presented in Table 1. The data give evidence that the fulfillment of the commonly accepted criterion, namely, the

absence of the direct line-of-sight between the cathode and the substrate, is not always sufficient to provide the required degree of MP's removal from the plasma. The rebound of solid particles (as in the graphite case) necessitates

the application of additional measures to have the particles absorbed (ref. 14).

Table 1. Filtering quality Nex/Nent and carrying capacity for magnetic plasma filters

Type of filter Knee(45 °) Tor.(45 °) Knee (90 °) Rectang.

(ref.) (ref. 12) (ref. 15) (ref. 14) (ref. 11)

1,7 25,0 0 17,0

Ii rid, % 3,0 2,5 3,8 2,5

Dome

(ref. 16)

1,7

2,5

HCA +tor. (60 °) Radial

(ref. 17) (ref. 18)

0

6,5 8,4

TRANSPORT MAGNETIC FIELDS AND THEIR CORRECTION

Another important characteristic of the system forming a clean erosion plasma (in addition to the degree of

cleaning) is the carrying-capacity of the filter, which is dependent on the level of losses of plasma in its travel along

the filtering channel. From general considerations one should expect that the efficiency of transport would be the

higher, the shorter is the path L from the injection site (active cathode surface) to the exit, and the wider is the plasma

guiding duct (the larger is the radius R in the cylindrical plasma duct case). Really, considering that the average displacement of charged particle across the magnetic field for the time t is equal to Ax _ _ t_,, the condition for

the particle to pass through the plasma guide can be written as

R-r>_ DffD_t_ ,_ (1)

where r is the radial coordinate of injection, lx is the time of particle drift from the cathode to the exit of the sysH

tem, D is the coefficient of diffusion across the magnetic field. If t_ = L ,/_,,, D± H 2a/T (ref. 19), then condition (1) can be written as

R _>r+ HR-,/T- " " (2)

Here n is the plasma density, H is the magnetic intensity, T is the temperature, Vo is the particle velocity along

the field. Apart from the diffusion of particles across the magnetic field to the walls, one should also take into account their departure due to the flow shift in the curvilinear field in the direction of the centrifugal force along the

unit vector v_ (centrifugal drift) and in the direction [H, grad H] (gradient drift). One must also take into account the

losses due to the reflection of certain particles from contractions of the flux of transporting magnetic lines. These

contractions can arise at the sites, where the fields of individual plasma guiding duct parts are joined, e.g., between

the plasma generator output and the toroidal plasma guide input (ref. 20), between plasma guide coils.

So, the greater is the plasma guide radius with respect to the cathode radius rc (at rc -< r), the shorter is the path

between the input and output of the system, the higher is the magnetic field (up to a certain limit (ref. 7 to 20)), the

smaller is the angle of plasma guide bend, and the smaller are the variations in the thickness of a transporting bundle

of field lines, the more favorable are the conditions of plasma passage along the plasma guide duct. It is evident that

the mentioned requirements relating to the geometry parameters are in contradiction with the requirements, the fulfillment of which provides the best filtering quality of the system (see the preceding sections). In view of this, the

choice of the geometry parameters R, rb L and (p (total angle of plasma guide bend) is determined by the permissible

compromise between the "conflicting" requirements. As for the ion component losses, they can be minimized (for

the given geometry of the system under design that provides the required degree of filtration) only by selection of

optimum intensity and topography of the transporting magnetic field. The optimum level of absolute field induction

NASA/CP--2001-210948 511 August6,2001
values in the sources under design can be assessed from the data found in publications on the topic (ref. 7, 8, 20).

Computer simulation appears to be a rather convenient way of selecting the field topography. For this purpose we

used the application program SUPERFISH.

Fig. 1 exemplifies the influence of certain magnetic system elements of the plasma guide system on the behavior

of magnetic field lines. The field lines are shown for the simplest system consisting of two rectilinear solenoids that

form the field bent at 90 °. Between the solenoids, there arises a region with a strongly asymmetric

1

4

3

b

!

c

d e

\\

!6

Fig. 1 Influence of magnetic system elements to transporting field geometry. 1- cathode, 2 -stabilizing coil, 3 -

solenoids, 4 - correctional coil, 5 - substrate surface, 7 - deflecting coil.

widening of the transporting magnetic tube (shaded) (fig. la). This region is characterized by great plasma losses:

the particles drift across the attenuated magnetic field towards the "bulging" part of the region. The mentioned widening of the tube can essentially be compensated by incorporation of the correction coil between the solenoids (Fig.

lb). The influence of the output coil is illustrated by Fig. lc. Figs. 1 d, e show the shapes of the transporting magnetic tube in the vicinity of the substrate in the presence of deflection coils placed behind the substrate on either side

of the axis of the system. As the coils are energized by alternating current, the system permits swinging of the tube

(and hence, of the plasma flow) by the law which ensures the assigned distribution of the condensate in thickness

over the substrate surface.

COMPARISON CHARACTERISTICS OF PLASMA SOURCES WITH TOROIDAL (90 °) AND RECTILINEAR FILTERS

The first version of the filter with a plasma guide shaped as a quarter of torus (Fig. 2) remains most popular to

the present day in both the laboratory practice and commercial production. At the same time, the disadvantages mentioned above (complexity, a comparatively low productivity and, as indicated by many investigators, imperfect cleaning of plasma by these filters) retard a further advance of vacuum-arc synthesis of DLC films in production practice.

Therefore, the problems of perfection of the existing filters and the search for alternative solutions still remain

urgent.

NASA/CP--2001-210948 512 August 6, 2001
8 8

a b c

Fig2. Vacuum-arc plasma source with curvilinear filter: scheme (a); macroparticles trajectories in the plasma

duct with platelike ribs (b) and with the ribs of triangular cross section (c).

1- cathode; 2 - anode with focusing solenoid; 3 - stabilizing coil; 4 - plasma duct with transporting coils;

5 - output focusing coil; 6 - scanner; 7 - ribs; outlet plane.

As an alternative to a curvilinear plasma filter, the so-called "rectilinear" filter have been developed, (Fig. 3).

Here, the main design requirement, i.e., the absence of the direct line-of-sight between the plasma source cathode and

the substrate, is fulfilled through the installation of a screen on the axis of the rectilinear cylindrical plasma guide.

The transportation of the plasma ion-electron component is here effected along the axisymmetric magnetic field lines

that cross the active cathode surface and bend around the screen. The device is essentially simpler than the curvilinear version of the filter; however, up to date, the issue of competitiveness of the rectilinear filter from the viewpoint

of productivity and degree of MP's removal from the plasma still remains unclear in relation to its "toroidal" prototype. The present study is an attempt to elucidate this issue by performing comparison tests of the two competing

systems. In order for the test data to represent objectively the basic potentialities of the systems under comparison,

the geometry of their plasma-guiding channels and the topography of transporting magnetic fields were preoptimized by the use of the above-described techniques. The two systems under study have previously been described elsewhere (ref. 6, 21).

The calculations have shown that the traditional toroidal-filter geometry (Fig. 2a) with a set of platelike ribs on

the inner plasma guide surface does not provide a complete removal of MP from the plasma flow (Fig. 2b). It appears

much more efficient to suppress MP's through the replacement of platelike ribs by the ribs of triangular cross section

(Fig. 2c). With due account for previous experience (ref. 7, 8, 23), the optimum magnetic fields in the system were

chosen with respect to the maximum ion current at the plasma guide outlet.

The computed trajectories of MP's in the rectilineal filter (Fig. 3a) for the most efficient geometric version of intercepting ribs are presented in Fig. 3b. Fig. 3c shows the optimum run of magnetic field lines in the rectilinear filter

(calculation).

b

4 5

8

/

j6

a 1/ "_ axis

C

Fig. 3. Vacuum-are plasma source with reetilineal filter:

scheme (a); maeropartieles trajectories (b); magnetic field lines (e).

1- cathode; 2 - anode; 3 - trigger; 4 - stabilizing coil; transporting coils;

6 - baffle; 7 - ribs; 8 - substrate.

NASA/CP--2001 210948 513 August 6, 2001
ThetestresultsarepresentedinTable2.HereV is the deposition rate of DLC coating, S is the coated spot area,

AV is the deviation of coating thickness from the average value, m is the defect (MP) density in the coating, d - dimensions of identified MP, h is the coating thickness, Ah is the thickness deviation. The DLC coatings were produced on polished copper specimens at an arc current of 120 A. The specimens were placed at a distance of 100 mm

from the exit end of the filter. The microhardness was measured with the Nanoindentor-2 at a load of 5 mN (-5G).

The computations of MP motion, even with the assumptions much overestimating the output quantity of particles

(]Vex) show that at certain geometrical parameters of the plasma guide the Nex value can be reduced to be negligibly

small (down to zero). However, our experiments as well the experience of other investigators do not confirm this

conclusion. This can be explained by the fact that in addition to multiple rebounds that in addition to multiple rebounds other mechanisms, not included in our calculations, play an appreciable role in the MP

Table 2. Test results.

Type of S, V, h, Ah, H,

filter cm 2 gm/h gm % GPa

Tor.(90 °) 250* 2 1.1 +5 40+2

Rectilinear 200 2.5 1.5 _+5 41_+2

*) circular scan of plasma flow

d, m, N,,,,/N,,nt,%(calc. )

cm 2

gm

1...2 5 105 0

1...3 1.6 10v 4.4

transport. Nevertheless, as it follows from the data presented here, the calculation of MP trajectories in the

plasma guide with due account for MP recoiling provides a true qualitative estimate of the degree of MP removal by

the filtering systems under comparison.

CONCLUSIONS

1. The possibility has been demonstrated for the perfection and optimization of filtering properties of the existing plasma filters, and also for prediction of these properties for the systems under design through computer simulation of MP motion in the plasma guide with due account for multiple MP rebounds from the walls.

2. The proposed method of estimating the efficiency of macroparticle flow suppression in combination with

the computation of magnetic field distributions in the transport channel of filters permits optimization of the filter by

two criteria: (i) maximum cleaning of macroparticles from the output plasma flow; (ii) minimum losses of ion component of the plasma to be filtered.

3. The quantitative difference between the obtained calculated and experimental data may indicate that, in addition to multiple partially elastic rebounds of MP from the walls, other mechanisms can play an essential role in the

MP transport along the filter. Yet, the calculation of MP trajectories with due account for MP recoiling provides a

true qualitative estimate of the degree of MP's removal by the filtering systems under comparison.

REFERENCES

1. Strel'nitskij, V.E., et al.: An X-ray study of metastable modification of cubic diamond (in Russian). DAN

UkrSSR, A,5 (1976) 459.

2. Strel'nitskij, V.E., et al.: The intermediate form of crystalline carbon (in Russian). DAN UkrSSR, ser.A, 8

(1977) 760.

3. Strel'nitskij, V.E.; Padalka, V.G.; Vakula S.I.: Some properties of diamond-like films produced at condensation

of carbon plasma flow with the use of HF potential (in Russian). Zhurn. Tekh. Fiz. 48, 2 (1978) 377.

4. Aisenberg, S.; Shabot, R.: Ion-beam deposition of DLC films. J. Appl. Phys., 42, 7(1971) 2953.

5. Golyanov, V.M.; Demidov, V.P.: A way to produce artificial diamonds (in Russian). USSR Authors' Certificate

N 411037, 1973.

6. Aksenov, I.I.; Belous, V.A.; Padalka, V.G.: A device to deposit coatings in vacuum (in Russian). ). USSR Authors' Certificate No. 605425, 1978.

7. Aksenov, I.I., et al.: A device to clean macroparticles from the vacuum-arc plasma (in Russian). Prib. Tekhn.

Ehksp. 5 (1978) 236.

NASA/CP--2001-210948 514 August 6, 2001
8. Aksenov, I.I.,etal.:Ahigh-efficiency sourceofapurecarbonplasma(Rus.).Zh.T.F.50,9 (1980) 2000.

9. Boxman, R. L.: Vacuum Arc Deposition: Early History and Recent Developments. Proc. of the XIXth ISDEIV,

Xi'an, China, Sept. 2000.

10. Anders, A.: Approaches to rid cathodic arc plasmas of macro and nanopar ticles: a review - Surf. and Coat.

Techn., 120-121 (1999) 319.

11. Gorokhovsky, V. I.: U. S. Patent No. 5,435,900. 1995.

12. Falabella, S.; Sanders, D. M.: U. S. Potent No. 5,279,723. 1994.

13. Schulke, T.; Anders A.: Velocity distribution of carbon macroparticles generated by pulsed vacuum arcs.

Plasma Sources Sci. Technol. 8(1999) 567.

14. Aksenov, I. I.; Zaleskij, D. Yu.; Strel'nitskij, V. E.: On the Efficiency of System for Filtered Catholic Arc

Deposition. 1-st Int. Congr. on Radiation Physics, High Current Electronics and Modification of Materials, Sept.

2000, Tomsk, Russia. Proceedings. Vol. 3, 130 - 138.

15. Martin, P. J.; Netterfield, R. P.; Kinder, T. J.: U. S. Patent No. 5,433,836, July 18, 1995.

16. Sanders, D. M.; Falabella, S.: U. S. Patent No 5,282,944, Febr. 1, 1994.

17. Witke, T.; Siemroth, P.: Deposition of Droplet-Free Films by Vacuum Arc Evaporation: Results and Applications. IEEE Trans. on Plasma Sci., 27, 4 (1999) 1039.

18. Aksenov, I.I.: A system for forming radial flows of erosion plasma (Rus.). Trudy XIV Mezhdunar. Konf. po

Fisike Radiats. Yavlenij i Radiatsionnomu materialovedeniyu. 12-17 June, 2000, Alushta, Ukraine, 270.

19. Artsimovich, L.A.; Luk'yanov, S.Yu.: Charged particle motion in electric and magnetic fields (in Russian).

Moscow, Nauka publ., 1972, 224.

20. Anders, S.; Anders, A.; Brown, I.: Focused injection of vacuum arc plasmas into curved magnetic filters. J.

Appl. Phys. 75 (10) (1994) 4895.

21. Aksenov, I.I., et al.: A rectilinear plasma filtering system for vacuum-arc deposition of diamond-like carbon

coatings, Diamond and Related Materials. 8 (1999) 468.

NASA/CP--2001-210948 515 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

AMORPHOUS DIAMOND DEPOSITED BY CATHODIC ARC WITH

CONTROLLED ASPERITY

Ming-Chi Kan 1, Jow-Lay Haung 1, Wen Ting Yeh 2, James Sung 2

1. Department of Materials Science and Engineering, National Cheng-Kung University, Tainan, Taiwan,

R.O.C.

2. Kinik Company, and National Taipei University, Taipei, Taiwan, R.O.C.

Abstract

Amorphous diamond is diamond-like carbon (DLC) made of pure carbon. It can be conveniently

deposited by cathodic arc that delivers a flux of ionized carbon atoms. The carbon ions are driven by an

electrical field to impinge upon a substrate. At the site of impingement, carbon atoms my pile up to form

amorphous diamond with distorted tetrahedral bonding. The sp3/sp 2 ratio of amorphous diamond, and the

asperity of the deposited film are strongly dependent on the electrical current formed by the flux of the

carbon ions.

Amorphous diamond is an excellent electron emitter for FEA applications. The threshold emitting

voltage can be greatly reduced; or the emitting current significantly increased by reducing the activation

energy (work function) of vacuum, and/or by increasing the tip sharpness (field enhancement factor) of the

emitter. The activation energy can be reduced by optimizing the sp3/sp 2 ratio of amorphous diamond. On

the other hand, the tip sharpness can be increased by adjusting the asperity height of deposited film.

Amorphous diamond was deposited on silicon substrate of n-type (100) at a bias of 20 V. The current

was varied at 30, 50, 80, and 100 A. The sp3/sp 2 ratio was measured by ESCA; and asperity height, by

AFM. It was found that the sp3/sp 2 ratio peaked at a value of about 50% when the current was at 80 A.

However, the average asperity height (Ra) increased linearly with the increasing current from

approximately 0.5 nm of 30 A to about 1.1 nm of 100 A. The maximum asperity height is approximately

tenfold of Ra value.

This research confirmed that the two key factors for enhancing electron emission, sp3/sp 2 ratio and

asperity height, can be controlled by varying the current of cathodic arc, and hence an optimized

amorphous diamond film for field emission can be fabricated.

Introduction

Cathode ray tubes (CRTs) has been widely used due to their good image quality, manufacturability

and economy in displays. Unfortunately, conventional CRTs are big in size and high power consumption

for TV sets and desktop computer monitors. They are therefore unsuited for compact and portable

applications. Flat panel displays (FPD) offer significant packaging and power advantages over CRTs.

Field emission display (FED) has attracted much attention in recent years due to its excellent

properties, such as low operation voltage, high emission current, high image quality and wide viewing

angle. Carbon-based materials used as FED cold cathodes have attracted attention for emission displays and

vacuum electronic devices [ 1].

Amorphous diamond (DLC) films are very attractive for field emission application [2-4].

Fowler-Nordheim [5] correlated the field emission with the work function (O) and filed enhancement factor

(F).

1.54X10-6(g]Vg)2 [ "3'2v(,)1 J = exp - 6.87x107

(I_ ;_]V_ ] (1)

NASA/CP--2001-210948 516 August 6, 2001
WhereO:theworkfunction(ineV),v(y):thecorrectionfactorsrelatedtotheelectricfield,13:thefield

enhancement facor,Vg:thevoltageappliedtotheemitterelectrode. Onecancontrolamorphous diamond

sp3/sp2ratioanditsasperityheighttoreduceworkfunctionandincreasefildenhancement factor[1,6-7].

Diamondandamorphous diamondhavegoodmechanical, chemical stabilityandnegativeelectron

affinity(NEA).If highsp3/sp2ratioandfieldenhancement facorcanbeobtained, betterelectronemission

propertiesthanplanarfilmscouldbeachieved.

Theobjectiveofthisarticleistoinvestigatetheeffectsofthestructureandsurfacemorphology of

amorphous diamondontheemissionproperties. Samplesweredepositedbycathodicarcatdifferentarc

current. Thesp3/sp2ratiowasmeasured byESCAandasperityheightbyAFM.

Experimental

N-type silicon substrates, (100) oriented wafer with a resistivity of 0.02-0.008 f_-cm were used in

experiments. The wafer was cleaned by RCA clean process, then dried by N2 blowing. Amorphous

diamond films were grown at room temperature with cathodic arc. The substrates were cleaned by Ar ion

for 20 minutes and C ion implanted for 1 minute before deposition. The water-cooled substrate holder was

used to avoid overheat on substrate. The amorphous diamond was made of pure carbon target. A carbon

plasma stream emerging from a vacuum arc discharge passed through a magnetic coil before depositing

onto substrate. The magnetic coil filtered out neutral carbon atoms and macroparticles. The carbon ions

were driven by an electrical field to impinge upon the substrate. At the site of impingement, carbon atoms

piled up to form amorphous diamond with distorted tetrahedral bonding. The films were deposited at a bias

of 20 V. The currents used were 30, 50, 80, and 100 A separately. Films thickness was in the range of 100 -

150 nm.

The surface morphology of deposited films was studied using an Atomic Force Microscope (AFM).

The bonding type of amorphous diamond film was determined by Electron Spectroscopy for Chemical

Analysis (ESCA, XPS). A Mg Kc_ X-ray source that has energy of 1253.6 eV, accelerated voltage of 12 keV

and current of 17 mA was used For ESCA analysis. Samples were cleaned by Ar ion for 5 minutes before

ESCA analysis.

Results and Discussion

The bonding type ( sp 3 , sp 2 ) of amorphous diamond films was determined by ESCA in binding

energy region of 276 to 296 eV. Carbon 1S peak measured by ESCA could contain sp 3 and sp 2 bonding

carbon, and carbon oxide contamination in film surface. The 1S peak was fitted by Gaussian function. The

hot filament CVD diamond film was used as standard, and the sp 3, sp 2 , carbon oxide composites was at

284 + 0.2 eV, 285 + 0.2 eV, and 286 + 0.2 eV respectively.

Fig.1 shows the results of carbon IS peak fitting at various currents. It was found that the ratio of

sp3/sp 2 was about 50% at 80 A. Gaussian integral was used to calculate sp3/sp 2 ratio with calculation results

show in Fig. 2. The sp3/sp 2 ratio increased linearly with the increasing current from 30 A to 80 A. However

the ratio reduced when a high current of 100 A was reached. Cathodic arc contains ionized carbon atoms

that are mixed with clusters of molten carbon. The latter freeze to from micro particles of amorphous

carbon. The carbon atoms in these micro particles are primarily in sp 2 bonding. With the increase of arc

current, the temperature will rise, so is the amount of micro particles. It is postulated that the reduction sp 3

content when the current was increased to 100 A was due to overheating of the arc.

The ESCA results Fig.3 indicated the central position of Cls signals at different currents. Fig.1

and Fig.3 suggested that the amorphous diamond films had higher sp 3/sp 2 ratio when the central position of

Cls is closed to the sp 3 carbon bonding of 284 + 0.2 eV.

NASA/CP--2001-210948 517 August 6, 2001
(a)

"' ',,sp

Carbon//ide0X ' "' _'\," ,',:

/ ' , #,

• . ( ." . \ .

Binding Energy (eV)

2_0

©

(c)

i -' ,_c£

24_

©

Carbon oxi/de ,'

Binding Energy (eV)

Figl. Carbon 1S multi-peak fitted. (a) current 30

100 A.

,sp 2'',

;, e \ ',

j, i

,/

:i_ Carbon d_de ", '::,

< ,, ,,,

o- I

i ii

e_, f_ 28_ 2_ 2B2

(b) Binding Energy (eV)

2_lgI_n / i '".

°°_ Carbon _?(ide / _;,

x

,o_ _ / ; \ '/, x

/

(d) Binding Energy (eV)

A, (b) current 50 A, (c) current 80 A, (d) current

.2

¢:a

60-

55-

50

4_2

4o2

30-

25-

20

_5-

_0-

0

2O 30 4O 50 60 70 80 90 1 O0 110

Arc Current ( A )

g"

d

_3

100 A

80A

50A

282 284 286 288 290

Binding Energy ( eV )

Fig2. Gaussian integral calculated results. Fig3. ESCA Cls signals of all sample.

The film surface morphology and asperity height was studied using AFM in a contact model. The

average asperity height (Ra) increased with increasing current. The maximum asperity height is

approximately tenfold of Ra value. Fig.4 shows AFM 3-D micrographs. Table 1 lists roughness analysis of

AFM. An optimum energy to achieve minimum asperity height and Ra could be obtained.

Results in Fig. 2 and table 1 suggested that the highest sp 3/sp 2 ratio (50%) and nano-scale tips

(maximum asperity height is about 12.612 nm) could be obtained at arc current of 80 A.

NASA/CP--2001-210948 518 August 6, 2001
(a) (b)

4

(c) (d)

Fig.4 AFM 3-D micrographs (scanning range 3 gm x 3 lain ) of amorphous diamond films deposited

using (a) current 3{1A, (b) current 50 A, (c) current 8{1A, (d) current 100 A.( X axis: 1 lain/div., Z

axis : 20 nm/div.)

Table 1. Roughness analysis of AFM.

Parameter Ra (nm) Rms (nm) Rmax. (nm)

Bias:20V, Arc current:30A 0.567 0.727 6.073

Bias:20V, Arc current:50A 0.448 0.577 4.751

Bias:20V, Arc current:80A 0.937 1.204 12.612

Bias:20V, Arc current: 100A 1.088 1.423 16.043

Conclusion

Amorphous diamond films contained highest sp 3/sp 2ratio deposited using arc current was 80 A from

ESCA measurement. The minimum asperity height and Ra was obtained at arc current was 50 A from

AFM micrograph. We can achieve amorphous diamond film contained highest sp 3/sp 2ratio and nano-scale

tips when arc current was 80 A. Thus, the sp3/sp 2ratio and asperity height can be controlled by varying the

current of cathodic arc.

Amorphous diamond is an excellent electron emitter for FEA applications. This research indicate that

p3/sp2 ratio and asperity height, can be controlled by varying the current of cathodic arc, and hence an

optimized amorphous diamond film for field emission can be fabricated. Optimizing the sp3/sp2 ratio of

NASA/CP--2001-210948 519 August 6, 2001
amorphous diamond can reduce the activation energy. On the other hand, adjusting the asperity height of

deposited film can increase the tip sharpness.

Acknowledgement

The authors would like to thank Kinik Company, Taiwan, R.O.C. for its financial support.

References

1. G.A. J Amaratunga, M. Baxendale, N. Rupesinghe, I. Alexandrou, M. Chhowalla, T. B. A.

Munindradasa, C. J. Kiley, L. Zhang and T. Sakai," Field emission from new of thin film amorphous

carbon having nanoparticle inclusions and carbon nanotubes," New Diamond and Frontier Carbon

Tech., Vol. 9, No. 11, (1999) 31-51.

2. V.G. Litovchenko, A. A. Evtukh, N. I. Klyui, Y. M. Litvin, S. Y. Kudzinovsky, A. G. Chakhovskoi

and T. E. Felter, " Electron field emission from undoped and doped DLC films," Mat. Res. Soc. Symp.

Proc., Vol. 558, (2000) 577-582.

3. D.A.A. Ohlberg, A. A. Talin and T. E. Felter, " Emission characteristics of arrays of diamond-coated

silicon tips," Mat. Res. Soc. Symp. Proc., Vol. 424, (1997) 393-397.

4. V.G. Litovchenko, A. A. Evtukh, R. I. Marchenko, N. I. Klyui and V. A. Semenovich, " Enhancement

of field emission from cathodes with superthin diamond-like carbon films," Appl. Surf. Sci., Vol. 111

(1997) 213-217.

5. Su-Hua Yang, " Fabrication and Characterization of field emission display," Department of electrical

engineering, Nation Chen Kung University Tainan, Taiwan, R.O.C., Dissertation for Ph. D, 1998.

6. O. Groning, O. M. Kuttel, R Groning and L/Schlapbach, " Field emission from DLC films," Appl.

Surf. Sci., Vol. 111 (1997) 135-139

7. S.E. Huq, R D. Prewett, J. C She, S. Z. Deng and N. S. Xu, " Field emission from amorphous

diamond coated silicon tips," Mater. Sci. and Eng. B74 (2000) 184-187.

NASA/CP--2001-210948 520 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ELASTIC CONSTANTS OF NANOMETRIC DIAMOND-LIKE CARBON FILMS BY

SURFACE BRILLOUIN SCATTERING

M.G. Beghi, C.E. Bottani, A. Li Bassi,

Nuclear Engineering Dept., Politecnico di Milano, Via Ponzio 34/3, 1-20133 Milan, Italy,

..mg.c..o..b_e.g_aa_o_li..mi_'.i_t, fax+39 02 2399 6309, phone+39 02 2399 6351

B.K. Tanner

Dept. of Physics, University of Durham, South Road, Durham DHI 3LE, UK

A.C. Ferrari, J. Robertson,

Engineering Dept, Cambridge University, Cambridge CB2 1PZ, UK

ABSTRACT

Carbon coatings of thickness down to 2 nanometers are needed to increase the storage density in magnetic hard

disks and reach the 100 Gbit/in 2 target. Methods to measure the properties of these ultrathin films, in particular of

hard films on softer substrates, still have to be developed. We show that the combination of Surface BrilIouin

Scattering and X-ray reflectivity measurements makes the elastic constants of such films measurable.

Tetrahedral amorphous carbon films were deposited on silicon by an S bend filtered cathodic vacuum arc,

achieving a continuous coverage on large areas free of macroparticles. Films were characterised for thicknesses

down to - 2 nm. X-ray reflectivity measurements provide film thickness and mass density. Densities above 3 g/cm 3

are found, indicating a significant sp 3 content. Evidence was also found of a density gradient, the density being

lower at both the internal and external surfaces. Surface Britlouin scattering measurements provide the dispersion

relations of surface acoustic waves at wavelengths of few hundred nanometers. Wave velocities can be computed, by

a continuum elastic model, as functions of the substrate properties, the wavevector and the film properties. The

measured dispersion relations are in agreement with those computed by a model which includes a single

homogeneous equivalent film, and allow the derivation of the film properties by the fit of the computed velocities to

the measured ones. Thickness and density being independently obtained by X-rays, the film elastic constants are

measured.

For a 8 nm thick film the Young's modulus E turns out to be around 480 GPa, with a shear modulus G lying in

the 140 - 260 GPa interval. For a 4.5 nm thick film E is around 320 GPa, with G lying in the 100 - 180 GPa interval;

these lower values' are probably due to the higher influence of the lower density close to the interfaces. Work is in

progress to obtain the elastic constants of films of thickness down to 2nm. These results show that carbon films can

be grown which have a thickness of few nanometers and still have the high density and the peculiar mechanical

properties of tetrahedral amorphous carbon. Such films are thus ideal candidates for the magnetic hard disk coating.

NASA/CP--2001-210948 521 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

STRUCTURE ANALYSIS OF TETRAHEDRAL AMORPHOUS CARBON FILMS

USING SYNCHROTRON RADIATION LIGHT SOURCE

Churl-Seung Lee, Kwang Yong Eun, Kwang-Ryeol Lee

Thin Film Technology Research Center, Korea Institute of Science and Technology, RO.Box 131,

Cheongryang, Seoul, 130-650, Korea

Ki-Hyun Yoon

Department of Ceramic Engineering, Yonsei University, Seoul, 120-791, Korea

E. Gullickson, H. Padmore

Lawrence Berkeley National Lab., University of Berkeley, Berkeley, CA 94720, USA

Tetrahedral amorphous carbon (m-C) film has been attracted much attention due to their excellent

physical and chemical properties. For example, extremely high hardness of the film larger than 60Gpa and

outstanding tribological properties were frequently reported. Furthermore, advantages in the deposition

process such as smooth film surface and low deposition temperature, makes the film an important candidate

for high performance coating of PCB microdrills, precision dies, and high density data storage devices.

However, as in hydrogenated amorphous carbon films, the structure and properties of ta-C film is also

strongly dependent on the deposition condition. In order to understand the deposition mechanism and to

optimize the properties of ta-C fihns, analysis of atomic bond structure of the film would be a prerequisite.

In rite present work, atomic bond structure of ta-C film was analyzed by near edge x-ray absorption

fine structure (NEXAFS). Using energetic condensation of carbon ions from filtered vacuum arc plasma,

about 100nm thick ta-C films were deposited on Si (100) wafers. During the deposition, adc bias voltage

ranging from 0 to -500V was applied to obtain films with various atomic-bond structures. The residual

compressive stress of the film was obtained from the curvature of film/substrate composite using the

Stoney equation. Nanoindentation, low energy EELS, and RBS were employed to characterize the

mechanical properties and atomic-bond structures of the films. NEXAFS spectrum was obtained by using

Advanced Light Source (ALS) at Lawrence Berkeley National Lab.

In the present work, we focused on the changes in NEXAFS spectrum in various deposition condition

and post annealing processes. The relationship between peaks observed in the spectrum and the structural

change in the film was identified. Based on these results, the suggested models on the atomic bond

structure of ta-C film will be discussed.

NASA/CP--2001-210948 522 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

NANO-STRUCTURAL PROPERTIES OF CARBON MATERIALS OBTAINED

FROM ORGANIC COMPOUNDS

Mutsumasa Kyotani a'b, Chiharu Yamaguchi a'b, Akiko Goto a, Katsuhiro Sasaki a'b, Hisaji Matsui b,

a Frontier Carbon Technology Laboratory(FCT), Japan Fine Ceramics Center(JFCC)

c/o National Institute of Materials and Chemical Research, 1-1 Higashi, Tsukuba, Ibaraki 305-8565, JAPAN

b Department of Research & Development, Osaka Gas Co., Torishima, Konohana-ku, Osaka 554-0051,

JAPAN

Kazuo Akagi

Institute of Materials Science and Tsukuba Advanced Research Alliance (TARA), University of Tsukuba,

Tsukuba, Ibaraki 305-8573, JAPAN

Yoshinori Koga

dNational Institute of Materials and Chemical Institute, Tsukuba, Ibaraki 305-8565

Corresponding Author: Mutsumasa Kyotani, Japan Fine Ceramics Center(JFCC), c/o National Institute of

Materials and Chemical Research, 1-1 Higashi, Tsukuba, Ibaraki 305-8565

Tel +81-298-61-9981, Fax +81-298-61-4796, e-mail; L__kyg_t'_a,nj@_h_(zI_n_e:._n_Ln_t,__g__)_._2

ABSTRACT

There are several ways of synthesizing carbon materials from organic compounds containing hydrocarbon

polymers. Nano-structural properties of carbon materials obtained from organic compounds might be

expected to be different from that of carbon ones from other ways such as a discharge and a chemical vapor

deposition (CVD), because carbon materials are usually produced from organic compounds at lower

temperatures. We have tried to synthesize carbon materials from some organic compounds with an

electrochemical reaction and a pyrolytic one. In the case of an electrochemical reaction, carbon specimens

were almost amorphous because the reaction temperature was room temperature. Electrochemical carbons

were heated at temperatures below 1000 °C to be crystallized. A pyrolytic reaction was also carried out at

temperatures below 1000 °C. Various shaped graphite structures in nanometer size were formed, depending

on chemical structures of original organic compounds and on conditions of carbonization. The formation

mechanism of nano-structures will be discussed.

Keywords: Carbonization, Pyrolysis, Electrochemical reaction, Organic Compounds,

Nano-graphite structure

Introduction

Organic compounds are often used as a precursor for carbonization(ref.1). There are several ways of

synthesizing elemental carbon from organic compounds containing organic polymers. Nano-structural

NASA/CP--2001-210948 523 August 6, 2001
propertiesofcarbonmaterialsobtainedfromorganicprecursors mightbeexpectedtobedifferent fromthatof

carbononesfromotherwayssuchasadischarge(ref.2) andalaserabrasion, becausecarbonmaterialsare

usuallyproducedfromorganicprecursors atlowertemperatures(ref.3). Physicalndchemical propertiesof

carbonmaterialsarebelievedtodependmainlyonsolidstructures ofthematerials.

Electrochemical preparationf elemental carbonhasbeenstudiedconsiderably inthelasttwodecades

mainlyinaviewpointofacademic(ref.4). Inmostcase,areactionofelectrochemical carbonization iscarried

outatroomtemperature(refs.4,5). Ontheotherhand,pyroliticcarbonizations well-knowninbothview

pointsofacademicandpractical(refsl,6). Mostorganiccompounds arecarbonizedattemperatures higher

than500°Cinaninertgasorinvacuum. Therearefewstudiesonnano-structural propertiesof carbon

materialssynthesized by meansof bothelectrochemical rectionandpyroliticone.Wehavetriedto

synthesizeelemental carbonfromfluorinatedaromaticcompounds andlinearchain-likepolymersbymeans

ofanelectrochemical rectionandapyroliticone.Structural propertiesof thesynthesized carbonmaterials

havebeenexamined.

Experimental

Precursors used in this work are fluorinated aromatic pitchs of naphthalene, methyl-naphthalene, anthracene,

and quinoline, and linear chain-like polymers, poly(tetrafluoroethylene)(PTFE) and polyacetylene(PA). The

fluorination of aromatic pitchs was performed by Osaka Gas Co., in Japan. A thin film of PA was polymerized

from acetylene gas using Ziegler-Natta catalysis. Fluorinated aromatic pitchs and PTFE were defluorinated

and carbonized by means of a reaction of electrochemical reduction at 25 °C in a solvent of

tetrahydrofuran(refs.7,8). The method of electrochemical reaction was described in a previous paper.

Carbonization of PA thin films was performed by a pyrolitic reaction at a temperature in the range of 500 °C

to 1000 °C in vacuum. Structures of the synthesized carbon materials were investigated using Raman

scattering spectroscopy, x-ray diffraction measurements, and scanning and transmission electron

microscopy(SEM and TEM, respectively). All of the specimens of synthesized carbon materials were very

small black particles or powders. In order to prepare specimens for observation of TEM, these carbon

materials were suspended in a solvent of ethanol.

Results and Discussion

Fig.1 illustrates Raman scattering spectra of carbon specimens defluorinated by means of the

electrochemical reduction. The precursors of the carbon specimens in Fig.1 are fluorinated anthracene,

quinoline and methyl naphthalene pitchs. All carbon materials have peaks of D-band(ca 1350cm -1) and

G-band(ca 1580cm -1) in the Raman spectra, while the peaks are very small in the case of quinoline pitch, as

seen in Fig. 1. In the case of a carbon specimen obtained from fluorinated naphthalene pitch, however, a very

small peak appears at 2100cm -1 together with both D- and G-bands in a Raman spectrum, as reported in a

previous paper(ref.8). This small peak was assigned to carbyne structure. From the results of x-ray

diffractions and TEM observation, these carbon specimens had little crystalline order. Probably, the reason

why these carbon materials are almost an amorphous state is that a temperature for the electrochemical

carbonization is too low to grow a graphite crystal.

Raman scattering spectrum of a carbon specimen annealed at 800 °C for 60min under an atmosphere of

nitrogen gas is shown in Fig.2. This Raman spectrum is very similar to that in Fig. 1. A crystalline graphite

was, however, recognized in this annealed carbon by x-ray diffraction patterns and TEM observation. Fig.3

NASA/CP--2001-210948 524 August 6, 2001
showsahighresolutionTEMshowingalatticeimagehavingadistanceof0.34nmassignedto(002)faceofa

graphitecrystal.Theseresultsrevealthatagraphitecrystalinnano-sizeisabletogrowfromamorphous

carbonwithannealingat800°C.

Wehavealreadyreportedthatamorphous carbonobtainedbythesamelectrochemical reductionasthatin

thisworkiscrystallizedintheformof aribbon-shaped graphitecrystalinnano-sizebyatreatment of an

electronbeamradiationat800°C.Therefore, theresultsmentionedabovesuggest thataheat-treatment at

800°C(ref.9)isabletopromoteamorphous carbontogrowanano-sizegraphitecrystal,butnotnano-ribbon

graphite.

Acarbonmaterial obtainedfromPTFEbytheelectrochemical reductionwasalsoalmostamorphous state.

It shouldbenotedthatRamanscatteringspectrumofthiscarbonshowsapeakatashiftwavelengthinthe

rangeof2000to2100cm-1whichisattributedtoatriplebondofcarbon/carbon, socalled,carbynestructure.

Theamorphous carboncontainingcarbynestructurewasalsoableto graphitizewithanelectronbeam

radiationat800°C.

Ontheotherhand,varioushaped-graphite structureswereobservedin carbonmaterialsobtainedby

pyrolisisofPAatatemperature of500-1000°C.Anexampleofgraphitestructures ofpyroliticarbonfrom

PAis exhibitedin Fig.4,showingasemi-sphere-like graphitecrystal.Theorderof thegraphitecrystal,

however, isnotsomuchperfectbecausethe(002)faceisbent,asseeninFig.4.ThisresultrevealsthatPAis

abletographitizeinvariousformsatatemperature below1000°C.

Anthracene

Methylnaphthalene

'| 579.8

1579.8

1586.4

2400 2200 2000 1800 1500 1400 1200 1000

RarnanShifl[cm-ll

Fig.1 Raman scattering spectra of carbons obtained by electrochemical defluorination of

fluorinated anthracene, methyl naphthalene, and quinoline pitchs as indicated in the figure.

NASA/CP--2001-210948 525 August 6, 2001
5

............................................ :-2........................

,r ,

, ? \

/

/

/

/

r

20_W

'\

Fig.2 Ram_.q ,_c:itlcxi_tF _pc_;h_ of cd_b_ vt_t_!.It;c:, by eJ_cLro_hemJ_al d_,?h!eliuadc_l ci"

:::::::::::::::::::::::::::::::::::::::::::::::::: " ::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::: iiiiiiii ii iii i iiii i iiiiiiiii iii}iiiiiii 

iiii !i !!!!! !!i ! ii! i !ii i ii !i!i iiii!iiiiiiii!ii i ii iiii!i !!! !!!!!!! iii iiiii iii !i 

...................................................................... .o_o_,ooo°o

!!ii!i!!iii!!iii!!iiiii!!i!iiii!iiii!i!iii!ii_iii_i!iii!iiiii!i!!iii!i!iiiii!iiiiiiii!_iii!i!iii!ii!iiii!iiiiiiii!!i!i!iiii!!i_!iiii!iii_!iiii!iii!_

::::::::::::::::::::::::::::::::::::::_i_i_i_i:i_i_i_!_!_i_!_i_i_i_i_i_i_i_i_!_b.m_!_i_i_i_i_i_i_i_i_i_i_i_i_i_i_i_i_i_i_!_!_i_i_!_!_!_i

Fig.3 "tIR'I.*EM of carbon obtained by _leclrechemical deflmniaatien of :['hmdoated ,:al_hihalenc

pilch, a,M theft annealed at 800,'(_; for 60mb_ under niirogen gas atm{_spbere. A lattice image

of (002) face of graphite structure can be seen.

Fig,4 f_-RTEM of carbon obtaii_ed by pyrotisis ot polyacetylene thin film a_ 1A)0{)_C for filtmin

Semi-s_3he]ical a]ld hollow smmtm e {ff g_aphite cm_ be seen.

NASA/CP--2001-210948 526 August 6, 2001
Acknowledgments

This work is part of the national research project, Frontier Carbon Technology(FCT) sponsored by the New

Energy and Industrial Technology Development Organization (NEDO).

References

1. Jenkins G.M, Kawamura K.,: Polymeric carbons-carbon fiber, glass and char. Cambridge University Press,

London, UK, 1976.

2. Saito Y., Carbon Nanotubes Produced by Arc Discharge. New Dia. & Fron.Carbon Tech., Vol.9, No.l, 1999,

1-30.

3. Yudasaka M., Kikichi R.,: Supercarbon-Synthesis, Properties, and Applications., Yoshimura S., Chang

R.RH., Eds., Springer, Berlin, 1998, 99-105.

4. Kavan L., Electrochemical Carbon. Chem.Rev.,Vol 97, 1997, 3061-3082.

5. Kastner J, Kuzmany H., Kavan L., Dousek F.R, Kutti J., Reductive Preparation of Carbyne with High Yield.

Macromolecules., Vol 28, 1995, 344-353.

6. Kaburagi Y., Hishiyama Y., Highly Crystallized Graphite Films Prepared by High-Temperature Heat

Treatment from Carbonized Aromatic Polyimide Films. Carbon, Vol 33, No.6, 1995, 773-777.

7. Kawase N., Yasuda A., Matsui T., Yamaguchi C., Matsui H., Carbyne-like carbon electrochemically

prepared with sacrificial anode. Carbon, Vol 36, 1998, 1234-1235.

8. Yamaguchi C., Kyotani K., Murase H, Nishino H., Matsui H., Koga Y., Fujiwara S. Preparation of Polyyne

Containing Carbonsby means of Electrochemical Reduction of Fluorinated Cyclic Carbons. Proceedings of

ADC/FCT'99, NEDO, 1999, 567-570.

9. Kyotani H., Yamaguchi C., Goto A., Sasaki K., Matsui H., Koga Y., Structure Formation of Nanoribbon

Graphite from Carbyne-like Carbons. Synthetic Metals, in press.

NASA/CP--2001-210948 527 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CARBONIZATION OF POLYACETYLENE BY PYROLYSIS

• 1 1 • • 12

Akiko Goto 1,Mutsumasa Kyotanl, Kazuo Tsugawa, Yoshlnorl Koga '

1FCT Research Department, Japan Fine Ceramics Center, c/o Research Center for Advanced Carbon Materials,

AIST Tsukuba Central 5, 1-1 Higashi, Tsukuba, Ibaraki 305-8565, Japan

2Research Center for Advanced Carbon Materials, AIST Tsukuba Central 5, 1-1Higashi, Tsukuba, Ibaraki 305-8565

Kazuo Akagi, Guangzhe Piao

Institute of Materials Science, Center for TARA, University of Tsukuba, Tsukuba, Ibaraki 305-8573

ABSTRACT

Polymeric materials are often used as a precursor for synthesized carbon materials. The pyrolytic reaction at

high temperatures is effective to carbonize polymeric materials. Polyacetylene (PA) is a linear hydrocarbon polymer

and has a simple conjugated molecular structure containing sp 2 hybridized orbital. The purpose of this study is to

carbonize PA with a pyrolytic reaction and to investigate structures of carbon material obtained by the reaction. PA

thin film was synthesized using Ziegler-Natta catalyst and then was washed in a solvent of toluene. Carbonization of

PA was carried out by a pyrolytic reaction at various temperatures below 1000°C for one hour under a vacuum of

100Pa. The dehydrogenation of PA was confirmed with an elemental analysis. Structural properties of the carbon

material obtained by the pyrolysis were investigated using both scanning and transmission electron microscopy

(SEM and TEM, respectively), Raman spectroscopy, and x-ray diffraction measurement. Two peaks appeared at

wavenumbers of 1350cm -1 and 1580cm -1 in a Raman spectrum of the carbon material. A selected area electron

diffraction showed a strong reflection having a distance of 0.34nm. TEM observation showed that the carbon

material has various shaped graphites in nanometer size and also an amorphous structure.

Formation of carbon structures obtained by the pyrolysis of PA will be discussed.

Keywords: carbonization, Polyacetylene, pyrolysis, nanosize graphite

INTRODUCTION

It is well known that carbonization with pyrolytic reactions of organic compounds containing organic polymers is

one of methods for synthesizing carbon materials. Many organic compounds have been tried to carbonize using

pyrolytic reactions at high temperatures under a certain atmosphere such as an inert gas (ref. 1).

Polymeric materials are often used as a precursor to be carbonized by means of pyrolytic reaction (ref. 1). In

particular, a graphite structure has been formed by pyrolysis of some aromatic polymers such as polyimide (ref.2,3)

and poly(p-phenylene vinylene) (ref.4). After graphitization of these polymers, highly oriented and well-crystallized

graphite have grown with a heat treatment at a temperature of above 2000°C in an inert gas (ref.3,4). There is,

however, few paper of carbonization from a hydrocarbon polymer having a linear carbon bond structure. Fortunately,

PA has the same sp: bond in the molecular structure as aromatic polymers. Therefore, it is able to expect that a

graphitic structure will be yielded by carbonization with pyrolysis of PA. Chuvyrov et al have performed

carbonization of PA at 627 K for 6 hours and then graphitization at above 1227 K. They have reported only an

electron diffraction pattern as a structural assignment of graphite (ref.5).

We have tried to carbonize a thin film of polyacetylene by a pyrolytic reaction at a lower temperature, and to

investigate the structure of the carbons obtained by that reaction.

EXPERIMENTAL

PA thin film was synthesized using Ziegler-Natta catalyst and then was washed in a solvent of toluene. Thickness

of the PA thin film was about 100g m. A tubular furnace containing a quartz glass tube of 50mm in diameter was

used for pyrolysis. The PA thin film left on a quartz glass boat was put in the quartz glass tube. The pyrolytic

reaction was performed at a temperature in the range of 500°C to 1000°C for 60min in a vacuum of 100Pa. A

heating rate was 8°C/min. The furnace was cooled slowly to room temperature after the pyrolysis. A small amount

NASA/CP--2001-210948 528 August 6, 2001
ofblackmaterial canbefoundonthequartzboatafterthepyrolysis. Moreover, innerwallsofthequartzglasstube

inthefurnacebecameblack.Thisresultmeansthatalargeamount ofcarbonizedPAwasdispersedinalldirections.

Theblackmaterial onthequartzglassboatisapartofpyrolyzedPA.Theweightandtheelemental analysis(Perkin

Elmer2410II)ofthepyrolyzedPAontheboatwasexamined. Structural propertiesofcarbonmaterialsobtainedby

pyrolysisofthePAthinfilmwereinvestigated usingaRamanscatteringspectroscope (JascoNRS-2100), x-ray

diffractionmeasurements (Rigaku),SEM(HitachiS-5000), andTEM(Jeol2000FXII). A lightsourceofRaman

scatteringwasawavelength of514.5nmofanargon-ionlaser.

RESULTS AND DISCUSSION

Table. 1 and Figure. 1 show changes in the weight and the elemental composition of the pyrolyzed PA on the

quartz glass boat with pyrolysis temperature. A degree of dehydrogenation of PA thin film increases with increasing

pyrolysis temperature. In particular, the content of hydrogen is below lwt% at pyrolysis temperatures above 800 K.

However, an yield, which means a percentage in weight of the pyrolyzed PA on the quartz glass boat to an original

PA thin film, decreases with increasing pyrolysis temperature, as can be seen in Table 1 and Figure. 1. These results

reveal that carbonization of PA is almost finished at temperatures above 800 K and that the content of carbonized PA

which is dispersed during pyrolysis increases with increasing pyrolysis temperature. The weight loss of about 85wt%

after the pyrolysis may be due to the many kinds of the product boiled during heating and keeping at each

temperature.

Raman scattering spectra of the pyrolyzed PA on the quartz glass boat are shown in Figure.2. Two large peaks at

about ll50cm -1 and at about 1500cm -1 appear for the original PA thin film. This spectrum is almost the same as the

typical one of trans PA (ref.6). On the other hand, the pyrolyzed PA at different temperatures has two large peaks at

about 1350cm -1 and at about 1580cm -1, while the width of the peaks depends on the pyrolysis temperature. The peak

at 1350cm -1 can be assigned to the D-band originated from an amorphous carbon structure. Another peak at

1580cm -1can be assigned to the G-band originated from a graphite structure.

A lot of fibrous structures were observed for the original PA thin film by SEM. After the pyrolysis reaction,

however, these fibrous structures disappeared. From the results of TEM observations, various shaped graphite

structures in nanometer size such as a hollow sphere and a hollow ellipse were observed with an amorphous structure

for carbon specimens obtained by the pyrolysis of PA. Figure.3 and Figure.4 show TEM images of a carbon

specimen pyrolyzed at 1000 K with a selected area electron diffraction pattern (SAED)(Figure.4) which illustrates a

ring reflection having a spacing of 0.34nm. This spacing is attributed to that of the (002) face of a graphite crystal.

ACKNOWLEDGEMENT

This work is a part of the Frontier Carbon Technology Project supported by NEDO.

Table 1.Changes in the weight and the elemental composition of the pyrolyzed PA on the quartz glass boat

with the temperature.

Temperature of

experiment [K]

Weight of the

original PA thin film

[g]

Weight of the

pyrolyzed PA left on

the quartz boat [g]

Yield [%]

Results of elemental analysis [wt%]

C H Others

500 0.1603 0.0343 21.4 80.55 3.21 16.24

600 0.1519 0.0284 19.7 88.38 2.45 9.17

800 0.1579 0.0272 17.2 93.19 0.66 6.15

1000 0.1550 0.0240 15.5 94.12 0.16 5.72

Calculated hydrogen content of PA is 7.6 _wt%.

NASA/CP--2001-210948 529 August 6, 2001
-v

>.

25

2O

15

10

m

"_H content

"_"..., .....

8

7

5 0

600 800 1000 1200 1400

Temperature K

6

5 _

t43

cO

3 o

2

Figure 1.Changes in the yield and H content of pyrolyzed PA with pyrolysis temperature.

2000 ! 800 1600 1400 1200 1000

Wavenumber cm -1

Figure 2. Raman scattering spectra of PA thin film and pyrolyzed PA at defferent temperatures. The

pyrolyzed temperature is included in the figure.

NASA/CP--2001-210948 530 August 6, 2001
Figure 3. TEM image of a h.Itow sphere graphite structure.

...... T ::::::::::::::::::::::::::::::

20nm

Figure 4. TEM image of a holh)w ellipse graphite structure.

NASA/CP--2001-210948 531 August 6, 2001
REFERENCES

1. Jenkins G.M.; Kawamura K.: Polymeric carbons-carbon fibre, glass and char. Cambridge University Press,

Cambridge, 1976

2. Kaburagi Y.; Hishiyama Y.: Highly crystallized graphite films prepared by high-temperature heat treatment

from carbonized aromatic polyimide films. Carbon, vol. 33, no. 6, 1995, 773-777

3. Ohnishi T.; Murase I.; Noguchi T.; Hirooka M.: Preparation of graphite film by pyrolysis of polymers.

Synth. Met., 18, 1987, 497-502

4. Ohnishi T.; Murase I.; Noguchi T.; Hirooka M.: Highly conductive graphite film prepared from pyrolysis of

poly(p-phenylene vinylene). Synth. Met., 14, 1986, 207-213

5. Chuvyrov A.N., et al.: Structure and conductivity of new carbonaceous films obtained by polyacetylene

pyrolysis. Synth. Met., 41-43, 1991, 1443-1446

6. Arbuckle G.A., et al.: Spectroscopic properties of polyacetylenes synthesized via three modifications of

Ziegler-Natta catalytic system. Synth. Met., 79, 1996, 183-188

NASA/CP--2001-210948 532 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

A FLORID CARBON STRUCTURE

Y. K. Yap, M. Yamaoka, M. Yoshimura, Y. Mori and T. Sasaki

Department of Electrical Engineering

Osaka University, 2-1 Yamadaoka, Suita, Osaka 565-0871, Japan.

Tel: +816-6879-7707 Fax: +816-6879-7708

Email: yap(o_>ssk.pwr.eng.csaka-.a.ac.jp

ABSTRACT

In this work, we report on a flower-like carbon structure: florid carbon. Florid carbon structures were grown on

Si substrate that pretreated with Fe(N03)3 solution. These structures are grown by a RF-plasma assisted pulsed-laser

deposition (PLD) technique at 800 °C with pure oxygen as the working gas. The florid carbon has 2 to 10 branches

and appeared like a flower as viewed from electron and optical microscopes. All these branches are originated from

the center of the florid structure and spread outwards in a random direction. The overall structure is 40-60 gm in

diameter with each branches a few gm in width and 20 gm in length. The formation of these florid structures

required oxygen atmosphere, Fe203 particles and ion bombardment that induced by the substrate bias voltages.

According to Raman spectroscopy, the florid carbon structures are constructed from high-order graphitic

microstructure.

Keywords: Carbon, Fe203, RF-plasma, pulsed-laser deposition

INTRODUCTION

The carbon atoms allow sp, sp 2 and sp 3 hybridization and construct many materials like graphite, diamond,

diamond-like carbon films, and carbon nanotubes. It is interesting to prepare carbon materials by non-conventional

synthesis techniques like those involved transition metal catalysts. An outstanding example from such kind of

approach has been the carbon nanotubes (CNTs) (ref. 1). The CNTs, both single wall and multi-walled can be

prepared by arc discharge evaporation method. Catalysts like Fe, Co and Ni were shown to promote the productivity

of the single all CNTs (refs. 2 to 4). Further, synthesis of aligned multi-walled CNTs was recently established by

using various types of chemical vapor deposition techniques by using transition metal catalysts (refs. 5 to 7). In this

work, we report on a flower-like carbon structure: florid carbon.

Florid carbon structures were grown on treated Si substrates. The florid carbon has 2 to 10 branches and

appeared like a flower under the images of field emission scanning electron microscopy (FESEM). All these

branches are originated from the center of the florid structure and spread outwards in a random direction. All

branches are ended with a dome-like cap that appeared in a brighter image as compared to other parts of the florid

carbon. The overall structure is 40-60 gm in diameter with each branches a few gm in width and 20 gm in length.

The synthesis and structural properties of florid carbon are described in this paper. Results from FESEM, energy

dispersive X-ray (EDX), X-ray image and micro-Raman spectroscopy are discussed.

EXPERIMENTAL DETAILS

The formation of these florid structures required ion bombardment and substrate pretreatment by Fe203 particles.

The florid carbon was synthesized by means of RF-plasma assisted pulsed-laser deposition. Our deposition system

consist of a solid-state UV laser (5c0, fifth harmonic generation of Nd:YAG lasers at a wavelength of 213 nm) with

pulse duration of 3 ns (ref. 8). The RF plasma coupled to a Mo substrate holder induced negative self-bias voltage on

the substrate (Vb) and initiated ion bombardment on the growth surface. A laser intensity of 0.5 GWcm -2 was focused

on the rotating graphite target (99.99 at. %). Similar approach was used to prepare carbon nitride films with

predominantly sp3C-N bonds (ref. 9). The carbon plume generated from such ablation is the only carbon source for

the deposition on the treated Si substrates attached on a heater 4 cm away from the target.

NASA/CP--2001-210948 533 August 6, 2001
ThesubstrateusedfordepositionwascoatedbyFe(N03)3solutioniethanol anddriedinairatRTbeforebeing

mountedintothevacuumchamber. All thesesubstrates werethentreatedunderthe02RF-plasmaandVbof- 400V

forabout15minutesat800°C.AllFe(N03)3particlesontheSiweredissociated intoFe203atsubstratetemperatures

above-47°C.Aftertheplasmatreatment, laserablationwascarriedoutat800°Cfor2hoursunderVbof-300V.

RESULTS AND DISCUSSION

All deposited samples were examined by both optical and electron microscope. At low magnification view,

clusters of Fe203 can be observed within a size of < 10 gm. These particles are irregular in shapes. We were

surprised to observe some flower-like structures as we increased the magnification of microscopes. The flower-like

structures as view from FESEM are shown in Figure 1. As shown, the flower-like structure has about 10 branches

spread outwards from the center of the structure in a random direction. All branches are ended with a dome-like cap

that appeared in a brighter image. A magnified view at the tips of the branches is also shown together. These tips

appeared to be thicker as compared to the branches. These structures cannot be formed when Fe203 particles are

absent, or when no bias voltage is applied. Oxygen ambience is also necessary since formation of florid carbon was

failed when hydrogen or Ar were used as the working gas. Note that, under the synthesis condition of these structures,

no carbon films can be deposited on the untreated Si substrates due to the etching of oxygen plasma.

Figure 1. Typical appearance of the flower-like structures as viewed from FESEM (top). A magnified view at

the tips of the branches is shown at the bottom.

NASA/CP--2001-210948 534 August 6, 2001
Thecomposition ofthesestructureswasanalyzedbyusingtheenergydispersiveX-ray(EDX).Becauseofthe

irregularstructureandthelowatomicmassofthecarbon, ordinaryEDXisdifficulttomakeconclusiveinformation

aboutthestructures. WethenusedtheX-rayKc_imagingtechnique(EDX-option) forourpurpose. Theresultsare

showniFigure2.Imagesasshown(fromlefttoright,toptobottom)areSEMimageofthe

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:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::: :.:.:.:-:.:.:.:-:-x.z.x..x.:.:.:.:-:.x.:.x.:.:.x.x.:.:.:.:.:-:.:.:.:.:.:

.:.:.:.:.:.:.x.:_,x.:.:.:.:.:.:.:.:.:.x.:.x.x.:.x.x.:.:.:.:.:.x.:.:.:.

iiiiiii_.@.._...:._!_i_iiiiiiiiiii_iiii_i_ii!iiiiiii!iii!iiii_!i_i_iiiiii!i_iii!ii

_:!iii!i_i_i!_!i_!_!:i:i:i:i:i:i:i:i:i:i:i:iiiiiii_iii_i!:_!iiiiii!i!iiiii!i_ii:i:!

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::: ,%'o°oOo°°',o,o,O,°oO,'o'o'o'°Oo°.'._-O_O.62_1 I .°.%°,'oO,Oo'oO°*o°oO.'60°°°OoO,O.

Figure 2. SEM image of florid carbon and X-ray images of florid carbon at selected elements: Fe, Si, and C as

indicated at the upper left corners of individual images.

NASA/CP--2001-210948 535 August 6, 2001
flower-likestructures, X-rayKc_imagesofthestructures forseveral selectedelements (Fe,SiandC).TheX-ray

imageofCisalsoshownwithartificialtracingoftheoutlineofflower-likestructures.

Thepresent oftheselectedelements canbeindicatedbythedottedsignalsofrespectiveimages. Accordingto

theX-rayimageofFe,itisdifficulttomakeanyconclusionaboutthedistributionofFe.ThedottedsignalfromFe

seemstoappearuniformlyeverywithintheexaminedarea.However, thedensityofthedotsislowandthusis

assumedastheimpurityonthesurface. AsfortheSi,becauseoftheuseofSisubstrate, high-density Sisignals

(dots)areindicatedaroundtheflower-likestructures. Thepatternoftheflower-likestructures canberecognizedas

theemptyarea.Thismeans, thesignal fromtheSisubstratebelowtheflower-likestructures isshadowed bythe

deposit andthusappearedasthedot-lessarea.Wethinkthatsimilarconditionishappened forFebutisnotclearly

shownfromtheX-rayimageduetothelowcontrast betweentheareawithflower-likestructures andtheSisurface.

Thus,Fe203films/ particles are supposed to be embedded below the flower-like structures (see also discussion

about the Raman spectroscopy). Likewise, because of the thick deposition of carbon, signal from the C elements

appeared in the pattern of flower-like structures (see the trace at the last image as well). From these results, we

concluded that the flower-like structures are carbon materials: a florid carbon structure.

To supplement our result, micro-Raman spectroscopy is employed to examine the microstructure of the florid

carbon. As shown in Figure 3, a strong Raman scattering is detected at a Raman shift of 1600 cm -1. This peak is

recognized as the G band of graphite. Thus, the florid carbon is supposed to have a graphitic carbon microstructure.

Beside the G band, a shoulder is observed at a Raman shift of about 1360 cm -1. This shoulder is usually referred as

the D band of amorphous carbon. As judged from the high intensity ratio between the G and D bands (IG/ID), the

florid carbon is supposed to have high graphitic crystalline order. Note that no Raman signal can be detected from Si

surface that is not covered by Fe203 particles. This means, carbon deposits of the florid structures are stabilized on

the Fe203 without which, will be etched away by the oxygen plasma.

._.a

I,Illl,l,lll,llllllll:l,l,llll,lllll,I

,, , i I ,, ,, i ,, ,I i ,, ,I , ;,, I, , i , I, , i , I,, i

Gi DI '

1800 1700 1600 1500 1400 1300 1200 1100 1000

Raman Shift (cm -])

Figure 3. Raman spectra detected from the florid carbon.

CONCLUSION

We have described on the synthesis of a flower-like carbon structure. The formation of these florid carbon

required Fe203 particles and ion bombardment that initiated by the RF-plasma induced negative dc substrate bias

voltages. No florid carbon can be prepared when no bias voltage is applied. Further, florid carbon cannot be form

when hydrogen or Ar gases were used for deposition. According to micro-Raman spectroscopy, the florid carbon is

supposed to be constructed from high order graphitic carbon.

NASA/CP--2001-210948 536 August 6, 2001
REFERENCES

1. Ijima S.: Helical microtubules of graphitic carbon. Nature 354, Nov. 1991, 56-58.

2. Ijima S.: Single-shell carbon nanotubes of 1-nm diameter. Nature 363, Jun. 1993, 603-605.

3. Bethune D. S.: Cobalt-catalysed growth of carbon nanotubes with single-atomic-layer walls. Nature 363,

Jun. 1993, 605-607.

4. Thess A. et al.: Crystalline ropes of metallic carbon nanotubes. Science 273, Jul. 1996, 483-487.

5. Li W. Z., et al.: Large-scale synthesis of aligned carbon nanotubes. Science 274, Dec. 1996, 1701-1703.

6. Ren. Z. E, et al.: Synthesis of large arrays of well aligned carbon nanotubes on glass. Science 281, Nov.

1998, 1105-1107.

7. Tsai. S. H., et al.: Bias-enhanced nucleation and growth of the aligned carbon nanotubes with open ends

under microwave plasma synthesis. Appl. Phys. Letts. 74, June 1999, 3462-3464.

8. Yap Y. K. et al. High-power fourth- and fifth-harmonic generation of a Nd:YAG laser by means of a

CsLiB_O10. Opt. Letts. 21, Sept. 1996, 1348-1350.

9. Yap Y. K. et al.: Influence of negative dc bias voltage on structural transformation of carbon nitride at

600 °C. Appl. Phys. Letts. 73, Aug. 1998.915-917.

NASA/CP--2001-210948 537 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

SOLID LUBRICANT PROPERTIES OF CARBON ONIONS

PREPARED FROM DIAMOND CLUSTERS AND PARTICLES

A. HIRATA, M. IGARASHI and T. KAITO

Department of Mechanical Sciences and Engineering, Tokyo Institute of Technology

2-12-1, O-okayama, Meguro-ku, Tokyo 152-8552, JAPAN

ABSTRACT

We have investigated tribological properties of carbon onions prepared by heat treatment of diamond clusters

or particles. A new type of preparation technique was proposed in order to obtain large amount of carbon onions in

short time. Ten milligrams of diamond clusters used as source material were heated with an infrared radiation

furnace up to approximately 1730°C in argon ambient. As a result, heating for one minute at 1730°C transformed

diamond clusters into carbon onions. Near-round and multi-layer structure of the carbon onions was observed and

confirmed by high resolution TEM. The particle size of these carbon onions ranged from 5 to 10 nm that

corresponds to the size of diamond clusters. This preparation technique also applied to diamond particles less than

0.5 gm in diameter to produce larger carbon onions. Tribological properties of the carbon onions were examined by

ball-on-disk type friction testing using a silicon wafer and a bearing steel ball. The carbon onions were distributed

on the silicon wafer without adhesive. The carbon onions exhibited lower and stable friction coefficients in

comparison with graphite powders both in air and vacuum at room temperature. The wear rates of steel balls

sliding on the silicon wafer on which carbon onions are distributed were extremely lower than that sliding on the

wafer over which graphite powders spread. Moreover, it is found that the carbon onions display its superior

tribological property on harder disk materials with smoother surface. Larger carbon onions prepared from diamond

particles have maintained low friction property on rough surfaces of silicon disks.

Keywords: carbon onion, diamond, cluster, solid lubricant, tribology

INTRODUCTION

Concentric graphitic shell structured particle, which is generally called carbon onion, is one of fullerenerelated materials together with C60 and carbon nanotube. This carbon onion has several highly potential properties

suitable for solid lubricant that works as micro bearings, for example, quasi-spherical structure, weak

intermolecular bonding as well as high mechanical strength. In order to investigate tribological properties of

carbon onions, at least milligrams order of carbon onions are required. Several synthesis methods such as highenergy electron bombardment of carbon soot in TEM (ref. 1, 2), annealing of ultra-disperse diamond (ref. 3) and

carbon-ion implantation at high temperature into copper (ref. 4) have already been reported, however, these

methods have not demonstrated production of numerous carbon onions applicable to friction testing. Thus the

research of tribological properties of carbon onions has made little progress, while the tribological properties of C60

have investigated by many researchers after finding the way of production of macroscopic quantities of C60 by W.

Krfitschmer et al. in 1990 (ref. 5). This work focuses on the preparation of milligrams of carbon onions from

diamond clusters or particles by heat treatment and the study oftribological properties of the carbon onions.

PREPARATION OF CARBON ONIONS

Uniform carbon onions in size and shape are desirable for study of the tribological properties. In addition to

this, a mass preparation technique should accomplish its process in short time. In order to meet these requirements,

we have proposed a new preparation technique to obtain numerous carbon onions by infrared radiation heating

treatment of diamond clusters or particles. The transformation of diamond to carbon onions have confirmed by

TEM observation.

NASA/CP--2001-210948 538 August 6, 2001
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Figure 1. TEM observation of diamond clusters.

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Figure 2. TEM observations of carbon onions prepared from (a) diamond clusters and (b) diamond

particles less than 0.1 gm.

How to prepare numerous carbon onions

We investigated macroscopic and rapid preparation of carbon onions by heating diamond clusters by infrared

radiation in inert ambient at atmospheric pressure. Figure 1 shows TEM observation of the diamond clusters used,

of which average diameter is approximately 5 nm. Heating of diamond clusters was carried out with an image

furnace that rapidly elevates and precisely controls the temperature of sample. The experimental procedure we used

is as follows. The cylindrical holder made of graphite filled with ten milligrams of diamond clusters was placed in

the furnace. After evacuation of inside of the furnace at approximately 1.3 Pa, argon gas was introduced into the

furnace. The diamond clusters were heated in argon ambient at 1.5x10 s Pa. The highest temperature at 1730°C

was attained by about 40 minutes heating and was maintained for one minute. After that, the diamond clusters

were cooled in the argon flow. The same procedure was applied to three kinds of diamond particles less than 0.1,

0.25 and 0.5 gm except for maintaining maximum temperature for sixty minutes.

Confirmation of carbon onion formation

Diamond clusters and particles given heat treatment were observed with high resolution TEM to identify their

microscopic structure. Figure 2(a) shows that diamond clusters are transformed into quasi-spherical particles

ranging 5-10 nm in diameter with closed concentric shells. The distance between the shells is regularly

approximately 0.35 nm that closely matches the lattice parameter do02( 0.34 nm) of bulk graphite. Particles with

multi-layered structure are also formed from diamond particles less than 0.1 gm as shown in Fig. 2(b). These

observations indicate the heat treatment we proposed is effective to transform diamond clusters and particles into

carbon onions. Furthermore, size distribution of the carbon onions prepared by our method corresponds to the size

of raw material and nearly equals to the particle size of diamond clusters and particles.

NASA/CP--2001-210948 539 August 6, 2001
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Figure 3. Friction coefficient variations of carbon onions, graphite and C60/70 on Si(100) at a normal force

of 0.95 N and sliding speed of 6.3 mm/s in (a) air and (b) vacuum.

TRIBOLOGICAL PROPERTIES OF CARBON ONIONS

Tribological properties of carbon onions as solid lubricant were studied by ball-on-disk type friction testing.

Friction coefficients and wear rates were measured by the testing in air and vacuum using several sliding disk

materials with a certain range of surface roughness. These tribological properties of carbon onions were compared

to that of other carbon materials such as graphite and C6o/7o.

Experimental

Friction testing was performed normally using a bearing steel ball at 1/8 in. diameter and a (100) silicon wafer

with mirror surface. Ten milligrams of carbon onions were spread over a wafer without adhesive. Graphite

particles at average diameter of 25 gm and mixture of C6o (75-90 wt.%) and C7o particles were used as reference

materials of solid lubricant.

NASA/CP--2001-210948 540 August 6, 2001
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Figure 4. Wear marks formed on bearing steel balls after friction testing in vacuum using (a) no

lubricant, (b) graphite, (c) C_0/70,(d) carbon onions from diamond clusters and (e) carbon onions from

diamond particles less than 0.1 gm.

Sliding behavior in air and vacuum

Friction was measured in air at relative humidity of 55% and vacuum at 1.3x10 3Pa at room temperature.

Figure 3 shows friction coefficient variations as a fimction of sliding distance at a normal force of 0.95 N and

sliding speed of 6.3 mm/s. In air, friction coefficients between steel and silicon are reduced using graphite and

carbon onions. On the contrary to this, C_0/70works an ineffective lubricant as already reported by V.L. Kuznetsov

et al. (ref. 6). Carbon onions, in particular that are prepared from diamond clusters, show lower friction coefficients

under 0.1 than graphite. Figure 3(b) shows carbon onions maintain low friction property even in vacuum, whereas

graphite exhibits no lubrication ability as generally known. The low friction coefficients of carbon onions in

vacuum seem to be more stable than in air.

Figure 4 shows optical microscope observation of the contact area of bearing steel balls after friction testing in

vacuum, and it is found that extremely small wear marks for carbon onions are formed comparing to that for other

lubricants. Wear rates determined by this observation are 10 _°-10 _2 mm3/N/m for carbon onions while 10 _-10 7

mm3/N/m for graphite and C60/70.

Sliding behavior on silicon disk with a certain range of surface roughness

It is required to determine the effects of surface roughness of sliding area on the lubrication properties of

carbon onions because surface of sliding components made with metal or ceramics have a certain range of surface

roughness. Moreover, this issue is important to discuss the lubrication mechanism of carbon onions. Thus friction

was measured using silicon wafers scratched with alumina powders to give a certain range of surface roughness.

Figure 5 shows relation between surface roughness of silicon disk and friction coefficients obtained in air and

vacuum using carbon onions prepared from diamond clusters and particles. In air, carbon onions from diamond

particles exhibit friction coefficients of 0.1-0.2 from smooth surface to rough surface up to 0.3 btmRa. This low

friction property maintains on the surface with larger dimension of roughness than particle size of carbon onions,

which is believed to occur by cohesion of carbon onions in the ambient with water vapor. On the contrary to this,

friction coefficients for carbon onions from diamond clusters increases with surface roughness because the carbon

onions are smaller so that they would hide in valleys of surface and exist outside of the contact area between a ball

and a disk. In vacuum, smaller carbon onions from diamond clusters as well as particles less than 0.1 btm exhibit

high friction on rough surface because the carbon onions are released from cohesion by nonexistence of water

vapor and stay at the bottom of surface profile. Larger carbon onions from diamond particles less than 0.5 btm still

keep low friction on rough surface up to 0.3 btmRa.

NASA/CP--2001-210948 541 August 6, 2001
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Figure 5. Dependence of friction properties of carbon onions on surface roughness of sliding disks in (a)

air and (b) vacuum.

o

,m

o

o

C

0

kl.

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o6t

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Figure 6. Friction coefficient variations as a function of sliding distance at a normal force of 0.05 N and

sliding speed of 6.3 mm/s on Si, silica glass, Pyrex glass and stainless steel.

Sliding behavior on various disk materials

Frictions between a bearing steel ball and disks of silica glass, Pyrex glass and stainless steel were determined

in air when carbon onions prepared from diamond clusters were spread on the disks. Surface of the disks were

polished to finish its surface roughness below 10 nmRa where the carbon onions exhibit low friction property on

silicon disk as shown in Fig. 5(a). Friction testing was conducted at a normal force of 0.05 N and sliding speed of

6.3 mm/s. Figure 6 shows friction coefficients as a function of sliding distance on the three kinds of disks and

silicon wafer. Carbon onions show different lubrication property with disk materials, and friction coefficients

increase in the order of silicon, silica glass, Pyrex glass and stainless steel. Relation between friction coefficients

and Vickers hardness of the disks is shown in Fig.7. It is found that friction property of carbon onions depends on

hardness of sliding materials, and carbon onions exhibits lower friction coefficients on harder materials.

NASA/CP--2001-210948 542 August 6, 2001
0.8

0

cO

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kl.

0.6

0.4

0.2

Stainless steel

Pyrex glass

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0 500 10100 15100 2000

Vickers hardness

Figure 7. Friction coefficient variations as a function of Vickers hardness of sliding disks of silicon, silica

glass, Pyrex glass and stainless steel.

CONCLUSIONS

Attempts to prepare numerous carbon onions were carried out by heat treatment of diamond clusters or

particles, and tribological properties of the carbon onions were studied by ball-on-disk type friction testing. The

results obtained in this work are as follows: (1) Carbon onions similar in size to raw materials of diamond clusters

and particles are prepared by heating at around 1730°C in argon ambient at 1.5x105 Pa; (2) Carbon onions exhibit

the low friction coefficients and the wear rates of sliding materials in both air and vacuum; (3) Larger carbon

onions prepared from diamond particles maintain low friction property on rougher surfaces; and (4) Low friction

property of carbon onions occurs on harder sliding materials.

REFERENCES

1. Ugarte, D.: Curling and Closure of Graphitic Networks under Electron-beam Irradiation. Nature, vol. 359,

no. 22, Oct. 1992, 707-709.

2. Ugarte. D.: Formation Mechanism of Quasi-spherical Carbon Particles Induced by Electron Bombardment.

Chem. Phys. Lett., vol. 207, no. 4-6, May 1993, 473-479.

3. Vladimir L., et al.: Onion-like Carbon from Ultra-disperse Diamond. Chem. Phys. Lea., vol. 222, May

1994, 343-348.

4. Cabioc'h T. et al.: A New Technique for Fullerene Onion Formation. J. Mater. Sci., vol. 30, 1995, 4787-

4792.

5. Krfitschmer et al.: A New Form of Carbon. Nature, vol. 347, Sep. 1990, 354-358.

6. Blau RJ. and C.E. Haberlin: An Investigation of the Microfrictional Behavior of C60 Particles Layers on

Aluminum. Thin Solid Films, vol. 219, 1992, 129-134.

NASA/CP--2001-210948 543 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

AB-INITIO DESCRIPTION AND PREDICTION OF PROPERTIES

OF CARBON-BASED AND OTHER NON-METALLIC MATERIALS

D. Bagayoko, G. L. Zhao, and S. Hasan

Department of Physics, Southern University and A&M College

Baton Rouge, Louisiana 70813, USA

Corresponding Author: Diola Bagayoko, P. O. Box 11776, Baton Rouge, LA 70813, USA

E-mail: Bagayoko@aol.com, Telephone: (225) 771-2730, Fax: (225) 771-4341

ABSTRACT

We have resolved the long-standing problem consisting of 30%-50% theoretical underestimates of the band gaps of

non-metallic materials. We describe the Bagayoko, Zhao, and Williams (BZW) method that rigorously circumvents

the basis-set and variational effect presumed to be a cause of these underestimates. We present ab-initio,

computational results that are in agreement with experiment for diamond (C), silicon (Si), silicon carbides (3C-SiC

and 4H-SiC), and other semiconductors (GaN, BaTiO3, A1N, ZnSe, ZnO). We illustrate the predictive capability of

the BZW method in the case of the newly discovered cubic phase of silicon nitride (c-Si3N4) and of selected carbon

nanotubes [(10,0), and (8,4)]. Our conclusion underscores the inescapable need for the BZW method in ab-initio

calculations that employ a basis set in a variational approach. Current nanoscale trends amplify this need. We

estimate that the potential impact of applications of the BZW method in advancing our understanding of nonmetallic materials, in informing experiment, and particularly in guiding device design and fabrication is simply

priceless. Acknowledgments. This work was funded in part by the Depaxtment of the Navy, Office of Naval

Research (ONR), through the Timbuktu Academy (Grant No. N00014-98-1-0748), and by NASA, through

PIPELINES (Grant No. NAG5-8552 and Cooperative Agreement No. CC 13-00010).

Keywords: Prediction, Energy or Band Gaps, BZW Method, Raleigh Theorem, the Hamiltonian

INTRODUCTION

The fundamental understanding and device applications of carbon-based and other non-metals have suffered

from a pervasive inability of ab-initio, theoretical calculations to produce accurate, unoccupied energy levels or

bands. The 30%-50% underestimates of the band gaps of these materials (refs.1 to 2) are just a symptom of

presumed, theoretical errors in the unoccupied energies. The inability of these calculations to reproduce measured,

optical properties of semiconductors is another symptom. Using the Rayleigh theorem (refs. 3 to 4) and the physical

content of the Hamiltonian, we describe a basis set and variational effect (refs. 5 to 10) that is a source of errors in

calculated, unoccupied energies of atoms, molecules, clusters, semiconductors, and insulators. We present the

Bagayoko, Zhao, and Williams (BZW) method (Refs. 5 to 7) that rigorously avoids the above effect and provides a

predictive capability for variational calculations that utilize a basis set.

In the remainder of this manuscript, we first discuss the persistence of the noted underestimation problem from

1900 to 1998, along with various attempts to solve it. This background is followed by a detailed description of the

basis set and variational effect, a presentation of our general method and of the Bagayoko, Zhao, and Williams

method that resolved the problem. We then present electronic properties for diamond (C), silicon (Si), silicon

carbide (3C- and 4H-SiC), barium titanate (BaTiO3), gallium nitride (GaN), zinc oxide (ZnO), and aluminum nitride

(A1N). Unlike most previous attempts, particularly the ones utilizing local density functional potentials, our BZW

results agree with experiment not only for the unoccupied energy levels and the band gaps, but also for the optical

transitions, the dielectric function, and the effective mass. We present our predictions for the electronic and related

properties of a newly discovered cubic phase of silicon nltride (c-Si3N4) and of selected carbon nanotubes. Our

conclusion, in light of the growing trend toward the nanoscale, underscores the need for the BZW method in

calculations that employ a basis set in a variational approach of the Rayleigh-Ritz type. It also notes the unavoidable

nature of the BZW method for descriptive or predictive computations of properties of nanostructures. This

conclusion finally intimates the huge time and money savings that could result from this work--in terms of guiding

the design and fabrication of molecular, nanoscale, and other devices.

NASA/CP--2001-210948 544 August 6, 2001
THEORETICALUNDERESTIMATES OF UNOCCUPIED ENERGIES

The Problem: The Unoccupied Energy Catastrophe

"Calculated Electronic Properties of Metals," (ref. 11) has been available since 1978. In this book, Momzzi,

Jauak, and Williams provided the calculated electronic energy bands, charge densities, and related properties of

metals. Their results, obtained with density functional potentials, generally agree with experimental measurements

and with other theoretical findings. We are not aware of a similar book on semiconductors and insulators. One

reason for the absence of such a reference manual stems from the facts that from the dawn of quantum mechanics in

1900 to 1998, there existed no easy way to calculate accurately the unoccupied energy levels of bands of

semiconductors or insulators. Indeed, most density functional calculations of electronic properties of these materials,

including diamond and other carbon-based systems, led to band gaps that were 30% to 50% smaller than the

measured values (refs. 1 to 7). This band gap catastrophe is actually a symptom of a more general problem

stemming from the inaccuracy of the calculated, unoccupied energy levels or bands. Such a problem,

understandably, reverberates throughout other calculated properties of these systems, including optical transition

energies and the dielectric function. The seriousness of this problem may be best understood by noting that most

device-utilization of semiconductors or insulators involves excited states. The unoccupied band catastrophe has

therefore been a major stumbling block to efficient device design and fabrication. Even though it is not as wellknown as the band gap catastrophe, the above underestimation of unoccupied energies also affected theoretical

results for atoms, molecules, and clusters. Additionally, it is pervasive in calculated properties of nuclei in the

framework of the shell model!

In the case of semiconductors and insulators, the above problem has been ascribed to limitations of density

functional potentials, with emphasis on the local ones. These potentials (refs. 12 to 16), from the early 1970 to

present, continue to be extensively utilized in theoretical descriptions of many body systems. Other approaches to

the representation of the potential, i.e., those based on Hartree Fock formalism, did not resolve the problem and are

often accompanied by other difficulties. For instance, the limitations of the Hartree Fock method are described by

Callaway (ref. 17); they include inordinately large band widths.

Attempts to Resolve the Problem

Several approaches have been utilized in an attempt to remedy the theoretical underestimation of unoccupied

energy levels or bands. They included the use of density functional potentials that are non-local, as done in the

generalized gradient approximation (GGA) (ref. 14). Some approaches applied the "scissors approximation" to

unoccupied bands obtained with local density functional potentials (ref. 18). This approximation essentially consists

of rigidly shifting upward the calculated, unoccupied energy levels or bands (ref. 18). The self-interaction

correction (SIC) and the self-interaction with relaxation correction are other approaches (ref. 2) that have been tried

with limited success. Perhaps one of the most widely used approach that clearly goes beyond density function theory

is the dressed Green function (G) and screened Coulomb (W) interaction formalism that employs non-local, energydependent, and non-Hermitian operators. This GW method (Ref.1) is perhaps the most successful of the ones

referenced above, even though it also has its limitations. In particular, the GW method tends to give band-gaps that

are slightly larger than the experimental values. Many GW calculations are not self-consistent; the self-consistent

ones are reported (ref. 19) to increase the overestimation of the band gaps of semiconductors and insulators.

In 1998, we had obtained calculated energy bands that reproduced experimentally measured properties,

including the band gap, optical transition energies, and the dielectric function of barium titanate (refs. 5 to 6). The

extent of this agreement with experiment, for all calculated properties, warranted a search for an explanation; luck

and over-complete or incomplete basis sets were not satisfactory explanations_n light of the pervasiveness of the

agreement with experiment. Our prior experiences with calculations with several basis sets of different sizes (Ref.

9) led us to the explanation of the relative positions of the conduction bands as compared to the occupied bands. In

particular, Bagayoko (Ref. 9) had utilized the Rayleight theorem to explain the rigid shift of band energies, up or

down, depending on whether the basis size is reduced or enlarged, respectively. We identified and characterized a

basis set and variational effect that explains most of the unoccupied energy catastrophe. This effect is described

below in the next section.

An Explanation of the Problem: A Basis Set and Variational Effect

As explained in Reference 7, in a typical ab-initio, self-consistent calculation that employs the linear

combination of atomic orbitals method, electronic eigenfunctions are expanded using basis sets derived from atomic

NASA/CP--2001-210948 545 August 6, 2001
calculations (Refs.9and17).Chargedensitiesandpotentialsareconstructed using only thefimctions belonging to

occupied states (ref. 17). The Hamiltonian matrix is generated and diagonalized. Key output quantities are discrete

energy levels or continuous bands and related wave functions. The resulting output wave functions for the occupied

states are employed to generate a new charge density and the computations are repeated. The process of using the

output of iteration (n) -Jar occupied states - to construct the input for iteration (n+l) continues until self-consistency

is reached. Various measures are utilized to define self-consistency, i.e., when basic quantities, including charge

densities, potentials, eigen-energies, etc., are respectively unchanged from one iteration to the next.

The trial basis sets are found by various authors by augmenting the atomic orbitals with polarization and diffuse

orbitals (ref. 7) whose numbers, until the work of Bagayoko, Zhao and Williams (refs. 5 to 7), had no particular

limit. In fact, from a completeness standpoint, the larger these numbers the better, provided no "catastrophic sinking

of energy levels" or negative Millican population numbers occur for occupied states (ref. 20).

The above arbitrariness in the selection of the size of the basis set, by virtue of the Rayleigh theorem and the use

of the wave functions for the occupied states to generate the Hamiltonian, leads straightforwardly to the basis set

and variational effect noted above. Indeed, the Rayleigh theorem states that when an eigenvalue equation is solved

using a basis set of dimension N and one of dimension N+ 1, such that all the orbitals in the first basis are included in

the second, larger one, then the eigenvalues resulting from the calculation with (N+I) orbitals will be respectively

lower or equal to their corresponding ones from the calculation with N orbitals. Specifically, if the eigenvalues in

each calculation are ordered from the lowest to the highest, then E1N+I < E1 N, E2 N+I < E2 N, E3 N+I _<_ E3 N, etc.

Physically, these inequalities indicate that a given variational eigenvalue approaches the corresponding exact

eigenvalue from above.

Clearly, the above theorem suggests that calculations should employ basis sets that are as large as possible. This

fact and the preoccupation with ensuring the quantum mechanical completeness of the basis set partly explain the

reason that the following effect was not quantitatively circumscribed for a hundred years. (The rather vague

expression of "basis set effects" is not to be confused with the basis set and variational effect as defined and

circumscribed herein.) A fact that places a limit on the size of basis sets in variational calculations, however, stems

from the utilization of the wave fimctions of the occupied states to generate the charge density, the potential, and

the Hamiltonian matrix. This matrix is contains the physics of the calculations. We define a minimal basis set as one

that contains just enough orbitals to account for all the electrons in the system under study. Beginning with a

minimal basis set, if one performs several calculations with basis sets that are each obtained by augmenting the

previous, smaller one, then one is expected to reach a size such that the charge density, the potential, and the

Hamiltonian no longer change. One does. Past this point, the use of larger basis sets, by virtue of the Rayleigh

theorem, will still lead to a lowering of some unoccupied energy levels or bands, even though the occupied ones no

longer change. There lies the basis set and variational effect, i.e., the extra lowering of some unoccupied energy

levels of ban&; by virtue of the Rayleigh theorem (reJ_'. 3 to 4 and 9), when unnecessarily large basis sets are

employed in self consistent, variational calculations of the Rayleigh-Ritz type.

To facilitate the forthcoming presentation of the Bagayoko, Zhao, and Williams (BZW) method for avoiding

this effect, we define the optimal basis set of dimension N as that of the first of the calculations, with increasing

basis sets, such that the next calculation produces the same charge density, potential, Hamiltonian, and occupied

energy levels or bands. This optimal basis set defines a new form of convergence, i.e., that of the Hamiltonian and

of the occupied energy levels or bands vis avis the dimension of the basis set. Reference 7 showed that calculations

with basis sets larger than the optimal set, if otherwise done correctly, do not change the charge density, the

potential, the Hamiltonian, or the occupied energies. They do continue indefinitely to lower some unoccupied

energies, given that the fundamental theorem of algebra dictates the appearance of new eigenvalues upon the

increase of the size of the basis set or of the dimension of the Hamiltonian matrix.

METHOD

General Method: the Potential and the Formalism

The general method employed in our calculations has been extensively described in the literature (Ref. 5 to

10). Our self-consistent, computational formalism consists of the linear combination of atomic orbitals (LCAO).

We employed the electronic structure package of the group of Professor Bruce Harmon at the Department of Energy

(DOE) Laboratory at Iowa State University (Ref. 21 to 22). We utilized Gaussian orbitals and therefore refer to our

NASA/CP--2001-210948 546 August 6, 2001
approachasalinearcombinationofGaussianorbitals(LCGO).Indescribingmanybodysystem,thechoiceofthe

potential isacriticalone.Inallthecalculations dicussed below,weusedtheCeperley-Alder type(ref.15)oflocal

densitypotentialsaparameterized byVosko,Wilk,andNusair(ref.16).

TheBagayoko, Zhao, and Williams (BZW) Method

We introduced this method in 1998 for completely circumventing, from first-principle, the above identified basis set

and variational effect. The BZW method basically requires a minimum of three self-consistent calculations that

utilize basis sets of different sizes. It begins with the minimum basis set, i.e., the basis set needed to account for all

the electrons of the atomic or ionic species that are present in the molecule, cluster, or solid under study. Completely

self-consistent calculations are carried out. For the second calculation, the minimal basis set is augmented with one

or more atomic orbitals that belong to the next and lowest lying energy levels in the atomic or ionic species. The

self-consistent bands from calculations I and II are compared, graphically and numerically. In general, there will be

qualitative (shape and branching) and quantitative (numerical values) differences between the occupied bands from

calculations I and II. A third calculation is performed, using the basis set for calculation II as augmented with

orbitals representing the next lowest-lying atomic energy levels.

The above process is continued until the comparison of the occupied energy levels leads to no qualitative or

quantitative difference. When the results (i.e., occupied energies only) from calculation N and those from

calculation (N+I) agree within the computational error, then the optimum basis set is that of calculation N. This

selection rests on the fact that a lowering of unoccupied levels, as in calculation (N+I), after the occupied levels

converged with respect to the size of the basis set, as in calculation N, is dictated by the Rayleigh theorem and does

not necessarily reflect the physics of the system. Such a lowering may not be the manifestation of any fundamental

interaction, but rather the expression of the basis set and variational effect identified above. In fact, Reference 7

showed that calculation (N+I) and others with larger basis sets do not change the self-consistent charge density or

potential from their values as obtained in Calculation N that utilized the optimal basis set. In multi-species systems

like GaN, two or more orbitals may be added at a time if the affected atomic levels are close in energy. While we

utilized the occupied energy bands for the determination of convergence with respect to the size of the basis set,

other parameters could be employed. They may include the charge density for occupied states and particularly the

potential.

RESULTS: CALCULATED ELECTRONIC PROPERTIES OF NON-METALS

We list in the table below the calculated band gaps of diamond, 3C- and 4H-SiC, and of several other

semiconductors. Space limitation does not allow us to delve into the details of these results and of many others that

are available in our previous publications (refs. 6 to 10). For each material listed below, the reader should refer to

these sources (ref. 6 to 10 and 23 to 25) for details. Other results in these papers include the electronic energy bands,

the density of states, effective masses, charge transfers, contour plots of charge densities, and optical transition

energies. For BaTiO3 and 3C-SiC, we provided the dielectric function and the total energy curves, respectively.

The verified theoretical predictions of the group of Dr. Carter T. White (refs. 26 to 27) at the Naval

Research Laboratories (NRL) on properties of carbon nanotubes partly inspired our interest in these materials. We

report below the band gaps for carbon nanotubes (10,0) and (8,4). Carbon nanotubes (10,0) belong to a symmorphic

0

group. Its diameter is 7.83 A and there are 40 atoms in the unit-cell. Carbon nanotube (8,4) belongs to a

0

nonsymmorphic group. The diameter of (8,4) nanotube is 8.29 A. The chiral angle of (8, 4) relative to the zigzag

direction is 19.10. There are 112 atoms in its unit-cell. While (10,0) has a direct band gap of 0.95 eV at the gamma

point, the 0.9 eV gap of (8,4) is off the F-point. The highest energy state of the valence band and the lowest state of

the conduction band are located at about 0.07X, where X= (1,0,0)/_. These nanotube band gaps, and that of 3.68

a

eV for the newly discovered cubic phase of silicon nitride (c-Si3N4), are predictions-given that we are unaware of

any available experimental values, to date.

NASA/CP--2001-210948 547 August 6, 2001
TableI. Calculated electronic properties (i.e., band gaps) of selected carbon-based and other semiconductors

versus measured values. All the calculations employed the BZW method. The results below are therefore

obtained with the optimal basis sets for the affected materials. Data for c-Si3N4 and for carbon nanotubes

(8,4) and (10,0) are predictions. Further details are available in the identified publications. Eg stands for the

band gap. Effective masses, M*, have symmetry and directional (parallel and perpendicular) labels as

subscripts. The free electron mass is m0. Energies are in electron volts (eV). Reference 7 explains the concept

of practically measurable band gap, as per tail structures in the density of states. When a range is provided

for the theoretical band gap, the lower limit indicates the theoretical minimum gap while the upper limit

shows the largest experimental value, depending on the sensitivity of measurement instrument and method

and of subsequent analysis.

(ref. 6)

Eg

M p,ll

M'p,±

M_,,,ll

M£±

GaN (ref. 7)

Eg

Si

(ref. 7)

Eg

Band Width

(valence)

M_t

ZnO

Eg

AIN (ref. 24)

Eg

Carbon

nanotube

(8, 4)

Eg

Calculation

2.6 eV

7.5 mo

1.2 mo

3.4 mo

1.2 mo

3.4 eV

1.02 eV

12.1 eV

3.2 eV

5.5-6.2 eV

Measurements

2.8, 3.0 eV

very anisotropic

1.0 - 1.5 mo

-3.4 eV

0.2 -t- 0.02 mo

1.14, 1.17 eV

12.5 eV

0.19 mo

0.98 mo

3.4 eV

3.9-6.2 eV

0.90 eV Unknown to us

Diamond (C)

(ref. 7)

Eg

Wv (valence

Band width)

Mn,ll

Mn,ll

3C-SiC

(ref. 8)

Eg

m*x

M* xw

4H-SiC

(ref. 8)

Eg

n,1

Mn,ll

M_x I

c-Si3N4

(ref. 23)

Eg

ZnSe (ref. 25)

Eg

Carbon

Nanotube

(10, O)

Eg

Calculation Measurements

5.05 eV 5.3, 5.48 eV

21.35 eV 21 -t- 1 eV and

24.2 eV

1.1 at- 0.2mo 1.4mo

0.30 at- 0.03 mo 0.36 mo

2.24 eV 2.2, 2.4 eV

0.72 at- 0.04 mo 0.677 at- 0.015 mo

0.22 -t- 0.02 mo 0.247 -t- 0.011 mo

3.11 eV 3.2, 3.3 eV

0.41 at- 0.02 mo 0.42 mo

0.31 at- 0.02 mo 0.33, 0.29 mo

0.62 -t- 0.03 mo 0.58 -t- 0.01 mo

0.27 -t- 0.02 mo 0.31 -t- 0.01 mo

3.68 eV (Not yet available)

2.6-3.1 eV 2.8 eV

0.95 eV Unknown to us

NASA/CP--2001-210948 548 August6,2001
CONCLUSION

As pertheaboveresultsandothers,thedescribedBZWmethodresolvedthelongstandingtheoretical

underestimation, by30%-50%,ofthebandgapsofsemiconductors andinsulators, includedcarbon-based materials

andcarbonnanotubes. Assuch,themethodprovidesatoolforsignificant contributions oftheoreticalcalculations

in (a)thecorrectdescriptionof non-metallicmaterials, (b)thepredictionsof electronic, optical,andrelated

propertiesofsemiconductors, (c)industriallyrelevant bandgapengineeringandrelateddesignandfabricationof

devices, and(e)thesearchfornovelmaterials.Inparticular, current trendstowardthenanoscale, inlightofthe

prevailingnatureofquantumeffectsattheatomicandmolecular levels,seemtodictatetheuseofthismethodfor

computations thatutilizebasisetsinavariational approachoftheRayleigh-Ritz type.

REFERENCES

1. Rubio,A,Corkill,J.L.,Cohen, M.L.,Shirley,E.L.,andLonie,S.G.,Phys.Rev.B48,1993,11810.

2. Vogel,D.,Kriiger,P.andPollmann, J.Phys.Rev.B55,1997,12836.

3. Gould,S.H.:Variational MethodsJor Eigenvalue Problems. University of Toronto Press, Toronto, Canada,

1957, Chapter 2.

4. Mikhlin, S. G.: TheNumericalPerJbrmance of Variational Methods. Wolters-Noordhofl, Groningen, 1971,

Chapters 1, 2, and 7.

5. Bagayoko, D., Zhao, G. L., Williams, T. D. LDA prediction of electronic properties of BaTiO3. Bull. APS, 43,

No. 1, 1998, 846.

6. Bagayoko, D. et al.: Ab-initio calculations of the electronic structure and optical properties in ferroelectric

tetragonal BaTiO3. Journal of Physics: Condensed Matter, 10, 1998, 5645.

7 Zhao, G L, Bagayoko D. and Williams T. D.: Local density functional prediction of electronic properties of

GaN, Si, C, and RuO2. Physical Review B 60, 1999, 1563-1572.

8. Zhao, G. L., agayoko, D.: Electronic structuJce and charge transfer in 3C- and 4H-SiC, New Journal of Physics

2, 2000, 16.1-16.12 (http://www.nip.org).

9. Bagayoko, D.: Contraction of Gaussian Basis Sets and the Total Energy of FCC Copper. Int. J. of Quantum

Chemistry: Quantum, Chemistry Symposium, 17, 1983, 527.

10. Bagayoko, D., Zhao, G. L.: Predictive Ab-Initio Computations of Properties of Ferroelectric Materials.

International Journal of Modem Physics B Vol. 13 Nos. 29, 30 & 31, 1999, 3767-3773.

11. Momzzi, L., Janak, J. F., and Williams, A. R.: CaleulatedElectronic Properties of Metals. Pergamon, New

York, NY, 1978.

12. Hohenberg, H. and Kohn, W. Phys. Rev. 136, 1964, 3864.

13. Kohn, W. and Sham, L. J. Phys. Rev. 140, 1965, Al133.

14. Perdew, J. andZunger, A. Phys. Rev. B 23, 1981, 5048.

15. Ceperley, D. M. and Alder, B. J. Phys. Rev. Lett. 45, 1980, 566.

16. Vosko, S. H., Wilk, L., and Nusair, M. Can. J. Phys. 58, 1980, 1200.

17. Callaway, J.: Quantum Theory of the Solid State, Second Edition, Academic Press, Boston, USA, 1991.

18. Johnson, K. A. and Ashcroft, N. W. Phys. Rev. B 58, 1998,15548

19. Sch6ne and Egniluz, A. G. Phys. Rev. Lett. 81, 1998, 1662.

20. Bagayoko, D. et al.: 3d Impurities in Aluminum. Phys. Rev. B 54, 1996, 12184.

21. Harmon, B. N., Weber, W., and Hamann, D. R. Phys. Rev. B 25, 1982, 1109.

22. Zhao, G. L. et al. Phys. Rev. B 40, 7999(1989)

23. Bagayoko, D. and Zhao, G. L.: Predicted Electronic Properties of Cubic Silicon Nitride (c-Si3N4). Accepted for

Publication in Physica C, 2001.

24. Luo, Y.: Ab-Initio Computation of the Electronic Structure of Wuxtzite Aluminum Nitride. Master's degree

thesis in Physics, Southern University and A&M College, Baton Rouge, Louisiana, 70813, USA, 2000.

25. Torrence, L.: The BZW Method and the Electronic Properties of Zinc Selenide (ZnSe). Master's degree thesis

in Physics, Southern University and A&M College, Baton Rouge, Louisiana, 70813, USA, 2001.

26. White, C. T. and Mintmire, J. W.:Density of States Reflects Diameter in Nanotubes. Nature 394, 1998, 29-30.

27. Mintmire, J. W. and White, C. T.: Universal Density of States for Carbon Nanotubes. Phys. Rev. Lett. 81, 1998,

2506-2509.

NASA/CP--2001-210948 549 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

INVESTIGATION OF DEFECT STRUCTURE IN MULTILAERED CAGE CARBON

USING X-RAY EMISSION SPECTROSCOPY.

L.G. Bulusheva, A.V. Okotrub

Institute of Inorganic Chemistry SB RAS, pr. Ak. Lavrenteva 3, Novosibirsk 630090

V.L Kuznetsov, Yu.V. Butenko

Boreskov Institute of Catalysis SB RAS, pr. Ak. Lavrentcva 5, Novosibirsk 630090

A. Fonseea

Laboratoire de Resonance Magnetique Nuclealre, Facultes Universimires Notre-Dame de la Palx, Rue de Bmxelles

61, 5000 Namur, Belgium

M.I. Heggie

The School of Chemistry, Physics and Environmental Science, University of Sussex, Falmer, Brighton BN1 9QJ,

UK

ABSTRACT

The cage of carbon nanotubes and onions might incorporate various defects, the portion and kind of which is

likely to depend on the synthetic conditions. To reveal the effect of structural defects on the electronic state of

carbon, multiwall carbon nanotubes and onion-like graphitic particles were examined by X-ray emission

spectroscopy, which measures carbon 2p partial density of occupied states. Multiwall carbon nanotubes were

synthesized in an electric arc and by decomposition of acetylene at 700 °C over Co/Fe catalyst. The samples of

onion-like carbon were produced by the annealing of ultra-dispersed diamonds at 1600°C and 1800°C. Transmission

electron microscopy showed that nanoparticles (tubes or onions) formed at the higher temperature have more perfect

structure. So, the arc-discharge carbon nanotubes are predominantly straight, while the catalytic ones are often bent

and curved. The layers of carbon onions prepared at 1800°C are better graphitized and ordered.

CKo_ spectra of the high-temperature produced nanotubes and onions were found to be very similar to the

spectrum of polycrystalline graphite indicating that these carbon materials are characterized by about the same

portion of defects, which mainly concentrated at the particles boundary. Whereas CKc_ spectra of the sample

produced at lower temperatures exhibited an increase of density of n occupied states in about 1.5 times for

nanotubes and 2 times for onions compared to that for graphite. The observed perturbation of the density of states is

the most likely to be due to defects, which disrupt a uniformity of the conjugated n-system.

To estimate the influence of the structural defects on the density of n-states a number of caged graphitic models

incorporating non-hexagonal tings., spa-hybridized atoms, or incomplete bonds were calculated using quantumchemical semi-empirical method AM1. Comparison between the experimental spectra and theoretical ones indicated

that the localization of occupied n-electrons increases somewhat when two former kinds of defects are occurred in

the tubular and spherical cage and more considerably for the holed structures formed in carbon deficient conditions.

This work was supported by tNTAS (Project No. 97-1700), the Russian Foundation for Basic Research (Projects

No. 00-03-32463a and No. 00-03-32510a) and the Royal Society (Joint Project grant).

Keywords: carbon nanolubes, onions, electronic structure.

Lyubov G. Bulusheva

Institute of Inorganic Chemistry

Pr. Ak. Lavrenteva 3

Novosibirsk 630090

Russia

Fax: +7 (3832) 344489

E-mail: bul@che.nsk, su

NASA/CP--2001-210948 550 August 6, 2001
Proceedingsof the SixthAppliedDiamondConference/Second

FrontierCarbonTechnologyJointConference(ADC/FCT2001)

Y. Tzeng,K. Miyoshi,M. Yoshikawa,M.Murakawa,Y. Koga,

K. Kobashi,G.A.J.Amaramnga(Editors)

HRTEM study of carbon aggregates under irradiation

J.C. ARNAULT, NL ROMEO, G. EHRET, et F. LE NORMAND

Groupe Surfaces-lnterfaces, Institut de Physique et Chimie de Strasbourg, IPCMS-GSI,

UMR 7504, Bat 69, 23, rue du Loess, 67037 Strasbourg

F. BANHART, P. KOHLER-REDLICH

MPI Stuttgart, Seestr. 92, 70174 Stuttgart

Under CVD conditions of carbon activated by a hot filament, well organized graphitic planes and

carbon aggregates (onions, polyhedral graphite) have been observed on a copper substrate [1]. More, graphitic

aggregates with a negative curvature called tripods were also found. This particular geometry theoretically

predicted, was never observed under CVD conditions. In order to better understand the formation mechanisms

of these aggregates, HRTEM observations have been performed using the JEOL ARM-1250 kV microscope

(MPI Stuttgart). This high energy microscope is rather adequate to study the stability and the reactivity far from

equilibrium of carbon aggregates using an intense electron beam (200A I em _) under annealing at 770 °C [2].

Hence, the direct transformation of graphitic shells to diamond phase was directly monitored [3].

We determined the radial distribution of the inter-shell distances was by HRTEM image analysis. In

this way, the structural organization of the carbon shells under irradiation could be followed. More, this

analysis is sensitive to the presence of sp3 defects in the aggregate structure and may be useful to determine the

sp _ / sp3 ratio. Several steps were observed: the apparition of bended planes with a sixfold symmetry, a diamond

phase at the centre after a strong irradiation, then, the evolution to the structure of an onion healed from its

defects. The tripod borders surrounding the negative curvature were also investigated and dislocations were

localized. The organization of these dislocations with antiphase boundaries can explain the negative curvature.

References

[ 1] F. Le Normand, L. Constant, G. Ehret and C. Speisser, J. Mater.Res., 1998.

[2] F. Banhart, Rep. Prog. Phys., 62, 1181-1221, 1999.

[3] F. Banhart and P.M. Ajayan, Nature, 382,433, 1996.

Keywords Carbon aggregate Irradiation

Microscopy

Transmission Electron

Corresponding author

Jean-Charles ANNAULT

IPCMS / GSI

23, rue du Loess

67037 STRASBOURG Cedex

France

Tel: (33) 3 88 10 71 02

Fax: (33) 3 88 10 72 48

e-maih Jean-Charles.Amault@ipcms.u-strasbg.fr

NASA/CP--2001-210948 551 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

SYSTEMATIC STUDY OF CONDITIONS FOR DEPOSITION OF VARIOUS

CARBON STRUCTURES PREPARED BY ION BEAMS

N. G. Shang, F. Y. Meng, Q. Li, X. F. Duan, Y. Lifshitz, C. S. Lee, S. T. Lee, L Bello °

Center Of Super-Diamond & Advanced Films (COSDAF) and

Department of Physics & Materials Science,

City University of Hong Kong, Hong Kong, China

ABSTRACT

The article presents a systematic study of various carbon structures prepared by an ion beam

method and reports detail conditions for their deposition. The ion beams were extracted from variable

mixtures of methane/hydrogen/argon plasmas in a Kaufman ion source and impacted on the silicon

substrates.

The deposited carbon structure includes amorphous and gmphitic carbon phases, beta silicon

carbide, and nanocrystallites of diamond in their hexagonal, cubic and and rombohedral configurations as

revealed by high resolution transmission electron microscopy. In addition to diamond nanocrystallites

confined in an amorphous carbon phase, the nanocrystallites also developed epitaxial relationships with the

silicon substrates at specific deposition conditions. The deposition products, and nanostructured diamond

crystallites depended very much on the ion beam parameters, however, they were strongly affected by the

substrate temperature.

This work supports our previous experimental results acquired from a bias assisted hot filament

CVD method and further elucidates diamond nucleation and heteroepitaxial growth.

*Corresponding author

Mailing address:

Department of Physics & Materials Science,

City University of Hong Kong, Hong Kong,

Kowloon Tong

Kowloon,

China

Tel.: +852-2788-7891

Fax: +852-2784-4696

E-mail address: apibello@ciWu.edu.hk (I. Bello).

NASA/CP--2001-210948 552 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

ENERGY-CONTROLLED DISCONTINUOUS STRUCTURES IN CARBON

MATERIALS

D.Shvets

Institute for sorption and problems of endoecology of NAS Ukraine

13, Naumova str., 03164, Kiev, Ukraine

E-mail: dshvets@ispe.kiev.ua

The study of features of forming of model discontinuous films (reduced ions of metal) on matrixs

of a various nature has given a number of fundamental results stipulated by revealing of their

mobility, anomality of optical, thermoelectric and issuing properties. On carbon matrixs the

formation of such structure proceeds basically on the diffusive dodge, with the help which one there

is a key opportunity of realization energy-controlled of forming process of discontinuous films

assigning type such as with dedicated, for example by issuing properties.

The purpose of our work was finding - out by the help of a scanning submicroscopy method for

formation conditions of discontinuous structures at energy-controlled reactings of processes

initiation by their forming from adsorbed in porous layers of carbon matrixs of model solution

HAuC14. The process was conducted as at processing the most carbon material by a stream of

ionizing radiating (e, gamma-quantums), and at beaming all heterogeneous system. The detailed

analysis of received obtained microphotographs demonstrates, that the forming of discotinuous

films starts in into-surface layer of a carbon fragment, and then is distributed deep into matrixs, i.e.

the secondary discontinuous films will be derivated new.

Is marked, that at particular modes of processing of a matrix the fissile process of growth of

discontinuous structures takes place, their sizes achieve magnitudes 2-10 microns and then the

process of growth is retarded. That fact merits notice also, that at "weak" modes of matrix

processing by a stream of radiatings the discontinuous films will be derivated by the way of surface

crystal grains.

Is retrieved, that at augmentation of stream power the process of forming of discontinuous films is

significally changed, proceed volumetric diffusion resulting in to forming of volumetric cluster

structures more actively. The discontinuous films, characteristic for low-energy processes of

forming, at middle- and high energy processes are enlarged, becoming clusters of the various

geometrical form.

Is showed, that in discontinuous structures of gold at fields some volt on centimeter is watched a

deviation from a Ohm's law, that is rather essential at obtaining conductance at major entered

powers and issue, observed at it, of hot electrons, and also excited by a radiation current of a light

by discontinuous films.

The new technology of obtaining of items is offered on the basis of discontinuous structures of a

given type, sizes and various geometrical constitution. The centres of issues reshaped on offered

technology, will have very small scatter on energies thanking high-density of metal in nonuniform

structure, that in turn will allow to reshape very major emission currents (200 mkA and more), that

is much higher than arguments of insular frames, for example, from molybdenum.

NASA/CP--2001-210948 553 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

POSSIBILITIES OF IDENTIFICATION OF ADVENTITIOUS CARBON ON A

CARBON SURFACE BY XPS AND XAES

A. P. Dementjev, K. I. Maslakov, A. V. Naumkin, O. O. Zabusov

IRTM, RRC Kurchatov Institute, Kurchatov sq. 1, Moscow 123182, Russia

ABSTRACT

The main question of the characterization of any carbon surfaces by surface sensitive methods is whether

adventitious carbon (AC) presents on the samples. The inherent structures of the CKVV Auger spectrum for AC on

non carbon surface have been determined. There are two features of identification of AC surface: 1) the sp3-character

of the CKVV Auger spectrum and 2) the left part of the CKVV spectrum that originates from electron energy loss

lies below that of polyethylene and graphite but above that of methane. These results allow wide use of the CKVV

spectra with X-ray excitation for study carbon samples prepared in-situ or ex-situ and identification of AC on it.

Keywords: Adventitious Carbon, AES, XPS

INTRODUCTION

It is well known that AC is always observed on oxide and metal surfaces prepared ex-situ. The AC on sample

surfaces may originate from the atmosphere air, sample handling and/or contamination in the analysis chamber by

vacuum pump oil. The Cls XPS peak of AC is used as the binding energy standard for studying of insulators. The

nature of AC as the base for charge reference in the XPS method was investigated by many authors (ref. 1-3).

The complexity of chemical analysis of AC by XPS on such samples as graphite, diamond, diamond-like

materials and organic samples is caused by the small range of the Cls binding energy (ref. 4,5).

The Auger decay in carbon atoms involves valence band electrons and therefore the AES is often capable of

yielding information concerning the chemical environment of atoms in gas-phase molecules or in near-surface

region of solids (ref. 6,7). Obviously, the Auger spectra can provide useful information about carbon atoms in AC.

With this in mind we have studied the XPS and the X-ray excited AES (XAES) spectra of AC on SiO2, A1203, oxide

metals and native silicon oxide surfaces

The aim of this paper is to develop the potential of the X-ray excited Auger electron spectroscopy (XAES)

combined with the X-ray photoelectron spectroscopy to solve the problem of identification of carbon chemical states

on surfaces of real films.

EXPERIMENTAL

The XPS and the XAES data were obtained using the MK II VG Scientific spectrometer with oil pump.

Photoelectron and Auger processes were excited with an A1 Kc_ X-rays source. The vacuum in the analysis chamber

was 1.10 -7 - 5* 10 -8 Pa. The spectra were collected at the constant analyser energy mode with the pass energy of 20

eV for the XPS, and 50 eV for the XAES. The Cls spectra were acquired with 0.1 or 0.05 eV step size; and 0.25 eV

step size was used for the CKVV spectra. All the spectra were referenced to Fermi level using the Cls binding

energy at 284.8 eV.

We can't vary AC concentration on sample surfaces so we look over many samples to obtain different

concentrations.

Auger spectrum shape is shown to be dependent on AC concentration on a surface but practically independent on

a substrate. So we present the results of the AC native silicon Auger spectra shape investigation on SiO2.

The samples were being stored in air from 1 hours up to some years. Carbon concentrations on surfaces were

measured by the Cls, Si2p and Ols XPS lines. Some samples contain 10 - 15% others about 30-80% of AC. The

Cls, CKVV, wide, Si2p spectra were acquired for 40 minutes. Then the sample surfaces were cleaned in-situ with

Ar ions at 3 keV and Ip -- 25 mkA/cm 2 for 30 seconds. The noticeable concentration of carbon on the sample

surfaces was accumulated only after 2-hour storage in the analysis chamber. These data allowed suggestion that the

AC had been formed due to absorption outside the spectrometer.

RESULTS AND DISCUSSION

Fig. 1 shows the comparison of Wide spectra of polyethylene with that of AC on native silicon oxide (17%) and

SiO2 (80%). One can see the significant difference in intensity of Cls and Si2p peaks for the last samples. The Si2p

NASA/CP--2001-210948 554 August 6, 2001
peakintensityinnativesiliconoxidechangesslightlyafterremovingofACbyArsputtering. Thisfactallowstodo

suppositionaboutsmallthicknessofAC.WedenotetheACasmonolayer (ML).TheintensityofClspeakonSiO2

isbiggerthanthatonnativesiliconoxide.WedenotetheAC(TL)asathicklayer.Accordingourestimationbased

ontheequationfrom(ref.7)andontheSi2pXPSintensitywithandwithoutoftheAC,athicknessoftheoverlayer

is equaltoabout8monolayers. Theinformationdepthof theCKVVemissionaccordingtoestimationonbase

calculation(ref.8)isequal2-3monolayers. So,thethicknessofAC(TL)considerably exceedstheinformationdepth

oftheCKVVemission.

69

t4

69

tt1.0

0.8

0.6

0.4

0.2

Wide, XPS

Si2s

!

II

Cls Ols

:' i

:: ....... polyethylene ,

I! .... AC(ML) I',

I_i -- AC(TL) ;_

I

• II

II

II

II

Ill I!

.............. . ..... ..

Si2p ;: ,, /_::

, ' ._.._I-'_ ,

, I , I , I ,

100 200 300

0.0

I , I

0 400 500

Binding energy, eV

Figure 1. The wide XPS spectra with different concentration of AC.

Fig. 2. shows the comparison of the CKVV spectra of polyethylene with that for two concentrations of AC: ML

and TL. Two fold-polyethylene films were split in air and then were put in a chamber. Any cleaning procedures

were not performed in the chamber. The CKVV of AC (TL) practically coincides with that of polyethylene, but

there are differences in the slope and the structure of the left part of spectra. The differences in the parts of the

spectra are shown in frame. The CKVV width of AC(ML) is more narrow and this spectrum has more remarkable

incline of the left part. Thus, the shape of CKVV spectra depends on the concentration of AC on surfaces.

1.0 CKVV _ AC (TL)

,/o.-_ .... Polyethylene

0.8 if/" %_,_ Z//..' '..\ ...... AC (ML)

(/)

4t- 0.6 . .-"" .... """ .. . ° ",,,,

>' 0.4 "'" ,t'- _ ,

_c 0.2

0.0 231 238 24 _ _.

I , I , I , I , I , I

230 240 250 260 270 280

Kinetic energy, eV

Figure 2. The comparison of the CKVV spectra for polyethylene with that for two concentration of AC.

NASA/CP--2001-210948 555 August 6, 2001
1.0 - CKVV, XAES

SJ

0.8

._,

o.6

_>' i"I IIS

•_ 0.4

0.2

0.0

220

III

24O

x i i

S°o I

" "" IV

"'-- AC(ML)

....... Methane

--- Polyethylene

I , I ,

260 280

Kinetic energy, eV

Figure 3. The comparison of the CKVV Auger spectra of polyethylene, methane and AC(ML).

In order to have a better understanding of the structure CKVV spectra, we have compared spectra of AC and gas

phase molecules from (ref. 11).

The structure of the CKVV spectra of TL(AC) coincides very closely with that of polyethylene (Fig. 2). This

clearly demonstrates the character of sp3-bonds in carbon contamination. But there is a very important difference at

the left part of the spectra that due to energy loss of Auger electrons as they escape from a surface [6]. The left part

of the CKVV spectrum of AC(ML) (Fig. 3) lies below that of polyethylene but above that of gaseous molecules of

methane, neopentane and normal pentane (ref. 9). The comparison of the CKVV Auger spectra allows suggestion

that the molecules of AC (ML) are at the position between the gas molecules and solid: one part of the molecules is

connected with solid and the rest dangle outside the surface. This supposition are confirmed by the CKVV of

AC(TL) in Fig. 2. The spectrum practically coincides with that of polyethylene. It is necessary to mark increasing

width of the CKVV spectra with increasing of a number atoms in gas molecules (ref. 9) CH4, C2H6, CH3CH2CH3,

CH3(CH2)3CH 3 . So, the width of the CKVV spectra pentane CH3(CH2)3CH 3 is close to that of AC(TL) and

polyethylene spectra, but the left part of gas phase Auger spectra do not change with increasing of a number atoms

in gas molecules. The feature I shows the left tail of spectra that is due to the distinction in the electron energy loss.

There are the inherent features in II and III for each spectrum. The feature IV shows the difference in the width for

every spectrum.

The information depth of the CKVV Auger emission is equal to 2-3 monolayers. So, when AC layer thickness

exceeds this value spectra will look like the spectrum of TL (AC) on fig.2.

Three cases are possible:

a) an Auger spectrum of any carbon surface has the typical shape of sp3-bonds. In order to identify chemical bonds

on the surface it is necessary to analyze in detail the left part spectrum (as shown in the frame of Fig.2.).

b) an Auger spectrum of any carbon surface has the typical shape of sp2-bonds. It means the absence of any AC on

the surface.

c) an Auger spectrum of any carbon surface does not coincide with the typical shape of sp2-or sp 3- bonds and the left

part has a slope. It means that the spectrum is a sum of that of substrate and AC(ML).

NASA/CP--2001-210948 556 August 6, 2001
5 ,,-_ r"/ _ ......... AC(ML)

%.

I , , / \ , -- Sum

4 ....... _ "''''1/ _ '

HOPG

'_ %IIi%%i

C4 3

¢- | -. .- .... • ". %.

"_ ..° "° %.

0 , I , I , I

220 240 260 280

Kinetic energy, eV

Figure 4. The simulation of the CKVV Auger spectrum of graphite with AC(ML)

Fig. 4 shows a simulation of the CKVV Auger spectrum of graphite with AC(ML) with equal intensity spectra.

The summation shows a modification of the total spectrum. This qualitative data give an insight of what must

happen when the surface of graphite is covered by AC. As one can see, the detection of AC(ML) on a surface with

sp 2-bonds produces no problems.

2.0 --Sum

1.8 /_-_ ....... AC(ML)

1.6 ,," ",_-, ,y _ .... Polyethylene

,, %. I

1.4 ........ _ ''

C

1.2

..d

1.0 _...:-.,.-,

>_ 0.8 "'''--- "'" "_

c0.6 . ..-- ............. ._._

0.4 .............. " _.

0.2 _",%.

0.0

I , I , I

220 240 260 280

Kinetic energy, eV

Figure 5. The simulation of the CKVV Auger spectra of polyethylene with AC.

Fig.5 shows the simulation of the CKVV Auger spectra of polyethylene with AC with equal intensities. The

detection of AC on surface with sp 3 -bonds presents problems. In the case one should to analyze both the shape of

the main feature and left tail spectrum.

So, Fig. 4,5 display the transformation of the Auger spectra of carbon surface sp 2 and sp 3- bonds due to the

presence of AC(ML). We have analyzed a lot of carbon films with sp % and sp 3- bonds prepared ex-situ and we

observed sometimes the CKVV spectra as the sum in Fig. 4,5. We observed the CKVV Auger spectra similar to that

of AC(TL) only on CNx (sp 2 - bonds) and melamine C3NdHd (sp 2 - bonds) surfaces.

NASA/CP--2001-210948 557 August 6, 2001
CONCLUSION

TheAugerspectrastructurecorresponds top3-structure anddependsonAClayerthickness. Therearethree

featuresforidentificationfAConacarbonsurface:

a) thesp3-character oftheCKVVAugerspectrum;

b) thestructureoftheleftpartspectrum;

TheidentificationfAConcarbonsurfacewithsp2-bondsdonotpresentsproblems. TheidentificationfACon

carbonsurfacewithsp3-bondspresentssomedifficultiessomoreaccurateworkisneededwithleftpartofthe

CKVVAugerspectra. NoteverycarbonsurfacecontainsACasopposedtometalandoxidesurfaces. Forexample,

HOPG,fullerene, quaterphenil andpolyethylene surfacespreparedex-situ have not contained AC. The row N(E)

CKVV spectrum is powerful method of the study of AC on carbon surfaces.

REFERENCES

1. Brandt E.S., Untereker D.F., Reilley C.N., Murray R. W., J. Electr. Spectr. Relat. Phen., 14 (1978) 113

2. Swift P., Surf. Inter. Anal., 4 (1982) 47

3. Barr T. L., Seal S., J. Vac. Sci. Tehnol., A 13 (1995)1239.

4. Belton D.N., Schmieg S.C., J. Vac. Sci. Technol. A 8 (1990) 2353.

5. Wagner C.D.,. Naumkin A.V, Kraut-Vass A., et al., NIST X-ray Photoelectron spectroscopy Database. NIST

Standard Reference Database 20, Version 3.0 (Web version)

6. Ramaker D. E., J. Vac. Sci. Technol. A 7 (1989) 1614.

7. Ramaker D. E., Critical Reviews in Solid State and Materials Science, 17 (1991) 211

8. Practical Surface Analysis by Auger and X-ray Photoelectron Spectroscopy, Ed. by Briggs D. and Seah M. P.,

John Wiley & Sons, 1983, Chapter 5, p. 211, eq.5.27.

9. Seah M.P., Dench W.A., Surf. Interface Anal.,l(1979)2.

10. Rye R. R., Jennison D. R., Hutson J. E., J. Chem. Phys., 73 (1980) 4867.

NASA/CP--2001-210948 558 August 6, 2001
7

NITRIDES, CARBIDES, AND B-C-N MATERIALS

Proceedingsof the SixthAppliedDiamondConference/Second

FrontierCarbonTechnologyJointConference(ADC/FCT2001)

Y. Tzeng,K. Miyoshi,M. Yoshikawa,M.Murakawa,Y. Koga,

K. Kobashi,G.A.J.Amaramnga(Editors)

SYNTHESIS OF cBN FILMS BY PLASMA-JET CVD IN THE Ar-N2-BF3-H2 GAS

SYSTEM

s. Matsumoto*, W.J. Zhang**, and J. Yu

Advanced Materials Laboratory, National Institute for Materials Science

1-1 Namiki, Tsukuba-city 305-0044, Japan

ABSTRACT

cBN films have been synthesized on Si substrates from the Ar-N2-BF3-H2 gas system by bias-assisted plasma jet

CVD. The BN films obtained show a strong IR absorption of the cBN peak and the cBN content was estimated to be

over 90 %. Crystal quality was as high as sharp XRD peaks of cBN and Raman peaks of TO and LO phonon modes

were observed. Crystal size observed by TEM was 0.2-0.5 gm in width and nearly film thickness in length. Residual

stress in the films was evaluated to be 1-2 GPa and large film thickness of over 20 gm was achieved. The deposition

rate was as high as 16 gm/h. Three factors seem to contribute to this success: the use of fluorine, substrate biasing,

and a dc arc plasma jet. The characteristics of this synthesis method, properties of the cBN films obtained, and the

shortcomings of this method to be overcome were reviewed and the roles of the above three factors were discussed.

Keywords: cubic boron nitride, plasma CVD, dc arc jet, fluorine, bias

INTRODUCTION

Owing to its unique properties such as high hardness, high thermal conductivity, low dielectric constant (for high

frequency), largest band gap among all the covalent materials so far known, dopability to n- and p-type

semiconductors, and low reactivity to iron-family metals, cubic boron nitride (cBN) is expected to be a potential

material for high-temperature high-frequency electronics and for short wavelength optical and optoelectronics, and

as hard coatings for cutting ferrous materials. After the success of chemical vapor deposition (CVD) of diamond,

low-pressure synthesis of cBN has been also tried for about 20 years by using various methods, but the performance

has been far behind that of diamond CVD (ref. 1-2). These methods can be classified into three categories, i.e.,

physical vapor deposition (PVD), low-pressure CVD, and medium-pressure CVD as summarized in Table 1. The

first two apply ion bombardment to substrates by using ion beams or substrate bias under low gas pressures. The

third one does not use ion bombardment but has not yet yielded reproducible results. The first two can give cBN

films even with high phase purity, but owing to the effect of ion bombardment on substrates, the crystal size was

limited to a few to several tens nm and the films have high residual stress, which limited the film thickness below 1

gm. By using a plasma jet and a fluorine-containing gas system, we have succeeded in synthesizing cBN films with

large crystal size, low stress, and film thickness over 20 gm (ref. 3-12). In this paper, the experimental procedure

and results are reviewed, and characteristics of the method, film properties, and the problems to be overcome will be

discussed.

Table 1. Comparison of deposition conditions between the present methods and the methods used so far.

PVD Low-pressure CVD Medium pressure CVD Present method

Method or plasma BS, IBAD+EB, IBD, Helicon ECR, rfCCP, rfICP, Helicon Hot-ilL+if, Microwave DC-jet

Source and _as hBN or B, +Ar or Ar+N: BeH6(TMB, BZ)+Ar+N: BeH6(NaBH3)+NH3+H: BF3+Ar+Ne+H:

Pressure (Torr) 104_10 2 104-10 2 0.1-100 50

Bias voltage, Ion energy 10_-10 _ 50-10 _ - -85V

Bias Current (mA/cm 2) 10 2-1 10 1-1 - 100

Rows 4 and 5 were summarized after ref. 2. BS: bias sputtering, IBAD: ion-beam-assisted deposition, EB: electron beam

evaporation, IBD: ion beam deposition, TMB: trimethyborazine, BZ: Borazine

DEPOSITION METHOD

* Corresponding author, e-mail:matsumoto.seiichiro@nims.go.jp

Present address: Department of Physics and Materials Science, City University of Hong Kong, Kowloon,

Hong Kong

NASA/CP--2001-210948 561 August 6, 2001
Thesetupofthedepositionapparatus isshownelsewhere(ref.3).AnAr-N2-BF3-H2 gasmixturewasusedasa

plasmaandreactant gasunderachamber pressureof50Tort.Apieceof(001)Siwafer(14x14x0.5mm)wassetas

asubstrateonawater-cooled substrateholderandplacedintheflameofthedc-arcplasmajet.TheSisubstratewas

coveredwithaBNcoverwithaholeinthecenterinordertoconcentrate thbiascurrentontotheSisubstrate. The

substrateholderwasnegativelybiasedwithadcpowerora400kHzrf power.Thesubstratetemperature was

measuredfromthebacksidewitha fiberthermometer (OptonicoOT100).Otherdepositionparameters are

summarized inTable2.Thephasecomposition, structure, andcrystal qualityofthefilmswereinvestigated byFTIR,

XRD,SEM,TEM,andRamanspectroscopy.

Ar(slm)

N2(slm)

BF3(sccm)

H2(sccm)

Pressure(Torr)

Jetpower(kW)

Substratetemperature (°C)

Biasvoltage(V)

Duration(min)

Table 2. Deposition conditions

dc bias

20

1.5

3

5-3

50

3-6

930 - 1150

-85

10- 180

rf bias

20

2.5 - 0.0025

3- 10

0 -30

50

3-6

840- 1080

-130--175 (self bias)

10

CHARACTERISTICS OF THE DEPOSITION METHOD AND PROPERTIES OF THE cBN FILMS

A. The characteristics of the deposition method.

1. Fluorine is necessary for cBN growth. B2H6 or BC13 instead of BF3 did not give cBN under the similar deposition

conditions. There is an optimum window in the BF3/H2 ratio for cBN deposition, i.e., 3/5 to 3/1.5. If the ratio is

higher, etching predominated and if lower, hBN deposits.

2. There is a window for substrate bias voltage also, about -85 V for dc bias and around -150 V for rf selfbias. The

crystal quality of cBN was much better for films obtained with dc bias than that with rf bias. The properties of

cBN films described below are those of the films by dc bias.

3. The chamber pressure during deposition used in the present experiments is 50 Torr, which is more than 3 orders

higher than those used in conventional CVD and PVD methods using ion bombardment. The deposition pressure

seems to affect bias current, deposition area, deposition rate, crystal quality, and controllability of the process.

Therefore, it should be chosen depending on what performance one wants.

4. Substrate temperature is rather high, 930 - 1150 °C for the dc bias. High substrate temperature was favorable for

obtaining a high-quality cBN film, probably because of high bias current through an insulator film like BN, high

reactivity on the surface, and annealing effects during deposition.

B. The properties of the cBN films obtained.

1.cBN content estimated by IR absorption spectroscopy was over 90 % but the strong absorption of the cBN peak

makes the evaluation of the percentage difficult. However, hBN peaks always appeared even if they were very

weak. Examples of IR spectra are shown in Fig. 1 for the films deposited for 10 and 60 min.

2. SEM and TEM observation revealed the columnar growth of crystals and the crystal size was 0.2-0.5 gm in

lateral width and about the film thickness (a few gm) in length. Clear Raman peaks of cBN phonon modes were

observed for the first time for a cBN film from the gas phase. An example of Raman spectra for the cBN film

obtained is shown in Fig. 2 together with the spectrum of commercially available cBN crystals made by a high

pressure high temperature method.

3. Compressive stress in the films was estimated to be 1-2 GPa by peak shifts of XRD and Raman spectra. This

value is lower than the value 2-20 GPa reported so far and a film of over 20 gm in thickness can be obtained.

4. Deposition rate of 0.3 btm/min, which is 3-5 times higher than the highest value so far reported for cBN film

growth.

NASA/CP--2001-210948 562 August 6, 2001
5.ThereisaninterfacelayerofSi/aBN/hortBN/cBNforthefilmbythismethodalsoasalwaysobservedforcBN

filmsbyPCDandCVDmethodsusingionbombardment. Thethicknessoftheinterfacelayerisrelativelyhigh

(several tensto200nm).

inn t (a) 10 min f-_ /-_

=o _'1 "-'

_ ..........

6o- (b) 60 rain S

20-

4000 35'o03o'o0 25'o02o'o0 15'00 10'00 500

Wave numnber (cm 1)

Fig. 1 IR absorption spectra of cBN films on Si deposited at

1150°C under -85V dc bias for (a) 10 min and for(b) 60 min.

20"

1..5-

10-

0..5-

5

v00

"_ 3-

_'oo1d00 1_oo IAO 11oo 1_oo

(b)

'4o

1500 1400 1300 1200 1100 1000 900

Raman shift (cm 1)

Fig. 2 Raman spectra of (a) a cBN film deposited at 1150°C

for 1 h and (b) a 0.4ram-size crystal commersiaJly available.

C. However, the crystal quality and the growth rate are still far behind those achieved for diamond film growth.

1. As far as IR spectra are concerned, usually phase pure cBN films cannot be obtained as mentioned above, while

phase pure diamond films can be obtained easily by many CVD methods.

2. The crystal size of the cBN films is small and crystal quality is still low. Single crystal diamond with the size of

several hundred gm can be obtained from the gas phase and the FWHM of about 2 cm- for the diamond Raman

peak has been achieved, whereas the smallest FWHM of the cBN peaks is about 10 cm -1 at the present stage. The

background level of the Raman spectra by fluorescence is high even for one of the lowest cases as shown in Fig.

2, while it is very low for high-quality diamond films.

3. Film surface is usually rough and crystal faces are not observed. Figure 3 shows SEM surface morphologies of

cBN films deposited (a, b) under most usual conditions (-85V bias, 3 sccm BF3, 3.5 sccm H2) and (c, d) by threestep process (-85V, 3 sccm BF3, 3.5 sccm H2, -85V, 3 sccm BF3, 1.2 sccm H2; -70V, 3 sccm BF3, 1.2 sccm H2).

The lower hydrogen flow rate and lower bias gave a more flat surface (Fig. 3c, 3d) but the surface is composed of

small particles of several tens nm in size. On the other hand, the surface of diamond films are usually covered

with crystal facets or smooth if textured or epitaxial growth occurs. Appearance of smooth surfaces is reported for

amorphous or nonocrystalline diamond films.

4. Residual stress is still exists which prevents good adhesion of thick cBN films on thick substrates, which means

that the stress is relaxed partly by plastic deformation of the Si substrate at high substrate temperatures during

deposition.

On the process of the present deposition method, there are also some problems to be overcome.

5. The deposition rate of cBN films is still low, only one fiftieth of the highest rate of diamond deposition by dcjet CVD; a deposition rate of about lmm/h has been reported.

6. The substrate temperature used is too high to deposit cBN on various substrates.

7. There are some shortcomings for DC-jet CVD, i.e., high cost of the operating gases and the electric power, and

lower controllability of the process due to its high gas temperature.

DISCUSSION

Three factors seem to contribute to the success of cBN synthesis by this method: the use of fluorine, substrate

biasing, and a thermal plasma. Fluorine has high reactivity in forming strong chemical bonds with many elements

and it has also high etching ability as many fluorine compounds are being used as etchants in semiconductor

NASA/CP--2001-210948 563 August 6, 2001
,'

............ •.... ,',',',°,',_','J_,_,,f,'-_:_'_.°o'." .......................... ,_,;,',N,N,,;°°"o,',,',I,'-,N_, ","

. •. ,_o,,, ,, _,o,_,_o;o_.o ,-°,°,, ,,,,, ,,,,,,. °,,,,°.......,°,,,, ,.. ,eep...,

_..f_._'..:_:::._x-_x-:°;.:-x-:°:.: ,'._,.._.._x+:_c.a_'eye'..:..'_'.."...+'e4f4.::_.:_+, ,_@.:,

• °

_, °

_°°

Fig. 3 SEM morphologies of cBN films deposited under the 85 V bias (a) and 70V bias (c). (b)

and (d) are the enlarged images of (a) and (c), respectively. See text for more details.

technology. It must react with unsaturated BN bonds such as in hBN and forms sp 3 B-F and N-F bonds or

preferentially etches hBN as demonstrated by Kalss et al (ref. 13). It must also react with cBN and formed surface

B-F and N-F bonds, which stabilize cBN surfaces. However, the surface bonds with fluorine on cBN, particularly BF bonds, seem to be stable and prevent the further growth of cBN (ref. 14). Therefore, substrate bias during the

deposition process must be necessary to grow a cBN film. In our method also, cBN did not deposit without substrate

bias. However, because the gas pressure used is as high as 50 Torr, the effective bombardment energy should be

much lower and the effects of substrate bias is probably different from those in conventional PVD and CVD

methods. The process at the very near the film surface seems more important as mentioned above. Owing to its high

ion and radical density, thermal plasmas can deposit diamond in high growth rate. For cBN synthesis, high-density

plasma contributes not only to high concentration of reactive species but also to high ion fluxes to the substrate. The

bias current in this method was in the order of 100 mA, which is about 2 orders higher than that used so far for cBN

synthesis.

REFERENCES

1. P.R. Mirkarimi, K.F. McCarty and D.L. Meddlin: Mater. Sci. Eng. R21 (1997) 47.

2. T. Yoshida: Diamond Relat. Mater. 5 (1996) 501.

3. S. Matsumoto and W.J. Zhang: Jpn. J. Appl. Phys. 39 (2000) L442.

4. W.J. Zhang and S. Matsumoto: Appl. Phys. 71 (2000) 469.

5. W.J. Zhang and S. Matsumoto: Chem. Phys. Lett. 330 (2000) 243.

6. W.J. Zhang and S. Matsumoto: J. Mater. Res. 15 (2000) 2677.

7. S. Matsumoto and W.J. Zhang: to be published in Diamond Relat. Mater. 10 (2001).

8. W.J. Zhang, S. Matsumoto, K. Kurashima and Y. Bando: to be published in Diamond Relat. Mater. 10 (2001).

9. W.J. Zhang, M. Tansho, S. Matsumoto and T. Mori: J. Appl. Phys. 89 (2001) 1734.

10. W.J. Zhang and S. Matsumoto: Phys. Rev. B 63 (2001) 073201.

11. W.J. Zhang and S. Matsumoto: submitted.

12. W.J. Zhang and S. Matsumoto, Q. Li, .I. Bello and S.T.Lee: The paper submitted to this conference.

13. W. Kalss, R. Haubner and B. Lux: Diamond Relat. Mater. 7 (1998) 369.

14. K. Larsson and J.-O. Carlsson: J. Phys. Chem. B 103 (1999) 6533.

NASA/CP--2001-210948 564 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

STRUCTURE VARIATION OF CUBIC BORON NITRIDE FILMS IN A TWO-STEP

PROCESS BY CHANGING THE BIAS VOLTAGE AND GAS COMPOSITION

W.J. Zhang, ° S. Matsumoto

Advanced Materials Laboratory, National Institute for Materials Science,

1-1 Namiki, Tsukuba, 305-0044, Japan

Quan. Li, S.T. Lee, and I. Bello

Center Of Super-Diamond and Advanced Films (COSDAF) and Department of Physics and Materials Science,

City University of Hong Kong, Kowloon, Hong Kong

ABSTRACT

To have a comprehensive understanding of the effects of experimental conditions on the film quality, and to

improve the crystallinity of cubic boron nitride (cBN) films deposited by chemical vapor deposition, a two-step

process by changing the bias voltage and gas composition was designed. The film structure and crystallinity were

investigated by Raman spectroscopy, scanning electron microscopy, and transmission electron microscopy. A

critical bias voltage for the growth of the cBN films was revealed; the potential to extend the lateral size of crystal

columns by this two-step process was demonstrated. In addition, the etching behavior of fluorine and its dependence

on the substrate bias voltage was studied.

Keywords: cubic boron nitride, columnar growth, crystal size, crystallinity, TEM.

INTRODUCTION

The excellent physical and chemical properties of cubic boron nitride (cBN), i.e., its hardness being second only

to diamond, high thermal conductivity, chemical inertness, high refractive index and wide band gap in connection

with p- and n-type dopability, make it an extremely promising material for applications in many practical fields.

During the recent years, much effort has been put into improving the preparation of cBN films, different kinds of

deposition were developed. According to the growth mechanism, the deposition methods can be sorted to be

dominated by i) chemical and ii) physical effects. For the deposition utilizing chemical effects, the same effects of

hydrogen to preferentially etch the sp 2 components and to stabilize the sp 3 structure as that in the deposition of

diamond films by chemical vapor deposition (CVD) were expected because boron nitride (BN) forms both cubic and

hexagonal (hBN) phases in many ways analogous to carbon. The deposition of thin films containing cBN by CVD

methods in hydrogen-based plasma was reported by several authors (ref. 1). However, careful and detailed

characterization of the films was still needed to confirm the reliable formation of cBN phase from hydrogen-based

plasma without applying substrate bias. The successful methods which can lead to the deposition of nearly phase

pure cBN films make use of physical effects in spite of physical vapor deposition (PVD) or CVD performed (refs. 2

to 3). The ion-bombardment of the substrate plays an essential role for the formation of the cubic phase by physical

effects. By involving the low-energy ion bombardment during film growth, the deposited cBN films generally

exhibit high compressive stress, which is thought necessary to obtain cubic phase (refs. 4 to 6). This stress results in

a restricted maximum film thickness of up to several hundred nanometers and the poor adhesion of the cBN filmsto

substrates. Currently, the limitation in thickness and crystal size (up to tens ofnanometers), and poor crystallinity of

cBN films deposited by both physical and chemical vapor deposition methods represent a serious obstacle in the

further development and application of cBN films.

Another important aspect in the current deposition process is to optimize the experimental parameters separately

for nucleation and growth of cBN films. Although a layered structure of amorphous/hexagonal/cubic BN is normally

observed in the deposition of cBN films, and the formation of such an interface was considered to correspond to the

incubation of cBN nucleation, the same experimental condition was used for both nucleation and growth in nearly

all previous works. Litvinov et al (refs. 7 to 8) observed an improvement of film crystal quality as well as a decrease

of stress with decreasing ion energy by reducing the bias voltage in a two-step process of ion-assisted sputtering,

* C orre sponding Author; pre sent address: Center Of Super-Diamond madAdvancext Films (COSDAF) and Department of Physics

and Materials Science, City University of Hong Kong, Kowloon, Hong Kong; electronic mail: ap_vjzh@clty_e&_i,_ Fax:+852-

2788 7830; Tel: +852-2788 9596.

NASA/CP--2001-210948 565 August 6, 2001
whichsuggeststhatit maybearealisticwaytoimprovethecrystallinityofthecBNbyoptimizingthenucleation

andgrowthconditionsseparately. Inthispaper,wereportthedepositionofcBNfilmsbyDCjetplasmachemical

vapordepositionusingagasmixtureofAr-N2-BF3-H2. Atwo-stepprocessbychangingthebiasvoltageandgas

compositionwasdesigned. TheformationofcBNwasidentifiedbyglancing-angle x-raydiffraction(XRD)and

Fourier-transformed infraredspectroscopy (FTIR).Thevariationof crystallinityandfilm structurewas

systematically studiedbyRamanspectroscopy, scanningelectronmicroscopy(SEM),andtransmission electron

microscopy (TEM).A criticalbiasvoltageforthegrowthofcBNfilmswasrevealed. In addition, theetching

behavioroffluorineanditsdependence onthesubstratebiasvoltagewasstudied.

EXPERIMENTAL DETAILS

ThedepositionfcBNfilmswasperformedina dcjetplasmaCVDreactorbyusingagasmixtuxeofAr-N2-

BF3-H2. Theexperimental setuphasbeendescribedelsewhere(ref.9). Pre-scratched siliconwafersof(001)

orientationwereusedassubstrates. Thedepositionprocesswasdividedintotwosteps.Theexperiment parameters

foreachsteparelistedinTable1.Inthefirststep,cBNfilmsweredepositedunderahighsubstratebiasvoltageand

ahighhydrogenconcentration. In thesecondstep,thehydrogenflowratewasdecreased, andthesubstrate

temperature wasincreasedto 1150°Ctoincreasethecrystallinityandminimizetheresidual fimstress. Aftereach

step,thephasepurityandthestructureofthefilmsampleswereanalyzed.

Table1 The experimental conditions for the two-step deposition process of cBN films by dc jet plasma CVD.

Step one Step two

Ar (slm) 20 20

N2 (slm) 1.5 1.5

BF3 (sccm) (10% in Ar) 30 30

H2 (sccm) 5 0~2

Substrate temperature (o C) 1000~1040 1150

Pressure (Tort) 50 50

Bias voltage (V) -85 +40 ~ -70

Duration (rain) 10~60 30~50

RESULTS AND DISCUSSION

During the deposition of cBN films in the Ar-N2-BF3-H2 gas system, the fluorine has been demonstrated to act

as an effective etchant that can preferentially etch the hBN phase. If an in-situ etching process by switching off

hydrogen and substrate bias is performed after the deposition, the deposited cBN films will be etched back, and the

phase purity and crystal quality of the cBN films will be improved (ref. 10). To distinguish the origin of the etching,

the in-situ etching was performed under the negative, zero and positive substrate biases after a thick cBN film was

deposited . The SEM cross-section morphologies of the films before and after etching were shown in Fig. 1. Fig. 1

(a) is the film deposited for 1 hour at the gas flow rates of 30 sccm for BF3 and 5 sccm for H2. The film thickness is

about 12 _tm. After the above deposition, the hydrogen flow rate was decreased to zero, and the substrate bias was

switched to zero [Fig. l(b)], 85 V [Fig. l(c)], and +40 V [Fig.1 (d)] for 30 min, respectively. Etching rates of 0.17

_lm/min for no substrate bias, 0.13 _tm/min for negative bias and 0.2 _xm/min for positive bias are obtained from Fig.

1. It can be seen that the etching rate is the highest for the positive bias and the lowest for the negative bias,

indicating that negatively charged fluorine ions might dominate the etching. The fact that the etching rate at zero

substrate bias is higher than that at negative bias suggests that the ion-bombardment induced sputtering by heavy

ions like argon and nitrogen does not contribute to the etching. This observation can also rule out the proposed

growth model of selective sputtering (refs. 11 to 12) in our process.

When hydrogen was added to the gas phase in the second step, both growth and etching occurred depending on

the bias voltage applied to the substrate. Fig. 2 (a) shows the SEM cross-sectional image of the cBN film deposited

for 10 min at the same condition for the sample as shown in Fig. 1 (a). The film thickness is measured to be about 3

_tm. Then hydrogen flow rate ,aas decreased to 2 sccm, and the depositions were performed for 50 min at the

substrate bias voltage of 70, -60, -50, -40, -30, -20 and 0 V, respectively. For the bias voltage of 70 and 60 V, a

further growth of the film was observed, and the growth rate for 70 V was higher than that for 60 V. Fig. 2(b)

NASA/CP--2001-210948 566 August 6, 2001
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the bias voltage of 0 V (b), -85 V (c), and +40 V (d). The hydrogen flow was switched off during the etching.

i"°i::'::::i:')i::S '''" ""i:i:i:_'4"_""_:i_':':_':::'

N.° ._,

._i _::'::" "_

(a) (b)

(c) (d)

Fig. 2. SEM cross-sectional morphologies of the films (a) deposited for 10 min at the bias voltage of-85 V

and the hydrogen flow rate of 5 sccm, (b), (c) and (d) after performing the second-step process at the

hydrogen flow rate of 2 sccm and the bias voltage of-70 V, -50 V and -40 V on the film as shown in (a),

respectively.

NASA/CP--2001-210948 567 August 6, 2001
showsthecross-sectional imageofthefilmdepositedat 70Vforthesecondstep.Athicknessofabout5_tmwas

obtainedafterthetwo-stepgrowth.Thismeansthathefilmgrewforabout2 _tminthesecondstep.Whena

substratebiasof 50V wasappliedduringthesecondstep[Fig.2(c)],thefilmthicknessnearlykeptconstant,

demonstratinghatnofurthergrowthoccurredorthegrowthrateequalstoetchingrateduringthesecondstep.When

thebiasvoltageislowerthan 50V, onlyetchingwasobservedinthesecondstep.TheSEMcross-sectional

morphologyofthefilmdepositedatthebiasvoltageof 40VforthesecondstepisshowninFig.2(d).Thefilm

wasetchedfrom3_xmbackto2_tmduringthisstep.

Toinvestigatethevariationofphasepurityandcrystalqualityaftereachstep,Ramanspectroscopy wascarried

out.Fig.3showsthecorresponding RamanspectraofthesamplesshowninFig.2.Forthesampledepositedbythe

firststep,twointensivepeakslocatedatabout1050.5and1302.7cm-1axeobserved, whichareassignedtoscattering

bythetransverse(TO)andlongitudinal optical(IX))phononmodesofcBN.Thepeakpositionobservedhereagrees

wellwiththatobtainedfromcBNsinglecrystalssynthesizedbyhigh-pressure, hightemperature (HPHT)method 1

(1056 and 1304 cm- for the TO and LO phonon modes, respectively). The full-width at half maximum (FWHW) of

TO and LO phonon modes is measured to be about 25.6 and 18.3 cm-, respectively. The peak positions and FWHM

of the TO and LO modes of the cBN films deposited by the two-step process at different substrate bias voltage are

listed in Table 2. For the sample deposited at the substrate bias of 70 V, the FWHW of both TO and LO modes

decreases. It is known that, the line width and peak position of the peaks depend directly on the crystal size and

defect states in the film (ref. 13). Therefore, the decrease of FWHM demonstrates the improvement of crystallinity

of the film, i.e., increase of crystal size or decrease of defect states in the film. For bias voltage equal or below -50 V

where the films were actually etched back during the second step, the FWHM of both TO and LO modes also

decreased. In this case, the decrease of FWHM due to the increase of crystal size can be ruled out. The decrease of

FWHM should be due to the decrease of defective sites during etching.

"-7

=

Si 0 __J

F-- Z

I i I i I i I i I , I ,

900 1000 1100 1200 1300 1400 1500

Raman Shift (cm 1)

Fig. 3. Raman spectr a of the boron nitride films.

Spectra (a)-(d) correspond to the films shown in Fig. 2 (a)-(d) respectively.

In our previous report, a critical bias voltage was observed for the formation of cBN (ref. 10). Below this value,

only hBN can be deposited. The critical bias voltage decreases as the hydrogen flow rate decreases. Actually, this

critical value can be considered for the nucleation of cBN on hBN. In the present experiments, a critical bias voltage

for the growth of cBN after nucleation is also observed. The deposited cBN films are etched away when the bias

voltage is below this value. On the contrary, the cBN film can go on growing when the bias voltage is above this

value, and the crystallinity or crystal size are evidenced to be improved by Raman spectroscopy.

NASA/CP--2001-210948 568 August 6, 2001
Table 2 Peak position and AFWHM of both Raman TO and LO modes of the samples deposited under

different substrate bias at the second step.

Bias voltage a) Mode Peak Position AFWHM

(v) (cm-1) (cm-1)

One step TO 1050.5 25.6

IX) 1302.7 18.3

-70 TO 1054.4 22.0

IX) 1305.3 12.4

-50 TO 1053.1 24.6

IX) 1304.0 14.1

-40 TO 1051.8 22.7

IX) 1303.0 13.4

a) Substrate bias voltage applied during the second step.

:::::::::::::::::::::::::::::::: ":i::i:.:: ::::::::::::::::::::::::::::::::::::::

';

iiiiii !!ii!:!ii.ll_'. _:.i!iiii_2_ iii!!ii!!iii!iiii:iiii!iii_!!iii_!iii_ili_!!iiiii

:::200:nm::": :::::::::::::::::::::::::: i.i"':':'" " "

Fig. 4. Cross-secti onal TEM images of the films deposited at the substrate bias voltage of (a) -40 V and (b) -

70 V at the second step; (c) SAED pattern taken at the top the column shown in (b). The electron beam size of

about 1 I_m was used for the diffraction.

A columnmc growth behavior was normally observed by both SEM and TEM, and each column was confirmed

to be a single crystal with a vertical length almost equal to film thickness (ref. 14). To investigate the effects of bias

voltage and hydrogen concentration on the film structure, TEM was performed to study the variation of columns

deposited by the two-step process. Fig. 4 (a) and (b) shows the TEM cross-sectional images of the films deposited

by the two-step processes with bias voltages of _40 V and 70 V applied on the substrate during the second step,

respectively. For the bias voltage of 40 V, the columns extend slightly in lateral size from bottom to top, and the

average column lateral size at the film surface is about 200 nm. Compared with our TEM cross-sectional

observation of the film deposited by the first step, no obvious changes of the columnar lateral size can be observed.

For the bias voltage of 70 V, a stronger increase of the columnar lateral size can be observed from the height of

about 3 gm of the column (as indicated by the arrow). Compared with Fig. 2 (a) and (b), this position should

correspond to that the process changed to second step. The column ar lateral size at the film surface is over 1 gm, the

selective area electron diffraction (SAED) pattern taken at the top site of the column verify the nature of single

crystal of the column, which is believed the largest crystal size deposited from vapor phase so far. We also observed

NASA/CP--2001-210948 569 August 6, 2001
thecross-sectional structureofthefilmdepositedbyonestepatconstant biasvoltageandhydrogenconcentration

[sampleasshowninFig.1(a)].Thecolumnlateralsizewasfoundtobeabout500nm,whichisalmostahalfofthe

sizeasthatshowninFig.4(b).Thisconfirmsthatheextraextensionofthecolumniscausedbythegrowthwith

decreased biasvoltageandhydrogenconcentration.

CONCLU_ONS

InthedepositionfcBNfilmsintheAr-N2-BF3-H2 gassystem, thenegativelychargedfluorineionsdominantly

contributeotheetching.Thefactthattheetchingrateat2erosubstratebiasishigherthanthatatnegativebias

suggeststhatheion-bombardment inducedsputteringbyheavyionslikeargonandnitrogenisnotoperativeinthe

process. AcriticalbiasvoltageforthegrowthofcBNafternucleationispresented. ThedepositedcBNfilmsare

etchedawaywhenthebiasvoltageisbelowthisvalue.Onthecontrary, thecBNfilmcangoongrowingwhenthe

biasvoltageisabovethisvalue,andthecrystallinityorcrystalsizeareshowntoimproveasevidencedbyRaman

spectroscopy andTEMobservations.

REFERENCES

1. Konyashin, I.;Bill,J.;andAldinger, F.:Plasma-Assisted CVDofCubicBoronNitride.Chem. Vap.Deposition,

vol.3,no.5,Oct.1997,239,andthereferences therein.

2. Yoshida, T.:VaporPhaseDepositionfCubicBoronNitride.DiamondRelat.Mater., vol.5,no.3-5,Apr.

1996,501,andthereferences therein.

3. Mirkarimi, P.B.;McCarty, K.F.;andMedlin,D.L.:ReviewofAdvances inCubicBoronNitrideSynthesis.

Mater.Sci.Eng.R,vol.21,no.2,Dec.1997, 47,andthereferences therein.

4. McKenzie, D.R.,etal.:Compressive Stress-Induced FormationfCubicBoron-Nitride. DiamondRelat.Mater.,

vol.2,no.5-7,1993, 970.

5. McKenzie, D.R.:Generation andApplications ofCompressive StressInducedbyLowEnergyIonBeam

Bombardment. J.Vac.Sci.Technol. B,vol.11,no.5,Sep.1993,1928.

6. Caxdinale, G.F.,etal.:AnalysisofResidual StressinCubicBoronNitrideThinFilmsUsingMicromachined

Cantilever Beams. DiamondRelat.Mater., vol.5,no.11,Nov.1996,1295.

7. Litvinov,D.;andClaxke, R.:ReducedBiasGrowthofPure-Phase CubicBoronNitride.Appl.Phys.Lett.,vol.

71,no.14,Oct.1997,1969.

8. Litvinov,D.;TaylorII,C.A.;andClaxke, R.:Semiconducting CubicBoronNitride.DiamondRelat.Mater., vol.

7,no.2-5,Feb.1998, 360.

9. Matsumoto, S.;andZhang,W.J.:High-RateDepositionfHigh-Quality, ThickCubicBoronNitrideFilmsby

Bias-Assisted DCJetPlasmaChemical VaporDeposition. Jpn.J.Appl.Phys., vol.39,no.5B,May2000,L442.

10.Zhang,W.J.;andMatsumoto, S.:TheRolesofHydrogenandFluorineinTheDepositionfCubicBoron

NitrideFilmsinAr-N2-BF3-H2 System. Chem.Phys.Lett.,vol.330,no.3-4,Nov.2000, 243.

11.Reinke, S.;Kuhr,M.;andKulisch, W.:Mechanisms inIon-Induced C-BNGrowth. DiamondRelat.Mater., vol.

3,no.4-6,1994, 341.

12.Reinke, S.;Kuhr,M.;andKulisch, W.:CriticalTestofThec-BNSputter Model.Surf.Coat.Technol., vol.

74/75, part2,Oct.1995, 723.

13.Zhang,W.J.;andMatsumoto, S.:Investigations ofCrystallinityandResidual StressofCubicBoronNitride

FilmsbyRammaSpectroscopy. Phys.Rev.B,vol.63,no.7,Feb.2001,73201.

14.Zhang,W.J.,etal.:StructureAnalysisofcBNFilmsPreparedbyDCJetPlasmaCVDFromAnAr-N2-BF3-H2

GasSystem. DiamondRelat. Mater.(inpress).

NASA/CP--2001-210948 570 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ION BEAM DEPOSITION AND SURFACE PROCESSING

OF CUBIC BORON NITRIDE

C. Ronning, H. Feldermann, H. Hofs_iss

II. Institute of Physics, University of G6ttingen, Bunsenstr. 7-9, D-37073 G6ttingen, Germany

ABSTRACT

Cubic boron nitride (c-BN) thin films were grown by mass selected ion beam deposition (MSIBD). For the

nucleation of c-BN we find sharp threshold values of 125 eV for the ion energy and 150 °C for the substrate

temperature. The phase diagram for the nucleation by MSIBD significantly differs from that obtained by ion-beam

assisted deposition (IBAD). In addition, the initial nucleation parameters and the subsequent growth parameters

are not identical for both techniques; i.e. for MSIBD subsequent growth of c-BN takes place over a wider range of

deposition parameters. In order to remove contaminants as well as the sp R-bonded BN-surface layer, which always

covers c-BN films, we investigated the effects of dry etching with low energy fluorine ions as well as of sputter

cleaning with 1 keV argon ions to c-BN surfaces. The sputter cleaning process eliminates the contaminants, but

destroys the spg-bonded BN, whereas the results on dry etching are vice versa. These findings will be discussed in

respect to other surface processing techniques published.

Keywords: cubic boron nitride, ion beam deposition, surface treatment, growth model, etching.

INTRODUCTION

Regarding mechanical and electronic applications, cubic boron nitride (c- BN) may leave diamond far behind,

because it is chemically inert against iron and oxygen at high temperatures [ref. 1] and because it can be doped

p- and n-type [ref. 2]. So far, low-pressure synthesis of c-BN thin films is only possible with a variety of ion

assisted PVD and CVD techniques [ref. 3]. The microstructure of these films usually displays a three phase proression from the substrate to (i) an amorphous layer which contains a mixture of substrate, boron and nitrogen

atoms, (ii) a textured sp2-bonded boron nitride (t-BN) layer and (iii) a textured nanocrystalline (crystallite size: 5-

50 nm) c-BN layer [ref. 3]. Attempts to grow c-BN by chemical processes alone failed so far [refs. 3 and 4]. It is

therefore generally accepted that ion bombardment is necessary for c-BN nucleation and growth; i.e. without ion

irradiation, only h-BN can be formed.

In recent years many investigations have been made to determine the process parameters for the growth and

nucleation of c-BN, which will be reviewed [see also refs. 3, 5 and 6]. The second part of this paper introduces new

results concerning the processing of c-BN surfaces. These experiments were performed because the surface

preparation of c-BN or other substrates is crucial for the realization of direct and continued growth of c-BN.

NUCLEATION AND GROWTH OF c-BN

Several authors have studied the influence of different deposition parameters in ion-beam assisted deposition

(IBAD) on c-BN formation [refs. 7 to 10]. In IBAD of BN films, the growing film is bombarded with nitrogen ions

and additionally with noble gas ions like Ar. It was even stated that Ar ion bombardment is necessary to achieve

c-BN growth or at least enhances the growth significantly [ref. 10]. However, c-BN films have been obtained with

IBAD using solely N ion bombardment [ref. 7] or even neutralized energetic N atoms alone [ref. 11]. In a detailed

systematic study of IBAD deposited BN films by Kester et al. it was found that the h-BN/c-BN growth regime can

be described by the mean momentum transfer per deposited target atom, given in units of _/eV. a.m.u, as a

function of substrate temperature and the B:N arrival ratio [ref. 7]. This dependence is schematically shown in

Figure 1 (solid lines) for B:N 1. For momentum transfer values between 200 and 400 the cubic phase is formed,

whereas for values below 200 only h-BN is formed. Films grown with deposition conditions in between these

regions contain both phases. Above a value of 400 no film growth was observed due to completeresputtering of the

deposited material. The second important parameter is the substrate temperature. No c- BN growth is observed for

NASA/CP--2001-210948 571 August 6, 2001
temperatures below150°C;whereas, filmsdepositedathighertemperatures containanincreasingcontent ofthecBNphase.ArecentstudybyLitvinovandClarke[ref.12]showsthatc-BNfilmscanbegrownattemperatures up

toover1000°C.WithIBAD,filmsareusuallygrownusingevaporated boronatomsandadditional ionbomardent,

typicallywithamixtureofN2+,N+andAr+ions.IBADgrowthisthereforeathercomplexanddifferentprocesses

suchascondensation andthermaldesorption, implantationofions,recoilimplantationf atomsdepositedonthe

surfaceandsputteringhaveto beconsidered. Thec-BNgrowthregimeis thusa complexfunctionof three

parameters: substratetemperature, ionenergyandfluxratioofionstoneutral(boron)atoms[refs7and13].

Inpreviousstudieswecouldshowthatc-BNfilmscanbegrownsolelybythedirectdepositionofenergetic

boronandnitrogenionsusingmassselectedionbeamdeposition(MSIBD)[ref.14and15].In thiscase,the

depositionparameters, likeionenergy,ionfluxratioofdifferentionspeciesandthesubstratetemperaturearewell

definedandindependently controllable[ref.16].Asopposedto IBAD,bothnitrogenandboronaredepositedas

singlychargedenergeticionsandnonoblegasorotherionsnorotherneutralatomsormoleculesareinvolved.

ThisrelativelysimpledepositionprocessmakesMSIBDtotheidealtooltostudytheinfluenceofthedeposition

parametersonc-BNgrowth,leadingto abetterunderstanding of thenucleationandgrowthmechanism. We

observec-BNformationaboverathersharpthresholdvaluesofTs 150°CandElon = 125 eV [ref. 15]. The c-BN

nucleation regime extends up to at least 2.5 keV ion energy. We included the MSIBD nucleation regime of c-BN in

Figure 1 (dotted lines). The most striking differences between our data and the published phase diagrams

describing IBAD boron nitride film growth are the following: With MSIBD

(i) the grown phase depends solely on two well-defined and independently adjustable deposition

parameters (ion energy and substrate temperature),

(ii) the thresholds for c-BN nucleation are sharp and no mixed h- BN/c-BN phase nucleation regime

exists,

(iii) growth of c-BN at high energies is not restricted by a resputter limit,

(iv) the observed threshold energy for nucleation is much lower [ref. 7 and 13].

Considering the complex growth mechanism in IBAD and particularly that IBAD growth takes place at conditions

close to the resputter limit due to the intense heavy ion bombardment, these differences are not surprising. The

corresponding phase diagrams are unavoidably influenced by sputtering processes.

:5--, 600

500

;<

>" 400

o

_, 3oo

q=

200

i/" / / / / /Regi_no 2" /

!' // / / ,/ /Re_utt_'rin_ ,

b/ // ,: / Z/

r z ,7 T z T

.._ / / _ / , / , /

_100 " / • • / I /

:_111 £Ill[llllDllllllDllllll_EIIIllellllln_llllln'elllllnE ,Tucleal_ot7 regime / / /

o

=iexagonal_ , i

o i00 200 300 400 500

substrate temperature [°C]

Figure 1. Comparison of the IBAD (solid lines, after ref. 10) and MSIBD (dashed area, after ref. 15)

h-BN/c-BN nucleation regimes as a function of the average moment transferred to the film atoms and of the

substrate temperature.

Experiments by Hahn et al. and more recently by Litvinov and Clarke show that the energy thresholds for

nucleation of c-BN on the one hand and for its growth on the other are not the same [ref. 12 and 17], which is in

good agreement with our results [ref. 18]. Once initiated, c-BN growth can be maintained at ion energies down to

about 60 eV. Furthermore, a study by McCarty et al. [ref. 19] demonstrates that, once nucleated, continued growth

NASA/CP--2001-210948 572 August 6, 2001
ofc-BNtakesplaceattemperatures aslowas75°C.Evenroomtemperaturegrowthof c-BNhasbeenrecently

provenbyourgroup[ref.20].Therefore, c-BNgrowthisalmostindependent ofthesubstratetemperature.

Thesubstratemperatureof over150°Caswellastheionenergyabove125eVis thereforeonlyof

importanceforestablishingtheinitialboundaryconditionsnecessaryforc-BNnucleation. Theorientedtextured

t-BNintermediatelayer,whichisformedbytheaccompanying highcompressive stressduetotheiondeposition

process, providestheseboundaryconditionsforc-BNnucleation[ref.21].Therefore, c-BNnucleationwithno

intermediatelayershouldbepossibleonsuitablecrystallinesubstrates havingagoodlatticematchandwhichdo

notundergoamorphization underionirradiation. Aluminumnitride(w-A1N)hassuchpropertiesandwerecently

provedthatdirectnucleationandgrowthofspg-bonded BN-phases(c-BNaswellasw-BN)takeplaceonw-A1N

[ref.22].Anothersuitablesubstrateis,ofcourse, c-BNitself;however, c-BNthinfilmsarealwayscoveredwitha

thinsp2-bonded BN-surfacelayeraftergrowthduetothenatureofthedepositionprocesswithenergeticparticles

[ref.5and23].Thethicknessofthissp2-bonded surfacelayerisinthenm-rangeandcorresponds tothemeanion

range(i.e.penetrationdepth)oftheparticlescontributingtothegrowth.If suchafilmisexposedtoairorstored

underbadvacuumconditions(>10-6mbar)thesp2-bondedsurfacelayerrapidlytransformsintoa thicker

BOxCyNz-layer duetotheveryreactiveboron.Aseamless andcontinuedre-growthofthec-BNlayerisnotposible

withsuchasurfacelayerandthecompletenucleationprocesswithitst-BNintermediatelayerhastoberepeated

(ref.[24]).

SURFACEPROCESSING

In order to remove the sp2-bonded BN- or the BOxCyNz-surface layer suitable chemically reagents, sputtercleaning processes, or dry etching methods are necessary. Such a surface process should remove all contaminants

from the surface including the components of the media used, should not destroy the spg-bonding of the c-BN

structure underneath, and should retain the 1:1 BN-surface stoichiometry. In this study, we investigated the effects

of dry etching with low energy fluorine ions (50 eV) as well as sputter-cleaning with 1 keV argon ions. The energy

of 1 keV was used due to the high sputter coefficient. We grew films by MSIBD using 500 eV B+- and N+-ions at

Z

<

(a) as deposited

, I , I , I , I , I ,

to air for 1 week I

O

Ar

/

after 1 keV 4°Ar + bombarde _ment

0 100 200 300 400 500 600

Anger electron energy [eV]

i

n'r r

(d) EELS

• as deposited

o after 1 keY 4°Ar÷

bombardment

500 1000 1500 2000

primary electron energy [eV]

Figure 2. Auger electron spectra of a c-BN film grown by MSIBD recorded (a) directly after growth, (b)

after exposure to air for one week, and (c) after 1 keV argon sputter cleaning. The right plot (d) shows the

plasmon energy as a function of primary electron energy measured with EELS on the same c-BN film

directly after deposition (closed symbols), and after the argon ion bombardment (open symbols).

NASA/CP--2001-210948 573 August 6, 2001
200 °C. These parameters are well within the c-BN nucleation region. Samples were taken out of the deposition

system, stored in air for a certain time (days to weeks), and transferred back into the UHV-system. The c-BN

samples were irradiated with fluorine ions or with argon ions perpendicular to the surface at pressures below 1.10 -s

mbar. After each irradiation dose, the samples were analyzed in-situ with Auger electron spectroscopy (AES) and

electron energy loss spectroscopy (EELS), without breaking the UHV-conditions.

Figure 2(a) shows an Auger spectrum recorded directly after the deposition of the c-BN film. Clearly visible

are the two Auger signals of stoichiometric B and N, respectively, but no features of C or O are present. Therefore,

the grown films are contamination-free. We performed also EELS on this sample with different primary electron

energies. In figure 2(d) the measured _ plasmon energy (closed symbols) is plotted as a function of the chosen

primary electron energy (PEE) and thereby as a function of the film depth. For 2 keV primary electron energy, we

obtain a plasmon energy of 30-32 eV. This high value can be associated with the presence of c-BN with its high

mass density at this depth. EELS-measurements with a PEE of 2 keV on h-BN result into values around 25 eV due

to the much lower mass density [ref. 5, 25, and 26]. For the c-BN film, figure 2(d) shows that the plasmon energy

drops under 25 eV for low primary electron energies down to 100 eV. This is due to the sp2-bonded surface region.

After exposure to air for one week, the sample was transferred back into the UHV-chamber and AES- as well as

EELS-measurements were repeated. The AES-spectrum, displayed in figure 2(b), demonstrates the presence of

high amounts of oxygen and carbon on the surface, and the plasmon energies (not shown) dropped by 0.5 eV for

high PEEs. This indicates that the c- BN film is covered with a low density BOxCvNz-layer, which is thicker than

the sp2-bonded BN-surface layer directly after growth. Figure 2(c) shows the AES-spectrum of the c-BN film

directly after bombardment with 1 keV Ar-ions with a dose of 1.1017 cm -2. It is evident that this procedure results

into a cleaning of the surface. The amounts of O and C have almost disappeared, but residual amounts of Ar could

be detected. However, the EELS-results show that the c-BN phase in the surface region has been destroyed by the

Ar-irradiation. The bulk plasmon energy measured with a PEE of 2 keV has only the h-BN value of 26,5 eV (see

open symbols in Figure 2(d)). Therefore, such a sputter-cleaning procedure is not suitable for uncovering the c-BN

surface. Significantly lower ion energies should be used; however, low sputter coefficients must be taken into

account in this case. Furthermore, nitrogen should be used instead of argon in order to maintain the cleanness and

stoichimetry of c-BN surface. A successful re-growth of c-BN using a mixture of low energy nitrogen and argon

ions as a sputter-cleaning tool has been recently published [ref. 27].

We performed the same experiment with low energy (50 eV) 19F+ ions instead of the argon sputtering, because

l

.d

la

<

C

(b) exposure to air for 1 h

, I , I , I , I,l,I,

F

(c) after 5.0eV '9F+@ e!chmg .....

100 200 300 400 500 600

Auger electron energy leVI

32r

(d) EELS

• as deposited

o after 50 eV 19F+

ions

700 500 1000 1500 2000

primary electron energy leVI

Fignre 3. Anger electron spectra of a c-BN film grown by MSIBD recorded (a) directly after growth, (b)

after exposure to air for one hour, and (c) after 50 eV fluorine dry etch. The right plot (d) shows the

plasmon energy as a function of primary electron energy measured with EELS on the same c-BN film

directly after deposition (closed symbols), and after the fluorine dry etch (open symbols).

NASA/CP--2001-210948 574 August 6, 2001
recent studies suggest that halogen containing plasmas can efficiently etch III-V nitrides [ref. 28 and 29] and

because such low energy ions do not create significant damage. Figure 3(a) shows the AES-spectrum directly after

deposition, and figure 3(d) demonstrates that we started again with a c- BN film, which was covered with C- and Ocontaminants after the exposure to air. Comparing to the argon sputtering, the fluorine dry etch reduced the

amounts of O and C, but did not eliminate them. Furthermore, the fluorine signal appeared at 630 eV in the AES

spectrum, as shown in figure 3(c). On the other hand, the EELS-results (figure 3(d)) clearly prove that the

thickness of the sp2-bonded layer was significantly decreased after the fluorine surface processing and that the sp gbonding structure of the c-BN phase was not destroyed. However, a re-growth of c-BN on top of such prepared

surfaces has failed up to now, maybe due to the residual C- and O-contaminants.

Finally, we show that fluorine is indeed an effective caustic for boron nitride. We grew an about 2-3 nm thick

sp2-bonded BN-layer by the deposition of 0.01 C B+ and N+-ions. In the AES-spectrum, shown in figure 4(a),

signals of B, N, and the underlying Si are visible. Both boron and nitrogen signals completely disappeared with

increasing fluorine etching, as seen in figures 4(b-d). The BN-layer is removed after a total deposited charge of

0.05 C fluorine ions and only a fluorinated silicon surface remained. If we compare the deposited BN and F charge,

we can conclude that about 4-5 fluorine atoms are necessary to remove one BN unit. We suggest that during the

etching process the most stable and volatile BFg-compound is formed, whereas nitrogen disappears as N2-

molecules.

CONCLUSIONS

Cubic boron nitride (c-BN) thin films can be deposited with a variety of different methods and these films

contain c-BN with an amount of up to 90 %. Cubic boron nitride formation is only observed if the growing film is

bombarded with energetic ions or neutral atoms, and it is generally accepted that the impact of energetic ions is

crucial to achieve c-BN. The second important nucleation parameter is the substrate temperature. Both ion energy

and substrate temperature do strongly influence the nucleation but not the growth of c-BN. The substrate temperature of over 150 °C as well as the ion energy above 125 eV is only of importance for establishing the initial

boundary conditions necessary for c-BN nucleation. Therefore, c-BN nucleation with no intermediate layer should

be possible on suitable crystalline substrates having a good lattice match and which do not undergo amorphization

under ion irradiation. Direct nucleation and growth of sp 3-bonded BN-phases (c-BN as well as w-BN) take place

N 42%

B 55%

100 200 300 400 500 600 700 100 200 300 400 500 600 700

, - , - , - , - , - , - , - , , , - , - , - , - , - , - , - ,

F 25%

N 22%

C 5% 0 3%

Si 23%

(a) as deposited "(c) after 0.035C, 50eV 19F+ irradiation

p,_ , - , - , - , - , - , - , - , L, - , - , - , - , - , - , - ,.

"_ N 32%

°= Si 3% / / <

B 44% F 18%

19

(b) after 0.02C, 50eV F ÷ irradiation

100 200 300 400 500 600 700

Si 56% F 26%

. (d) after 0.05C, 50eV 19F+ irradiation

100 200 300 400 500 600 700

Auger electron energy [eV]

Figure 4. Auger electron spectra of a BN film grown by MSIBD recorded (a) directly after growth, (b)

after 0.02 C, (c) after 0.035 C, and (d) after 0.05 C fluorine dry etch (50 eV).

NASA/CP--2001-210948 575 August 6, 2001
onw-A1N[ref.22].Ontheotherhand,c-BNthinfilmsarecoveredwithathinsp2-bonded BNsurfacelayerafter

depositionduetothenatureofthedepositionprocess, andafterexposuretoairc-BNfilmsarecoveredwithanlow

densityBOxCyNz-surface layer.Uncoveringc-BNsurfacesi ratherdifficultandweintroducedtwopossible

methodstoachievethisgoal.Thesputtercleaningprocesswith1keVargonionseliminatedthecontaminants, but

destroyedthespg-bonded BN;whereas, theresultsondryetchingwithlowenergyfluorineionswereviceversa.Up

tonow,there-growthofc-BNontopofc-BNwasonlyrealizedwithamixtureoflowenergynitrogenandargon

ions[ref.27].Onefurtheralternativehasrecentlybeenpublished, whichwewouldliketo mention:cleaningand

etchingcanalsoberealizedwithanexposureof c-BNfilmsto hydrogen-plasmas [ref.23,30to 32].No

experiments towardsre-growthhavebeenconductedafterthisprocedure; however, c-BNsurfacesshowanegative

electronaffinity,whichmaybeduetotheH-covering.

ACKNOWLEDGEMENTS

WewouldliketothankS.Eyhusenforreadingthemanuscript andvaluablediscussions. Financial supportby

theDeutscheForschungsgemeinschaft (DFG)undergrantRo1198/4-1,2isgratefullyacknowledged.

REFERENCES

1. Pouch, J.J.; Alterovitz S.A. (eds.), Synthesis' and Properties of Boron Nitride, Mat. Sci. For. 54&55

(1990).

2. Mishima, O., et al., Science 238 (1987) 1818.

3. Mirkarimi, R; McCarty, K.F.; Medlin, D.L.; Mater. Sci. & Eng. R 21 (1997) 47.

4. Bohr, S.; Hanbner, R.; Lux, B.; Diam. Relat. Mater. 4 (1995) 714.

5. Feldermann, H., et al., in 'Itard coatings", A. Kumar (ed.), TMS conference proceeding (1998) p. 143.

6. Ronning, C., et al., Diam. Rel. Mater. 9 (2000) 1767.

7. Kester, D.J.; Messier, R.; J. Appl. Phys. 72 (1992) 504.

8. Ichiki, T.; Momose, T.; Yoshida, T.; J. Appl. Phys. 75 (1994) 1330.

9. Ikeda, T.; Appl. Phys. Lett. 61 (1992) 786.

10. Wada, T.; Yamashita, N.; J. Vac. Sci. Technol. A 10 (1992) 515.

11. Lu, M., etal., Appl. Phys. Lett. 64 (1994) 1514.

12. Litvinov, D.; Clarke, R.; Appl. Phys. Lett. 71 (1997) 1969.

13. Reinke, S.; Kuhr, M.; Kulisch, W.; Diam. Relat. Mater. 4 (1995) 272.

14. Hofsfiss, H., etal., Appl. Phys. Lett. 67 (1995) 46.

15. Hofsfiss, H.,etal., Phys. Rev. B 55 (1997) 13230.

16. Hofsfiss, H., etal., Diam. Relat. Mater. 3 (1994) 137.

17. Hahn, J., etal., Surf. Coat. Technol. 92 (1997) 129.

18. Hofsfiss, H.; Eyhusen, S.; Feldermann, H.; Ronning, C. ; to be published.

19. McCarty, K.F., etal., Diam. Relat. Mater. 5 (1996) 1516.

20. Feldermann, H., etal., Appl. Phys. Lett. 74 (1999) 1552.

21. Hofsfiss, H., et al., Appl. Phys. A 66 (1998) 153.

22. Feldermann, H., et al., submitted to J. Appl. Phys.

23. Reinke, R; et al., J. Appl. Phys. 88 (2000) 5597.

24. Hofsfiss, H.; etal., Nucl. Instr. & Meth. B 106 (1995) 153-158.

25. Widmayer, R, Ziemann, R Boyen, H.G.; Dia. Rel. Mater. 7 (1998) 385.

26. Widmayer, R, et al., Phys. Rev. B 59 (1999) 5233.

27. Boyen, H.G., et al., Appl. Phys. Lett. 76 (2000) 709, and presented at the spring meeting 2001 of the

German Physical Sociaty (DPG).

28. Tempez, A., etal., J. Vac. Sci. Technol. 17 (1999) 2209.

29. Kalss, W.; Hanbner, R.; Lux, B.; Diam. Relat. Mater. 7 (1998) 369.

30. Powers, M.J., etal., Appl. Phys. Lett 67 (1995) 3912.

31. Harris, S.J., etal., Appl. Phys. Lett. 67 (1995) 2314.

32. Konyashin, I., etal., Diam. Relat. Mater. 8 (1999) 2053.

NASA/CP--2001-210948 576 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

ELECTRON FIELD EMISSION OF CUBIC BORON NITRIDE THIN FILMS

P. A. Enearnaeidn*, and Yo Sheng

Harrison M. Randall Laboratory of Physics, The University of Michigan, Ann Arbor, MI 48109-1120

A. Daniel, and R. Clarke

Applied Physics Program, The University of Michigan, Ann Arbor, MI 48109-1120

EXTENDED ABSTRACT

Field emission results from an on-going study of thin films of polycrystalline cubic boron nitride (cBN) are

presented. Cubic BN has a high thermal conductivity, dielectric breakdown strength, and perhaps the largest

bandgap of the III-V Nitrides. These properties make it an attractive material for use in high power and/or high

temperature electronic and optical devices. The physical similarities of cBN with Diamond in band gap type

(Indirect), resistivity, and the density. Provide further encouragement to investigate the potential of the cubic boron

nitride thin film as, among other things, a field emitter. Samples under study have been synthesized using the

reduced-bias ion-assisted sputtering technique (ref. 1), which enables the growth of 100% cubic phase up to 2gm in

thickness.

For thin film deposition, a custom designed ultra high vacuum chamber, with an in-situ reflection high energy

electron diffraction system for analysis of the film structure, is utilized. Substrates are introduced into the system

individually, through a turbo-pumped load-lock chamber, which typically reaches pressures of - 1 x 10 -7 Torr before

the substrate is transferred to the adjacent UHV growth chamber. The growth chamber is cryo-pumped and can

10 -1° achieve a base pressure of 9 x after bakeout. The UHV environment allows for the thermal desorption of any

silicon substrate oxide prior to growth, and also ensures a clean environment for film deposition.The reactive

sputtering of a hexagonal BN target of 99.99% purity performed using a UHV 2-inch planar RF magnetron source

mounted in a source-up configuration on the base flange assembly supplies source material for the growth of the

cubic thin films. RF power at 13.56 MHz is supplied to the sputtering source. A high field Neodymium Iron Boride

magnet provides additional electron confinement to enable source operation at pressures down to -8 xl0 -4 Torr in

an Argon/Nitrogen gas environment. A Tectra Ion Plasma source is also mounted on another port of the base flange,

at an optimal angle and distance from the substrate to immerse it in the Nitrogen plasma it generates while being well

removed from the plasma confinement region. This additional source of Nitrogen ions to the growth environment

helps to maintain the proper stoichiometry of boron and nitrogen for optimum cBN formation. In addition, the

substrate is resistively heated up to 1100 °C, and a negative DC potential of 100 V is applied to the substrate to

optimize the incident energy of the ions.cBN films grown by this technique have been shown to exhibit a (100) outof-plane texture (ref. 2), which differs from that of cubic boron nitride films grown using other techniques (refs. 3 to

5). Various substrates have been utilized for this experiment: Si (100), etched polysilicon, and Si nanotip field

emitter arrays. Reflection High Energy Electron Diffraction (RHEED) was used to characterize the surface

morphology of the deposited films in situ; RHEED analysis of the films grown is performed using a Vieetech 30 keV

electron gun and a phosphor screen. A CCD detection and analysis system (k-Space Associates, Model KSA 300) is

used for diffraction pattern image acquisition and analysis. The image is created on the phosphor as the electron

beam is incident along a high symmetry direction of the crystal surface, creating a diffraction maxima at a position

on the phosphor which satisfies the Laue diffraction condition. Typical information which can be obtained includes

the orientation of the crystal, or a selected grain of a mosaic crystal, the crystal structure and the morphology of the

sample surface. Generally speaking, angular spread (or rings) in the diffraction pattern corresponds to IN-PLANE

variations of the crystal, which are typically observed with a polycrystalline thin film surface. Radial extent of the

diffraction maxima generally correspond to out-of-plane variations in the crystal structure of the surfaceScanning

Electron Microscopy (SEM) and Atomic Force Microscopy (AFM) data of the cBN thin films were obtained outside

the growth environment. With the use of a Digital Instruments Nanoscope Atomic Force Microscope in contact

mode, we obtained a Mean Roughness of 5.435 nm over an area of l_m 2, at a Height Scale of 144.4 nm. The aspect

ratios of the features would indicate significant field enhancement for Field Emission Measurements done on these

types of films.

NASA/CP--2001-210948 577 August 6, 2001
FourierTransformInfraredSpectroscopy (FTIR)providedcrystalphaseandcompositioninformationof the

boronnitridefilms,withthehighintensitypeakcorresponding tothecubicBNzonecenterTOphononabsorption,

typicalofthetetrahedral sp 3 bonding of the cubic zinc-blende structure, being observed. A Nicolet Magna 550

Series II spectrometer was employed in the acquisition of the FTIR data, which uses interferometry as opposed to the

conventional dispersive spectrometer technique. The spectrometer was configured for mid-IR operation, with a

deuterated tri-glycerine sulphate detector, and a KBr beam splitter. Background spectra was collected using a bare

subsection of the Si substrate on which the film was grown. In this manner, the effects of the silicon substrate can be

removed from the sample spectra. Both background and sample spectra were collected using a 0.06cm -1 resolution

over the spectral region of 400 cm -1 to 4000 cm -1. 1000 scans were typically averaged to obtain the final spectrum.

Samples were mounted on a 2-Circle rotation stage so that measurements may be be performed at oft-normal

incidence. The spectrometer was purged continuously with dry nitrogen to remove water and carbon dioxide

absorption from the spectra.The field emission of the BN films grown on the different substrates was examined using

a customized parallel plate configuration at pressures of 10 .8 Torr or lower. Preliminary data of the films grown on

Si (100) substrates suggests a low emission threshold of -2.75 V/_m, with a current of-10 _A at -3.5V/_m.

ACKNOWLEDGEMENTS

This work is supported in part by ONR Grant N00014-91-J-1398.

REFERENCES

1. D. Litvinov, C. A. Taylor II, and R. Clarke, Diamond Relat. Mater. 7 (1998) 360.

2. D. Litvinov and R. Clarke, Appl. Phys. Lett. 47 (1999) 955.

3. H. Hofs_iss, C. Ronning, U. Griesmeier, M. Gross, S. Reinke, and M. Kuhr, Appl. Phys. Lett. 67 (1995) 46.

4. A. Ballal and L. Salamanca-Riba, Thin Solid Films 224 (1993) 46.

5. H. Ltithje, K. Bewilogua, S. Daaud, M. Johansson, and L. Hultman, Thin Solid Films 257 (1995) 40.

Keywords: Cubic boron nitride; Field emission; Ion-assisted reduced-bias RF Sputtering; RHEED; FTIR

spectroscopy

*Corresponding author. E-mail: pedroe@urrfich.edu. Fax: (734) 764-2193. Phone: (734) 764-7446.

NASA/CP--2001-210948 578 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

RECTIFYING CHARACTERISTICS OF W/P-TYPE C-BN FILMS

K. J. Liao, W. L. Wang, C.G.Hu, J.Xu

Department of Applied Physics, Chongqing University, Chonqging 400044, R R. China

ABSTRACT

Cubic boron nitride thin films were grown on Si substrate using r.f. plasma CVD thermally assisted by a tungsten

filament. The films obtained were characterized by scanning electron microscopy, Infrared absorption spectra and

electron energy loss spectroscopy. P-type semiconducting c-BN films were obtained by introducing Be powder into the

deposition chamber. The rectifying properties of W/p-type c-BN films were observed in the 200-580 °C range. The

experimental results have shown that the diode characteristics of c-BN films strongly depended on the deposition

conditions such as filament temperature. This is due to the content of c-BN in the films was changed with the variation

of the filament temperature.

Keywords: c-BN films; Semiconducting properties; Plasma CVD; Mobility, Rectifying properties

INTRODUCTION

Cubic boron nitride (c-BN) is similar to diamond with outstanding physical and chemical properties. The special

advantages of c-BN are its inertness, even at a high temperature, which can be used for ferrous metal working tools and

other wear applications. In contrast, these are not ideal for diamond (1-4). Experimental studies showed that there are

possibilities of p-and n-type doping (5). There are many potential applications including protective coatings, coated tools,

optical windows, heat sinks, and high temperature electronic devices. C-BN can be prepared by various chemical and

physical vapor deposition, such as ion-beam-assisted physical vapor deposition, r.f sputtering, activated reactive

evaporation, MW plasma CVD, ECR plasma chemical vapor deposition, ion-induced CVD, r.f plasma CVD, inductively

coupled plasma and d.c plasma CVD . But, to date, the application of c-BN films in the electronic industry has made

slow progress. This is due to that there are a lot of defects and grain boundaries in the c-BN films, resulting in low

carrier mobility (6).

In this paper, the microstructure and semiconducting properties of c-BN films were investigated by infrared

absorption spectra, electron energy loss spectra. The experimental results showed that the microstructure of c-BN films

has a considerable effect on its semiconducting properties.

EXPERIMENTAL DETAILS

The c-BN films were grown by r. f. plasma CVD thermally assisted by a tungsten filament (7). The deposition

system consisted of two main parts, i. e. plasma chamber and a substrate-heating chamber with a tungsten filament set

above the substrate. Be material was put in a ceramic boat in the plasma chamber for p-type doping. The chamber was

evacuated to a pressure of lxl0-3pa by the diffusion pump system. B_H_ and NH3 diluted with H2 were used as reactive

gases. Their concentration in hydrogen was 1%, and B_H_ to NH3 ratio was 1:3. Substrate material was P-type Si ( 100 )

with the resistivity of 0.01_-cm. The reactive gases were injected into the chamber when the filament and substrate

temperatures were 2200°C and 800°C, respectively. The gases were excited into the plasma state by r. f induction of

13.65 MHz and 100-200W. The gas flow was 100 sccm. The excited plasma was further thermally activated by the

heating of the tungsten filament, and beryllium materials were evaporated by heating. The Be-doped c-BN films were

deposited on the silicon substrate. The growth time was lh, and the film thickness was about 5500A.

The c-BN films obtained were characterized by scanning electron microscopy, infrared absorption spectra and

electrical measurements.

RESULTS AND DISCUSSION

Figure 1 shows SEM image of c-BN containing films. As can be seen, there is not a flat surface on the films, which

exhibit a cauliflower structure. The films consisted of fine grain with disorder form. The morphologies of the films were

changed little with Be doped concentration in our observed ranges. This may be ascribed to the low resolution of our

scanning electron microscopy. It was found by AES analysis that the predominant composition in the films is boron and

NASA/CP--2001-210948 579 August 6, 2001
nitrogen. Therearealsolittlepuritiesuchascarbonandoxygenearthefilmsurface. Themeasurement bySIMS

showedthathepuritycontent inthefilmsislowerthan0.1ppm.

Figure2showstheinfraredabsorptionspectraofBNfilmsdepositedatdifferent filamenttemperatures rangedfrom

1500to2200°C. Ther.f.powerandgaspressurewerekeptat150Wand70Pa,respectively. Thec-BNandh-BNhave

strongabsorptionbandsat1080cm1and1380cm1(c-plane) respectively. TheIRabsorptionintensitywaschangedwith

thefilamenttemperature. Thec-BNcontent inthefilmswasrapidlyincreasedwithincreasingfilamenttemperature. This

isduetovariouschemical speciesintheplasma, whichwerepromotedtodissociatethemintoatomichydrogenand

hydroboron(hydro-nitride) radicalsathightemperature, andenhancingchemical reactiontoformc-BNcontainingfilms

onthesubstratesurface. Ingeneral, atomichydrogenplaysasignificant roleduringthegrowthprocessofc-BNfilms.

Theexperimental studiesshowedthatheatomichydrogenwasetchingpreferentially thehexagonal BN.

t

I.. , ¢,o_"c "_'/_

, \ V \

iSO0 15011 IZOO 91)0 6BO

Wavenumber (cm -_)

Figure 1. SEM image of c-BN films Figure 2.The infrared absorption spectra of BN

films at different filament temperatures

Be doped c-BN films were obtained by introducing Be

powder in the CVD deposition chamber. During the growth

process, Be was etched by hydrogen plasma to form Be

hydrides. The Be hydrides entered the plasma, and Be was

incorporated into the c-BN films. The doping level in the

films was about 17-120 ppm measured by secondary ion

mass spectroscopy (SIMS).Hall effect measurement yields

carrier-type concentration and mobility. The experimental

results indicated that the Hall mobility in Be doped samples

was decreased with increasing carrier concentration as

shown in Fig.3. The Hall mobility with the carrier

concentration of 4x1018 cm 3 was 215 cm2v is 1 at room

temperature. From Fig.3, it can be seen that the p-type

semiconduting c-BN films could be achieved by beryllium

doping.

The diode characteristic of c-BN containing films was

investigated. Schottky diodes were formed by depositing

aluminum or tungsten on c-BN films through a shadow

mask. The contact area was about 0.85mm 2. The rectifying

proper ties of W/p-type c-BN films were observed in the

200-580°C range. The low level of leakage current was

_E

:U

300_,,, Pilomeai temperature •

?sol ""+N ooo

260

24O x_,

220_

200 , "............................... :

lO is 1016 10 I: lOl_

Hole Concentration (cm "3)

Figure 3. The Hall mobility in Be doped c-BN

films as a function of carrier concentration

observed for the reverse biases up to 100V as shown in Fig.4. The actual breakdown voltage exceeded 120V at room

temperature, and 90-110V at 170°C. The activation energy was about 0.7eV according to Arrhenius law for the increase

of the reverse current with temperature.

NASA/CP--2001-210948 580 August 6, 2001
O

["..................... (a) (b) (c)

-600f l I /

itt/

, 1 , j .

t211 gfl 411 .-40 -80

Voltage _'1

Figure 4. I-V characteristics of a W contacts

to Be-doped c-BN containing films

grown on P-type Si Substrate (a)

500°C; (b) 170°C; (c) 25°C.

-611Ui

-4001

-211(q:

i

t

](el // i. -/"

2001 (IJ_ (a)

400[ .............

120 gO

(z) tb) re)

1 /

1 t

// I

//, I

/ *

/

i [,'/ !i

-4O -80

Vcdtage (V)

Figure 5. The changes of the diode

characteristics of c-BN films with

filament temperature, filament

temperature; (a) 1700°C; (b)

1900°C; (c) 2100°C.

Fig.5 shows the changes of the diode characteristics of c-BN films with the filament temperature. The leakage current

of the diode was increased with decreasing filament temperature. This is owing to the structure of c-BN films which

changed with the filament temperature. Figure 6 shows IR absorption intensity ratio as a function of the filament

temperature. The IR absorption intensity ratio is intensity ratio ICON (1080cml) / Ih BY

3.0

2:

Figure 6.

2.5

2.0

1.5

1.0

Filament Temperature ( °C )

The infrared absorption intensity

ratio IC-BN/Ih-BN as a function of

the filament temperature.

3O

.-. 20t

8

L) 10;

i; ,,_( _perailJle

0.0

180

Figure 7.

" _00 :,'0 240

E_ergy (eV)

The electron energy loss spectra

of c-BN containing films at

different filament temperature.

NASA/CP--2001-210948 581 August 6, 2001
(1380cm1).All depositionconditionsremainedunchangedexceptthefilamenttemperature. It is apparent thatthe

content ofh-BNwasincreasedwithdecreasing filamenttemperature. Figure7represents theelectronenergylossspectra

(EELS)ofc-BNfilmsatdifferent filamenttemperatures. ThetestsamplesarethesameasthatinFig.6.The_z*peakis

characteristics ofh-BNinFig.7.The_z*characteristic peakofh-BNbecameverystrongwithdecreasingthefilament

temperature.

CONCLUSION

Wehaveinvestigatedhesemiconducting propertiesofc-BNfilms.Thesefilmsweregrownbyr.fplasmaCVD

thermallyassistedusingatungstenfilament. P-typefilmswereachievedbyintroducingBepowderintothedeposition

chamber. Schottkydiodeswereformedbydepositingaluminumortungstenonc-BNfilmsthroughashadowmask.The

lowlevelofleakagecurrentwasobservedforthereversebiasesupto100VTherectifyingpropertiesofW/p-typec-BN

filmswereobservedinthe200-580°Crange. Theexperimental resultshowedthatheeffectofthedepositionconditions

onthesemiconduting propertiesofc-BNfilmsisconsiderable. Toachievepracticalpplications ielectronicdevices, it

isnecessary togrowhighqualityc-BNfilms.

REFERENCES

1. Saiton H, Hirotsu Y, Ichinose Y: Conditions for formation of BN thin films by thermally assisted RF plasma CVD.

J.Japn.Inst.Mater. 54(2), 1990, 186-192.

2. Saito H, Hirose T, Matsui H, et al. Synthesis of BN films by the plasma CVD with vapor solids: BH3NH3, H2BO3

and NaBH4. Surf.coat.Technol. 39/40, 1989, 265-273.

3. Kuhr M, Reinke S, Kulisch W: Nucleation of cubic boron nitride (c-BN) with ion-induced plasma enhanced CVD.

Diamond Relat.Mater. 4, 1995, 375-380.

4. Yap Y.K, Aoyama T, Wada Y, et al. Growth of adhesive c-BN films on a tensile BN buffer layer. Diamond

Relat.Mater. 9, 2000, 592-595.

5. Hun S.M: Progress in c-BN semiconducting materials. J Chengdu Univ. Sci.& Technol. 45(3).Mar. 1989, 25-28.

6. Reinke S, Kuhr M, Kulisch W, et al : Recent results in cubic boron nitride deposition in light of the sputter model.

Diamond Relat Mater. 4, 1995, 272-283.

7. Wang W,L. Liao K.J, Wang S.X, et al. Microstructure and semiconducting properties of c-BN films using r.f.

plasma CVD thermally assisted by a tungsten filament. Thin Solid Films. 368, 2000, 283-286.

NASA/CP--2001-210948 582 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

SOME PROPERTIES AND CUTTING PERFORMANCE OF POLYCRYSTALLINE

CUBIC BORON NITRIDE WITHOUT ADDITIVES

Tadakazu Ohashi, Hidetoshi Nakajima, Yoichi Hamada, Katsuyoshi Omino and Kazuo Yamamoto

Mitsubishi Materials Corporation Central Research Institute, 1-297 Kitabukuro-cho, Omiya, Saitama 330-8508,

Japan.

ABSTRACT

Cutting tools made from the cubic Boron Nitride (cBN)-ceramics or metal composites are being widely used for

cutting hardened steels and others. It is expected that products containing higher volume fraction of cBN will show

better cutting performance. In this regard, polycrystalline cBN solids (PCBN) with no additives have been developed

by many techniques, but their cutting performance is not well evaluated yet. In this study, effects of cBN contents for

cutting performance of PCBN's were examined by comparing the cutting performance of chips made from PCBN's

with those of chips made from conventional cBN-ceramics or metal compacts. Wear resistance is proportionally

related to the cBN content of tool material for cutting gray cast iron (FC300, cutting velocity Vc=20m/s), Co based

alloy (Stelite No.6, Vc=0.83m/s) and Ni based alloy (Inconel 718, Vc=l.67m/s). PCBN shows excellent

performance in cutting gray cast iron (FC300, Vc=20m/s), Co based alloy (Stelite No.6, Vc=0.83m/s) and Ni based

alloy (Inconel 718, Vc= 1.67m/s). On the other hand, PCBN has inferior performance to the conventional cBN

composite in cutting hardened steel (SCM415).

Keywords: cBN, property, cutting, super-alloy, cast-iron.

INTRODUCTION

The cubic Boron Nitride (cBN) is the second hardest material next to the diamond, and possesses many

excellent physical and chemical properties similar to those of the diamond. The cBN-ceramics or metal composites

produced under pressures form about 4 to 6GPa and at temperatures from around 1300 to 1800K have superior

mechanical properties, and cutting tools made from them are being widely used for cutting hardened steels, cast iron

and others. In order to improve cutting performance of cBN's tool tips, tips made from products of different cBN

contents, and different crystalline grain sizes, and different kinds of binders are being produced by several

manufactures. Mechanical properties of cBN are better than those of binder phases. Therefore, it is expected that

products containing higher volume fraction of cBN will show better cutting performance. In this regard,

polycrystalline cBN solids with no additives have been developed by many techniques (ref. 1 to 4), and some of their

cutting performances were evaluated (ref. 5 to 6). But effects of cBN contents for cutting performance are not

obvious. Recently we have developed a manufacturing technique of polycrystalline cBN with no additives (hereafter

PCBN) in a disk form. In this process, disks of pBN (stable form at ambient temperature) were processed under high

pressures and at high temperatures. The starting pBN used was prepared by the chemical vapor deposition (CVD)

method. In this study, effects of cBN contents for cutting performance of PCBN's were examined by comparing the

cutting performance of chips made from PCBN's with those of chips made from conventional cBN-ceramics or metal

compacts.

EXPERIMENTAL PROCEDURE

PCBN were prepared under a pressure of 6.8GPa and at temperatures of 2073K, 2373K, 2573K and 2773K

respectively, and the heat-treatment time was 1800s. The thickness of starting disks was 1.5mm, and the highpressure press used in this study was developed by NIRIM (National Institute for Research in Inorganic Materials).

After the production of disks, their upper and lower surface were ground by a diamond grind wheel (#400), and then

rapped by a diamond paste of particle size below 2,um. Compositions of all samples were analyzed by a X-ray

diffract meter (XRD), and their fractured surface and rapped surface were etched by a molten NaOH for 60s and

these surface were examined by a scanning electronic microscope (SEM) and a transmission electron microscope

NASA/CP--2001-210948 583 August 6, 2001
(TEM)respectively, andtheirVickershardness wasalsomeasured. Thesynthesis conditionofPCBNsuitablefor

cuttingtestwasdetermined fromtheseresults.

PCBN,85vol%cBN-metal composite, 60vol%cBN-ceramics compositeand50vol%cBN-ceramics composite

wereevaluatedfortheircuttingperformance. AndthenVickershardness (load--19.6N), thermal conductivity and

fracturetoughness weremeasured. Fracturetoughness wasmeasured bytheIndentation-Fracture method(load--49N).

Afterthesemeasurements, sampleswereheatedinvacuumattemperatures of1073K,1273K,1473Kand1673Kfor

1800s. Afterheattreatments, Vickershardnesswasmeasuredagain(load=9.8N).

Cobasedalloy(SteliteNo.6),Nibasedalloy(Inconel 718),andgraycastiron(FC300)andhardenedsteel

(SCM415,HRc=61)wereusedasworkpieces. Thecompositions ofworkpieceweregiveninTable1.Andthe

cuttingconditionwasgiveninTable2.Thewearresistancewasdetermined basedonthewidthofwearoftheflank

face.

Table1. The compositions of work piece (wt%).

Co Cr W C

Co based alloy

Bal 28 4 1

Ni Fe Ti Mo A1 Nb+Ta

Ni based alloy

Bal 19 0.9 3 0.5 5.1

Hardened steel

Fe Cr Mo C

Bal 0.9 to 1.2 0.15 to 0.3 0.13 to 0.18

Si Mn P

0.15 to 0.35 0.6 to 0.85 0.03<

Cast iron

Table 2. The cutting condition.

Cutting velocity (Vc) Cutting depth (ap) Feed rate (f) Coolant

(m/s) (mm) (mm/rev.)

Cast iron 10 and 20 0.1 0.1 DRY

Co based alloy 0.83 and 3.33 0.15 0.05 WET

Ni based alloy 1.67 and 3.33 0.1 and 0.3 0.12 WET

Hardened steel 0.33 to 3.33 0.1 0.1 DRY

Fe based and other compositions are not determined. Grade of iron is determined as

tensile strength by JIS. Tensile strength of FC300 is equal or greater than 300 N/mm 2.

RESULTS AND DISCUSSION

SOME PROPARTIES OF PCBN

A diffraction peak due to the compressed BN (a type of BN phase trapped in PCBN under a high pressure (ref.

2)) was found in the sample prepared at pressure of 6.8GPa and at temperature of 2073K. Above 2373K, only the

cBN phase alone was produced. The SEM observation showed that the sample prepared at 2373K is a unique phase,

i.e., the grain size is below 1/_m. In the sample prepared at 2573K, a partial grain growth was seen, and grains grow

remarkably at 2753K. The TEM observation showed that twinned grains were produced as grains grew. Few defects

were seen before the grain growth, but some defects were also induced in cBN grains after the grain growth. Vickers

hardness of this sample is 52GPa (reduction from 60GPa). Based on the results above, the sample prepared at 2373K

was selected for cutting tests.

NASA/CP--2001-210948 584 August 6, 2001
7O

AJ 60 ,,.

_" 50

40

N 30 .............................. I _ ]" "" "" )_" ..... .... .......... --. -,.

> 20 -_ ,_.

\

X \'O

/ "It

"!

10

0

0 200 400 600 800 1000 1200 1400 1600 1800

Treatment temperature (K)

Figure 1. Vickers hardness of PCBN and conventional cBN composites after heat treatment in vacuum: (Q)

PCBN; (I) BN-H; (0) BN-M; (&) BN-L.

Some properties of PCBN were compared with those of the conventional cBN composite (Table 3). PCBN has

the highest hardness, the highest fracture toughness and the highest thermal conductivity among all samples. In this

study, a sample having higher volume content of cBN possessed higher hardness and larger fracture toughness. After

the heat treatment in vacuum at 1673K, hardness of PCBN remained almost the same as before the heat treatment.

On the other hand, hardness of conventional cBN composites was reduced after the treatment at 1673K (Figure 1).

Table 3. Properties of PCBN and conventional cBN composites.

CBN content (vol%)

Binder phase

PCBN

100

BN-H

85

Co

BN-M

60

TiC

BN-L

50

TiN

Vickers hardness (GPa) 60 40 32 28

Fracture toughness (MPa. m 1/2) 6.82 6.79 5.59 4.93

CUTTING PERFORMNCE OF PCBN

Wear resistance in cutting the gray cast iron was almost same for all sample chips tested at the cutting velocity

(Vc) of 10m/s. At the Vc of 20m/s, difference of flank wear became larger in the chips, and the sample chips having

higher content of cBN showed better wear resistance (Figure 2). PCBN possessed the best wear resistance. In cutting

Co based alloy at the Vc of 0.83m/s, chips of higher cBN content showed better wear resistance, and PCBN

possessed the best wear resistance. At Vc of 3.33m/s, BN-M showed the best wear resistance (Figure 3). In cutting

Ni based alloy at theVc of 1.67m/s, chips of higher cBN content showed better wear resistance, and this result is the

NASA/CP--2001-210948 585 August6,2001
same as for Co based alloy. PCBN possessed the best wear resistance. At Vc of 3.33m/s, BN-L had the best wear

resistance (Figure 4). In cutting hardened steel, wear resistance of PCBN was lower than that of the conventional

composite at all cutting velocities (Figure 5).

0.14

0.12

0.1

_0.08

._ 0.06

0.04

0.02

BN-L BN-M

[] Vc=10m/s • Vc=20m/s

BN-H

Sample

PCBN

0.1

0.08

,_ 0.06

_0.04

0.02

Figure 2. Flank wear in cutting of cast iron.

(a) Vc= 0.83m/s

e _

I I I I

0 200 400 600 800

0.25

Cutting time (s)

0.2

_0.15

0.1

0.05

1000 0

(b) Vc= 3.33m/s

.J

f

I

,,mr

f

i

/,,Jr

/_, ..O

,O'"" .O

/ ,0""

I I I I I I

20 40 60 80 100 120

Cutting time (s)

Figure 3. Flank wear in cutting of Co based alloy: (0) PCBN; (•) BN-H; (0) BN-M.

NASA/CP--2001-210948 586 August 6, 2001
0.8 0.35

c= 1.67m/s, 0.1mm

0.7 0.3

i

_-, 0.6 0.25

0.2

/' / /

0.2 [i m/ _ / 0.1

0.1 _ 0.05

0 0

0 200 400 600 800 1000

Cutting time (s)

(b) Vc= 3.33m/s, ap= 0.3mm

f

,!i. )

I

0 200 400

Cutting time (s)

Figure 4. Flank wear in cutting of Ni based alloy: (O) PCBN; (•) BN-H; (0) BN-L.

600

>

©

>

0

100

A

I

i

i

n

I

4

%

%,

i I I I i i i i n I i i i I i i i ! i i I i

1000 10000

Cutting time when flank wear is equal to 0.2mm (s)

Figure 5. Flank wear in cutting of hardened steel: (0) PCBN; (A) BN-L.

NASA/CP--2001-210948 587 August 6, 2001
Basedontheresultsabove, it isthought thathemechanical propertyofthetooledge,thethermal stabilityand

theaffinitytoworkpiecesdeterminethecuttingperformance. It isalsothought thathecontent ofcBNinatool

material isthemosteffectiveforimprovement ofthemechanical propertyofthetooledgesincecBNhasbetter

mechanical propertyhanthebinderphase. IncuttingraycastironatVcof20m/s, wearresistanceisproportionally

relatedtothecBNcontent intoolmaterials. Theseresultsindicatehathemechanical propertyofthetooledgemay

bethemostimportant factorforwearresistance. Theforceaddedonthetooledgebecomes largerasthecutting

velocityisincreased. IncuttingCobasedalloysandNibasedalloys,thewearresistanceisproportionally relatedto

thecBNcontent alowervelocity, buthewearresistanceathighervelocitydoesnotdependonthecBNcontent in

thetoolmaterials. Thetemperature ofthetooledgeduringcuttingmayincreaseinproportionwithhardness ofwork

piece.If theincreaseoftemperature ofthetooledgeisrelatedwithhardnessofworkpiece,thetooltemperature will

bethehighest incuttinghardenedsteelandit willbethelowestincuttingcastironinthisstudy.Basedonthis

assumption, temperature riseoftooledgeincuttingCoandNialloywillbehigherthanthatofcuttingcastiron.In

thisregard, theaffinityofthetoolmaterialstoworkpiecesmaybemoreimportant thanthemechanical propertyof

tooledgeforthetoolife.Incuttinghardenedsteel,it isthought thathetemperature riseisthehighest inthisstudy.

Affinityoftoolmaterialstoworkpieceisthought asamainfactorfordeterminingtheperformance ofcutting

hardenedsteelevenatlowestcuttingvelocity.

CONCLUSION

1. Wearresistanceisproportionally relatedtothecBNcontent oftoolmaterial forcuttingraycastiron

(FC300, Vc=20m/s), Cobasedalloy(SteliteNo.6,Vc=0.83m/s) andNibasedalloy(Inconel 718,

Vc=l.67m/s).

2. PCBNshowsexcellent performance incuttingraycastiron(FC300, Vc=20m/s), Cobasedalloy(Stelite

No.6,Vc=0.83m/s) andNibasedalloy(Inconel 718,Vc=l.67m/s).

3. PCBNhasinferiorperformance totheconventional cBNcompositeincuttinghardenedsteel(SCM415).

REFERENCES

1. M. Wakatsuki; K. Ichinose; and T. Aoki: Synthesis of Polycrystalline Cubic BN, Mater. Res. Bull., Vol. 7,

1972, 999-1003.

2. F.R. Corrigan: High pressure science and technology, Vol. 1, edited by K. D. Timmerhaus and M.S. Barber,

Plenum, New York, 1979, 994-999.

3. A. Onodera; K. Inoue; H. Yoshihara; H. Nakae; T. Matsuda; and T. Hirai: Synthesis of cubic boron nitride

from rhombohedral form under high static pressure, J. Mater. Sci., 25, 1990, 4279-4284.

4. M. Akaishi; T. Satoh; M. Ishii; T. Taniguchi; and S. Yamaoka: Synthesis of translucent sintered cubic boron

nitride, J. Mater. Sci. Letters, 12, 1993, 1883-1885.

5. T. Ohashi; K. Yamamoto; Y. Hamada; and T. Tanase: Some properties and cutting performance of

polycrystalline cubic boron nitride with no additives, Int. J. of Refractory Metals & Hard Materials, 16,

1998, 403-407.

6. S. Uesaka; and H. Sumiya: Mechanical Properties and Cutting Performances of High Purity Polycrystalline

cBN Compact, Manufacturing Sci. and Engineering, MED-Vol. 10, 1999, 759-766.

NASA/CP--2001-210948 588 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaramnga (Editors)

Cubic Boron Nitride Composite Coating for Advanced Machining Application

Ajay P. Malshe l'*, S.N. Yedave 1, W.D. Brown I and William Russell 2

1Materials and Manufacturing Research Laboratories (MRL), University of Arkansas,

Fayetteville, AR

ZMaterials Research Development (MRD), Valenite Inc., Troy, MI

* Contact Information- Tel (501) 575-6561 E-mail <apm2@engr.uark.edu>

Cubic boron nitride (cBN), known as the second hardest material, has good thermal

stability and chemical inertness for machining ferrous alloys. Such coating, unlike PCBN

bits, on cemented carbide substrates provides the hardness and chemical resistance of

cBN in combination with the toughness of cemented carbide body. Most research groups

working on cBN deposition employ either high-energy plasma or ion bombardment for

the stabilization of the cubic phase. These methods, however, offer the following

challenges in the growth process: (1) excess compressive stress, (2) non-homogeneous

phase synthesis, (3) limited stoichiometry control, (4) inadequate adhesion With substrate,

and (5) phase stabilization on different substrates. These issues are amplified when

considering manufacturing on large scale. One of the important alternatives is to develop

cBN films in combination with other materials to resolve the phase stability and intrinsic

stress issues. The result will be composite films with physical properties dependent on

the material makeup of the system. In a novel successful approach reported here, instead

of attempting to grow cBN from the vapor phase, sub-micron size cBN powder particles

are applied as a conformal coating. Then porous cBN particle matrix is infiltrated using

TiN to synthesize composite coating. We have observed excellent cBN-to-TiN and

coating-to-carbide substrate adhesion, and outstanding machining results.

NASA/CP--2001-210948 589 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

SUPERHARD COMPOSITES BASED ON CUBIC BORON NITRIDE

Anatoly Bochko

Institute of Metallurgy and Materials of National Academy of Science of Russia

Guerold Bobrovnitchii

Laboratory of Advanced Materials of State University of the North Fluminense,

Av. Alberto Lamego, 2000, CCT, Horto, Campos dos Goytacazes, RJ, Brazil,

CEP: 28015-620, Tel/Fax + 55-24-7263733, e-mail: ,_uerold@2,lenEbr

ABSTRACT

During past decade's cubic boron nitride (cBN) and wurtzite boron nitride (wBN) have received considerable

attention. These dense modifications of boron nitride possess hardness approaching to diamond and thermal stability

better than that of diamond. They are chemically inert and do not react with iron. Superhard materials made on their

base found wide application in many areas of metal cutting operations.

The compacts sintered from cBN or wBN powder have a polycrystalline composite structure either of cBN or

cBN plus wBN. In final industrial form they also have high fracture toughness. The cutting tools made of this

material show better service life in interrupted cutting of cast irons and hardened steels.

It is essential to have a proper understanding of this class of superhard material in order to realize its full

potential not only in cutting applications but due to their unique physical properties, also in electronics.

The present paper shows a comparative analysis on modern methods of producing superhard composites based

on cubic boron nitride and wurtzide-like boron nitride. Wide gaining of these materials will help to find for them

various fields of application.

Key Words: high pressure, cubic boron nitride, wurtzite boron nitride.

INTRODUCTION

During past decade's cubic boron nitride (cBN) and wurtzite boron nitride (wBC) have received considerable

attention. These dense modifications of boron nitride possess hardness approaching to diamond and thermal stability

better than that of diamond. They are chemically inert and do not react with iron. Superhard materials made on their

base found wide application in many areas of metal cutting operations.

The compacts sintered from cBN or wBN powder have a polycrystalline composite structure either of cBN or

cBN plus wBN. At last case they have high fracture toughness. The cutting tools made of this material show better

service life in interrupted cutting of cast irons and hardened steels.

It is essential to have a proper understanding of this potential not only in cutting applications, but due to their

unique physical properties also in electronics.

BORON NITRIDE AND ITS MODITICATION: CRISTAL CHEMISTRY CHARACTERISTIES

Boron nitride exists in hexagonal graphite-like form (gBN), hexagonal wurtzite-like form and cubic zinc blend

form or cubic nitride form. Under ambient conditions gBN is a stable phase and wBN and cBN are the metastable

phases. The crystal chemical characteristics of various phases of boron nitride are given in table 1.

NASA/CP--2001-210948 590 August 6, 2001
Table 1. Crystal chemical characteristics of different modifications of boron nitride (ref.1)

Coordination Lattice Parameter Density Modification N ° atoms/unit cell

number " a " , nm c, nm g/cm 3

gBn 3 4 0.2504 0.661 0.29

wBN 4 4 0.2550 0.423 3.50

cBN 4 8 0.3615 ..... 3.51

METHODS OF SYNTESIS OF cBN POLYCRYSTALS

Known processes can be generally classified in four categories:

a) Catalytic conversion process, one-step process in which the catalyst, metal or alloy, aids in the transition of

gBN to cBN simultaneously with the formation of the compact.

b) Bonding medium process, two-step process in which the first step comprises the conversion of gBN to cBN

and the second one comprises the formation of the compact from cleaned cBN crystals mixed with the metal

alloy which aids in the bonding of the cBN to the compact.

c) Direct sintering processes, a two-step process which is the same as the process (b) except that compact is

formed without addition of metal or alloy to aid in bonding cBN crystals.

d) Direct conversion process, one-step process in which substantially pure gBN is directly transformed to a

cBN compact without the aid of the catalyst and/or bonding medium.

Hard phase BN compacts are of two general types: cluster compact and composite compact.

A cluster compact is defined as a cluster of abrasive crystals bonded together either in a self-bonded relationship

(1) or by means of some combination of between the crystal (2) or by means of some combination of (1) and (2). For

example, Borason (ref.2), Elbor-R (ref.3), Belbor (ref.4), Hexanite-R (ref.5), Amborit (ref.6), Kiborit (ref.7) and

others are those cluster compacts.

A composite compact is defined as a cluster compact bonded to the substrate material, such as cemented

tungsten carbide or cemented titanium carbide. The examples of such composite compacts are Compax BZN (ref.8),

DB 50 (ref.9), Sumiboron (ref. 10), Wurzin (ref. 11), BPK (ref. 12), Composit 10D (ref. 14) and others.

The catalytic and bonding medium processes are generally disadvantageous because catalyst and bonding

medium are lower in hardness than cBN and retention in the resultant mass reduce the hardness and abrasive

resistance of the masses.

The direct conversion process, which is theoretically and practically possible, has been found to have high losses

in practice because it is difficult to achieve consistently the sufficient number of crystal to crystal bonds distributed

uniformly within the compact. Without this, the strength and density of the compact are less than it is necessary. And

the no less so significant progress in this direction was reached.

The considerable amount of high pressure sintering of cBN in presence of various solvent catalysts has been

done at the General Electric Company, and then at the De Beers Company. Many different composites were created

in the old USSR. The most famous composites, which were created in "J.E." and De Beers are Borazon and

Amborite. Aluminum has proved to very effective solvent-catalyst for manufacturing of cBN composites. The

residual binder phases in the Borason and Amborite range of products are aluminum nitride and aluminum diboride.

The sintering of these composites is done over a pressure range 5-7 GPa, at temperature range 1500-1200 ° C.

Polycrystalline cBN, tool blanks consist of the layer 0.5-0.7mm of cBN crystals bonded to one another on a

cemented carbide substrate. This composite structure is achieved through high temperature high pressure process,

resulting in an extremely high wear and impact resistant product with very consistent physical properties from blank

to blank. As the cobalt in the WC-Co cemented carbide substrate melts at high pressure-high temperature process,

begins to infiltrate the voids in the compact present in cBN, thus completing the cementation process within the

existing cBN crystals.

Physical properties of hard and superhard materials are presented in table 2.

NASA/CP--2001-210948 591 August 6, 2001
Table 2. Physical properties of hard and superhard materials

Property

Density, 10 kg/m 3

WC+Co

ISO

(ref. 9)

14.7

Syndite

PCD25

(ref. 6)

3.86

Amborit

PCBN

(ref. 6)

3.42

Kiborit

PSBN

(ref. 12, 13)

3.4

Composite

05 PCBN

(ref. 14)

4.3

Compressive strength, GPa 4.5 7.61 2.73 2.9 2.3

Bending Strength, GPa Fracture 2.7 1.19 0.57 --- 0.47

Toughness, MPa m 10.8 8.89 6.30 105 6.7

Knoop hardness, GPa 13 50 32 36 18.8

Young' s modulus, GPa 620 810 680 880 620

Poisson's ratio, GPa 0.21 0.07 0.22 0.16 0.16

CTE, 10 K 5.4 4.6 4.9 ......

Thermal conductivity, w/m.k 100 560 100 100 ---

DIRECT CONVERSIONS gBN - cBN AND gBN - wBN

The direct conversion under static pressure condition of gBN to more dense wBN and CBN phases at pressure of

10 GPa and above described in details by Bandy and Wentorf (ref.5). More recently, numerous reports and patents

have been concerning the direct conversion of gBN to cBN cluster compacts under pressure below 10.0 GPa. In the

works of N.N. Sirota and A.M. Mazurenko (ref.4, 16) it was claimed that superhard material, which might be

synthesized under pressure of 6.0 GPa and higher over the temperature range from 1800°C to 3000°C. Practically at

the same time the works of the Japanese scientists were published (ref.7). In the publication (ref.8) report direct

conversion occurs under pressure more than 5.0 GPa (preferably 6.0 GPa and above) and sintering temperature more

than ll00°C. The material, created by technology (ref.4) received Belbor name and now is wide known in the

industry.

Publication (ref.19) reports a pyrolitic boron nitride (pBN) as the starting material for the synthesis of cBN

cluster compacts in a direct conversion process practiced under pressure of 6.9 GPa and at temperature between 1800

and 1900°C. The resulted product was characterized as the soft mass having large amount of unconverted gBN.

Publication (ref.17, 18) also reports the synthesis of wBN and cBN. There were no reported results of the

successful formation of either wBN or cBN. And only the publication (ref.20) on F.R. Corrigan reported the results

of the successful formation of cBN cluster composite in a direct conversion pBN_BN process under pressure and

temperature between 2190°C and 2400°C without a trace of gBN. Thermal conductivity of this material was equal to

470-700 w/m k (at 100°C) F.P.Corrigan proposed to use this cluster composite as a heat sink. This result was

confirmed later in (ref.21) where thermal conductivity was found to be equal to 200-575 w/m .k.

The direct conversion process of gBN to wBN by dynamical method was fixed for the first time at the end of the

60 Th. (ref.22). Further on this process was examined rather good and at present time is considered to be finally fixed

and that it goes on by martensitic mechanism (ref.23).

COMPOSITES ON BASE OF wBN

The first composite with the use of wBN was the composite (PHBN), made on the bases of the mixture of

wBN+cBN powder (ref.24). The synthesis of the composite on the base of only wBN was the Problem sometime,

until it was solved by the synthesis of Hexanite-R (ref.25).

Converted mechanism of wBN to cBN during synthesis of the composites on the base of wBN, as it is shown in

(ref.23) carries not martensitic, but diffusion character. Physical properties of these composites are given in table 3.

NASA/CP--2001-210948 592 August 6, 2001
Table 3. Physical properties of superhard materials ( direct synthesis).

Elbor-RM Belbor PHBN Hexanite-R

Property (ref. 3, 14) (ref. 4, 14) (ref. 14, 24) (ref. 14)

Density, 10 kg/m 3 3.42 3.45 3.4-3.42 3.1-3.44

Compressive strength, GPa 2.7 4.0-6.5 3.4-4.9 3.5-4.0

Bending Strength, GPa Fracture ......... 4.0-1.2

Toughness, MPa m 4.2 --- 13-17 15-17

Knoop hardness, GPa 35 38 40.6-40.0 41.8

Young's modulus, GPa 840 800 800 750-820

Poisson's ratio, GPa 0.16 ...... 0.16

CTE, 10 K ......... 1.8

Thermal conductivity, w/m.k 60-80 100 70 25-30

DISCUSSION

The comparison of the mechanisms of formation of the composites with the presence of the catalyst of or

bonding binders and composites made by direct conversion of wBN to cBN and also the values of fracture

toughness, presented in table 2 and 3, allows making analogy between hard alloys and superhard composite

materials.

In this one and another case we have hard (wc) or superhard (cBN) particles, surrounded by less hard and more

plastic matrix Co, AIN, AIB2, when meeting with it the cracks, achieved critical size and capable to lead to the

failure of the whole composite, brake.

In case of Hexanite -R this analogy is less obvious than for example, in the case of PHBN (ref.24), where wBN

transform into cBN completely, but in this case such approach is also correct. Increased resistance of Hexanite-R to

impact loading, so very high fracture toughness, can be also explained, by the fact that within structures with

wurtzite structure wBN differs by increased extension of coordination tetrahedron along the axis "c" (c/a =1.6). The

last circumstances allows to consider wBN not only as superhard material similar to cBN, but more plastic

comparatively to cBN. Except the grains of wBN are smaller than the than the grains of recrystallized cBN, what

usually renders them higher fracture toughness.

CONCLUSION

Presented overview of the composite on the base of cubical and wurtzite boron nitride touched only the main

picture about this class of materials.

Performed analysis of their formation mechanism and physical properties allowed to identify those peculiarities,

which make them irreplaceable in metal machining and also notice those, which may be used in electronics.

REFERENCES

1. Kurdjumov A.V., Pilyankevich A.N. Phase transformations is carbon and boron nitride (Russian),

Nauk. Dumka, Kiev.1979, p.198.

2. WentorfR.H., Delai A.J. US Patent N o 3 233 988, 1966.

3. Lysanov V.S (1978) Elbor in mechanical engineering (Russian), Ed. Mashinostroenie,

Leningrad. 1978, p.132.

NASA/CP--2001-210948 593 August 6, 2001
4. MazurenkoA.M.Theinvestigationofthekineticspolymorphtransitionhexagonal BNincubic

modificationaddevelopment equipment forsuperhighpressure(Russian), Thesis, Minsk,1972,

p.158.

5. Francevich I.N.etal.Superhard materials(Russian). Ed.Nauk.Dumka, Kiev.1980, p.202.

6. BossonP.K.Ind.DiamondRev.,50,1990,p.228

7. NovikovN.V.Thesynthesisof superhardmaterials(Russian). Ed.Nauk.Dumka,Kiev.1986,

p.277.

8. Wentorf R.H.andRoccoW.USPatent No3745623.1973.

9. HooperR.M.etal.Ind.DiamondRev.,49.1989.

10.Tabuchi, HaraA.,YaznS.etal.Performance ofSumiboronBN2000, SumitomoElectricTech.

Review, 81978,p.57-65.

11.ArakiM.,SawaokaA.Cuttingtoolsinsertsofhexagonal boronitride,J.Ind.Explos.Soc.Jap.,v.

44,4,1984,p.230-235.

12.ShulzhenkoA.Aetal.Synthesis, sinteringandpropertiesof cubicboronnitride(Russian),

Ed.Nauk. Dumka, Kiev.1993, p.193.

13.GrischenkoE.I.etal.Theturningofnickelalloyswiththetoolofcubicboronnitride(Russian),

Ed.Nauk.Dumka, Kiev.1993, p.178.

14.BochkoA.V. et al. Technological peculiaritiesof mechanical treatmentwith tool of

polycrystalical superhard materials(Russian), Ed.Nauk.Dumka, Kiev.1991,p.181.

15.BundyF.P.,WentorfR.HDirectransformation of hexagonal boronnitridetodenserforms

J.Chem. Phys,38,1963,p.1144-1149.

16.SirotaN.N.Processforproductionfcubicboronitride.BritishPatent No131716.1973.

17.Wakatsuki etal.Synthesis ofpolycrystalline cubicboronnitride,Mat.Res.Bull,7, 1972,p.999-

1004.

18.Ishinoseetal.Synthesisofpolycrystalline cubicBN,Proc.of 4-thInt.Conf.onHighPressure,

Kyoto,Japan, 1974,p.436-440.

19.Wakatsuki etal.Synthesis ofpolycrystalline cubicboronnitride,Proc.of4-thInt.Conf.onHigh

Pressure, Kyoto,Japan,1974,p.441-445.

20.CorriganF.R.Cubicboronnitride(cBN)compact anddirectconversionprocessformakingsame

frompyroliticboronnitride,Patent GB2002333.1978.

21.OsitinskayaT.D.etal.Thethermal conductivityandsomeperculiarities of cubicboronnitride

polycrystals (Russian), Sverkhtverdye materially, 6,1989,p.23-26.

22.DreminA.N.et al.Theformationof wurtzite-likemodificationof boronnitrideat shock

compression (Russian), Dokl.Acad.NaukSSSR, v.172,5,1967,p.1066.

23.KurdjumovA.V.etal. Polymorphicmodificationsof carbonandboronnitride(Russian),

Handbook, Ed.Metallurgy, Moscow. 1994,p.202.

24.VereschaginL.F.Synthetic damondsandhydroextrusion (Russian), Ed.Nauka,Moscow.1982,

p.195.

25.AlexeevskyV.P.,BochkoA.V.,DzhamarovS.S.etal.Methodforproducingofpolycrystalline

boronnitride.USAPatent No3876751, 1975.

NASA/CP--2001-210948 594 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

REACTIVE MAGNETRON SPUTTERING OF HARD C-N AND Si-C-N FILMS AND

THEIR PROPERTIES

J, Vlcek, lVl. Kormunda, J. Cizek and V. Hajek

Department of Physics, University of West Bohemia, Univerzitni 22, 306 14 Plzen, Czech Republie

E-mail: vlcek@kfy.zcu.cz, Fax: +420-19-7422825, Tel.: +420-19-7423136

ABSTRACT

C-N and Si-C-N films were deposited on Si(1,0,0) substrates using dc magnetron sputtering with a

negative bias voltage on the substrate.

In the case of C-N films a high-purity graphite target was sputtered in pure nitrogen at a substrata

temperature of 600 °C. The film characteristics were primarily controlled by the pressure, p, (0.05 - 5 Pa), the

discharge current on the magnetron target, In, (0.5 - 3 A), and the rf induced negative substrata bias voltage,

Ub, (-300 to -1200 V). The films, typically 1-2 gm thick, were found to be amorphous, and they possessed the

N/C atomic concentration ratio (determined by ERD) up to 0.35, the hardness up to 40 GPa, the elastic

recovery up to 85 %, and good adhesion. In order to clarify the complex relationship between the process

parameters and the film characteristics, we systematically studied, using optical emission spectroscopy and

electrical measurements under the same conditions, the density of N atoms and CN radicals near the substrata,

and the kinetic energy and flux of ions bombarding the magnetron target and growing films. Good correlation

between the N/CN concentration ratio in front of the substrate and the N/C atomic ratio in the films was found

over a wide range of the investigated process parameters. We have shown that the C-N films have a high

hardness only when the energy and flux of the nitrogen ions are sufficiently high for effective incorporation of

nitrogen into the layers (and for ion-induced desorption and resputtering of hydrogen-containing species from

them); such conditions require the pressures around 0.5 Pa and lower, and the Ub values between -500 and -

700 V.

Si-C-N films of various compositions were produced by reactive dc magnetron co-sputtering of silicon

and graphite in nitrogen-argon gas mixtures using a single sputter target with different Si/C area ratios. The

total pressure and the discharge current on the magnetron target were held constant at p = 0.5 Pa and Im= 1 A,

and the substrate tempera_re was adjusted at 600 *C by an Ohmic heater or it was in the range from 135 to

210 _C without the heater. The negative substrate bias voltage was varied from a floating potential of about

-20 V to -500 V. The film compositions (determined by RBS and ERD), and hence the surface bonding

structure (XPS), the surface morphology (Ab2vi) and mechanical properties (see our measurements of hardness,

effective Young modulus, elastic recovery and friction coefficient), were primarily controlled by the Si fraction

in the magnetron erosion track area (5 - 80 %) and by the Ar fraction in the gas mixture (0 - 75 %). Depending

on the process parameters, ternary SixCcN z films within the composition range 9 _ x ___35 at %, 6 _ y < 6l at %

and 26 _<z <_52 at % were prepared possessing the hardness between 20 and 40 GPa.

Keywords: earbon-nitride films, silicon-carbon-nitride films, reactive magnetron sputtering, controlled film

composition, hard protective layers.

NASA/CP--2001-210948 595 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

The Speculation of Superdiamonds

James Sung

KINIK COMPANY, Taipei, Taiwan, R.O.C.

Taipei University of Technology, Taipei, Taiwan, R.O.C.

Abstract

Diamond has the highest bond energy per unit volume of all known materials, and hence it is assumed to possess the

highest hardness. Diamond's hardness comes from its small atoms that each of them forms 4 covalent bonds. To make a

structure harder than diamond, its atoms must be smaller than carbon, and/or these atoms form at least 4 covalent bonds.

The first consideration would strike off all elements with period number higher than 2. The second criterion would

eliminate all elements lighter than carbon. Hence, only C, N, O, F, and Ne are possible candidates of superdiamond.

However, in order to become a superdiamond, these elements must form mono-atomic structures with coordination

number higher than 4. Moreover, no lone pair electrons are allowed, so all their valence electrons must be involved in

single covalent bonds.

The number of valence electrons in simple cubic carbon is less than the coordinated number of 6. As a result, the

bonds may turn metallic, so it may not be as hard as diamond. Potential superdiamond structures include diamond-like

nitrogen, simple cubic oxygen or fluorine, and body center cubic (BCC) neon. If these elements can form single

covalent bonds that involve all their valence electrons, they could become superdiamond. Otherwise, diamond's

hardness may be as insurmountable as the speed of light.

The above hypothetical structures of superdiamond may be synthesized by directing collimated beams of

single ions that are converging from the intended directions aiming at a common center. Such a technique was

developed by Nobel Laureate Y. T. Lee decades ago. The possible instantaneous formation of the predicted

hypothetical structures, even though they may be highly metastable, can be studied in-situ by laser strobe light flashed

at femtoseconds (10 a5 see). Such femtochemistry has already been invented by Dr. Ahmed Zewail, the latest Nobel

Laureate.

Key Words: Superhard Materials, Carbon Nitride, Simple Cubic Carbon, Simple Cubic Oxygen, Body-Centered

Neon, Superdiamond, Femtochemistry.

NASA/CP--2001-210948 596 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

SYNTHESIS AND DIELECTRIC PROPERTIES OF CARBON NITRIDE THIN FILMS

William C. Lanter

Innovative Scientific Solutions, Corp., 2766 Indian Ripple Rd., Beavercreek, OH 45440

David C. Ingrain

Edwards Accelerator Laboratory, Ohio University, Athens, OH 45701

Charles A. De Joseph

Air Force Research Laboratory, Wright-Patterson AFB, OH 45433

Richard LC. Wu

K Systems Corp., 1522 Marsetta Dr., Beavercreek, OH 45432

ABSTRACT

Thin films of amorphous Carbon Nitride (CN_) films have been directly deposited onto silicon (Si), quartz (Qz),

Coming 7059 Glass (G1), aluminum (AI), and glassy carbon (GC) substrates utilizing a RF ion beam source.

Varying concentrations of methane and nitrogen were used as the feed gas. The effects of RF discharge power

(100-600W), ion beam deposition energy (100-1750eV), total gas flow (6-36seem), and the N2/CI-h gas ratio

(0.2-3.0) on the film growth rate and the corresponding dielectric properties of the films were studied. Parallel flat

plate capacitors were constructed on the films grown on aluminum substrates to determine the dielectric constant,

dielectric strength, resistivity, and the dissipation factor of the grown fihns

Corresponding Author

William C. Lanter

ISSI Corp.

2766 Indian Ripple Rd.

Beavercreek, OH 45440

e-mail: kv.i.l.lj..a...m_:J.a_.n.Le r._}!lL._..:...af..:..mi.l_

Fax: (937) 656-4095

Voice: (937) 255-5179 x34

Keywords: Carbon Nitride, Thin Films, Amorphous, Dielectric

Posting online: Your abstract received in MS Word format may be posted online on the conference web site. If you

do not want your email address to be shown along with this abstract, please indicate so here.

Your abstract should not exceed one page.

Electronic files will be accepted in Microsoft Word format only. Submit your Word file as an e-mail

attachment to KazuhisaMivoshi@grc.nasa.gov with a copy sent to ADC-FCT-2001@,eng.aubum.edu. All others

please fax or mail one hard copy to the address shown below. Abstracts must be received on or before the

announced deadline.

Kazuhisa Miyoshi

NASA Glenn Research Center

21000 Brookpark Rd.

Mail Stop 106-5

Cleveland, OH 44135-3191 USA

Fax: (216) 433-5544

Tel: (216) 433-6078

NASA/CP--2001-210948 597 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Structural properties of Carbon Nitride Films Deposited by Reactive Pulsed

Laser Deposition Technique

A.R.Phani and J.E.Krzanowski

Department of Mechanical Engineering, University of New Hampshire, NH

Abstract

Carbon nitride films have been deposited by reactive pulsed laser deposition technique

by ablating carbon in the nitrogen atmosphere at different substrate temperatures and

back ground pressures. Si(1 11) substrates are used in the present investigation. Deposited

films are uniform and shows good adhesion to the substrates The deposition rates depend

on laser fluence, back ground pressure, and target substrate distance. The nitrogen

concentration in the films increases with increasing back ground nitrogen gas pressure

and laser fluence. Fourier transform infra red spectroscopy has been employed to

distinguish the C-C and C-N bonds. X-ray Photoelectron spectroscopy has been used to

study the composition of the deposited films. X-ray Diffraction and Transmission

electron microscopy techniques revealed that the deposited films have shown oriented

microcrystalline structure at relatively high laser fluence. This could be due to the high

kinetic energy of the radicals in the laser produced plasma plume. Electronic, mechanical

and tribological properties &these films have also been discussed.

NASA/CP--2001-210948 598 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

MICRO-WEAR PROPERTIES OF

ION-PLATED CARBON NITRIDE THIN FILMS

Shuichi Watanabe*, Shojiro Miyake and Masao Murakawa

Nippon Institute of Technology

4-1 Gakuendai, Miyashiro-machi, Saitama 345-8501, Japan

*Corresponding Author: E-mail: shwata@nit.ac.jp,

Fax: +81-480-33-7745, Phone: +81-480-33-7729

ABSTRACT

Carbon nitride films have been synthesized by several kinds of vapor phase techniques including plasmaCVD, sputtering and ion-beam technique, etc., because of considerable interest in super-hard cubic or beta-type

C3N4 compound. We also prepared carbon nitride films by means of an ion-plating. Carbon nitride films were

deposited on Si substrate by reactive evaporation in carbon vapor and nitrogen plasma. The micro-wear

properties of the films were examined by means of the micro-scratch test together with a nano-indentation test,

utilizing an atomic force microscope. The results show that a suitable partial pressure ratio for Ar/N2 was shown

to be required for C-N bonding and hardness for films formed under this partial pressure ratio was shown to be

somewhat higher than hardness for other films. Hardness was also shown to increase as substrate bias decreased.

In the nanoindentation evaluation, the Ar/N 2 partial pressure was shown to influence film hardness. Finally,

micro-wear was shown to be small for films with high hardness.

Keywords: Carbon Nitride Film, Reactive Ion Plating, Micro-Wear

INTRODUCTION

Research into the synthesis of crystalline carbon nitride thin films has been stimulated by the suggestion of

the possibility of bulk moduli greater than that of diamond by Liu et al. (ref. 1), and a number of tests using

various vapor phase synthesis methods have been reported. Although there are some reports of success in

synthesizing very fine crystalline carbon nitride (ref. 2), most such films are amorphous. The development of a

method capable of generating consistent crystalline thin films is therefore eagerly awaited.

The authors are researching the use of the magnetically enhanced plasma ion plating method (ref. 3), with

the goal of synthesizing crystalline carbon nitride thin films. The mechanical properties of films generated under

various conditions were surveyed and are reported here as a first step in this research. Specifically, the

nanoindentation hardness and micro-wear properties were investigated.

FILM FORMATION

The magnetically enhanced plasma ion plating method, which has been used successfully to form cubic

boron nitride films (ref. 3), was used as the film formation method. Details regarding the equipment used in this

process can be found in existing literature (ref. 3). Carbon nitride films were deposited on Si substrate by

reactive evaporation in carbon vapor and nitrogen plasma. Argon/nitrogen mixed gas was flowed into the

chamber and the partial pressure ratio of Ar/N2 was varied in 1/0, 3/1, 2/1, 1/1, 1/2 and 0/1, when the total gas

pressure during deposition was set up at 6.6 x 10.2 Pa. An Si wafer (100) was used as the substrate and

99.999% pure Gr was used as the evaporation source. RF bias (Vdc) varied in the range of 10 - -60. The

electron beam output, discharge potential, and discharge current were fixed at 250 mA, 60 V, and 16 A,

respectively. A film formation time of 1800 seconds was used.

STRUCTURE AND COMPOSITION OF THE FILMS

First, in order to investigate the C-N bonding states, IR absorption spectra were measured by FT-IR.

Figure 1 shows the IR spectra for films formed (using a Vdc of -50 V) under various partial pressures. The C-N

stretching vibration mode with the most remarkable difference (wave number: around 1120 cm -1) is highlighted

in the figure. From the figure, absorption due to C-N stretching vibration becomes prominent as N2 is added to

the reactive gas. In addition, as the N2 partial pressure is increased, the absorption peak shifts to high wave

numbers, and as the Ar partial pressure decreases, this absorption becomes less apparent. From these results, the

presence of Ar is thought to influence C-N bonding states and bond formation (probably due to the Ar ion

bombarding effect).

Next, a composition analysis of the film was performed by Auger Electron Spectroscopy (AES). Although

NASA/CP--2001-210948 599 August 6, 2001
noresultsareprovidedhere,theN/CatomicratioforafilmformedunderanAt/N2partialpressureratioof3/1

wasconfirmedtobeapproximately 0.6.However, smallquantitiesofOandWweredetectedinthisfilm.TheW

isthoughttohaveoriginatedfromthehotelectronemissionfilament(madeofW)usedtogeneratedtheplasma.

TheOisthoughttohaveoriginatedfromsmallquantitiesofwatercontainedinthesurroundinggas.Reducing

theseimpuritieswillbeathemeoffutureresearch. TheN/CvaluesweresmallerforfilmsformedunderAt/N2

ratiosotherthan3/1.

(Ar/N_F1owRatio)

1200 1100 1000

W_venumber(cni 1)

Figure1. IR spectra for films formed under various gas flow ratios. (Vdc: -50 V)

FILM HARDNESS MEASUREMENT

The surface hardness (Vicker's hardness: 5 mN load) was measured for films formed under various

conditions (Vdc fixed at -50 V). The results are shown in Fig. 2. Film hardness was somewhat higher for films

formed under flow ratios of 3/1 and 2/1, which clearly showed C-N bonding in Fig. 1. However, the differences

were not remarkable. For most samples the average Hv value was 2300 ~ 2800. The reason for the extremely

low hardness results measured for the film formed under a flow ratio of 1/3 is still not known.

Next, the hardness of films formed at various RF substrate biases was measured. A flow ratio of 3/1,

which afforded the highest hardness in the above investigation, was used. The results are shown in Fig. 3. As

shown in the figure, film hardness increases as bias decrease; an Hv of 3400 was obtained at a Vdc of-10 V.

NANOINDENTATION AND MICRO-WEAR MEASUREMENT OF THE FILMS

Since the hardness evaluation described above used Vicker's hardness with a 5 mN load, substrate

influences are included in the measured values. In order to evaluate the film itself, the nanoindentation hardness,

which uses small loads, was evaluated. Micro-wear was evaluated at the same time in order to estimate the wear

resistance of the film. Micro-wear was evaluated by a micro-scratch test (ref. 4) using an atomic force

NASA/CP--2001-210948 600 August 6, 2001
=Y.

T-j

35OO

3000

2500

2000

1500

1000

5OO

0

1/0 5/1 3/1 2/1 1/1 1/3 0/1

Flow ratio (Ar/N2)

Figure 2. Surface hardness values of films formed under various gas flow ratios. (obtained by Vicker's

hardness test, load: 5 mN, Vdc fixed at -50 V)

o_.-

_rb

.3>

4000

3500

3000

2500

2000

1500

1000

500

0

_-. 41,

I I I

0 -20 -40 -60 -80

Vdc(V)

Figure 3. Surface hardness values of films formed at various RF biases (Vdc). (obtained by Vicker's

hardness test, load: 5 mN, gas flow ratio fixed at 3/1)

microscope (AFM). The same instrumentation was capable of simultaneously executing nanoindentation

measurement.

An equilateral triangle stylus tip made of diamond was used as the indenter. Loads of 10, 20, and 30 gN

were applied, and the surface of a 200 x 200 nm area scanned for scratches. The wear depth was measured and at

this time, allowing for evaluation of wear resistance properties of the film itself without including influences

from the substrate. A schematic representation o f the micro-scratch evaluation is shown in Fig. 4.

FiRure 4. Schematic representation of micro-scratch evaluation.

NASA/CP--2001-210948 601 August 6, 2001
The nanoindentation measurement results for films formed in reactive gas containing Ar/N2 mixed gas

flow ratios of 1/0, 3/1, and 0/1 are shown in Fig. 5. These results were obtained at a load of 100 _LN. A

maximum value of approximately 32 GPa was obtained when the flow ratio was 3/1. The hardness is clearly

high compared to films formed without nitrogen gas flow and films formed using only nitrogen gas flow. The

N/C with a flow ratio of 3/1, at approximately 0.6 according to AES analysis, was also higher than values

obtained under other conditions. The above results clearly demonstrate that film hardness is increased through

the inclusion of nitrogen.

0 )..........................

I i

35 ............................................................ 1

z........ ............... I

..... 25 t ;:.::::W.::x_N 1

t....................... ;::::':':+:::' ',i_:i:i

15 ' ....................... " "" "N:, :': ...................... 1

0 t...................... ':":*:':" N i!iL.'iiiiiNi.;i,:....... ":" ''

',,'f . .%°.%:,_

170 371 0/1

Flow ratio (Ar/Nz)

Figure 5. Na_oindentation hardn_s values of film_ formed under _m'imzs gas flow ratios.

(load: 100 /_N. Vdc fixed at -50 V)

The micro-wear measurement results for the films are shown in Fig. 6. The wear due to scratching under

various loads was observed by depth using the same indenter tip and a small load of 1 micro-N. The wear depths

were derived from the results. Although increased wear was observed for all films as loads increased, wear on

the film formed under a flow ratio of 3/1 was clearly smaller than wear on the films formed under other

conditions. This result agrees with the nanoindentation measurement result and shows that differences in

hardness influence wear rates.

_L

9

8

7

6

5

4

3

2

1

0

0

/_ q/u

o/1

3/1

I I I

10 20 30 40

Load (11 N)

Figure 6. Variations of micro-wear depths of films formed under various gas flow ratios on loads.

CONCLUSION

Carbon nitride films were synthesized using the magnetically enhanced plasma ion plating method. A

suitable partial pressure ratio for Ar/N_ was shown to be required for C-N bonding and hardness for films formed

under this partial pressure ratio was shown to be somewhat higher than hardness for other films. Hardness was

NASA/CP--2001-210948 602 August 6, 2001
alsoshowntoincreaseassubstratebiasdecreased. Inthenanoindentation evaluation, theAr/N2partialpressure

wasshowntoinfluencefilmhardness. Finally,micro-wear wasshowntobesmallforfilmswithhighhardness.

REFERENCES

1.A.Y.LiuandM.L.Cohen:Science, 245(1989)841.

2.S.L.Sung,T.G.Tasi,K.P.Huang, J.H.Huang,andH.C.Shih:Jpn.J.Appl.Phys.,37(1998)L148.

3.M.MurakawaandS.Watanabe: Surf.Coat.Technol., 43/44(1990)128.

4.S.Miyake,R.KanekoandT.Miyamoto: DiamondFilmsandTechnol., 14(1992)205.

NASA/CP--2001-210948 603 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

DEMONSTRATION OF GROWTH OF ULTRA-HARD CARBON NITRIDES FILMS BY

HOT-FILAMENT CHEMICAL VAPOR DEPOSITION

(ADC/FCT 2001)

Saibal Mitra, Justin Hartman

Department of Physics and Engineering Physics

University of Tulsa, Tulsa, OK 74104

and

Tobias Hanrath

Departrrent of Chemical Engineering

University of Tulsa, Tulsa, OK 74104

ABSTRACT

Development of novel ultra-hard materials is of great interest to the material science community. In particular,

carbon nitride films have attracted a lot of interest due to its extreme properties. The hardness of pure 13-C3N4 has

been predicted to be at least equal, and perhaps greater, than that of diamond. This paper describes the growth of the

ultra-hard phase of carbon nitride using hot-filament chemical vapor deposition (HFCVD). Carbon nitride films were

grown on (100) oriented crystalline Si substrate by activating a mixture of CH4/NH3 over a hot tungsten filament. The

films were deposited over a wide range of conditions. The ratio ofNH3/CH4 was held between 5 and 10 while the

reactor pressure was varied between 80 mT to 550 torr and the filament-to-substrate distance between 0.4 to 0.5 cms.

Once deposited, these films were characterized primarily by AFM and XRD. AFM micrographs show the presence of

facetted crystallites. The observed XRD peaks were compared with theoretical predictions for the c_-C3N4 and 13-C3N4 -

the two ultra-hard phases of carbon nitride. The XRD results clearly suggest the presence of both c_-C3N4 and 13-C3N4.

Our results also indicate that the surface morphology of the deposited films is sensitive to growth conditions,

particularly reactor pressure and filament-to-substrate distance. This paper will also compare the effects of using

nitrogen or ammonia in the gas phase on the material quality. Clearly, additional work, that would systematically span

a wide range of deposition conditions, is needed to optimize conditions that would lead to the growth of primarily the

ultra-hard phases of carbon nitride.

Keywords: carbon nitride, CVD, ultra-hard films, AFM, XRD.

INTRODUCTION

Following the theoretical predictions by Liu and Cohen (1) that a novel material [_-C3N4, whose structure is similar

to 13-Si3N4, should have hardness comparable or even exceeding that of diamond, carbon-based nitrides have

attracted widespread attention. In addition to 13-C3N4, several other possible structures of carbon nitride, such as c_,

cubic and graphite, were also predicted (2,3). Consequently, there has been intense experimental interest and

numerous groups have attempted to synthesize the ultra-hard phase of carbon nitrides. The results of these studies

have been mixed. In many cases amorphous carbon nitride (a-CN) films were synthesized (4 6). Other groups have

reported the presence of small C3N 4 crystallites embedded in _CN films (7,8). However, recently Chen et al. have

reported the growth of well-facetted o_-Si×CyNz crystals and a mLxture of c_- and [_-C3N 4 crystals (9,10). Bursill et al. (11)

have demonstrated the existence of an fcc phase of carbon nitride with a unit cell parameter of about a 6.3 A, and

Nguyen and Jeanloz (12) have pointed out the presence of a cubic carbon nitride phase when the material is grown

under high pressure and temperature.

Though the growth of pure and perfect 13-C3N4 still remains a challenge, one promising method is growth by hotfilament chemical vapor deposition (HFCVD). There are reports in the literature that describe the successful growth of

13-C3N4 and other ultra-hard phases of carbon nitride using HFCVD and some of its variation (9,13-15). In this work,

our objective has been to demonstrate the growth of carbon nitride films by activating a mixture of methane (CH4) and

ammonia (NH3) in a HFCVD reactor and characterize them using atomic force microscopy (AFM), scanning electron

microscopy (SEM) and x-ray diffraction (XRD). Using NH3 instead of nitrogen gas (N2) as a source of nitrogen has a

NASA/CP--2001-210948 604 August 6, 2001
numberof advantages. Table1liststhebondenergiesofN2,hydrogen(H2)andNH3andotherrelevant radicals

typicallyfoundinthegasphaseinaHFCVDreactorwhenamixtureofNH3andCH4isactivatedoverahotfilament.

Table 1. A comparison of bond strengths of some common gases used in a HFCVD process (ref.16).

Bond type

_H

N-N 226.8

H-NH 7

Bond Strength

(kcals/mole)

104.2

110

H-N 103

C-N 67

C-H 99

The energy required to break a hydrogen atom off an ammonia molecule is about 110 kcals/mole which is less

than half the energy required to break the nitrogen molecule. Thus, using NH3 instead of N2 is preferable because it

is relatively easy to strip a hydrogen atom from the ammonia molecule. This hydrogen atom has two major roles in

promoting the growth of carbon nitride films. It helps the formation of sp 3 carbon bonds by etching away the

graphitic sp 2 phase. It is well known that the sp 3 C-N bond is more stable than either the sp 1 C---N or the sp 2 CN

bonds (16). Also, nitrogen is incorporated into the growing surface more effectively when the H from the -CH bond

on the surface is abstracted by the -NH radical forming a -CN bond and an H2 molecule.

EXPERIMENTAL DETAIIS

The carbon nitride films were grown in a home-built hot-filament chemical vapor deposition (HFCVD) reactor

where a mixture of NH3/CH4 was activated by passing the gas over a single hot tungsten filament whose temperature

was estimated to be ~2000°C. Each of the gases were introduced through separate outlets into the reactor. While CH 4

was introduced just above the filament, NH3 was introduced from top of the glass jar itself. The filament-substrate

distance, df_, was nominally set to 0.5 cm for samples A and B and 0.4 cm for sample C and was measured both before

and after deposition to ensure that the conditions did not change during the course of a run. The growing film was

heated directly by the hot filament. All samples were deposited on (100) oriented silicon substrates.

The chamber was pump ed down to a base pressure of ~25 mTorr using a mechanical pump. Initially, CH4 was

introduced and its flow rate controlled using a mass flow controller and its partial pressure was set to the desired

value. Once the partial pressure of CH4 had stabilized, NH3 was added to the reactor. The [N]/[C] ratio was controlled

by adjusting the partial pressure of NH3. The final pressure of the reactants was set by adjusting the outlet valve

located before the mechanical pump. The gas mixture was then activated using a hot-filament.

Table 2 describes the growth parameters for the three samples used in this work. While the ratio of nitrogen-tocarbon was held constant, the flow rate of CH 4 was varied between 5 - 10 sccm. Finally, the pressure in the reactor

was varied from 220 torr to 550 torr to change the microstructure and the "quality" of these films. Once deposited, the

films were characterized by AFM for surface morphology and XRD for crystalline microstructure.

Table 2. Deposition conditions for the growth of carbon nitride in a hot-filament CVD

reactor. The methane flow rate was 10 sccm for all samples.

Sample # Filament-substrate Pressure (torr) Deposition rate (mg/min.)

distance, df_ (cm)

A 0.5 220 6.2 X 10 -3

B 0.5 300 6.2 X 10 -3

C 0.4 550 1.3 X 10 -3

D 0.7 350 not available

NASA/CP--2001-210948 605 August 6, 2001
RESULTS

Ashasbeenmentionedarlier, themicrostructure ofHFCVDfilmsissensitivetothevariousgrowthparameters.

Carehasbeentakentoensurethatheconditions, duringdeposition, remainedstable.SamplesAandB,whichwere

depositedatconditionsveryclosetoeachother,hadsimilardepositionrates.SampleC,whichwasdepositedata

higherpressure, hadasmallerrateofdeposition. DepositionrateforsampleDisnotavailable. Clearly, moreworkis

neededtoestablishanycorrelationbetweengrowthratesandreactorpressures.

ThesurfacemorphologyofthedepositedfilmswasstudiedbyAFMmadebyDigitalInstruments. TheAFM

imageswereobtainedwithahorizontal headoperatinginairatroomtemperaturesingetchedtungstentips.The

scansizewas5.0_tmwhilethescanratewas0.3Hz.Figure1showstheAFMmicrographoftwotypicalsamples.

SampleBisawellfacettedfilmwithan~l_tmgrainsizewhilethecrystallitesizeinsampleDisconsiderably smaller.

(a) (b)

Figure1. AFM of two carbon nitride samples deposited at (a) Sample B, 300 torr, dt_= 0.5 cm, and (b) Sample D,

350 torr, d¢_= 0.7 cm.

The XRD spectra were measured by using the Cu Kc_l radiation using a standard 0 - 20 scanning. Table 3 lists all

the observed interplanar d-spacings for the three samples described in this work. The measured values of the

interplanar d-spacings are compared to the theoretical values as calculated by Guo and Goddard (17) and Teter and

Hemley (3) for c_- and ]3-C3N4, respectively. The observed interplanar d-spacings that have a close match with either

c_-C3N4 or ]3-C3N4 - the ultrahard phases of carbon nitride - are indicated in bold.

All three samples exhibited the strongest peak at d 1.35A. This reflection has been attributed to the (400) plane

from the Si substrate. Unfortunately, another XRD reflection with d 1.33A is also expected from from the (221) plane

of ]3-C3N4 (3). Thus, even if this peak was present, it would be completely overwhelmed by the Si (400) relfection.

Interestingly in sample A, a closer inspection reveals six reflections with interplanar d-spacings of 1.359 A, 1.358 A,

1.356 A, 1.355 A, 1.354 A and 1.356 A. In this particular case, only one of the six reflections may be assigned to Si (440).

At least one other peak may be assigned to reflection from the (221) plane of _-C3N 4 while the rest of the other peaks,

at this time, are of unknown origin. It is possible, the presence of these additional peaks is due to the presence of

other phases of carbon nitride that are yet to be fully accounted for. Two other peaks that have been associated with

the Si substrate have d-spacings of 1.64 A and 1.25 A. All three samples exhibit a reflection with d 1.68 A which may

be attributed to graphite (004) plane. However, it should be noted that reflection from graphite (002) plane is absent in

all samples. The intensity from the (004) plane in graphite is 4% of the (002) plane (ASTM card 41-1487). Thus, the

intensity of reflection from the (002) plane is expected to be at least an order of magnitude larger than the (004) plane

and hence it is unlikely that this peak was from graphite. Consequently, this peak has been reassigned to the

reflection from the (301) plane of c_-C3N4. The peaks that have been assigned to the ultrahard phases of carbon nitride

((z-C3N4 or _-C3N4) in sample A have d-spacings of 3.19 A, 2.08A, 1.90A, 1.68A, 1.59A, 1.52A, 1.50A, 1.46A, 1.35A (one

peak), 1.26A. Similarly, for samples B and C, there axe a number of peaks with d-spacings that are in good agreement

with the predicted theoretical values for either c_-C3N4 or ]3-C3N4 and all of these peaks have been listed in Table 3.

These XRD results are similar to those obtained by Zhang and co-workers (15). They deposited crystalline carbon

NASA/CP--2001-210948 606 August 6, 2001
nitridebyactivatingamixtureofmethane, ammoniaandhydrogenbyrf-plasmaassistedHFCVD. Theseworkershad

reportedthepresenceof13-C3N4 (thoughnotc_-C3N4) intheirfilms.

Table 3. Comparison of the experimental measurements of interplanar d-spacings with theoretical values for a- and

I$-C3N4 (ref. 15 and 3, respectively).

_-GN4

d-spacings (hkl) d-spacings

5.50 (100) 5.50

3.17 (110)

2.75 (200)

2.25 (101)

2.08 (210)

1.95 (111)

1.83 (300)

1.81 (201)

3.55

3.17

2.75

2.37

2.14

2.08

c_-C3N4

(hid)

1.60 (200)

1.59 (211)

1.52 (310)

1.47 (301)

Interplanar d-spacings in (A)

1.33 (221)

Sample A Sample B Sample C

(100) 4.97

3.82

(101) 3.49 3.64

(110) 3.19 2.98

(200) 3.03 2.72 2.71

(201) 2.53

(102)

(210) 2.08

1.90 (211)

1.87 (112)

1.77 (202)

1.70 (301)

1.49 (103)

1.45 (311 )

1.90 1.91 1.90

1.68 1.69 1.68

1.66 1.66 1.65

1.63 1.64 1.63

1.59 1.61

1.52 1.51 1.52

1.50 1.51 1.50

1.46 1.42 1.42

1.41 1.41

1.40 1.40

1.351 1.36 1.35

1.31 (222)

1.26 (320) 1.27 1.26 1.26

1.26 1.25 1.25

1Multiple peaks are observed. The strongest peak is attributed to Si (400) reflection. See text for details

CONCLUSIONS

Carbon nitride films have been grown by HFCVD on (100) oriented Si substrates. AFM measurements show that the

films are well facetted while XRD measurements indicate the presence of c_-C3N4 and 13-C3N4 the ultrahard phases of

carbon nitride. However, there are a number of XRD peaks that are unaccounted for indicating the presence of other

phases. We have suggested that using NH3 as a source gas for nitrogen has distinct advantages since it is easier to

break the ammonia molecule and nitrogen is easily incorporated in the growing surface. However, the -CN bond

energy is a relatively low 67 kcals/mole (see Table 1) and is easily broken. Consequently, the ratio of NH3/CH4 should

NASA/CP--2001-210948 607 August 6, 2001
berelativelyhightoensurehighconcentration ofitrogeninthegasphase.Inthiswork,wehavenotstudiedthe

effectoftheadditionofhydrogenoncarbonnitridefilms.Theroleofhydrogenisimportant andwillbeaddressed in

futurework.It isevidentthatHFCVDis anattractiveapproachthatcanleadtothefabricationofsingle-phase

ultrahardcarbonnitridefilms.Clearly, additional workneedstobedonetoachievethisgoal.

_CES

1. A.Y.LiuandM.L.Cohen: PredictionfNewLowCompressibility Solids. Science, 25,(1989), 841

2. A.Y.LiuandR.M.Wentzcovitch: Stabilityofcarbonnitridesolids. Phys.Rev.B,50(1994) 10362.

3. D.M.TeterandR.J.Hemley, Science, 27153(1996).

4. C.S.Torngetal.:StructureandbondingstudiesofC:Nfilmsbyrf sputteringmethod.J.Mater.Res.,5 (1990)

2490.

5. D.Li,Y.W.Chang,, andW.D.Sproul: Nano-indentation studiesofultrahighstrengthcarbonnitridefilms:J.

Appl.Phys., 74(1993) 219.

6. D.Martonet.al:Carbonnitridedepositedusingenergeticspecies: Atwo-phasesystemPhys.Rev.Lett.,73118

(1994).

7. T.Y.YaaagndC.P.Chou:Growthandcharacterization ofcarbonnitridethinfilmspreparedbyarc-plasma jet

chemical vapordeposition. Appl.Phys.Lett.,67(1995) 2801.

8. C.Niu,Y.Z.Liu,andC.M.Leiber:Experimental RealizationoftheCovalent SolidCarbonNitride.Science, 261

(1993)34.

9. L.C. Chenetal.:Si-containing crystallinecarbonnitridederivedfrommicrowaveplasma-enhanced chmical

vapordepositionThinSolidFilms,30366(1997).

10.Y.Chenetal.:Synthesis ofcrystallineC3N4filmsandthenewC-Nphases. Mat.Res.Soc.Proc., 441(1997)717.

11.L.A.Bursillet.al:Carbonnitridefilmsproducedbyhighenergyshockplasmadeposition: J. Mater.Res.,10

(1995) 2277.

12.J.H.NguyenandR.Jeanloz: Initialdescriptionfanewcarbonnitridephasesynthesized athighpressures and

temperatures. Mater.Sci.Engg., A20923(1996).

13.A.Leonhardt etal.:Preparation ofCNx-phases usingplasmaassistedandhotfilamentchemical vapordepostion.

ThinSolidFilms,332,(1998), 69.

14.G.Wang, Y.Chen,andL.Guo:Synthesis andCharacterization ofPureCrystallineC-Nfilm.PhysicaScriptica, T69,

(1997), 108.

15.Y.Zhang, H.Li,andQ.Xue:Synthesis ofCrystallineCarbonNitridebyChemical VaporDeposition. Mat.Res.Soc.

Symp. Proc., 410,(1996), 317.

16.D.R.Lide,CRCHandbook ofChemistry andPhysics(CRC,BocaRaton)75aEd.,(1994).

17.Y.GuoandW.A.Goddard: IsCarbonNitrideHarderThanDiamond?No,ButItsGirthIncreases WhenStretched

(NegativePoissonRatio).Chem. Phys.Lett.237(1995), 72.

NASA/CP--2001-210948 608 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Process Parameters for Synthesis of Carbon Nitride Thin Films

using Electron Cyclotron Resonance Plasma Sputter Method.

Eiji Kamijo, Yoshifumi Aoi, Tatsuya Hirahara, Takashi Tochimoto

Faculty of Science and Technology, Ryukoku University,

Otsu, Shiga 520-2194, Japan.

TEL:+81-77-543-7465, FAX:+81-77-543-7483

E-MaJl:kami] ou @nns._/ukoku.ac.j p

ABSTRACT

Carbon nitride thin films were synthesized by electron cyclotron resonance (ECR) plasma

sputtering method with a carbon target and under various process parameters. Especially, plasma

atmosphere, microwave input power and substrate bias voltage as the process parameters were

investigated. The nitrogen concentration (N/C atomic ratio) and chemical bond structure (N-spC,

2 3 .......

N-sp C, N-sp C) of nitrogen and carbon atoms in synthesized thin films were discussed with XPS,

FT-IR, Raman spectroscopy. The nitrogen concentration, chemical bond structure and other

properties were clearly dependent on the process parameters. Especially, in the nitrogen

concentration and the chemical bond structure of carbon nitride thin films, there are remarkable

differences between the effect of positive and negative bias voltage applied to the substrate. These

results were implied that the most important process parameters in synthesis of carbon nitride thin

films are to irradiate the deposited film surface with higher flux density of low energy nitrogen ion

below 40eV. And also, these results are supported the sub-implantation model for synthesis of

metastable phase such as diamond and cubic boron nitride.

Keyword: carbon nitride film, electron cyclotron resonance plasma sputtering, chemical bond

structure, low energy nitrogen ion, sub-implantation model.

INTRODUCTION

Since Liu and Cohen in 1989 theoretically predicted the crystalline 13-C3N4 compound with a hardness

equal to that of diamond, many researchers have tried to synthesize this material[ 1,2]. There are a lot of

reports on synthesis of carbon nitride films, nevertheless no clear evidence of 13-C3N4 thin films has been

reported so faac[3-7]. The details of chemical and structural characteristics of the carbon nitride compounds

are still unknown.

A sputtering type electron cyclotron resonance plasma deposition method, hereafter called ECR

plasma sputtering method, has the advantages of depositing the films under irradiation with a low energy

ion, high density plasma stream onto the substrate in order to the low plasma gas pressure (under 10 -1 Pa).

The energy level of the activated species in ECR plasma is a few tens eV, which is much lower than that

produced by other ion beam sources. It is considered that irradiation of the substrate with high density

plasma is favorable for the formation of a metastable phase. Furthermore, the use of energetic particle with

low energy is expected to reduce structural damage. Consequently, ECR plasma sputtering method is useful

to synthesize metastable phase such as hydrogen-free diamond-like carbon (DLC), microcrystalline

diamond and CNx films[8,9].

We had reported on the synthesis of crystalline carbon nitride films deposited with the ECR plasma

sputtering method[9]. However, the nitrogen content in the crystalline carbon nitride films deposited at

600°C was very low. So, we tried to define the process parameters of ECR plasma sputtering system for

increasing nitrogen content in the films.

In order to increase the nitrogen content in the films, the influence for nitrogen content and chemical

bond structure of the synthesis process parameters, especially, plasma atmosphere, substrate bias voltage

and microwave input power in the ECR plasma sputtering system were studied. For investigate the

relationship between the nitrogen content, the chemical bond structure of carbon to nitrogen and atomic,

molecular, and ionic species in the ECR nitrogen plasma, optical emission spectroscopy (OES) and ion

density were measured.

NASA/CP--2001-210948 609 August 6, 2001
EXPERMENTAL PROCEDURE

Carbon nitride thin films were synthesized

using the ECR plasma sputtering system on Si(100)

substrate under various conditions, nitrogen partial

pressure in plasma gas atmosphere (Nz/Nz+Ar), bias

voltage of the substrate, and microwave input

power.

The schematic diagram of the ECR plasma

sputtering apparatus is shown in Fig.1. Before

deposition, the vacuum chamber was evacuated to

1.0xl0 -4 Pa. A 2.45 GHz microwave power,

magnet coil power for magnetic field of 875 Gauss

was applied to the ECR system. The mixed gas

composed with Ar and N2 was introduced into the

vacuum chamber to generate ECR plasma. The total

pressure was kept at 1.3x10 -1 Pa. The nitrogen

partial pressure was varied from 0.1 up to 1.0. A

cylindrical pure carbon taxget was placed at the

plasma-extracted window as surrounding the plasma

stream. A negative target voltage of 370V for the

sputtering of carbon was applied to the target holder

coated with A1203 via DC power supply. A substrate

Microwave

( ; 45GHz,

Magnet c__

Target DC Power

I _ I Optical emission

II I spect ..... py

Substrate/ t.........__ I i[_..._l

.insP...... pplr-"-tJ\ 1

Langmu'r probe

Fig. 1 Schematic diagram of the ECR plasma

sputtering apparatus.

bias voltage was applied to the substrate holder with another DC power supply, and was varied between

+40V and -40V. The substrate temperature was ambient temperature in this all experiments. The substrate

was exposure to the ECR plasma extracted by the inclined in magnetic field.

The deposited films were evaluated with x-ray photoelectron spectroscopy (XPS), FT-IR, Raman

spectroscopy.

EXPERIMENTAL RESULTS AND DISCUSSION

Nitrogen contents.

Nitrogen atom to carbon atom ratio (N/C) in the films deposited at several nitrogen partial pressures

was evaluated with measured XPS spectra as shown Fig. 2. The total plasma working gas pressure and

substrate bias voltage was 1.3x10 -1 Pa, -30 V, respectively. As the nitrogen partial pressure (Nz/Nz+Ar)

increased, the N/C atomic ratio increased lineally from 0.2 up to 0.35. The films deposited at ambient

temperature with the maximum N/C atomic ratio were obtained at 1.0 of the nitrogen partial pressure,

namely pure nitrogen atmosphere. From the result, it is implied that Ar gas has a bad effect for increasing

of N/C atomic ratio in the films.

0.5

0.4

"_ 013

L)

0.2

0.1

0 0

0

0 o

o "o._, 'o.:_ "o._,'o._ "1'

N 2 / (N2+Ar)

O

Fig. 2 N/C ratio in the film deposited at

several nitrogen partial pressures.

0.5

0.4

©

_ 0.3

20.2

0.1

0

O0

0 o

0

40 20 0 -20 -40

Substrate bias / V

Fig. 3 N/C atomic ratio in the films deposited

at various substrate bias voltages.

NASA/CP--2001-210948 610 August 6, 2001
N/Catomicratiointhefilmsdeposited atvarioussubstrate biasvoltages andatpurenitrogenpressure

of1.3x10-1Pawasevaluated withXPSasshownFig.3.TheN/Catomicratioincreasedwithdecreasing

thesubstrate biasvoltage, andtheN/Catomicratioincreased lineally. However, theN/Cratiomaybe

saturated atlessthan-30Vsubstratebiasvoltagerange. Thethinfilmsdidnotdeposit asubstrate bias

voltagelssthan-50Vinordertore-sputtering ofthedeposited films.Andalso,N/Catomicratiointhe

filmsdepositedatvariousmicrowaveinputpowerandatpurenitrogenpressureof 1.3x10-1Pawas

evaluated. TheN/Catomicratiogradually increased withincreasing themicrowave inputpower.

Thesersultssuggest thathenitrogenpartialpressureandsubstrate biasvoltageisimportant process

parameters forsynthesis ofcarbonnitridefilmswithhigherN/Catomicratio.Theeffectsofmicrowave

inputpowerissmaller thantheotherprocess parameters foynthesis offilmswithhighN/Catomicratio.

Chemical bond structure.

Chemical bond structure of carbon and nitrogen atom was evaluated with XPS spectra. The XPS

spectra of N ls electrons of films was seen the shifts of the main XPS peak from the binding energy values

corresponding to simple elements 402eV for N ls peak. The energy shift toward to lower values 399eV is

typical for nitrogen incorporation in carbon nitride films[10]. Furthermore, N ls and C ls spectra axe an

asymmetric broadening toward higher binding energy. These suggested that the spectra comprise several

features of varying binding energy levels corresponding to different bonding structures of N and C atoms.

Many attempts have been made to deconvolute N 1s spectra in previous reports[ 11-15].

N sp3C

• 398eV

; Ns o/

N spC

NN, NO

405 400 395

Binding energy / eVo @

Fig. 4 Typical deconvolution result of the Nls

spectra of the films deposited at substrate

bias voltage -30V and microwave power

200W.

ioo

_8o

o

_60

¢9

r_

2O

O'

• ,:_-sp3C

• £N-sp2C |

e_?pc _ !

0.2 0.4 0.6 0.8

N2!(N2+Ar)

Fig. 5 XPS peak area ratio on each chemical

bond structure of the films deposited at

various nitrogen partial pressures.

o

4 4

_2

L)

z

O

O

0 0

0

0

o 05 o.:_ o._, ok 1'

N 2 / (N2+Ar)

Fig. 6 Peak area ratio of (N- sp3C)/(N-sp2C)

vs nitrogen partial pressures.

o4

L)

_3

1

O0

o Oo

40 20 0 -20 -40

Substrate bias / V

Fig. 7 Peak area ratio of (N- sp3C)/(N-sp2C)

of the films deposited at various substrate

bias voltage.

NASA/CP--2001-210948 611 August 6, 2001
Astandardnumerical deconvolution procedurewithGaussian fittedcurveshasbeenappliedtoN ls

spectrawithfourcomponents. Thebindingenergyofnitrogencombinedwithsp 3 hybridized carbon (N3 2 . . 2 . .

sp C), sp hybridized carbon (N-sp C), sp hybridized carbon (N-spC), and N-N and N-O bond could be

assigned to 398.0eV, 400.0eV, 399.5eV, and 402.0eV, respectively. Typical deconvolution result of the N ls

spectrum was shown in Fig. 4.

The XPS peak area ratio on each chemical bond structure of the films deposited at various nitrogen

partial pressures was shown in Fig. 5. The peat: area ratio of (N-sp3C)/(N-spZC) calculated from Fig. 4 was

shown in Fig. 6. The peak area ratio of (N-sp3C)/(N-spZC) of the films deposited at various substrate bias

voltage and microwave input power was shown in Fig. 7 and Fig. 8.

o4

:')3

1

O

O

O

O

0 50 100 150 200

Microwave power / W

Fig. 8 Peak area ratio of (N- sp3C)

/(N-sp2C) of the films deposited at

various microwave input power.

"3

300

i i i i

2nd Positive System )

1st Positive System

(

I

400 500 600 700 800

Wavelength / Inn

Fig. 9 Typical emission spectrum of N2 ECR

plasma observed with OES.

These results indicate that the nitrogen atoms combined with sp 3 hybridized carbon (N-sp3C) increased

with increasing the N/C atomic ratio, and the nitrogen atoms combined with sp2 hybridized carbon (Nsp2C) decreased with increasing the N/C atomic ratio. The change in the peak area assigned to N-N or N-O

bonds did not observed. The (N-sp3C) /(N-sp2C) peat: area ratio rapidly increased with increasing the

nitrogen partial pressure of 1.0. In the concerning of microwave power, the (N-sp3C)/(N-sp2C) peat: area

ratio rapidly increased at more than 150W compared with less than 100W. Furthermore, the (N-sp3C)

/(N-sp2C) peak area ratio sharply increased at minas range in substrate bias voltage.

Plasma diagnostics were performed with optical emission spectroscopy (OES) and Langmuir probe

method to define the mechanism of above-mentioned results. Typical emission spectrum of N2 ECR plasma

observed with OES was shown in Fig. 9. The emission bands in the UV region, centering at 316, 337, 357,

and 380urn, axe attributed to the second positive N2 band[16-19]. The bands in the visible region, centering

at 540, 580, 650, and 750nm, are attributed to the first positive N2 band[16-19]. The emission band at 391

N2+(3 91 ran)

o:m:)

s; -:d0 :s'0 -2d0 -2s0

Microwave Power / W

Fig. 10 Emission intensity in N2 ECR plasma

as a function of microwave input power.

[]1o8

ioo

50

10

Ys

s; :do :s'0 2o'0250

Microwave power / W

Fig. 11 Relationships between ion intensity

and microwave input power.

NASA/CP--2001-210948 612 August 6, 2001
and425nmcorrespond tothefirstnegativesystemofN2÷ion[16-19]. Theemissionintensityat391nm

(N2÷ion)and580nm(N2radical) inN2ECRplasmaasafunctionofmicrowave inputpowerwasshownin

Fig.10. Therelationships betweenionintensitymeasured withLangmuir probeandmicrowave inputpower

.... 3 2 . .

was shown m Fig. 11. The relation between the peak area ratio of (N- sp C)/(N-sp C) and microwave input

_ower was shown in Fig. 12.

o4

_3

O

o

O

50 100 150 200

Microwave power / W

Fig. 12 Relationship between the peak area

• 3 2 •

ratio of (N- sp C) /(N-sp C) and microwave

input power.

D bai;d_ i' '_...----G band 'i'

+30 V i

ground

-30 V

ib06 ' is6"0 " 2666 "

Raman shift / cm 1

Fig. 13 Typical Raman spectra of the films

deposited at various substrate bias voltages.

These results indicated that the intensity of N2+ ions and ion density rapidly increased at more than

150W in microwave input power. And also, an increase of the N2+ ion density contributed to an increase in

N/C atomic ratio and N-spJC chemical bond structure of the films. It was suggested that the N/C ratio and

the peak area ratio of (N-sp3C) /(N-sp2C) concerned with chemical bond structure in deposited carbon

nitride films strongly depend upon an ion density in the ECR plasma and the substrate bias voltage of less

than -50V.

Typical Raanan spectra of the films deposited at various substrate bias voltages were shown in Fig. 13.

The peak area ratio of D band (disorder: 1350 cm -1) to G band (grafite: 1550 cm -1) increased with decreasing

the substrate bias voltage. It was suggested from these experimental results that the preparation and

characterization of metastable carbon nitride films with higher nitrogen content are depend upon the

substrate bias voltage less than -50V, correspond to ion energy, and ion flux density irradiated onto the

growing film surface.

A growth model for formation of metastable phase in comparison between the these experimental

results and subimplantation model reported by Lifshitz et al. and Robertson et al. were discussed[20,21].

Our experimental results were implied that the subimplantation model could be applied qualitatively for the

formation of metastable phase carbon nitride films.

CONCLUSIONS

Amorphous carbon nitride films were deposited on the Si(100) substrate by the novel ECR sputtering

method. The N/C atomic ratio reached to 0.35 at pure nitrogen plasma, -40V in substrate bias voltage,

200W in microwave input power. The films do not deposit at less than -50V of substrate bias voltage,

because of re-sputtering the deposited film.

Nitrogen atoms in the films axe chemically bonded mainly with sp 3, and sp 2 hybridized carbon.

Increasing the nitrogen partial pressure and microwave input power in process parameter leads to the peak

• 3. 2 . . .

area ratio of (N-sp C) /(N-sp C). Decreasing the substrate bias voltage in process parameter leads to the

• 3 2 . +. . .

peak area ratio of (N-sp C)/(N-sp C). These results were affected the increase of the N2 1on density in the

ECR plasma.

NASA/CP--2001-210948 613 August 6, 2001
Acknowledgements

This work was partly supported by the High-Tech Research Center Program of the Ministry of

Education, Sports, and Culture, Japan. The authors would like to thank Mr. M. Sasaki and Mr. K. Sakayama

of the Industrial Research Center of Shiga Prefecture for Ra_nan spectroscopy measurements.

References

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2. A.Y. Liu, M. L. Cohen : Phys. Rev. B 41(1990) 10727.

3. J.H. Kaufman, S. Mertin, D. D. Saperstein : Phys.Rev. B 39(1989) 13053.

4. D. Li, Y. W. Chung, M. S. Wong, W. D. Sproul : J. Appl. Phys. 74(1993) 219.

5. K.M. Yu, M. L. Cohen, E. E. Haller et al. : Phys. Rev. B 49(1994) 5034.

6. T. Yen, C. Chou : Appl. Phys. Lett. 67 (1995)2801.

7. Y. Chen, L. Guo, E. G. Wang : Philosophical Magazine Lett. 75(1997) 155.

8. E. Kamijo, T. Nakamura, Y. Tani : Nucl. Instr. and Meth. B 121(1997) 110.

9. Y. Tani, Y. Aoi, E. Ka_nijo : Appl. Phys. Lett. 72(1998) 1652.

10. K.J. Boyd, D. Maxton, S. S. Todorov, et al. : J. Vac. Sci. Technol. A13(1995)2110.

11. D. Maxton, K. J. Boyd, A. H. A1-Bayati et al. : Phys. Rev. Lett. 73,1 (1994)118.

12. H. Sjostron, L. Hultman, J. E. Sundgren et al. : J. Vac. Technol. A 14, 1(1996) 56.

13. J.M. Ripalda, I. Montero, L. Galian : Diamond Relat. Mater. 7(1998)402.

14. I. Gouzman, R. Brener, A. Hoffman : Surf. Sci. 283(1995)331.,

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1999, p. 112.

NASA/CP--2001-210948 614 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

OPTICAL EMISSION STUDIED TO THE CARBON NITRIDE FILM DEPOSITION BY ECR

PLASMA SOURCE ENHANCED PLANAR MAGNETRON SPUTTERING

XU Jun, MA Tencai, ZHANG Jialiang, MEI Xianxiu, DENG Xinlu

State Key Lab for Materials Modification by Laser, Electron and Ion Beams, Dept, of Physics,

Dalian University of Technology, Dalian 116024, R R. China

ABSTRACT

Carbon nitride (CNx) thin films with proper chemical stoichiometry as well as excellent properties

were deposited by electron cyclotron resonance (ECR) plasma source enhanced planar magnetron

sputtering. The optical emission spectra from both the negative glow discharge zone near the

target surface and the plasma zone formed by ECR discharge and magnetron discharge were

studied separately. The dominant chemical species in both zone were found to be N2÷, N2", and

CN radicals. But the intensity of CN radical in plasma zone was greater than that in negative zone,

where a emission band of C2 was detected. The dependence of the intensities of N2+ and CN

radical on the deposition conditions such as working gas pressure and sputtering voltage were

investigated, and the formation of CN radicals, which could be vital in the CNx film deposition

process, was discussed. It was shown that a possible mechanism could be the reaction of N2÷ with

C2 in gas phase.

Key Words: Carbon nitride; ECR plasma; Sputtering; Optical emission spectra

Corresponding author: XU Jun

E-mail: xujun@dlut.edu.cn

Fax: +86-411-4708389

Tel: +86-411-4708384(0)

4708867(h)

NASA/CP--2001-210948 615 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

An XPS study of pulsed laser deposited CN x films

F. Llg- NORMAND, J. HOMMET

IPCMS, UMR 7504 CNRS, BP 20, 67037 Strasbourg Cedex 2, France

T. SZORENYI 1'2, C. FUCHS 1, E. FOGARASSY 1

CNRS-PHASE, UPR 292 CNRS, BP 20, 67037 Strasbourg Cedex 2, France

2

Research Group on Laser Physics, 6701, Po Box 406, Szeged, Hungary

We present an XPS study of CN x films prepared by pulsed laser deposition (PLD) at room temperature.

The Cls and Nls core levels display many contributions whose assignment is controversial in the literature, due

to the lack of solid reference compound and to the widely varying configuration states of both carbon and

nitrogen. We therefore develop a procedure that enables to simply determine the binding energies of each

individual lines of the C I s and N1 s core levels in a given carbon and nitrogen configuration, involving a close

comparison with literature references of solid polymeric compounds containing nitrogen and solid carbon

references. It is believed that polymeric compound involving nitrogen may provide more pertinent references

than molecules as they account more adequately of the screening of the ionized state of the photoemission

process. This screening effect is even more pronounced in the case of aromatic or cyanogen-type bonds where

delocalized rc electrons strongly affect the screening of the core hole. We therefore propose an assignment of

these lines that is tested on CNx films when changing the most important parameters of the PLD process (laser

fluence, nitrogen pressure, target-to-substrate distance) and of subsequent post-treatments (vacuum thermal

treatments and sputtering by argon ions). We find that the main environment is formed by N(pyramidal)-

C(trigonal) and N(colinear)-C(linear) configurations which can be explained by the trend to privilegiate sp 3 and

sp 2 hybridisation states around carbon and nitrile configuration around nitrogen.

Keywords

Pulsed Laser Deposition Carbon ni_ides X-my Photoemission Spectroscopy

Corresponding author

Frangois LE NORMAND

IPCMS / GSI

23, rue du Loess

67037 STRASBOURG Cedex

France

Tel: (33) 3 88 10 70 98

Fax: (33) 3 88 10 72 48

e-mail: Franc°is'Le'N°rmand@ipcms'u'strasbg'fr

NASA/CP--2001-210948 616 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

C3N4 Controversials and the Chinese Connection

James C. Sung

KINIK Company, Taipei, Taiwan, R.O.C.

Taipei University of Technology, Taipei, Taiwan, R.O.C.

Abstract

C3N4 research has ushered the way toward a new fronterior of science. It combines exotic idea with

theoretical modeling. It also triggered creative experimentals. This research has discovered that CN films can

be much harder and more stable than common DLC films. Hence CN films have been implemented in the

production of hard drives. The C3N4 research also prompted the synthesis of (Si,C)3N4 crystals that may rival

the hardness of cubic boron nitride. It is recommended that this material be mass Noduced using the V'LS

method. If this is accomplished, the economical (Si,C)3N4 superabrasive can replace much more expensive

industrial diamond and cubic boron nitride.

NASA/CP--2001-210948 617 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Improved Inversion Channel Mobility for 4H-SiC MOSFETs

Following High Temperature Anneals in Nitric Oxide

G.Y. Chung a), C. C. Tin a), J.R. Williams a), K. McDonald b), R.K. Chanana b), R.A. Weller c),

S.T. Pantelides b), L.C. Feldman b), O.W. Holland e), M.K. Das e) and J.W. Palmour e)

a) Physics Department, Auburn University, AL 36849

b) Department of Physics and Astronomy, Vanderbilt University, Nashville, TN 37235

c) Department of Electrical Engineering and Computer Science, Vanderbilt University, Nashville, TN 37235

a_ Solid State Division, Oak Ridge National Laboratory, Oak Ridge, TN 37831

e) Cree Research Inc., Durham, NC 27713

In the family of wide band gap materials (silicon carbide, the group III nitrides and diamond), SiC is the only

semiconductor that has a native oxide, and metal-oxide-semiconductor field effect transistors (MOSFETs) have been

fabricated using both 4H- and 6H-SiC. The 4H polytype has higher bulk carrier mobility [1], and is hence the

polytype of choice for power MOSFET fabrication. However, reported channel mobilities for 4H n-channel,

inversion mode devices are substantially lower than for 6H-MOSFETs. For power device applications, the

advantage provided by 4H-SiC of lower epilayer resistance for a given operating voltage is compromised by the low

channel mobility. Schorner, et al. [2] attribute the poorer performance of 4H devices to a large, broad interface state

density located at approximately 2.9eV above the valence band edge in both polytypes. More of these states lie in

the band gap for 4H-SiC (Egap - 3.3eV) compared to 6H-SiC (Egap - 3eV) where they act to reduce channel mobility

through field termination, carrier trapping and Coulomb scattering. Afanasev, et al. [3] proposed that interface

states in SiC/SiO2 structures result from carbon clusters at the interface and defects in a near-interface sub-oxide that

is produced when the oxidation process is terminated. The large interface trap density near the conduction band

edge proposed by Schorner, et al. has been observed experimentally for both n-SiC [4,5] and p-SiC [6]. Li, et al. [7]

originally reported improvements in the electrical performance of dry oxides on 6H-SiC annealed in nitric oxide

(NO). We have grown oxides on 4H-SiC using standard thermal techniques [8] and conducted post oxidation

anneals in NO [9]. We find that the interface state density near the conduction band edge in n-4H-SiC can be

reduced to levels comparable to the interface state density near the conduction band edge in 6H-SiC. Furthermore,

the effective channel mobility for inversion-mode 4H-SiC MOSFETs improves significantly following high

temperature anneals in nitric oxide.

Results are shown in Fig. 1 for standard hi-lo (quasi-static) C-V measurements on n-4H-SiC MOS capacitors

following NO passivation anneals of 2hr at 1175°C. "Re-ox" refers to samples that were not annealed in NO. "Reoxidation" is a standard wet oxidation termination step [10] that is carried out at several hundred degrees below the

oxidation temperature of 1100°C in order to reduce the interface state density near mid-gap for p-SiC/SiO2 [4,10,11].

As shown in Fig. 1, NO passivation reduces the interface state density significantly. The trap density at 0. leV

below the conduction band edge decreases by approximately one order of magnitude (from about 2X1013 to

2xl012eV-lcm -2) after NO annealing. AC conductance measurements performed at Cree Research, Inc. for similar

samples confirm the C-V data presented in Fig. 1.

The results of mobility measurements for a lateral 4H-SiC MOSFET fabricated with a dry oxide are shown in

Fig. 2. The device has a peak channel mobility of approximately 30cm2/V-s, and shows a x14 improvement in peak

mobility as a result of the NO passivation anneal. Similar results are obtained for wet oxide MOSFETs. Figure 3(a)

shows a plot of drain current versus gate voltage that was used to determine a peak effective mobility of 5cm2/V-s

for a device fabricated with Cree's standard dry-wet oxide. Figure 3(b) shows the improvement in mobility that can

be achieved with NO passivation. Drain current - gate voltage characteristics are also plotted in Fig. 3(b) for the

standard and passivated MOSFETs. The standard device exhibits the soft turn-on, high threshold, and single digit

mobility [Fig. 3(a)] typical of 4H-SiC MOSFETs. However, the MOSFET annealed in NO has a 5V threshold,

relatively sharp turn-on, and a peak channel mobility of 37cm2/V-s.

The characteristics of both the dry and wet oxide devices are consistent with a lower interface state density near

the conduction band edge. Also, the mobility versus gate voltage behavior for both the wet and dry oxide devices is

similar to that of silicon MOSFETs [12] with only a 15-20% reduction in mobility at the higher gate voltages where

the devices will typically be operated. However, unlike silicon, the mobilities reported here for SiC are a much

smaller fraction of the bulk mobility (-10% for SiC compared to - 50% for Si). This difference can be attributed to

the fact that, even after NO passivation, the interface state density near the conduction band edge is still ten times

NASA/CP--2001-210948 618 August 6, 2001
higherforSiCcomparedtoSi.Fullwafertestingforthewetoxidedevicesresultedinayieldof90%andan

averageeffectivemobilityof33cm2/V-s. Furthermore, everyworkingdeviceonthe3.5cmdiameter waferhada

peakmobilitygreater than30cm2/V-s - anindicationthatheNOpassivation processwasveryuniform.

Measuredchannel mobilitiesforlateral 4H-SiCMOSFETsfabricatedwithstandardepilayersandstandard

thermal oxidationtechniques aretypicallywellbelow10cm2/V-s. Anumber ofattempts- includingtheNO

passivation processreportedherein- havebeenundertaken inanefforttoimprovetheeffectivechannel mobility.

Sridevan, etal.[13]reportedn-channel mobilities as high as 165cm2/V-s for deposited oxides subsequently annealed

in wet N2 and Ar. Ogino, et al. [14] observed significant improvements in mobility (as high as 99cm2/V-s)

following low dose N implants into the channel region that were designed to adjust device threshold voltage. Yano,

et al. [15] fabricated MOSFETs with wet thermal oxides grown on the (1120) face of 4H-SiC and reported effective

channel mobilities of around 30cm2/V-s. However, to our knowledge, our results are the first to show substantial

improvement in the inversion channel mobility for lateral n-channel MOSFETs fabricated with standard thermal

oxidation techniques and standard 4H-SiC. Whether used separately or in conjunction with other techniques such as

low dose implantation, the NO passivation process represents significant progress in 4H-SiC MOSFET

development.

ACKNOWLEDGEMENTS: This work was supported by DARPA Contract # MDA972 98-1-0007 and EPRI

Contract # W0806905. Research at Oak Ridge National Laboratory sponsored by the Division of Material Sciences,

U.S. Department of Energy, under contract DE-AC05-00OR22725 with UT-B attelle, LLC.

[1] J.B. Casady and R.W. Johnson, "Status of silicon carbide (SIC) as a wide band gap semiconductor for high

temperature applications: a review," Solid State Electronics, vol. 30, no. 10, pp. 1409-1422, 1996.

[2] R. Schorner, P. Friedrichs, D. Peters and D. Stephani, "Significantly improved performance of MOSFETs on

silicon carbide using the 15R-SiC polytype," IEEE Electron Device Lett. vol. 20, no. 5, pp. 241-244, 1999.

[3] V.V. Afanasev, M. Bassler, G. Pensl and M. Schulz, "Intrinsic SiC/SiO2 interface states," Phys. Star. Sol. (A),

vol. 162, pp. 321-337, 1997.

[4] G.Y. Chung, C.C. Tin, J.H. Won and J.R. Williams, "The effect of Si:C source ratio on SiC/SiO2 interface state

density for nitrogen doped 4H and 6H-SiC," Materials Science Forum, vols. 338-342, pp. 1097-1100, 2000.

[5] M.K. Das, B.S. Um and J.A. Cooper, Jr., "Anomalously high density of interface states near the conduction

band in SiO2/4H-SiC MOS devices," ," Materials Science Forum, vols. 338-342, pp. 1069-1072, 2000.

[6] N.S. Saks, S.S. Mani, and A.K. Agrwal, "Interface trap profile near the band edges at the 4H-SiC/SiO 2

interface," Appl. Phys. Lett., vol. 76, no. 16, p. 2250-2252, 2000.

[7] H. Li, S. Dimitrijev, H.B. Harrison and D. Sweatman, "Interracial characteristics of N20 and NO nitrided SiO2

grown on SiC by rapid thermal processing," Appl. Phys. Lett., vol. 70, no. 15, pp. 2028-2030, 1997.

[8] G.Y. Chung, C.C. Tin, J. H. Won, J.R. Williams, K. McDonald, R.A. Weller, S.T. Pantelides and L.C.

Feldman, 'Interface state densities near the conduction band edge in n-type 4H- and 6H-SiC," Proc. 2000

IEEEAerospace Conf, vol. 1, pp. 1001-1005, 2000.

[9] G.Y. Chung, C.C. Tin, J.R. Williams, K. McDonald, M. Di Ventra, S.T. Pantelides, L.C. Feldman and R.A.

Weller, "Effect of nitric oxide annealing on the interface trap densities near the band edges in the 4H polytype

of silicon carbide," Appl. Phys. Lett., vol. 76, no. 13, pp. 1713-1715, 2000.

[ 10] L.A. Lipkin and J.W. Palmour, "Improved oxidation procedures for reduced SiO2/SiC defects," J. Electronic

Materials, vol. 25, no. 5, pp. 909-915, 1996.

[ 11] M.K. Das, J.A. Cooper and M.R. Melloch, "Effect of epilayer characteristics and processing conditions on the

thermally oxidized SiO2/SiC interface," J. Electronic Materials, vol. 27, no. 4, p. 353-357, 1998.

[ 12] S.C. Sun and J.D. Plummet, "Electron mobility in inversionand accumulation layers on thermally oxidized

silicon surfaces," IEEE Trans. Electron Dev., vol. 27, no. 8, pp. 1497-1508, 1980.

[13] S. Sridevan and B.J. Baliga, "Lateral n-channel inversion mode 4H-SiC MOSFETs', IEEE Electron Device

Lett. vol. 19, no. 7, pp.228-230, 1998.

[14] S. Ogino, T. Oikawa and K. Ueno, "Channel doped SiC MOSFETs', ," Materials Science Forum, vols. 338-

342, pp. 1101-1104, 2000.

[15] H. Yano, T. Hirao, T. Kimoto, H, Matsunami, K. Asano and Y. Sugawara, "Anisotropy of inversion channel

mobility in 4H- and 6H-SiC MOSFETs on (1120) face", Materials Science Forum, vols. 338-342, pp. 1105-

1108, 2000.

NASA/CP--2001-210948 619 August 6, 2001
Fig.1. Interfacestatedensityforn-4H-SiC MOS

capacitors before and after anneals in NO at 1175°C

for 2hr. The hi-lo C-V measurements were made at

room temperature and confirmed with AC

conductance measurements. "Re-ox" refers to

samples that were not annealed in NO. In agreement

with the suppositions of Schorner, et al. [2], the

passivation anneal is noticeably more effective for

4H-SiC.

1014

%-

> 1013

E

o

1012

£3

10_

2.4

• reox 6H

[] NO

• reox 4H

• NO

r_

q.

[]

2.6

i ,

Ec (4H)

Ec (6H) ,,"

A

k

A

A

A

A

d? • _

v v

v

2.8 3.0 3.2 3.4

E - Ev (eV)

Fig 2. Field effect mobility for dry oxide 4H-SiC

MOSFETs fabricated with and without an NO

passivation anneal at 1175°C for 2hr.

30

>

E

o 20

>,,

.I3 10

o

VDS -- 50mV Dry Oxide + NO

5 10 15 20

Gate Voltage (V)

Fig. 3. (a) Drain current as a function of gate voltage

for a MOSFET fabricated with Cree's standard, drywet oxidation procedure. Analysis (0Id/0Vg, Vds =

50mV) resulted in a peak effective mobility of

5cm2/V-s. (b) Mobility versus gate voltage and Id-Vg

characteristics for a device fabricated with a wet

oxide grown at Auburn and an NO anneal performed

at Vanderbilt. Id-Vg characteristics for the standard

device are included for comparisons that show

reduced threshold voltage and significantly improved

turn-on behavior for the passivated MOSFET.

0.25

<

"_ 0.2

E

,_ o.15

tO o.1

•_ o.o5

D

o

2.5

,, ,i .... i .... i .... i ....

VDS = 50mV (a)

M°bility = 5°m2N-s /_,/_

5 10 15 20 25

Gate Voltage (V)

.... t .... t .... i .... i ....

2 0 VDS -- 50mY (b) ,

1.5

._. 1"014

o5 - / / I Cree

° I / l

or..-.4 .... _---_-_'_-...- , ....

0 5 10 15 20

Gate Voltage (V)

40

30o

orem. ,-i.

,.<

20 ,_-

B

lo ,<

o

25

NASA/CP--2001-210948 620 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

STUDY OF GAP STATE DISTRIBUTIONS IN HYDROGENATED AMORPHOUS

SILICON CARBIDE GROWN USING ECR-CVD

K. Chew 1, Rusli, S. E Yoon, J. Ahn, Q. Zhang and V. Ligatchev

Microelectronics Division, School of Electrical and Electronic Engineering, Nanyang Technological University,

Singapore 639798, Republic of Singapore

E. J. Teo, T. Osipowicz and E Watt

Research Centre for Nuclear Microscope, Physics Department, National University of Singapore, 10 Kent Ridge

Crescent, Singapore 119260

ABSTRACT

The optical characteristics of silicon-rich hydrogenated amorphous silicon carbide (a-Sil_xCx:H) films with

varying carbon composition (x) are investigated using photothermal deflection spectroscopy (PDS). It is found that

the optical gap and the degree of disorder in the films are strongly affected by carbon alloying. The deconvolution of

the PDS spectra enables us to deduce the density of states (DOS) distribution of the films and the sub-gap features

observed are discussed in detail.

Keywords: density of states (DOS), photothermal deflection spectroscopy (PDS), hydrogenated amorphous silicon

carbide (a-Sil_xCx:H), absorption coefficient (cQ

INTRODUCTION

Hydrogenated amorphous silicon carbide (a-Sil_xCx:H) has attracted great interest due to its attractive potentials

in electronic applications, such as solar cells, optoelectronic devices, high-temperature engineering materials and

coating materials. There are several technological advantages in alloying carbon with hydrogenated amorphous

silicon (a-Si:H). Among them, carbon alloying increases the bandgap energy, while retaining its ability to be doped

n- or p-type. However, it also increases the gap state densities in the material, causing a degradation in its

photoelectric properties. Therefore, an investigation of the density of states (DOS) of a-Sil_×C×:H films with varying

carbon fraction is important towards the understanding of their electronic properties. In this work, the distributions of

gap states in Si-rich a-Sil_xCx:H (x < 0.4) films are investigated based on their sub-gap optical absorption measured

using photothermal deflection spectroscopy (PDS).

EXPERIMENT

The a-Sil_xC_:H films are prepared using an electron cyclotron resonance chemical vapor deposition (ECRCVD) system, the detail of which can be found elsewhere (ref. 1). The microwave was set to 900W and the upper

and lower magnetic coil currents were fixed at 100 A and 120 A respectively. The deposition pressure was

maintained at 20 mTorr and no intentional heating was applied. The hydrogen flow rate was maintained at 100 sccm,

while the gas ratio of methane to silane was varied to produce films with changing carbon-to-silicon (C/Si) ratio. The

carbon fractions x of the films were deduced from Rutherford backscattering spectrometry (RBS) and their optical

band gaps using the transmittance and reflectance measurements based on a dual beam Perkin-Elmer Lambda 16

UV-Vis spectrophotometer. PDS was used to measure the optical absorption coefficients (c 0 of the samples. The

pump beam was a 1000 W tungsten lamp and the probe beam was a 632.8 nm He-Ne laser. The PDS spectra are

normalized to the absorption spectra obtained from UV-Vis spectroscopy.

1 Address: Clean Room & Characterization Laboratory, S 1-B2C-21, School of Electrical and Electronic Engineering,

Nanyang Technological University, Nanyang Avenue, Singapore 639798. Email: __-a3_4_2_8_'O5"_._@_n_t___e___u_.,_sg, Phone:

(65)7904528, Fax: (65)7933318.

NASA/CP--2001-210948 621 August 6, 2001
RESULTS AND DISCUSSIONS

Figure 1 shows the normalized absorption spectra obtained from PDS. The experimental data are shown in

symbols, whereas the solid lines represent the spectra calculated based on the deconvolution procedure proposed in

ref. 2 and will be elaborated further in the later part of the discussion. The optical parameters that can be derived

from the spectra include the optical gap Eg and the B parameter in the Tauc's model _ = B(hv - Eg ), where ot is

the energy dependent absorption coefficient and hv is the photon energy.

¢0

©

Ea

10 8 ,

10 7.

10 %

10 5.

10 4,

10 _ •

O

10 0

10 112 114 116 116 210 212 214 216 216

Photon Energy, hv (eV)

Figure 1. Experimental (symbols) and calculated (solid lines) PDS spectra based on the DOS

deconvolutions of a-Sil.xCx:H with varying composition x. The spectra are displaced vertically by a decade

for clarity.

The Tauc gaps Eg are shown in Fig. 2, where an increase with x is seen, as expected from the results shown in

Fig. 1. For a-Siz_xCx:H, there are uncertainties in deducing Eg as the linear region needed for the linear extrapolation

to ot -- 0 is narrow, due to their large band gaps. Therefore, another parameter E04, which represents the photon

energy at which c_ reaches 104 cm -1, is also determined. It has a same trend as Eg, though consistently higher by

approximately 0.3 eV. The increase in the optical gap with carbon incorporation has already been reported for a-Siz_

xCx:H (refs. 3 and 4). The reason behind the increase is elucidated by calculations performed by Robertson (ref. 5)

for crystalline silicon carbide (c-SiC), which can serve as a first approximation to its amorphous counterpart. It was

shown that for Si-Si bonds, the cy-like states are situated at 0.5 eV below the valence band edge (Ev), whereas the cy*-

like states are situated at 1.7 eV above the conduction band edge (Ec) of c-SiC. Therefore, broadening of the Si-Si cy

states into a band in a-Siz_xCx:H will affect the DOS near Ev, but with little effect at Ec. As x is decreased, a large

number of Si-Si bonds are present, prompting the Si-Si cy states to broaden into a band, which raises E_ closer to Ec

and thus lowers the energy gap.

A comparison is also made between the optical gaps of our samples that are highly hydrogen diluted with those

of samples deposited by cosputtering, which used two separate silicon and carbon targets in an argon and hydrogen

gas atmosphere (ref. 6). As can be seen in Fig. 2, the Tauc optical gaps for the cosputtered films, denoted by Egcosp,

are consistently lower than that of our films. This difference can be due to the different hydrogen contents in the

films. The highest hydrogen content in the cosputtered films measured by thermal evolution was about 30%, which

is much lower than that in our films, deduced from elastic recoil detection analysis (ERDA), that can reach -50%.

Such high hydrogen contents are consistent with those reported for typical glow-discharge films (ref. 6 and 7). In Sirich samples, hydrogen tends to passivate (Si) dangling bonds, resulting in a lattice of Si-H sites, or break strained

NASA/CP--2001-210948 622 August 6, 2001
Si-Sibonds,toformstronger Si-Sibonds,therebyrecedingthevalencebandandleadingtoasharper andmore

abruptvalencebandedge.Consequently Egwillincrease, ascanbeconfirmedfromFig.2.

2.5-

2.4-

2.3-

2.2-

;>

d3

2.1-

_'_ 2.0-

.

"_ 1.9

©

1.8-

1.7-

1.6

--A-- Eo 4

----e--Eg I

--I-- ggcosp ._

010 011 012 013 014

Carbon fraction, x

i

0.5

Figure 2. The dependence of Eg and E04 on carbon composition x. Plots for Egco_p for cosputtered films (ref. 6)

is also shown for comparison.

800 -

750 -

700 -

650 -

-> .

¢_ 600-

550-

500 -

450-

400

i -,-B

'2°4 _ I --e--E

_,,0_ _ .'--. i u

e,ooJ _";'--."4-. I---°--%

80 _

4404_0 4_0 _0 @ @ _0 _0 _0

i i , i , i , i , i , i , i

0.00 0.05 0.10 0.15 0.20 0.25 0.30 0.35

Carbon fraction, x

16O

15O

140

>

130 ._

120

_=

110

100

90

80

0.40

Figure 3. The dependence of the parameter B and Urbach energy, E. and Ew, on carbon fraction, x. Inset:

the correlation between B and E. and Ev.; the solid lines represents the linear relation between them.

Figure 3 plots the Tauc parameter B and the Urbach energy Eu as a function of x. The former represents the joint

optical density of states whereas the latter is a measure of the degree of disorder in the films obtained from (z --

(Zoexp(hv/E,_), where (zo is a pre-exponential factor. The valence Urbach energy, E_u, deduced from the

deconvolution procedure of the spectra shown in Fig. 1, is also included. A decrease in B and an increase in Eu and

E_u with x can be seen. For x _< 0.22, B values are quite high, indicating that the films are Si-rich (ref. 3). For such

samples, the band edges are Si-like, so it is expected that their characteristics are similar to a-Si:H, where the tail

NASA/CP--2001-210948 623 August 6, 2001
widthiswideratthevalencebandthanattheconductionband(ref.2and5).Assuch,Eucanbetakentorepresent

thevalencebandtailwidth.Indeed, thisisalsotheassumption madefortheDOSdeconvolution ofourPDSspectra

todeducethevalenceUrbachenergy, E,_,whichshowsastrongcorrelationtoEu.Furtherdiscussions ontheDOS

deconvolution resultswillbepresentedshortly.Thebandtail statesareconsideredtobeweakSi-Sibonds,arising

fromthestrainpresent inthedisorderednetwork(ref.8). Carbonalloyingtendtoinducemorestructural disorder

andstraininthenetworkbecauseofthedifferent bondlengthsandbondstrengthsofcarbonandsilicon(ref.9).

HenceEuwillincreasewithcarbonalloying,asseeninFig.3.It isthoughtthatBisrelatedtotheoverallstructural

disorder, whileEuisrelatedtolocalmicrostructural disorderresultingfromcarbonalloying.Alinearrelationis

foundbetweenBandEu,whereEu(meV)=279- 0.32B(cmeV)-1/2,asshownintheinsetofFig.3. Thisindicates

thathemicrostructural defectsarisingfromcarbonalloyinghaveapronounced effectontheoveralldefectstructure

ofthefilms.AsimilartendisalsoobservedforF_.

ToobtaintheDOSdistributionfromthePDSspectrashowninFig.1,wefollowthedeconvolution procedureof

ref.2. Generally, theDOSdistributionfamorphous semiconductors canbedividedintothreeregions: apower-law

regionattributedtotransitions involvingextendedstates, anexponential absorption(Urbach) edgewhichencroaches

intothebandgapanddefectsstatesdeepinthepseudogap. All possibletransitionsbetweenthethreeregionsare

consideredinthedeconvolution process. Thefittingparameters involvedinthedeconvolution includetheconstant

M,whichisassumedtobeindependent ofenergyandproportional totheaveragevalueofthemomentummatrix

elementsinTauc'sopticalabsorptionmodel,theopticalgapEg,thevalenceUrbachenergyE,_,thefilleddefect

densityNdanditspositionEpk,andtheemptydefectdensityNd,anditspositionEpk,inthepseudogap. Asinthe

caseof a-Si:H,thefilledandemptydefectsareassumedtobeoriginatedfromdanglingbonds.Althoughmany

modesof transitionsarepossible, butnotallof themwill haveaninfluential effectontheoverallabsorption

coefficient c(.Fromthedeconvolution procedure, it isfoundthathecontributions oftheemptydefectstatesNd*,

whicharisemainlyfromtransitionsfromthevalenceband,tothetotalabsorptionareobscuredbythemuchlarger

contributionfromthetransitionsbetweenthevalencebandtailstatesandtheconductionband,duetothemuch

highervalencebandtailstatedensities. This,coupledwiththenoisiersignalatlowerphotonenergies, resultsina

largeuncertainty indetermining Nd,.

©

*d

1E19- 1E19

s_1 E18

1E18-. _E_77'O- " 8'o " 9'o " _;o " _o " _o " GO

ii_m/m

--m-- N d

1E17

i / u i u u i i

0.00 0.05 0.10 0.15 0.20 0.25 0.30 0.35 0.40

Carbon fraction, x

Figure 4. The dependence of filled defect densities Nd and empty defect densities Nd* on carbon fraction x.

Inset: the relation of Nd and Nd, with N_u. The solid line shows the correlation between Nd and Ev,.

NASA/CP--2001-210948 624 August 6, 2001
ThebestfittedfilleddefectdensitiesNddeducedforthefilmsareshowninFig.4asafunctionofx. The larger

Nd seen at increasing carbon fraction x can be interpreted as an increase in the dangling bond defect densities,

concomitant with an increase in the valence band tail slope, E,_,, discussed earlier. The variation in the defect and

band tail state densities can be explained by generalizing the weak bond to dangling bond conversion model of

Stutzman (ref. 10). This model attributes tail states in a-Si:H to weak bonds, which when lie beyond a certain

demarcation energy above the valence-band edge, are converted to dangling bonds by bond breaking so that they can

minimize their energy locally. Therefore, it is expected that a wider band tail tends to give rise to higher defect

densities. Indeed, if we plot the defect densities Nd (believed to be associated with dangling bonds), against E,_,, as

shown in the inset of Fig. 4, a correlation can be seen, where N d oc exp (E,_/13 meV). This tend to suggest the

disorder-induced fields and strains responsible for controlling the slope of the Urbach edge are caused by dangling

bond defects (ref. 11). On the other hand, the empty defect densities Nd* remains constant for x -- 0 and x -- 0.18, and

the type of correlation seen for Nd is not found, probably due to the large uncertainty in the determination of Nd* as

mentioned before.

©

Z

o

1E25. i I_ xl000

1E24-_ _1_. i _._x,

1E23- -_ ........ _ _,_............. _ ...... _ ...... _.__"_"___.... .h__,

1E22- ......... l---- ,_c_- J', - - - - zr_z_i_L --

1E21 - - _7._'_ '_ 'a_xl '_/d/_]/i_ i

,

1E20. _ _ '_ O " '

uu %_'_ _x A',X J/LJ', ,,

:k __ o_'°6 \ t'g/ I ,

1E19- __n__X= 0 "b-'_'d/"°n 0 [OZ'_/u _ [

1E18 _ --o--x=0.18 ____'_auNS_ 4/ i ', 'ln rl-J 'q" P ll_ q:J / l l

--a--x=0.22 -t_ _:1 I u [] ' l

1E171 ' _/ I '

--v--x= 0 36 , Lu , ,

i i

1ElS-I , i , I , i

-3.0 -2.5 -2.0 -1.5 -1.0 -0.5 0.0 0.5

Energy, s (eV)

Figure 5. The density of states (DOS) distribution calculated from the deconvolution of the PDS spectra

shown in Fig.1. The spectra are displaced vertically by a decade for clarity.

The DOS distributions derived from the deconvolution of the PDS spectra are shown in Fig. 5. The origin for

the energy scale, s, refers to the conduction band edge, Ec. In order to obtain a good fit between the experimental

and calculated spectra, the peak energy for the mid-gap defect distribution for x < 0.22, Epk is taken to be located at -

1.0 eV below Ec. However, for x -- 0.36, a good fit can only be obtained if Epk is shifted to -1.2 eV below Ec. One

plausible reason is that as the carbon fraction x increases, the Si dangling bonds are no longer the major contributor

to the deep defect densities and C dangling bonds may also play a role. The significance of C dangling bonds is

supported by calculations performed for the density of defects in a-Sil_xCx:H (ref. 5 and 12). As evidenced from ESR

(electron spin resonance) measurements (ref. 6), carbon alloying in a-Si:H produces resonance g values between that

of Si (g = 2.0055) and C (g = 2.003), indicating that the spin densities are made up of Si and C singly occupied

dangling bonds, Si3°and C3°, where the subscript represent the coordination number and the superscript represent the

charge state, with symbols + for positive, - for negative and 0 for neutral. On the other hand, LESR (light induced

electron spin resonance) measurements (refs. 13 and 14) found that, in addition to the neutral dangling bonds, the

spin densities are also contributed by charged defects of Si and C. Owing to the higher electronegativity of C, the Si3

level lies above the C3 level, and there will be a charge transfer to give C charged defects with a configuration

C 3 and Si charged defects, Si_ (ref. 5). The Si3 level is repelled upwards as C is introduced into Si while the C3

level is repelled downwards as Si is introduced into C. If Si3 outnumber C3 centers, as in Si-rich alloy, this will

NASA/CP--2001-210948 625 August 6, 2001
produce Si ,Si0and centers, where isrepelled downwards towards Evwhile isrepelled upwards

towards Ec by the interaction among the Si and C atoms. It is probable that at higher x, C 3 outnumbers Si 0 , and

that accounts for the observed shift of the defect level closer towards the valence band.

CONCLUSIONS

The absorption spectra of a-Sil_xCx:H with varying carbon fraction x, are studied by means of photothermal

deflection spectroscopy (PDS). From the spectra, it is found that the optical gap, parameter B and Urbach energy are

well correlated with x. The possible reasons for the correlation among the parameters are discussed. From the

deconvolution of the PDS spectra, the DOS of the samples with varying x can be found. Variation in the defect peak

densities and their energy levels are discussed in term of carbon alloying.

REFERENCES

1. Rusli, et al.: Effect of pressure on the deposition of hydrogenated carbon films using the electron cyclotron

resonance chemical vapor deposition. J. Vac. Sci. Tech. A, vol. 16, no. 2, Mar/Apr 1998, 573.

2. Curtins, H.; and Favre, M.: Surface and bulk states determined by photothermal deflection spectroscopy.

Amorphous silicon and related materials, World Scientific, Singapore, 1988, vol. 1,329 - 363.

3. Conde, J., et al.: Optoelectronic and structural properties of amorphous silicon-carbon alloys deposited by lowpower electron-cyclotron resonance plasma-enhanced chemical-vapor deposition. J. Appl. Phys., vol. 85, no. 6, Mar

1999, 3327-3338.

4. Desalvo, A., et al.: Optoelectronic properties, structure and composition of a-SiC:H films grown in undiluted and

H2 diluted silane-methane plasma. J. Appl. Phys., vol. 81, no. 12, Jun 1997, 7973-7980.

5. Robertson, J.: The electronic and atomic structure of hydrogenated amorphous Si-C alloys. Philos. Mag. B, vol.

66, no. 5, 1992, 615-638.

6. Palsule, C., et al.: Photoluminescence characterization of a-Sil_xCx:H alloys prepared by cosputtering. Phy. Rev.

B, vol. 48, no. 15, Oct 1993, 10804-10814.

7. Robertson, J.: The electronic and atomic structure of hydrogenated amorphous Si-C alloys. Philos. Mag. B, vol.

66, no. 5, 1992, 616 and reference thereof.

8. Unold, T.; Hautala, J.; and Cohen, J. D.: Effect of carbon impurities on the density of states and the stability of

hydrogenated amorphous silicon. Phy. Rev. B, vol. 50, no. 23, Dec 1994, 16990.

9. E1 Khakani, M. A., et al.: Composition and thermal-annealing-induced short-range ordering changes in

amorphous hydrogenated silicon carbide films as investigated by extended x-ray-absorption fine structure and

infrared absorption. Phy. Rev. B, vol. 51, no. 8, Feb 1995, 4903-4914.

10. Street, R. A., and Winer, K.: Defect equilibra in undoped a-Si:H. Phy. Rev. B, vol. 40, no. 9, Sep 1989, 6236-

6249.

11. Amer, N. M. and Jackson, W. B.: Optical Properties of Defect States in a-Si:H. Semiconductors and

Semimetals, Academic Press, 1984, vol. 21, part B, chap. 3.

12. Alvarez, E; Sebastiani, M.: Equilibrium density of defects in hydrogenated amorphous silicon carbon alloys. J.

Appl. Phys., vol. 71, no. 15, Jun 1992, 5969-5975

13. Xu, X., et al.: Thermal-equilibrium defect in undoped hydrogenated amorphous silicon, silicon-carbon, and

silicon-nitrogen. Phy. Rev. B, vol. 41, no. 14, May 1990, 10054.

14. Shimizu, T., et al.: Nature of localized states in hydrogenated Si-based amorphous semiconductor films

elucidated from LESR and CPM. Jpn. J. Appl. Phys., vol. 28, no. 4, Apr 1989, 590.

NASA/CP--2001-210948 626 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

The Carbon Implantation in Silicon by Cathodic Arc

Wen Ting Yeh 1, Ming-Chi Kan 2, Sea-Fue Wang 3, James Sung 1

1. Kinik Company, and National Taipei University, Taipei, Taiwan, R.O.C.

2. Department of Materials Science and Engineering, National Cheng-Kung University, Tainan,

Taiwan, R.O.C.

3. National Taipei University, Taipei, Taiwan, R.O.C.

Abstract

Carbon atoms were implanted into silicon substrates by cathodic arc at 1000 V

with a range of currents that were varied from 20 A to 100 A. AES analyses

confirmed that the silicon surface was clean without oxygen contamination. The

implanted carbon atoms penetrated below the surface with a rapid decrease of

concentration. It implied that the diffusion of carbon atoms was minimal so the

implantation was achieved primarily by physical means.

The depth of the implanted carbon atoms increased with the arc current, but it

peaked at 80 A at a depth of approximately 100 nm. ESCA disclosed a similar

profile of carbon bonding for the amorphous diamond accumulated on the surface.

The maximum sp 3 content was about 50% that occurred when current was at 80 A.

Cathodic arc contains ionized carbon atoms that were mixed with clusters of

molten carbon droplets. The latter froze to form micro particles that lodge on the

deposited film of amorphous carbon. The carbon atoms in these micro particles

were primarily held by sp 2 bonding. With the increase of arc current, the

temperature rose, so was the amount of micro particles. It is postulated that the

reduction of sp 3 content when the current was increased to 100 A was due to the

overheating of the sample by the electrical arc. Moreover, as the speed of ionized

carbon atoms was reduced by the shielding effect of these micro particles piled up on

the surface of the amorphous diamond. As the consequence, the implantation depth

became shallower.

Introduction

Diamond-like carbon (DLC) coating has numerous superb properties such as

high hardness, low friction, and extreme inertness. Moreover, it can cover the

surface of a variety of materials, even with complicated shapes. Hence DLC has

become the premium coating material for many advanced applications, such as hard

drives, razor blades, and acid barriers.

DLC is a composite material that encompasses diamond, graphite, and polymer

(hydrogen source). DLC that contains pure carbon is called amorphous diamond.

NASA/CP--2001-210948 627 August 6, 2001
Thismaterialis composedof randomlypackedcarbonatomsthatarebondedin

distortedtetrahedral coordination(sp3)andtrigonalayout(sp2). Thehigherof the

ratioamongthesetwoendmembers, themorediamond-likethepropertiesof the

coating.

Amorphousdiamondcanbeconvenientlydepositedby cathodicarc(Fig.1).

Thearccontainscationsof carbonthatarevaporizedfroma graphiteor carbon

cathode.Thesecarboncationscanbedriventowardasubstratehatisbiasedata

negativelectrical potential. Thekineticenergyofthecationsoftendeterminesthe

fateof thebondingcarbon. If theenergyislow,theimpingingcarbonatomscannot

formsp3bondingsotheybecomeamorphousgraphite. Ontheotherhand,if the

energyis sufficientlyhigh(e.g.,50-100eV),amorphousdiamondwill beformed.

However,if thecarboncationsaremovingtoofast,theycouldpenetratebeneaththe

surfaceofthesubstrate, wheretheyarelodgedasimplantedatoms.

Magnetic Co I[" Ar °r H21nle_

Topum °

Water

Fig. 1: The schematic of cathodic arc (left) and the appearance of such an arc (right).

The most difficult problem to apply amorphous diamond is to make the coating

stick to the substrate. In order to increase the adherence strength, the surface of the

substrate must be thoroughly cleaned to avoid any contamination. Furthermore, it is

often desirable to implant the substrate with carbon atoms before actually coating the

surface. In this way, amorphous carbon coating can be anchored by having this

"root."

The amount of implanted carbon atoms in a given time can be increased by

enhancing the negative bias of substrate and/or the electrical current (cation flux) of

the arc. In this research, the concentration profiles of implanted carbon atoms were

studied under a range of arc current. The substrate chosen was silicon, the most

compatible material for applying amorphous diamond coating due to its tendency to

form carbide (SIC) and its perfect match of thermal expansion with diamond.

Experimental

The substrates were square chips of N-type silicon. They were cut from a

wafer that was oriented with (100) face parallel to the surface. The resistivity of

these samples were in the range of 0.02_0.008 _2-cm. All samples were submerged

in HF (10%) contained DI water for 20 seconds. They were subsequently

ultrasonically agitated in DI water for 15 minutes. After cleaning, all samples were

stored in ethanol solution (95%) until use. Before conducting the experiment, each

sample was blown dry using nitrogen gas.

During the experiment, the chamber was maintained at a vacuum level of

2.0x10 -5 Torr. The bias of the substrate was set at 1000 V. The surface of each

sample was first etched by ions of argon and hydrogen for 20 minutes to remove any

NASA/CP--2001-210948 628 August 6, 2001
remainingcontaminant.It wasthenbombardedbycarboncationsfor1minutewith

currentcontrolledat20,40,60,or 80ampers.Duringtheion bombardment, the

temperatureof thesampleroserapidly. Topreventthesamplefromoverheating, it

wasattachedto a stagethatcontainedcirculatedwaterfor cooling. Evenso,

temperaturecouldstillriseabovethestabilityofamorphousdiamond(about500°C).

Hence,thearcwasturnedonintermittentlyin every12secondswith1-minuterest

timeinbetween.

TheimplantedsampleswereanalyzedbyAugerelectronspectroscopy (AES)to

revealthecarbonprofileasafunctionof diggingtimebytheelectronbeam. The

penetrationtimeis proportionalto thedepthbeneaththesubstratesurface.In

addition,thebondtypesandtheirratio(sp3/sp2)oftheamorphousdiamonddeposited

onthesamplesurfacewereanalyzedbyESCA.

Results and Discussions

Figure 2 shows the carbon profiles of different samples bombarded at various

currents of cathodic arc. The carbon concentrations were complementary to silicon

concentrations. Based on these profiles, the original depth of the substrate surface

could be determined by noticing the appearance of silicon atoms at the expense of

carbon atoms. The maximum depth of implantation may also be identified by the

disappearance of carbon atoms. In addition, the intermediate depth where carbon

and silicon atoms had equal abundance was located by the cross over point of the two

opposite concentration profiles for carbon and silicon. These three characteristic

depths of carbon implantation for each sample were summarized in Table 1, and they

were plotted graphically as shown in Figure 3.

i

/

_'---- ) " ,0 _. 5- : ' - G-.... &....

20 A 40 A

',. ,/

-.. ,.

,z'_L

60 A 80 A

Fig. 2: Concentration profiles of samples bombarded by cathodic arc at various levels

of currents.

NASA/CP--2001-210948 629 August 6, 2001
Tablel: ImplantationDepths(inpenetrationtimeofseconds)

Arc current (A) 20 40 60 80

Original Interface 0 0 0 85

Cross Over Point 10 sec 30 sec 56 sec 235 sec

Maximum Depth 100 sec 130 sec 180 sec 370 sec

200 _..-..-..-..-..-..-..-..-..-..-..-..-..-..-..-..-..-..-..-._._._._._._._._._._._._._._._+5

100 ..................................... I

0 ................................ I

SIK2 100 _ _

2oo i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:U_5._i

3OO

F.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'.'." ..... '.'.'1

20 40 60 80

Arc cu-K_ (A)

Fig. 3: Carbon implantation depths as a function of arc current.

Figure 4 shows the sp3/sp 2 ratio of carbon bonds as determined by ESCA.

100

80

60

4O

20

0

........................................ I

£i£i£i£i£i£i£i£i£i£i£i£i£i+i

iiiii_ii_iiiiiiiiiiii_i_iiiiiiiiiiiii_iiiiiiiiiiii_iiiii'. In sp3_a_%)]

33 50 80 100

Arc cu-K_ (A)

Fig. 4: The sp3/sp 2ratio of amorphous diamond deposited at various arc currents.

The above data indicated that the penetration depth of carbon atoms increased

rapidly with the surging of arc current. Although the penetration depth may be more

dependent on the kinetic energy of the carbon atom and hence the voltage of the

negative bias at the substrate best controls it, the increase of the flux of carbon atoms

had allowed the spread of kinetic energies of carbon cations. As a result, more

carbon cations had acquired higher energy to penetrate deeper. However, the

penetration depth was peaked at approximately 100 nm (370 seconds) when the

current reached at 80 A. Further increases of the arc current drastically raised the

temperature of the electrical arc. The consequence was the melting of graphite

cathode

De and the splashing of micro droplets of molten carbon. The piling up of

these carbon nano-particles on the substrate surface had slowed down the speed of

impinging carbon cations. The result was a decrease of implantation depth.

NASA/CP--2001-210948 630 August 6, 2001
Thesp3/sp2ratiosoftheamorphousdiamondcoatingrevealedthesametrendthat

peakedatanarccurrentof 80A. Thedecreaseof diamondcontentinthethinfilm

athigherarccurrentwaslikelyduetotheoverheatingofthesubstrate.

Conclusion

The carbon penetration depth in silicon is determined by the speed of impinging

cations. With the increasing of the cation flux, more carbon atoms can acquire a

higher kinetic energy to reach a greater depth. At a bias electrical potential of 1000

V, the maximum penetration depth is attained at approximately 100 nm when the arc

current reaches at 80 A. Further increases of the arc current tend to produce micro

particles that may pile up on the surface of the amorphous diamond. The shielding

effect of these micro particles can reduce the speed of impinging carbon atoms and

hence the implantation depth of carbon atom becomes shallower.

Acknowledgement

The authors would like to thank KINIK Company, Taiwan, R.O.C. and

Department of Materials Science and Engineering, National Cheng-Kung University,

Tainan, Taiwan, R.O.C for its financial support.

References

1. Roy, R. 1992, Press conference for diamond synthesis by the low-pressure

solid-state source, Penn State University, U. S. A.

2. Bohr-Ran Huang et al.: The surface properties on both side of the isolated

diamond film, Material Science and Engineering B64 (1999) 187-191

NASA/CP--2001-210948 631 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ELECTROMAGNETICALLY ACCELERATED PLASMA SPRAYING FOR BORON

CARBIDE COATING

J. Kitamura*

Joint Research Consortium of FCT, Japan Fine Ceramics Center, c/o AIST

1-1 Higashi, Tsukuba, Ibaraki 305-8565, Japan.

Tel.: +81-298-61-9982; Fax +81-298-61-4796; E-mail: j unya@mmc.go,jp

S. Usuba, Y. Kakudate, H. Yokoi, K. Yamamoto, and S. Fujiwara

National Institute of Advanced Industrial Science and Technology (AIST),

ABSTRACT

A new plasma spraying method using a coaxial cylindrical electrode plasma gun is presented to synthesize hard

and dense ceramics coatings with high adhesion. Pulsed arc plasma is generated between center anode and outer

cathode by supplying pulsed large electric current. The plasma is compressed and accelerated by electromagnetic

force generated by the large electric current. The plasma heats and accelerates source powder materials injected into

the electrode using a pulsed powder injector with a large flow rate. The pulsed powders injection method is also

presented. The heated and accelerated source powders deposit onto substrates setting at the exit of the electrode. The

measured maximum electric current is 100kA. The speed of the accelerated plasma is ranged from 1.5 to

3.0x 103naJsec and the estimated maximum electromagnetic pressure is approximately 1MPa. Dense and hard boron

carbide (B4C) coatings with high adhesion and high crystallinity are formed on mirror polished SUS304 substrates.

Keywords: plasma spraying, electromagnetically accelerated plasma, pulsed powder injection, boron carbide.

INTRODUCTION

Boron carbide (B4C) is an attractive material as a super hard coating under high temperature and high heat flux

(ref. 1). Sintered bulk B4C is one of the hardest materials, which is only surpassed by diamond and cubic boron

nitride (ref. 2). It also has a low density and a good resistance to chemical agent. Several methods have been

investigated to synthesize the boron carbide coating such as radio-frequency plasma chemical vapor deposition

(CVD) (ref. 3), microwave plasma CVD (ref. 4), laser assisted CVD (ref. 5), plasma jet CVD (ref. 6), sputtering (ref.

7), gas conversion technique (ref. 8) and plasma spraying (ref. 9). In the case of using the B4C coating under high

temperature and heat flux, it seems reasonable to use thick coatings (-100_tm or thicker) to protect the base substrate

materials from the ambient gas. In order to make such a thick coating, it is considered that the plasma spraying

method is suitable. However, the density and the adhesion of coatings formed using the plasma spraying are

generally low because of the low speed powder jet. The substrate coated such a sparse coating is damaged by the

ambient gases passing through a pore and the coating is easily detached. The density of the coatings formed using

high velocity oxygen-fuel flame spraying (HVOF) and Detonation gun (D-gun) spraying are high in the thermal

spraying because high-speed powder jet is generated (ref. 10). However, these spraying methods have a limit to heat

source powder to high temperature. This means that it is difficult to use these methods for coating formation of high

temperature melting materials such as boron carbide. Thus, the development of the spraying method having ability to

generate high-speed powder jet keeping high temperature is important for dense and thick coating formation of boron

carbide.

Dense and high adhesive coatings of tungsten carbide, tantalum carbide and zirconium boride are produced

using electrothermally exploded powder spraying (ELTEPS) method proposed by Tamura et al (ref. 11). The velocity

of the fastest particles jet accelerated by the electrothermal explosion is estimated to be 900m/sec. However, this

method is only applicable to conductive materials because the source materials are heated and accelerated by

applying pulsed large electric current. Railgun plasma spraying method (ref. 12, 13) has similarity to the ELTEPS

method in terms of using pulsed large electric current to heat source materials. Source conductive materials setting

between two parallel electrodes are accelerated by electromagnetic force in this method. The large current also

generates arc plasma that heats the source materials simultaneously. The velocity of the fastest source materials in a

plasma state is 2km/sec. Dense coatings of tungsten carbide, tungsten and TiA1 with high adhesion are produced.

Tahara et al. reported coatings formation of insulators such as A1203 and ZrO2 using magnet-plasma-dynamic (MPD)

NASA/CP--2001-210948 632 August 6, 2001
arcjetgenerator (ref.14).Inthismethod, insulatingsourcematerialsplacednearcathodeareablatedbyconcentrated

plasmajetandsprayedonthesubstrates.

Wehavebeendevelopinganewtypeofelectromagnetically accelerated plasmaspraying(EMAPS)system

havingcoaxial cylindrical eectrodes. Thepurposeofdevelopment istoconduct intenseaccelerating andheatingof

powdersbylargecurrent arcplasmawithoutlimitationinelectricconductivity ofpowders. Forthispurpose, alarge

flowratepulsedpowderinjectorisalsodeveloped. Thispaperpresentsbasiccharacteristics ofthiscoatingmethod

includingtheoperationofthepowderinjectorandthebehaviorfthearcplasmaandpreliminaryresultofboron

carbide(B4C)spraycoatingsonstainlessteel(SUS304) substrates.

EXPERIMENTAL

Illustrations oftheplasmasprayingprocessinatypicalexperimental conditionareshowninfigure1.(a)Mirror

polishedsubstrates (SUS304, 10x10mm2)aresetattheexitoftheelectrode. Sourcepowdersaresetatthepowder

injectorandtheArgasispressurized totheinjector.Thecondenser bank(3mF)ischargedupto 5.9kV.(b)B4C

powder(60_tm, 0.1-1.0g)with the high-pressure Ar gas (0.7MPa) are injected into the inner electrode by the pulsed

powder injector. (c) The pulsed large electric current is applied to the electrode from the condenser bank while the

injected powders and the gas are filled in the electrode. The pulsed current is concentrated to metal lead wires (the

plasma starting point) at first. Next, the wire is exploded to establish arc plasma. The large current generates large

magnetic field at the same time. (d) The plasma is accelerated and compressed by the electromagnetic force. The

powders are accelerated and heated by the accelerated plasma. (e) The powders deposit onto the substrates.

The advantage of this method is to be able to conduct intense accelerating and heating of powders by large

current arc plasma without limitation in electric conductivity of powders. The valve of powder injector is considered

to be one of the key devices of this method, because the response and reproducibility of injection must be high

enough to set precise delay time from the powder injection to the initiation of the large current discharge, which is

considered to be an important factor to control characteristics of the coating. It is however impossible to employ a

conventional electromagnetic valve widely utilized for a fuel injector because of its small flow rate and insufficient

reproducibility of the response. For this reason, a simple method for powder injection is newly developed as shown

in figure 2, where (a) the powder is initially set in a pressurized gas reservoir whose nozzle is closed by a diaphragm

made of thin resinous film, (b) at a

given time, the fine metallic lead

wire located near the diaphragm is

exploded by pulsed current of about

lkA, (c) the wire explosion causes

instantaneous rupture of the

diaphragm and subsequent injection

of the powder with pressurized gas

into the space inside the cylindrical

electrodes. With this method, the

pulsed powder injection with large

flow rate and high reproducibility is

expected to be achieved.

RESULTS AND DISCUSSION

The optimal condition of the

pulsed powder injector is

investigated. The parameters

selected in this study are the

diameter of the lead wire, Ar gas

pressure, the thickness of the film

and the capacitance. The space

inside the outer electrode is

evacuated in this optimization and

following B4C coating formation.

powder

injector i_[j_ (a) ,,o ,, der i---g _ ....... (a) 0 ........................

• -4 ". • ..

...................... _.................. _ \" .........

_., lcad wire L..... - .... ! lead wae --a'F:L--J.'_s_tch,

(b) i_ ___- _i:' .... II the ilmer eleclrode II

•. • .. _ _

(c) , wn-e ,_ :..: .::. T

I.......... 7

" the inner electrode I

[.................. J

(a)

;-.2_5_:.V.V.-...V .......

. :'.'.V'_;.'.'.'.'.'.'.'.'.'." ....... -'- (c) ,-............. ,

,

I.......... ] , t oie,:t,o :

i t

Figure 1. Illustrations of the

spraying process.

Figure 2. Illustrations of the

large flow rate pulsed powder

injector and their injection

process.

NASA/CP--2001-210948 633 August 6, 2001
Thetypicalwaveformof thepressurechangein theinjectorin the

optimalconditionis shownin figure3. Theamountof theinjected

powderis1.5g.Figure3(b)isamagnification ofthefigure3(a),where

thebeginningofthepressurechangeisclearlyseen.Wedefinethetime

betweenthetriggertimeandthetimeofstartofthepressurechangeas

theactingtimeasshowninfigure3 (b).Fiveshotsweredoneatthe

sameconditiontoinvestigatethereproducibility oftheactingtime.The

actingtimeanditsreproducibility inthisoptimal conditionare320_tsec

and±10_tsec respectively. Theincreaseofthepressurechange, inother

word,theincreaseoftheArgasjetvelocityisseenat1800_tsec (white

arrowinfigure3(b)).Thissuggests thealmostpowderswhichwasthe

obstructionfthegasjethavebeeninjectedtotheelectrodefromthe

powderinjectorat thistime.Thus,theflowrateof thepowderis

estimatedtobe1g/msec.

Typicalwaveformof measuredelectrical currentandthepulsed

plasmaemissionsdetectedbythethreeopticalsensorsduringthe

sprayingprocessareshowninfigure4wherethedelaytimeis3.75msec.

Inordertoobtainthemaximumcurrentof 100kA,theappliedvoltage

ofthecondenser bankisadjustedto5.9kV.Thedischargedurationis

approximately 300_tsec. Theaveragespeedoftheplasmafromthepoint

A(10cm)topointC(36cm:theexitoftheelectrode) is2.2km/sec and

the estimatedmaximumelectromagnetic pressureis 1MPa.The

measuredplasmavoltageisrangedfrom100to150Vandtheestimated

appliedenergyisrangedfrom1.0to 1.5kJ.Whenthedelaytimeis

changedfrom3to20msec, thespeedoftheaccelerated plasmaisalso

changedrangingfrom1.5to3.0x103m/sec. Thehighreproducibility of

thecoatingformationis expectedbecausethereproducibilityof the

actingtimeofthepowderinjectorisapproximately 1/10atthedischarge

time.

BoroncarbidecoatingisformedonamirrorpolishedSUS304

substrate. Thegrowthrateisapproximately 1_tm/shot whentheamount

of thesupplyingpowdersis 0.1g.Thegrowthrateincreasesto

approximately 2_tm/shot asthesupplyingpowdersareincreasedto

0.5g.Growthrateissaturatedatthisamount ofthesupplyingpowders.

Ascanningelectronmicroscopic(SEM)imageofanas-sprayed boron

carbidecoatingisshownifigure5.Boroncarbidesplashingisclearly

seen. It showsthatheboroncarbidepowdersmeltduringthespraying

process. Thethicknessof thecoatingincreasesalmostlinearlyby

furthersprayingontothecoating.Figure6 showsa crosssectional

SEMimageoftheboroncarbidecoating,whichissynthesized by9

shots.Thethicknessofthecoatingisapproximately 20_tm. Thisfigure

showsthatheporositiesinthecoatingarelowandtheuniformityis

veryhigh.Novacancyis observedbetweenthecoatingandthe

substratethatindicatesthehighadhesionfthecoating.

l° t (a)

"E"

0 I .... '_ ,_ ..... i ...........

a_ 0 0.01 0.02 0.03

Time [sec]

1o ,_actingW_,,,

time

co 8

6

4

.E

'_ 0

0 0.001 0.002 0.003 0.004 0.005

Time [sec]

Figure 3. (a) Pressure change in the

injector during powders injection

process. (b) Initial stage of the

pressure change.

120

_" lOO

== 80

60

in

_3

40

20

o

0.00

Figure 4. Typical

discharge current

plasma emissions.

A: 10cm

B: 20cm

C: 36cm

(fromthe plasma

0.05 0.10 0.15 0.20 0.25 0.30

Time [msec]

waveform of

and measured

An X-ray diffraction pattern shows that the original crystalline structure of the powder is preserved in the

coating as shown in figure 7. Composition ratio of boron and carbon of the powder (B/C-4) is also preserved in the

coating from electron probe micro analysis (EPMA). The Vickers hardness of the coating obtained by ten tests

ranged from 2200 to 3200kgf/cm 2 (2600kgf/cm 2 average), when the load and loading time are 50gf and 10sec. These

values are smaller than sintered B4C which Vickers hardness is ranged from 3100 to 3700kgf/cm 2 (3450kgf/cm 2

average) at the same loading condition. These smaller values of the coating result from the mixing ofcupper from the

electrode and/or insufficient electromagnetic force.

NASA/CP--2001-210948 634 August 6, 2001
• ..::?:.

p • %°

Figure 5. SEM image of surface of the BoC

coating.

.'. -,,.-..,'.'." "...'."",:,:.'..'.0o_'.'_::;-:.'-_ .'-"_'.

__,,'.'¢1_-_o:"_°° "°'" " *_'* " • "° " ° "*'•" "°" " "_' " '_" ' ....._°°'_°%°%q;

....,...._. ....... _ ......... .,,. :_ ..- _:_¢

gs!?i :i iii::iii:.;i :: ii)ii::siii;i iiisiii ::i iiiiiiiii.

.% * .. • • ......................°.............°.....................-......1.., •:;_0_3_ .4:.:.:-:.:-:.:.:.'.'.:.:.:.:.:.:.:.'.'.:.:.:.:.:.:.: .:.:.:.:.:.:.:.:.:.:-:.:.:-:-:.:

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::: ;;f,_:o:.'.:.

..... ......................_...:._._._.:......._._._._._...._._._..._...:.:._..._._._._._._._._._._.....

Figure 6. Cross-sectional SEM image of the

BoC coating on SUS304 substrate after 9

shots.

CONCLUSION

Electromagnetically accelerated plasma spraying (EMAPS)

system using a coaxial cylindrical electrode is developed. A large

flow rate pulsed powder injector is also developed and combined to

the spraying system to select source powder materials without no

limitation such as conductivity. The measured maximum discharge

current is 100kA. The discharge duration is approximately 300_tsec.

The speed of the accelerated plasma is ranged from 1.5 to

3.0x 103naJsec and the estimated maximum electromagnetic pressure

is approximately 1MPa. The high reproducibility of the coating

formation is achieved from the high reproducible acting time and the

large flow rate of the powder injector. Dense and hard boron carbide

(B4C) coatings with high adhesion and high crystallinity are formed

on mirror polished SUS304 substrates.

REFERENCES

=.

• B.C _'

15 20 25 30 35 40

2theta (deg)

Figure 7. X-ray diffraction pattern of

the boron carbide coating on the

SUS304 substrate.

1. Buzhinskij, O. J., Semenets, Yu. M., Fusion Engineering and Design, 45, 1999, 343

2. Riedel, R., Adv. Mater. 6, 1994, 549.

3. Lee, S., Mazurowski, J., Ramseyer, G., et al., J. Appl. Phys. 72, 1992, 4925.

4. Lee, K.W., Harris, S., Diamond Relat. Mater. 7, 1998, 1539.

5. Oliveira, J. C., Oliveira, M. N., Conde, O., Surf. Coat. Technol. 80, 1996, 100.

6. Postel, O. B., Heberlein, J. V. R., Diamond Relat. Mater. 8, 1999, 1878.

7. Jankowski, A. F., Schrawyer, L. R., Wall, M. A., J. Appl. Phys. 68, 1990, 5162.

8. Gotoh, Y., Yamaki, T., Ando, T., et al., J. Nucl. Mater. 196-198, 1992, 708.

9. Bolt, H., Araki, M., Linke, J., et al., J. Nucl. Mater. 233-237, 1996, 809.

10. The Science and Engineering of Thermal Spray Coatings, Pawlowski, L., John Wiley & Sons 1995.

11. Tamura, H., et al., J. Thermal Spray Technol. 7, 1998, 1472.

12. Usuba, S., Fujiwara, S., Imura, T., T. IEEE Japan, Vol. l14-A, 1994, 603, (in Japanese).

13. Lund, C. J., Resad, C., Eliezer, Z., Wear 119, 1987, 251.

14. Shibata, T., Tahara, H., Yasui, T., et al., Proc. 26th Int. Electric Propulsion Conf., 1999.

NASA/CP--2001-210948 635 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

DEPOSITION OF BORON CARBIDE THIN FILMS AND THEIR HARDNESS

Fumio Kokai* and Manabu Taniwaki

FCT Project/JFCC, c/o AIST, Tsukuba Central 5, 1-1-1 Higashi, Tsukuba, Ibaraki 305-8565, Japan

Masatou Ishihara and Yoshinori Koga

Research Center for Advanced Carbon Materials, AIST, 1-1-1 Higashi, Tsukuba, lbaraki 305-8565, Japan

Keywords: Boron carbide, Thin films, Pulsed laser deposition, Hardness

Covalently bonded solids based on boron, carbon, or nitrogen are the hardest materials, and B4C is the third

hardest after diamond and cubic boron nitride. In addition to its hardness, B4C has a high melting point,

semiconductivity, and high resistance to chemical reagents. These properties make thin films of boron carbide (BC)

an attractive candidate for room- and high-temperature applications in the fields of coating materials and electronics.

Among numerous preparation techniques, pulsed laser deposition (PLD) has been successfully used to form thin

films of various materials. In this method, a laser beam incident on a solid target induces an explosive evaporation of

atomic and molecular species. The highly energetic species expand in vacuum and are deposited on a suitable

substrate. As to the deposition of BC thin films, few PLD studies have been reported. In this work, we report the

deposition of amorphous BC films by PLD using a fourth harmonic beam from a Nd:YAG laser. We discuss the

deposition of the BC films depending on laser fluence and the relationship between energetic ablated species and

nanoindentation hardness of the deposited films. The values of hardness of tetrahedral amorphous carbon (ta-C) films

are also measured and compared.

In a stainless steel chamber with a base pressure of <10 -_ Torr, a commercially available sintered B4C target

containing 5% oxygen was mounted and rotated at 10 rev./min. The laser beam (266 nm, 4-5 ns pulse width) from a

Q-switched Nd:YAG laser, operated at a repetition rate of 10 Hz was focused on the target through a lens. The spot

diameter of the laser beam was 0.8-1.5 mm, and the laser fuence was 1-3 J/cm 2. The films were deposited on Si (100)

substrate at room temperature. The substrates were set parallel to and 30-50 mm away from the B4C target.

Deposition times were 30-60 min. After deposition, the films were removed from the chamber and their thicknesses

were measured with a depth profiler. The deposition rate obtained at the center of the substrate ranged from 2 to 12

nm/min. The morphology of the surface and cross section of the films was examined using scanning electron

microscopy (SEM). Particulates in the deposited films were observed and counted. The films were also characterized

with x-ray photoelectron spectroscopy (XPS). The hardness values of the films were determined by nanoindentation

using a continuous stiffness measurement technique. The continuous stiffness measurement allows the hardness to be

measured as a function of indentation depth. For comparison, ta-C films were also formed by PLD using a graphite

target at laser fluences of 1.3-3.5 J/cm 2.

The surface SEM images of deposited films showed the presence of round and irregularly shaped particulates

embedded in the smooth films. Higher magnification images of the films showed the particulates most likely

originated from melt droplets from the B4C target surface. We found that the particulates size ranged from 0.25 to 5

gm and that their average size was about 1 gm. We also found that with an increase in fluence, the number of

particulates decreased significantly and the number of round particulates increased relative to that of irregularly

shaped ones. Although particulates, an undesirable product, exist in the film, the cross section image of the film

showed that the deposited film itself has a dense structure without peeling off from the Si substrate.

XPS analysis showed boron, carbon, and oxygen atoms in the surfaces of BC films deposited at different

fluences. The atomic compositions of the films were determined by integrating the areas of the respective Cls, B ls,

and Ols XPS spectra after correcting for their sensitivity factors. With increasing fluence, the B/C ratio increases

from 1.8 to 3.2 and the O/C ratio decreases from 0.7 to 0.2. Although a fairly large change is observed in the B/C

ratio, the B/C ratios of all the films are smaller than that of the bulk target (3.6).

To compare the type of chemical bonds present in the films, the Bls spectra were deconvoluted into three

components through reference to previous studies. The main component with a peak at 188.6 eV is assigned to B

atoms bonded to B. The other two components at 190.6 and 193.0 eV are assigned to B atoms bonded to C and to O.

The Cls spectra were also deconvoluted into three components. The main component at 282.9 eV is assigned to C

atoms bonded to B. The other two components are assigned to C atoms bonded to C and to O. At larger fluences the

ratios of the two main components at 188.6 and 282.9 eV increased significantly, which is consistent with the

increase in the B/C atomic ratio with increasing fluence. The Raman and FTIR spectra of the deposited BC films

showed only broad bands with peaks at around 1100-1200 cm -1, indicating an amorphous structure. The dominating

factor in the formation of BC films, which are dense and adherent to the substrate, appears to be the arrival of

NASA/CP--2001-210948 636 August 6, 2001
energeticBandCspeciesfollowedbytheircondensation onthesubstrate. AsinPLDofothermaterialsthroughthe

useofanultraviolet laser,themajorarrivingspeciesarethoughttoconsistof atoms,ions,andsmallmolecular

speciescomposedof mainlyBandC.In ourprevioustime-resolved emissionmeasurement intheablationplume

generated fromaB4Ctarget,B,C,andtheirionswerefoundtoexpandwithvelocitiesof1-30km/s.TheBdeficient

composition ofthedepositedfilms,comparedtothetargetone,mayhavebeenduetothechemical bondstabilityin

thefilm.Thebondstrengthsare607,448, and297kJmo1-1 forC-C,B-C,andB-B,respectively.

ThehardnessofBCfilmswiththicknesses of- 500nmwasmeasured inthepartsofthefilmswhicharefreeof

particulates. ThevaluesofhardnessareplottedinFigure1asafunctionof laserfluence.TheYoung'smodulus

valuesofthefilmsarealsoplotted.Theeffectoflaserfluencecanbeobservedfromtheresultsobtained. Although

thehardnessvalueisabout14GPaforthelowestdegreeof laserfluence, increasingthefluenceleadstoarapid

increaseinfilmhardnessupto32GPa.Themodulusvariesfrom175to313GPainparallel tothehardness, whichis

a general trendobservedin hardandelasticceramicsfilms.Themeasuredhardnessvaluesof 14-32GPaare

comparabletothoseforBCfilmsdepositedbyDCmagnetronsputtering(10-35GPa),RFmagnetronsputtering

(19.5-30GPa),andion-beamsputtering(31GPa),andapproachthoseforbulkB4C(3000-4000inVickershardness).

Forta-Cfilms,thehighervaluesofhardnessrangingfrom24to65GPawereobtained. Similartothedepositionf

BCfilms,thehardness valuesofta-Cfilmsincreasedaslargerfluenceswereused.

Webelievethathedominant factorintheobservedincreaseinBCfilmhardnessistheformationofdenser

filmsbythearrivalofmoreenergeticionsandatomsgenerated bylaserirradiationathigherfluences. It isgenerally

acceptedthatahighlaserfluenceplaysacrucialroleinthedepositionfdenseamorphous carbonfilmswithalarge

numberof sp3hybridizedcarbonatoms.In theabove-mentioned sputteringtechniques, BCfilmswithhigher

hardnesswereobtainedundersuitablesubstrate-bias conditions, indicatingtheroleofionsinthefilmdeposition. I

ourdepositedfilms,atomiccompositionandbondingstatecomposedofB andC atomsignificantlyvaried

dependingonthelaserfluence.Furtherstudiesarerequiredtobetterunderstand therelationshipbetweenbonding

statesinthefilmsandtheirhardness.

ThisworkwassupportedbytheFrontierCarbonTechnology Project, whichwasconsignedtoJFCCbyNEDO.

a.

o

e-

"I"

40 600

30-

20-

10-

0

0

I I I

• • O

O "--_

[] []

I I I 0

1 2 3 4

Fluence (J/cm 2)

5OO

400

m

300

0

E

200

0

lOO

Figure 1. Variation of hardness and Young's modulus of BC films as a function of laser fuence.

*Corresponding author:

E-mail: k_-kai-f@aist.go.jl', Fax: +81-298-61-4474, and Tel: +81-298-61-4573

NASA/CP--2001-210948 637 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

THE SPATIAL-RELATIVE JOINT DENSITY OF STATE

FOR AMORPHOUS CARBIDES

Z.Q.Ma

Functional Materials Laboratory, Department of Physics, Xinjiang University,

Urumqi 830046, P.R.China.

Extended Abstract

Naturally, the carbides of pure in composition and complete in atomic lattice are not present whether in

crystalline or in amorphous states. Most of natural and synthetic diamond contains nitrogen, boron and other rare

impurities as "structural defects", even if amorphous carbon (a-C) films manufactured through PVD (Physical Vapor

Deposition) techniques, such as filtered cathodic vacuum arc technique, laser ablation and ion beam deposition.

However, in the practical application of industrial fields, we usually treat the diamond and graphite crystals as

complete lattice sites, while other allotropes of carbon with sp 3, sp 2 and sp 1 configuration mixing in solids are known

as amorphous structure. Their physical and chemical properties are somewhere different from diamond partner.

Initially, the interest in amorphous carbon thin films was as a mechanical material that showed 'diamond-like'

properties [1,2]. Currently, there is renewed interest in a-C, as well as its doping counterparts, as electronic or

opto-electronic materials. The investigation showed that these properties were correlative to the distribution of

electronic states, as well as energy bands in solids, which was determined by optical absorption spectrum.

The joint density of electronic states (JDOS) in amorphous semiconductors has been previously studied by Tauc

et al. [3], and a simple definition was given as:

J (h(o ) = [Nc(E)N_(E - hgo)dE (1)

J

N c (E) and N v(E) representing the conduction band and valence band one-electron density of states (DOS)

functions, respectively.

Later, O'Leary [4] suggested an elementary formalism for the optical absorption spectrum which does account

for the spatial correlations between the conduction band and valence band potential fluctuations and does allow one

to relate the distribution of electronic states with this absorption spectrum. The obvious feature is local JDOS

function in amorphous solids

j zoc (hco) = I N_°_ (E)N_°_ (E - boo)dE (2)

N_ °_ (E)

total JDOS in amorphous material is an average of local JDOS, i.e.,

J (hco ) = (j;oc (h(o ))

In amorphous solids, one-electron potential profiles, V C(R) and

band are not constant, and the local density of states can be expressed as:

and N_°_(E) denoting the local conduction band and valence band DOS functions, respectively. The

(3)

Vv (R), in conduction band and valence

NASA/CP--2001-210948 638 August 6, 2001
,_2-m *3/2

c 4E - Vc (R) (4) N_°_(E)- _ 2h3

-v/Em *3/2

N_°_(E) - 2 v3 _/vv(R)- E (5)

Jr h

Where mc and my are the effective masses of carriers. Under a proper approximation, the JDOS of amorphous

semiconductor is obtained to be:

[ [(ha) - Eg o )2 + o_2,

J(h(o) ----> _ [ h(o- E go 2 ]

?p[ l,

ho)> E

go

h(o _< Eg o

(6)

Where Eg o is the average width of band gap, cy the standard error. The optical absorption coefficient is derived as:

a(hco) = D 2(ho)). J(hco) (7)

Where D(ho)) denotes the optical transition matrix element:

( fl_[._ or e 2 / 1/2

D (h_0) = - 7- - 1_ _ c_(k)[

n/_ o (1)

To a fixed system and known optical transition,

(8)

D(ho)) ---->1/ho), thus, the relationship between absorption

coefficient (c 0 and indirect band gap (Eg) for amorphous semiconductor can be derived as:

( a h co ) 1/2 = B ( h co - E _ ) (9)

For hO)_> Eg, B is a constant. However, owing to the disorder characteristic of amorphous semiconductors

introduces distributions of tail states which encroach into the otherwise empty gap region, the optical transitions

involving the tail states are responsible for the tail in the optical absorption spectrum observed for photon energies

below the gap (hO) _< Eg ).

For the amorphous carbides, such as diamond-like amorphous carbon (DLC) thin films with the electronic state

configuration of sp3+sp2+sp I or surface doping modification of diamond crystals by ions, the spatial-resolution

optical absorption and fluorescence spectra show that the JDOS in amorphous carbon is of spatial dependence in

long distance [5]. The joint density of states decides the photoelectron transition and the electric conductivity of a-C

solid.

NASA/CP--2001-210948 639 August 6, 2001
UsingTaucapproximation andanewdefinitionofbandgap,therelativevalueofopticalabsorptioncoefficient

andbandgaphavebeenderived. Fromtheopticaltopographies intransmission mode,thenon-uniformity ofatomic

configurationadelectronicstatesirevealedinanytypeof amorphous carboncomplex. Theindicationhasbeen

appliedtothesyntheticdiamondirradiatedwithlowenergyionatR.TandLN2ambient.

The"hoping"conductionof carrierandtheselectiveabsorptionof localstatesin carbonallotropesare

interpretedthroughthevariationofelectronicstates. Thelikelyenergybandstructureofamorphous carbonsolidsin

short-distance orderwithin'nanometer' scopehasalsobeendiscussedasfigure1.

In ourrecentinvestigation, the"localfeatures"of opticaldensityof thoseimpuritycontainedand/ or

irradiation damaged solids were experimentally resolved with a special optical topographical transformation and a

surface scanning method applied to tetrahedral covalent bond of carbon. The difference of relative optical gap over

all of bulk is associated with "local structure" of atomic bonding, which causes optical transition between localized

states, and some "color centers" are formed in dispersed sites. The photoluminescence and cathodeluminescence

emitted from modified "nano-volume" of insulator (SiO2, A1203, ZrO, etc.) is ascribed to localized band gap. The

conductivity of amorphous structure being larger than its host in crystal is taken as the decrease of tunneling

potential height, resulting in the increase of efficient carrier transport and the narrow of forbidden bands.

The impurity and defect in solids can induce a spatial distribution of electronic states as well as band gap (as

figure), so as to lead to the variation of optical and electrical properties.

4.5

4.0

A

3.5

I.IJ 3.0

O.

2.5

c 2.0

1.5

'_ 1.o

I.kl 0.5

0.0

Fig. 1

E 2(r) E 3(r)

I i I i I i I i I i I

0 2 4 6 8 10

Space site in length (Angstrom)

The distribution of optical gap for amorphous carbide.

NASA/CP--2001-210948 640 August 6, 2001
REFERENCES

1. Aisenberg, S.; and Chabot, R., J. Appl. Phys. 42 (1971) 2953.

2. Tsai, H.; and Bogy, D.B., J. Vac. Sci. Technol. A 5 (1987) 3287.

3. Tauc, J.; Crigorovici, R.; and Vancu, A., Phys. Star. Sol. 15 (1966) 627.

4. O'Leary, S.K.; Zukotynski, S.; and Perz, J.M., Phys. Rev. B 51 (1995) 7795.

5. Ma, Z.Q.; Liu, B.X.; and Naramoto, H., et al., Vacuum 55 (1999) 207.

Keywords: Joint density of electronic state, Amorphous semiconductor, Carbides, Optical absorption

NASA/CP--2001-210948 641 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ELECTRICAL AND OPTICAL PROPERTIES OF TERNARY BCN THIN FILMS

Guanghua Chen Shengjun Zhang Xuemei Song

College of Materials Science and engineering, Beijing Polytechnic University, Beijing 100022, China

Jinshun Yue

Department of Physics, Lanzhou University, Lanzhou 730000, China

ABSTRACT

The close similarity of the atomic arrangement between graphite and hexagonal boron nitride (h-BN) indicates

the possibility of synthesizing BCN compound with probable properties variable between the two materials. The

ternary boron carbonitride (BCN) thin films were deposited by radio frequency (rf) reactive sputtering method. The

structure and composition of the films were determined by X-ray diffraction (XRD), Fourier transform infrared

(FTIR) and X-ray photoelectron spectroscopy (XPS), and the optical properties was measured by Ultraviolet-visible

spectoscopy. The bandgap measured varies from 0.27eV to 0.87eV with carbon source (methane) in the precursor

changing from 20% to 0. The electrical and optical properties were investigated by measuring the temperature

dependence of the conductivity. It is suggested from the experiments that there are two possible states that the

carbon may exist, namely as the dopant in h-BN when the carbon content in the films is lower, and the BCN alloy

when the carbon in the precursor is high.

Keywords: BCN films; Optical bandgap; Electronic properties

INTRODUCTION

The attraction of seaching for materials with hardness close to or higher than that of diamond, the hardest

material in the world, and the interest in the investigation of novel functional devices materials, open the door to the

study of the composite of the first elements in group III, IV and V, i.e. B, C and N. It is well known that diamond is

the hardest material in the world and cBN the second. Although theoretical prediction indicates that compound of

carbon and nitrogen, [_C3N4, may be even harder than diamond, there are still disagreements on this material (ref. 1

to 3). Firstly, most experiments on the attempt of preparing C3N4 obtained no stochiometric composition, the claims

of stochiometric materials are absent of confirmly evidence. Secondly, it is much difficult to acquire the crystalline

superhard phase. Thirdly, the experimental hardness was not such high. In addition, although physical properties

such as valumetric lattice energy, bulk modulus and shear modulus, ionicity (the harder the material, the higher the

covalency fraction in the bond), melting point and band gap are helpful to predict hardness (refs. 3 to 4), the

mechanical hardness appears to be in close relation to shear moduli. So the bulk moduli based prediction of C3N4

may be not such accuracy (ref. 3).

The ternary of boron carbonitride (BCN) is both scientifically and technologically interesting to material

researchers as cBN is isoelectronic material with diamond, and hBN is the isoelectronic with graphite, as well as the

atomic arrangement of cBN and diamond, hBN and graphite being similar. There are two modification of BCN, the

cubic phase (cBCN) and the heragonal one (hBCN). It is suggested that the cubic phase is supposed to be superhard,

with hardness similar to or over cBN. It's a favorable prospective superhard material as diamond can be alloyed with

NASA/CP--2001-210948 642 August 6, 2001
ferrouselementsinhightemperature leadingtoitsdisabilityincuttingtheselements, andthecompressive stressin

cBNfilmsis muchtoolargecausingitseasilyexfoliationfromthesubstrate. However, fabricationofBCNis

extremelyhard.Ontheotherhand,thehexagonal phasehasmanyinterestingproperties. Asiswellknown,graphite

issemimetal. Thereisaoverlapof0.04eVbetweenthevalencebandandtheconductionband.hBNisainsulator

withawidebandgapof largerthan3.8eV.hBCNissuggestedbyreplacingcarbonbyboronorbynitrogen,

respectively. Thegoodpropertiesof hBCNincludesemiconductivity, goodemissioncharacteristic (ref.5)

photoluminescence, andelectroluminescence, thermoelectricity, rectification, intercalation, resistant to oxides,

whichleadingtotheprobableapplications oflight-emittingdevices, thermoelectric conversiondevices, negative

electrodeof secondaryLi batteries(ref.6)andwear-resistant coating.ThecompositionfB-C-Nsystemmay

writtenasBCxN,of wherex maybe0, 9,2, 2.5,3, 4, 7.(ref.7)BCNcompoundwasfirstsynthesized by

Badzianetal (ref.8),usingCVDprocesswithBC13, CC14, N2andH2asthestartingmaterials. Kouvetakis etal(ref.6)

synthesized BC2NusingCVDtechniquewithBC13andCH3CNasthestartingmaterial. Ontheotherhand,sincethe

reportofcarbonnanotubein1991(ref.9),BCNattractedmuchinterest andmanyattemptshavebeenmadetothe

goodpropertiesofBCNnanotubes. TheBCNnanotubesareconsideredpriorto thecarbononesasbandgap

independent ofdiameter, chiralityandthenumberofwalls,whilecarbonnanotubecanbemetalorsemiconductor

bycontrollingofthediameter andchirality. Liuetal(ref.10)predictedthreepossiblestructureofBC2N,thatistype

I whichismetal,typeII withbandgapof1.6eVandtypeIII withbandgapof0.5eV.Kawaguchi (ref.11)reviewed

thepreparationmethodsofsomeBCNcompounds.

Weherereporttheelectrical andopticalpropertiesoftheBCNfilmspreparedbyrf sputteringtechnique. The

structureandcompositionf thefilmsweredeterminedbyX-raydiffraction(XRD),Fouriertransforminfrared

(FTIR)andX-rayphotoelectron spectroscopy (XPS),andtheopticalpropertieswasmeasuredbyUltraviolet-visible

spectroscopy.

EXPERIMENTAL

TheBCNfilmsweredepositedbyrf sputteringtechnique(ref.12).ThetargetwasasinteredhBNdiskwitha

diameter of 10cm,thepurityofthetargetis99.99%. Fusedsilicaandsinglecrystallinesi(100)waferwereusedas

thesubstrate. Thesubstrates werecleanedinarigidprocessbeforedeposition. Thesubstrates werecleanedwithan

organicsolventandthenetchedwith5%HFfor5min,thenrinsedwithde-ionizedwater.Thecleanedsubstrates

weredippedinthewater-freedethanol. Beforeputintothedepositionchamber, thewafersweredriedbyN2blowing.

TheworkingaseswereAr andCH4.Thegasflowwascontrolledbythemassflowcontrollers. Theworking

pressurewaskeptat1Pabythevalve.Thesubstratetemperature was450°C,thedensityofrfpoweris3.82W/cm2.

Thegrowthrateofthefilmwastypicallyaboutllnm/min.

RESULTS AND DISCUSSION

The XPS spectra show that there are more than one type of bonding scheme for B, C and N atoms, which are

the B-N, B-C, N-B, N-C bonds (ref. 13). (See Table 1)

XRD results indicate that the diffraction patterns are coincident with that of the reported BC2N. The lattice

constants are indexed to be a--2.4 and c--6.93.

NASA/CP--2001-210948 643 August 6, 2001
Table1.Bonding scheme for B, C and N atoms from XPS results

B1S C1S N1S

Literature (ref. 13) 189.5 284.4 398.7

Experimental 190.2 284.2 398.5

FWHM 3.2 2.8 2.6

Fitting 189.5 190.8 284.2 285.6 398.4 399.4

Assigning B-C B-N C-B+C-C C-N N-B N-C

Electrical properties

The resistivity was measured by the four-probe method with variation of temperature, evaporated A1 electrodes

of 3xlmm 2, with a distance about 2mm between each other. The relations of conductivity and temperature of two

samples, where the C contents are 31% and 38% respectively as determined by XPS, were shown in Figure 1. The

films were deposited at 5% CH 4 partial pressure, while the flow rate of Ar is 20sccm and 15sccm, respectively. The

curve in figure 1 indicates that the conductivity variation of both the films in the measured temperature coincident

with the Arhenlus equation. And at the identical temperature, the conductivity of the films with more C content

1E-5 - _... .....

1E-6 _ "_x"_e"¢-'_" _'_- _ _., b 1F,-7

I_ 1F,-9

Ilim

1F,-1(} J "_'.r,.,,u" a

2.4 2.6 2.8 3.0 3.2 3.4

1000/T (K "l)

Figure 1. The relations of conductivity and temperature

(a: C content is 31%; b: C content is 38%)

(curve b) is higher than that with less C content (curve a). The conductivity at room temperature are 1.1x10-8(_2cm) -1

and 3.5x10-6(_2cm) -1, respectively. The activation energy calculated is 0.7eV for curve b and 0.8eV for curve a. Yao

et al (ref. 14) reported that when C content up to 50% in the BCN films, the resistivity at room temperature is lx10 -2,

and the activation energy is 0.11eV. It seems that the conductivity of the films increases with the C content. It is

suggested that when the C content is very low, the structure of hBN is not destructed, carbon in the films is dopped

to change the Fermi level of hBN, i.e. p-type when C substitute N and n-type when C substitute B. When the C

content is high enough to change the structure of hBN, then BCxN films will be formed. As the bandgap of hBN is

larger than 3.8eV, that of BC2N is 1.6 eV, the latter is semiconductive. That is to say, carbon in the films will change

the bandgap of the films, on the other words, the bandgap of the BCN films is adjusted by C content.

The Hall effect measurement shows that the films shown in figure 1 is p-type. The hole motivity and hole

concentration at different temperature was done on films deposited on fused silica substrate. It is shown in figure 2

that the hole concentration rises with temperature while the mobility decreases with temperature in large.

NASA/CP--2001-210948 644 August 6, 2001
g-,

'E

1020

1019

10 is

1017 I I I I I

0 1 2 3 4

100/T (K "1)

101

"E

10 o *

m

10 "1

Figure 2. The dependence of hole concentration and hole motivity on temperature

Optical property

The films were deposited at different CH4 partial pressure. The CH4 partial pressure of Sample a, b, c, d and e

is 0, 5%, 10%, 15% and 20%. The FTIR spectra of films were shown in figure 3. When the CH4 partial pressure is

larger than 10%, there appears a peak at ll00cm -1, accompany with the two peaks of hBN, which are the B-N-B

bending mode at 780cm -1 and B-N stretching mode at 1376cm -1. The 1100cm -1 peak maybe assigned to the vibration

of C-B bond is in the range of 1250-1380 cm -1, which is overlap to the B-N stretching vibration.

The absorption coefficient ct vs. Photo energy was plotted in figure 4. The bandgap can be acquired from the

curve c, d, e is 0.87, 0.70 and 0.27eV. It is clearly that the bandgap of the films decreases with the increase of carbon

.=

content.

2.5

2.0

1.5

1.0

0.5

0.0

500 1000 1500 2000 2500 3000 3500 4000

Wavenumber (cm "1)

Figure 3. The FTIR spectra of films deposited at different CH4 partial pressure

CONCLUSION

We deposited the films with various carbon contents. The electrical measurement shows that conductivity of

the BCN films increases with the carbon contents, and the BCN films is semiconductive. Hall effect measurement

NASA/CP--2001-210948 645 August 6, 2001
indicatesthefilmsareofp-type.Thebandgaps ofthefilmsdecreasewiththeincreaseofthecarboncontents. The

bandgapmeasuredissuggested thathecooperationcarbonintheBNfilmsinfluencesthepropertiesofthefilms

greatly.Whenthecarboncontent islower,thedopedhBNfilmsacquired, whilewhenthecarboncontent becomes

larger,thestructureofhBNwillbesubstituted byBCN,andthebandgapbecomesnarrowsharply.

'E

L.

1.2x10 5

1.0xl0 5

8.0x10 4

6.0x10 4

4.0x10 4

2.0x10 4

0.0

0.0 0.5

/ /b

1.0 1.5 2.0 2.5 3.0 3.5 4.0

Photo Energy (eV)

Figure 4. The relation between photo energy and absorption coefficient

ACKNOWLEDGMENT

This study was supported by the National Natural Science Fund and Fund of Beijing Natural Science

Foundation.

REFERENCES

1. Muhl, S.; Mendez, J. M.: A review of the preparation of carbon nitride films. Diam. Relat. Mater., vol. 8, 1999,

1809.

2. Badzian, A., et al.: Silicon carbonitride, a new hard material and its relation to the confusion about harder than

diamond C3N4. Thin Solid Films, vol. 354, 1999, 148.

3. Teter, D. M.: Computational alchemy: the search for new superhard materials. MRS Bull., vol. 23, 1998, 22.

4. Gilman, J. J.: Chemical and physical "hardness". Mat. Res. Innovat., vol. 1, 1997, 71.

5. Sugino, T.; Hieda, H.: Field emission characteristics of boron carbon nitride films synthesized by

plasma-assisted chemical vapor deposition. Diam. Relat. Mater., vol. 9, 2000, 1233.

6. kouvetakis, J.; Sasaki, T.; Shen, C.; Hagiwara, R.; Lerner, M.; Krishman, K. M.; Barlett, N.. Synth. Met., vol.

34, 1989, 1.

7. Watanabe, M. O., et al.: Electrical properties of BC2N thin films prepared by chemical vapor deposition. J.

Appl. Phys., vol. 78, 1995, 2880.

8. Badzian, A. R., Niemyski, T., Appenheimer, S., Olkusnik, E.. Proc. 3rd Int. Conf. On Chemical Vapor

Deposition, Glaski E A. (Americal Nuclear Society, Hinsdale, IL, 1972), vol. 3, 1972, 747.

9. Iijima, S.: Helical microtubules of graphitic carbon. Nature(London), vol. 354, 1991, 56.

NASA/CP--2001-210948 646 August 6, 2001
10. Liu,A.Y.;Wentzcovitch, R.M.;andCohen, M.L.:Atomicarrangement andelectronicstructureofBC2N.

Phys.Rev.,vol.B39,1989,1760.

11. Kawaguchi, M.:B/C/Nmaterialsbasedonthegraphitenetwork. Adv.Mater., vol.9,1997,615.

12. Yue,J.,etal.:TernaryBCNthinfilmsdepositedbyreactivesputtering. ThinSolidFilms,vol.375,no.1-2,

2000,247.

13. Deshpande, S.V.,etal.:Filament activatedchemical vapordepositionfboroncarbidecoatings. Appl.Phys.

Lett.,vol.65,1994,1757.

14. Yao,B.,etal.:Amorphous B-C-Nsemiconductor. J.Appl.Phys., vol.84,1998,1412.

NASA/CP--2001-210948 647 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Monatomic structures of B, C, N, and O: first-principle study of relative stabilities and

bulk moduli

James C. M. Sung

Department of Materials and Mineral Resources, National. Taipei University of Technology, Taipei 104,

Taiwan, R.O.C.

B. R. Wu and S. L. Lee

Department of Chemistry, National Chung Cheng University, Chia-yi 621, Taiwan, R.O.C.

M. F. Tai

Department of Physics, National Chung Cheng University, Chia-yi 621, Taiwan, R.O.C.

Abstract

Relative stabilities and bulk moduli of monatomic structures with coordination numbers 4 (cubic

diamond or ed), 6 (simple cubic or sc), 8 (body-centered cubic or bee), and 12 (face- centered cubic or fcc) for

elements B, C, N, and 0 were calculated based on the first- principles. The structure for each element may be

stabilized by the geometrical matching between the coordinated positions and the orbital shape. Thus, all

elements favor loosely packed cd structure at low pressures, and they prefer moderately packed sc structure at

high pressures. However, denser phases with bcc and fcc structures may become more stable at much higher

pressures.

B has only three valence electrons that are insufficient to fill up all sp3 orbitals in a cd structure.

Moreover, these three p electrons can only flU half of the three mutually perpendicular bonds in sc structure. As

the result of reduced stabilities of cd and sc structures, relative stabilities of the four structures for B are less

distinctive.

C has four valence electrons that match perfectly with four tetrahedral coordinated atoms. Hence, its cd

structure is greatly stabilized by the extensive overlap of bonded electrons. As the result of this expanded region

of stability, the equilibrium pressure between cd and sc structures is pushed up substantially.

Both N and O have sufficient p electrons to align with octahedrally coordinated atoms. Hence, their sc

structures are greatly stabilized and the equilibrium pressures between cd and sc structures for N and O are

much reduced relative to that for B and C. The electron density maps of ed coordination for N and O suggest

that these monatomic structures may be isotropic superconductors.

Bulk modulus of a symmetric structure may be determined by the average concentration of electrons in

overlapped orbitals. C has the smallest atomic volume of the four elements studied, so it is uniquely capable to

form the cd structure with the highest bulk modulus (447 GPa) of all structures. Although C's p orbitals are not

fully occupied for effective bonding in octahedral coordination, its small atoms still make the bulk modulus of

the se structure the second highest (337 GPa) among various phases of the four elements studied.

Except for diamond, the structures discussed above are all hypothetical. However, they may be

synthesized by the bombardment of collimated beams of atoms focussed to a common center. Moreover, the

structures so formed, even though they may be metastable, can be studied by the laser spectroscope fired in

femtosecond (10 "_ second) bursts. Such femtochemistry has been developed by Dr. Ahmed Zewail, the latest

Nobel laureate.

Key Words: Diamond, Superhard .Materia!, Btqk .\!edu!us, Boron, Nitrogen, Oxygen

NASA/CP--2001-210948 648 August 6, 2001
8

CARBON NANOTUBE PRODUCTION, PROPERTIES, AND

APPLICATIONS

Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

RECENT PROGRESS IN DEVELOPMENT OF LARGE SCALE SYNTllESIS OF

MULTI-VvNLLED CARBON NANOTUBES OF TIlE FRONTIER CARBON PROJECTS

M_too Yumura, Satoshi Ohshima, Hiroki Ago and Kunio Uchida

National Inslitute of Materials and Chemical Research, Tsuka:ba, Japan

ltitoshi Inoue,Toshiki Komatsu

Japan Fine Ceramics Center, Nagoya, 456-0023, _]apa_t

ABSTRACT

We have been developing the mass producuon methods of nmlti walled carbon nanotubes (MWNT) We have :tevetcped

a method to produce several 10 reams of M\_'NT per day by continuous arc method which h_s the higher prodnctivib'

compared ruth the arc disct_rge method usecl from 1996. Furthermore, aiming at proceeding research and development

for the MWNT as the material for industry, we started to develop Ihe new productie_ method to produce severat kg eflhe

IVYO;NT. We have developed a cllemical syntheses method of MWNT by using tow boiling poit_l hydrocarbon as rite row

material and ultra fine metal particle as the catalyst We used ultra fine metal panicle as a catalyst metal and benzene as a

carbon source. Advantages of this method are selective production of MV_T and cost of mass production will be low. It

is quite easy for tlfis method to scale up the MWIVr production. We have developed l!te pilot p"lant uhieh is capable to

produce seveIal kg of MWNT per day.

In 1999. We started to construct the large scale experimental continuous reactor in cooperation with Shown Denko Co.

Ltd. as the Frontier Carbon Technology Project o[MinisttT of International Trading m:d l.}ldtk';lry (MITI), _itd _e

succeeded in producing 200g or-carbon naoolube per hour Now, we are cartTing out rite experiment usi:tg t!tis plant

(Figure 1) to m_e sm'e its possibility for industrial utilization and we are developing several industrial applications of

MWNT.

This work is supported by the Frontier Carbon Teehnclo_' Project of tl:e MIT1.

Keywordsl carbon nanotubes, mass productim:, synthesis

Figaro 1. The pilt_t plant [gr mass prod,action ofM_ ,rNT

NASA/CP--2001-210948 651 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CATALYTICALLY GROWN CARBON NANOTUBES

AND THEIR APPLICATIONS

Chunming Niu, Bob Hoch and Howard Tennent

Hyperion Catalysis International, 38 Smith Place, Cambridge, MA 02138

ABSTRACT

Recently, catalytic growth of carbon nanotubes has attracted considerable attention for its potential in

commercial production. We have been producing catalytically grown carbon nanotubes in large quantity for more

than a decade. Our process is based on chemical conversion of hydrocarbon feedstocks catalyzed by supported

transition metal catalysts. It is a simple petrochemical technology, producing high purity nanotubes of ten

nanometers or less in diameter and tens of microns in length. Due to their structural uniformity, our nanotubes are

ideal template and building blocks for new nanostructured materials. We have prepared highly crystallized SiC

nanofibrils by "topotactic" conversion of carbon nanotubes. The average diameter of SiC nanofibrils is 15 nm. A

novel carbon structure has been engineered from nanotubes through a wet chemical process. This new carbon has an

open frame porous structure, excellent thermal and chemical stability; and more significantly, it is free of

micropores. The later property is important for the development of catalysts with high selectivity and energy storage

devices with high power. The new carbon has been prepared in various physical forms, including particles,

extrudates and membranes (sheet), to meet specific requirement of different applications. Electrochemical capacitors

based on nanotube membrane electrode showed unprecedented power performance. ARC constant of 7 milliseconds

was measured from a ten-cell device, which is approaching the value of electrolytic capacitors. We have prepared

nanotube-polymer nanocomposites by in situ polymerization. Electrical measurement showed that a conducting

network can be formed in polymer matrix at very low nanotube concentration. A resistivity of 650 _-cm was

measured from a nanotube-polystyrene nanocomposite containing 0.1% nanotubes. Note that, to reach the same

conductivity, a considerably higher loading is required if current commercial conductive additives, such as carbon

blacks, carbon fibers and steel fibers, was used.

INTRODUCTION

Since the discovery of arc-grown carbon nanotube in 19911, considerable efforts have been made to the

development of methods for its large-scale production. To date, the arc-grown method has not yielded large

quantities of nanotubes uncontaminated by other forms of carbon. The other method, laser ablation of graphite

target, had only limited success in production of gram-quantity of single-walled nanotubes. Recently, attention has

been attracted to catalytic growth method.

We have been producing catalytically grown carbon nanotubes in large quantity for more than a decade 2'3.

This has made it possible for us to develop various applications, which require bulk quantity of nanotubes. In this

presentation, we will discuss properties of Hyperion carbon nanotubes and several new nanomaterials prepared from

nanotubes.

HYPERION CARBON NANOTUBES

Hyperion carbon nanotubes are produced by catalytic decomposition of hydrocarbons. It is a simple

petrochemical technology, producing high purity carbon nanotubes of ten nanometers or less in diameter and tens of

microns in length. The microstructure of Hyperion nanotube is similar to that reported for buckytubes, i.e.

cylindrical graphitic layers arranged parallel to the tube axis with a distinctive hollow core (insert in Fig. 1).

Hyperion carbon nanotubes are produced in several different forms of aggregates. Fig. 1 shows one type of

aggregates in which nanotubes are arranged semi parallel to each other to form bundles with diameters from one to

two microns and lengths about twenty microns. The tubes can be dispersed in suitable solvents. A TEM micrograph

from dispersed tubes is showed as an insert in Fig. 1. Table 1 compares properties of Hyperion nanotubes to those of

arc-grown multiwalled buckytubes. In several ways, properties of Hyperion nanotubes are unique. These properties,

including high purity, uniform diameter, high concentration of the edge carbons on the surface and loose powder

morphology, are ideal for new nanomaterials synthesis.

NASA/CP--2001-210948 652 August 6, 2001
•"-'-'-'."'-,'-'.".'.'+'._'_-_'_:'.o ".'--_-".'._.'.'.x_. ...-x..._-._

.....°°........° :::... .'1-'_:::::_...'J'_l_. -°, " :_.:::':'@::_'::i

• o%%%°°%°°°°%°.°°o°°°o° o... ,°o • o.°° °o_1, o .o°oO °. o°°

•"::.':'2::_ ".::;'.::':

.:.'.:.:.:.-.'.o°'.'.:o°.'.°.°.:°:°.-.:°:.:°:°°.:.'.'.:°'..:.:.'.' ?.:_ • /..._.;..°:+..v.?...:..°....

.°°°°° ,..°°°o°*°°.°°0.° *o°°°.°.o • .,° °o.°°°°¢,°°.°0. o°o°.,_ , ° ° ° ° °. • _°.*° .o._,o..... °°.*, ...,..............,°.,.......,.,........°.....:.,.....,.....;. _.$_£_ _°:,.._:_,:...._..........,_

......................°............°..°...°.°°.......°°°..-...._.._.-,_-,_°-....._.-°... I.T.:'I_-.._ •._ :i:?i.i,i:i'_:i'i:i:i:i:i:)i'i:i:i:!:i:!'i,'i:i:i:i:??!:!:_ _'_:__!:!';'!.'";_.'_:

:.:.:.:°:.:.:':-:.:-;.:.:"-:-:°:.:.Z.:.:+Z°:':°Z".:g l_'il_{e._5.;';°:':':°:+" . .'_'_':'."_'_'i

:°:.:°:.-°:. ":"°:'"':_°:';':":':':':'":"'"°"°'"':" :.".:°:.:.:._.:+;:'.:::::"':5:" .''"'""'°" ":':"..::1.___¢ ""°'""".'_'i "'"°' __"""_'° _"_--_ "°'"

Fig. 1. SEM micrograph of aggregates of Hyperion

nanotubes. Multiple tubes are arranged semi-parallel

to each other to form bundles with diameters between

1 and 2 gm and lengths around 20 gm. Inserts: High

resolution TEM micrograph of a section of a single

nanotube; TEM micrograph of dispersed nanotubes.

Average diameter is -8 nm.

Table 1. Comparison of properties of Hyperion nanotubes to those of multiwalled buckytubes

Hyperion nanotubes Multiwalled Buckytubes

Produced in loose powder form (0.05-0.1

g/cc), can be dispersed

• High tube purity, nearly 100% carbon is in

tube form.

• Uniform diameter, average around 8 nm.

• Curved and bent

• Crystallized with lots of defect

• Easy to scale up, can be produced at low cost

• Produced as a hard cathode deposit, very

difficult to disperse

• Contain substantial amount of other carbons

• Wide diameter distribution

• Straight

• Highly crystallized, nearly defect free

• Difficult to scale up at low cost

NEW CARBON STRUCTURE FROM CARBON NANOTUBES

For many applications in the areas of catalysis, separation and energy storage, it is highly desired to have

carbon structures with properties, including mesoporosity, narrow pore size distribution, high surface area, defined

surface functionality, and high mechanical and chemical stability. Using Hyperion carbon nanotubes as starting

building blocks, we have successfully prepared a carbon structure, which meet such requirements. The new carbon

can be prepared in forms of extrudates for catalyst support, particles for chromatographic separation medium, and

sheet electrode for high power energy storage devices. Fig. 2A shows SEM micrographs of extrudates.

°.._ -._,*.°._ ,),...o... ,...,.,..-.O..o..O.O.-.°°.°.....°_o°.......°.-..°..-_

_, ¢:-e,.,4_2,_ ,_.:.::.:.:+:.:.:.:.:.:.:.:.:.: :.:.:.:.:.:...: .:_ 1

_,4,'_N:'_ _:':':':':':':':':':':':':':':':':::':':°:':':':':.'._'.'.'_J"

__',,"-_'_ _ii!!i!iiiiiiiiiiii!i!ili!_ii!i!i_ii::ii:ii!!!.:!i:i!!!:'!i!iii::i

_, ,,.,_ , ::::::::::::::::::::::::::::::::::::::::::::. !_;::°... _ 08

0.6

N .......... '....

.:.:.:.._:.:..:.:.:.:.:.._- .° . °° • ..:.:.:.:. . • :...: "°:.:'._

•:.:.:.:.-'.:':.:.!._:.:'_.:.:.:.:°N.N-N_.T_I._.:.._::::;I,_I_.:,:.:.I.Ig.: -:.:.: 0.2

: ........_.:.:..-,,.-....,....-. ... °....1... ...:.....*,-....;.:. •.....pj ..... -...v°-..e_..'R_:.':-'{...:::::_.':::_:::...°_-.-..... _.

_. ..... :...:.:.,.. ........,....,- ...._......: ..._ • :

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

/

f

10 100 1000 10000

Pore Diameter (A)

Fig. 2A&B. SEM micrographs of bimodal pore structure of extrudates prepared from nanotubes, inserts: 2 mm

diameter extridates and pore structure of grains; BJH desorption cumulative pore volume of extrudates prepared

from nanotubes.

NASA/CP--2001-210948 653 August 6, 2001
Theextrudates haveabimodal porestructure, macropores betweengrainsandmesopores inidethegrains.The

mesopores arewell-connected spacesinthenetworkofintertwinednanotubes; hencetheyhavelowpore-resistance

forchemical transport. Fig.2BshowsaBJHcumulativeporevolumeplot.Itcanbeseenthathestructureisnearly

freeofmicropores. Thesurfaceareaoftheextrudates is341m2/g.Thetotalporevolumeis0.87cc/g.

Fig.3AshowsaSEMmicrographofnanotubesheetelectrode. Incontrast toextrudates, it waspreparedto

havemesopores only.Electrochemical capacitors 4-8fabricatedwithnanotubeelectrodesshowedunprecedented

powerperformance andfrequencyresponseproperties. ARCconstant of7milliseconds wasmeasuredfroma10-

celldevicewithsulfuricacidastheelectrolyte. Thedevicehadacapacitance of 0.22F,anequivalent series

resistance(E.S.R).of0.033_ andanoperational voltageof 10V.It hadexcellent energystorageefficiency; all

storedenergywasavailablefordischarging uptofrequencyof189Hz(Fig.3B).

-0.10

_" -0.08

o

N _.04

-0.02

/

0.00 ' ' ' _ ' ' /' _ ' ' ' I ' ' ' _ ' ' '

0.00 0.02 0.04 0.06 0.08 0.10

Z' (Ohms)

Fig. 3A&B. SEM micrograph of a nanotube electrode; Complex-plane impedance plot of a 10 cell electrochemical

capacitors fabricated with nanotube electrodes.

NEW NANOMATERIALS PREPARED USING NANOTUBES AS TEMPLATES

Because of their unique properties, Hyperion carbon nanotubes are ideal templates for new nanomaterials

synthesis, in particular, high temperature ceramic materials. We have showed that large quantity of SiC nanofibrils

(Fig. 4A) with diameters around 15 nm can be prepared by direct reaction of Hyperion nanotubes with SiO vapor at

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V'.:.'._'. _ "-'.'." ...'.'.......,............-.'..+....'.-.'.............'.... -_2 • _ _ :_'; ". " "_ ._ :i._,' ....:_........:......... =======================================

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... .. ..................,..._;.,..k_, .....,,:.. , • • .

'.".,'.'.'.'.'." .'..... .'.'-'. ":'.'-'." .'.'-.x-;,.,_....:' " "." .. " _ _" .y................. ..,. _..,.,.......,....... . +.... , .,....×_. • ._ . . .

'.,.,.'.'.'.:.:'." ::.:,..: .,:.'_£_i.._,:,:..

•.t,...,,.,°.....,:':.'.._,,:':,IRe,,NN,,I,;.;_..'°.:.,.:...;_.;,,_,,,.:...:.._;:_,. .; :.-.- 4,,l_5,,_.,_...-'o_,._z_'_..:_._N.;c:_e.2:.N..if.-..

.... :::::::::::::::::::::::::::::::::::::: .................. __. "".._'J__:x " :::::::::::::::::::::::::::::::::::::::::::::: ................................................... ":_::.'::.-_;'_:::_Z:.';c ._ .......... ;'_'4":x- ,_....'.-_,c "":'!:_';"_;:'-."_;';_i::: ..... ,_,o,:..............

Fig. 4A&B. SEM micrograph of SiC nanofibrils prepared by reaction of Hyperion nanotubes with SiO vapor, insert:

TEM micrograph of SiC nanofibrils, the average diameter is -15 nm; SEM micrographs of Si3N4 nanowhiskers,

insert: electron diffraction pattern of a single whisker. The zone axis is [2110].

NASA/CP--2001-210948 654 August 6, 2001
high-temperature 7. Thereactionwascontrolledsuchthatdiffusionwasprimarilylimitedto thedirection

perpendicular totheaxesofnanotubes toachieveapsuedo-topotactic conversion. TheSiCnanofibrilshaveahighly

crystallizedmultigrainstructure, whichwereformedbybondingsinglecrystalneedles(grains)atends.Wehave

alsopreparedo_-Si3N4 singlecrystalnanowhiskers (Fig.4Banditsinsert)bycarbothermal reduction-nitridation of

SiOusingnanotubes asstartingcarbonmaterials 8.Diametersofwhiskerswerefrom20nmto200nm.Webelieve

thatstructural natureofournanotubes andtheiraggregates playedanimportant roleincontrolofthediametersof

nanowhiskers.

NANOCOMPOSITES BASEDONNANOTUBES

Fig.5showsaschematic preparationprocedurefornanotube-polystyrene composites. Composites with0.1

to12%nanotubes werepreparedbydispersinganotubes instyrenemonomer first,followedbyin situ bulk

STYRENE MONOMER & HYPERION NANOTUBES

Dispersing

A dynamic Nanotube

network filled with

styrene monomer

Polymerrization

Polystyrene chains Nanotubes

Nanotube-polystyrene

composite in which

nanotubes are highly

dispersed.

Fig.5. Schematic diagram of preparation procedure for nanotube-polystyrene composites.

Table 2. Resistivities of nanotube-polystyrene composites

PS-Nanotubes (f_-cm)

Nt (%) in situ Polymer Solution Mix Latex Mix

0.10 650

0.25 101.8 ---

0.5 15.8

1.0 8.8

2.0 ---

4.0

12 0.18

IM

>108 ......

>108 ......

14950 263 >108

330 ......

Nylon 6-6/CF PC/CF Nylon 6-6/CB

>108(7%) 104 (10%) 106(3%)

104(15%)

NASA/CP--2001-210948 655 August 6, 2001
polymerization. TGAanalysisofcomposites showedthatanear100%monomer topolymerconversionwas

achieved. Glasstransitiontemperatures fom97to99°Cweremeasured byDSC,whichsuggested highmolecular

weightwasobtained. Table2compares resistivitiesofnanotube-polystyrene compositesothosemeasured from

composites preparedbysolutionmixing,latexmixingandinjectionmolding. Resistivities ofseveral nylon66-

carbonfibers/carbon black,polycarbonate-carbon fibercomposites arealsolisted.It showsthatnanotubes have

greatconductingetworkformingability.However, totakeadvantage ofit,agreateffortisneededtodispersethem.

It isremarkable thatataloadingleverof0.1%,aresistivityof650_-cmwasmeasured.

REFERENCES

1. S.Iijima,Nature, 354,56(1991).

2. H.Tennent, USPatent, 4663,230, (1987).

3. C.Synder, etal,WO89/07163, (1989).

4. C-M.Niu,etal,"Highrateelectrochemical capacitors basedoncarbonnanotubeelectrodes", inproceedings of

39thPowerSourceConference, CherryHill,NJ,June,2000.

5. C-M.Niu,etal,Appl.Phys.Lett.,70(11),1480,1997.

6. C-M.Niu,etal,"NanoCapVM: Highpowerelectrochemical capacitors", inproceedings of8thinternational

seminar ondoublelayercapacitors andsimilarenergystoragedevices, DeerfieldBeach,Florida,1998.

7. C-M.Niu,etal,Mater.Res.Soc.Symp.Proc.,410,179(1996).

8. C-M.Niu,etal,USPatent 5,814,290, (1998).

ACKNOWLEDGEMENT

Electrochemical capacitor andcompositeprojectswerepartiallyfundedbyaphaseII SBIRgrantfromBMDOanda

phaseI SBIRgrantfromDARPA,respectively.

NASA/CP--2001-210948 656 August6,2001
Proceedings oftheSixthApplied Diamond Conference/Second Frontier Carbon Technology Joint Conference (ADC/FCT 2001) Y.Tzeng, K.Miyoshi, M.Yoshikawa, M.Murakawa, Y.Koga, K.Kobashi, G.A.J. Amaratunga (Editors)

NUMERICAL ANALYSIS ON THE DISPERSION PROCESS OF CARBON CLUSTERS

IN GAS EVAPORATION SYNTHESIS USING D.C. ARC DISCHARGE

S. Usuba, Y. Kakudate, H. Yokoi, H. Suzuki and S. Fujiwara

National Institute of Advanced Industrial Science and Technology

1-1 Higashi, Tsukuba, Ibaraki 305-8565, Japan

Phone: +81-298-61-4789, Fax: +81-298-61-4783, E-mail: s-usuba@aist.go.jp

ABSTRACT

For the purpose of understanding the mechanisms of buffer gas effect on the yields of carbon clusters in the gas

evaporation synthesis using D.C. arc discharge, a residence time of cluster in high temperature state was calculated in

a simplified process where a single cluster was emitted from the arc and dispersed in the buffer gas. The results for

C60 dispersion in helium showed that in the case where the thermal convection could be excluded such as a synthesis

under micro-gravity condition, the residence time increased monotonously with the gas pressure, while in presence of

thermal convection such as in the case of synthesis on the ground, the residence time became maximum at the

pressure range of 0.2 - 0.3 atm and decreased with pressure in higher pressure region. Since the residence time was

considered to be a good index of attainable clusters growth stage, it was suggested that under high pressure condition

where the thermal convection became significant, the progress of growth processes of particularly large size clusters

was greatly inhibited by the effect of thermal convection. Based on this idea, a simple model was proposed to explain

experimentally observed buffer gas effects on the clusters yield.

Keywords: fullerenes, nanotubes, arc, convection, micro-gravity.

INTRODUCTION

It has been known that the yields of carbon clusters in the gas evaporation synthesis using D.C. arc discharge

were remarkably influenced by both species and pressure of the buffer gas (ref. 1). Understanding the mechanisms of

this buffer gas effect has been required to improve the efficiency of synthesizing various types of carbon clusters

including fullerenes and nanotubes. For this purpose, the effects of convective transport on the clusters yields were

experimentally examined by using micro-gravity facilities (refs. 2, 3). In this paper, we calculated the temperature

histories of the clusters in a simplified vapor dispersion process where single cluster was emitted from the arc and

dispersed in the buffer gas. The purpose of this calculation was to examine the buffer gas effect from the viewpoint

of a residence time of the cluster in the high temperature state which was necessary for clusters growing processes to

be promoted. Calculated residence time was considered to be a good index of attainable clusters growth stage, which

strongly influenced the cluster compositions of final products.

MATHEMATICAL METHOD

Calculations were performed in the following two steps. Firstly, velocity

and temperature distributions of the buffer gas in the gas evaporation vessel

were calculated under the boundary conditions of the cylindrical vessel of

100mm in the diameter and 100mm in height including the D.C. are which was

expressed as a heat source of 5000K as shown in figure 1. Calculations for both

helium and argon as typical buffer gases were conducted with a computer code

developed by using SIMPLE (Semi Implicit Method for Pressure Linked

Equations) technique (ref. 4). This code can deal with the time dependent

axisymmetric laminar flow of viscous fluid with variable density. With this

code, solutions of velocity, temperature and density are obtained at 40 x 20

grid points on the plane including the axis of rotation.

In the second step, the temperature histories were calculated for two

I t00mm [

Figure 1. Boundary conditions

of the gas evaporation vessel

using D.C. arc discharge.

NASA/CP--2001-210948 657 August 6, 2001
different sizedclusters, C2andC60, whichwererepresentingypicalclustersexistinginthecarbonvapor,ina

simplifiedvapordispersion processwherethoseclusterswereemittedfromtheareattimet=O and dispersed

independently in the buffer gas. The probability density _pof cluster location at a certain time t was calculated by

O_o __Vr_____+ Vz ___.___:div(Dgrad_), (1)

_t or oz

with a condition of cluster preservation,

fIj dv=l, (2)

where vr and vz were r and z components of steady state gas velocity calculated in the previous step, dv was a

volume element. Integration in equation (2) was done over the whole volume in the vessel. D was a diffusion

constant of cluster in buffer gas under the solid body collision approximation (re£ 5) given by,

3

where T and P were temperature and pressure of the gas, N_ and/ca were Avogadro's and Boltzmann's constant, rg

and mg were radius and atomic weight of buffer gas atom, and rc and me were radius and molecular weight of cluster

approximated by spherical shape. Then, the cluster temperature Tct_,e,, which was 5000K at t=O, was calculated by

using the probability density as

Tr,.... = IIIT._.dv" (4)

Integration in equation (4) was done over the whole

volume in the vessel. Tct_t,r was a decreasing function

of time. Finally, a residence time _'was defined as the

time required for Tct,_t_ to change from 5000K to

1500K which represented the lower limit of high

temperature state which was necessary for the clusters

growth to be promoted (ref. 6).

RESULT AND DISCUSSION

Figures 2(a) and 2(b) showed the typical results

of velocity and temperature distributions in a steady

state under three different buffer gas conditions. The

velocity results showed strong dependency of

convection strength on the density of enclosed gas.

The maximum convection velocities which appeared

near the arc surface in cases of helium 0.1 atm,

helium 1.0 arm and argon 1.0 atm were calculated to

be 0.21, 0.94 and 1.13 m/s, respectively. In the same

order, the radii of high temperature zones which

could be approximated by spherical shape as shown

in figure 2(b) were 13.5, 8.8 and 7.3mm, respectively.

Figure 2(c) showed the probability densities of C60

location at t=_. It was shown that the probability

densities under three different buffer gas conditions

had different distributions although they gave same

values of Tcluster.

The pressure dependencies of the residence time

"t'of C2 and C60 were shown in figure 3, where results

of diffusion-only case which corresponded to the

micro-gravity condition where the effeet of thermal

convection could be excluded were also plotted. In

the diffusion-only case, v was almost linearly

increasing function of pressure. This indicated that

without the thermal convection, the cluster growth

Heat ::=::==:-::7::::

sooroo i!iiiiii

asAre 'i!ii[l_i

\ iilii (a)

................ !!

High

Temp.

Zone

1500-

5000K

(b)

x=33ms

(c) __.OOl

He 0.latin He latin Ar latin

Figure 2. Calculated results : (a) flow velocity;

(b) temperature; (c) probability density of C6o

location cluster for T¢t,,_,,= 1500K at t=-_;

NASA/CP--2001-210948 658 August 6, 2001
100

oo

E

I-- 10

Helium ,' " .'" '

oO °

0.01 0.1 1 10

Pressure /(atm)

100

10

0.01

Argon '

°

°°,,•"

0.1 1 10

Pressure / (atm)

Figure 3. Pressure dependencies of the residence time _"

of Cz and Cso in helium and argon.

could simply progress to higher stage as the gas pressure increased. On the other hand, in presence of thermal

convection, the value of r showed remarkable shift to smaller value compared with diffusion-only case in the

pressure region more than 0.1 atm for helium and 0.02atm for argon. Furthermore, in the ease of C60, rbeeame a

decreasing function of pressure in the pressure range more than 0.3atm for helium and 0.05atm for argon. Although

the physical meanings of the differences between helium and argon in the pressure dependencies of rwere not clear,

the results suggested that under particularly high pressure condition, the progress of clusters growth process

involving large size clusters was greatly inhibited by the effect of thermal convection.

To relate above considerations with experimentally observable pressure dependencies of the clusters yields, a

simple model of the pressure dependency of cluster size distributions in the final products was developed. Here, the

cluster size meant the approximated radius of the cluster. Following two assumptions were introduced. The first was

that the size distribution was described by a normal distribution function with a constant relative variance, o/d=const,

where o and d were the standard deviation and the mean cluster size, respectively with reference to experimentally

observed size distributions of fine carbon particles produced by the thermal decomposition of benzene (ref. 7). The

second assumption was that the mean cluster seize simply increased as the growth process proceeded. Using those

assumptions, the contour maps of the clusters yield as a function of both cluster size and gas pressure were

schematically plotted in figures 4(a) and 4(b). Figure 4(a), corresponding to diffusion-only case, showed that the

mean cluster size increased monotonously as the gas pressure increase. On the other hand, figure 4(b) showed that in

presence of thermal convection, the mean cluster size shifted to smaller size compared to the diffusion only case and

became a decreasing fimetion of pressure in high pressure region. From this model, yield of certain cluster could be

derived as a function of gas pressure. Characteristic results for two different sized clusters were shown in figures 4(c)

and 4(d). In the case of cluster of size dl, the yield became maximum at almost the same pressure in both diffusiononly case and convection existing case. In higher pressure region, the yield decreased monotonously with pressure in

diffusion-only case. On the other hand, in convection existing case, the yield kept higher level and increased again

/,.,oi.u.iono 

d2 .................

>-

(c) dl Yield

Diffusion+

_:_,. Convection

BufferGas Pressure

Figure 4.

(b) Diffusion+ Convection

(d) dzYield

Diffusi%nlY

Diffusion÷ ............. _.

Convecti%_ ...........

Buffer Gas Pressure

Simple model of the pressure dependency of cluster size distributions in

the final soot products and derived clusters yields as a function of pressure.

NASA/CP--2001-210948 659 August 6, 2001
with pressure. Experimentally observed pressure dependency of C60

yield in helium was shown in figure 5 in comparison with that obtained

in micro-gravity condition using free-fall facility (ref. 2). It was found

that the observed difference between 1-g and 0-g condition in the

pressure dependencies of yield was consistent with our model as shown

in figure 4(c) except that the second increase of the yield in convection

existing ease was not confirmed in the ground experiment shown in

figure 5.

Figure 7(d) showed another important result in the case of larger

cluster of size d2 that the maximum yield in convection existing case

was extremely small compared to that in the diffusion-only case. This

meant that the growth of clusters of particularly large size was strongly

inhibited by the effect of thermal convection. To generalize this result,

it was suggested that the thermal convection might be a serious obstacle

for effective synthesis of all clusters whose growth processes required

long residence time of their precursor clusters.

0.0

' I ' I ' I ' I ' I ' I

• Ground (l-g)

0 Free fall (O-g)

°°•0_•

• OOOOl

0 •0

0

0

, I , I , I , I . I , I

0.2 0.4 0.6 0.8 1.0 1,2

Initial Helium Pressure/(arm)

Figure 5 Observed pressure

dependencies of C_o yields in helium by

ground (normal gravity) and free-fall

(micro-gravity) experiments (ref. 2).

CONCLUSION

By calculating a residence time of cluster in high temperature state in a simplified process of gas evaporation

synthesis using D.C. arc discharge, it was revealed that the thermal convection had an effect in obstructing the

progress of clusters growth particularly involving large size clusters. A simple model of the pressure dependency of

cluster size distributions in the final products developed based on those results could explain experimentally

observed effect of micro-gravity on the pressure dependence of C6o yield. According to this model, it was concluded

that for the effective synthesis of clusters whose growth processes were considered to require long residence time of

their precursor clusters, such as large size fullerenes, endohedral metallofullerenes and single-wall nanotubes, the

thermal convection might be one of the important factors to be controlled.

REFERENCES

1. Saito, Y., et al.: Yield of fullerenes generated by contact arc method under He and Ar., Chem. Phys. Lett,

vol.200, no.6, (1992), 643

2. Usuba, S., et al.: Synthesis of Carbon Clusters from D.C. Arc Plasma under Reduced Convective Transport

Condition., Proe., ADC/FCT'99, (1999), 219

3. Mieno, T.: Effective production of endohedral metallofullerenes in repetitive gravitation-free are discharge

using a vertical swing tower., Jpn. J. Appl. Phys. vol. 37, (1998), L761

4. Patanker S. V. and Spalding D. B: A calculation procedure for heat, mass and momentum transfer in threedimensional parabolic flow., Intl J. Heat Mass Transfer, vol. 15, (1972), 1787

5. Williams, F.A.: Combustion theory- the fundamental theory of chemical reacting flow systems, second ed.,

The Benjamin/Cummings Publishing Company, Inc., MA, (1984)

6. Haufler, R. E., Mat. Res. Soc. Symp. Prec., 206, (1990), 627

7. Siega, D.C. and Smith, GW.: Particulate carbon., Pergamon Press, Oxford, (1981), 152

NASA/CP--2001-210948 660 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

An electrically conducting boron-doped carbon nanotube polystyrene composite

P.C.P. Watts, W.K. Hsu*, H.W. Kroto, D.R.M. Walton

School of Chemistry, Physics and Environmental Science

University of Sussex, Brighton BN1 9Q J, UK

G.Z. (;hen, D.J. Fray

Department of Materials Science and Metallurgy,

University of Cambridge, Pembroke Street, Cambridge, CB2 3QZ, UK.

Abstract

A boron-doped multi-walled carbon nanotube/polystyrene composite exhibits relatively low electrical resistance,

accompanying by minor variations in conductivity, when compressed or bent.

Introduction

Recently, boron-doped multi-wailed carbon nanotubes (BMWCNs), have been produced by an arc discharge process. The

boron content is ca. 1% by wt, as revealed by EELS[I,2]. Two features readily distinguish multi-walled carbon nanotubes

(MWCNs) from BMWCNs. (a) The conductivity of BMWCNs is greater than the MWCNs, and (b) the BMWCNs are

longer[I,2]. It is believed that boron acts as a catalyst in the arc, networking preferentially at the tube tip. In this paper,

conducting composites of MWCNs/or BMWCNs and polystyrene were produced and the resistance of both composites was

found to be relatively lower than that of a carbon black-polystyrene composite; in particular, the minor variations in

conductivity were apparent when both composites were subjected to mechanical loading.

Experiruental

MWCNs and BMWCNs, produced in the arc discharge, were subject to mild oxidation and filtering in order to remove

polyhedral particles[I,2]. TEM revealed that carbon particle content was relatively low, compared with unoxidised materials.

Nanombe-polymer composites were produced by mixing MWCNs and BMWCNs respectively with a solution of polystyrene

(2g, Aldrich, UK) in toluene (20-30 cm 3) in various percentages (i.e. 2.5, 5, 7.5, 10 and 12.5 %, polymer carbon content, by

weight). The toluene solutions containing nanotube and polymer suspensions were ultra-sonicated (10 rain), then transferred

to a Petri dish (9 × 1.3 cm) and the dish was subjected to a vacuum (10 2 torr) in order to remove the toluene. The residual

MWCN and BMWCN-films (ca. 0.3 mm thick) were detached from the dish using a scalpel. A hot compact carbon

black-based composite polystyrene (HCCBCP, 1 mm thick) was purchased (Goodfellow, UK) in order to mmke comparison

with nanotube-polymer composites. The carbon-to-polymer ratio in this commercial product is possibly 1:1 (wt), which is

commonly found in industry. A 1 mm thick BMWCN-film (12.5 %, polymer carbon) was also made for comparison with the

HCCBCP-film. Due to the elongated nanotube structure, the differences in morphology between the MWCNs (or BMWCNs)

and the polymer can be easily distinguished by SEM and TEM. For AC impedance measurements (Solartron 1260 Frequency

Response Analyzer) at room temperature, two- and four-probe geometries were employed. The composite films and

electrodes were assembled so that the film was sandwiched between two Al-discs (1 cm, diam) suitable for a two-probe test

(Fig la). A ca. 20 kg load was then applied to the discs. For the four-probe test, the film was sandwiched between two

fibre-plastic plates (Tufnol, RS, UK), one of which contained four aligned needle-like Cu screws (ca. 0.5 cm apart). The

plates were tightened and the Cu-electrodes were screwed down so as to maximize contact with the film (Fig lb). For two- or

four-probe tests, the samples were examined under loading (vide supra). The film bending test was carried out using silver

wire electrodes (3 cm apart), attached by means of silver paint to the films (1 × 5 cm), so that each film could be bent

manually and reversibly during multiple meter monitoring. The bending angle lay in the 0-45 ° range.

NASA/CP--2001-210948 661 August 6, 2001
+.e+) .,X(-C_+_.,,_.

+

+,+,,+, - +++ ,:o,o,+,+,++,,+_+*tl)+O

Fig 1. (a) Two-probe measurement device,(b) four-probe measurement device.

Results and Discussion

Fig 2 shows an SEM image of a BMWCN-film (cross-section, 0.3 mm). The film thickness is fairly uniform (Fig 2a) and its

internal structure consists of small dark voids, surrounded by a white composite matrix (Fig 2b). Fig 3 shows typical SEM

and TEM images of BMWCNs in the polymer. The tubes are arranged at random within the matrix (Fig 3a). Tube-tube

contact, either via tube bundling or tube-tube crossing, is evident (e.g. Fig 3b), and the electrical conducting path is therefore

established. TEM also reveals the presence of polymer-coated nanotubes (arrows, Fig 3c), the coating influencing the

tube-tube contact structure. First, if the individual tube surface coatings are established after tube-tube contact, the tube-tube

contact structure will be strongly maintained. Second, if the nanotubes are coated before tube-tube contact, the coating

prevents electrical contact between adjacent tubes.

".'°_°'_-_--°'°-°'_'_'.;._._.;.;.;._._.;.:._._.:.:._._.;.;._.;.;.;.;.;.;.:.:.;.;.:._.:.:._._.:._. • , ,:, • ..;.:...2+;.:.:.:.:+:.:.:.:.:.:.:.:.:.: . •. -

-+ • ...... . .......... ....

•: _ _ _ i ?,+,-,-,.'.'.:.'.:.'.'.:.'.'.'.'.'.'.'.'.'.'.:.:.:.:.'.'.:.'.:.:.'.'.:.

_'.%'.'.'.'°°+°+++'° .'.+'.%%%'N°.%_°"iI_'_-_+.%_.,+._++%=_'++.'. + : ,_J'.'_+,_d" _ " "o'_".'i_,.°,_.'Jl%'i't"_i ,t."_.;.'.'...'.-.-.+.-.,"..°....-.........',.-,,-.,-...+. .. ...-.-.,-+......1...-.. + , • _...,e.-_...,'+"e. ',t- .'.'.'.'.'.'.'." .'.'. °. ,'-'.'-'. ".""'_ .'.'.'-'.'.'.'.'.+- .... "." ::::::::::::::::::::::::::::::::::::::

:.:.:.:.:.:.;-:.:.:;:-:...:-:.:....R..:.:::... .'" • ". :.5 .'_ :::::::::::::::::::::::::::::::::::::::::::

.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:.:...v.:.:...:.i:_-.::::....-:::...:.:...:.,: _ao:,g__.'oz'..,-_:: - ..

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::: ..:": ....... ...................................................................:.:..-...., -.N

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::_:i iiii i : . iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii 

Fig 2. (a) SEM images of BMWCN-f'dm cross section, (b) Enlarged SEM image of (a).

Various electrical contacts exist in the polymer matrix, including tube-tube, tube-particle and particle-particle. However, the

tube-tube contact dominates, because the quantity of particle present is limited by oxidation. The resistivity along the c-axis is

ca. 3-4 orders of magnitude greater than along the in-plane direction (10 6-10a fl.m) in graphite, due to the difficulties in

promoting electron flow across the 3.4 A layer separation. Nanotubes in contact (i.e. basal plane contact) implies that the

electrons flow within the polymer matrix via routes of lowest resistance. Fig 4 shows the surface morphologies of

BMWCN-films (0.3 mm thick), top and bottom faces respectively, for various concentrations of nanotubes. Based on Fig 4, a

few distinguishing features became apparent as the electrical resistance changed. (a) At low concentrations (i.e. 2.5-5 %), the

difference in resistance between the top and bottom faces is ca. one order of magnitude, which is also reflected in the

different morphologies seen in Figs 4a and c (top faces) and Figs 4b and d (bottom faces). In Figs 4a and c, the surface

morphology consists of white islands, surrounded by dark regions. The largest island is ca. 0.5-1.2 _tm wide. The bottom face

morphology (Fig 4b and d) is similar to sand-paper surfaces. We did not observe nanotubes protruding from the top and

bottom faces; in fact the nanotubes are likely to be embedded in the polymer with either tube bodies or tube tips partly

exposed to air.

NASA/CP--2001 210948 662 August 6, 2001
A B

Fig 3. (a) SEM image of BMWCNs in polystyrene; (b) TEM image of BMWCNs in polystyrene; (c) TEM image of

polymer coated MWCNs (arrows). Insert: enlarged TEM image of coated MWCNs. Large arrow: polymer coating.

Arrow A: central core of carbon nanotube. Arrow B: carbon nanotube surface. The fringe contrast is darker in the

coated carbon nanotube body than in the coating polymer.

The formation of sand-paper surface-like structures (bottom face) possibly reflects the presence of denser nanotube

aggregates, due to nanotube settlement during film formation. The white and dark regions seen in the top faces (Fig 4a and c)

are either tube-rich or tube-poor, however we were unable to locate tube-rich regions by SEM. The surface morphology of

both top and bottom faces is essentially the same when the BMWCN concentration exceeds 7.5 % (Fig 4e-h). In other words,

nanotubes are evenly spread throughout the polystyrene matrix. The 1 mm thick BMWCN-film exhibits a similar

phenomenon when present to the extent of 7.5 %.

3. Impedance measurements

a. Compression

Fig 5a-c shows two-probe impedance spectra of MWCN and BMWCN-films (0.3 mm thick) under 20 kg loading, together

with 1 mm thick BMWCNs and HCCBCP-films. The HCCBCP impedance varies within the 103 - 105 _ range for all

frequencies (100-10,000 Hz) (Fig 5a). Variation begins at ca. 12,000 _ (10,000 Hz, arrow 3, Fig 5c), and gradually

NASA/CP--2001-210948 663 August 6, 2001
_,-...°.°....-.. • °.° .,.°-°oo_,_.o_,_-o°,_.,...............°°°. _,_,. ,- ,_ °....°...._....._.:._.:::._........-............. :....

_%-_oo. _.Oo-_Oo_Oo_p_. • ._.o_°.d,,%_.... ... . ... °°..°. •

_i!i!i!_i_i_.xx.ii!!!!i!iii!!!!!!!!iii!i!iiiii I

:-:i.i.-.i:i-::::::.:-.i:i:::-:::':'.'.i:'.'..':i:'..i:i:i:..:?'{._._:

:.:.:o:.:.:.:.:.x.......:.:.:.'.-....'...........:...-.:.-.-..-.:..'.....

x.-.x.-...-.'.:.:.:...:.:...:.'.'.....-.:...-.'.:.'.....:.:.'.......,

:._:!:i-i:!:?_:i:_:i-::;.i._:?:.::':i:::i:::::-.i.::i--:::i.i.i.-.:.:.:_ _!&

•:.x.:.:.:.:.:.:o:.:.:.:...: .-.:.:.:.:.:.:.'.:.:.:.:.:.:.:.'.:.-.'.:

_i:i:i:i:!:i:?i:i:i:i:::i:ii:_?i:.:i'i:i.i":?i::'"!:i"'i: :'_

!_!!!!!!'i_: i:i:_!i:_:i: "'"'!_iiiii:iiiiiii.'?iiii

 iiiiiiiiiiiii iiiiiiii!iiiiii!iiiiiii!!iiiiii!i iiiiiiiiiiiiiiiiiiiiiiii

_i:i.!:i:i:!:i::::':':::'i:'":':':'::::?i:"::'::::'i:::::":.'i:'.':: i ":-":i:i:i:i_:_:i

..:ii!ii_iiil

O°°;.'°:°;°:.

..:.°.O.....

........ 4_:_.:.x.:"

_i_ii!iii!i_<._i:i.'°:°.::.:;'.::.::_::_:::_:°_°..';;:_:_'_::_:_::_:_:::

°..°°°°°°° .1_°.° ° °°°°;°-.°°°.:°°.°°.°-.:°.°°°°..+.°..:...°..

_.,..-o°°°°°° Oo°°°-._o°...°..°_l.o_o...o°....°o°°°_-.°°°°o°°,..°°,

:'._::::_:::::.o.::::::._..'::.,-_._::_._:::::::::::::::::::

,_.1...: °....° ° .O-°... .._ .:.10_.:.:.-.:...'.:

_2ff.. _o',.:i:!::::_.::::':::::_:._::._':.'::::-::::::

Fig 4. SEM images of top and bottom face morphologies with various BMWCN

content: (a-b) 2.5 %; (c-d) 5 %; (e-f) 7.5 % and (g-h) 10 %.

increases to 38,300 _ at 260 Hz (arrow 4, Fig 5c). The impedance eventually reaches 36,600 _ (arrow 5, Fig 5c).

Nanombe-based composite films exhibit relatively low impedances (2-5 × 102 _) without significant variation in conduction

(Fig 5a and b). The BMWCN-films, either 0.3ram (175 _) or 1 mm (450 _) thick, exhibit an impedance one order of

magnitude lower (Fig 5b) compared to MWCNs (3,000 _). Small humps are distinguishable in the MWCN-films (arrows 1

and 2, Fig 5c). Unfortunately, four-point impedance measurements could not be carried out on the thin MWCN and

BMWCN-films (0.3 mm thick), due to film penetration by the Cu needle electrodes. Fig 5d shows the results of four-probe

NASA/CP--2001-210948 664 August 6, 2001
measurements onthe1mmthickBMWCNandHCCBCPfilms.Theimpedance oftheHCCBCP-filmvariesdrastically, in

markedcontrast totheBMWCN-film(700_,corresponding to3.2_.mresistivity).

b.Bending

ThevariationinHCCBCPresistance becomes significant whenthefilmisbent,heamplitude beingonetotwoordersof

magnitude. ThevariationinBMWCNresistance lieswithinthe10-30_ rangewhenthefilmisbent.

.... :- - .+ ........ . _. _ +++ _+_+_++_ * _ +,++,+++..-..,. o_ +'

5 _+,u_o4 3.50E+_ "

+

- p.Or,F._,j _\

5._E*_ " "+"

O 0¢E+_3 ; ............................

I

i i

i i+ _

/+" %' t _+ L,

,,_ • ,+ '%++

l_¢qu_} ( !o_ It_.t Fr_ti_t_ne) (l#g [Izl

Fig 5. (a) Two-probe impedance spectrum of carbon materials labeled. Iz lis the impedance module. (b) Enlarged

impedance spectra of two BMWCNs-films of (a). (c) Enlarged impedance spectra of HCCBCP and MWCN-films of

(a). (d) Four-probe impedance spectra of HCCBCP and BMWCN-films (both 1 mm thick).

(b)

5 Oct-,-(;+'...............................

0 2 5r;_O_

-- 2,ooL_,o@

I _912+ 02

+,OCI= e_2

c

Network formation, via tube-tube bundling and crossing within the polymer matrix, results in relatively low resistance at low

polymer carbon percentages (10-12.5 %), as compared with spherical carbon particle-based conducting polymers (30-50 %

commonly found in industry). This conducting network is robust and is less segregated than the spherical carbon particles

when loaded. We assume that the pressure applied (20 kg) to the composite film induces small polymer creep during

impedance tests. The spherical carbon particles segregate and aggregate simultaneously in the presence of creep, resulting in

conductivity variation (HCCBCP-film, Fig 5a and d). Tube-tube detachment or separation is difficult. For example, the

surface coating (insert, Fig 3c), maintains tube-tube contact to some extent in the presence of loading. According to previous

reports, tube bending in the polymer has been observed (Fig 3b) and prevails when the polymer matrix is loaded. Meanwhile,

when the robe-polymer composite is compressed, transference of the load from the matrix to the robes results only in outer

shell tube stress[3]. In other words, the tube-tube is able to maintain contact with minor-distorted outer shells. The drastic

variation in the HCCBCP impedance profile (Fig 5c and d) mainly arises from the presence of a capacitor component: (a)

Carbon black consists of graphitic domains within the particle walls, which lower the electrical resistance. In other words, the

electrical resistance in the particle will be much lower than the particle-particle contact resistance (ca. two/three orders of

magnitude difference). Consequently, a capacitor-like structure is produced, i.e metal/insulator/metal (particle/contact

resistance/particle). The presence of the capacitor component will influence the impedance profile. According to the

following equation, the impedance Z = [R: + (XL-Xc):] m, (R = resistance, XL = inductive reactance, Xc = capacitive

NASA/CP--2001-210948 665 August 6, 2001
reactance). Ifthecomposite flmcontains bothcapacitor andinductor components, i.e. XL ¢ 0, Xc ¢ 0, then the impedance

profile is non-linear. (b) Carbon-black usually contains heteroatoms (e.g. H, O and S), which form metallic or insulating

islands within the particles, or in the carbon lattice. If metallic islands (e.g. odd-alternate g-type radicals) are formed, the

charge carriers travel via tunneling from one metallic island to another[4]. Accordingly, the capacitor-like structure (metallic

island/contact resistance/metallic island) is also generated within aggregated particles. For the MWCN- and BMWCN-films,

the presence of a linear and horizontal impedance profile is indicative of the intrinsic resistors, i.e. XL= Xc = 0, and Z = R.

The geometry of two- and four- probe measurements may also give rise to variation in the impedance profile, due to contact

resistance between film and electrodes. In the two-probe test, the film surfaces were entirely covered by an A1 disc, which

minimizes the contact resistance. In the four-probe tests, four aligned Cu needle electrodes were brought into contact with the

film. If the distribution of the carbon component around the Cu electrodes is non-uniform, the contact resistance increases.

This contention is supported by the presence of less variation in two-probe impedance profiles than in the four-probe system

(HCCBCP-film, Fig 5a, d). However, the BMWCN-film exhibits linear and horizontal impedance profiles, either in a two- or

four- probe test, implying the presence of uniform robe distribution in the polymer. The lower impedance in the

BMWCN-film, compared to the MWCN-film, arises because boron-doping of carbon nanotubes results in increases in charge

density on CNs. The BMWCNs exhibit lower resistivity, in bulk and in polymer, indicating that electrical conduction within

the polymer is actually established via individual tube-tube contact, not via electron hopping from tube to tube. As described

above, nanotubes are intrinsic resistors, therefore the conducting network in the matrix can be expressed as a combination of

resistors in parallel (1/R_ff = l/R1 + l/R2 .... i.e. tube bundling) and in series (Reff = R1 + R2 .... i.e. tube-tube crossing); R1,

R:... represent individual nanombe resistors. If electron hopping from robe to robe is present as a major conduction

mechanism, then the intrinsic resistance of the tube (i.e. Rn, n = 1, 2...) will not significantly contribute to the Reff value,

either in parallel or in series. In other words, the R_ff value of the BMWCN-film will not differ significantly from that of the

MWCN-film, as long as the carbon content of both composites is identical. In practice, the Reff value of the BMWCN-film is

one order of magnitude smaller than the Reff value of the MWCN-film, indicating that individual tubes contribute to the

overall R_ff. The BMWCNs are much longer than the MWCNs. The longer tubes lead to preferential networking and

percolation in the polymer at relatively low concentrations. The presence of a small variation in the MWCM-film (arrows,

Fig 5c) possibly arises as follows. The bulk MWCNs contain both semiconducting and metallic tubes, the former exhibiting a

0.1-1 eV band gap. The MWCN tube-tube contact in the polymer is established via combinations of (a)

metallic-semiconducting tubes, (b) metaJlic-metaJlic tubes and (c) semiconducting-semiconducting tubes. When the current

flows between (b) and (c), a capacitor-like structure is produced.

W.K. Hsu: Corresponding Author: w.hsd@sdssex.ac.dk Tel: ++44 1273 877375, Fax: ++44 1273 677196

Acknowledgements

We thank the Leverhulme Trust (UK) for the research funding.

References

[1]. W.K.Hsu et al, J. Mat. Chem. 2000, 10, 1425.

[2]. W.K. Hsu et al, Chem. Phys. Lett. 2000, 323, 572.

[3]. L.S. Schadler, S.C. Giannaris, P.M. Ajayan, Appl. Phys. Lett. 1998, 73, 3842.

[4]. L.J. Dunne., et al., J. Phys. Condens. Matter 1996, 8, 2127.

NASA/CP--2001-210948 666 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

MAGNETIC COUPLING OF CARBON NANOTUBES

Trinh Vo 1,Tai-Chou Lee 1,Behnaz Payandeh

IDepartment of Chemical Engineering, Rice University

6100 Main Street, Houston, TX 77251

Marc Robert 1.2.3

1Rice Quantum Institute

2Center for Nanoscale Science and Technology,

3Chemical Engineering Department, Rice University

ABSTRACT

The magnetic coupling between carbon nanotubes filled with magnetic elements, which is assumed to be of the

indirect type, is shown to lead to ferromagnetic ordering for arrays of both metallic and semiconducting carbon

nanotubes filled with elements of magnetic transition and rare-earth metals. The critical temperature and

spontaneous magnetization are determined.

The most significant findings of this study are that metallic and semiconducting carbon nanotubes filled with

magnetic elements can be turned into magnetic materials which, depending on the nature and concentration of the

magnetic elements, can be ferromagnetic at room temperature. This creates the possibility of new, artificial

magnetic materials that possess the remarkable mechanical electronic properties of pure carbon nanotubes.

Key words: Carbon nanotubes, direct and indirect magnetic coupling

Corresponding author: Trinh Vo, Department of Chemical Engineering, Rice University, 6100 Main, Houston,

TX 77251, email address .v.9.t(_..fi.c...e.:e_d.u_, fax number (713) 348-5478, phone number (713) 348-2787.

NASA/CP--2001-210948 667 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

"LITHIUM STORAGE IN NANOMATERIALS"

O. Zhou, U. of North Carolina, Chapel Hill, NC

In this talk we present recent results on the effects of processing on the Li storage

capacity of SWNTs. Electrochemistry and solid state nuclear magnetic resonance

measurements show that the reversible Li storage capacity &purified SWNTs

increased from LiC6 to LiC3 after chemically etching. All the SWNTs become

metallic upon Li intercalation with the electronic density of state at the Fermi

level increasing with increasing Li concentration. We will also present results

from electrochemical, x-ray and Raman studies of Li reaction with nanostructured

Si (n-Si). The energy barriers for formation of Li-nSi alloys are significantly

reduced. Reversible electrochemical reaction takes place at room temperature,

400K below what is possible for bulk Si. The high Li concentration and low

electrochemical potential with respect to Li/Li ÷ make n-Si attractive for Li storage

applications.

Work done in collaboration with B. Gao, H. Shimoda, X.P. Tang, A.

Kleinhammes, S. Sinha, L. Fleming, J. Owens, L.E. McNeil and Y. Wu.

NASA/CP--2001-210948 668 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ELECTROCHEMICAL HYDROGEN STORAGE IN SINGLEWALLED CARBON

NANOTUBES

Won Seok Kim, Kay Hyeok An, Young Soo Park, Kwan Ku Jeon,

Hee Jin Jeong, Jeong-Mi Moon, and Young Hee Lee

Department of Physics, Institute of Basic Science, Sungkyunkwan University, Suwon 440-746, South Korea

ABSTRACT

We have investigated key factors for electrochemical hydrogen storage in singlewalted carbon nanotubes

(SWCNTs) to improve storage capacity by controlling the amount of transition metals during synthesis and postannealing process. The influence of morphology and crystallinity of SWCNTs on the storage capacity is also

studied. With increasing the amount of transition metals, the capacity of hydrogen storage in SWCNT is improved,

due to transition metals acting as conducting agent. The plateau region in the discharging voltage profile extends

with increasing the oxidation time at 300 °C in air. The elevation of annealing temperature in Ar also extends the

plateau region, resulting from increase of purity; crystallinity, and electric conductivity of electrode. The maximum

discharging storage capacity of 450 mAh/g (1.63 H-wt%) was obtained at a cutoff voltage of 0.8 "V\ We observed

that the plateau region in the discharging voltage profile, which is of critical importance for practical application,

was maintained during the constant-current discharge mode, particularly at high metal content, long oxidation

time in air, and the elevation of annealing temperature in Ar. The CV measurement also shows redox reaction

peak near 1 V during discharge, in good agreements with the charge-discharge experiment,

Keywords: SinglewaUed carbon nanotubes, Raman spectroscopy, TGA, charge-discharge experiment, CV

measurement.

INTRODUCTION

A new energy storage material with high capacity, light mass, and long cycle life has been strongly demanded

for portable electronics and fuel cells. Large empty space inside the single-walled carbon nanotubes (SWCN_fs) and

large surface area opens a possibility for hydrogen storage applications. Until now, it has been reported that

hydrogen can be stored in SWCNTs from high-pressure approach (reE 1 to 3) that has limitation in batte D'

application. Electrochemical hydrogen storage, however, can be directly applicable to the secondary battery [5-8].

Nuzenadal et al. (ref. 4) showed that the SWCNTs produced by arc-discharge could store hydrogens of

ll0mAh/g, corresponding to hydrogen storage 0f0.39 wt %. Lee et aL (refi 5) also showed that the maximum

capacity of hydrogen in SWCNTs reached 140mAh/g, slightly higher than that of the previous value. However, no

plateau region in the voltage was shown during the constant-current discharge, implying an impractical usuage.

Fortunately, Qin et at. (ref. 6) ignited recently the effort of hydrogen storage in multiwalled (MW) CNTs, where

the MWNTs were synthesized by catalytic decomposition of a methane gas on a Ni catalyst/La203, They obtained a

high discharge capacity of 200 mAh/g as well as high rate capability and long charge-discharge cycle life of 100

charge-discharge cycles with the capacity loss of less than 12 %. They also measured cyclic voltammograms (CVs)

with different sweep rates which proved clear redo× reactions in the MWCNTs. Recently Rajalakshmi et at. (ref. 7)

also obtained the net reversible discharge capacity of hydrogen with SWCNTs of up to 800 mAh/g, where the

capacity was recalculated by taking account of the amount of the SWCNTs (about 10 wt%) contained in the sample.

In this report, we have investigated key factors for electrochemical hydrogen storage in SWCNTs to determine

storage capability by controlling the amount of transition metals during synthesis and post-annealing process. The

influence of morphology and erystallinity of SWCNTs on the storage capacity is also studied.

EXPERIMENTAL

NASA/CP--2001-210948 669 August 6, 2001
SWCNTsaresynthesized byDCarc-discharge underabetiumpressltreof100torr.Agraphiterod(diameter:

dram),anaccede, withaconcentricbole(diamete,: 4ram)isfilled_ithgraphitepowdermixed wi_b 20 t_r 5 wt%

of Ni, Co, and FeS (l:1:1), respectively. The graphite rod is are-discharged at 22 V and 55 A (ref. 8). Another

graphite rod is used as a counter-electrode. We limit the weight percenl of the transition metals to 5 wt% (lower

timit) and 20 wt% (upper limit) during tbe discharge in order to investigate the effect at'transition metals on the

hydrogen storage capacity. Although the yield might be lowered with large amount of transition metals, uniform

distribulion of the transition metals in the sample may enhance the conductivity of lhe electrode anti eventually

increase the storage capacity. The samples here are collected from two different parts. The one is web-llke

SWCNTs (called as web from now oo) obtained on the wall of the ch_nber, and the other is _aken from the

cathode part (called as cathode),

To make the electrode in a form of pellet, two parts of SWCNTs powder with different amount of trmtsitioo

metals are selected m first and ground with a mortar. A blnder_ PVA (pnly vinyl alcohol), and SWCN'Is p_;wder ia

the ratio of !:l are mixed and casted in a pellet with a pressure of 1000 psi. The pellet is then dried at 100 °C for

12 hours to get rid of the remaining moisture in the vacuum oven. I_ is heat-treated in Ar atmosphere m evaporate

the binder (PVA), which is easy to diss(_lve in water, and to improve crystallini b' at" SWCNT_. However, i1 sbnuld

be preheated at 300°C in Air tbr an hour to prevent the pellet fi'om being cracked during baking in Ar atmosphere.

As _he last step for preparing the working electrode of a celt, the working electrode is immersed in 6 M Knit for

t2 hours to react with water during charge-discharge test and CV test.

A unit cell tbr the electrochemical storage of hydrogen is fabricated with SWCNTs electrode as working

electrode and nickel hydroxide as a counting electrode. The unit ceil is soaked in 6 M KOH aqueous solution as an

electro}_e and separated by a thin polymer.

The elecirochentieal measurement with the unit cell is carried t_ut using WBCS-3000 galvanostat.pntentiostat

model, with the help of a t_o-electrode syst_n using SWCNTs as a _orking electrode. Charge-dicharge with

constant current and CV with constant voltage during test is used for collecting data reported here.

RESULTS AND DISCUSSIONS

The SWCNTs grown by DC arc-discharge usually contain amorphous carbons and transition metals

sltnoundecl by amorphu_s carbons, l he SWCNT bundles are shown in Fig. l(a) with a diameter of a few tenths nm

and with a length of a few hanth'eds _m,, where the dimneter otml individual SWCNT is 1.4 rim. The web sho_s

better quality, than the cathode, l-lowever, we choose the cathode for the electrochemical storage, since SWCNTs

from this part usually contain more transilion metals compared to the _eb, which plays an important role in

enhancing the slorage capacity ot'lbe secot_dary haltery and this wilt be discussed later.

it is noted that the cathode witb 5 wl% of transition metals in Fig. l(b) has higher yield than that with 20 wt%

in Fig, I(c). The conducting materials contained in the pristine SWCNTs is found to be around 9 _%, 12 wt%,

and 31 wt%, which were measured by the thermogravimet_'ic analysis (TGA) tbr _he sample (a), (b)_ and (c),

respedi rely.

,

T............ 1_1,. _

F!GURE 1. Scanning electron microscope (SEM) images of as-prepared SWCNTs : (a) SWCNTs (web) with 5

wt% trat_sition melals, (b)SWCNTs (cmht_de) with 5 wt% transition metals, and (c) SWCN'Fs (cotbode) whb 20

v_t% lrat_sitlon melMs.

The SWCNT electrodes are tested electrochemically under static current of 50-,10f) mA/g and the lower cutoff

w_llage or" 0.4 V tier a reference tes_ _ith above, lhree smnptes. It shows lhat the pellet with large amount _t

NASA/CP--2001-210948 670 August 6, 2001
lransitk_nme_als, sample(c),canreservemosthydmge_uptoaround140mAh/gatcut=off voltage of 0.4 V. as

sh_wn in Fig. 2(c), where the mass introduced here implies the total weight of electrodes from SWCNTs,

amorphous carbons, and transition metals. Theoretical estimation on SWC,WFs shows that the metallic SWCNIs

become insulator with hydrogen adsorption tret. 9, t0). This suggests thai the additional conducting materials

compensate the electron transport in the electrode and eventually increase the storage capacity. The maximum

capacity down at zero vohage reaches 450 mAh/g. Optimization of the maximum storage capacity wilh the amounl

of transition metals is further necessary in the timwe

,_ _;_ ,'--' ' _......

spe¢i¢'_ cap_.c,{v (mA_Jgl

Figure Z. The charge--discharge tests of the as-prepared sample of (a) SWCNTs (web) with 5 wt% transition

metals, (h) SWCNI's (cathode) with 5 wt% transition metals, and (e) SWCNTs (cathode) with 20 wt%

transition metals.

h_ _rder to remove tire PVA binder from the electrode, we heated the sample by two steps. We first preheated

the sample in air at 300 C tbr five hours. This oxidizes the PVA, The degree of oxidalion is als_ importat_l, as

shown in Fig. 3(a). The plateau region in the 'collage extends with blcreasing anne',cling hours. 'Ihus a complete

oxidation is importani in order to remove the PVA more effectively in the second step annealing. This preheat

treatment further prevents the sample from being cracked by the second step aunealing. In the second-slep

annealing, most of PV,ks are burned and some are transformed into carbon particles. This high temperature

mmealing also improves the crystal|inity of the pristine SWCNTs. We annealed the sample at different temperature

in Ar atmosphere for 5 hours. The elevation of annealing temperature in Ar extends the plateau region, as shown

in Fig. 3(b). At temperature greater than 900 °C in Ar atmosphere, the electrode is cracke& resulting in the

decrease of the storage capacity of hydrogen and the plateau regi_nl. We emphasize that lhe preheat treatment is an

essential step in order to remove the PVA binder.

,_,i(a)!

._ -4 ...... ?g- ..... ?.2...... :;

t JI!_llo_y Im Art_,) Otpac_'_ lmA_l

FIGURE 3. The specific charge.discharge curves in terms of (a) the annealing time in air with the sample

coolainlng 20 wt% transition otetai_ and (b) lhe a_mealing temperature in Ar with the sample containing 5

wt%.

NASA/CP--2001-210948 671 August 6, 2001
Figure4 shows the SEM images taker_ before and after the heat treatment of SWCN_f grown with 20 w1% of

transition metals. We clloose tbe preheat trealment at 300 "C in air tbr tlve hours and then the sec:_nd step

amxealing at 900 "C in Ar _r five houls. 'Ihe SEbl images and Raman spectroscopy evidenlly show thal the heat

treatmem enhanced the erystal]inity of the SWCNTs.

.,5" ": '._ _ ".

g . ,. :;:..

,

m

(e)

.', .', ¢, d. ,,', ,,'; ;% 2.

Figure 4. SEM figure of S3_'CNTs with 2_ wt% transitbm metals of (a) the pristine sample, (b) al_er pre-heat

lreatment at 300 °C in air tbr llve hirers fallowed by tbe heat treatment at 900 °C io Ar atmosphere tbr 5

hours, and (e) Raman spectra with respect to Fig, 4(a) and Fig. 4(b),

Therntogravimeo'ic an',dysis (TGA) :lluslrates that tbe beal treatment Noeess gels rid _ff moorphoos carbons

and finally around 40 _t% of trmtshion metals rmnain in the elemrode emnpared to rite pristine sample with

around 30 wt% of transition metals as sho_m in Fig 5. The amoum of bydr,gen storage in the transition metals

and amorpbo_s carbons are negligible (ret_ I). One may delermine rite compositions of SWC_Ts and (abe] parts

by integrating the dig'erenliated TGA curve. We esliomate rite hydrogen slorage Cal',ad_y from pore S%'C,_Ts to

reach up to 450 mAh/g with a large plateau region at a cut-off voltage of 0.8 V, corresponding to 1.63 wt% of

hydrogen, a_ shown in Fig. 6(a).

- ......................... . .... . _ ¢*za

; : ....... ' "Q.:a i0_

temp_ralur_ ¢ *C_

................................ :ctc

.... ::: !i (b) _,',_

i x,,....

U.mger_t,trc [ "(t)

Figure 5. Thermogra_imetrie analysis (TGA)of cathode with 20 wt% transition metals of (a) tbe raw sample,

(b) after heat treatment in open air and Ar for 5 hours respectively.

We emphasize that 450 mAh/g corresponds to the plateau region, which is applicable te pra_icat seconda)y

battery. CVs with different sweep rates are shown in Fig: 6(b). CVs are widety used to investigate the

electrochemical process of the electrode. Figure 6(b) shows tbat tbe adsorption of hydrogen into the working

electrode is observed between t.33 X, and 1.37 V, while the desorption of hydrogen _mr, the working electrode

takes place between 1.02 \,' and 1.08 V. l_he peak position of the current [br adsorption and desorpfion is

proportional Io the sweep rate, even if the specific storage capaci b' of the electrode is not changed.

NASA/CP--2001-210948 672 August 6, 2001
1.o.

i(a)

,_.: ,% o_O

f

Figure 6. (a) Charge and discharge p_file of the optimized annealed sample with --40 wt% of SWCNTs and

(b) cyclic voltammogram test of the sample. The adsorption and desorption voltages of hydrogen of both

graphs agree well with each other.

SUMMARY

We have invesl;gated key tkcmrs far eJectrocbemical bydmgen storage in SWCNTs m determine storage

capability by c_ma'olling the amount ef trans;tion metats dm'ing synthesis and post-aone-,ding process. The

influence of morphology and crystal linity of SWCNTs on the storage capacity is also studied. After heat treatment

in mr and Ar atmosphere, it is m_ted that the trao_i6on metals in the sample increased with increasing the time of

heat treatmo,.t, and the amount of SWCNTs m:d eryst-,dlintty was also m_proved. On 1he r_thm band, the remained

amorphous carbon decreased, TGA, gEM and Ramm,. spectroscopy demonstrated these results. Eventuafly,

SWCNTs can slate hydrogen electrochemically up to 450 mAh/g corresponding to 1.63 H-wt%. The CV

measorement also showed des_;q_tion peak and adsorplioo pc',& near 1.05 V and 135 V during discharge anti

charge process respectively, %hich was in toed agreement with the plateao ef the charge-discharge experiment.

ACKNOWLEDGMENYS

'fhis p_',!iect _as st_plmrled by the MO_T through NRL program and in parl b_¢BK21 pragrmn. The amtmrs

are gratefid to Kwang-Ju branch of Korea Basic Science Institute for 4he asst._tamze '- .... _,_FE-SEM.

REFERENCES

The correspomiing anther: '(nag I-tee Lee, Deparlmenl ofphysics_ Stmgkyt,nkwan k;niversity, Suwon 440-

• " k • _ 0 " 746, Kyunggi, South Korea, e-mail: _o_w.,_:_'_in:.skkt,.a¢.k,'. fax: _ 8_-.,l-2.0-59.,4. phone: ,82-31-

290-?053

1. Dilion, A, C. el ol : Storage of Hydr,,_gen in Single-Walled Carbtm Nanotubes. Nature: eel. 386, Mm_,:h

t 997, 377.

2. Li,_, C el a/.: H:¢dn,gen Storage ;n Single-Walled Carh_m Nanotubes at Room Temperature Science, eel

280, Noveml;e, 1099, , ,a..

3. Ye, Y. et al.: H>drogen Adso%,tion and Cohesi,e E_zergy of Sing;e- Walled Cm bun Nanohtbes. Appl Phys

Lelt., eel. 74, no. 16, April 1999, 2307.

4. Na_eaabel. C. et al. : Electn)chemical Slorage of Hydrogen in Nanoutbe Materials. Electrochem. SolidState Lett., vt_l..2_ m_. ,,_ 1099, 30.

5. Lee, S. M. et M.: _ydrogen Adsorption and Storage in Cm'bon Nm_otubes. Syrthetiv Metals_ voi. 113,

2000, 209.

6. Qio. X. e_ a/• l_lectrocbemical Hydrogen Storage _ffMultiwalled Carbon Nanombes. Electrochem. Solid

Staie l_,elt., vet. 3, 2060, 532.

NASA/CP--2001-210948 673 August 6, 2001
7. Rajalakshmi, N.et at.: Electrochemical Investigation of Single-Walled Carbon Nanotubes for Hydrogen

Storage. Electrochemica Acta, vol. 45, 2000, 4511.

8. Park, Y. S. et al. unpublished.

9. S.M. Lee, et aL: Hydrogen Storage in Carbon Nanotubes. accepted to Synthetic Metals, July 2000.

10. S. M. Lee, et aL: A Mechanism of IIydrogen Storage in Single-Walled Carbon Nanotubes. submitted to J.

Kor. Phys. Sot., Jan 2001.

NASA/CP--2001-210948 674 August 6, 2001
Proceedings oftheSixthApplied Diamond Conference/Second Frontier Carbon Technology Joint Conference (ADC/FCT 2001) Y.Tzeng, K.Miyoshi, M.Yoshikawa, M.Murakawa, Y.Koga, K.Kobashi, G.A.J. Amaratunga (Editors)

HIGH CAPACITANCE SUPERCAPACITORS USING SINGLEWALLED CARBON

NANOTUBE AND POLYPYRROLE HYBRID ELECTRODES

Kay Hyeok An, Jeong-Mi Moon, Won Seok Kim, Young Min Shin, and Young Hee Lee*

Department of physics, Institute of Basic Science, Sungkyunkwan University, Suwon 440-746, South Korea

Young Soo Park, Kwan Ku Jeon, Hee Jin Jeong, Seoung yol Jeong, Seong Chu Lim, and Dong Jae Bae

Department of Semiconductor Science and Technology, Jeonbuk National University, Jeonju 561-756, South Korea

ABSTRACT

Singlewalled carbon nanotube (SWNT) and Polypyrrole (Ppy) hybrid electrode was fabricated to improve the

specific capacitance of the supercapacitor. To characterize the SWNI'-Ppy hybrid electrodes, charge-discharge

cycling test for measuring specific capacitance, cyclic voltammogram, and AC impedance test were executed. The

individual nanotubes were uniformly coated with Ppy by in-situ chemical polymerization of pyrrole. In comparison

to the respective pure Ppy and SWNT electrodes, The SWNT-Ppy hybrid electrode showed very high specific

capacitance by 5 ~ 10 times of that of individual component. We obtained a maximum specific capacitance of 265

F/g from the SWNT-Ppy hybrid electrode containing 15 wt% of conducting agent. In cyclic voltammogram data of

the SWNT-Ppy hybrid electrodes, we observed both Faradic redox behavior from the Ppy and an electric double

layer behavior from the SWNT, indicating that the type of a supercapacitor using the SWNT-Ppy hybrid electrodes

was combination of an electric double layer type and Faradie redox type.

Keywords: carbon nanotube, polypyrrole, supercapacitor, hybrid electrode.

INTRODUCTION

Recent applications of carbon nanotube (CNT) as an electrode material for supercapacitor have ignited

significant worldwide investigation on their microscopic and macroscopic porous structures and electrochemical

behaviors (refs. 1 to 10). The CNT electrodes exhibit unique pore structure and high usage efficiency of specific

surface areas (rcf. 1, 3). The CNT electrodes tbr the electric double layer type of supercapacitor have excellent

absorption characteristics due to the accessible mesopores formed by the entangled individual CNTs (ref. 1, 6).

However, there is a limitation to widen an effective surface area of the CNTs, although some processes for breaking

and solving the bundle structure of the CNTs may enhance the utilization efficiency of their surface area.

In a previous paper, we have reported that supercapacitors using singlewalled carbon nanotube (SWNT)

electrodes have a maximum specific capacitance of 180 F/g and a measured power density of 20 kW/kg at the

energy density of 7 Wh/kg in a solution of 7.5 N KOH (ref. I, 7). We have also revealed a possibility of the carbon

nanotubes (CNTs) application to the supercapacitor. Yet there are plenty of rooms to improve the performance for

the supercapacitor using the SWNT electrodes. However, the specific capacitance of supercapacitors using CNT

electrodes is smaller than that of the supercapacitors using activated carbons and activated carbon fibers (2000

3000 m2/g), which are commonly used as electrode materials for supercapacitors. Therefore, it is necessary to

reform the SWNTs for tile application to have high capacitance.

A supercapacitor is an electrochemical energy device with high power density, which could be used in

applications such as pulse power devices and electric vehicles, although the energy density of supercapacitor is

smaller than that of the secondary battery. This small specific energy density of supercapacitor is a serious

drawback to be applied to the practical devices, and consequently recent the main interest of supercapacitors is

their capability to deliver high specific capacitance and energy density. In the energy storage by supercapacitors,

there are two fundamentally different approaches: (1) for the electric double layer type of supercapacitor (EDLC),

the charge storage process is non-Faradaic, i.e., ideally no electron transfer takes place across the electrode

interface. The charge accumulation is achieved electrostatically by positive and negative charges residing on two

interfaces separated by a molecular dielectric. Because the charging and discharging of such EDLC involves no

chemical phase and composition changes, such capacitors have a high degree of cyclability on the order of 10 s

times and a high specific power density, although the specific energy density is small. The typical electrodes of

NASA/CP--2001-210948 675 August 6, 2001
suchEDLCareactivatedcarbonsandCNTs,whichhavehighsurfacearea.(2)Inthesupercapacitor basedon

pseudoeapacitanee (redoxtypeof supercapacitor), the essential process is Faradaic, i.e., the charge storage is

achieved by all electron transfer that produces redox reaction (Faradaic reaction) in the electroactive materials

according to Faraday's law. The supercapacitors based on pseudoeapacitance have higher specific capacitance than

EDLC, due to redox reaction like a battery, however, redox reaction gives rise to high internal resistance in

supercapacitors, resulting in decrease of specific power density. The typical electrodes of the supercapaeitors based

on pseudocapaeitance are metal oxides (i.e., RuO2, IrO2, Co304) and conducting polymers (i.e., polypyrrole,

polyaniline, polythiophene)

In the present study, SWNT and Polypyrrole (Ppy) hybrid electrodes are fabricated to improve the specific

capacitance of the supercapacitor. The combination of two types, the electric double-layer and the redox system,

have been developed for enhancing the performance of supercapacitors.

EXPERIMENTAL

The SWNT powder was synthesized by dc arc discharge tinder a helium pressure of 100 torr, where a graphite

rod (diameter: 6 ram) having a concentric hole (diameter: 4ram) filled by a mixture of the graphite powder with 5

wt% of Ni, Co, and FeS (l:l:l), was discharged in a bias of 22 V and a current of 55 A (ref. 11). SWNTs were

collected from only the chamber wall. Coating of Ppy on the SWNT walls was carried out using in-situ chemical

polymerization of pyrrole monomer with SWNTs. A Ferric chloride (FeCI_. 10 mmole) and a sodium ptoluenesulfonate (CH3C6H4SO3Na, 20 mmole) as oxidants for polymerization of pyrrole were added to the

suspended SWNTs (0.4 g) in 100 mL of DI water. After ferric chloride and sodium p-toluenesulfonate were

completely dissolved, pyrrole monomer (50 mmole) was slowly added dropwise into the above SWNT suspension

solution with constant magnetic stirring for four hours at room temperature. After the polymerization was over, the

SWNT-Ppy powder formed was filtered and rinsed with DI water and methanol until the filtrate was colorless. The

prepared SWNT-Ppy powder was then dried in vacuum at 60 °C for 24 hours.

The procedure of the electrode preparation was as follows. The SWNT-Ppy powder was mixed with 5 wt% of

the suspended PTFE (po[ytetrafluoroethylene) as a binder in isopropyl alcohol. The mixture was kneaded and

rolled by a bar-coater to form a sheet whose thickness is about 250 I_n. The prepared sheet was cut to a size (3 cm

x 3 cm) of a supercapacitor electrode. The working electrode for supercapacitor was prepared by pressing the cut

sheet on Ni-foam with 120 pores per inch under 1000 psi. The prepared working electrode was then dried in

vacuum at 60 °C ['or 24 hours. A unit cell fbr the capacitor was fabricated with two SWNT-Ppy hybrid electrodes

separated by a thin polymer (Celgard) in 7.5N KOH aqueous solution as an electrolyte.

The unit cell was charged at a constant voltage of 0.9 V for 60 sec - 2 hrs, and then discharged at a constant

current density of 10 ~ 100 mA/g. To characterize SWNT-Ppy hybrid electrodes, charge-discharge cycling test for

measuring specific capacitance, cyclic voltammogram, and AC impedance test were executed. The AC impedance

of a unit cell in the frequency range from I00 kHz to I0 mHz was measured by an Electrochemical Impedance

Analyzer (Jahncr Elcctrik IM6, Germany). Cyclic voltammetric measurement was performed with a PotentioGalvanostat (EG&G 273A, USA). Cyclic voltammogram is measured by conventional three-electrode beaker cell

in 7.5N KOH aqueous solution as an electrolyte. A platinum plate (3 cmx 3 cm) served as a counter electrode and

Ag/AgCI was used as a reference electrode.

RESULTS AND DISCUSION

Figure 1 shows the field-emission scanning-electron-microscope (FE-SEM, Hitachi $4700) images of asgrown SWNTs, pure Ppy and SWNT-Ppy powder formed by in-situ chemical polymerization. The as-grown

SWNTs are randomly entangled and cross-linked, and some carbon particles arc also observed. The purity of asgrown SWNTs was roughly estimated to be about 25-30 %. From Figure l(b), one can see that the pure Ppy

synthesized without SWNT shows a typical granular morphology. The granule size of pure Ppy is about 2-3 prn.

Figure l(c) reveals the individual nanotube bundles to be uniformly coated with Ppy, indicating that in-situ

chemical polymerization of pyrrole is effectively achieved. Compared to the as-grown SWNTs, the SWNTs coated

by Ppy are very thick (300 -500 nm), and carbon nanopanicles contained in the as-grown SWNT sample are also

uniformly coated by Ppy, resulting in the increase of the particle size.

NASA/CP--2001-210948 676 August 6, 2001
Acyclicvoltammetrie masurement is helpful to understand macroscopic electrochemical surface reaction at

tile electrode of the supercapacitor during the charging and discharging process. I'heretbre, cyclic vollammetry _m

tile as-grown S\_¢NTs, pure Ppy and SWNT-Ppy hybrid electrodes is performed to investigate possible dominant

mode of electrochemical energy storage on the three types electrodes. Figure 2(a)-(c) present the cyclic

voltammetric behaviors of the as-grown SWNTs, pure Ppy, and SWNT-Ppy hybrid electrode at the sweep rate of 5

mV/s, respectively In the case of the as-grown SWNT electrode (Fig. 2(a)), lhe charging and discharging cyclic

voltamnlogram is close rectangular shape and there is no current peak caused by the redox reaction, indicating a

typical electric double layer (EDL) behavior, while the pure Ppy electrode clearly shows faradaic redox reactions,

which are observed at 0.47 V and 0.4 V witb respect to the reference electrode of Ag/AgCI, representing oxidation

and reduction processes. For the SV_rNT-Ppy hybrid electrode, we observe both a behavior of tile electric double

layer which is caused by the SWNT in the voltage range of 0 - 0.35 V and the Faradaic redex behavior which

resulted from the Ppy in the range of 0.35 - 0.6 V. We note that the current scales in Fig. 2(a) and 2(b) are much

smaller than that of the hybrid type in Fig. 2(c).

Fig. I. The FE-SEM intages of as-grown SWNT, pure Ppy, and SWNT-Ppy powder.

Thc neutral Ppy reacts with a hydroxyl ion at the surthce during charging by releasing one electron to the

electrolyte, forming a dipole layer at the Ppy surface, as shown below (Eq. (I)). This dipole layer eventually forms

an electric double layer, which will increase further the charging current, Because the eft%ctive surface area was

significantly increased by the presence of the SWNTs in lhe hybrid eleclrode, as shown in Fig.t(3), the

pseudocapacitance by the iedo._ reaction and the capacitance by EDL wele simultaneously enhanced, as shown in

Fig. 2{c). The reduction peak at 0.4 V again indicatcs the similar redox reaction to pure Ppy electrode (Eq. (2)).

ppym) + OH- --- charging -- Ppy_'OH- * ePpy+*OH-- + e--- discharging -- ppyi0} + OH....... (i)

....... t2)

0.04---- ....................... _-

,,.o](a) //,

< 0.02 J

-o-1 L.............. -

-0.02 ] t

0.0 0.1 0.2 0_] 0,4 0.5 0.6 0.7 0,$ 0.9

Vollage (V _,_. AglAg _)

0.0_

11.04

0.0_

0.02

0;Ol

(b) t

/

//'? . /:f

0:00 "\jr

-,0,O I

ii.I 0_,* 03 0.4 0,_ 0._ 6.7

+

VoRage (V vs Ag/Ag )

o.2.d....................... i- .=.t(¢) -7

i

o.is]

= o.JoI /

L. 0GSt _

c_ 0.00 4 _ ....... "\ /'

-0.10 {

_.15

o.o 0.1 0.2 0_'_ 0.4 0.5 0.b 0.7

Vonage (V vs Ag/Ag +)

Fig. 2.The cyclic voltammetrie behaviors of (a) the as-grown SWNTs, (b) pure Ppy, and (c) SWNT-Ppy

hybrid elerlrodes at the sweep rate uf 5 mV/s.

NASA/CP--2001-210948 677 August 6, 2001
Figure3showsthespecificcapacitances oftheas-grownSWNTs, purePpy,andSWNT-Ppy hybridelectrodes

asafunctionofdischarge current density, wherethedataweretakenfromthesamples chargedat0.9Vfor10min.

Incomparison tthepurePpyandas-grownSWNTelectrodes, theSWNT-Ppy hybridelectrodes showsveryhigh

specificcapacitance by5 ~ 10times,althoughtespecificcapacitance ofSWNT-Ppy hybridelectrodecreased

graduallywithincreasingdischargecurrentdensity. Thisvery large specific capacitance seems to be caused by

uniformly coated Ppy on the SWNTs, signifying that active sites on Ppy chains are increased. We emphasize that

the capacitance by EDL is also enhanced by the enlarged surface area of Ppy.

250-

_,_ 200-

150-

¢o

t00-

N

50

0

0

e_'o-_ SWNT-Ppy hybrid

_'_-O_ O _Q

pure Ppy

A\ as-grown SWNT

\,

2'0 ' 40 60 80 100

Discharge current density (mA/g)

Fig. 3. The specific capacitances of the as-grown SWNTs, pure Ppy, and SWNT-Ppy hybrid electrodes as a

function of discharge current density at a charging voltage of 0.9 V for 10 min.

The specific capacitancc of the as-grown SWNT electrode slightly increases and saturates to the maximum

value with increasing discharge current density, while both of pure Ppy and SWNT-Ppy hybrid electrodes show a

gradual decrease in the specific capacitance with increasing discharge current density. For pure Ppy electrode, the

specific capacitance at a discharging current density of 100 mA/g drops by about 95 % of that at 10 mA/g, while

the capacitance of SWNT-Ppy hybrid electrode drops only by about 25 % even for large discharging current density.

The deterioration of specific capacitance of pure Ppy electrode is caused by the large internal resistance of the

electrode compared to that of SWNT electrode. The large internal resistance (or equivalent series resistance, ESR)

causes a large IR drop at high discharging current density, resulting in the decrease of the specific capacitance and

power density. The magnitude of the ESR can be clearly shown in the complex-plane impedance plots, as shown in

Fig. 4.

10

8

6

(a)

v

as-grown SWNT

SWNT-Ppy hybrid

O_o/O/s

1200

1O00-

800-

v 600-

(b)

400-

200-

OJ

0

j&/--

//

PPY /

&:,-

/,

/

&:

2O0

• t • L • i , , •

10 400 600 800 1000 1200

z'(n) z'(n)

8-

7-

6-

5-

4-

2-

1-

0  ̧

0

(c) Sup-P: 0 wt%

_. Sup-P: 5 wl%

• Sup-P: 10 wt%

• Sup-P: 15wt%

53 Sup-P: 20 wt%

&

_j 4: :4::

2 3 4 5 6 7

z' (_)

Fig. 4. The complex-plane impedance plots for (a) as-grown SWNTs and SWNT-Ppy hybrid electrodes,

(b) pure Ppy, and (c) SWNT-Ppy hybrid electrode containing 5 ~ 20 wt% of the conducting agent

NASA/CP--2001-210948 678 August6,2001
Figure4presentsthecomplex-plane impedanceplotsfortheas-grownSWNTs, purePpyandSWNT-Ppy

hybridelectrodes inthefrequency rangeof from100kHzto I0mHz.Theimpedance behaviorftheas-grown

SWNI'electrodeis closeto anidealEDLC.Accordingly, theas-grownSWNTelectrodeshowsmuchlower

impedance thanthepurePpyandSWNT-Ppy hybridelectrodes, givingrisetoasmallIRdropfortheas-grown

SWNTelectrode. Theinternalresistanceof thepurePpyelectrodesis verylarge,resultingin thelarge

deterioration ofthespecificcapacitance atparticularlylargedischarging currentdensityof 100mA/g.Although

theinternal resistance oftheSWNT-Ppy hybridelectrodeismuchsmallerthanthat of the pure Ppy electrodes, the

internal resistance of the SWNT-Ppy hybrid electrode is still large compared to that of the as-grown SWNT

electrode. In order to reduce internal resistance of the SWNT-Ppy hybrid electrode, 5 - 20 wt% of the conducting

agent (Sup-P) is added into SWNT-Ppy hybrid electrode, when the electrodes are prepared. The complex-plane

impedances are shown in Fig. 4(e). With increasing the amount of conducting agent, the phase angle of tan (R_ma_

/ Rf_z) at low frequency region is close to 90 degree, indicating a significant reduction of the internal resistance in

the SWNT-Ppy hybrid electrode. The internal resistance of the SWNT-Ppy hybrid electrode with the conducting

agent above 15 wt%, become smaller than that of the as-grown SWNT electrode. These suggest that the conducting

agent obviously affects the specific capacitance of the supercapacitor in the hybrid electrode.

270-

260-

"_ 250-

_a

240

230-

220-

210-

o 200-

190-

e_

180-

170

O

/ ,\

,/ \\

10 mA/g /_ •

• / \ ,

0 _0 /' "O

O_ \/

/

/

/ \\

// ,• 100 mA/g e'

/

//

/

./

0 5 10 15 20

Addition amount of Sup-p (wt%)

Fig. 5. The effects of conducting agent on the specific capacitance of the SWNT-Ppy hybrid electrode as a

function of discharging current density at a charging voltage of 0.9 V for 10 min.

The effects of conducting agent on the specific capacitance of the SWNT-Ppy hybrid electrode are shown in

Fig. 5. The maximum specific capacitance appears at 15 wt% &conducting agent in electrode, where the data

were taken from the samples charged at 0.9 V for 10 rain. The effect of the conducting agent is clearly seen

particularly at high discharging current density. The specific capacitance significantly increases with increasing the

amount of the conducting agent up to 15 wt% and drops at 20 wt%. We obtain a maximum specific capacitance of

265 F/g from the SWNT-Ppy hybrid electrode containing 15wt% &the conducting agent. This value is larger than

the previously reported value of 113 F/g from multiwalled carbon nanotubes (ref. 5) and our previous record of 180

F/g from the heat-treated SWNT electrode (ref. 1). The conducting agent itself does not directly contribute to the

specific capacitance but by reducing the IR drop. However, an excessive addition of the conducting agent

diminishes the effective portion of the active electrode material, resulting in the reduction of the specific

capacitance. The difference of a specific capacitance between discharging current density 10 mA/g and 100 mA/g

is larger at small amount of conducting agent due to the larger internal resistance.

SUMMARY

We have introduced the SWNT-Ppy hybrid electrode to improve the specific capacitance of the supercapacitor

by combining the EDL and the redox reaction. In comparison to the pure Ppy and SWNT electrodes, the SWNTNASA/CP--2001-210948 679 August 6, 2001
Ppyhybridelectrodes showsveryhighspecificcapacitance by5 - I0times.Wereveal thathisverylargespecific

capacitance iscausedbyuniformlycoatedPpyontheSWNTs,ignifyingthatactivesitesonPpychainsare

increased. Weemphasize thatthepseudocapacitance bythredoxreactionandthecapacitance byEDLwere

simultaneously enhanced bytheenlargedsurfaceareaofPpy.Weobtainamaximumspecificcapacitance of265

F/gfromtheSWNT-Ppy hybridelectrodecontaining15wt%oftheconductingaent. The addition of conducting

agent into the SWNT-Ppy hybrid electrode effects increasing specific capacitance by reducing internal resistance of

supercapacitor.

ACKNOWLEDGEMENTS

This project was supported by the Korean Ministry of Science and Technology (MOST) through National

Research Laboratory (NRL) program and in part by Brain Korea 21 (BK21) program.

REFERENCES

* The corresponding author: Young Hee Lee, Department of physics, Sungkyunkwan University, Suwon

440-746, Kyunggi, South Korea, e-mail: leeyoung@yurim.skku.ac.kr, fax: +82-31-290-5954, phone: +82-

31-290-7053

1. An, K.H., et al.: Supercapacitors Using Single-Walled Carbon Nanotube Electrodes, Advanced Materials,

vol. 13, no. 7, April 2001,497-500.

2. Frackowiak, E.; 136guin, F.: Carbon Materials for The Electrochemical Storage of Energy in Capacitors,

Carbon, vol. 39, no. 6, 2001, 937-950.

3. Ma. R.Z., et al.: Carbon Nanotube Electrodes for Fabricating Supercapacitors, Journal of Tsinghua

University, vol. 40, no. 8, August 2000, 7-10.

4. Ma, R.Z. et at.: Study of Electrochemical Capacitors Utilizing Carbon Nanotube Electrodes, Journal of

Power Sources, vol. 84, no. 1, November 1999, Pages 126-129.

5. Niu, E.K. et al.: High Power Electrochemical Capacitors Based on Carbon Nanotube Electrodes, Applied

Physics Lettcrs, vol. 70, no. 11, March 1997, 1480-1482.

6. Diederich, L., et al.: Supercapacitors Based on Nanostructured Carbon Electrodes Grown by Cluster-Beam

Deposition, Applied Physics Letters, vol. 75, no. 17, October 1999, 2662-2664.

7. An, K.H., et al.: Electrochemical Properties of High Power Supercapacitors Using Single-Walled Carbon

Nanotube Electrodes, Advanced Functional Materials, 2001, in press.

8. Peigney, A., et al.: Specific Surface Area of Carbon Nanotubes and Bundles of Carbon Nanotubes, Carbon,

vol. 39, no 4, 2001, 507-514.

9. Bacsa, R.R., et al.: High Specific Surface Area Carbon Nanotubes from Catalytic Chemical Vapor

Deposition Process, Chemical Physics Letters, vol. 323, no. 5-6, June 2000, 566-571.

10. Eswaramoorthy, M.; Sen, R.; Rao, C.N.R.: A study of Micropores in Single-Walled Carbon Nanotubes by

the Adsorption of Gases and Vapors, Chemical Physics Letters, voL 304, no. 3-4, April 1999, 207-210.

1I. Park, Y.S., et al.: Low Pressure Synthesis of Singlewalled Carbon Nanotubes by Arc Discharge, Synthetic

Metals, 2001. In press.

NASA/CP--2001-210948 680 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

HIGH-PRESSURE POLYMERIZATION

OF SINGLE WALL CARBON NANOTUBES.

M. Popov, M. Kyotani and Y. Koga

Joint Research Consortium of Frontier Carbon Technology, JFCC, c/o NIMC,

Higashi, Tsukuba, Ibaraki, 305-8565, Japan.

R. J. Nemanich

Department of Physics, North Carolina State University, 408A Cox, Box 8202, Raleigh, NC, 27695-8202, USA.

ABSTRACT

A new superhard material composed of polymerized single wall carbon nanotubes (P-SWNT) has been

synthesized which exhibits a bulk modulus exceeding or comparable with diamond and hardness exceeding or

comparable with cubic BN. The P-SWNT were synthesized by applying a shear deformation under load in a

diamond anvil cell; the procedure of stress tensor variation. After intermediate phase transitions at 14 and 19 GPa,

single wall carbon nanotubes are transformed to P-SWNT at a pressure of 24 GPa. The process of polymerization is

accompanied by irreversible changes in the Raman spectra. The nanotubes do not collapse at least up to pressure 55

GPa (maximum pressure of the study). Bulk modulus of 462 to 546 GPa was found out for P-SWNT from the

comparative study of pressure dependence of the Raman modes of P-SWNT (high energy mode -1590 cm -1) and

diamond (1333 cm-1). This value exceeds the bulk modulus of diamond (420 GPa for single diamond crystal). PSWNT are characterized by high value of hardness (62 to 85 GPa) exceeding or comparable with cubic BN.

Hardness measurements were performed using nanoindentation technique.

Keywords: single wall carbon nanotubes; superhard materials; high pressure; hardness; Raman scattering.

INTRODUCTION

Graphite sheet properties determine the interesting mechanical, electrical etc. characteristics (e.g. sp 2 bonding in

the graphite sheet is stronger than sp 3 bonding in diamond). Creation of superhard carbon nanocluster-based

materials may be provided by possibility of the forming of sp 3 bonding between curved graphite (sp 2 bonded) sheets.

Thus, theoretical and experimental studies give the value of Young's modulus for single wall carbon nanotubes

(SWNT) (ref. 1) in the tera-Pascal range (ref. 2, 3). Polymerized SWNT (P-SWNT) via sp 3 inter-tube bonding are

expected to have highly anisotropic mechanical and electrical (metallic and semiconductor) properties while

exhibiting superhard material characteristics. Possibility of polymerization of nanotubes and different P-SWNT

structures were studied theoretically in ref. 4.

Polymerization of nanotubes and transfer of mechanical properties of single nanotube into P-SWNT is possible

by analogy with ultrahard fullerite (ref. 5 to 7), since the nanotube is a graphite sheet curved into a cylinder, while

fullerene is that curved to sphere. The striking mechanical properties of ultrahard fullerite (ref. 8) with hardness,

wear and bulk modulus exceeding diamond are attributed to the properties of the C_0 molecule (ref. 9) and the

possibility of sp 3 intermolecular bonding (ref. 10) which leads to the 3D fullerite polymer (ref. 5 to 7).

The process of formation of sp 3 bonding between graphite sheets at room temperature is sensitive to the stress

tensor in the specimen. Under hydrostatic conditions graphite does not transform to diamond at least up to 80 GPa

(ref. 11), while stress tensor variation (by applying shear deformation under load) leads to direct transformation of

graphite to diamond at a pressure of 17 GPa at room temperature (ref. 12). Success in ultrahard fullerite synthesis

using stress tensor variation in a shear diamond anvil cell (SDAC) demonstrates the advantages of this procedure

(ref. 5). At room temperature 3D polymerization of C_0 via sp 3 intermolecular bonding begins under non-hydrostatic

pressure at 18 GPa (ref. 5). In the present study this procedure was successfully applied for synthesis of P-SWNT.

EXPERIMENTAL DETAILS

The SWNT bundles used in this study were obtained from Tubes@Rice (Rice University in Houston, TX). The

nanotubes are composed of a nearly random mixture of arm-chair, zig-zag and chiral helicities. The mean tube

NASA/CP--2001-210948 681 August 6, 2001
diameter of1.2nmisobtainedbythemanufacturing procedure(ref.13)andisappropriates forourRaman

measurements.

InSDACcontrolledsheardeformationsappliedtoaspecimenunderpressurebyrotationofoneofanvils

aroundanaxisofload.Theapplicationfsheardeformation(stresstensorvariation) decreases thehysteresis ofa

phasetransformation andmakesitpossibletoobtainahomogeneous phase(seeforexamplerf.5andreferences

therein). Atthephasetransition, astep-likeanomalyoftheradialpressuredistributioninthesample(duetovolume

andelasticmodulusjump)appearsforthecaseofnon-hydrostatic compression. Accordingtotheprocedure

describedinref.5,theseanomalies weredetectedforpressures of14+1,19+1and24+1GPa.Thepressurewas

measured bytheruby-fluorescence procedure. Rubyparticles(1to3gminsize)weredepositedontotheworking

surfaceofthediamondanvil.Specimens wereloadedinthegasket withoutapressuremedium.

ThesamplestatewhileunderpressurewasregisteredbyRamanscattering. TheRamanspectrawererecorded

usingaJascoNRS-2100triplematespectrometerquippedwithacharge-coupled detector, Ramanmicroscope and

excitedbythe514.5nm(2.41eV)lineofanargonionlaser.Unfortunately, the1344cm-1bandofthenanotubes are

obscuredbytheRamanscatteringfromthediamondanvilsforpressureupto15GPa.Undermorehighpressureit is

difficulttoseparateshifted1594and1344cm-1bandswithhighenoughprecisionbecauseoflinesbroadening and

lowintensityof1344cm-1band.

Pressuredpendence ofdiamondRamanbandwasmeasured forcomparison. First-order Ramanspectrawas

collectedfrom3gm- sizeregionofdiamondanvilclosetothesample. Theexperimental procedureisdescribedin

detailnref.14.

Hardnessmeasurements ofP-SWNTafterpressuretreatment wereperformedusingNanoindentation System

(MTSSystemsCorporation) equippedwithBerkovichdiamondindenter. Becauseofsmallthicknessofsample

(about20gm)andnumerous cracksaroseinthesampleuponpressurerelease, thehardness wasmeasured at

penetration depthabout40nm.Toincreasevalidityandprecisionofdata,cubicBNfaces(100)and(111)were

measuredatthesamepenetration depthforcomparison. Force-depth indentationcurvesofP-SWNT, cubicBNfaces

(100)and(111)andfusedsilicaareplottedinFig.1.Thepossibilityandpeculiarities forcorrect hardness

measurements ofsuperhardmaterialsatthissmallpenetration deptharediscussed inref.15.

3.3

3.0

2.7

2.4 Z

2.1

1.8

E

1.5

tO 1.2

0 0.9

d

0.6

0.3

0.0

0 5 10 15 20 25 30 35 40 45

Displacement into surface (nm)

Figure 1. Force-depth indentation curves of P-SWNT (marked by solid circles), cubic BN faces (100)

(marked by hollow circles) and (111) (marked by crosses) and fused silica (marked by hollow triangle) are

plotted.

NASA/CP--2001-210948 682 August 6, 2001
RESULTS AND DISCUSSION

Polymerization

By analogy with fullerene C_0 (ref. 5 and 7) the pressure-induced polymerization of the SWNT should lead to

broadening and re-distribution of bands intensities in the Raman spectra. Shown in Fig. 2 are spectra appropriate to

virgin SWNT (Fig. 2a) and to P-SWNT (Fig. 2b). The spectra of P-SWNT after unloading do not depend on

pressure treatment (24 to 55 GPa). Farther to the assumed polymerization it is possible to conclude from the Raman

spectra that the nanotubes do not collapse under pressure of at least 55 GPa: since the SWNT band positions are not

changed. The peculiarities of spectra also differentiate it from other known types of graphitic and amorphous carbon

materials (ref. 16). Because of bands broadening, instead of 4 bands in high frequency region for SWNT (at 1551,

1568, 1592 and 1598 cm -1, Fig. 2a) it is possible to resolve only 1561 and 1598 cm -1 bands for P-SWNT using

Lorentz multi-peaks fit (Fig. 2b; spectra of SWNT is plotted for comparison). These bands give maximum at 1594

cm -1 in both spectra. The most distinctive evidence of existence of nanotubes after 55 GPa pressure is presence of

low-frequency radial breathing mode of SWNT (ref. 17) at 188 cm -1 in the spectra of P-SWNT (Fig. 2c).

%-

4

cO

tt12.

6.

a %-

.m

c

4

c

............. i ...... 1............................... w....

1550 1600 1650

Raman shift cm-1

1,0-

0,5"

b

1000 15'00

Raman shift (cm -1)

160 24-0 320

Raman shift (cm -1)

2600 28'00 30'00 32'00

Raman shift (cm 1)

Figure 2. Raman spectra of SWNT (a), P-SWNT (b) (spectra of SWNT is plotted in (b) for comparison),

(c) is low-frequency region of P-SWNT (overhead) and SWNT (below).

Essential changes in the Raman spectra are associated with the 1344 cm -1 band: the relative intensities of the

intense bands (188, 1568, 1598 and second-order 2681 cm -1) are the same within 40% after the compression, while

the relative intensity of the 1344 cm -1 band increased by an order of magnitude. A decreasing of the Raman intensity

(by a factor of 2 or more) after relatively low (5.2 GPa) pressure was mentioned in ref. 18. The second order spectra

of P-SWNT and SWNT are plotted in Fig. 2 (d).

The changes of relative intensities of the bands around 1340, 1580 and 2680 cm -1 (in Fig. 2b and 2d) resemble

those observed in spectra of polycrystalline graphite and glassy carbon (ref. 16 and 19): the spectra of the first is

analogous to SWNT, while the second is analogous to P-SWNT. These changes in polycrystalline and disordered

carbon are partly attributed to the break down of light scattering selection rules (ref. 16, 19, 20 and references

NASA/CP--2001-210948 683 August 6, 2001
therein). Inviewofthedirectrelationship betweenthephonondensityofstatesforgraphiteandSWNT(ananotube

isformedbyrollingupagraphenesheet toacylinderandthatgivesthepossibilityofapplyingthezone-folding

methodforcalculationfthephononmodesofnanotubes (ref.21))itispossibletoattributehechangesinthePSWNTspectratothebreakdownofSWNTselectionrules.Theperturbation ofthenanotubestructureiscausedby

internanotube bondingasresultofpolymerization. Thepossiblemechanismofthepolymerization asmentioned

aboveissp3internanotube bonding(ref.4).Thismechanismisknownotonlyforfullerite(ref.7)butforgraphite

sheets, too(ref.22).

Bulkmodulus

Afterpolymerization, themechanical propertiesoftheSWNTwillbecharacteristic ofa3Dcarbonnetwork, and

asaresult,hebulkmodulusofthismaterial willincrease. Important informationaboutthebulkmodulusofPSWNTcanbederivedfromthepressuredpendence forthemodesinhigh-energyangeoftheRamanspectra(i.e.

thebandsaround1590cm-1). InFig.3thepressuredpendence w/w0fortheWo= 1594cm-1band(asmentioned

above, itconsistsfromseveral bands)isplotted.ThebulkmodulusBiscalculatedfromtheequation

T_=(B/Q)be)/aP (1)

where 7_ is the Gruneisen parameter for a quasi-harmonic mode of frequency _ (ref. 23). In general, the scaling

parameter 7-_ 1 for the covalently bonding group IV semiconductors (ref. 23) and particularly for graphite 7 = 1.06

(ref. 24) to 1.11 (ref. 25) and for diamond 7 = 0.96 (ref. 26). Recently, the Gruneisen parameter of 7 = 1.24 was

calculated for SWNT (ref. 27). It is important to note, that 7-_ 1 is expected only for the mode -1590 cm -1 that is

analogous to the graphite E2g(G) mode. This mode is characteristic of sp 2 bonding (ref. 20). At the same time for the

disorder induced breathing mode of graphite, the Alg(D) band at -1340 cm -1, 7 = 1.9 according to ref. 25. The

pressure dependence of this mode for the nanotubes also differs from the 1594 cm -1 band. Under pressure the mode

shifts to a frequency nearer to the characteristic sp 2 band.

#

e.o

E

rr

rr

1.08.

1.06.

o • °'. ell

db£°o°_ • * •

O O ,# 80 • O •

1.04" ;# _ ._ : •

1.00. , . . • .

0 10 20 3'0 4'0

.,,.-i'.

' 5'0 ' 6'0

Pressure (GPa)

Figure 3. The pressure dependence of the 1594 cm 1 band of SWNT (hollow circles), P-SWNT (solid

circles), diamond (stars).

In Fig. 3, the pressure dependence of the 1594 cm -1 band is plotted over the entire pressure range. Data

appropriate to SWNT and the first and second SWNT phases are marked by hollow circles. Changes in slope

(hollow circles) indicate phase transitions at pressures of 14+1, 19+1 and 24+1 GPa. The pressure dependence about

6 cm-1/GPa for SWNT 1594 cm -1 band under pressure below 14 GPa is consistent with data reported in ref. 28 for

SWNT (5.7 cm-1/GPa): the shift agrees with the pressure dependence of the high-energy modes of isolated nanotubes

under compression (ref. 18 and 28).

For determination of Om/OP and B of the P-SWNT, we have used data appropriate to sample treated by pressure

above 25 GPa and shear deformation. These data are marked by solid circles in Fig. 3. No hysteresis is seen in the

pressure dependence of P-SWNT. This is direct evidence that polymerization of the nanotubes is irreversible.

NASA/CP--2001-210948 684 August 6, 2001
Toapplyequation(1)forcorrectbulkmoduluscalculationfP-SWNT, it isnecessary tostudystresstensorin

thesample(weapplynon-hydrostatic compression). Toperformthisprocedure, weplottedw/w0pressure

dependence fordiamondRamanbandw0--1333cm-1measuredatthesameconditions. IntheFig.3appropriate to

diamondpointsaremarkedbystars.Fortunately, thepressuredpendence forthediamondRamanbandandPSWNTareverysimilarwithhighaccuracy. InthiscasebulkmodulusofP-SWNT, Bpisproportional tothatof

diamondBd:Bp--"{p/TdBd .Where yp is Gruneisen parameter for P-SWNT and y,: is diamond Gruneisen parameter.

Consequently, the bulk modulus of P-SWNT exceeds bulk modulus of diamond 1.1 to 1.3 times, depending on

choice of 7: experimental data for graphene sheet or calculated value for SWNT.

According to ref. 29 bulk module of single diamond crystal is 420 GPa. Thus, for P-SWNT bulk module is 462

to 546 GPa.

Hardness

To increase precision and validity of hardness data, a procedure of calibration against material with known

hardness (standard) was used. This procedure was elaborated in details in ref. 8 and 30.1 mm-size cubic BN single

crystals (faces (100) and (111) (ref. 31) and fused silica were used as standard materials. Hardness (H) of P-SWNT

was calculated with the relation H=P/S. P is load to the indenter and S is the projected area of contact. If we compare

indentations in different materials (with hardness H1 and H2) with the same value of S, then H1-- (Pl/P2) H2, where Pl

and P2 are appropriate loads on indenter. Results of hardness measurements of P-SWNT, c-BN and fused silica are

plotted in Fig. 1.

Hardness calibration measurements of c-BN and fused silica were performed at indentation depth about 1 btm.

The hardness values are: c-BN (111) is 66_+ 1 to 73_+ 1 GPa depending on orientation of indenter relatively to

crystallographic axis of sample (so-called hardness anisotropy); c-BN (100) is 62_+1 GPa (effect of anisotropy is not

so strong as for (111) face; fused silica is 12_+0.5 GPa.

Hardness of P-SWNT determined according to the procedure of calibration against c-BN belongs to interval 62

to 85 GPa. Load-displacement curves for P-SWNT appropriated to the confines of the range are plotted in Fig. 1.

The lower curve of P-SWNT (minimum hardness) corresponds with the curve of c-BN face (100). Hardness data

scattering of P-SWNT depends on microcracks in the sample. As the sample was formed under high pressure and

upon unloading, large internal stresses should led to cracking. Hence, the real hardness of P-SWNT is expected

higher then measured range 62 to 85 GPa.

CONCLUSIONS

A new superhard material composed of polymerized single wall carbon nanotubes has been synthesized which

exhibits a bulk modulus (462 to 546 GPa) exceeding or comparable with diamond and hardness (62 to 85 GPa)

exceeding or comparable with cubic BN.

This work was supported by the Frontier Carbon Technology Project, which was consigned to JFCC by NEDO.

We thank Dr. T. Taniguchi for samples of cubic BN and Dr. Y. Ohkubo (JASCO) for providing with preliminary

Uaman measurements.

REFERENCES

1. S. Iijima, T. Ichihashi, Nature (London) 363, 603 (1993).

2. D.H. Robertson, D. W. Brenner, J. W. Mintmire, Phys. Rev. B 45, 12592 (1992).

3. M.M.J. Treacy, T. W. Ebbesen, J. M. Gibson, Nature 381, 678 (1996).

4. L.A. Chernozatonskii, Chem.Phys.Lett., 297, 257 (1998).

5. V. Blank et al., Physics Lett. A 188, 281 (1994).

6. V. Blank et al., Physics Lett. A 205, 208 (1995).

7. V. Blank et al., Carbon 36, 319 (1998).

8. V. Blank et al., Diamond Relat. Mater. 7, 427 (1998).

9. R.S. Ruoff, A. L. Ruoft, Nature 350, 663 (1991).

10. A.M. Rao et.al., Science 259, 955 (1993).

11. A. F. Goncharov, Sov. Phys. JETP 71, 1025 (1990).

NASA/CP--2001-210948 685 August 6, 2001
12.V.D.Blank,G.A.Dubitsky, S.A.Plotyanskaya, inHigh Pressure Science and Technology. Proceedings of the

Joint XI AIRAPT and XXXIII EHPRG Intern. Conf., edited by W. A. Trzeciakowski (World Scientific,

Publishing Co. Pte. Ltd, Singapore, 1996), pp. 325-327.

13. A. G. Rinzler etal., Appl. Phys. A, 67, 29-37 (1998).

14. M. Hanfland and K. Syassen, J. Appl. Phys. 57, 2752 (1985).

15. A. Richter et al., Diamond Relat. Mater., 9, 170 (2000).

16. D. S. Knight, W. B. White, J. Mater. Res. 4, 385 (1989).

17. A. M. Rao et al., Science, 275, 187 (1997).

18. U. D. Venkateswaran et al., Phys. Rev. B 59, 10928 (1999).

19. R. J. Nemanich, S. A. Solin, Phys. Rev. B, 20, 392 (1979)

20. A. C. Ferrari, J. Robertson, Phys. Rev. B 61, 14095 (2000).

21. R. A1-Jishi, L. Venkataraman, M. S. Dresselhaus, G. Dresselhaus, Chem. Phys. Letters 209, 77 (1993).

22. V. Blank et al., Diamond Relat. Mater. 8, 1285 (1999).

23. B. A. Weinstein, R. Zallen, in Light Scattering in Solids, edited by M. Cardona, G. Guntherodt (Springer, Berlin,

1984), Vol. IV.

24. M. Hanfland, H. Beister, K. Syassen, Phys. Rev. B 39, 12598 (1989).

25. A. F. Goncharov, V. D. Andreev, Sov. Phys. JETP 73, 140 (1991).

26. M. Hanfland, K. Syassen, S. Fahy, S. G. Louie, M. L. Cohen, Phys. Rev. B 31, 6896 (1985).

27. S. Reich, H. Jantoljak, C. Thomsen, Phys. Rev. B 61, 13389 (2000).

28. C. Thomsen et al., Appl. Phys. A 69, 309 (1999).

29. V. M. Prokhorov, V. D. Blank, S. G. Buga, V. M. Levin, Synthetic metals 103, 2439 (1999).

30. V.Blank, etal., J. Mater. Res., 12, 3109 (1997).

31. T. Taniguchi, S. Yamaoka, J Cryst Growth, 222, 549 (2000).

NASA/CP--2001-210948 686 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

COVALENT SIDEWALL FUNCTIONALIZATION OF SINGLE WALL CARBON

NANOTUBES

I.W. Chiang, R. K. Saini, E. T. Mickelson, W. E. Billups, R.H. Hauge and J. L. Margrave

Department of Chemistry, Center for Nanoscale Science and Technology

Rice University, Houston, TX 77005

ABSTRACT

Progress of fluorination of single-wall carbon nanotubes is being reported. Covalent attachment of alkyl

groups including methyl, n-butyl and n-hexyl groups to the sidewalls of single wall carbon nanotubes (SWNTs) has

been achieved. Quantitative measurement of the alkylation was done by thermal gravimetric analysis. FTIR,

Raman and UV-Vis-NIR were used to characterize these alkylated SWNTs. Application of these nanotubes are

being investigated--fibers, composites, batteries, lubricants, etc.

INTRODUCTION

Among the wide range of nanometer scale structures prepared to date, carbon nanotubes and, in particular,

single-wall carbon nanotubes (SWNTs) stand out as unique materials for fundamental research and emerging

applications. Since the initial discovery in 19911 , and subsequent development of large-scale syntheses of

buckytubes, various methods for their synthesis, characterization, and potential applications have been pursued 2' 3.

In some cases, chemical modification will be required to reach full potential on nanotube application. However, this

area of research presents enormous challenges and results have been slow to emerge.. Mickelson et. al 4,5were able

to fluorinate the sidewalls of carbon nanotubes in 1998. In addition to that astonishing preparation of

fluoronanotubes, several derivatizations of carbon tubes were reported _' 7. The fluorinated materials can be further

functionalized by attaching alkane chains to the sidewall of SWNTs, as demonstrated by Boul et al. 8 In this paper,

we present an up-to-date account of recent results on the fluorination of single-wall carbon nanotubes and the use of

fluorinated SWNTs as intermediates for sidewall derivatization.

EXPERIMENTAL PROCEDURE

Fluorination

Precursor nanotubes used in this study are laser-grown-oven SWNTs from tubes@rice, or HiPco SWNTs also

from Rice University. The single-walled carbon nanotubes, about 1 nm in diameter, were produced by the dual laser

vaporization of Co/Ni doped graphite rods at tubes@rice 9 and by high pressure CO gas phase reaction 1°. After the

usual nitric acid cleaning procedures 9, laser-oven-grow tubes were subjected to a 2-stage oxidation process 11 to

further remove the metal catalysts and amorphous carbon presented in the samples. HiPco tubes were obtained asproduced from the reactor. They underwent similar multi-stage gas-phase purification procedures 12.

Several reactions were carried out to establish optimal fluorination conditions (i.e. temperatures and reaction

times). The goal is to reach saturation, C2F, without destroying the tube structure 4. A fluorine/HF/He mixture (2

cc/min, 0.2 cc/min and 100 cc/min respectively) was passed through a temperature-controlled Monel flow reactor

containing the SWNTs samples. HF gas was turned off immediately after the reaction, while fluorine gas was kept

flowing until the chamber temperature for the fluorination reactor dropped below 100°C to avoid the additional HF

adsorption onto the nanotubes.

Alkylation

10 mg of the fluorinated tubes were added to a flame dried 100 ml three necked round bottomed flask under

an atmosphere of argon. Dry hexane (5ml) was then added and the contents were sonicated for 5 min. Alkyllithium

( methyl-, butyl- and hexyl-lithium) in hexane (8 ml) was added dropwise to the resulting suspension and the

mixture was sonicated (VWR scientific model 50 HT)for 20 min. The solution was stirred overnight under argon

and then cooled in ice. The unreacted hexyllithium was quenched by the slow addition of ethanol with stirring.

After quenching, water (50 ml) was added and the solution was stirred to dissolve salts formed. Nanotubes were

then filtered on a PTFE filter paper (0.45 gm) and washed with water. Collected nanotubes were again suspended in

NASA/CP--2001-210948 687 August 6, 2001
3NHC1,sonicatedfor15minandagainfilteredandwashedwithplentyofwater.Finallythetubeswereplacedin

hotethanol andsonicatedfor15minandfiltered.Finalproductsweredriedundervacuumovernight.

Instrumentation

Thermogravimetric (TGA)datawereobtainedusingaTAinstrument model2960system.Raman, IR,

masspectraandTGAexperiments werecarriedoutwithSWNTsintheformofbuckypaper. SamplesforUV-VisNirwerepreparedbysonicatingSWNTsfor- 10 min in 0.15 wt % Triton-X D20 solution with a cup sonicator.

Raman spectra were obtained with a Renishaw micro-raman operating with a 780 nm laser. UV-Vis-Nir spectra

were obtained with a Shimadzu UV-3101 PC spectrometer. FTIR spectra were obtained using a Nicolet

spectrometer with ATR accessory.

RESULT AND DISCUSSION

Fluorination

Laser-Oven-Grown SWNTs

HF is known as a catalyst for fluorination of graphite 13. It is also suggested that with the addition of HF in

the fluorination, C-F bonds of more covalent in nature will be produced. Figure la shows the FTIR spectra of

nanotubes fluorinated with or without HF addition at 200°C. Apparent upshift of the C-F stretching frequency is

observed with HF addition to the fluorine atmosphere. A higher degree of fluorination is also observed with higher

relative C-F intensity. Some shifting effect is also seen when raising the fluorination temperature (Figure lb).

1176.0

,q

(a) _ .

I II1

; ii

(b) ,

! .',', ',il ll0S.0

4000 3500 3000 2500 2000 1500 1000

1176.0

1200.7

i (a) _] \

4000 3500 3000 2500 2000 1500 1000

Figure l(a).IR spectra of fluorinated SWNTs Figure l(b).IR spectra of fluorinated SWNTs

a. With HF addition a. 200°C

b. Without HF addition b. 250°C

1 1

*Negative absorption at 2000 cm and 2200 cm is due to the diamond crystal on the ATR accessory

Because of the nature of the nanotubes, black in color, light does not transmit through therefore we could

not take IR spectra of tubes in transmission modes. The initial FTIR spectra were taken using a diffuse reflectance

accessory. No peaks were observed at C-F bond region (-1200 cm-1). When mixing fluorotubes with KBr and

subsequently heating on the hot plate at - 120°C, one can immediately see the evolution of gas which is yellow.

KBr must have reacted with Fluorotubes: KBr + C2F --) KF + Br2 + Cn This also indicates the weakness of the C-F

bond on fluorotubes. Similar KBr tests were done on fluorinated C_0 and fluorinated graphite. While heating,

fluorinated C60 changes its color from white to yellow, indicate the loss of fluorine. C-F bonds in fluoro-graphite

are much stronger. Under the same experimental conditions, fluorine remains attached to the graphite. The same

set of experiments was repeated using NaI instead of KBr. Purple gas, which is expected to be I2, evolved during

heating fluorotubes with NaI.

Figure 2 and Figure 3 shows the FTIR and Raman spectra of defluorinated nanotubes that were originally

fluorinated under different temperatures. At temperature above 250°C, the intensity of defects/disorders, as

indicated by the 1298 cm -1 peak, is almost the same intensity of the nanotube tangential modes at 1592 cm -1. In

order to keep the tube structure and obtain a saturated C/F ratio (-2) at the same time, we decided to use 250°C as

NASA/CP--2001-210948 688 August 6, 2001
ouroptimalfuorinationtemperature. TheminimumreactiontimetoreachC2Fis-5 hoursat250°C.Electron

microprobeanalysis(EMPA)andIR indicatedthedesiredcarbon-fluorine ratioandC-Fstretchingfrequency.

ThesefluorinatedtubesaresolubleinDMFandalcohol solvents5.

i -"-

i "--x...

' r',

] (a) , ,

• I\

3500 3000 2500 2000 1500 1000

Wavenumbers (cmFigure 2.IR spectra of defluorinated SWNTs

(a) Fluorotubes (b) after defluorination

(e)

-_T................ ._...'--_......................

(d)

(c)

.................. :.....!i....................

(b)

Ca)

500 1000 1500 2000 2500

Figure 3.Raman spectra of defluorinated SWNTs

(a) Pristine SWNTs (b) T=250°C (c) T=315°C

(d) T=350°C (e) T=375°C

HiPco SWNTs

SWNTs produced by HiPco process have smaller tube diameters, 1.1 nm, when comparing to laser grown

SWNTs, 1.3 nm. Therefore, HiPco tubes are more reactive to the fluorination. A lower fluorination temperature

was used to fluorinated HiPco nanotubes. C2F stoichiometry can be reached at 150°C for fluorinated HiPco

SWNTs. Fluorination at 250°C apparently has caused major destruction of the tubes as evident also by Raman

spectra. Physical observation also indicate the discoloration of nanotubes sample, i.e. SWNTs become yellowish

when fluorinated at 250°C. These fluorotubes are also soluble in alcohols , DMF and THF. Overall, lower

temperature should be used when fluorinating HiPco tubes.

Alkylation

UV-Vis-Nir serves as an excellent monitor for sidewall perturbation of nanotubes n. This may be attributed

to rehybridization at carbon (sp2 to sp3) since the g electrons in the highest occupied molecular orbitals (HOMOs)

that are used to form new bonds are no longer available, and the van Hove transitions characteristic for unperturbed

SWNTs vanish. The UV-Vis-Nir spectra of pristine and hexylated SWNTs are displayed in Figure 4. Thus the

absence of electronic transitions for the hexylated SWNTs supports the assumption that sidewall functionalization

by n-hexyl groups has occurred.

Infrared spectra of SWNTs alkylated by different groups are shown in Figure 5. n-hexyl- and n-butyl

substituents exhibit typical C-H stretching absorptions at 2957, 2924 and 2855 cm -1. The relative intensity of the CH stretching absorptions is higher when the SWNTs are functionalized by n-hexyl groups. Each material exhibits a

peak at - 1578 cm-1 that can be assigned as the active carbon stretching mode of the nanotubes. Assignment of

infrared bands in the 1000 cm -1 to 1500 cm -1 region may arise from the bending modes and umbrella (rocking)

modes of the -CH3 groups.

NASA/CP--2001-210948 689 August 6, 2001
i

(b)

i'+,.. /-,,

400 6_0 £00 1000 I -'__00 ;4O0 1_00

!_ ewiength (rim)

Figure 4. UV-vis-Nir Spectra of hexyl-SWNTs

(a) Pristine SWNTs (b) HexyI-SWNTs

'_"t_,..._._ _._ 2920.6

] (a %....

-" _ 121412155.28

b_

",. & 2925.44

(b) -..

" , 121141_.08

7

"__ _ "_ _ _" _ j4t,_ 2923.8

(c _

1211.2

"+000 5._00 3000 2.=00 2000 1.=00 10t)5

Wav6n;Jn/r_;s f err,-

Figure 5.IR spectra of Alkylated SWNTs

(a) Methyl-SWNTs (b) n-Butyl-SWNTs

(c) HexyI-SWNTs

Thermal gravimetric analysis of degassed (80°C) derivatized SWNTs was used to evaluate the extent of sidewall

functionalization. Figure 6 shows the typical TGA data from alkylated SWNTs. One can clearly notice a loss of

weight over two temperature ranges. Region 1 corresponds to the loss of alkyl groups. The calculated carbon-toalkyl group ratio (C/R) for methyl, n-butyl and n-hexyl SWNTs from TGA weight loss data are listed in Table 1.

The C/R ratios of n-hexyl-SWNTs to n-butyl-SWNTs is about 1.2, similar to their molar mass ratio of 1.5. Similar

studies indicate a higher degree of functionalization when methyl lithium is used as the alkylating reagent.

..... r.............

\ = "- (a)

8>

_Eo

_s

70.

e,

_o.

?eL, ion I _' (b)"" ":

Peg!u_

m _o eo Eo 1co 12o 1_ leO lEO

_rn.)

Table 1. Weight % of alkyl groups (R) on the sidewall

of nanotubes

•R% is obtained from thermal gravimetric analysis.

• The sample was heated up tp and held at 250°C in At.

*C/R=carbon to alkyl group ratio by weight

R R% C/R

Methyl 3.65 27

n-Butyl 9.16 36

Hexyl 11.64 42

Figure 6. Weight loss after heating functionalized SWNTs under Ar

(a) Methylated SWNTs (b) n-butyl-SWNTs (c) n-hexyl-SWNTs

• Weight losses were obtained by thermal gravimetric analysis. It was programmed in three segments : (1)

ramping from room temperature to 80°C and hold for 30 min (2) ramping from 80°C to 250°C and hold

for 30 min (3) ramping from 250°C to 500°C and hold for another 30 min

The spectral signatures of pristine SWNTs return after heating n-hexyl-SWNTs in Ar at 800°C,

demonstrating that at the higher temperature, the alkylation is completely reversible (Figure 7). Further support for

NASA/CP--2001-210948 690 August 6, 2001
reversibilityisprovidedby Ramanspectra.Thusuponheating,therelativeintensitiesof thetangentialnd

breathingmodesreturnandincrease.

(a)

i

' "47d ' '6'00 ' '80d ' "1bob' '1)00' '1407' 'i;0'0' ' '

Waveleng_l

Figure 7a. UV-vis-Nir spectra of Ar heated

hexylated SWNTs

(a) Hexyl-SWNTs (b) Hexyl-SWNTs after heating

{ (a_ I

(b

i,

500 1000 1500 2000 2500

Wavenumb_ (cm 1)

Figure 7b. Raman spectra of Ar heated

hexylated SWNTs

(a) Hexyl-SWNTs (b) after Ar heating

Steric effects were evaluated by comparing the relative reactivites of n-butyl lithium and tert-butyl lithium.

Whereas the van Hove transition totally disappears for the products from the reaction using n-butyl lithium, it is

somewhat lower in relative intensity, yet observable, for tert-butyl lithium reaction(Figure 8). As expected, the

weight loss of t-butyl SWNTs is much less than n-butyl SWNTS after heating SWNTs in argon, confirming that tertbutyl is not effectively reacted with SWNTs since both groups have the same molecular mass.

i t-BuLi 3.22

!' \ n-BuLi 9.16

{',

{ *_... "-.._ (b)

aoo _o0 £00 ino0 1200 I_*00 _GOn

Wave!eng31 _r.n.

Figure 8. UV-vis-Nir spectra of butylated SWNTs

(a) SWNTs reacted with t-Butyl lithium

(b) SWNTs reacted with n-Butyl lithium

R % is obtained from thermal gravimetric analysis

The effect of solvent on these alkylation reactions was also investigated. Fluorotubes are soluble in THF,

but not in hexane or ether. Table 2 shows the weight % gain when SWNTs were reacted with n-hexyl lithium at

room temperature in various solvents. The effect of temperature is illustrated in Table 3. The effect is negligible in

THF, whereas a significant increase of alkyl attachment is observed with n-hexane. Overall, we can obtain an

optimized yield of alkyl-SWNTs by lowering the reaction temperature in hexane solution.

NASA/CP--2001-210948 691 August 6, 2001
Table2.Weightpercentof attached n-hexyl groups

with reaction done in various solvents.

Table 3. Weight percent of n-hexyl groups on

SWNTs with different reaction T and

solvents

Solvent R%

THF 10.0

Ether 7.0

Hexa_e 11.7

TMEDA 6.54

R%

THF

Hexa_e

*R % is obtained from TGA

Room T -40°C

10.2 % 10.0%

11.7% 16.6%

These results are consistent with a multi-step process that is initiated by one-electron transfer from the alkyl

lithium reagent to the nanotube. Expulsion of fluoride from the resulting radical anion would lead to a radical site

on the SWNT. Recombination of the alkyl radical from the lithium reagent with the SWNT would lead to the

alkylated nanotube. Crowding of alkyl groups suggests that not every fluorine would be replaced by an alkyl group.

Steric effects probably account for the observation that more extensive alkylation occurs when sterically less

demanding methyl lithium is used. Additional strong support for the electron transfer process comes from the fact

that phenyl lithium which is known to react by a two electron process does not react with fluorinated SWNTs.

CONCLUSION

The reaction chemistry of solvated Fluorotubes could be an important route for preparation of various

functionalized tubes with a variety of useful properties. In addition, Fluorotubes can be used in several applications,

including battery and composite formation. We have also demonstrated that alkyl lithium reagents may be used to

attach alkyl groups to the sidewalls of nanotubes. Thermal gravimetric analysis combined with UV-Vis-Nir

spectroscopy provides a qualitative measurement of the degree of functionalization. The pristine nanotubes can be

recovered after thermolysis demonstrating that the alkylation is reversible. Evidence suggests that some of the alkyl

groups are attached to sites forming less stable C-R bonds. Chemical modification on the sidewall by alkylation

leads to derivatives that are soluble in common organic solvents such as tetrahydrofuran and chloroform. Future

studies will focus on cross-linking of SWNTs as a route to new materials that might be useful in the manufacture of

nanotube fibers and composites.

ACKNOWLEDGEMENT

The authors would like to thank the supports from Advance Technology Program of Texas and the Robert

A. Welch Foundation, NASA and the National Science Foundation for support of this work.

REFERENCES

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2.

3.

4.

5.

6.

7.

Iijima, S.;Helical microtubules of graphitic carbon.Nature (London).,354(6348), 1991, 56.

Ebbesen, T. W. and Ajayan, P. M.;Large-scale synthesis of carbon nanotubes.Nature.,358(6383),1992, 220.

Heben, M. J. and Dillon, A. C.;Room-temperature hydrogen storage in nanotubes.Science.,287(5453),2000,

593.

Mickelson, E. T.;Huffman, C. B.;Rinzler, A. G., et al.;Fluorination of single-wall carbon nanotubes.Chem.

Phys. Lett.,296(1,2),1998, 188.

Mickelson, E. T.;Chiang, I. W.;Zimmerman, J. L., et al.;Solvation of fluorinated carbon nanotubes.J. Phys.

Chem. B.,103,1999, 4318.

Hamon, M. A.;Chen, J.;Hu, H., et al.;Dissolution of Single-Walled Carbon Nanotubes.Adv.

Materl.,11(10 ),1999 , 834.

Wong, S. S.;Joselevich, E.;Woolley, A. T., et al.;Covalently functionalized nanotubes as nanometer-sized

probes in chemistry and biology.Nature.,394(6888), 1998, 52.

NASA/CP--2001-210948 692 August 6, 2001
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9.

10.

11.

12.

13.

Boul, P. J.;Liu, J.;Mickelson, E. T., et al.;Reversible sidewall functionalization of buckytubes.Chem. Phys.

Lett.,310(3,4), 1999, 367.

Rinzler, A.;Liu, J.;Dai, H., et al.;Large-scale purification of single-wall carbon nanotubes: process, product

and characterization.Appl. Phys. A..,67,1998, 29.

Nikolaev, P.;Bronikowski, M. J.;Bradley, R. K., et al.;Gas-phase catalytic growth of single-walled carbon

nanotubes from carbon monoxide.Chem. Phys. Lett.,313(1,2),1999, 91.

Chiang, I. W.;Brinson, B. E.;Smalley, R. E., et al.;Purification and characterization of single-wall carbon

nanotubes.J. Phys. Chem. B.,,2000, 105(6), 2001:11157-1161.

Chiang, I. W.;Brinson, B. E.;Margrave, J. L., et al.;Purification of single-wall carbon nanotubes produced

by high pressure CO process.J. Phys. Chem. B.,,2001, submitted.

Nakajima, Y.,Fluorine-carbon and fluoride-carbon materials : chemistry, physics, and applications,New

York, 1995

NASA/CP--2001-210948 693 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

LOW-TEMPERATURE FLUORINATION OF CARBON NANOTUBES.

A.V. Okotrub, N.E Yudanov, L.G. Bulusheva, A.I. Romanenko

Institute of Inorganic Chemistry SB RAS, pr. Ak. Lavrenteva 3, Novosibirsk 630090

A.L Chu_ilin

Boreskov Institute of Catalysis SB RAS, pr. Ak. Lavrenteva 5, Novosibirsk 630090

ABSTRACT

Two samples of carbonaceous material produced in an electric arc were fluorinated. The first one was cut from

the inner part of cathode deposit grown in helium gas of 800 Torr and contained roughly 80% of multiwall tubular

and polyhedral particles, the second one was synthesized by co-evaporation of graphite and cobalt-nickel powder

and involved about 20% of single-wall carbon nanotubes. Fluorination of these materials was performed at room

temperature using a volatile BrF3. Elemental composition of the fluorinated samples contained multiwall and singlewall nanotubes, which was determined by X-ray photoelectron spectroscopy, corresponds to CFo.3 and CFo_

respectively. Under the same synthetic technique graphite is fluorinated up to C2F stoichiometry.

Fluorinated materials were examined by transmission electron microscopy. Only two or three outer shells of

multiwall nanoparticles were found to be usually fluorinated, while the inner part retained intact. In the rare cases,

when the particles were conglomerated, they had up to 15 fluorinated layers with the estimated average separation of

5.21 A. The low-temperature fluorination was detected to reduce the content of multiwall nanotubes in

the sample that may be explained by peeling of the outer layers or tube unfolding. If so, the most of the arcproduced multiwall carbon nanotubes are composed of the nested scrolls surrounding the cylindrical inner shells.

Fluorination also decreases the interaction between single-wall nanotubes composed a rope leading to they splitting.

Conductivity of pristine materials and fluorinated ones was measured by four-point technique. Upon the

fluorination the room temperature conductivity of the samples was found to decrease more than in 104. Similar

temperature dependence of conductivity for fluorinated and non-fluorinated multiwall carbon nanotubes at T< 100K

suggests a retention of intertube contacts in the former sample. The electric transport in the fluorinated single-wall

carbon nanotubes differs from that in the pristine material and is characteristic of narrow-gap semiconductors.

Field emission properties of pristine single-wall nanotubes and defluorinated ones were measured. The latter

sample showed the increase of electron emission in the low-voltage region of I-V dependence.

This work was supported by INTAS (Project No. 97-1700), the Russian scientific and technical program

<<Actualdirections in physics of condensed states>>on the <cFullerenes and atomic clusters>> (Projects No 98055) and the

<<Surface atomic structures>> (Projects No 4.14.99) and the Russian Foundation for Basic Research (Projects Nos. 00-03,

32463a, 00-03-32510a and 00-02-17987).

Keywords: carbon nanotubes, fluorination, electric transport.

Alexander V. Okotrub

Institute of Inorganic Chemistry

Pr. Ak. Lavrenteva 3

Novosibirsk 630090

Russia

Fax: +7 (3832) 344489

E-mail: spectrum@che.nsk.su

NASA/CP--2001-210948 694 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

SYSTEMATIC STUDY FOR PURIFICATION OF SINGLEWALLED CARBON

NANOTUBES

Jeong-Mi Moon, Kay HyeokAn, Chulsu Jo, Young Hee Lee*

Department of Physics, Institute of Basic Science, Sungkyunkwan University, Suwon 440-746, Kyunggi,

South Korea

Young Soo Park, Hee Jin Jeong, Dong Jae Bae, Seong Chu Lim

Department of Semiconductor Science and Technology, Jeonbuk National University, Jeonju 561-756, South

Korea

Gyeong-Su Park

Analytical Engineering Laboratory, Samsung Advanced Institute of Technology, Suwon 440-600, Kyunggi,

South Korea

ABSTRACT

We present a systematic study for purification process of singlewalled carbon nanotubes (SWCNTs), which

was developed by combining two-step processes of thermal annealing in air and acid treatment. The process

involves the thermal annealing in air with the powders to be rotated at temperatures of 470 °C for 50 rain, which

burns out the carbonaceous particles, and an acid treatment with HC1 for 24 hours, which etches away the catalytic

metals. Control of the annealing temperature and rotation of the sample are crucial for high yield. Our

reproducible optimal purification process provides a total yield of about 25 ~ 30 wt% with less than 1 wt% of

transition metals, which was confirmed by the thermogravimetric analysis (TGA). We also find that bundling and

length control depend on the different acid treatments. In order to disperse the SWCNT bundles, the hydrochloric

treated samples were again boiled in nitric acid for six hours. Scanning electron microscope (SEM) and

transmission electron microscope (TEM) indicate the well-dispersed SWCNTs, although SWCNTs were broken

into small pieces. We observe that some multiwalled CNTs were formed during the nitric acid treatment. The

origin of such a formation mechanism will be discussed.

Keywords: two-step purification processes, singlewalled carbon nanotubes, multiwalled carbon nanotubes.

INTRODUCTION

Traditionally arc discharge was used for mass production (refs. 1 to 3). More recently laser vaporization

technique provided relatively highly purified singlewalled CNTs (SWCNTs) (ref. 4). Most approaches in general

produce powders containing not only CNTs but also other carbonaceous particles such as amorphous carbons,

fullerenes, nanocrystalline graphites, and transition metals which were introduced as a catalyst during the synthesis

(refs. 1 to 7). These sometimes hinder an accurate analysis of the CNT characteristics and limit the best

performance of the CNT applications to new functional devices.

Several attempts have been tried to purify the CNT powders. Gas phase reaction or thermal annealing in air or

oxygen atmosphere has been attempted, although the yield of final product was relatively low (ref. 8,9). The key

idea with these approaches is a selective oxidative etching processes, based on the fact that the etching rate of

amorphous carbons is faster than that of CNTs. Since the edge of the CNTs can be etched away as well as

carbonaceous particles during the annealing, it is crucial to have a keen control of annealing temperatures and

annealing times in order to obtain high yield, although the yield is also dependent on the purity of the original

sample. Liquid phase reaction in various acids has been tried to remove the transition metals (ref. 7,10,11). This

process involves repeated steps of filterings and sonications in acidic solution, where the transition metals were

melted into the solution. CNTs are severely damaged with strong acid solution, although the wall structures could

be recovered by the subsequent annealing. Therefore the choice of acids, the immersing time, and temperature are

the key factors to have high yield, while maintaining the complete wall of CNTs. These two processes are

complementary with each other to remove the carbonaceous particles and transition metals.

NASA/CP--2001-210948 695 August 6, 2001
In this Letter, we adopt these two purification steps to SWCNTs synthesized by catalytic arc discharge. Special

cares are taken to have high yield during annealing process and acid treatment. The whole processes are

reproducible and completely robust for mass production. The total yield of 25 ~ 30 wt% was easily obtained with

the metal content of less than 1 wt%.

EXPERIMENTAL SECTION

Singlewalled carbon nanotubes (SWCNTs) were prepared by conventional catalytic arc discharge (ref. 12). The

total amount of catalyst in a graphite powder was fixed to be 5 wt%, with the ratio of the transition metals (Ni :

Co : FeS = 1:1:1), where sulfur is added as a promotor (ref. 12). The raw sample contains carbonaceous particles

and transition metals. The systematic purification procedure is as follows. This cloth-like raw sample was grinded

mechanically and transferred to the heating chamber, where the CNT powder was heated at 470 °C for 50 min in

air in order to remove carbonaceous particles.

In order to remove the catalysts, the annealed powder was immersed and filtered in 6 M hydrochloric acid

(HC1) for various times. To unbundle SWCNTs, the sample obtained at the last step was sometimes boiled in 30 %

nitric acid (HNO3) for 4 to 6 hours. After rinsing and drying, a grayish black, thin mat comprised of SWCNTs was

finally obtained.

RESULTS AND DISCUSSION

Figure l(a) shows a SEM image of the grinded raw soot collected from the whole chamber. In addition to the

bundle of SWCNTs, the white spot, carbonaceous particles, mainly amorphous carbons are shown in the figure.

Furthermore, we can see that the metal particles were mainly embedded in larger amorphous carbons, as evidenced

in Fig. 2(a). This sample was thermally annealed at 470 °C for 50 min in air to remove selectively the amorphous

carbons. The SEM image of oxidized SWCNTs is shown in Fig. l(b). The figure shows a dramatic decrease of

amorphous carbons after oxidation. As a result of oxidation, the weight of the soot was reduced to about 40 wt% of

the original sample.

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Figure 1. SEM images of (a) the raw SWCNT sample, (b) heat-treated sample, (c) HCl-treated sample, and

(d) HNO3-treated sample.

NASA/CP--2001-210948 696 August 6, 2001
TheTEMimageoftheoxidizedSWCNTsshowninFig.2(b)suggeststhatamorphous carbonscoatedonthe

surfacesofSWCNTbundlesandthosesurroundingthecatalyticmetalswereremoved. Nowthetransitionmetals,

asindicatedbytheblackspotsinFig.2(b),areleft.Thetransitionmetalscanbeetchedawaybyanacidtreatment.

Figurel(c)showstheSEMimageoftheacid-treatedsample, wheretheannealedsamplewasimmersedin 6M

hydrochloricaidfor24hours.Thecorresponding TEMimageinFig.2(c)clearlyshowsthathetransitionmetal

particleswerewellmeltedaway.Yetsomecarbonaceous particles, whichareleftduringtheannealingprocedure,

stillremaininthesample, asillustratedinFig.l(c).Special careshouldbetakentoobtainbetter-purified samples.

Thisacidtreatment gavetheyieldof70wt%ofthethermallyannealedsamples. Thusthetotalyieldoftwostepsis

about25-30wt%.InordertodispersetheCNTbundles, theabovesampleswereagainboiledin30%nitricacid

forsixhours.Someopaquebackgrounds represent thewell-dispersed SWCNTs, althoughSWCNTswerebroken

intosmallpieces, asshowninFig.l(d).It isinterestingtoobservesomemultiwalledCNTs(MWCNTs), asshown

inFig.2(d).

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Figure 2. TEM images of (a) the raw SWCNT sample, (b) heat-treated sample, (c) HCl-treated sample, and

(d) HNO3-treated sample.

The MWCNTs were formed during the nitric acid treatment. We tried several different boiling times in nitric

acid. The MWCNTs rarely appeared at boiling time less than 2 hours. Figure 3(a) is the high-resolution TEM

image of the sample boiled in 30% nitric acid for 2 hours. The SWCNT bundles and graphitic sheets are shown in

the figure. These graphitic sheets are believed to be produced by the part of SWCNTs. No graphitic sheet was

observed in the as-grown SWCNTs, as shown in Fig. 2(a). The sample treated by 30% nitric acid for 6 hours

showed mostly MWCNTs and graphitic sheet, as shown in Fig. 3(b). The observed MWCNTs are very defective

and bamboo-typed, similar to those grown by chemical vapor deposition at low temperature. It may be conjectured

that small graphitic flakes are intermediate phases during the transformation from SWCNTs to MWCNTs. Since

this is low-temperature (100 °C) process, this hypothesis of transformation is hardly acceptable. Catalytic

decomposition of SWCNTS and formation of MWCNTs should be involved in order to promote such reactions in

the electrolyte. Yet, the origin of such a formation mechanism is not clear at this moment.

NASA/CP--2001-210948 697 August 6, 2001
i:i:!:!:?i::_.$?.i:i¢i:i:i:i:i_'_i?.i:i:!:i

Figure 3. HRTEM images of HNO3-treated sample for (a) 2 h and (b) 6 h.

Figure 4 shows the thermogravimetric analysis (TGA) graphs and the differentiated TGAs in air for the raw

soot and the purified SWCNTs. In the raw sample, the weight starts reducing near 300 °C. This decomposition

temperature is relatively low but catalyst-assisted dissociation presumably takes place during the process. The

SWCNTs are completely burned near 900 °C. The remaining materials are the transition metals, which is about 8

wt%. This demonstrates the existence of at least three phases in the raw sample. One phase that shows a broad

peak position at 470 °C is assumed to be amorphous carbons, although small peak appears near 550 °C, which may

be conjectured as crystalline nanoparticles. Another phase at 800 °C is the SWCNTs. The larger integrated area of

the first and second peaks compared to the third one indicates that more carbonaceous particles than SWCNTs in

fact exist in the sample, reflecting the purity of our powder to be low. Our choice of the annealing temperature at

470 °C seems to be reasonable from this analysis. The TGA of the heat-treated sample (dashed line) shows

relatively a simple monotonic decrease with increasing temperatures. The amorphous carbons are mostly removed,

although a small amount of amorphous carbons still remain in the sample, as indicated by the small peak from the

differentiated TGA curve [Fig. 4(b-(ii))] near 400 °C. The remaining metal content is high in this case (18 wt%),

as expected from the relative composition. Figure 4(a-(iii)) shows that the transition metals are nearly melted away

to the TGA detection limit (< lwt%) for acid-treated sample followed by the thermal treatment. The differentiated

TGA peak (dotted line) is narrower in the purified sample, indicating mostly burning of the SWCNTs. Yet, a small

bump near 550 °C exists, suggesting that carbonaceous particles still remain in the sample. The estimated purity of

the sample from the integrated area of the differentiated TGA is about 96 wt%. After the nitric acid treatment, the

burning temperature is lowered, as shown in the dot-dashed line, and the CNT peaks are separated [Fig. 4(b-(iv))],

indicating some other new phases are generated during the nitric acid treatment, as suggested in the TEM images.

t (iii) a

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0 260 4;0 6;0 800 1000 0 200 400 600 800 1000

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Figure 4. (a) TGA and (b) differentiated TGA graphs of (i) the raw SWCNT sample, (ii) heat-treated sample,

(iii) HCl-treated sample, and (iv) HNO3-treated sample, respectively.

NASA/CP--2001-210948 698 August 6, 2001
CONCLUSION

We have obtained high-yield SWCNTs by combining two-step processes of thermal annealing in air and acid

treatment. The optimal conditions are the thermal annealing in air with the powders at temperatures of 470 °C for

50 min and HC1 acid treatment for 24 hours. Our reproducible optimal purification process provides a total yield of

about 25 ~ 30 wt% with less than 1 wt% of transition metals, and 96 wt% of purity, which was confirmed by the

thermogravimetric analysis. We observed a phase transformation from SWCNTs to MWCNTs in the nitric acidtreated sample

ACKNOWLEDGMENTS

This work was supported in part by the MOST through NRL and New Frontier programs and in part by the

BK21 program. The authors are grateful to Kwang-Ju branch of Korea Basic Science Institute for the assistance of

FE-SEM.

REFERENCES

Corresponding author. Young Hee line: Department of Physics, Institute of Basic Science, Sungkyunkwan

University, Suwon 440-746, Kyunggi, South Korea, E-mail: leeyoung@yurim.skku.ac.kr,

fax: +82-31-290-5954, phone: +82-31-290-7053

1. Journet, C., et al.: Large-Scale Production of Single-Walled Carbon Nanotubes by The Electric-Arc

Technique. Nature, vol. 388, August 1997, 756.

2. Saito, Y., et al.: High Yield of Single-Wall Carbon Nanotubes by Arc Discharge Using Rh Pt Mixed

Catalysts. Chem. Phys. Lett., vol. 294, no. 6, September 1998, 593-598.

3. Shi, Z.J., et al.: Mass-Production of Single-Wall Carbon Nanotubes by Arc Discharge Method. Carbon,

vol. 37, no. 9, 1999, 1449-1453.

4. Thess, A., et al.: Crystalline Ropes of Metallic Carbon Nanotubes. Science, vol. 273, 1996, 483.

5. Kataura, H., et al.: Optical Properties of Single-Wall Carbon Nanotubes. Synth. Met., vol. 103, no. 1-3,

June 1999, 2555-2558.

6. Flahaut, E, et al.: Synthesis of Single-Walled Carbon Nanotubes Using Binary (Fe, Co, Ni) Alloy

Nanoparticles Prepared in situ by The Reduction of Oxide Solid Solutions. Chem. Phys. Lett., vol. 300, no.

1-2, January 1999, 236-242.

7. Shi, Z.J., et al.: Purification of Single-Wall Carbon Nanotubes. Solid State Commun., vol. 112, no. 1,

August 1999, 35-37.

8. Ebbesen, T.W., et al.: Purification of Nanotubes. Nature, vol. 367, February 1994, 519.

9. Zimmerman, J.L., et al.: Gas-Phase Purification of Single-Wall Carbon Nanotubes. Chem. Mater., vol. 12,

no. 5, July 2000, 1361-1366.

10. Tohji, K., et al.: Purification Procedure for Single-Walled Nanotubes. J. Phys. Chem. B, vol. 101, no. 11,

March 1997, 1974-1978.

11. Liu, B., et al.: Synthesis and Characterization of Single-Walled Nanotubes Produced with Ce/Ni as

Catalysts. Chem. Phys. Lett., vol. 320, no. 3-4, April 2000, 365-372.

12. Park, Y.S., et al.: Low Pressure Synthesis of Singlewalled Carbon Nanotubes by Arc Discharge. Synth.

Met., 2001. In press.

NASA/CP--2001-210948 699 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Carbon nanotube biosensors for space molecule detection

and clinical molecular diagnostics

Jie Han

NASA Ames Research Center

Both space molecule detection and clinical molecular diagnostics need to develop ultra sensitive

biosensors for detection of less than attomole molecules such as amino acids or DNA. However

all the electrode sensor systems including those fabricated from the existing carbon nanotubes,

have a background level of nA (nanoAmp). This has limited DNA or other molecule detection to

nA level or molecules whose concentration is much higher than attomole level. A program has

been created by NASA and NCI (National Cancer Institute) to exploit the possibility of carbon

nanotube based biosensors to solve this problem for both's interest. In this talk, I will present our

effort on the evaluation and novel design of carbon nanotubes as electrode biosensors with

strategies to minimize background currents while maximizing signal intensity. The fabrication

of nanotube electrode arrays, immobilization of molecular probes on nanotube electrodes and in

vitro biosensor testing will also be discussed.

Dr, Jie Han

Technical Lead, Nanotechnology

NASA Aries Research Center

han@nas.nasa.gov

NASA/CP--2001-210948 700 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

SEPARATION OF CO2 FROM FLUE GASES BY CARBON-MULTIWALL CARBON

NANOTUBE MEMBRANES

Rodney Andrews, Marit Jagtoyen, Eric Grulke

University of Kentucky Advanced Carbon Materials Center

and

Center for Applied Energy Research

2540 Research Park Drive, Lexington, KY 40511

Andrews@caer.uky.edu, ph 859-257-0265, fax 859-257-0220

Ki-Ho Lee, Zugang Mao, Susan B. Sinnott

Department of Materials Science and Engineering, University of Florida

154 Rhines Hall, PO Box 116400, Gainesville, Florida 32611-6400

ABSTRACT

Multiwalled carbon nanotubes (MWNT) were found to be an effective separation media for removing CO2 from

N2. The separation mechanism favors the selective condensation of CO2 from the flowing gas stream. Significant

uptakes of CO2 were measured at 30°C and 150°C over the pressure range 0.5 to 5 bar. No measurable uptake of

nitrogen was found for this range of conditions. The mass uptake of CO2 by MWNT was found to increase with

increasing temperature. A packed bed of MWNT completely removed CO2 from a flowing stream of CO2/N2, and

exhibited rapid uptake kinetics for CO2.

Keywords: multiwall carbon nanotubes, gas separations, graphitization, molecular dynamics simulations.

INTRODUCTION

The feasibility of using a novel carbon-multiwall carbon nanotube membrane for separating CO2 from the flue

gas of a power generation plant is being studied. Such an innovative membrane system offers unique advantages

over existing technologies: refractory carbon-carbon membranes are resistant to temperature and chemical attack, the

multiwall nanotube derived pores in the membrane are mono-disperse, the pore size can be controlled, and the rapid

kinetic and diffusion rates will yield high permeate fluxes. As the first step toward design and construction of a

working carbon-carbon nanotube based membrane, specific goals have included evaluation of the separation

mechanism, either diffusive or adsorptive, and a test to demonstrate proof-of-concept separation. Experimental work

has determined uptakes and separation efficiencies for CO2 and N2 mixtures by open ended multiwall carbon

nanotubes.

As-produced multiwall nanotubes (MWNT) are open at one end, but contain small particulate inclusions of iron

catalyst blocking access to the core. We have found that a simple graphitization step (heating to 1800 °C in an inert

atmosphere) removes this iron contaminant, greatly enhancing adsorption within the nanotube core. We have

produced and graphitized sufficient quantities of MWNT to allow for bench scale studies on these materials.

EXPERIMENTAL

Production of MWNT

We have previously described the synthesis of MWNTs by reacting hydrocarbon vapor over a dispersed iron

catalyst that is deposited in situ in a quartz tube reactor within a multi-zone furnace (ref. 1). A xylene-ferrocene

feedstock was continuously injected via syringe pump into a preheat section operated at -250 °C. The xyleneferrocene vapors were swept into the reaction zone of the furnace by an Ar/10% H2 carrier gas that also maintained a

partial pressure of 32 mbar carbon inside the reactor. The reaction zone was held at 725 °C, with an Ar/H2 flow rate

of 6 L/min (STP). The operating procedure was: the quartz tube and substrates were installed into the furnace and

then purged with Ar gas; the preheater and furnace heaters were ramped up to achieve the desired stable

temperatures; the liquid and gas feeds were started and run for the desired reaction times. The MWNTs grew on

NASA/CP--2001-210948 701 August 6, 2001
boththequartztubewallandflatquartzplateaddedforadditional surfacearea,formingthickmatsofwell-aligned

nanotubes thatcouldbereadilyharvestedbybrushingthesurfaces.

Graphitization ofMWNT

Forgraphitization, thebulkMWNTsampleswerecenteredwithinahorizontal electricresistance tubefurnace.

Afterpurgingthesystemwith dry nitrogen (Air Products NF grade) and maintaining slightly above atmospheric

pressure, the samples were heated from ambient to 1000 °C at 20 °C/min. Once 1000 °C was reached, the furnace

was put under automatic control (optical pyrometer) and heated at 12.5 °C/min to the set point. Samples were

treated to 2800 °C, and held at temperature for 45 min.

Mass Uptake of CO2 and N2 on Multiwall Nanotubes

The mass uptake of CO2 and N2 on graphitized MWNT was measured using the Hiden Intelligent Gravimetric

Analyzer (IGA). The nanotubes were loaded into the sample side of the apparatus, and the chamber sealed and

evacuated to 10-6 mbar. High purity CO2 or N2 was introduced into the chamber and the pressure raised in 500 mbar

increments to a maximum of 5 bar. At each pressure step, the IGA waits for the sample mass to equilibrate and then

records this value versus the samples initial value. This procedure was performed for both CO2 and N2 gases at 30°C

and 150°C.

Separation of Flowing Stream of CO2 and N2 on Packed MWNT Bed

Apacked bed ofnanotubes was used to separate a flowing stream of CO2 and N2. The gas composition was

50% CO2 and 50% N2 with a volumetric flow of 100 SCCM total. The bed consisted of 0.26 g of graphitized

MWNT packed into a fixed column 6 mm in diameter and 80 mm long. Ablank column of the same dimensions was

installed as a bypass line, and the outlet from this system (diluted with a carrier stream of He) was sent to a mass

spectrometer. Figure 1 shows a schematic of the fixed bed apparatus. Once a baseline spectra was determined by

flow on the bypass side, the gas flow was switched to the MWNT bed, the mass spectrometer was used to detect

CO2, indicating breakthrough. After saturation of the bed, the flow was switched to the bypass, and the time taken to

fill the empty bed was measured. The empty bed time was used to calculate the proportion of time for CO2

adsorption on the MWNT prior to saturation and the time due to simple flow dynamics. As the flow time on the

empty bed would in actuality be longer than for a packed bed, the measured CO2 uptake value on the packed bed

should be considered a minimum value and could be higher than that measured.

CO 2 Mass ]

Flow

Controller] I

N 2 Mass ] _

c°Fltr°ller _ L_

 iiiiiiiiiiiiiii! i! !i!!!iii

He Sweep

Figure 1. Packed bed flow separation apparatus.

ToSpec. Mass

NASA/CP--2001-210948 702 August 6, 2001
TheMolecular DiffusionandDynamicFlowofCO2,N2and02throughCarbonNanotubes

Thesimulations wereperformedonaCompaqAlpha64Workstation. Thekeycomponents ofworkcompleted

orunderwaytosatisfythemodelingrequirements are:

(1)Searching Lermard-Jones parameters

Forthecalculationfintermolecular potentialsandtotalenergy, Lermard-Jones potential parameters ofnewly

appliedatoms, i.e.nitrogenandoxygen, hadtobeobtained. Whileparameters hadbeencomputationally compared

withthebondlengthsandbondingenergies, andeachother,wechosethebestparameters.

Atomic# AtomicMass[g/mol] JkB[K]* o[A]

C 6 12.0 28.129 2.757(ref.2)

N 7 14.0 37.300 3.310(ref.3)

O 8 16.0 80.507 3.033

(2)Coulombinteraction

Applyingpartiallychargedcarbondioxide, wehavetoaddedpotentialsduetoCoulombpotential intototal

energy[4].

UqC*qC*=+6.105/rC*C* [eVA/A]

UqC*qO*=+3.053/rC*C* [eVA/A]

UqO*qO*=+1.526/rC*C* [eVA/A]

(3)Covalent potentials

Theparameters forcovalent potentials(Rebopotentialsinthissimulation) ofN2,02,andCO2arenotprepared

yet.Untiltheparameters forthegaseswillbeobtainedexperimentally orcalculatedbytrial-anderror,weusethe

parameters assumedtobethesamevaluesasthoseforH2.

(4)Modificationtoexistingcode

Inadditionofsomeparameters, wearemodifyingourprogramfornewelements becausethecurrentlyused

programisadjustableforthesimulationofhydrocarbon systemthatconsistsofjusttwoatomtypes.

RESULTS AND DISCUSSION

N2 and CO2 Adsorption on MWNT

The as-produced MWNT were graphitized prior to use to ensure removal of any blockages caused by residual

iron catalyst from the core oflhe nanotubes. No significant mass uptake of N2 was measured on the MWNT in

either static or flowing mode. The mass uptake of nilxogen was unaffected by temperature (30°C or 150 °C).

The mass uptake of CO2 on MWNT was measurable at both 30°C and 150°C. As shown in Figures 2 and 3, the

mass uptake was low (-3% by weight) but was significant. The uptake was higher at higher temperatures, indicative

of a chemisorption mechanism. A screening study showed this uptake to be reversible, therefore full desorption

isotherms will also be measured to determine if this mass uptake is reversible or not. This finding will significantly

affect the design of a MWNT membrane in future work.

The mass uptakes increased with increasing gas pressure (to 5 bar) as would be expected. The significant mass

uptake of CO2 onto the nanotubes at 5 bar and 150°C signifies that the MWNT are an effective medium for

separation of CO2 and N2. The goal of this project was to determine what separation mechanism would be viable for

the design of a nanotube membrane for CO2 removal from flue gases, either diffusion driven or condensation driven.

As the uptake of CO2 is significant while the uptake of N2 is negligible, the condensation mechanism appears to

dominate. However, full desorption isotherms at higher temperatures and continued modeling of the system are

required to fully understand the operating limitations of this system and to verify that condensation is the controlling

mechanism.

Apacked bed of MWNT was used to separate a flowing stream of 50 SCCM CO2 and 50 SCCMN2. The

packed bed configuration was employed to provide a more realistic comparison of the behavior of a membrane based

on MWNT. The packed bed also provides valuable data on the uptake kinetics relative to gas velocities and

concentrations. The bed was able to take up all the CO2 for 1.4 minutes (Figure 4), while the blank bed retained CO2

NASA/CP--2001-210948 703 August 6, 2001
160 ..............................................................................

140

120

100 -

60

60-

4O

2O

04

j*

-,-e,-,,VG B 42 ,7"

-X- VGB S2 "/"

./

./

.O f

1000 2000 3000 4000 5000

Pressure (mbar)

6000

Figure 2. Mass uptake of CO2 at 30 °C. Sample weights of 37 mg (VGB 42) and 54 mg

(VGB-S2).

for only 1 minute. Therefore the MWNT took up CO2 prior to breakthrough for at least 0.4 minutes (see

Experimental section as to why this is a minimum uptake value). This uptake value is in excess of 10% by mass and

indicates that nanotubes are an effective separation media for removing CO2 from N2. This result indicates that a

satisfactory separation of CO2 and N2 could be designed around a simple packed bed of MWNT.

1200

1000

800

__ 600

400

2OO

.......................................................................... 1

1000 2000 3000 4000 5000 6000

Pressure (mbar)

Figure 3. Mass uptake of CO2 at 150 °C. Sample weight of 54 mg.

NASA/CP--2001-210948 704 August 6, 2001
o_

,'T

2.00E-09 -

1.50E-09 -

1.00E-09

5.00E-10

0.00E+00

--Nitrogen

_Carbon dioxide

..,,-

f-

/

/

/

/

i

/

Empty bed

breakthough

On column /

MWNT packed bed

breakthrough

11 12 13 14 15 16 17

Time (min)

Figure 4. Selective uptake of CO2 on packed bed of MWNT. Gas composition: 50 SCCM N2 and

50 SCCM CO2. Packed bed: 0.26g of MWNT.

Modeling of Separation Mechanism

Classical molecular dynamics simulations are used to investigate the diffusive flow of pure molecules and binary

molecular mixtures. Standard Lermard-Jones potentials are used to model the intermolecular interactions. Both Hterminated and C-terminated open nanotube ends have been considered. The specific molecules that are being

examined include methane, ethane, nitrogen, oxygen, carbon dioxide, methane/ethane, methane/n-butane,

meflaane/isobutane, nitrogen/oxygen, nitJcogen/carbon dioxide and oxygen/carbon dioxide (nitrogen/carbon dioxide

are specifically considered for flais project). The simulations predict which binary mixtures separate as a result of this

diffusive flow and which remain mixed. They also indicate how these results depend on the nanotube properties such

as diameter and helical symmetry. The simulations provide information about how flae structure and size of the

molecules in the mixtures influence the results. For example, n-butane and isobutane are predicted to have

significantly different separation behaviors when they are mixed wifla methane molecules. In addition, molecules

with non-spherical aspect ratios exhibit different behavior than spherical molecules that affects both their diffusion

mechanisms and their separations in mixtures. The study includes individual nanotubes and nanotubes in bundles and

flae degree of separation predicted depends on whether or not the nanotubes are in bundles.

We are working to incorporate the appropriate potential for characterizing N2/CO2 mixtures into the molecular

dynamics program. In some studies these molecules are treated as single, spherical particles (ref. 5). While this

simplifies the simulation considerably, our previous work has demonstrated that molecular shape can have a

significant influence on molecular diffusion through nanopores (ref. 6). Therefore the approach that is being taken is

one where the atoms are treated explicitly and the molecules are characterized with Coulomb and Lermard-Jones (L J)

potentials (refs. 7 and 8). The atoms in the carbon nanotube walls will be characterized with a many-body reactive

empirical bond-order potential that is coupled to LJ potentials (ref. 9).

CONCLUSIONS

MWNT have been shown to be an effective media for the separation of CO2 from N2. The results indicate flaat

flaese materials should be effective for separating CO2 from flue gases at elevated temperatures and pressures. The

data indicate that a MWNT membrane system could be designed based on the mechanism of CO2 condensation

within the pores of the nanotubes. As flais uptake of CO2 has been found to increase with temperature, these

materials seem ideal for use at the elevated temperatures found in a flue stream. Concurrently, the use of a packed

NASA/CP--2001-210948 705 August 6, 2001
bed of MWNT has been show to be an effective separator for CO2 from N2 and does offer an alternative technology

for performing this separation.

REFERENCES

[1] Andrews R., et al.: Chem Phys Lett 1999; 303:467-474.

[2] Murthy, C.S.; O'Shea, S.F.; and McDonald, I.R.: Mol. Phys., 50 (1983) 531.

[3] Harris, J.G.; and Yung, K.H.: J. Phys. Chem., 99, 12021(1995).

[4] Cui, S.T.; Cochran, H.D.; and Cummings, RT.: J. Phys. Chem. B 1999, 103, 4485.

[5] Gao, G.-H.; Tan, Z.-Q.; and Yu, "L-X.: "Calculation of high-pressure solubility of gas in aqueous electrolyte

solution based on non-primitive mean spherical approximation and perturbation flaeory",Fluid Phase Equil.

165, 169-182 (1999)

[6] Mao, Z.; and Sirmott, S.B.: "A Computational Study of Molecular Diffusion and Dynamic Flow Through

Carbon Nanotubes",J. Phys. Chem. B 104, 4618-4624 (2000)

[7] Travis, K.R; and Gubbins, K.E.: "Transport diffusion of oxygen-nitrogen mixtures in graphite pores: A

nonequilibrium molecular dynamics (NEMD) study", Langmuir 15, 6050-6059 (1999)

[8] Nakanishi, K.: "Molecular simulation studies in supercritical fluid and related regions", Fluid Phase Equil.

144, 217-224 (1998)

[9] Sirmott, S.B., et al.: "Mechanical Properties of Nanotubule Fibers and Composites Determined From

Theoretical Calculations and Simulations", Carbon 36, 1-9 (1998)

NASA/CP--2001-210948 706 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

New Emitter Techniques for Field Emitter Displays

N.S. Lee, D.S. Chung, S.H. Park, S.J. Lee, J.H. Kang, Y.S. Choi, H.Y. Kim, M.J. Yun,, N.S. Park, LT. Han,,

J.W. Kim, and J.E. Jung, and Jong Min Kim*

Display Lab., Samsung Advanced Institute of Technology

P. O. Box 111, Suwon 440-600, KOREA

*Tel: 82-31-280-9311, Fax: 82-31-280-9349, email: jongkim@sait.samsung.co.kr

J.H. You, C.G. Lee, S.H. Jo, C.J. Lee, K.S, Choi, E.J. Chi, and H.G. Park

1FED team, CRD, Samsung SDI, 575, Shin-Dong, Paldal-Gu, Suwon 442-391, Korea

2Display Lab., Samsung Advanced Institute of Technology,

R O. Box 111, Suwon 440-600, KOREA

Carbon nanotubes (CNTs) have attracted considerable attention as a field emiller material due to their

excellent field emission characteristics, strong chemical stability, and high mechanical strength. Despite a great deal

of effort to realize applications of CNTs, it is expected that the first application of CNTs with a large commercial

market would be electron emitters for field-emission displays (FEDs). bEDs 'are characterized by superior display

performances such as fast response time, wide viewing angles, wide operation temperatures, cathode ray tube (CRT)

like colors, ultra-slim features, low cost, and low-power consumption. The application of CNTs to FEDs necessitates

their vertical alignment on cathode electrodes for better electron emission. At present, CVD technologies to

synthesize well-aligned CNTs over a large area of glass substrates at low temperatures have not been well developed.

Thus, we have focused upon the development of low-cost, simple, and scalable screen-printing technology for the

fabrication of cathode plates with CNT emitters. It is expected that CNTs will enable the production of a low cost,

low-power consumption, and scalability.

Our team fabricated a field emission display (FED) panel with 4,5, 9, and 15 inch diagonals using singlewalled carbon nanotubes. Our studies, however, were restricted to diode structures. For full gray scales and high

brightness of FEDs, a triode structure is required. This study present triode-type FEDs with CNT emitters, including

under-gate triode and normal-gate triode structures. For the under-gate triode structure FED, gate electrodes are

placed underneath cathode electrodes with an in-between insulating layer, On the other hand, the normal-gate triode

emitter structure resembles the Spindt-type cathode except a larger gate hole diameter. For both triode emitter

structures, electron emission from the carbon nanotube emitters is modulated by changing gate voltages. This talk

will introduce their fabrication processes and emission characteristics in detail. Simplicity of structure and

fabrication processes of our triode structure CNT-FEDs seems to enable them to possess a high potential of practical

applications.

Keywords: carbon nanotubes, FEDs, new emitter materials, triode structure, screen printing

NASA/CP--2001-210948 707 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FUNDAMENTAL PROPERTIES OF NANOTUBE FIELD EMITTERS FOR LARGE

AREA ELECTRON SOURCES

Kenneth A. Dean a, Babu R. Chalamala, Bernard F. Coil, A. Alec Talin, Johann Trujillo,

Yi Wei, and James E. Jaskie

Motorola Inc., 7700 S. River Parkway, Tempe AZ 85284

ABSTRACT

Building successful devices based on carbon field emitters requires a fundamental understanding of emitter

behavior during field emission. Detailed investigations of the field emission behavior of individual single-walled

and multi-walled nanotubes have shown several properties that are particularly important for devices. We describe

these properties, which include adsorbate interactions, current saturation, field evaporation, and environmental

degradation, and we discuss their impact on the design of field emission devices. We then describe challenges and

recent progress in the development of large area field emission sources.

INTRODUCTION

In recent years field emission electron sources based on the Spindt tip process have been successfully applied as

broad area electron sources. The Spindt tip process uses relatively expensive processes and tools to lithographically

define, deposit, and control emitter geometry within a few tens of nanometers across a large area panel. Highresolution displays have been built and sold using the Spindt technology) Yet there may be a less expensive way to

produce a large area electron source using field emitters. Carbon nanotubes are naturally sharper than Spindt tips,

and they can be produced with better control of nanometer-scale geometries than Spindt tips without the use of

expensive tools and processes.

At Motorola, we have investigated the feasibility of large electron sources built with carbon nanotube emitters.

Since electron sources and display designs must satisfy numerous and sometimes conflicting criteria, it is important

to understand in detail the behavior of the electron source, so that optimal designs can be implemented. In this

work, we present results obtained from a detailed investigation of the fundamental field emission behaviors of

carbon nanotubes. We have applied this understanding to the development of prototype electron sources for field

emission displays. We present our recent progress toward a large area field emission display (FED).

NANOTUBE ELECTRON SOURCE FUNDAMENTALS

Designing a practical device using carbon nanotubes as the electron source requires a detailed understanding of

the nanotube current-voltage behavior and the nanotube lifetime. Unfortunately, nanotube field emission behavior is

not as simple as depicted by mathematical models like the Fowler-Nordheim expression. Complicated physical

phenomenon and interactions with the environment occur under conditions used in device designs. The goal of this

investigation was to obtain the necessary understanding of nanotube field emission by thoroughly characterizing

nanotube field emitter behavior over a much wider range of operating conditions than would be encountered in any

given application.

In recent years, the field emission properties of broad area films have been measured using a broad area

phosphor or a small diameter metallic probe. While these techniques provide engineering information about emitter

uniformity and current density, they contribute little to the understanding of field emitter physics. The primary

shortcoming is that the cleanliness of the field-emitting surface can neither be measured nor controlled. The

electrode geometry has poor gas conductance and the emitters and anode are generally not thermally cleaned. The

local environment at the field emitter is dominated by outgassing from electron-bombarded surfaces, rather than by

the pressure of the measurement chamber. Consequently, the field-emitting surface is unknown.

Recognizing that the surface cleanliness of the nanotube was an important experimental issue, we chose field

emission microscopy (FEM) as our experimental technique. The geometry of a field emission microscope magnifies

the electron beam as much as 1,000,000 times. Spatial variations in the electron beam are readily observed, and

these are related to surface contamination, adsorption and desorption events, and the electronic and atomic structure

of the field-emitting surface. In contrast to using a broad-area electrode technique, the field emission microscope

a Ken.Dean@motorola.corn

NASA/CP--2001-210948 708 August 6, 2001
usesitsgeometrytoenhancetheelectricfield.Theanodecanbespacedfromthecathodebymorethan5cmwhile

maintainingoperatingvoltagesintherangeof500to5000V. This,incombinationwiththeabilitytobakeouthe

anode, cathode, andvacuumsystem, provideawell-controlled vacuumenvironment. Thecathodeisalsoaffixedto

aheatingelement, allowingthetemperatureofthesubstratetobevariedfromwellbelowroomtemperature to>

2000°Cwithinamatterofseconds. Withathermal ceaningprocedure, controlofthevacuumenvironment, andthe

abilitytomonitortheemittersurface, important fieldemissionvariablescanbecontrolledprecisely.

Fieldemissionmicroscopyexperiments wereperformedinanultrahighvacuumchamber, whichincludeda

smallfield-emittingsamplepositionednominally2cmfromaphosphor-coatednode.Adetaileddescriptionfthe

apparatus andexperimental methodswerepresented elsewhere. 2Singlewallednanotubes (SWNT)wereattachedto

aWheaterfilament,whichallowedustoremovesurfaceadsorbatemoleculesthroughflashheatingof theW

filament.Nanotubes wereattachedtothefilament(withoutabinder)byrubbingthefilamentinnanotubepowder.

About25%ofpreparedsamplesemittedfromasinglenanotubeovertheentiremeasurable currentrange.Field

emissionpatternswererecordedusingacommercial CCDcameraandvideorecordingsystem. TheSWNTsamples

usedinthesexperiments originatedfromeitherRiceUniversity(90%puresingle-walled nanotubes afterfurther

purificationatMotorola) orMaterialsandElectrochemical ResearchCorporation(5%single-walled nanotubes by

volumeinamatrixof carbonsoot).Wefoundthefieldemissioncharacteristics fromthetwosourcestobe

indistinguishable.

Fieldemissionmicroscopy hasbeenusedtostudymetalsurfacesforover60years.Generally, theimage

producedbyfield-emittedelectronsonthephosphor screenproducesanintricatepatternwiththesamesymmetryas

themetalcrystal.Contrast arisesintheimageasaresultofdifferences intheelectronworkfunctionsofthecrystal

planesbisectedbythesurfaceandinthefieldenhancement differences betweenatomicstepsinthecrystalplanes. 3'4

Inthecaseofnanotubes trminatedwithcapstructures, TamuraandTsukadahavecalculatedthelocalstatesexistat

thepentagons inthenanotubecapwhichmightcreateenhancetunnelingattheselocations. 5 Theexistenceofend

stateshasbeenverifiedbyCarroll,etal.usingscanningtunnelingspectroscopy. 6 Consequently, onemightexpect

thatFEMimagesofnanotubes willproduceapatterncorresponding tothespatialdistributionoftheselocalcap

states.Inshort,nodesintheFEMimagemaybemapsofthepentagonlocationsonthecaps.Themostcommon

capstructuredepictedintheliteratureis the(10,0)cap,whichcontainsacentral pentagonsurroundedbyfive

symmetrically distributedpentagons. However, thereareatleastsevendifferent capstructures thatcanexistatthe

endofthe(10,0)nanotube, andthereareasmanyas7700possiblecapstructuresforSWNTswithdiameters smaller

than16A.7'8WeexpectedtoobservealargevarietyofnanotubeFEMimages, eachshowingatleastsixnodes.

OurfirstFEMexperiments withsingle-walled nanotubes producedimageswithone,two,andfournodes,with

twonodesbeingthemostcommon, andfournodesbeingrare(Fig.l).Overtime,theseimagesflickered, rotated,

andchangedtheirnumberofnodes.Alargepercentage oftheseFEMimageswerelopsided, aswouldoccurif the

axisofsymmetryofthefieldemittingsurfacewasnotalongtheaxisofsymmetryofthenanotube. Thistypically

occursinchiralnanotubes, wheretheapexisnotalongtheaxisofsymmetryofthenanotube. Lopsidedimages

wereexpected. However, thesmallnumberofnodes, theloworderofsymmetry, andtheinstabilityoftheimages

werenotconsistent withtheexpectedbehaviorfacappednanotubesurface.Intheliterature, nearlyidentical FEM

imageswereproducedonmetalsurfaceswhenmoleculesadsorbedontothesurface,suggestinga similar

phenomenonin naotubes 3'9Thebehaviorfadsorbates onmetalsurfaceshadbeenstudiedextensively, andseveral

otherbehaviorshadbeenidentified:(1)theimagenodeschangeovertime,(2)therateofchangeoftheimage

increaseswithtemperature, (3)theimageswereremovedbydesorbingthemolecule, leavingtheimageofaclean

metalsurface, and(4)theadsorbateimagesre-adsorbed ontoacleannanotubesurfacebysupplyingadsorbate

molecules. WeverifiedallofthesebehaviorsintheFEMimagesanddetermined thatadsorbatemoleculeswere

responsiblefortheone,twoandfour-lobedFEMimagesof ournanotubes 2Moreover, tunnelingthroughan

adsorbatemoleculeproducesapproximately 100timesthecurrent astunnelingthroughacleananotubesurface, but

thecurrentfluctuates.Removal oftheadsorbateresultsinloweremissioncurrentandhighlystabletemporal

behavior.

Thisinvestigationdemonstratedhatnanotubefieldemissionoccurspredominantlythroughadsorbate

moleculeswhentheoperatingpressureisabove10-9torr,thevacuumsystemandphosphor havenotbeenoutgassed,

theelectrodes arecloselyspaced, orthecathodehasnotbeenoutgassed above600°C.Theseconditionsarenotmet

inmostnanotubeinvestigations reportedintheliterature.It isalsonotclearthatheseconditionscanbemetin

practicaldevices.Consequently, field emissionfromadsorbatesis animportantaspectof bothnanotube

characterization andnanotube-based devices.Inaddition, thedominanceofadsorbatefildemissionundertypical

conditionsisnotlimitedtoSWNTsandMWNTs.Weseefieldemissionthroughadsorbatemoleculesinallother

carbonmaterialsthatwehaveinvestigatedin thismanner,includinggraphite,diamond-likecarbon,and

nanodiamond (Fig.2).

NASA/CP--2001-210948 709 August6,2001
D D

A,o.SO_ bo.t,_

Clean NT U

Figure 1. FEM images resulting from adsorbate molecules on SWNTs, as compared to the more complicated

pattern produced by a clean SWNT.

SWNTs Micro-crystalline graphite Nanocrystalline graphite film DLC Nano-diamond

Figure 2. Adsorbate images observed from various carbon materials.

We found that field emission through adsorbates leads to increased field emission current and increased noise,

but it also causes current saturation behavior. Collins and Zettl 1° and Bonard, et al., 11 reported current saturation in

the I-V characteristics of carbon nanotubes. Using the field emission microscope, we were able to measure the I-V

behavior of a single SWNT while recording the FEM image. We found that current saturation occurs in the I-V

behavior of an individual SWNT, and the onset occurs for currents between 100 nA and 300 nA per nanotube. The

combination of FEM images and I-V measurements detail the process (Fig. 3). Initially, field emission occurs

through an adsorbate. As the current increases, the adsorbate desorbs at a current of about 100 nA, leaving the

nanotube surface clean, as seen in the FEM image. The I-V behavior now follows that of a clean nanotube. The

composite behavior of the two curves produces current saturation. The case in Fig. 3a clearly illustrates the

mechanism, but it is not the most common scenario. Generally, the adsorbate moves around on the surface before

desorption, resulting in several states prior to desorption, each with a lower current. The resulting current saturation

curve looks much smoother than that in Fig. 3a. Under ideal vacuum conditions, a hysteresis occurs in the I-V curve

when the adsorbate is removed because the down I-V sweep tracks the 'clean' I-V curve. However, when a single

nanotube is operated in a poor vacuum environment, such as 10 -6 torr with unbaked vacuum system and phosphor,

re-adsorption is rapid and no hysteresis occurs for slow I-V sweeps (fig. 3b). In Fig. 3b the error bars represent the

range of current fluctuation measured at each data point. The magnitude and frequency of current fluctuation

increase for currents in the range of 100 to 300 nA. This occurs because the high current density condition excites

the adsorbate molecule in a manner similar to thermal excitation, causing surface diffusion of the adsorbate. With

enough excitation, the adsorbate molecule desorbs and the current becomes stable.

lO00

_l G Adso_ate!V1

i" -*"-- c:_an :v I t _ _

F ,,' // i v

_ i _ , :_ ._. ol ...... J / ..... I

400 s_/_oo 700 5o0 _oo\_,o_ 1_0o o

! Voltage (V)

a) ovTgg,,75g,,.,0gv 0v ,oov (b)

10 _

I........................ ;.................

{1

10 o i II

10 _u [ |

f

10 '_ _-

1,9 _2 i %_

10 _;

10 14 _ , ,

i000 i200 "i _,00 I,_0

- - i t,,, vo,tage tv)

1800 2tYa0

Figure 3. The field emission current of an individual SWNT saturates at ~ 100 nA because the adsorbate

molecule desorbs (a). The I-V curve follows the dashed line until desorption, where it then follows the solid

line for all voltages. Under non-ideal vacuum conditions, the up I-V sweep and the down I-V sweep both show

current saturation due to adsorbate surface diffusion, desorption, and readsorption (b).

NASA/CP--2001-210948 710 August 6, 2001
Therearemanyotherphenomena thtcausecurrentsaturationinothersystems, includingcontact resistance,

transport limitation,internal spacecharge, andvacuumspacecharge. Thefactthathe'clean' I-V curve does not

show current saturation to the destruction limit of the nanotube, even at much higher currents than 100 nA,

demonstrates that none of these other mechanisms are significant in this measurement. Calculations have shown

that vacuum space charge is not an issue and analysis of the electron energy distribution verified that other

mechanisms are insignificant.

Adsorbates can be removed from the surfaces of SWNTs and MWNTs by heating the emitter in ultrahigh

vacuum above 600°C for a few seconds. This procedure produces a clean nanotube surface with highly stable field

emission behavior. The current-voltage behavior is stable and repeatable over a large current range. 12 The I-V

characteristics of an individual SWNT produce a Fowler-Nordheim plot that was found to be linear over a large

current range (Fig. 4). The resulting field emission images of SWNTs generally show 4 to 12 nodes and the kinds

of complex patterns that could be consistent with the spatial distribution of electronic states in nanotube caps (Fig.

5a). The clean images of MWNTs generally show between four and six nodes, but the structure of the nodes are

quite different. The nodes are not solid and round as in the case of SWNTs, but rather they are pentagonal donuts

(fig. 5b). Interference fringes are also observed between the MWNT nodes. Another difference between clean

SWNT and MWNT images is that one adsorbate molecule tends to obscure all the nodes on a SWNT, but an

adsorbate molecule can exist on any or all MWNT nodes without affecting the others. This observation suggests

that nodes of MWNTs have a greater physical separation than the SWNT nodes. Because the shape and number of

MWNT nodes closely agree with the expected image arising from surface states at the pentagons, Saito, et al., have

proposed that these nodes are electronic states associated with cap pentagons. 13 The interpretation of SWNT FEM

images remains less clear.

106

v_ 107

10 s

8

10 9

10 lO

1400

e

1600 1800 2000 2200 2400

Voltage (V)

F_ e 1013

#e e

e.

_-10 t .,,.

1016 L

4104 5104 6104 7104

1N (V1)

Figure 4. The I-V and Fowler-Nordheim plots for an individual SWNT show stable behavior consistent

with a tunneling process.

.o. :iJBiiiiiiiiiiiiii : 

 i!i!i!!ii@iiii iiiiiiiiiiiiii!iiiiii ii! !ii iii i i i i i i ! !iiiii ,:,, i i! iiiiiiiiiiiiiiiiiiiiiiiiiiiiiiii

Figure 5. Field emission microscope images of clean SWNTS (a) and MWNTs (b).

While the SWNT emission current is stable over several orders of magnitude, we observed a current threshold

beyond which the emission characteristics degrade rapidly. Under high current conditions, the Fowler-Nordheim

plot deviated from linearity, the field emission patterns became diffuse or blurred, and rings formed around the field

emission image (Fig. 6). These phenomenon all indicate that the local temperature of the nanotube tip is increasing

substantially. 14 We estimate that when the current from a SWNT reaches the range of 300 nA to 2 _tA in a room

temperature experiment, the local emitting surface temperature exceeds 1600 K. At this point, the nodes in the FEM

image spin rapidly. Prolonged spinning degrades the current-voltage behavior of the nanotube, consistent with a

NASA/CP--2001-210948 711 August 6, 2001
shorteningprocess.Ourresultsindicatethatthemaximumfieldemissioncurrentthatcanbeextractedfroman

individual SWNTislimitedb_thermally-assisted fielevaporationof thenanotube.Weestimatehemaximum currentdensitytobeinthe10to 109A/cm2range,basedonournanotubedimensions. Incontrast tometals, the

thermal fieldevaporation processisgentleandit doesnotcreateadestructivearcdischarge.

Thegasesintheoperatingenvironment alsoimpactthestabilityofSWNTfieldemitters.Inultrahighvacuum,

wefindalmostnodegradation. Forexample, inFig.7,tennanotubes fieldemitwitha100%dutycyclefor>350

hours.Inanothersetofexperiments, weoperatedthesameSWNTemittersfor40hoursinvacuumenvironments

containing10-6torrH2,10-7torrAr,10-7torrH20,and10-7torr02.Nodegradation wasdetectedduringoperationin

H2orAr. Somedegradation occurredinwater,andoperationin10-7torr 02permanently reducedtheemission

current by75%.Weconcludethatoxygen-containing speciespermanently degradeSWNTfieldemissionbehavior.

Whilethismakesthefabricationofdeviceswithlonglifetimesmorechallenging, otherinvestigators hadindicated

thatMWNTsshowbetterstabilitythanSWNTsintheirmeasurement environments.

Fromourinvestigations ofthefundamental fiedemissionbehaviorsofcarbonnanotubes, wederiveseveral

designrulesforlargeareaelectronsources.First,adsorbates willcontributetofieldemissioncurrentfluctuation,

eveninsealedvacuumdevices.Designsshouldincorporateamultiplicityofemittersperelement tostatistically

averageoutthefluctuation, therebyproducingmorestablecurrent.Currentfluctuationwillalsobesmallerinan

electronsourcedesignedsotheSWNTsdonotemitcurrentintherangeof100nAto300nA. Clearly,electron

sourcedesignscannot calluponindividual SWNTstoemitmorecurrent than300nAto2_tApernanotubewithout

irreversibledamagefromfieldevaporation. Thiscriterionsetsalimitfortheminimumnumberoffieldemitting

nanotubes necessary toachievecurrentdensityrequirements. Finally,oxygenandoxygen-containing speciesmust

beremovedfromtheoperatingenvironment ofSWNTfieldemitterstoensurelonglifetimes.

<

t--"

v

off

,¢

-m

O

1000

100

10

1

0.1

500

.... i .... i .... i .... i .... i ....

600 700 800 900 1000 1100

Voltage (V)

1030V 1050V 1050V

Figure 6. The current-voltage behavior of an individual SWNT at high currents shows thermal field emission

and the FEM image blurs. Still higher currents induce movement of the nodes in the FEM image and cause

irreversible degradation of the I-V characteristics, consistent with field evaporation. Arrows in the 1050 V

FEM image depict the motion.

5

4

3

2

1

0

0 50' 100' 150' 200' 250' 300' 350

Time (hours)

Figure 7. The field emission current from 10 SWNTs shows no long-term degradation in ultrahigh vacuum

under continuous (non-pulsed) operation. Current fluctuation results from adsorbate interactions.

NASA/CP--2001-210948 712 August 6, 2001
LARGEAREADEPOSITIONOFNANOTUBES

Fieldemissioncathodesinthefieldemissionmicroscopeproducecurrentinthemicroampere range.Most

cathodeapplications requirealargeareacathodecapableofproducingcurrentsintherangeof10mAtomorethan1

A. Inthissection, wedescribercent progresstowardsalargeareafieldemissionsourceforfieldemissiondisplays

thatiscapableofproducingsuchcurrents.Forfieldemissiondisplays, thereareseveral otherspecifications i

additiontothecurrent levelthatmustbesatisfiedsimultaneously. Theelectronsourcesmustbebothcosteffective

andmanufacturable ovrlargeareas,andtheymustbespatiallyuniformtowithin3%betweenadjacent pixels.

Economically viablelargeareadisplaysmustbefabricatedonglassubstrates, andthisglassmustbeoneofthe

typesalreadymass-produced forLCDdisplays, windows, orplasmadisplays.Theglasseshavestraintemperatures

between~ 530°C(sodalimesilicate)and~ 666°C(aluminoborosilicate), but deformmeasurablyat lower

temperatures givensufficient time.Thechallengeistodeposit nanotubes ontoalargeareasubstratewithoutusing

excessivetemperatures.

AtMotorola, wehavedevelopedtwomeansofdepositingfield-emittinganotubes overlargeareas.Either

methodcanbeusedtoplaceemittingnanotubes onthemetalizationpatternthatwedevelopedforfieldemission

displayswitha2.5"diagonal dimension. TwoofthesepatternsareincorporatedoneachOA10glasssubstrate

(fromNipponElectricGlass).In thefirstmethod, wegrownanotubes usingchemical vapordepositionontoa

patternedsubstrate. Acatalyst layerisplacedontothesubstrateinthepixelregionstocontrolthelocationofthe

grownnanotubes. TheCVD-grownnanotubes havediameterssmallerthan3nmandlengthsexceeding1.5_tm

(Fig.8b). Fieldemissionisuniformoveralargearea,asseeninalargeareadiodetest(Figs.8a,c).Noballast

resistorsareincorporated inthisteststructure, sothefieldemissionuniformityreflectsthelengthuniformityofthe

nanotubes. Indiodemode,wetypicallymeasurecurrentdensitiesxceeding1.4mA/cm2ontheanodeor8mA/cm2

fromthenanotube-containing areasofthecathode. Thiscurrentdensityissufficient toproducehighvoltagefield

emissiondisplaysforofficeandindirectsunlight applications. Thepreliminaryexperimental systemsupportsa6"

growtharea,butthisareacanbescaledsignificantly.

Wedevelopeda secondmethodfordepositingfield-emittinganotubesoverlargeareasubstrates in which

nanotubes areincorporated intoabindersystem.Thenanotube/binder mixturecanbeprintedforloweresolution

geometries, adescribedbyLee,etal.15Sinceourtestvehicleisrelativelyhighresolutionanditrequirespatterned

dimensions lessthan 40 _tm, we incorporated the nanotubes into a photosensitive binder and patterned them to

produce discrete pixels. The nanotubes, as patterned, do not field emit well. We believe that they are pinned against

the surface by the binder or by surface forces. An activation step is required to produce the desired field emission

behavior. Others have reported the need for an activation process. 15 We developed an activation process that is

completely compatible with the fabrication of our triode and allows nanotubes to be deposited at any point in the

process. We have an in-house capability for activation over a 50" diagonal display.

The nanotube-in-a-binder method for producing large area electron sources promises to be a very inexpensive

and scalable process. A primary cost factor, that of capital equipment, is insignificant when compared to the

equipment required to make Spindt tip field emission displays or liquid crystal displays. Consequently, we expect

this nanotube method to be viable technology for displays with diagonal dimensions greater than 25 inches.

Figure 8. CVD-grown nanotubes on OA10 glass (a). Image produced on a phosphor screen during a field

emission measurement of the sample in (a) where the emitting region is a 2.5" diagonal area patterned with ~

49,000 pixels (b). Close-up view of patterned diode array showing pixelation (c). The diode has no ballast

resistors.

NASA/CP--2001-210948 713 August 6, 2001
NANOTUBEDISPLAYDEVICES

Oncethetechnologyexiststo depositlowcost,largeareananotubefieldemitters, thenextchallengeisto

integratetheseemittersintoa displaystructurewitha gatingelectrode.Whilesimplyfabricatingthegating

electrodeinawaycompatiblewithnanotubefildemissionischallengingenough, alltheotherpartsandprocesses

ofthedisplaymustbecompatiblewiththisprocess.Theincorporation ofspacers, getters,ealingmaterials, and

highvoltagetechnologyallplacelimitationsontheprocesses thatcanbeusedtofabricatethedevice.

Thedesignofthegatingelectrodewithnanotubes oftheproperdimensionhassignificant implicationsforthe

overallcostoftheFED.Thisdesignaffectsthedrivingelectronicsthatwillbeusedtoproduceanimageinthe

display.Thevoltagethatisrequiredtoturnapixelfromits'on'to'off' states, orswingvoltage,determines the

requiredsemiconductor technology, andhencethecostofthedriverelectronics. Plasmadisplays, forexample,

swingmorethan150V,andtheyrequiresemiconductor technologywithdielectricsolation.Driverelectronics for

anHDTV(720x 1280x3)plasmadisplaycurrentlycosthemanufacturers thousands ofdollars.16Inthefuture,

drivercostsmightdecreaseto$1200.Thus,evenif thecapitalequipment, laborandmaterialsrequiredtofabricate

theFEDarefree,themanufacturing costofadisplaywitha150Vswingwillstillcostmorethanathousanddollars

tomanufacture, andthesalepricewillbesignificantlylargerthanthat.Wehavedetermined thatwithaswing

voltageofabout50volts,wecanusedriverelectronicsdevelopedatMotorola,ndforthesameHDTVresolution,

thedriverelectronics willcostlessthan$100toproduce.

Forourfirstattempttointegratenanotubes intoanFEDtriode,weutilizedexistingSpindttiptechnology.

Catalyst waspatternedatthebottomofthewellsofSpindt-tipcathodesinplaceofthemetaltip. Nanotubes were

grownintothedeviceusingtheCVDtechnologydescribedabove.AnexampledeviceisshowninFigure9. The

128x 128pixelarrayproducedananodecurrent densityof1mA/cm2witha30Vgateswing.Thiscurrent density

wassufficient tomakeahighvoltagefieldemissiondisplayforofficeandindirectsunlightapplications. Thelow

voltageswingcouldaccommodate lowcostdriverelectronics. Whilethisprocessdemonstrated thecompatibilityof

nanotubes withatriodestructure, itdidnothaveamanufacturing costthatwaslessthantheSpindttipfieldemission

display.Thecapacitance oftheSpindtdeviceisalsoobjectionable forlargearea(>25"displays).

Wedesignedatriodestructureandafabricationtechnology especiallyfornanotubefildemitters.Theprocess

isscalabletolargeareasandthecostofcapitalequipment issubstantially smaller.Wesubstituteelectroplating for

someexpensivethinfilmprocesses. Thedevicecapacitance hasbeenreducedtotolerablevelsforlargearea

displaysbyusinga10_tmcathodeplane-to-gate electrodespacing. Thetriodestructurewasimplemented withboth

theCVD-grownnanotubes andthenanotubes mixedwithbinder.AsmallregionofatriodewithCVD-deposited

nanotubes isshowninFig.10toillustratethepixelation.Comparetothedioderesultsinfigure8c.Noballast

resistortechnologyhasbeenusedinthisprototype.Thenanotubes-in-binder technologywasappliedtocreatea

fully-sealedfieldemissiondisplay,completewithspacersandRGBcolorphosphors.Witha70V swing,the

displayswitchedcurrentbetween0.01mA/cm2and3mA/cm2. Thisis adequateto produceindirectsunlightreadable, high-voltage fieldemissiondisplayswithlowcostdriverelectronics. Theseprototypeswereoperatedwith

ananodevoltageof5KV.

CONCLUSIONS

Rigorousinvestigationf thefundamental fieldemissionbehaviorof carbonnanotubeshasimprovedour

understanding ofthefieldemissionprocess. Withthisunderstanding, wehavegenerateddesignrulesthatallowus

toimprovecurrent stability, provideanenvironment forlonglifetimeoperation, andengineer deviceswithsufficient

currentdensityatlowvoltages. Prototypetriodedeviceshavebeenbuiltthatoperatewithlowswingvoltagesand

produceenoughlightfortypicaldisplayapplications.

NASA/CP--2001-210948 714 August 6, 2001
D ...................................

:" 3 / v'_ "

¢,-

,= 2 / _ ;

0:_ _ __.__ . . . '.

40 50 60 70 80 90

(a) (b) vg

Figure 9. A 128 x 128 triode array with CVD -grown carbon nanotube emitters. The image on a phosphor

screen was taken with a 3 mm anode-cathode gap (a). The current-voltage curve shows a 30 V swing is

required to produce sufficient current density.

................................. 1

J

25 50 75

Vg

100

Figure 10. Close-up area of triode array with CVD-grown nanotubes (a). I-V curve for a triode using the

nanotubes-in-binder technique (b).

REFERENCES

1. Chalamala, B.R.; Wei, Y.; and Gnade, B. E.: "Fed up with fat tubes". IEEE Spectrum, 42, 35 (1998).

2. Dean. K.A. and Chalamala, B.R..: "Field emission microscopy of carbon nanotube caps". J. Appl. Phys. 85,

3832 (1999).

3. Dyke, W.P.; and Dolan, W.W.: "Field emission". Advanced Electronics and Electron Physics, 8, 89 (1956).

4. Gomer, R.: "Field emission and Field Ionization "'. American Institute of Physics, NY, 1993.

5. Tamura, R.; and Tsukada, M.: "Electronic states of the cap structure in the carbon nanotube". Phys. Rev. B 52,

6015 (1995).

6. Carroll., D. L. et al.: "Electronic Structure and Localized States at Carbon Nanotube Tips". Phys. Rev. Lett. 78,

2811 (1997).

7. Astakhova, T. Yu. et al.: "Numerical generation ofnanotube caps". Sci. and Technol. 7,223 (1999).

8. Osawa, E. et al.: "Analysis of the growth mechanism of carbon nanotubes by C2 ingestion". Sci. Technol. 7,239

(1999).

9. Morikawa, H. et al.: "Current Voltage Characteristics of Field Emission through Phthalocyanine Molecules on a

Tungsten field emitter". Jpn. J. Appl. Phys. 36, 583 (1991).

10. Collins, P.G.; and Zettl, A.: "Unique Characteristics of cold cathode carbon-nanotube-matrix field emitters".

Phys. Rev. B 55, 9391 (1997).

11. Bonard, J.-M. et al.: "Field emission properties ofmultiwalled carbon nanotubes". Ultramicroscopy ,73, 7

(1998).

12. Dean, K.A.; and Chalamala, B.R.: "Current saturation mechanisms in carbon nanotube field emitters". Appl.

Phys. Lett. 76,375 (2000).

13. Saito. Y.; Hata. K.; and Murata, T.: "Field emission patterns Originating from pentagons at the tip of a carbon

nanotube". Jpn. J. Apl. Phys., Part 2, 39,271 (2000).

14. Dyke, W.P., et al.:"The field initiated vacuum arc. I. Experiments of arc initiation". Phys. Rev. 91, 1043 (1953).

15. Lee, N.S., et al.: "Carbon Nanotube-based Field Emission Displays for Large-Area and Full-Color

Applications". J. Appl. Phys. p.1 12B, 7154 (2000).

16. Park, K. Y.: "Why do PDP's cost so much". Information Display 7/98, 10 (1998).

NASA/CP--2001-210948 715 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

MULTIWALLED CARBON NANOTUBES GROWTH IN NANO-TEMPLATE AND

THEIR APPLICATION TO A FIELD EMISSION DEVICE

Soo-Hwan Jeong, Hee Young Hwang and Kun-Hong Lee*

Electrical and Computer Engineering Division and Department of Chemical Engineering, Pohang University of

Science and Technology, San31, Hyoja-Dong, Nam-Ku, Pohang, Kyungbuk, 790-784, Korea

Won-Ki Cho and Kwang-Young Kim

LG Electronics Inc., 16 Woomyon-Dong, Seocho-Ku, Seoul 13%724, Korea

ABSTRACT

Carbon nanotubcs (CNTs) have drawn much attention because of their unique physical properties and wide

variety of applications. One potential application is the field emitter tip. Advantages of the CNT as a field emitter

tip include a small radius of curvature, a high aspect ratio, high chemical inertness and mechanical strength. Many

researchers have fabricated densely packed CNTs on the electrode for this purpose. However, this approach has a

negative effect on field enhancement due to the field screening effect.

Recently, the growth of aligned CNTs into the pores of anodic aluminum oxide (AAO) using bulk aluminum is

reported. AAO can be prepared by anodizing AI in various electrolyte solutions using DC current. By varying the

anodizing conditions, one can control the density, the diameter, and the length of pores. Therefore, specific

properties of CNTs can also be controlled by growing CNTs with AAO templates. Furthermore, fabrication of the

Spindt-type triode structure is straightforward and achieved with minimum effort.

Two types of aluminum samples were used for the starting material. One was a 12gin-thick sputtered A1 on Si

wafer and the other was a sputtered A1 film of 400 nm thickness on Nb-coated Si wafer. AAO templates were

fabricated by anodizing Al films. After the Co catalyst had been electrocliemically deposited into the bottom of the

AAO template, CNTs were grown by the catalytic pyrolysis of C2H2 or CO at 650°C. Overgrowth of CNTs on the

AAO templates was observed. The growth of CNTs in an AAO template is significantly influenced by the presence

of H:. With no hydrogen added, CNTs did not grow over the AAO template, while CNTs were overgrown with H2.

In fact, the growth of CNTs in the AAO template with a Co catalyst is believed to be a complex process, involving

competitive hydrocarbon decomposition by both cobalt particles at the bottom of pores and the AAO template itself.

The density of our samples is approximately 107-10 s tips/cm 2. The diameter of the CNTs strongly depends on the

size of the pores in the AAO template. The electron field emission measurements on the samples showed a turn-on

field of 1.9-2.1 V/p.m and a field enhancement factor of 3360-5200.

To avoid a field screening effect, a low density of CNTs was favored. Our observation concerning the low turnon field and high field enhancement factors is explained in terms of a low field screening effect. This approach could

open new opportunities for application of CNTs in field emission device.

Keywords: Carbon nanotube, Fieldemission, Nano-template, Anodic aluminum oxide, Anodization

*Corresponding author, Mailing address : Electrical and Computer Engineering Division and Department of

Chemical Engineering, Pohang University of Science and Technology, San31,

Hyoja-Dong, Nam-Ku, Pohang, Kyungbuk, 790-784, Korea

e-mail address : ce20047@postech.ae.kr

Phone number : +82-54-279-2720

Fax number : +82-54-279-8298

NASA/CP--2001-210948 716 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FIELD EMISSION FROM CARBON NANOTUBES GROWN ON PLASMA

TREATED NICKEL/SILICON SUBSTRATE

W. K. Wong, Frederick C. K. Au, L Bello, C. S. Lee, and S. T. Lee*

Center of Super-Diamond & Advanced Films (COSDAF) and Department of Physics and Materials

Science, City University of Hong Kong, Kowloon Tong, Hong Kong

ABSTRACT

Carbon nanotubes were synthesized on silicon substrates coated with thin nickel film using

microwave plasma-enhanced chemical vapor deposition. Hydrogen plasma treatment was used to

vary the surface morphology of the thin nickel films. Scanning electron micrographs showed that the

size and distribution of the nickel clusters varied with the time of plasma treatment and bias voltage.

The nickel clusters were uniformly distributed over the silicon surface and expanded in their size as

the hydrogen plasma treatment progressed in time. Both the growth and field emission properties of

carbon nanotubes were strongly influenced by the surface morphology of the thin nickel films. The

field emission properties of carbon nanotubes were characterized by using a scanning probe and

diode-structured field emission systems. In addition, the structure of carbon nanotubes was studied

by high-resolution transmission electron microscopy.

Keywords: carbon nanotubes, chemical vapor deposition, hydrogen plasma treatment, field emission, and

electron microscopy.

*Corresponding author

E-mail address: apannale@cityu.edu.hk

Mailing address:

Department of Physics and Materials Science,

City University of Hong Kong,

Kowloon Tong,

Hong Kong

Fax: +852-2788-7830

Tel: +852-2788-9606

NASA/CP--2001-210948 717 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FIELD EMISSION PROPERTIES OF ULTRA-FINE MULTI-WALLED CARBON

NANOTUBES GROWN BY MICROWAVE PLASMA CVD

Weihai Fu, Bing Xiao, and Sacharia Albin*

Microelectronics Laboratory

Department of Electrical and Computer Engineering

Old Dominion University, Norfolk, Virginia 23529

ABSTRACT

Multi-walled carbon nanotubes (MCNTs) were grown using microwave plasma enhanced chemical vapor

deposition in a gas mixture of H2 and CH4. Thin NiCr films deposited by RF sputtering were used as catalyst on

thermally oxidized silicon substrates. MCNTs were characterized using scanning electron microscopy, highresolution transmission electron microscopy, Raman spectroscopy, and I-V measurement. Vertically aligned growth

of straight and defect-free MCNTs with outer diameters as small as 7 nm were observed. The growth of ultra-fine

MCNTs was due to the small catalyst particle size, which was in the range of 10 nm based on scanning tunneling

microscope study. The emission current from an area of 4x4 mm 2 reached more than 2nA at an applied voltage of

100V with an anode-cathode separation of 100 gm. The current increased exponentially with the applied voltage

and reached 1 mA at 400V. The high emission current dropped and stabilized from 1.07mA to 0.95mA within a 3-

hour long-term emission stability test. However, at lower emission current of 52 gA, no such decay was found over

15 hours continuous operation. The low turn-on field and stable emission were attributed to small diameter of

MCNTs that were vertically aligned.

Keywords: carbon nanotubes, field emission, MPCVD.

INTRODUCTION

Since the discovery in early 1991(ref. 1), carbon nanotubes (CNTs) have been extensively studied for their

unique properties and potential applications in flat panel displays (ref. 2), hydrogen storage (ref. 3), nanoprobe and

nanoelectronic devices (refs. 4 and 5). The unique geometric structure of CNTs combined with high electrical

conductance makes them an ideal choice for cold-cathode electron emission sources, where an applied electrical

field instead ofthermionic heating provides the energy for electron emission. Turn-on field lower than 1V/gm and

emission current density of 1A/cm 2 have been achieved from multi-walled CNTs (MCNTs) (ref. 6). MCNTs in this

study were produced by catalyst assisted thermal chemical vapor deposition (CVD) on conductive TiN substrate. A

high current field emission electron source made of aligned CNTs was successfully fabricated by De Heer et al

(ref.7). Cathode ray tube lighting elements made of aligned MCNTs have shown life times exceeding 10,000 hours

without significant decay (ref. 8). All these results indicate that CNTs are very promising for field emission

applications.

Low turn-on field, high current density, and stable emission in less stringent vacuum environment are desired

properties of CNTs for field emission applications. Single-walled CNTs (SCNTs), due to their fine tube diameters

and defect-free walls, generally display low turn-on field (ref. 9). However, this advantage over MCNTs will be

partially offset by the fact that these fine tubes are often put together forming a larger tube bundle. The screen effect

induced by these closely packed tubes will reduce the field enhancement factor of each tube. The thin tube wall and

cap are prone to damage caused by residual gas ion bombardment and emission current overheating. The

consequences on emission current are large fluctuation and decay with time. MCNTs, on the other hand, are more

stable but have higher turn-on voltages mainly due to their large diameters and defective tube walls. SCNTs are

mostly produced by arc discharge and laser ablation, where the resulting tubes are non-selective on substrate and

often in bundles entangled together (ref. 10). Extensive post-deposition processes are required to make SCNTs

useful for field emission applications. However, MCNTs grown by CVD using catalyst assistance can be patterned

and grown vertically aligned to the substrates to form arrays of CNT emitters (refs. 11 and 12). Considering both

* Corresponding Author. Phone: 757-683-4967, Fax: 757-683-3220, E-mail: salbin@odu.edu

NASA/CP--2001-210948 718 August 6, 2001
processsimplicityandelectronemissionproperties, it isdesirabletoproducetheCNTswithemissionsproperties

comparable toSCNTsandyetcanberealizedbythemethodsuedtoproduceMCNTs.

In thispaper,wepresentdetailsof CNTsgrownbymicrowaveplasmaCVD(MPCVD)usingNichrome

(Ni:Cr80:20)asacatalyst andtheirfieldemissionproperties. Ourmethoddemonstrates bothlowturn-onvoltage

andstablemissionathighcurrent densitymainlyduetothefineanddefect-freeMCNTsverticallyalignedtothe

substrate.

EXPERIMENT

CNTssamplesweregrownonoxidizedsiliconsubstratehavinga 100nmthickthermaloxidelayer.A

Nichromefilmof40nmthickdepositedbyRFsputteringonthesubstratewasusedascatalyst forCNTgrowth.The

surfacemorphologyofthecatalyst filmwascheckedbyacanningtunnelingmicroscope(STM).Siliconsubstrates

wereeitherscribedinto4><6mm2rectangular chipsorpatternedusingconventional photolithographyrocess

followedbyconcentrated HC1etchingto formarraysofNiCrsquares. Substrates wereloadedintoaMPCVD

system(ASTeX).Afterinitialpumpdownbyaturbomolecular pumpbackedtwostagepumpingsystem, the

graphitesubstrateholderwasinductivelyheatedupto600°C.Researchgradehydrogen(99.995%) wasleakedinto

thevacuumchamber ataflowrateof300sccmandthechamber pessurewasmaintainedat20Torr.Microwave

plasmawasignited,andtheprocessgas,CH4,atflowrateof90sccmwasadmittedtoraisethechamber pressureto

30Torr.CNTgrowthlastedfor10minutesatamicrowavepowerof800W. ToverifyCNTgrowth,wecleaved

samplesandmountedthemverticallyunderscanningelectronmicroscope(SEM)toviewthecrossection.CNT

structures anddefectswerestudiedusinghigh-resolutionransmission electronmicroscopy (HRTEM).CNTswere

detachedfromsiliconsubstrates bysubmergingethanol solutionwith ultrasonic agitation. The resultant CNTsethanol mixture was transferred onto standard copper TEM grids coated with holely carbon film. A Nicolet Fourier

transform (FT) Raman spectroscope with a 1064 nm Nd:YAG laser excitation source was used to characterize the

tube structure.

Field emission I-V characteristics of CNT samples were obtained under a vacuum of 10 -7 torr through a diode

type measurement setup. This arrangement gave a nominal overlay area of 4x 4 rnm 2, which was deemed as the

effective area for emission current density calculation during the I-V analysis. The gap between the anode and CNT

surface was adjusted to be 100 bun. A 8-M_ ballast resistor was used at the initial stage of I-V testing to prevent

current run-away. However, this resistor was replaced by a smaller one of 100 K_ after the sample bum-in. The

voltage drop on the ballast resistor was excluded from calculations based on I-V characteristics.

RESULTS AND DISCUSSION

Fig. 1 shows an SEM micrograph of the CNTs grown by our MPCVD process. Vertically aligned CNTs of 3

gm in length can be clearly discerned from the micrograph. The alignment of CNTs is due to the self-bias induced

field at the sample and plasma interface (ref. 13). Also from Fig. 1, fine tubes protruding from the film (shown by

the arrows) are observed. Fig. 2 is the top view of an array of CNTs grown on patterned Nichrome film using the

same growth conditions. No CNT growth is found at regions where catalyst film had been removed. The result

demonstrates that CNT growth is strongly catalyst dependent.

°o.°°°.%%.°°°,,...,.°°oO.,°°°°,°..%,,% ,,.o°....°°°o...%,j° °°°°°°%°°°°°,.,.

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

..e....,....,.,.._...,...;._.-.;.I,2.....................:,%-.._,.,.¢......%..,

_.'.4.<:N:...:.:._..........×......-_:.._:...-.........:.....:.-.-.

:%%..%%..%%1_.%_.%% .%%%...,%%%%%,,%%%_.,

:..,.'.,2:_2_%,.:%+:.:.x.:.N.>.'.._x,x.,,:._:. '°%?°*_%'o%_ °%?°'j'j?o'_'_°%%? °? °?°%?.'. _° • • ° °_

FIG. 1. SEM micrograph of CNTs grown by

MPCVD using NiCr as catalyst.

II::::.:::::::::-::_:-:.::_:::-8)l_.:._S:._:._:_.*,_

":'.".'.'.%':': ",',*:':'.°l' °l. ..°._;%.'°'.°° ° • ° •

..........°-.-.° ._**°,-,_ -.._°. ¢,°%_-, ,° . ,'._ %,

• .o%* °°°°%° • °,. % °o.%_ • °.. ° • ° _o_

• ..%..%.,,.

FIG. 2. Top view of patterned CNT array grown

by MPCVD using NiCr as catalyst.

NASA/CP--2001-210948 719 August 6, 2001
......,,......... +¢....o.°°.... °_, ...... I.-4_,..........°.: a575_ , °%°°°, ,°.%°°°°°° .*.OoOo.°°°o°°°°°oo%¶ p°°.% !_,%o%OoOoOoOo%O °°°°°°° _°oO°°°°°°°°°°o°°%°=°=OoO°%°o%°_ _o-°_%o° o%%%o°°°°_°

,:.:::.::::.:::::::::::.:.:.:+::?:.::_::.:::+'.-°:,:!_

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:::":'Y: :::::::::::::,':::::_':_::::::-"_::::::::'2i&_

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i_,_•%•°%1°%° • • ° • • • _ • • • ° ° °°°%°_%%%°_% o°_ ',% _ o • ° • ° ° • ° °

%•_•OoOo• %o %%%Oo-_°o,%%Oo°oO.O°Oo%%%%%%°,_o. _%, _ _. _*o%Oo°°Oo*°°.

:.,_°;+:._•:_°;°_:•'..•°;o:.:o'°:°:+•.;.',•_4o5,'_',_ ,_,;. X.;•:°:.;°:°:.:.:+

_o%%%°•Oo°oO_'_Oo•°Oo%%_%_%o_'_._'_o%•_o'_ ;%. _o•%°.°°o°O.°o%. °..

•,°,_°•o°•o•,,•,Jo°°°°°.° ,,°o,% _,,_°' .° °, , o. ° _.-....°.,_..

o= io°%_•_lJ o_oo °, ,_, ,°oo°._o. ..... ,

,°•%'°*,0°•• _. ,,'_ ,%• %,_, "_p_._, °'.'.°._°°.'.*_°.'.'.'.'_°°'.'

,_-.;.;.;•_._, ,,,°;.. ,_.;°:....°:°:.?:+:.:°:.x°:°x

_%o°_°% %°o%%%%°°°oO°°oO,%Oo °°°°.•°o • °°_%o° _°_ •

FIG. 3. TEM micrograph of CNTs produced FIG. 4. HRTEM of a fine CNT produced

by MPCVD using NiCr as catalyst, by MPCVD using NiCr as catalyst.

Unlike the CNT grown by arc discharge or laser ablation, where no tube alignment and site selection can be

achieved, our method shows that the growth can be well controlled and thus, the process is feasible for making CNT

based microelectronlc devices and sensors. Fig. 3 shows a typical TEM micrograph of our MPCVD CNTs at low

magnification. The two arrows labeled A and B indicate two different types of CNTs, which have been identified

among all our CNT samples.

Type A shows tube diameters mostly larger than 20 nm and included catalyst particles within the tubes. These

tubes are found to contain various growth defects and irregular shapes. Bamboo type structures and catalyst filling

are also observed among this type of tubes. However, Type B shows totally different CNTs with long straight tubes

that are free of defects along the length. Under HRTEM micrograph shown in Fig. 4, the tube outer diameter is

measured to be 7 nm with seven graphene layers separated at 0.34 nm. CNTs of such small diameter have not been

reported for MPCVD growth. Tubes similar to Type B have been found among those grown by arc discharge (ref.

1), where the tube growth rate is much faster than that in the MPCVD process. STM image of Nichrome catalyst

film after sputtering deposition is shown in Fig. 5. The average catalyst particle size is around 10 nm. Smaller size

particles can also be seen in the sputtered film image. The small NiCr particle size is due to low power density

(0.5W/cm 2) and short deposition time (4 minutes) used for sputtering, where substrate heating is minimized. It has

been demonstrated that CNT diameters were directly related to the catalyst particle size in MPCVD growth (ref. 14).

Therefore, we believe that the growth of ultra-fine CNTs is the result of small catalyst particle size.

Results from FT Raman spectroscopy, shown in Fig. 6, reveal two distinct peaks at wavenumbers of 1286 cm -1

and 1583 cm -1. The peak at 1583 cm -1 is due to the strong, high frequency, in-plane stretching mode (E2g2 or Gpeak) of crystalline graphite (ref. 15), which is believed to be the structure of the CNT walls (ref. 16). Another

Raman peak from our CNT sample at 1286 cm -1 is found to be deviated from the typical Raman peak of disordered

graphite (D-peak), which is at 1360 cm -1 obtained mostly with Ar or Kr laser excitation sources (ref. 15). The strong

D-peak is believed to be the result due to carbonaceous particles, defective tube walls and ends, and finite size of the

crystalline domains ofnanotubes (ref. 15). It has been found that the disorder-activated D-peak could have a strong

......._.....:.................:..............:.-.-.....

::::::::::::::::::::::::::::::::::::::::::::::::::::::::

.:..Oo_.o-...•;o:.;...., :_:°:.:.:.:.,.._.._.,°.;... :+::....b • ..:,. :,,..

::::::::::::::::::::::::::::::::::::::::::::: .;_;=;,_.::Q:::.:::

.....o.-,.-o-.OoO,,.,°°...°..,.°..°,.,_......,,, °,....°,-,,....-.,:,o.

..:.:..............:.. _........° °°...........:...°..,:....:.........:.: • °%% .%°°%°. _..°.. o .'.% °_o° °.°o°% .% °°o °

!:i:!.i:i::'!'..'NN:i:i:???!'_4!.'!:!'!:i:i:;i:?:i:i:!:!:i:?_.!:!:!:

.<::.v:::: !ig.if:.!g.igi::?::igg':ii'.i!iiiggili i i!!!:::iii

i.:iiiiix_,.'!?:iiiiiiil .,:::...::!:?i'?!:?:m2_ _m:

0.8

_0.6

0.4

0.2

0

1000

1286

1200 1400 1600 1800 2000

wave number (cm 4)

FIG. 5. STM micrograph of RF sputtered

NiCr film surface morphology.

FIG. 6. Raman spectrum of CNTs grown

by MPCVD using NiCr as catalyst.

NASA/CP--2001-210948 720 August 6, 2001
shiftupto50cmq/eV(ref.15)duetotheexcitinglaserphotonenergy, whichcouldaccount fortheobservedDpeakshiftfromoursample, sinceweusedanIRlaser.SimilarD-peakshifthasbeenshowninreference14for

CNTsgrownbyMPCVDwhenusinga1064nm(Nd:YAG)lasersourceforRamanexcitation.Thelargeintensity

ratioofD-peakoverG-peakfromourCNTsampleisexplainedduetothesmallplanarcoherencelengthofthe

disorderedgraphiticmaterial inthefilm,wheretheD/Gratioisinverselyproportional totheplanarcoherence length

overarangeof2.5-1000nm(ref.15).TheMPCVDprocessusedtogrowCNTsisbasicallythesameasthatused

fordiamondgrowthexcept thehighmethaneconcentration andtheuseofNiCrcatalyst (ref.17).Therefore, itisnot

surprisingto frednano-sizedcarbonparticleswithcrystallinestructureincludedinthegrownfilm,whichwould

contributetothestrongD-peakintheRamanspectrum.

FieldemissionfromMCNTscanbeanalyzedusingthetheoryproposedbyFowlerandNordheim(ref.18).The

simplifiedtheF-Nequationisasthefollowing(ref.19):

I=aV2exp(-b), (1)

where I is the emission current, V is the anode voltage, a and b can be written as the following equations:

1.56 xl0- 6o_fl 2 (10.4

a- 1.10 d2 exp[_-_-), (2)

b = 6.44xlO703/2d/fl (3)

Variants in Eq. (2) and (3) represent emission area (a), emitter work function (_, separation between anode and

emitter (d), and field enhancement factor (fl). From Eq. (1), a plot of ln(I/V:) versus (1/V), called the F-N plot, will

result in a straight line of slope (-b). For a given field emitter, the work function of the emitter, the distance to the

anode, and the field enhancement factor generally do not change during the I-V measurement. Therefore, slope (-b)

is a fixed value as can be seen from Eq. (3).

Typical Field emission I-V characteristics from our CNTs sample at low current are shown in Fig. 7(a). The

emission current increased exponentially from 2.7 nA to 48 gA as the electrical field was raised from 1V/gm to

2V/gm. Fig. 7(b) shows the F-N plot of the I-V curve. The linear slope of F-N plot confirms the current is due to

field emission. The I-V characteristics are found to be consistent from samples prepared separately but under the

same conditions. On increasing the applied field further to 4V/Mm raised the emission current to more than 1 mA, or

an emission current density of 6.25 mA/cm 2, as shown in Fig. 8(a). The high current F-N plot, Fig. 8(b), revealed a

slope change at higher field.

The low threshold for field emission from CNTs has been mainly attributed to the large field enhancement from

the small diameter and large aspect ratio of the tube (ref. 20). From the slope of the F-N plot in Fig. 7(a) and Eq.

(3), we estimate the field enhancement factor of our CNTs sample to be 5500. This value is much larger than the

range of 1000-3000, typical of MCNTs, but within the range of SCNTs that are between 2500-10000 (ref.21). It has

been shown that fine MCNTs similar to our samples but produced by arc deposition, have better emission

6O

5O

_ 40

3o

20

10

0

0 50 1 O0 150 200 250

Voltage (V)

c

-15

-20

-25

-30

0 0.005 0.01

1/V

0.015

(a) (b)

FIG. 7. I-V characteristics of CNTs at low emission current. (a) I-V plot, (b) F-N plot.

NASA/CP--2001-210948 721 August 6, 2001
1200

1000

_& 800

_, 6oo

t_ 400

2OO

-15

-20

-25

0 -30 _

0 100 200 300 400 500 0 0.005 0.01 0.015

Voltag e (V) 1N

(a) (b)

FIG. 8. I-V characteristics of CNTs at high emission current. (a) I-V plot, (b) F-N plot.

characteristics than both catalytic grown open end and defective MCNTs and SCNTs, even though the latter samples

have a much smaller tube diameter and hence larger field enhancement factor (ref. 22). This is explained as due to

the lack of surface states at the tube end and due to the screening effect of SCNT bundling (ref. 22). Comparing the

two types of CNTs in Fig. 3, we conclude that good field emission characteristics are most likely originated from

type B CNTs with fine tube diameters. Efforts are underway to distinguish experimentally the emission

characteristics from the two types of CNTs in our samples.

The long-term emission stability was also measured at both high and low emission currents as shown in Fig.9(a)

and (b), respectively. The emission current starting at 1.07 mA gradually decreases and stabilizes at 0.95mA, giving

a 13% decrease in a 3-houz period of stability test. The decay of the emission current may be due to permanent

damage caused by high current at those emitting sites having large field enhancement factor. The stabilization of the

emission current with the time shows that more emission site may be activated to share the current load and reduce

the damage to the individual site. No decay is observed from Fig. 9(b) for CNT samples emitting around 52 gA

over 15 hours. This fuzther confirms that the emitting sites among the MCNTs are quite dense and stable for field

emission.

1100

1050

1000

O 950

9OO i i i

0 50 100 150

Time (rain.)

200

53.0

52.5

52.0

51.5

51.0

0

i i i i

200 400 600 800

Time (min.)

(a) (b)

1000

FIG. 9. Long-term stability of field emission measured at (a) high current, and (b) low current.

NASA/CP--2001-210948 722 August 6, 2001
CONCLUSIONS

We have grown vertically aligned MCNTs by MPCVD using NiCr as catalyst. This method results in two

distinct types of CNTs. The samples have been characterized by SEM, TEM, Raman spectroscopy, and I-V

measurement. Tuxn-on field of 1V/gm at 2nA and emission current of lmA at 4V/Nn are achieved from samples of

4><4 mm 2 size. Good field emission characteristics obtained from the CNT samples are attributed to the unique

ultra-fine MCNTs with a tube diameter as small as 7nm. The small size of catalyst particle results in small tube

diameter. Stable, high current emission was achieved due to emission sites that are both dense and uniformly

distributed. These results demonstrate the potential of CNTs for various field emission applications such as flat

panel displays and vacuum electronic devices.

REFERENCES

1. S. Iijima, "Helical microtubules ofgraphitic carbon", Nature (London), vol. 354, Nov., 1991, p.56.

2. Q.H. Wang, et al.," A nanotube-based field-emission flat panel display", Appl. Phys. Lett. vol. 72, no. 22, Jun.

1998, p. 2912.

3. C. Liu, et al., "Hydrogen Storage in Single-Walled Carbon Natubes at Room Temperature", Science vol. 286,

Nov., 1999, p. 1127.

4. Y. Nakayama, et al., "Microprocess for fabricating carbon-nanotube probes of a scanning probe microscope",

J. Vac. Sci. Technol. B vol. 18, no. 2, Mar/Apr. 2000, p. 661.

5. Y.Y. Wei and G. Eres, "Directed assembly of carbon nanotuybe electronics circuits", Appl. Phys. Lett. vol.

76, no. 25, Juaa. 2000, p. 3759.

6. A.M. Rao, et al., "In situ-grown carbon nanotube array with excellent field emission characteristics", Appl.

Phys. Lett. vol. 76, no. 25, Juaa. 2000, p. 3813.

7. W.A. de Heer, A. Chatelain, and D. Ugarte, "A carbon Nanotube Field-Emission Electron Source", Science

vol. 270, Nov., 1995, p. 1179.

8. H. Muxakami, et al., "Field emission from well-aligned, patterned, carbon nanotube emitters", Appl. Phys. Lett.

vol. 76, no. 13, Mar., 2000, p. 1776.

9. W.B. Choi, et al., "Fully sealed, high-brightness carbon-nanotube field-emission display", Appl. Phys. Lett.

vol. 75, no. 20, Nov. 1999, p. 3129.

10. P. J. F. Harris, Carbon Nanotubes and Related Structure, Cambridge University Press, Cambridge, 1999, p. 6-

60.

11. Z.F. Ren, et al., " Synthesis of Large Arrays of Well-Aligned Carbon Nanotubes on Glass", Science vol. 282,

Nov. 1998, p. 1105.

12. S. Fan, et al., "Self-Oriented Regular Arrays of Carbon Nanotubes and Their Field Emission Properties",

Science vol. 283, Jan. 1999, p. 512.

13. C. Bower, et al., "Plasma-induced alignment of carbon nanotubes", Appl. Phys. Lett., vol. 77, no. 6, Aug. 2000,

p. 830.

14. Y. C. Choi, et al., "Effect of surface morphology of Ni thin film on the growth of aligned carbon nanotubes by

microwave plasma-enhanced chemical vapor deposition", J. Appl. Phys. Vol. 88, no. 8, Oct. 2000, p. 4898.

15. J.R. Dennison, M. Holtz, and G. Swain, "Raman Spectroscopy of Carbon Materials", Spectroscopy vol. 11,

no. 8, Oct., 1996, p. 38.

16. H. Hiura, et al.," Raman studies of carbon nanotubes", Chem. Phys. Lett. vol. 202, no.6, Feb. 1993, p. 509.

17. S. Albin, et al., "Diamond coated silicon field emitter array", J. Vac. Sci. Technol. A vol. 17, no. 4, July, 1999,

p. 2104.

18. R. H. Fowler and L. W. Nordheim," Electron Emission in intense Electric Fields", Proc. R. Soc. London A.

vol. 119, 1928, p. 173.

19. C.A. Spindt, et al., "Physical properties of thin-film emission cathodes with molybdenum cones", J. Appl.

Phys. vol. 47, no. 12, Dec. 1976, p. 5248.

20. Y. Saito and S. Uemuxa," Field emission from carbon nanotubes and its application to electron sources",

Carbon vol. 38, 2000, p. 169.

21. J.M. Bonard, et al., "Field emission from single-wall carbon nanotube films", Appl. Phys. Lett. vol. 73, no. 7,

Aug., 1998, p. 918.

22. J.M. Bonard, et al., "Field emission from carbon nanotubes: perspectives for applications and clues to the

emission mechanism", Appl. Phys. A vol. 69, July, 1999, p. 245.

NASA/CP--2001-210948 723 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CHARACTERISTICS OF CARBON NANOTUBES SYNTHESIZED BY

LASER ABLATION PROCESS

Hsiu-Fung Cheng "), Teng-Fang Kuo b), Tzai-Hsin Lai d), Chun-Horng Tsai ct and I-Nan Linb);

_/Department of Physics, National TaJwan Normal University, Taipei (117), Taiwan ROC;

b)Materials Science Center, National Tsing-Hua University, o/ Department of Engineering and System

Science, National Tsing Hua University, Hsinchu (300), TaJwan, ROC; d)Department of Physics,

Chung-Yuan Christian University, Chung-Li (320), TaJwan, ROC.

ABSTRACT

Since their first successful synthesis by the arc discharging technique 1, and the subsequent method

with the metal catalyst in an inert gas atmosphere 2, extensive investigations on characteristics of carbon

nanotubes have been pursued due to their unique physical properties 3 and potential technological

applications 4. While the carbon nanotubes were observed to emit electrons efficiently, the mechanism

of electron emission is unclear. How the characteristics of carbon nanotubes influence these properties

is also not well-studied. In this paper, we systematically investigated how the deposition parameters

in this process alter the characteristics of carbon nanotubes. The correlation between the electron

field emission properties of CNTs thus obtained with their characteristics was discussed. In

conventional laser ablation process, the target was placed perpendicular to the incident laser beams, and

was maintained at reaction temperature (1200 °C). The substrates were placed at exhausting end of

quartz tube (water cooled) to collected the carbon nanotubes formed in carrier gas streams. By

contrast, in modified pulse laser deposition process, the target surface was inclined to direct the laser

ejected plume toward the substrates, with both the targets and substrates maintained at deposition

temperature.

Carbon Nanotubes (CNTs) possessing good electron field emission properties can be successfully

synthesized using laser ablation process. CNTs can be formed in gas phase by either using a furnace

(-1050 °C) to trigger the reaction at high temperature or using the laser beams to re-excite the carbon

species at room temperature. The yield is high, -100%, but efficiency in these processes in low, even

when the substrates were biased a high voltage (Fig. 1). Modified laser ablation process, in which the

energetic carbon species were collected by a catalyst coated substrates immediately after they left the

targets, and not only efficiency is markedly increase the collection efficiency, to about lmm/h, but also

significantly lowered the substrate temperature (-750 °C). However, suitable pretreatment on

Ni-cataJyst to form nano-sized catalyst clusters is necessary in order to growth CNTs, and the diameter

of CNTs is closely related to the size of catalyst clusters. Thus obtained CNTs can be turned on at a

field as low as E0= 1.8 V/gin, attaining emission current density as large as Je=180 mA/cm 2 under

4 V/_m applied field (Fig. 2).

1)S. Iijima, Nature 354, 56 (1991); S. Iijima, T. Ichihashi, Nature 363, 603 (1993).

2)M. S. Dresselhaus, G. Dresselhaus, R. SaJto, Carbon 33, 883 (1995).

3)A. M. Rao et al., Science 275, 187 (1997).

4)R. Martel, T. Schmidt, H. R. Shea, T. Hertel, and Ph. Avouris, Appl. Phys. Lett. 73, 2447 (1998).

Keywords: Carbon Nanotubes, Field Emission, Laser Ablation

Corresponding author: I-Nan Lin

101 Kuang-Fu Rd., Sec. 2, Materials Science Center, National Tsing-Hua University, Hsin-Chu, TaJwan

300, ROC

_nlin (_'ra x.i:t hu. edu.t w

Fax: +886-3-5716977

Tel: +886-3-5742574

NASA/CP--2001-210948 724 August 6, 2001
Fig.1 SEM micrographs of carbon nanombes deposited an silicon substrates, which were collected at

exhausting end of the quartz with the application of bias voltage to substrates.

0018

0016

0014

A

012

._OlO

_008

_006

004

0 oo2

0 000

0 oo2

.,o"

I/E(l_m/v )

2

I , I , I , I , I

0 1 2 3 4

Elecric Field(VIII, m)

50torr Eo=1.8

,i,i

5 6

Fig.2 (a) Cross-ectiona3 SEM of CNTs deposited at 750 °C, under 50 torr of Ar:H2(5%) and (b)

electron filed emission properties of thus obtained CNTs.

NASA/CP--2001-210948 725 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FIELD EMISSION FROM CARBON NANOTUBES GROWN BY THERMAL CVD ON

NiCr COATED Si

Bing Xiao, Weihai Fu, and Sacharia Albin

Microelectronics Laboratory

Department of Electrical and Computer Engineering

Old Dominion University, Norfolk, VA 23529, USA

ABSTRACT

Multi-walled carbon nanotubes (CNTs) were grown by thermal CVD in C2H 2 mixed with diluting gas ofAr, H 2

or NH3. Nickel-chromium alloy (Ni:Cr 80%:20%) thin films (-50 nm thick) coated on n-type Si substrates were

used as catalyst. The substrate temperature was maintained at 800°C during CVD process at a pressure of 20 Tom

The CzH 2 flow rate was 30 sccm while the flow rate of the diluting gas was 300 sccm. No substrate pre-treatment

was used for CNT growth. The grown CNTs were multi-walled and were characterized using SEM and TEM. The

tube diameters of CNTs grown in CzH2/Ar and CzHz/H2 were 15-50 nm whereas samples grown in CzH2/NH3 were

much thicker, 25-100 nm in diameter. Field emission characteristics were measured through emission tests using a

diode configuration. The turn-on electric field was 2-3 V/gm for CNTs grown in CzHz/Ar and CzH2/H2 while it

increased to 3-4 V/gm for samples grown in CzHz/NH3. An emission current density up to 10 mA/cm 2 was obtained

from samples grown in CzHz/Ar and CzH2/H2 at a field of 4-5V/gm. CNTs grown in CzHz/NH3 had a higher turn-on

field and lower emission current density due to their larger tube diameters. Long-term emission stability tests showed

no decay during a 12-hour period when the emission current density was at -8.5 mA/cm 2. The results demonstrate

that the carbon nanotubes grown by thermal CVD can be used as a material for field emitters.

Keywords: carbon nanotubes, field emission, chemical vapor deposition.

INTRODUCTION

Carbon nanotubes (CNTs) are being investigated extensively for promising applications as field emitters (ref. 1

to 4) for flat panel displays (ref. 5 to 7) and vacuum microelectronic devices. CNTs have been prepared by various

methods (ref. 2, 4, and 8 to 11). Chemical vapor deposition (CVD) methods, such as microwave plasma CVD (ref. 4,

12, and 13) and thermal CVD (ref. 8 and 14 to 16), are commonly used to grow CNTs on substrates coated with

catalyst of Ni, Co, or Fe thin films. The catalyst film is required for nucleation and growth of CNTs. Compared to

microwave plasma CVD, thermal CVD is simpler and can be used to grow CNTs uniformly on large substrates.

Although microwave plasma CVD usually produces well-aligned CNTs, oriented CNTs can also be achieved by

thermal CVD (ref. 8, 15, and 16). However, some research groups reported CNT growth on insulating substrates (ref.

14 and 15); some methods even require porous substrates (ref. 8, 16) or pre-treatment on catalyst layer (ref. 14).

These methods may not be ideal for device fabrications. For thermal CVD, several gas mixtures have been used to

grow CNTs, such as acetylene-argon (ref. 14), acetylene-ammonia (ref. 16), and acetylene-hydrogen (ref. 15).

In this study, we present results of CNT growth by thermal CVD on NiCr coated Si substrates in acetylene

(C2H2) mixed with Ar, H2 or NH3 and their field emission properties.

EXPERIMENTS

CNTs were grown on n-type (100) Si substrates coated with -50 nm thick NiCr (80%:20%) film deposited by rf

sputtering. No substrate pre-treatments were done before the CNT growth. Thermal CVD growth of CNTs was

carried out in the vacuum chamber of a microwave plasma CVD system. The substrates were placed on an

inductively heated graphite heater with the NiCr film facing the heater. There was a gap of-1 mm between the heater

Author to whom all correspondence should be addressed; email: salbin@odu.edu, fax: (757)683-3220, phone:

(757)683-4967.

NASA/CP--2001-210948 726 August 6, 2001
andthesubstrate, thusgasmixtureshadaccesstothecatalyst film.Thevacuumchamber wasinitiallypumpeddown

toabasepressureblow10.4Torrwithoutgasfeed,thenthesubstrates wereheatedupto800°Cindilutingases,

Ar,H2orNH3.Eachtimeonlyonedilutingaswasused.Finally,thegaspressurewasraisedandmaintained

automatically at20TorrandC2H2wasintroducedtostarttheCNTgrowththatlastedfor10minutes. Thegasflow

ratesofC2H2anddilutingaseswere30sccmand300sccm,respectively. Thegrowthprocesswasstoppedby

shuttingoffC2H2flowandthesampleswerecooledowninArambient.

Properties ofthegrownCNTswereinvestigated byusingSEM,TEMandfieldemissionmeasurements. Aplaneto-planeconfiguration wasappliedtomeasurethedirect-current fieldemissionproperties. Theareaoftheanodewas

2.5mm2,whichwassmallerthantheCNTsampleinordertoavoidtheedgeffect(ref.17).Theseparation between

theanodeandCNTswas100gm.Fieldemissiontestswereperformedinavacuumbetterthan10.7TomA100k_

protectingresistorwasusedinthetestingcircuit.

RESULTSANDDISSCUSSIONS

(c)

Figure 1. SEM and TEM micrographs of CNTs grown on NiCr coated Si by thermal CVD in: (a)

C2H2/Ar, (b) C2H2/H2, and (c) C2H2/NH3.

NASA/CP--2001-210948 727 August 6, 2001
C2H2/NH3, thetubediameters arelargerrangingfrom25to100nm.Thisresultmaysuggest thatNH3canenhance

thethermal decomposition ofC2H2orfavorCNTnucleationlargercatalyst particles. Further, bothinnerandouter

tubediameters areirregularalongthetubeanddefectsarecommonlyseen.

12

10

a

A 8

6

C v

o 4

'_ 2

[JA

o

-18

-20

-22

-24

c

_J

-26

-28

-30

0.001

__ 11__._-*-AI

-,-B I

c I

i i i

3 4 5 6 0.002 0.003 0.004

Field (V/gm) 1N (1Nolt)

(a) (b)

0.005

Figure 2. Field emission characteristics for CNTs grown in C2H2/Ar (curve A),

C2H2/H2 (curve B), and C2H2/NH3 (curve C): (a)I-Vplot, (b) F-N plot.

Figure 2 shows typical I-V characteristics of the field emission from CNTs. From the above discussion, it is

expected that the field emission characteristics will show some independence on diluting gases used in CNT growth.

For samples grown in C2H2/Ar and C2H2/H2, the turn-on field (1 gA/cm 2) was 2-3 V/gm and emission current density

up to 10 mA/cm 2 was observed at a field of 4-5 V/gm. For samples grown in C2H2/NH3, the turn-on field was higher,

around 3-4 V/gm and the emission current density was much lower than the other two kinds of samples. This result

could be explained by considering a lower field enhancement due to larger tube diameters.

The results for long-term stability tests are displayed in Figure 3. Samples grown in C2H2/H2 had similar results

with samples grown in C2H2/Ar which is shown in Figure 3(a). The curves indicate no decay of the emission currents

during the 12-hr test period. The samples grown in C2H2/NH3 have a slightly larger current fluctuation than the other

two kinds of samples. This could be due to the fact that samples grown in C2H2/NH3 have larger tube diameters,

therefore they have fewer emission sites to stabilize the emission current. The short-term current fluctuations were

less than 1 percent.

g

13

A

c v

.£

[JA

9.5 2.00

8.5

u)

g 1.8o

I3

_ 1.6o

C

.£

1.40

[JA

7.5 ' ' 1.20 ' '

0 5 10 15 0 5 10

Time (hr) Time (hr)

(a) (b)

15

Figure 3. Long-term emission stability for CNTs grown in: (a) C2H2/Ar, (b) C2H2/NH3.

NASA/CP--2001-210948 728 August 6, 2001
CONCLUSIONS

By using thermal CVD, multi-walled CNTs have been grown on NiCr coated Si substrates in three different gas

mixtures, C2H2/Ar, C2H2/H2, and C2H2/NH3. No substrate pre-treatment was used for CNT growth. SEM and TEM

images demonstrate that CNT diameters are dependent on diluting gases used in the CVD process. CNTs grown in

NH3 have much larger tube diameters than CNTs grown in Ar or H2, resulting in a smaller field enhancement for

field emission. CNTs grown in C2H2/Ar and C2H2/H2 have better field emission characteristics than those of CNTs

grown in C2H2/NH3.

REFERENCES

1. de Heer, W.A.; Chatelain, A.; and Ugarte, D.: A Carbon Nanotube Field-Emission Electron Source. Science,

vol. 270, Nov 1995, 1179.

2. Kwo, J.L.; Tsou, C.C.; and Yokoyama, M.; et al.: Field Emission Characteristics of Carbon Nanotube

Emitters Synthesized byArc Discharge. J. Vac. Sci. Technol. B, vol. 19, no. 1, Jan/Feb 2001, 23.

3. Saito, Y.; Uemura, S.; and Hamaguchi, K.: Cathode Ray Tube Lighting Elements with Carbon Nanotubes

Field Emitters. Jpn. J. Appl. Phys., vol.37, part 2, no. 3B, 1998, L 346.

4. Murakami, H.; Hirakawa, M.; and Tanaka, C.; et al.: Field Emission from Well-Aligned, Patterned, Carbon

Nanotube Emitters. Appl. Phys. Lett., vol. 76, no. 13, Mar 2000, 1776.

5. Wang, Q.H.; Setlur, A.A.; and Lauerhaas, J.M.; et al.: ANanotube-Based Field-Emission Flat Panel Display.

Appl. Phys. Lett., vol. 72, no. 22, Jun 1998, 2912.

6. Choi, W.B.; Lee, Y.H.; and Lee, N.S.; et al.: Carbon-Nanotubes for Full-Color Field-Emission Displays.

Jpn. J. Appl. Phys., vol. 39, part 1, no. 5A, May 2000, 2560.

7. Wang, Q.H.; Yan, M.; and Chang, R.RH.: Flat Panel Display Prototype Using Gated Carbon Nanotube

Field Emitters. Appl. Phys. Lett., vol. 78, no. 9, Feb 2001, 1294.

8. Fan, S.; Chapline, M.G.; and Franklin, N.R.; et al.: Self-Oriented Regular Arrays of Carbon Nanotubes and

Their Field Emission Properties. Science, vol. 283, Jan 1999, 512.

9. Thess, A.; Lee, R.; and Nikolaev, P.; et al.: Crystalline Ropes of Metallic Carbon Nanotubes. Science, vol.

273, Jul 1996, 483.

10. Li, J.; Papadopoulos, C.; and Xu, J.M.; et al.: Highly-Ordered Carbon Nanotube Arrays for Electronics

Applications. Appl. Phys. Lett., vol. 75, no. 3, Jul 1999, 367.

11. Chen, Y.; Shaw, D.T.; and Guo, L.: Field Emission of Different Oriented Carbon Nanotubes. Appl. Phys.

Lett., vol. 76, no. 17, Apr 2000, 2469.

12. Bower, C.; Zhu, W.; and Jin S.: Plasma-Induced Alignment of Carbon Nanotubes. Appl. Phys. Lett., vol. 77,

no. 6, Aug 2000, 830.

13. Choi, Y.C.; Shin, Y.M.; and Lim, S.C.; et al.: Effect of Surface Morphology of Ni Thin Film on The Growth

of Aligned Carbon Nanotubes by Microwave Plasma-Enhanced Chemical Vapor Deposition. J. Appl. Phys.,

vol. 88, no. 8, Oct 2000, 4898.

14. Lee, C.J.; Kim, D.W.; and Lee, T.J.; et al.: Synthesis of Uniformly Distributed Carbon Nanotubes onA

Large Area of Si Substrates by Thermal Chemical Vapor Deposition. Appl. Phys. Lett., vol. 75, no. 12, Sep

1999, 1721.

15. Yoon, Y.J.; and Baik, H.K.: Synthesis of Carbon Nanotubes by Chemical Vapor Deposition for Field

Emitters. J. Vac. Sci. Technol. B, vol. 19, no. 1, Jan/Feb 2001, 27.

16. Sohn, J.I.; Lee, S.; and Song, Y.H.; et al.: Patterned Selective Growth of Carbon Nanotubes and Large Field

Emission from Vertically Well-Aligned Carbon Nanotube Field Emitter Arrays. Appl. Phys. Lett., vol. 78,

no. 7, Feb 2001, 901.

17. Choi, W.B.; Jin, Y.W.; and Kim, H.Y.; et al.: Electrophoresis Deposition of Carbon Nanotubes for TriodeType Field Emission Display. Appl. Phys. Lett., vol. 78, no. 11, Mar 2001, 1547.

NASA/CP--2001-210948 729 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ALIGNED AND OPENED CARBON NANOTUBE ARRAYS AND THEIR

FIELD EMISSION PROPERTIES

Zheng Wei Pan *, Zhong L. Wang

School of Materials Science and Engineering, Georgia Institute of Technology,

Atlanta, GA 30332-0245, USA

S.T. Lee

Center Of Super-Diamond and Advanced Films (COSDAF), Department of Physics and Materials Science,

City University of Hong Kong, Kowloon, Hong Kong

S.S. Xie

Center for Condensed Matter Physics & Institute of Physics, Chinese Academy of Sciences,

P.O. Box 603-92, Beijing 100080, P.R. China

ABSTRACT

Highly aligned multi-walled carbon nanotube arrays were synthesized in large scale by pyrolysis of

acetylene on film-like iron/silica substrates at 600 °C. Scanning electron microscope images show that the nanotubes

grow outwards separately and perpendicularly from the surface of the substrate to form an aligned nanotube array.

The nanotubes within these arrays are of uniform external diameter (-20-40 nm), with a spacing of about 100 nm

between the tubes. The length of the nanotube arrays increases with growth time, and reaches -2 mm after 48 hours

growth. The growth rate is -30-40 gm/h. The nanotube arrays can be easily stripped off from the substrates, and

transmission electron microscope observations reveal that the bottom ends of the tubes are naturally opened while

the top ends are closed and catalytic nanoparticles are encapsulated at the tips. The tip-guided growth of tubes might

be responsible for the formation of the opened carbon nanotubes.

A very low-field emission was achieved from the aligned and opened carbon nanotube arrays. Field

emission current densities of 10 gA/cm 2 were observed at applied fields of 0.6-1 V/gm, and current densities of 10

mA/cm 2 have been realized at applied fields as low as 2-2.7 V/gm. These results indicate that the aligned and

opened carbon nanotube arrays are superior field emitters.

Key words: aligned carbon nanotubes, opened carbon nanotubes, field emission

1. INTRODUCTION

Carbon nanotubes have been the focus of intense research interest since their discovery a decade ago (ref.

1) because of their unique physical and electronic properties. Carbon nanotubes can now be produced in large

quantities by the arc method (ref. 2), laser ablation (ref. 3) and thermal decomposition of hydrocarbon (ref. 4).

Controlled growth of aligned carbon nanotubes is important for both potential applications and detailed

characterizations of their properties. Up to now, several techniques based on thermal decomposition of hydrocarbon

have been developed to produce aligned multi-walled carbon nanotubes. For example, Li et al. (ref. 5) used

mesoporous silica substrates with aligned pores in which catalyst particles were embedded to produce aligned,

isolated, and dense carbon nanotubes. Terrones et al. (ref. 6) produced aligned tubes using laser-etched cobalt thin

film as a substrate. Pan et al. (ref. 7 and 8) successfully grew very long (up to 2 mm) aligned and opened carbon

nanotubes by pyrolysis of acetylene over film like iron/silica substrates. Ren et al. (ref. 9) and Fan et al. (ref. 10)

reported the aligned nanotubes grown on the glass and silicon substrates, respectively. The aligned carbon nanotubes

obtained by these researchers have a common feature: the nanotubes are approximately perpendicular to the surfaces

of the substrates from which they grew out, which is helpful in, for example, the investigation of some properties

and special applications of tubes. In particular, the successful synthesis of very long nanotubes (ref. 7 and 8) make

us possible to measure the tubes' properties by some conventional methods, and thus they are a suitable candidate

for studying their thermal (ref. 11), mechanical (ref. 12), and field emission properties (ref. 13).

Carbon nanotubes posses various superior properties for using as field emitters, such as sharp tips with a

nanometer-scale radius of curvature (ref. 1), high mechanical stiffness (ref. 14), high chemical stability (ref. 15), and

unique electrical properties (ref. 16). Indeed, recent field emission measurements show that carbon nanotubes

exhibit promise as field emitters in applications such as flat panel displays (ref. 17 to 22). The reported turn-on field

(Eto) and threshold field (Ethr) for electron emission, defined as the macroscopic fields needed to produce a current

*Corresponding author. Tel.: +1-404-385-0326; fax: +1-404-894-9140;

e-mail: zhengwei.pan@ mse.gatech.edu

NASA/CP--2001-210948 730 August 6, 2001
densityof 10gA/cm2and10mA/cm2,respectively, areintherangeof2-5V/gmand4-7V/gm.Thesevalues

shouldranknanotubes amongthebestelectronfieldemittersthatarenowavailable. However, theexperimental

measurements alsoindicatethatmuchlower-fieldemissioncanbereachedif thenanotubes arealigned(ref.18)

and/oropened(ref.21)Thus,bothalignedandopenedcarbonnanotubes areessential forbuildingtheverylow-field

nanotube-based emitters.

In thispaper,wewill describeourresultsin thesynthesisof verylongalignedandopenedcarbon

nanotubes arrays, aswellastheirfieldemissionproperties.

2.EXPERIMENTAL

Thesubstrates weusedtosynthesizeverylongalignedandopenedcarbonnanotubearrayswereprepared

byasol-gel processuingthefollowingtechnique.

Tetraethoxysilane (10ml)wasmixedwith1.5Mironnitrateaqueoussolution(15ml)andethanol (10ml)

bymagneticstirringfor20min.Afewdropsofconcentrated hydrogenfluoride(0.4ml)wasthenadded, andthe

mixturewasstirredforanother20min.Themixturewasthendroppedontoaquartzplatetoformafilmwith

thicknessof30-50gm.Aftergelationofthemixture,thegelwasdriedovernight a80°Ctoremovetheexcess

waterandothersolvent, duringwhichthegelcrackedintosmallpiecesofsubstrates withareaof5-20mm2.

Thesubstrates wereplacedin a quartzboatandwerethen

introducedintothechamberof atubefurnace.Thesubstrates were

calcinedat450°Cfor10hoursundervacuumandthenreducedat500

°Cfor5hoursinaflowof9%hydrogeninnitrogenunder180Torr.At

thisstage, largequantitiesofnanoparticles withsizeof5-50nmformed

evenlyonallsurfacesofthesubstrates(Fig.1).Energy-dispersive xrayspectra(EDS)takenfromtheseparticlesshowedthepresenceof

iron,siliconandoxygen, whichindicatedthathesenanoparticles were

iron/silicananocomposite particles. Theseiron/silicaparticleswill act

ascatalystsfornanotubegrowth.Subsequently, aflowof9%acetylene

in nitrogenwasintroducedintothechamber ata flowrateof 110

cm3/min, andcarbonnanotubeswereformedonthesubstratesby

depositionofcarbonatomsfromthedecomposition ofacetyleneat600

°Cunder180Torr.Thegrowthtimevariedfrom1to48hours.

Theas-grownnanotubes wereexaminedbyscanningelectron

•. ' 200 nm'

Fig. 1 SEM image of the substrate after

reduction, showing iron/silica nanoparticles

with size of 5-50 nm present evenly on the

substrate's surface.

microscope (SEM) and transmission electron microscope (TEM). The field emission measurements were carried out

in a vacuum chamber using a parallel plate configuration as that shown and described in Section 3.3.

3. RESULTS AND DISCUSSION

3.1 SEM and TEM results

We have used the method described above to prepare carbon nanotubes at a very high yield. After growth,

every surface of the substrate is covered with a nanotube array composed of large quantities of highly aligned

carbon nanotubes (Fig. 2a). The length of the nanotube arrays increases with growth time, and reaches up to 2 mm

after 48 hours growth (Fig. 2b), which is an order of magnitude longer (1 mm vs. 100 gm) than that described in

most previous reports. The growth rate is -30-40 gm/h. We believe that the length of the nanotubes could be even

longer if the growth time was further increased. The area of every aligned nanotube array is equal to the relevant

surface area of the substrate from which the nanotubes grow out; the largest area we have obtained is -15 mm 2.

High-magnification SEM image (Fig. 2c) shows that carbon nanotubes grow out separately and

perpendicularly from the substrate to form an array. The nanotubes within the array are of uniform external diameter

(20-40 nm) and spacing (-100 nm) between tubes. Most of the nanotubes in the array are highly aligned, although a

few of them appear to be slightly tangled or curved. We noted that no traces of particles or other graphitic

nanostructures are detected in both the bottom part and central part of the array, which indicate that the nanotubes

prepared in this study have very high purity.

The nanotube array can be easily stripped off from the substrate without destroying the array's integrity,

and the SEM image (Fig. 2d) taken from the bottom end of the array confirms that the nanotubes are highly aligned

and well separated. EDS collected from the bottom end of the array demonstrate the presence of carbon alone,

neither silicon nor iron could be detected.

NASA/CP--2001-210948 731 August 6, 2001
:_i:i' " _:" .:!::"i: " ":'_!:"'_ _.

_.:_::__:N,_._.':..:_::_" !:i:i: ._"i .::: .4" "!:: ..':" .._!

.................... ii

___!!i__'ii" .::i:" .:iiii iii'!:!i_::i:i:!':::.":" !:i:!:"

............................................ i:i'".:'.'i. ....

Fig. 2 SEM image of aligned carbon nanotube arrays. (a) A sample after 4 hrs of growth, showing nanotube arrays growing out

perpendicularly from the relevant substrate surfaces (indicated by arrows). (b) A sample after 48 hrs of growth, showing 2 mm

long nanotube arrays formed on both side of a film-like substrate. (c) Higher magnification SEM image of an aligned nanotube

array, showing nanotubes growing out separately and perpendicularly from the substrate surface. (d) Bottom structure of the

nanotube array showing well-separated, aligned nanotubes.

During SEM observations, we noted that the bottom ends .................._..........._.........._....._............._.........._...._'....___... iiiiiiii!iiiiii

of the nanotubes seem to be opened (see Fig. 2d). In order to verify : : : i i i

this resolution point, weTEMobserved observations the bottomof theends bottom of nanotubes end of intheTEM. nanotube Low-i[iiiiii;i;iiiiiii;i!!ii!iiiiiii_i!;i ,<i::iiiiii',ii!iii,,iiiiii

i:;[i!_i i

:i:i:i:i:i:i:i:!:i:i:i:i:i:i:!:i:i:!:::!:i::4._.:!:i:i:i:i:i:._i:_:::i:! !.i._i:i:i:!:

array reveal that no seamless caps exist at the bottom ends of the ::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

nanotubes and the tubes are indeed opened (Fig. 3a). That is to say, ::::::::::::::::::::::::::::_::::::::::::::::::::::::::_:::::::::::::::::::_.:::::

,...*,.o_._,,.°.*,..*.,o-.; .,.,.°.°°°°_.o..°,i..°o*,-_°,. _-_-,°°,_,o°°°°°...,,,&.,-_-

we have directly prepared open-ended aligned carbon nanotubes by ::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

,,°_,o%O°*°*o°°*_ °o °%%_,°%°°°°°°°°*,%,,*°°, _ _*°oo°_°,;_Oo°°_._',%%o° chemical vapor deposition without any treatment such as oxidation. ,.°.-...........-....::_.':°....e,....v............_°aI......,..........¢.......-.N

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

High resolution TEM observation (Fig. 3b) shows that the opened :.:.:._.4::':.:.:.:.:.:.:.:.:.:::°:,..:.:.:,:<4;>a.:.:;X4-:_._-:.;]_,(4_;.;,:

and aligned nanotubes are well graphitized and typically consist of ';':<°.':':':':'a':':':':':':'.':'>:':gg°:':':':':'e:':-_4+:':_;'_g"

10-30 concentric graphite layers. Contrast to the opened nanotubes ::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

obtained by oxidation methods (ref. 23 to 26), our opened and :._:_.:_:_:.:_:_:.:._.:.:.:._:._.:.:._.:.:_.:._';_;_;_:'_:_:':':_;_;':__.:_ _ ;:

aligned carbon nanotubes have a fine out-layer without any ::::::::::::::::::::::::::::::::::::::_:::::::::::::_:_:::::::::_:::_:_:::_:::_::::::::::

damaging, iiiiiii_iiiii:i:!:!:!:i:i:i:i:!:_:i:i:!:!:_:i:i:i:i:i:i:i:i:i:!:!:i:i:i:i:i:ii

We have also examined the top ends of the nanotube

arrays by both SEM and TEM. The SEM observations reveal that

large quantities of nanoparticles with size of 5-50 nm exist on the iiiiiiii!ii!!iiiii_!!_iii!i!!iiiiiii_!i!!!_i_i_iii!i_!!i!iiiii!!i!!i __a_"_

top ends of the nanotube arrays (Fig. 4a). The size and distribution i_iii__iiiii_iiiii_i_i!_iiiiii!iii_iiiii!i!i_iiii__ii_ii_iiii_i_iii_i!iiiii!iiii[iiii!iii !

,,o,,,,o,,ooo°o,o°o°o,o°,°oo°,°,°°°o°°°,° ....

of these particles are similar to the iron/silica nanocomposite Fig. 3 (a)Low-resolutionTEMimageofthe

particles formed on the surfaces of the substrates after reduction bottom end of a nanotube array showing aligned

(see Fig. 1). EDS collected from these particles show the presence and opened tubes. (b) High-resolution TEM

of carbon, iron, silicon and small amount of oxygen, which image ofthebottom end of an opened nanotube

indicates that these particles are iron/silica nanocomposite particles showing bowl-shape structure (indicated by an

covered with a thin carbon film due to carbon deposition, arrow) formed by the graphite sheets.

NASA/CP--2001-210948 732 August 6, 2001
Therefore, webelievethattheseiron/silicaparticlesarethose

present onthesurfacesofthesubstrates beforenanotubegrowth.

In orderto verifythisidea,weobservedthesubstratesfrom

whichthenanotubearrayshavebeenstrippedoft,noparticlewas

observedonthesurfacesof theresultingsubstrates. Thatis to

say,theiron/silicaparticlesthatpresentonthesurfacesofthe

substratesbeforenanotubegrowthwere"lift up" fromthe

substrates andtransferred tothetopendsofthenanotubearrays

bythegrowingcarbonnanotubes duringthegrowthprocess. It is

interestingto notethatsubsequent reductionof thesubstrates

afternanotubegrowthresultsin thetransportation of more

catalystparticlestothesubstratesurfacesfromwithinthebulk

substrate. Therefore, thesubstrates canbere-usedafterreduction

togrownewnanotubearrayswithcharacteristics similartothe

originalones.

TEMstudiesshowthatthetopendsofthenanotubes

areclosed,andencapsulated iron/silicananoparticlesare

consistently observedatthetopendsofthetubes(Fig.4b).

t:i:i:i:i.':"

.'_'_m...°..,,i.:,_,l...,...4,,,....

....... -',I,."..... _.., /_t_°'._;_,_:_ .....................................

Fig. 4 (a) SEM image of the tip structure of the

nanotube array showing iron/silica catalyst particles

with size of 5-50 nm. (b) TEM image showing the

encapsulated particle at the top end of the nanotube.

3.2 Growth mechanism

Several models have been proposed to explain the formation mechanism of carbon nanotubes prepared by

catalytic decomposition of hydrocarbon (ref. 27). It is generally accepted that the tubes grow by the extrusion of

carbon, dissolved in a metallic catalyst particle that is oversaturated in carbon at one part of the surface, and that the

catalyst particles promote "tip growth" or "base growth" depending on the contact force between the particles and

the substrate. For our case, the SEM and TEM results shown in Section 3.1 clearly show that the iron/silica catalyst

particles promote "tip growth" growth process and play an essential role in the formation of the opened nanotubes.

Fig. 5 schematically shows the tip growth model of our very long aligned and opened carbon nanotubes.

C b C C

Cata!ysl l_t.icle \\

(a) (9 (_) (¢

Fig. 5 Tip growth model of very long aligned and opened carbon nanotubes

After reduction, the iron/silica catalyst particles were formed on the surface of the substrates (Fig. 5a).

These particles are tangent with the substrate surface, so the combination force between the particles and the

substrate is very weak. During carbon deposition, carbon atoms resulted from the decomposition of acetylene

deposit on the surface of catalyst particles and then diffuse to the annular area between the particles and substrate

through surface diffusion and/or bulk diffusion (Fig. 5b), and form tube-like structure (Fig. 5c), i.e., carbon

nanotubes. Since the connection between the particles and substrate is very weak, the growing carbon nanotubes will

"lift up" the catalyst particles away from the substrate and initial the tip growth. The annular area between the

particles and substrate determine the inner diameter of the tubes. During carbon deposition, the catalyst particles

keep active and continuously absorb the carbon atoms from the vapor gas, so the carbon nanotubes can grow

continuously and thus their length increase with the growth time (Fig. 5d). We note from Fig. 5d that the hollow

cavity of the nanotubes is directly connected with the surface of the substrate, i.e., no seamless graphite cap exist at

the bottom end of the tubes, so the bottom end of the nanotube must be opened and should form a bowl-shape

structure if the nanotubes are removed from the substrate. The TEM image shown in Fig. 3 confirms this point.

It is not clear at this stage how such regularly aligned nanotube arrays form. In the case of the bulk

mesoporous substrate containing iron nanoparticles embedded in the pores, the template effect of the pores

constrains the carbon nanotubes to grow along the axis of the pores to form an aligned nanotube array (ref. 5). But in

the present case, no pore exists in the thin film substrate. Overcrowding growth (ref. 6) or competition growth

mechanism (ref. 28) may be responsible for the unique nanotube growth behavior reported here.

NASA/CP--2001-210948 733 August 6, 2001
3.3Fieldemissionproperties

Thefieldemissionmeasurements werecarriedoutinavacuumchamber atapressureof -5×10 .7 Torr at

room temperature. Fig. 6 shows the schematic diagram of the experimental setup. An aligned carbon nanotube array,

which was used as the cathode, was attached to a stainless steel substrate by silver paste, with the bottom ends of the

nanotubes facing upward. In this configuration, the nanotubes

were placed with their long axes perpendicular to the substrates,

and the bottom end of the nanotube array, which was composed of

high density, well separated, highly aligned and opened nanotube

tips, acted as the emitting surface. A copper plate with a diameter

of 1 cm, mounted on a precision linear feedthrough, was used as

the anode. A variable positive voltage up to 5 kV was applied to

the anode and the emission current (/) was indirectly determined

by measuring the voltage across a 500 k_ resistor. The distance

(d) between the emitting surface and the plate was determined by

first lowers the plate to the sample until electrical contact was

observed, then lifting the plate to a certain value. The

macroscopic electric field (E) was estimated by dividing the

applied voltage by the sample-anode separation (V/d). The

emission current density (J) was calculated from the obtained

emission current and the cathode surface area measured by SEM.

Four plots of emission current density versus

electric field (J-E plot) for a typical carbon nanotube array

with an emitting surface area of -4.62 mm 2 are displayed in

Fig. 7. These data were collected over four anode-sample

distances, 600, 550, 500, and 450 btm. The voltage was

raised from zero to 1500 V, then decreased to zero. After

the anode was moved 50 gm closer to the emitting surface,

the voltage cycle was repeated. It is clear from Fig. 7 that

the four cycles have almost the same J-E characteristics,

indicating a linear relation between the anode voltage and

distance. The uniform J-E characteristics also justify the

simple formula (E = V/d) used to estimate the macroscopic

field strength between the anode and cathode. Electron

emission is observed at an electric field as low as 0.5

V/gm. The electron emission turn-on field Eto and

threshold field Ethr are in the range of 0.75-0.89 V/btm and

2.14-2.32 V/gm, respectively, for the four J-E plots shown

in Fig. 7. Reproducibility tests from eight different

nanotube arrays consistently yield Eto and E,r in the range

of 0.6-1 V/btm and 2-2.7 V/btm, respectively, with average

values of 0.8 V/gm and 2.3 V/gm. Our values for both Eto

and E,r are more than two times lower than those obtained

for carbon nanotubes by other groups (ref. 18 to 20).

3o

.,25

v

"_15

to

Fig. 6 Schematic diagram of the experimental setup

for field emission measurement.

N

lIE [tm_V 4) _/

...... r r r .

0 I 2 3

Electri_ Field (V/txm)

Fig. 7 Emission J-E curves from an aligned and opened

carbon nanotube array with emitting surface area of 4.62

mm 2. (inset) The corresp-onding Folwer-Nordheim plots.

The J-E data shown in Fig. 7 were also analyzed by the Fowler-Nordheim theory (ref. 29) (see Fig. 7

inset). A linear relation was not observed in the ln(J/E 2) versus 1/E characteristics, suggesting that the current limit

is not dominated by the standard tunneling theory proposed by Fowler and Nordheim. Deviations from the FowlerNordheim behavior for metal tips are usually attributed to the space-charge-limited current (SCLC) mechanism (ref.

30). For our carbon nanotubes, although the total current density integrated over the entire anode area seems to be

too low for the SCLC mechanism, the current could, nevertheless, be limited by the SCLC mechanism if a much

higher local current density occurred at the cathode.

The stability of the aligned and opened nanotube samples is remarkable. During 24 hours of continuous

operation at 5 mA/cm 2, the current fluctuation was as low as +3% and the average current did not decrease over this

period. No changes in the morphology after 24 hours of continuous emission were observed by SEM.

NASA/CP--2001-210948 734 August 6, 2001
4. CONCLUSIONS

Very long aligned and opened carbon nanotubes were prepared at a very high yield by pyrolysis of

acetylene over film like iron/silica substrates. The nanotubes grew outwards separately and perpendicularly from the

surfaces of the substrates to form aligned nanotube arrays. The nanotubes within the arrays are of uniform external

diameter (20-40 nm), with a spacing of -100 nm between tubes. When the nanotubes were removed from the

substrates, opened nanotube arrays formed. These very long aligned and opened nanotube arrays show excellent

field emission properties with Eto and E,r in the range of 0.6-1 V/btm and 2-2.7 V/btm, respectively. Such extremely

low turn-on and threshold fields for a remarkably stable electron emission in the aligned and opened carbon

nanotube arrays offer excellent prospects for a low-field nanotube-based flat panel display.

REFERENCES

1. Iijima, S., Helical microtubules of graphitic carbon, Nature 354, 56-58 (1991).

2. Ebbesen, T.W., and Ajaya, P.M., Large-scale synthesis of carbon nanotubes, Nature 358, 220-222 (1992).

3. Thess, A., et al., Crystalline ropes of metallic carbon nanotubes, Science 273,483-487 (1996).

4. Endo, M., et al., The production and structure of pyrolytic carbon nanotubes (PCNTS), J. Phys. Chem. Solids

54, 1841-1848 (1993).

5. Li, W.Z., et al., Large-scale synthesis of aligned carbon nanotubes, Science 274, 1701-1703 (1996).

6. Terrones, M., et al., Controlled production of aligned-nanotube bundles, Nature 388, 52-55 (1997).

7. Pan, Z.W., et al., Very long carbon nanotubes, Nature 394, 631-632 (1998).

8. Pan, Z.W., et al., Direct growth of aligned open carbon nanotubes by chemical vapor deposition, Chem. Phys.

Lett. 299, 97-102 (1999).

9. Ren, Z.F., et al., Synthesis of large arrays of well-aligned carbon nanotubes on glass, Science 282, 1105-1107

(1998).

10. Fan, S.S., et al., Self-oriented regular arrays of carbon nanotubes and their field emission properties, Science

283, 512-514 (1999).

11. Yi, W., et al., Linear specific heat of carbon nanotubes, Phys. Rev. B 59, R9015-R9018 (1999).

12. Pan, Z.W., et al., Tensile test of ropes of very long aligned multiwall carbon nanotubes, Appl. Phys. Lett. 74,

3152-3154 (1999).

13. Pan, Z.W., et al., Very low-field emission from aligned and opened carbon nanotube arrays, J. Phys. Chem. B

105, 1519-1522 (2001).

14. Wong, E.W., Sheehan, P.E., and Lieber, C.M., Nanobeam mechanics: elasticity, strength, and toughness of

nanorods and nanotubes, Science 277, 1971-1975 (1997).

15. Ebbesen, T.W., et al., Electrical conductivity of individual carbon nanotubes, Nature 382, 54-56 (1996).

16. Frank, S., et al., Carbon nanotube quantum resistors, Science 280, 1744-1746 (1998).

17. Rinzler, A.G., et al., Unraveling nanotube: field emission from an atomic wire, Science 269, 1550-1553 (1995).

18. de Heer, W.A., et al., A carbon nanotube field-emission electron source, Science 270, 1179-1180 (1995).

19. Bonard, J.M., et al., Field emission from carbon nanotubes: perspective for applications and clues to the

emission mechanism, Appl. Phys. A-Mater. 69, 245-254 (1999).

20. Zhu, W., et al., Large current density from carbon field emitters, Appl. Phys. Lett. 75, 873-875 (1999).

21. Saito, Y., et al., Conical beams from open nanotubes, Nature 389, 554-555 (1997).

22. Wang, Q.H., et al., Field emission from nanotube bundle emitters at low fields, Appl. Phys. Lett. 70, 3308-3310

(1997).

23. Ajayan, P.M., and Iijima, S., Capillarity-induced filling of carbon nanotubes, Nature 361,333-334 (1993).

24. Tsang, S.C., Harris, P.J.F., and Green, M.L,H., Thinning and opening of carbon nanotubes by oxidation using

carbon dioxide, Nature 362, 520-522 (1993).

25. Ebbesen, T.W., et al., Purification of nanotubes, Nature 367, 519 (1994).

26. Tsang, S.C., et al., A simple chemical method of opening and filling carbon nanotubes, Nature 372, 159-162

(1994).

27. Amelinckx, S., et al., A formation mechanism for catalytically grown helix-shaped graphite nanotubes, Science

265, 635-639 (1994).

28. Li, D.C., et al., Structure and growth of aligned carbon nanotube films by pyrolysis, Chem. Phys. Lett. 316,

349-355 (2000).

29. Fowler, R.H., Norfheim, L.W., Electron emission in intense field, Proc. R. Soc. London A 119, 173-181 (1928).

30. Barbour, J.P., et al., Space-charge effects in field emission, Phys. Rev. 92, 45-51 (1953).

NASA/CP--2001-210948 735 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

MICROWAVE PLASMA ASSISTED BRAZING OF CARBON NANOTUBES AND

DEPOSITION OF CARBON FILMS ON IRON ELECTRODES FOR APPLICATIONS

AS ELECTRON FIELD EMITTERS

Yonhua Tzeng, Chao Liu, and Calvin Cutshaw

Alabama Microelectronics Science and Technology Center

Department of Electrical and Computer Engineering

Auburn University, Auburn, Alabama 36849 USA

ASTRACT

Deposition of carbon films including nanotubes as well as brazing of high-quality single wall carbon nanotubes,

that were pre-synthesized by other means, on iron electrodes as electron field emitters are reported. Microwave

plasmas in vapor mixtures of methanol-based liquid solutions were applied to fresh iron electrodes for carbon and/or

carbide deposition. In another application, microwave plasmas were used to heat iron electrodes to its iron-carbon

eutectic point around 1150°C so that single wall carbon nanotubes that were placed on the iron surface reacted with

iron to form low resistance nanotube-electrode contacts. Electron field emission characteristics of iron electrodes

treated by plasma enhanced chemical vapor deposition as well as plasma brazed carbon nanotubes on iron electrodes

will be presented.

Keywords: methanol, microwave plasma, electron emission, carbon nanotubes, brazing

INTRODUCTION

Cold and/or low-temperature low-electric-field electron emitters have been studied by a large number of

scientists and technologists because of their economic advantages and very attractive performance for many

commercial as well as special applications. Devices ranging from high-power microwave vacuum tubes to flat panel

displays for computer monitors all require electron emitters that are fast, operating at low temperatures, with high

and stable emission current densities at low electric fields.

A number of materials and structures have been studied and reported in an attempt to achieve high-performance

cold cathodes. The most effective means of achieving this goal are the fabrication of micro- or nano-structures that

exhibit a very high aspect ratio and the design and synthesis of coatings with low positive or even negative electron

affinities. The high aspect ratio, for example, a micrometer long carbon nanotube with a nanometer radius of

curvature at its tip, allows electric field enhancement to occur at the tip of the electron emitter so that low voltage

will be adequate to cause field emission of electrons from the tip of the emitter. Negative electron affinity allows

electrons to escape from a solid surface to the vacuum without needing additional forces or energy. For example,

diamond with proper surface terminations has been explored for cold cathode applications based on its negative

electron affinity. There are a good collection of references in the Proceedings of ADC/FCT '99 (ref. 1) about

various efforts, materials, and structures aiming at high-performance electron field emitters.

Several projects have been explored in the past by our research group in collaboration with scientists in the

related field. CVD diamond and carbon nanotubes were synthesized and characterized for their electron field

emission behaviors (refs 2-10). In an effort to achieve carbon nanotube electron field emitters for high current

density and stable electron emission applications, high quality carbon nanotubes may be synthesized separately and

then attached to an electrode forming low contact resistance and strong mechanical strength or be directly deposited

onto a selected electrode. In this paper, iron electrodes were used because iron serves as an effective catalyst for

carbon nanotube growth as well as forms good contacts with carbon materials at its iron-carbon eutectic temperature

around 1150°C, which is well below the melting point of iron, so that the electrode remains as a solid at the

fabrication temperature (ref. 11).

NASA/CP--2001-210948 736 August 6, 2001
Diamonddepositioninvapormixturesofmethanol-based liquidsolutionswasreportedinADC/FCT'99(refs.

12,13).Inthispaper, microwaveplasmasinvapormixturesofmethanol andethanol wereappliedtodeposit nanostructuredcarboncoatingsonironelectrodes usingironasacatalyst.High-power densitymicrowaveplasmasina

vaporofmethanol andethanol werealsousedtoheatandbrazesinglewallcarbonnanotubes ontoironelectrodes.

EXPERIMENTAL

Theexperimental apparatus ishowniFigure1.Amixtureofmethanol andethanol wasfedintoavacuum

chamber thatwasevacuated byamechanical pump.Athrottlevalveandamanometer pressuregaugecontrolledthe

chamber gaspressure. Arodantennawasusedtocouplemicrowavepowerfomarectangular waveguideintoa

cylindrical metalcavitythroughaquartzwindowthatseparatedheatmosphere fromthevacuumchamber.

metahnol-based

solution

coItainer

iron

microwave

__ tube

/ quartz

flow rate _ U // ___ window

controller

S .

f4\ o ptical

F

inlet

substrate

vacuum

chamber

Y

L

\,

pyrometer

J plasma

outlet to pump

, substrate

holder

Figure 1. Methanol-based microwave plasma apparatus.

The microwave power formed a plasma ball on the top of the substrate. Electrons in the plasma have very high

temperatures exceeding 10,000°C. The plasma heated the substrate to a preset temperature as well as heated the gas

mixtures for proper dissociation and reaction in the gas phase leading to carbon deposition on the substrate surface.

Iron from the substrate itself was used as the catalyst for the carbon coatings.

In the case of brazing of carbon nanotubes onto iron electrodes, single wall carbon nanotubes suspended in a

liquid solution was applied to the surface of an iron electrode. The iron electrode with carbon nanotubes on it was

used as the substrate to be exposed to the microwave plasma in a mixture of methanol and ethanol. The substrate

temperature was monitored by a dual-color optical pyrometer.

The carbon nanotube coated substrates were then loaded into a high vacuum chamber. Thin quartz plates were

used as spacers between an anode and the iron electrode, which serves as the cathode. A desktop computer

controlled the output of a high-voltage power supply for applying a voltage between the anode and the cathode. The

electric field is calculated by dividing the applied voltage by the gap spacing between the anode and the cathode.

The electron emission current was measured by a digital ammeter and recorded by the computer for further plotting

and calculation.

NASA/CP--2001-210948 737 August 6, 2001
RESULTS AND DISCUSSION

Microwave plasmas in methanol-based vapor mixtures with a low percentage of ethanol additives are very

oxidizing and can etch carbon nanotubes rapidly at elevated temperatures. Carbon nanotubes placed on

molybdenum electrodes were etched away after exposing to microwave plasma in a vapor from a solution of 98%

methanol and 2% ethanol by weight at 1250°C for a few minutes. In the case of iron electrodes, carbon nanotubes

dissolved into the iron electrode at a temperature above its iron-carbon eutectic temperature in addition to being

etched away by the plasma.

Shown in Figure 2 is the electron field emission current-electric field (I-E) characteristics for an iron electrode

with carbon nanotube on it before being exposed to the aforementioned microwave plasma for five minutes. The

field emission turn-on electric field shows a value above one volt per micrometer, indicating that the electron field

emission is not caused by carbon nanotube originally placed on the iron electrode. Indeed, what is shown in Figure 3

is a similar I-E curve for an iron electrode without carbon nanotubes being put on the top before being exposed to the

same plasma for the same period of time. SEM images of these two specimens show some sub-micrometer

structures on the iron surface that are assumed to be responsible for the measured electron field emission

characteristics.

1 0-4

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10-12

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Figure 2. I-E curve for an iron electrode after being exposed to a microwave plasma at 1250C for five

minutes. The iron electrode had carbon nanotube placed on it before being exposed to the plasma.

NASA/CP--2001-210948 738 August 6, 2001
Usingaliquidsolutionwith25%byvolumeofethanol in75%byvolumeofmethanol, amicrowaveplasmawas

appliedtoanironelectrodewithoutpre-synthesized carbonnanotubes onit at600°Cfortwohoursatapressureof

10 -4

10-6

_10 -8

0

" 10-1°

10 -12

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i i

0

l|ioeOeellii

0

itlile iili°i

i i i i i i i

0 0.5 1 1.5 2 2.5 3 3.5 4

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Figure 3. I-E curve for an iron electrode after being exposed to a microwave plasma

at 1250C for five minutes.

10 -4

10-6

0

i o..,14

illeieelilil li-4,

i_e e

|oOOOOiOO iiioil

-12 10 .....

0 0.4 0.8 1.2 1.6 2

E (Vlnm)

Figure 4. I-E curve for an iron electrode after being exposed to a microwave plasma in 25% ethanol and

75% methanol (by volume) at 600C for 2 hours.

NASA/CP--2001-210948 739 August 6, 2001
10Torr.Theturn-onelectricfieldshowninFigure4fortheplasmatreatedironelectrodewasmeasured tobe

slightlylessthanonevoltpermicrometer. With25%byvolumeofethanol intheliquidsolution,thecarbonto

oxygenratiowassufficientlyhightocausethecoatingsonthefreshironelectrode, whichledtothelowelectricfield

electronfieldemission.

Singlewallcarbonnanotubes havexcellent electronfieldemissionproperties. Althoughit hasalsobeen

reportedthatsinglewallcarbonnanotubes canbedepositedonsubstrates directly, itisdesirabletobeableto

optimizetheproductionprocessforsinglewallcarbonnanotubes independently whilethecoatingprocessfor

placingsinglewallcarbonnanotubes onelectrodes iseparately optimizedaswell.Brazingof highqualitysingle

wallcarbonnanotubes toelectedelectrodes toformapre-set patternorstructuremaybethebeststrategy.

Usingaliquidsolutionwith25%(byvolume) ofethanol in methanol, anironelectrodethathadsinglewall

carbonnanotubes placedonitssurfacewasexposedtoamicrowaveplasmaat60Torr.Thegreenishplasmaheated

theironelectroderapidlytoreachaboveitseutectictemperature withinalittlemorethanoneminute.A dual-color

opticalpyrometer wasappliedtomonitortheirontemperature. Theplasmawasturnedoffwhentheiron

temperature reached1160°C, whichisslightlyabovetheiron-carboneutectictemperature. ShowniFigure5isthe

I-Echaracteristics forthecarbonnanotubes brazedontotheironelectrode. Averylowelectronfieldemissionwas

measuredatmacroscopic electricfieldoflessthan0.5voltpermicrometer.

10-2 ....

10 -4

10-B

£ 10.8

• • ° o_ 4'Q6e'__

Q

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I i I

0.8 1.2 1.6

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Figure 5. I-E curve for an iron electrode with single wall carbon nanotubes eutectically brazed on it by

exposing to a microwave plasma in 25% ethanol and 75% methanol (by volume) until the electrode

temperature reached 1160C.

CONCLUSIONS

Methanol-based microwave plasmas were applied to deposit carbon coatings on iron electrodes as well as

brazing pre-synthesized single wall carbon nanotubes onto iron electrodes at temperatures near the iron-carbon

eutectic temperature. By means of iron-carbon eutectic brazing, electron field emitters are suitable for high current

density electron emission applications in which carbon nanotubes may be heated to high temperatures by the

emission current. The brazing process also allows the single wall carbon nanotubes to be produced by processes that

NASA/CP--2001-210948 740 August 6, 2001
arenotrestrictedbythecoatingprocess, makingitpossibleforachievingthebestperformance oflectronemitters

basedoncarbonnanotubes.

REFERENCES

1.

2.

3.

4.

5.

6.

7.

8.

9.

10.

11.

12.

13.

Yoshhikawa, Y. Koga, Y. Tzeng, C.-R Klages, K. Miyoshi, editors, Proceedings of ADC/FCT '99, pp.335-380 and

pp. 687-712, August 31-September 3, 1999, AIST-Tsukuba Research Center, Tsukuba, Japan.

N. S. Xu, R. V. Latham, and Y. Tzeng, "Similarities in the "Cold' - Electron Emission Characteristics of Diamond

Coated Mo Electrodes and Polished Bulk Graphite Surfaces," J. Phys. D: Applied Physics, 26, pp. 1776-1780,

1993.

N. S. Xu, R. V. Latham and Y. Tzeng, "Field-dependence of the Area-density of "Cold' Electron Emission Sites on

Broad-Area CVD Diamond Films," Electronics Letters Vol. 29, No. 18, 2 September 1993.

N. S. Xu, Y. Tzeng, and R. V. Latham, "A Diagnostic Study of the Field Emission Characteristics of Individual

Micro-Emitters in CVD Diamond Films," J. Phys. D: Appl. Phys. (Rapid Communication) 27, pp. 1988-1991,

1994.

N. S. Xu, R. V. Latham and Y. Tzeng. "Field-Induced Electron Emission from CVD Diamond Films on Planar Mo

Substrates," Diamond Films and Technology, Vol. 4, No. 4, pp. 249-258, 1994.

Yonhua Tzeng, Chao Liu and Zheng Chen, "Hot-Filament Assisted Fabrication of Carbon Nanotube Electron

Emitters,", Mat. Res. Soc. Syrup. Proc. Vol. 621. R7.5.1-R7.5.7, April 24-28, 2000, San Francisco, CA.

Zheng Chen, Yonhua Tzeng, Chao Liu, and Calvin Cutshaw, "Fabrication and Applications of Microstructured NiC Electrodes with High-Density Carbon- Nanotube Coatings," Mat. Res. Soc. Symp. Proc. Vol. 621. R7.4.1-

R7.4.4, April 24-28, 2000, San Francisco, CA.

Yonhua Tzeng, Chao Liu, and Zheng Chen, "Electron Emission from Microwave Plasma CVD Carbon

Nanotubes," Mat. Res. Soc. Syrup. Proc. Vol. 621. R7.3.1-R7.3.6, April 24-28, 2000, San Francisco, CA.

Y. Tzeng, C. Liu, and Z. Chen, "Low Temperature CVD Carbon Films on Glass Plates for Flat Panel display

applications," Mat. Res. Soc. Symp. Proc. Vol. 621. Q2.2.1-Q2.2.5, April 24-28, 2000, San Francisco, CA.

Tzeng, C. Liu, C. Cutshaw, and Z. Chen, "Carbon Nanotube Coatings for Electron Emitter Applications,"

Presented in the Seventh International Conference on New Diamond Science and Technology, July 24-28, 2000,

Hong Kong, China.

Y. Tzeng, "Rapid Smoothing of CVD White Diamond by Liquid Metal Films," Proceedings of the 3r_ International

Conference on the Applications of Diamond and Related Materials, NIST, Washington, DC, 1995.

Y. Tzeng, "Microwave Plasma Enhanced Chemical Vapor Deposition of Diamond in the Vapor of

Methanol-Based Liquid Solutions," Proceedings of ADC/FCT 1999, pp.20-24, August 31 -September 4, 1999,

Tsukuba, Japan.

Y. Tzeng, "Hot-Filament Assisted Deposition of Diamond in the Vapor of Methanol-B ased Liquid

Solutions," Proceedings of ADC/FCT 1999, pp.420-424, August 31 - September 4, 1999, Tsukuba, Japan.

NASA/CP--2001-210948 741 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

HYSTERESIS OF ELECTRON FIELD EMISSION FROM SINGLE-WALLED CARBON

NANOTUBES BRAZED ON IRON SUBSTRATES

Chao Liu, Calvin Cutshaw and Yonhua Tzeng

Electrical and Computer Engineering Department, Auburn University

200 Broun Hall, Auburn, Alabama 36849 USA

ABSTRACT

Hysteresis of emission current-electric field (I-E) characteristics of single walled carbon nanotubes brazed on

iron substrates by TiCuAg alloy is studied. Different behaviors of the I-E hysteresis with the substrate being held at

room temperature and that with the substrate being heated to an elevated temperature were measured. At room

temperature, the descending I-E curve was below the ascending I-E curve indicating that tunneling electron emission

enhanced by chemisorbed molecules on the surface of the cool carbon nanotube when the emission current first

started during the ascending cycle sifts the I-E curve to the upper left. The burn-out of defective or taller carbon

nanotubes may also cause the descending I-E curve to shift to the lower right. Both mechanisms may contribute to

the measured I-E hysteresis. On the other hand, the I-E hysteresis measured at elevated temperatures showed the

descending I-E curve being above the ascending I-E curve. The self-heating of carbon nanotubes by emission

current coupled by the less effective cooling with the substrate being at elevated temperatures during the descending

cycle, the contribution of thermionic electron emission from nanotubes, and the relatively easier bending and

straightening of carbon nanotubes at elevated temperature in alignment with the applied electric field may contribute

to the measured I-E hysteresis at elevated temperatures being different from that at room temperature.

Keywords: carbon nanotubes, electron emission, hysteresis, tunneling

INTRODUCTION

Carbon nanotubes are promising for cold cathode applications due to their excellent electrical and mechanical

properties, such as high aspect ratio, small tip and curvature, high mechanical strength, and high resistance to

chemical and physical attacks (refs. 1 to 4). Various experiments have been carried out and have demonstrated

attractively low threshold and turn-on electric fields for field emission of electrons, as well as high emission current

densities. Mechanisms for electron emission behaviors of carbon nanotubes are not adequately clear. In this paper,

the electron emission characteristics of carbon nanotubes are examined with emphasis on the hysteresis of the

emission current-electric field characteristics. The hysteresis of the I-E curves may be related to the stability of the

electron field emission in a long-term field emission operation as well as field emission operations under changing

environments. The carbon nanotube electron emitter exhibiting the least I-E hysteresis may also be the electron field

emitter that will provide the best long-term field emission stability.

EXPERIMENTAL PROCEDURE

The specimen used in this study was prepared by first putting a piece of brazing alloy (Wesgo Ticusil, containing

4.5% Ti, 26.7% Cu, and 68.8% Ag) sheet, which has liquidus phase above 850°C and solidus phase below 830°C,

onto a piece of iron with a flat surface. A drop of carbon nanotube suspended in a liquid was then applied onto the

brazing alloy sheet. After the liquid vaporized, the specimen was loaded into a vacuum furnace, which was pumped

to a pressure of 60 mTorr. The specimen was heated up to 850°C inside the vacuum furnace. The electron

emission characteristics of the specimen were measured in a high vacuum chamber under base pressure of lxl0 -e

Torr. A piece of quartz plate was used as a spacer between the specimen and a counter electrode to control the

distance between the anode and the cathode. A potential difference was applied between these two electrodes by a

computer controlled power supply. Emission currents were measured by a computer controlled picoammeter. The

specimen could be heated to various temperatures up to 860°C during the measurements by using a lamp with an

adjustable power supply.

NASA/CP--2001-210948 742 August 6, 2001
RESULTS AND DISCUSSION

The surface of carbon nanotubes appeared to have been coated with metallic nanoparticles as shown by the SEM

photograph (see Figure 1). Figure 2 demonstrates the electron emission hysteresis of the specimen at room

temperature. The applying electric field was firstly swept up, and the current was measured as a function of electric

field between the anode and the cathode. The electric field was then swept down, and the emission currents were

measured to be smaller than those corresponding to the same electric field during the ascending cycle.

Two mechanisms may have contributed to the measured I-E hysteresis. It is known that chemisorbed

molecules on a metal surface produce bonding states that have energy levels a few angstroms off the metal surface.

These states can increase the local tunneling current by forming a resonant tunneling of electrons (ref. 5). This

enhanced electron field emission occurs more significantly when the carbon nanotubes were cool and when the field

emission just started at a low level when the electric field is being swept up. It is less effective when the carbon

nanotube is cooling down from self-heating by emission current when the applied electric field is swept down. When

nanotubes emit electrons with a relatively high current density, the temperature of nanotubes may increase

dramatically to a high level (>1000K) due to the self-heating. When the temperature exceeds the desorption

temperature for the molecule, the binding states will then be removed. Re-adsorption during the descending cycle

may recover some of the chemisorbed molecules. The tunneling enhanced electron emission by chemisorbed

molecules, therefore, contributes to the shifting of the ascending I-E curve to the upper left direction.

The second mechanism that may contribute to the measured I-E hysteresis is the burn-out of sticking-out and

taller carbon nanotubes. Those taller carbon nanotubes were subjected to higher electric fields and emitted higher

electron current densities than those shorter or lower carbon nanotubes. With more current being carried through

fewer taller carbon nanotubes, some might be burned away when the applied electric field was swept to the

maximum value. When the electric field was swept down (decreasing with time), there were less tall carbon

nanotubes and therefore, the I-E curve for the descending cycle shifted to the lower right direction. Both the

tunneling effect and the burn-out mechanism appeared to support the measured I-E hysteresis at room temperature

with the descending I-E curve being below the ascending I-E curve. The dominant mechanism seemed to be the

enhanced tunneling effect due to chemisorbed molecules.

::"' "." ..:'f_!''"':!:!:i[:i:,.::i..,:::!:_':"

::::i }i ii

Figure 1. Scanning electron micrograph of brazed single-walled carbon nanotubes on an iron substrate.

NASA/CP--2001-210948 743 August 6, 2001
£9

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Figure 2. Electron field emission current as a function of applied electric field between the anode and the

cathode at room temperature. The dotted curve represents the cycle when the electric field was swept up.

The circled curve represents the other cycle when the electric field was swept down.

When the substrate was heated to elevated temperatures, however, the hysteresis behavior of the electron

emission I-E curve turned out to be the other way around, i.e., the I-E curve for the descending cycle with decreasing

electric field was measured to be above the I-E curve for the ascending cycle (see Figure 3). There should be a

thermionic emission component in the total emission current when the electric field was swept up to the maximum

level because of the high temperature of nanotubes due to the self-heating effect. When the substrate temperature

was at the room temperature, nanotubes could cool down from self-heated high temperature much quicker than when

the substrate was heated to an elevated temperature. The contribution by thermionic electron emission may thus be

higher during the descending cycle when the substrate was heated to an elevated temperature. The thermionic

emission current may have overwhelmed both the chemisorbed molecule enhanced tunneling current and the current

changing due to the burn-out of electron emitting nanotubes resulting in the hysteresis behavior shown in Figure 3.

Stepwise increases in electron emission current were sometime observed when the substrate was heated to

elevated temperatures (see Figure 4). This may be a result of the increasing flexibility and weaker bonding of

nanotubes leading to easier alignment of carbon nanotubes with the applied electric field causing the electron

emission current to change suddenly. As temperature increased, the bonds that kept nanotubes from being aligned

with the electric field became weaker. Therefore, when the temperature was high enough, it was possible that one

end of a CNT became free from the substrate surface and that this CNT was bent to align with the applied electric

field. The changing of the height of the tips of electron emitting carbon nanotubes might have contributed to the

stepwise increase and decrease of the emission currents. As shown in Figure 1, there was quite a coating of metallic

nanoparticles on the nanotubes. This coating may also contribute to the hysteresis of the I-E characteristics and the

step-wise changes of electron emission current. Further investigation is underway to explore the effects of coatings

on carbon nanotubes on electron field emission.

NASA/CP--2001-210948 744 August 6, 2001
10-4

_10 -6

(5

O

...1

10-8

0

cP

10-1o__ i i

0 0.3 0.6 0.9

E (V/urn')

i

1.2 1.5

Fig. 3. Electron emission current as a function of applied electric field between the anode and the cathode

with the substrate being heated to above 500°C. The dotted curve represents the cycle when the electric

field was swept up while the circled curve represents the cycle when the electric field was swept down.

10-5

10 -6

_10 -7

coOl 0-8

_.1

10 -9

(

10-lo

0

i i i

0.2 0.4 0.6

E_/_)

0.8

Fig. 4. Electron field emission current as a function of applied electric field between the anode and the

cathode with the substrate being heated to above 500°C. The dotted curve represents the ascending cycle

when the applied electric field increased with time. The circled curve represents the descending cycle when

the electric field decreased with time.

NASA/CP--2001-210948 745 August 6, 2001
CONCLUSIONS

Five mechanisms were adopted to explain the hysteresis of I-E characteristics of the electron field emission from

carbon nanotubes brazed by a TiCuAg alloy to an iron substrate. Those five mechanisms play different roles during

the electron emission of carbon nanotubes. At low ambient temperatures, the chemisorbed molecule enhanced

tunneling may be the dominant mechanism. The burn-out of nanotubes is negligible in most cases especially after

carbon nanotubes have been "conditioned", that is, defective or tallest carbon nanotubes have been intentionally

removed, because long term testing of emission current demonstrates a fairly stable emission current, indicating that

the damaging of emission sites are small (ref. 2). At higher ambient temperatures, however, because of the removal

of chemisorbed molecules on the nanotube surfaces, the contribution by thermionic emission and the relaxation of

nanotubes to align with the applied electric field might become dominant forces pushing the hysteresis of I-E curves

to a direction opposite that for substrate being at low temperatures. The sharp, sometimes stepwise, increases of

emission current might be the consequence of the changing shapes of the nanotubes in alignment with the electric

field as well as the absence of chemisorbed molecule enhanced electron emission by tunneling.

REFERENCES

1. Yahachi Saito and Sashiro Uemura: "Field emission from carbon nanotubes and its application to electron

sources", Carbon, 85, 7, 2000, p169.

2. W. Zhu et al.: "Large current density from carbon nanotube field emitters", Appl. Phys. Lett., 75, 6, 1999,

p873.

3. O. Gr/3ning et al.: "Field emission properties of carbon nanotubes", J. Vac. Sci. Technol. B, 18(2), 2000,

p665.

4. R. Saito, G. Dresselhaus, and M. S. Dresselhaus: Physical Properties of Carbon Nanotubes, Imperial

College Press, London, 1999 Journal of Applied Physics, Volume 85, Number 7, 1999, p3832.

5. Kenneth A. Dean and Babu R. Chalamala: "Field emission microscopy of carbon nanotube caps", J. Appl.

Phys., 85, 7, 1999, p3832.

NASA/CP--2001-210948 746 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

SECONDARY ELECTRON EMISSION OF MAGNESIUM OXIDE ON MULTIWALLED

CARBON NANOTUBES

Won Seok Kim, Young Min Shin, Hee Jin Jeong, and Young Hee Lee

Department of physics, Institute of Basic Science, Sungkyunkwan University, Suwon 440-746, South Korea

Whikun Yi, SeGi Yu, Taewon Jeong, Jeonghee Lee, Sunghwan Jin, Jungna Heo, and J. M. Kim

NCRI, Center for Electron Emission Source, Samsung Advanced Institute of Technology, P.O. Box 111, Suwon, 440-

600, South Korea

ABSTRACT

We have investigated the effects of electric fields on the SEE yield from the primary electron bombardment on

magnesium oxide (MgO) covering vertically aligned multiwalled carbon nanotubes (MWCNTs). We observe that the

SEE yield follows the general universal curve that the primary electron penetrating into the MgO sample should have

optimum primary electron energy to maximize the SEE. The SEE yield increases exponentially with increasing

negative bias applied to the sample. It is observed the SEE yield strongly depends on the thickness of MgO and

shows the maximum at some thickness. The SEE yield increases at least 22,000 at a thickness of 700A at a bias

voltage of - 1100 V and net primary electron energy of - 220 eV. The primary electrons bombarding the surface of

MgO film generate the secondary electrons in the film. The strong local field that is generated by the sharp tip of

vertically aligned multiwalled carbon nanotubes accelerates these electrons. Eventually, This gives rise to so called

Townsend avalanche effect, generating huge number of secondary electrons in the MgO film. These processes are

completely reproducible over the large area of the sample. Emission mechanism for such a high SEE yield will be

further discussed with energy spectrum analysis.

Keywords: Secondary Electron Emission, Multiwalled Carbon Nanotubes, Magnesium Oxide, Avalanche, Energy

Analysis.

INTRODUCTION

Since the advent of carbon nanotubes, they have been in the limelight and become one of the most promising

materials because of their structural, mechanical and electronic properties that have generated the wide range of

applications. The applications contain supercapacitor, hydrogen storages (ref. 1 to 3), field emission displays (FED)

(ref. 4, 5) and nano-devices (ref. 6 to 8). In particular, the mechanical strength, geometrical structure, and electrical

conductivity have motivated the application to the FEDs. The electric field at the apex of carbon nanotube becomes

extremely strong due to the high aspect ratio and such structural advantage can be applied to enhance the electric

field inside the dielectric materials as the external field is applied. The induced strong electric field inside the

material is beneficial to trigger the avalanche of the secondary electron emission, as the material is electronbombarded.

The emission of secondary electron from insulators has been of interest for a number of reasons (ref. 9). It is of

practical importance in devices such as electron multipliers and electron microscopes, as well as fundamental aspect

of electron-solid interactions. The secondary-electron emission (SEE) yield is defined as the ratio of emitted

secondary electrons to the primary bombarding electrons on the examined surface. The general process of emission

phenomenon is often viewed as three steps: (i) Bombarding electrons generate excited electrons in the solid by

mechanism of excited electron generation. (ii) The diffusive movement of excited electrons through the solid and

cascade multiplication are followed. (iii) Finally, the escape of some of the electrons into vacuum takes place.

In general, metals and semiconductors are relatively poor secondary-electron emitter, while insulators are good

secondary-electron emitters (ref. 10). In case of insulators, the minimum kinetic energy for a secondary electron to

escape is the electron affinity, which is the difference between the vacuum level and the conduction-band minimum.

The electron affinity is typically on the order of an electron volt for insulators. Since there are few conduction

electrons in insulators, the secondary electrons lose energy through the excitation of valence electrons into the

NASA/CP--2001-210948 747 August 6, 2001
conductionband.Thewidebandgappreventssecondary electronswithkineticenergylessthanenergygapfrom

participatingin suchelectron-electron cllision.Fortheseelectrons, electron-phonon andelectron-impurity

collisionsaremainlyresponsible fortheelectronenergyloss.Becauseoftheabsenceofelectron-electron scattering,

thesecondary electronlosesmuchlessenergyasit movesthroughthematerialndtheescapedepthbecomeslarge.

Therefore, ingeneral, theyieldsarehighininsulators. Forexample, themaximumSEEyield8maxis6.8forNaC1and

25forsingle-crystal MgO(ref.11).

Here,weobservetheSEEyieldfromtheMgOevaporated onMWCNTsiextremelyhighundertheelectric

field.TheSEEyieldstronglydependsonthethicknessof MgOandsaturateswithinsomelimitedrangesof

thickness. Eventually, themaximumSEEyieldexceedsavalueofatleast22,000atathicknessofabout700Aata

biasvoltageof - 1100V. Wealsomeasuretheenergyspectraof secondaryelectronsemittedfromthe

MgO/MWCNTs sampleforinvestigating themechanismoftheenormously highSEEyield.

EXPERIMENTAL

Tosynthesize MWCNTs, TiN(1000A)andnickel(1000A)aredepositedonsilicaglass(borondopedp-type)

withtheRFmagnetronsputter, whereTiNisusedasanadhesionlayerandnickelasacatalyst forthesynthesisof

theMWCNTs. TheMWCNTsarethermallygrownontheNiparticlesfromthereactionofseedinggasesofC2H2in

achemical vapordepositionreactor. MgOwithdifferent filmthicknessisdepositedontheMWCNTsbyelectron

beamevaporation.

TheSEEyieldismeasuredfromMgOonMWCNTswithourequipment, asshowninFig.1.TheMWCNT

substrateismountedonthesampleholderperpendicular toelectrongunthatcaneasilycontroltheamount andthe

kineticenergyofprimaryelectrons. BeforetheSEEmeasurement, thesampleisundernegativebias,whiletheother

remainingpartsaregrounded. Thechamber pressureisbelow1.0x 10.8torr.Voltagebalancingis carriedoutin

ordertoprevent secondary electronemissionbeforetheexperiment. Thekineticenergyofprimaryelectronsmustbe

higherthanthesamplebiastoovercomethepotential barrier.Theprimarycurrent(Ip)ismeasuredbyapplyinga

positivevoltageto thesampleinordertogatherincidentelectrons. Ratherthandirectlymeasuringthescattered

secondary electrons(Is),theelectroncurrentflowingintooroutof thesample(It)ismeasuredoverthestainless

sampleholder,sincethesecondary electroncurrent canbeeasilycalculatedbytherelationofIp=Is+It.Thus,SEE

yieldiscalculatedbythefollowingequation: 8=Is/Ip--1- (It/ Ip ).

Electron_ _

! gun _ J¢ Idaycup

Figure 1. Scheme for SEE experiment.

RESULTS AND DISCUSSION

Figure 2 shows the scanning electron microscope (SEM) images of MWCNTs grown at 650 °C for 25 min. Such

grown MWCNTs are well aligned vertically. The average diameter and length of MWCNTs grown at 650 °C is about

NASA/CP--2001-210948 748 August 6, 2001
50nmand20gmrespectively, asshownintheSEMimageofFig.2(a).Figure2(b)showstheMWCNTscovered

withMgOfilm.

:...:.._:::.::'.:._ ..........

._,.,...,,,..,_'I-.-... v.v. -.-........-..-.-.-.--.;,...,,.-..-'_:._

•:+:.:.:.;.:-:.:-:'..:.;.: ::.z.:.._..:-:.:.:_g..:.:.:-:.:-_.:._ (b)

::::::::: :::::::::::::::::::::::::::::::

iii:iiiiii::ii.?.!.i:i:i:i!!i_!!::_:_:i:iiiiii;!1ii:i:::i:?_!iii:/:!.!:i:iii!i:i:: ??Z:_

::::::::::::::::::::::::::::::::: .'::::U:;;"._:S::"" ::::::::::::::::::::::::: • •

:.:.::

•_: .. .. ....... ....,,,-_.... ._:1:_ ,'._ .....

': ..

Figure 2. The SEM images from (a) MWCNTs and (b) MWCNTs covered with MgO film.

SEE yield is investigated with the sample. Figure 3 shows the SEE yield as a function of the net electron energy,

which is the difference between the kinetic energy of primary electron and the potential energy of the sample. For

example, if the bias voltage of the sample varies from - 1500 V to 0 V under the fixed primary electron energy of

1500 eV, then the penetration energy into the sample from an electron gun varies from 0 eV to 1500 eV along the xaxis. In this experiment, we control two factors at the same time, which are the net electron energy and the applied

sample bias. From the experiment, the maximum SEE yield of -35 is obtained at 200 volt of net electron energy as

the primary electron energy is 1500 eV. The yield has stronger dependence upon the sample bias rather than the net

electron energy. Once the SEE yield reaches the maximum, it decreases exponentially with increasing the net

electron energy. This explains that although the primary electrons with higher kinetic energy penetrate deeper into

the sample and excite more electrons, those electrons within the escape depth can contribute to the SEE.

_i] _ --_--1500 eV

] / _ _ 1400 eV

/ --l 00eV

15_ ]z_. _, -- 1200 eV

"a 10 t ]7 t_Ax _ ---

100 0 100 200 300 400 500 600 700 800 900 100011001200130014001500

Net Electron Energy (eV)

Figure 3. The SEE yield of the MWCNTs as a function of net electron energy for different bias voltages.

There are many other parameters influencing the SEE such as diameter of MWCNTs, intra-tube space, and

thickness of MgO. We focus the effects of thickness of MgO on the SEE yield, since controlling the film thickness is

easier than other parameters.

MgO film is evaporated on the MWCNTs with a different thickness of 500, 700, and 900 ._. We observe that the

SEE yield shows the maximum with a thickness of MgO film of around 700, in which the SEE yield reaches up to

6500 under net electron energy of 220 eV in Fig. 4(b). Moreover, it is noted that such high yield is obtained as low

as -1000 V on the sample. We also see that the SEE yield above -1100 V exceeds the tolerance of our multimeter,

which is about 4 mA. We estimate the minimum SEE yield is as high as 22,000 at the sample bias.

NASA/CP--2001-210948 749 August 6, 2001
ThehighemissionyieldmentionedaboveisobtainedonlyfromMgO/MWCNTsample. NeitherMWCNTsnor

MgOyieldedsuchhighsecondary electrons. WhenwemeasuretheSEEyieldfromMWCNTandMgOseparately,

themaximumSEEvaluesreads1.2and4-5,respectively, whichagreecomparatively wellwithpreviousreport(ref.

12).Therefore, thehighyieldmaybeascribedtotheintegratedeffectsofgeometrical andelectricpropertiesof

MgO/CNT, i.e.,highaspect ratioofcarbonnanotubeandlowelectrondensityintheconductionbandofMgO.The

primaryelectronstransmittedintothesurfaceofMgOfilmgeneratethesecondary electronsinthefilm.Thestrong

localfieldgeneratedbythesharptipaccelerates theselectrons. Eventually, thisgivesrisetosocalledTownsend

avalanche effect,generatinghugenumber ofsecondary electronsintheMgOfilm.

70

6O

_o

_3o

c_

lO

10o / 1_o _o 3_o 400......... 500 600 700 800 900 100011001)301300

net electron energy [eV]

80(0

70(0

60(0

_5olo

_0(D

10(D

o

800[eVl

900[eVl

1000 [eV] \

50' 6 '50 '100"150' _I)" 250" N'0 ' 350" z_0' 450 ' 500" 550

netdeelron energy [eV]

lgCX),

10C_q

_SC_q

"_ 6C_q

1500 (eV)

_ I 1300(eV)

_00" 6" _0 " d_0 " 600 " 800 " 1_30 " 1_00 " 1_30 " 1600

primary deelron energy[eV]

Figure 4. SEE Yield of MWNTs coated by MgO thickness of (a) ~500 A, (b) ~700 A and (c) ~900 A.

The electron energy analyzer elucidates the above mechanism for the high SEE yields. Figure 5(a) shows the

secondary electron energy spectra for three different bias voltages at a fixed kinetic energy (1500 eV) of primary

electrons. The peak of over 1300 eV under the sample bias of - 1300 V is a back-scattered peak that occurs from

elastic scattering between primary electrons and MgO atoms. The maximum kinetic energy of the secondary electron

appears at 200 eV under the bias of-1300 V.

10x10 sIx0 10a "...................-1300 900 VV

_ -300v /\ L#

......

o 3oo 6oo 9oo 12oo

Kinetic Energy (eV)

.... l!

lsoo

(a)

:::::::izi::

[ ""2al;_e_" "l"""!!t'emn'sl]y_)

Figure 5. (a) Energy analysis of MWCNTs covered by MgO thickness of ~ 700 A, (b) Mechanism scheme of

the SEE on MWCNTs covered by MgO.

Each of the secondary emission peaks in spectra has a similar shape with a different position and intensity. Ihe

integrated intensity increases rapidly with increasing the bias voltage across the sample. This is the same result

obtained in Fig. 3 and 4, where the SEE yield is measured by the current flow (It) through the sample with respect to

the bias voltage. The peak position in the energy spectrum moves to higher energy at the higher bias voltage. The

energy band diagram, which explains the above spectral results, is drawn in Fig. 5 (b). A negative voltage (-V) is

applied to the nanotubes with respect to the analyzer, which is grounded. Under the influence of an electric field

across the MgO film, the secondary electrons generated inside the MgO film are accelerated to the surface, escape

from the surface, and are detected by the analyzer. At a higher bias voltage (-V2), the secondary electrons are

accelerated more strongly due to the high electric field across the MgO film which is generated by the sharp CNT

tips, resulting in a large peak located in a high kinetic energy region. The two facts that the general shape of each

NASA/CP--2001-210948 750 August 6, 2001
peakisalmost equivalent, regardless ofthebiaspotential, andthehighbiaspotential leadstoashiftofthepeaktoa

highenergy, stronglysupport theavalanchemechanisminducedbytheelectricfield.

CONCLUSION

WehaveinvestigatedheeffectsofelectricfieldsontheSEEyieldfromtheprimaryelectronbombardment on

magnesiumoxide(MgO)coveringverticallyalignedmultiwalledcarbonanotubes (MWCNTs). TheSEEyield

increasesxponentially withincreasingegativebiasappliedtothesample. It isobservedtheSEEyieldstrongly

dependsonthethicknessof MgOandshowsthemaximumatsomethickness. TheSEEyieldincreasesatleast

22,000atathicknessof700A atabiasvoltageof - 1100V andnetprimaryelectronenergyof- 220eV.The

primaryelectronsbombarding thesurfaceofMgOfilmgeneratethesecondary electronsinthefilm.Thestronglocal

fieldgenerated bythesharptipofverticallyalignedmultiwalledcarbonnanotubes accelerates theselectronsinthe

MgOfilm.Eventually, thisgivesrisetosocalledTownsendavalancheeffect,generatinghugenumber ofsecondary

electronsintheMgOfilm.Energyspectrumanalysisfurtherelucidates thismechanism.

ACKNOWLEDGMENTS

ThisprojectwassupportedbytheKoreanMinistryofScienceandTechnology (MOST)throughNational

ResearchLaboratory(NRL)programandinpartbyBrainKorea21(BK21)program.

REFERENCES

The corresponding author: Yong Hee Lee, Department of physics, Sungkyunkwan University, Suwon 440-

746, Kyunggi, South Korea, e-mail: leevoung@vurim.skku.ac.kr, fax: +82-31-290-5954, phone: +82-31-

290-7053

1. Dillon, A. C., et al.: Storage of Hydrogen in Single-Walled Carbon Nanotubes. Nature, vol. 386, March

1997, 377.

2. Lee, S. M., et al.: Hydrogen Adsorption and Storage in Carbon Nanotubes. Synthetic Metals. vol. 113,

2000,209.

3. Liu, C. et al.: Hydrogen Storage in Single-Walled Carbon Nanotubes at Room Temperature. Science, vol.

286, November 1999, 1127.

4. Dean, K. A. et al.: Current Saturation Mechanism in Carbon Nanotube Field Emitters. Appl. Phys. Lett, vol.

76, no. 3, January 2000, 375.

5. Collins, R G. et al.: Unique Characteristics of Cold Cathode Carbon-Nanotube-Matrix Field Emitters. Phys.

Rev. B, vol. 55, no. 15, April 1997, 9391.

6. Inganas, O. et al.: Carbon Nanotube Muscles. Science, vol. 284, May 1999, 1744.

7. Baughman, R. H. et al.: Carbon Nanotube Autuators. Science, vol 284, May 1998, 346.

8. Tans, S. J. et al.: Electron-electron Correlations in Carbon Nanotubes. Nature, vol. 394, August 1998, 761.

9. Schou, J. et al.: Ionization of Solid by Heavy Particles. Plenum, New York, 1993, 351-358.

10. Dekker, A. J.: Secondary Electron Emission. Solid State Physics, 6, 1958, 251.

11. Shih, A. et al.: Secondary Electron Emission Studies. Appl. Surf. Sci. vol. 111, 1997, 251.

12. Millet, J. M.: Secondary Electron Emission from Porous Solids. Phys. Rev. A, vol. 52, no. 1, July 1995, 433.

NASA/CP--2001-210948 751 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga, '

K. Kobashi, G.A.J. Amaratunga (Editors)

WORK FUNCTION AT THE TIPS OF MULTIWALLED CARBON

NANOTUBES

R.P. Gao, Z.W. Pan, Zhong L. Wang

School of Materials Science and Engineering, Georgia Institute of Technology,

Atlanta, GA30332-0245, USA

The unique geometrical structure of carbon nanotubes suggests that they are likely to be an ideal

field emitter for panel display application. Indeed, numerous field emission measurements show that the

carbon nanotubes exhibit superior field emission performance with required electric field much lower than

those of many other field emitters. For carbon nanotube emitters, most of the electrons are emitted from the

tips of the tubes, and it is the local work function that matters to the properties of the tube field emission.

Since the lack of suitable technical method to measure the work function of carbon nanotubes, the value of

work function of carbon nanotubes used in the literatures is mainly from the well-studied carbon or

graphite. Thus, it is necessary to experimentally measure the work function of carbon nanotubes.

In this study, we present experimental measurements of tip work functions of individual carbon

nanotubes. The experiment was done by a novel in-situ transmission electron microscopy technique [1,2].

Our results indicate that the tip work function show no significant dependence on the diameter of the

nanotubes in the range of 14-55 nm. Majority of the nanotubes (~75%) have a work function of 4.6-4.8 eV

at the tips, which is 0.2-0.4 eV lower than that of carbon. A small fraction of the tubes (~25%) have a work

function of-5.6 eV, about 0.6 eV higher than that of carbon. This discrepancy is suggested due to the

metallic and semiconductive characteristics of the tubes.

Key words: carbon nanotubes, work function, field emission

References:

[1] P. Poncharal, Z.L. Wang, D. Ugarte and W.A. De Heel Science 283, 1513 (1999).

[2] R.P. Gao, Z.W. Pan, Zhong L. Wang, Appl. Phys. Letts., in press.

Corresponding author:

Zhong L. Wang

School of Materials Science and Engineering

Georgia Institute of Technology

Atlanta, GA 30332-0245

e-mail: Zhong.wang@mse. gatech.edu

Tel: 404-894-8008

Fax: 404-894-9140

NASA/CP--2001-210948 752 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CHARGE TRANSFER, STRUCTURE, AND ELECTRONIC PROPERTIES OF KDOPED SINGLE-WALL CARBON NANOTUBES

Chulsu Jo and Young Hee Lee*

Department of Physics, Institute of Basic Science, Sungkyunkwan University, Suwon 440-746, Korea

ABSTRACT

Charge transfer, structure and electronic properties of K-doped single-wall carbon nanotubes bundle are

investigated by a first principles calculation. We choose (5,5) armchair single-wall nanotube bundle as pristine with a

diameter of 6.78 A, length of 2.46 A, and intertube distance of 3.35 A between the walls of nanotubes. The proper

numbers potassium atoms (n -- 1, 2, 3, and 5) are doped in the tube bundles. We considered that doped potassium

atoms are located on the interstitial channel sites and have symmetric configuration, which maximize the K-K

distance. Geometry of tube bundles is fully optimized. Intercalation of potassium atoms occurs lattice expansion and

the circular tubes are not nearly deformed. Binding energy is obtained by subtracting the total energy of the pure tube

bundle and self-energy of potassium atoms from the total energy of the K-doped system. Binding energies is negative

value up to K0.1C and vary to positive value from K0.15C. We also obtain total charge density distribution, amount of

charge transfer, and electronic band structure. For K0.05C, amount of charge transfer and Fermi level shift are 0.04e

per C atom and 1.5 eV, respectively. We can estimate the trend of conductivity enhancement induced by doping

through amount of charge transfer and Fermi level shift as function of K/C ratio.

Keywords: carbon nanotube bundle, binding energy, charge transfer, fermi level shift, conductivity.

INTRODUCTION

Single-wall carbon nanotubes (SWNT's) bundle can be doped with alkali metals, but little is known about the

exact positions of the dopants. In a SWNT bundle, the interstitial channels between tubes provide additional

intercalation sites(ref. 1, 2). Li- and K-doped SWNT tube bundles have been reported to enhance conductivity (ref.

3, 4) and hydrogen storage capacity (ref. 5, 6). Lee et al. found that the charge transfer induced by doping governs

the conductivity enhancement rather than the tube-tube interactions (ref. 3), while Ruzicka et al. concluded that the

tube-tube interactions influence conductivity by doping more strongly than the effect of the charge transfer (ref. 7).

It has been argued that whether the conductivity enhancement by doping is by an increase of charge transfer or the

tube-tube interactions. In recent experiments, one to one map between among Fermi level, charge carrier

concentration, and electrical conductivity was established. Fermi level shift by the redox reactions determined the

electrical conductivity of the Li-doped SWNT bulk material. The trends of Fermi level shift and conductivity were

very similar (ref. 8). These results indicate that the tuning of the Fermi level and charge transfer induced by doping

controls the conductivity. This interesting issue of conductivity enhancement, however, has not been clarified

theoretically up to now. The issues on the structures of K-doped SWNT bundles have been challenged by the

theoretical calculations. The doped K-atoms are located in the interstitial sites of tube bundles (ref. 9, 10), with the

same tube spacings to be maintained even after the doping, whereas the tube walls are severely distorted at K0.1C (ref.

10). Fermi level also shifts into a conduction band of pure tube bundle after Li-doping. The charge transfer occurs

from Li atoms to tubes, where Li-doping is saturated at Li0.6C (ref. 11). However, no systematic relation of the Fermi

level to the conductivity is given in terms of doping concentrations. In this report, we perform the first principles

calculations systematically for K-doped SWNT bundle in terms of doping concentrations. The bundle lattice expands

up to 8 % at K0.1C and no tube distortion is observed, in contrast with the previous reports (ref. 10). We find a

systematic relation between the amount of charge transfer and the Fermi level shift as a function of K/C ratio. We

further demonstrate that the conductivity of the doped tubes is enhanced by the charge transfer from potassium atoms

to the tubes, rather than by the tube-tube interactions.

NASA/CP--2001-210948 753 August 6, 2001
METHOD

Wecalculatetheelectronicstructures ofaK-dopedSWNTbundleusingdensityfunctional theory(DFT)within

thelocaldensityapproximation (LDA).Theionicpotentialsaredescribedby a norm-conserving nonlocal

pseudopotential generatedby Troullier-MartinsKleinman-Bylander form.Exchange-correlation functions

parameterized bythePerdew-Zunger schemeareused.Aplane-wave basisetwithanenergycutoffof55Ryisused.

IntegrationverBrillouinzoneiscarriedoutusing10k-pointssamplingwhichisgenerated bytheMonkhorst-Pack

scheme. Theconvergency criterionforstructureoptimizationisthatallforcesbe< 0.005 Ry/A. The energies are

converged to 10 -3 eV.

The tube bundle is modeled by a hexagonal unit cell with y=60 degree. We choose (5,5) armchair SWNT bundle,

where 20 carbon atoms are included in the primitive cell. The box size of the hexagonal unit cell is initially chosen to

be 10.13 A x 10.13 A x 2.46 A, and the tube diameter of 6.78 A and the intertube distance of 3.35 A between the

tube walls. The proper numbers of potassium atoms (K×C, x = 0.05, 0.1, 0.15, and 0.25) are doped in the tube

bundles. We consider that the doped potassium atoms are located at the interstitial channel sites and have a

symmetric configuration so as to maximize the K-K distance.

RESULTS AND DISCUSION

Figure l(a) shows the total energies of fully relaxed geometry as a function of the intertube distance. The

interstitial channel sites of doped K-atom are represented as solid circles. In our calculation, the intertube distance of

pure tube bundle is optimized to 3.2 A. The lattice constant increases with increasing potassium content, as shown in

Fig. l(b). The lattice constant expands by 8 % at K0.1C and saturates to 25 % at K0.15C. Their intertube distances

expand to 4.0 and 5.65 A, respectively. These values are comparable with the interlayer distance of 5.35 A at K0.125C

in graphite (stage I) (ref. 12). We note that the distortions of the tube wall were negligible for all cases up to K0.25C

with full relaxation, contrary to the previous reports that the tube walls were severely distorted by about 10 % - 20 %

for the doped (10,10) SWNT bundle (ref. 10, 11). This is perhaps due to the inappropriate treatments of the lattice

expansion by the doping.

(a) (b) 13.o

-3757.4 f

_', 12.5

_-37575[ __120

10.5

--I

-3757.81 10.0

i i = i = I

3.4 3.6 Intertube3..distance 4.0 ( A4.2) 4.4

i . i . i . i . i . i

0.00 0.05 0.10 0.15 0.20 0.25

x (KxC)

Figure 1. (a) The relaxed total energies as a function of the intertube distance for K0.05C. (b) Lattice constants

as a function of doping content.

Binding energies are also presented in Fig. 2. We define the binding energy as follows (ref. 9, 13),

E b = Et(CNT+ K) -Et(CNT) -nxEs(K ),

where Et (CNT+K) and Et (CNT) are the total energy of tube bundle with and without potassium atoms inside,

respectively. The Es (K) is the self-energy of potassium atom and nK is the number of potassium atoms. The binding

energies maintain negative values up to K0.1C and become positive from K0.15C, i.e., the doping of potassium atoms

up to K0.1C is exothermic process. Further increase of doping concentration requires an extra energy. In general, the

binding energy is correlated with the K-K separation distance. The K-K separation distance becomes shorter with

NASA/CP--2001-210948 754 August 6, 2001
increasingK-doping, whereasthebindingenergydecreases uptoK0.1C. Withfurtherincreasingthedopingcontent,

thebindingenergykeepsincreasing, followingsimilartrendstotheelectrostatic repulsiveenergy. Wenoteherethat

thebindingenergyisnotsolelydeterminedbytheK-Kseparationdistance. Atlowdopingcontent, therepulsive

energyisnegligibleduetolargeK-Kseparation distance. Instead, theattractiveinteractions betweentubewallsand

potassiumatomsgovernthebindingenergy. Furthermore thelectrostatic repulsiveenergyisminimizedbyreducing

theamount ofchargetransferfrompotassiumatomstothetube(ref.14),asshowniTable1.TheC-Kandthetubetubedistancesincrease, withincreasingdopingconcentration. Thissuggeststhattheconductivitymaynotbe

correlatedwithtube-tubeinteractions. Yet,K0.25Cadoptstoomanypotassiumatomsin theinterstitial sites,

increasingthebindingenergy. Thissuggests thatexcessivedopingmaynotbeadvantageous tthehydrogenstorage.

Incaseofgraphite, theamount ofhydrogengasstoredinK-intercalated graphitearesaturatedatK0.125C(ref.12).

1.0

o.8

>, 0.6

L,.

0.4

0.2

:._ 0.0

¢..

.m

nn -0.2

-0.4 O

= . | . | . | . |

0.05 0.10 0.15 0.20 0.25

X (KxC)

0.25

0<

0.20 o

o.15 v

v,- 0.10

Figure 2. Binding energies as a function of doping content.

K/C

K0.05C

K0.1C

K0.15C

K0.25C

Intertube distance (A) K-K (A) C-K (A) Mulliken charge (e)

3.95 10.73 (6.96) 2.96 0.81

4.00 6.20 2.98 0.78, 0.78

5.65 (5.35) 6.22 (4.92) 3.13 0.45, 0.49, 0.49

5.83 3.90 3.25 0.53, 0.47, 0.47, -0.20, -0.20

Table 1. Intertube distance, nearest neighbor K-K distance, nearest neighbor C-K distance, and Mulliken

charge of potassium atoms in K-doped tube bundle. The values in the parentheses are from intercalated

graphite (ref. 12).

It has been suggested that in Li-doped SWNT, the Fermi level shift is strongly correlated to the electrical

conductivity with Li/C ratio. The Fermi level shift and electrical conductivity have similar trends with increasing

Li/C ratio (ref. 8). In our calculations, we present the Fermi level shift and total charge transfer as a function of K/C

ratio, as shown in Fig. 3(a). The Fermi level shift increases abruptly with increasing K/C ratio up to K/C=0.05 and

saturates to near 2 eV. The similar trends are shown in charge transfer. This result strongly suggests that the

conductivity of the doped tubes is governed by the amount of charge transfer. Experimentally the normalized

conductivity of K-doped and Li-doped SWNT at the ratio of 0.04 is about 13 and 2, respectively (ref. 4, 8),

suggesting that the conductivity is enhanced more effectively by the K-doping rather than by the Li-doping at low

doping content. We note that a partial charge transfer is observed even at low doping concentration. The amount of

Mulliken charge transfer per carbon atom are 0.04e, 0.08e, 0.07e, and 0.05e for KxC (x = 0.05, 0.1, 0.15, and 0.25),

and the corresponding Fermi level shifts are 1.6 eV, 1.7 eV, 2.3 eV, and 2.0 eV. These values are in good agreements

with the previous estimation of the Fermi level shift of about 2 eV for charge transfer of 0. le per carbon atom (ref. 3).

Figure 3(b) shows the DOS at the Fermi level as a function of K/C atomic ratio. The DOS increases rapidly and

saturates at K0.05C, revealing similar trends to the Fermi level shift and the change of charge transfer, as shown in Fig.

NASA/CP--2001-210948 755 August 6, 2001
3(a).Thissuggests thathigherTccannot beexpectedvenwithhigheralkalimetaldopinginthetube,letalonethe

existenceofthesuperconductivity inthetube.

2.5 0.4

2.0 0.3 "'- ..-'"

-....

1.5 t 1"0_ "",e"

1.0 l o.5 g a 0.2

0.5 0.1

o.o o.o

0.0 ' '

0.00 0.05 0.10 0.15 0.20 0.25 0.00 0.05 0.10 0.15 0.20 0.25

K/C atomic ratio K/C atomic ratio

Figure 3. (a) The corresponding Fermi level shift and the amount of charge transfer of K-doped SWNT

bundle as a function of doping content. (b) The DOS at the Fermi level in terms of doping content.

SUMMARY

In summary, we have performed the first principles calculations on the total energy, charge transfer, and

electronic structures of K-doped SWNT bundles. We note that the distortions of the tube wall were negligible for all

cases up to K0.25C with full relaxation, contrary to the previous reports. The binding energy is not solely determined

by the K-K separation distance. The attractive interactions between tube walls and potassium atoms govern the

binding energy by the enhanced ionic bonds between potassium and nanotubes up to K0.1C. At higher potassium

doping the repulsive energies dominate the binding energy. Partial (not complete) charge transfer is occurred from

potassium atom to the tube and this trend becomes more severe with increasing doping concentration. The Fermi

level shift and the amount of charge transfer increase abruptly with increasing K/C ratio up to K/C=0.05 and

saturates to near 2 eV. This result strongly suggests that the conductivity of the doped tubes is governed by the

amount of charge transfer rather than the tube-tube interactions.

ACKNOWLEDGEMENTS

This project was supported by the Korean Ministry of Science and Technology (MOST) through the National

Research Laboratory (NRL) program and New Frontier program in Korea and in part by Brain Korea 21 (BK21)

program.

REFERENCES

Corresponding author. Young Hee Lee: Department of Physics, Institute of Basic Science, Sungkyunkwan

University, Suwon 440-746, South Korea, E-mail address: leeyoung@yurim.skku.ac.kr,

fax: +82-31-290-5954, phone: +82-31-290-7053

1. Lee, R.S., et al.: Conductivity Enhancement in Single-Walled Carbon Nanotube Bundles Doped with K and

Br. Nature, vol. 388, Jul. 1997, 255-256.

2. Rao, A.M., et al.: Evidence for Charge Transfer in Doped Carbon Nanotube Bundles from Raman

Scattering. Nature, vol. 388, Jul. 1997, 257-259.

3. Lee, R.S., et al.: Transport Properties of a Potassium-Doped Single-Wall Carbon Nanotube Rope. Phys. Rev.

B, vol. 61, no. 7, Feb. 2000, 4526-4529.

4. Claye, A.S., et al.: Structure and Electronic Properties of Potassium-Doped Single-Wall Carbon Nanotubes.

Phys. Rev. B, vol. 62, no. 8, Aug. 2000, R4845-R4848.

NASA/CP--2001-210948 756 August 6, 2001
5. Chen,R;Wu,X.;andTan,K.L.:HighH2UptakebyAlkali-DopedCarbonNanotubes underAmbient

PressureandModerateTemperatures. Science, vol285,Jul.1999,91-93.

6. Yang,R.T.:HydrogenStoragebyAlkali-DopedCarbonNanotubes-Revisited. Carbon, vo138,2000,623-

641.

7. Ruzicka, B.,etal.:Optical anddcConductivity StudyofPotassium-Doped Single-Walled CarbonNanotube

Films.Phys.Rev.B,vol.61,no.4,Jan.2000,R2468-2471.

8. Jouguelet, E.;Mathis,C.;andPetit,R:ControllingTheElectronicProperties ofSingle-wall Carbon

Nanotubes byChemical Doping.Chem. Phys.Lett.,vol.318,Mar.2000,561-564.

9. Miyamoto, Y.,etal.:IonicCohesionadElectronDopingofThinCarbonTubuleswithAlkaliAtoms.Phys.

Rev.Lett.,vol.74,no.15,Apr.1995,2993-2996.

10.Gao,G.,etal.:PositionofKAtomsinDopedSingle-Walled CarbonNanotubeCrystals. Phys.Rev.Lett.,

vol.80,no.25,Jun.1998,5556-5559.

11.Zhao,J.,etal.:First-principles StudyofLi-Intercalated CarbonNanotubeRopes. Phys.Rev.Lett.,vol.85,

no.8,Aug.2000,1706-1709.

12.Enoki,T.,etal.:Hydrogen-Alkali-Metal-Graphite TernaryIntercalation Compounds. J.Mater.Res.,vol.5,

no.2,Feb.1990,435-466.

13.Theintercalationenergy, similartothebindingenergy, maybedefinedbysubtractingtheenergyofapure

tubebundlefromthetotalenergyoftheK-intercalated tubebundle.See,forinstance, asshownireference

11.

14.Mullikenchargeslocalizedateachatomarecalculatedbyintroducingatomicorbitalbasisdensity

functional methodwithinlocaldensityapproximation.

NASA/CP--2001-210948 757 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ELECTRICAL RESISTIVITY OF ANISOTROPICALLY ALIGNED SINGLEWALLED

NANOTUBE FILM

Dong Jae Bae, Kae Hyeok An, Jeong Mi Moon, Seong Chu Lim and Young Hee Lee

Department of Physics, Sungkyunkwan University, Suwon, Kyunggi 440-746, Republic of Korea

Keun Soo Kim, Young Soo Park

Department of Semiconductor Science and Technology, Jeonbuk National University, Jeonju 561-756,

Republic of Korea

ABSTRACT

We present a simple way of preparing anisotropically aligned nanotube fihns by the simple mechanical rubbing on

film as a form of mat and the measured data on the electrical resistivity of aligned CNTs films. High anisotropy of 5

to 15 in the resistivity was obtained depending on the sample thickness and sample preparation conditions. In

degassed sample, a monotonic decrease of the resistivity with increasing temperature cannot be explained by either

the long-range Coulomb interactions or the inclusion of short-range interactions, which is different from an isolated

singlewalled carbon nanotube. Therefore we tried fittings using several theoretical and/or empirical formulas on the

temperature dependence of resistivity. We found that the measured resistivity showed neither metallic nor

semiconducting behavior. We found the best fitting formula with two terms, which combined nonlinear power law

term to the linear term in the temperature dependence. We obtained the power index of 0.446 that gave rise to the

Luttinger parameter, g=0.186, with bulk contact. The contribution of the linear term is less than that of power term

over the measured temperature range. Here in order to explain the transport mechanism we propose that the

temperature dependence of resistivity in the CNTs film would be similar to that of an isolated metallic singlewalled

nanotube, although we must include some contribution of linear temperature dcpcndencc from the Fermi liquids and

that the detailed structure of the CNTs film could be modeled by the simplified network of the effective resistors, in

which each resistor is connected by a CNT resistance and a junction resistance in series in mean field concept.

Keywords: nanotube, resisfivib; anisotropy, film, Luttinger liquids, Fermi liquids

Corresponding author : Young Hee Lee

E-mail : ]..e..e.y'._.!._@s.p..r..c..2. =_h9..n..b....u k.:.a..¢. :..lg..

Fax : +82-63-270-3585

Tel : +82-63-270-3336

NASA/CP--2001-210948 758 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

OXIDATION OF CARBON NANOTUBES AND CHANGE OF THEIR ELECTRONIC

AND GEOMETRIC SRTUCTURES

Seong Chu Lim, Hee Jin Jung, Dong Jae Bae, Eun Kyung Suh

Semiconductor Physics Research Center, Jeonbuk National University

JeonJu, Jeonbuk, 561-756, Republic of Korea

Young Min Shin, Kay Hyeok An, Young Hee Lee*

Department of Physics and Institute of Basic Science, Sungkyunkwan University

Suwon, Kyunggi, 440-746, Republic of Korea

Jaewu Choi

Center for Advanced Microstructures and Devices, Louisiana State University

Baton Rouge, L. A., 70806, U. S. A.

ABSTRACT

We have studied oxidation of vertically aligned multiwalled carbon nanotubes (MWCNTs) grown by thermal

chemical vapor deposition (CVD). Ultraviolet photoelectron spectroscopy (UPS) and transmission electron

microscopy (TEM) are applied to characterize electronic and structural changes. Vertically aligned CNTs are

cleaned by annealing about 700 °C for a few minutes inside an ultrahigh vacuum system (UHV). Oxidation with

molecular and atomic oxygen shows that CNTs are very resistant to oxygen exposures. However, when they are

exposed to massive amount of atomic oxygens, the UPS and TEM present that the electronic structure and

morphology of CNTs are utterly damaged. 2p4z electron states below 3 eV from the Fermi level have almost

disappeared, and 2p-_ electron states around 8 eV are significantly affected. Such variations in valance band

modify the metallic CNTs to be semiconducting. Although the annealing above 700 °C made a slight recovery in

the DOS near the Fermi energy and the valance band structure seems to be metallic, a complete recovery of

electronic structure is not achieved. We also noticed that the collective _+Tz plasmon excitation was observed at

about 23 eV. This value is lower than typical plasmon of 27 eV. The lowering of collective _+Tz excitation energy

implies that the oxidation severely distorted the ordering of the c-planes of CNTs and consequently, the

contribution of 2p4z electron to the collective _+Tz excitation became weaker. The TEM images illustrate that the

cap of CNTs is eroded completely and left open, whereas the body remains rather resistant to oxidation. Through

the cap opening process, nickel particles embedded inside the cap are exposed to vacuum and are observed from

the UPS. We propose that the emergence of nickel particles at the tip of CNTs is a serious impediment in achieving

the dependable operation of field emission array (FEA).

Keywords: Carbon nanotubes, Oxidation, Amorphous carbon, Cap opening, Nickel oxides.

INTRODUCTION

The reason that carbon nanotubes have attracted a great deal of interest is the fact that CNTs have very useful

properties and the electronic properties can be varied depending on chiral angle and diameter of CNTs (ref. 1 and

2). Therefore, the control of chiral angel and diameter of individual CNT during a growth is very essential for

useful applications. However, unfortunately, to our best knowledge, it is practically impossible up to now to

modulate their chiral angel and tube diameter during growing process. So, various ways to change the electrical

properties have been attempted such as gas-phase doping of alkali metal and adsorptions of gas adsorbates (ref. 3

and 4). In particular, oxygen among those gases has been used for purification by selectively burning other phases

of carbon materials and for the capillary filling of CNTs with metals by oxidatively opening the cap (ref. 5 and 6).

In addition, the exposure of the CNT emitters to oxygen leads to degrade considerably field emission currents (ref.

7). The conductivity and thermo-electric power change dramatically upon the adsorption and desorption of oxygen

(ref. 8 and 9). Since oxygen has profound influence on the electronic and geometric structures together with the

fact that the CNTs are naturally doped with oxygen during the sample preparation, understanding the effects of

NASA/CP--2001-210948 759 August 6, 2001
1500-

•,,i 1200-

'_ 900-

"_ 600-

300-

0

25 2'0 1'5 1'0 5 0

Binding energy (eV)

1580 L

380 L

200 L

140 L

100 L

OL

Figure 1. LIPS after thermal oxidation of CVD-grown MWCNTs for different oxygen doses.

oxygen on CNTs has significant meanings not only for the application and modulation of CNTs, but also for

perceiving their intrinsic properties.

EXPERIMENTS

In this study, we grew vertically aligned MWCNTs using thermal CVD. The growth parameters such as

substrate temperature, pressure, and flow rate of C2H2 were descried in detail in the literature (ref. 10). The CNT

film is transferred into the UHV system with a pressure below 2 x 10 -9 torr. We outgassed the CNTs above 1020 K

for three minutes. The pressure of the UHV system was increased to 3 xl0 -9 torr and dropped to a base pressure

immediately during the outgassing,. The UPS measurement has been conducted in a synchrotron radiation center,

Center for Advanced Microstructures and Devices (CAMD) at Louisiana State University. Monochromatic photon

has been obtained from dispersion of 3 m toroidal grating. The resultant combinational resolution of the

photoemission spectroscopy is 100 meV. The kinetic energy of photoelectrons is measured with an energy analyzer

with a maximum resolution of 10 meV. In order to see the change of morphology, we collected the oxidized CNTs

by simply scratching the substrates and prepared them for the TEM measurement.

We used oxygen gas with a research purity and precisely controlled the flow of oxygen gas with a leak valve.

1200

(a) 1200

"_ 900. _,'_,1000 -

,._ = 8oo

.

_[ 600. =_

,_ 600

_00]_30 _4ooj 2OOOL

20 1'5 1'0 ; ; 25 2'0 1'5 1'0 5 0

Binding energy (eV) Binding energy (eV)

Figure 2. (a) Oxygen-adsorbed CNTs near liquid temperature. (b) Oxygen-desorbed CNTs at elevated

temperature.

NASA/CP--2001-210948 760 August 6, 2001
The pressure of the vacuum system is brought up to lxl0 -6 torr after oxygen exposure and then we elevated the

substrate temperature above 1020 K. After each thermal oxidation, the heating was turned off and the vacuum

system was evacuated down to a base pressure again. Figure 1 presents the UPS spectra at various dosages of

oxygen. However, interestingly, there is no oxygen peak observed and no difference is seen either. It is reported that

the oxidation of the CNTs is more active on the cap than the body and the oxidation temperature is about 950 K

(ref. 11). In addition, the adsorption of oxygen molecule at the defect sites is exothermic (ref. 11). Oxygen

molecules adsorbed on the tube wall are physorbed and therefore these molecules do not change the tube structures

even at high temperature. Oxygen molecules adsorbed on the tub edge or the cap is chemisorbed. These molecules

will be desorbed in a form of CO or CO2 molecule at high temperature. Since MWCNTs have many concentric

shells, the outmost shell may be disintegrated with the latter reactions. However, we observe no change in the UPS

measurements. This suggests that although such oxidative erosion removes an outmost oxidized layer, this process

simultaneously brings a new inner layer from MWCNTs to vacuum. Consequently, the MWCNT structure appears

the same all the times. In order to see the effect of oxygen adsorption, the substrate was cooled down to liquid

nitrogen temperature. The substrate temperature was about 90 K. Exposure of CNTs at 90 K to oxygen gases is

shown in Fig. 2. As we keep supplying oxygen gases to the chamber, the oxygen peak is gradually developed near

13 eV.

So far, we have not seen a rigorous change in the DOS from oxidation. Atomic oxygen would be more effective

in inducing more serious destruction of the CNTs. Oxygen molecules are decomposed by a Mo filament, which is

located in front of the CNT film. The Mo filament was heated up to 1500 ~ 2000 °C. We applied the negative bias

of 15 volt in order to drive atomic oxygen to the CNT film. Since the separation between the filament and CNT

film is about 15 cm, we consider that the bias is not strong enough to physically sputter the CNTs. Although the

exact temperature of the substrate was not measured, the substrate temperature is higher than room temperature

due to the radiational heating from the Mo filament, but still much lower than the oxidation temperature. The

CNTs are exposed to an atomic oxygen to lxl0 -6 torr for 45 minutes and subsequently heated above 900 K for 10

minutes to remove oxygens from the surface. We saw no remarkable changes in the DOS, similar to the liquid

nitrogen experiment. These results suggest that oxygen dosage may not be enough to see the changes in the UPS.

We next introduce more oxygens in the chamber. At the same oxygen pressure, we increased an exposure time

of atomic oxygens to 70 minutes. After massive oxidation, considerable variations are observed as shown in Fig. 3.

From the pristine CNTs, the electrons contributing to 2~7.6 eV are assigned to 2p-_ electrons, which overlap with

the top of 2p-_ electrons (ref. 13). After the CNTs are massively exposed to oxygen,

5000-

4000-

om

.d 3000.

+" 2000-

1000-

.

60 50 40 30 20 10 0

Binding energy (eV)

Figure 3. UPS spectra of massively oxidized CNTs after annealing at various temperatures.

NASA/CP--2001-210948 761 August 6, 2001
iii iiiiii !!!iiiiiiiiii!i!i!!!!!!!!

iiiiii!i!!!iiiiiiiiiiiiiiiiiiiiii!iiiiiii__i_____________

:o

:i'i'i'i'i'i'i""""""i'i'i'i'i'i'i'i_"""""i'i'i'-i'i'i'i :i:

iiiiii i"liil "iiiiiiii

Figure 4. TEM images of CNTs (a) before and (b) after oxidation.

electronic structure is changed dramatically. The DOS of 2p-_ electrons near 3 eV almost disappears and that of

2p-cy is severely affected. The origin of those changes is the transformation of the graphitic layer to amorphous

carbons with 5_O 3 network (ref. 12). The effect of the increase of the amorphous carbons on plasmon excitation will

be discussed later. As the CNT caps are oxidatively opened, edges and defect sites are predominantly occupied with

the cybond not only between oxygen and carbon atoms, but also between carbon atoms. Therefore, before we

started annealing, the UPS spectra show strong 2p-cy electron states. However, with increasing the temperature of

CNTs, the oxygen atoms desorb from the surface of CNTs through the formation of CO and CO2. Annealing the

CNTs sublimates oxygens and reconstructs the morphology of CNTs. It is clearly shown in Fig. 3 that the slow

development of 2p-_ electron state is observed as the annealing temperature is raised.

Peaks appeared 23 and 46 eV are attributed to collective cy+_ plasmon excitation. It is reported that the typical

collective cy+_ excitation from CNTs is around 27 eV and this plasmon excitation contains information about

structures because cy+_ excitation can shift upon the diameter and the number of wall of CNTs (ref. 14). The shift

of collective cy+_ excitation energy to lower energy is correlated with the weakening of _ electron contribution.

These _ electrons in graphite participate in weak interlayer interaction. Therefore, the shift of collective cy+_

excitation to lower energy explains that the transformation of graphene sheet to amorphous carbon degrades the

ordering along c-plane. The destruction of c-plane ordering is well shown TEM images in Fig. 4(b). In addition,

amorphous carbon is seen around CNTs. Comparing to unoxidized CNTs shown in Fig. 4(a), the oxidation opened

the cap and damaged the body of CNTs, whereas the degree of oxidation is less significant on the body.

Strong peaks near 28 eV are attributed to nickel oxides. Nickel is usually used as catalyst and it migrates along

the tube during the growth. The appearance of nickel explains the cap is opened. For the operation of field

emission array (FEA), the existence of oxygen inside FEA is extremely dangerous not only because oxidation

exhausts DOS near the Fermi energy, resulting in increase of turn-on voltage, but also because oxidation exposes

nickel particles to vacuum, which would lead FEA to vacuum breakdown at high emission current. However, the

peak around 38 eV is not adequate for oxygen. We were not able to identify the peak at this moment.

CONCLUSION

We have studied the influence of oxygen on the electronic and geometric structures of CNTs. MWCNTs have

shown a metallic band structure initially. However, after they are heavily exposed to atomic oxygen, their band

structure becomes semiconducting, whereas it changes back to metallic band structure again when de-oxygenation

is carried out. These variations in DOS result from structure transformation. The outer wall of CNTs becomes

amorphous phase and cap structure is completely vanished. Surprisingly, nickel emerges to vacuum during the

oxidation. We propose that protruding of nickel at the apex of open CNTs is an obstacle to be overcome for the

dependable operation of FEA. In addition, the decrease of ordering along the c-planes contributes to the lowering

of the collective cy+_ excitaion.

NASA/CP--2001-210948 762 August 6, 2001
ACKNOWLEDGEMENT

Thisworkis supported by the Ministry of Science and Technology through the National Research Laboratory

program and, in part, by the Brain Korea 21 program.

REFERENCES

*To whom correspondence should be addresses; electronic mail: leeyoung@_/Tarim.skku.ac.kr, phone number: 82-

31-290-7053, fax number: 82-31-290-5954.

1. L.C. Venema, et al., Spatially resolved scanning tunneling spectroscopy on single-walled carbon

nanotubes. Phys. Rev. B. vol 62. no. 8, Aug. 2000, 5238.

2. D.L. Carroll, R Redlich and P. M. Ajayan, Electronic Structure and Localized States at Carbon Nantube

Tips. Phys. Rev. Lett. vol. 78, no. 14, Apr. 1997, 2811.

3. R.S. Lee, et al., Conductivity enhancement in single-walled carbon nanotbe bundles doped with Kr and

Br. Nature vol. 17, Jul. 1997, 255.

4. C. Zhou, er al., Modulated Chemical Doping of Individual Carbon Nanotubes. Science vol. 290, Nov.

2000, 1552.

5. R M. Ajayan, et al., Opening carbon nanotubes with oxygen and implications for filling. Nature vol. 362,

Apr. 1993, 522.

6. R M. Ajayan and S. Iilima, Capillarity-induced filling of carbon nanotubes. Nature vol. 361, Jan. 1993,

333.

7. S.C. Lim, et al. Effect of Gas Exposures on Field Emission Properties of Carbon Nanotube Arrays.

(unpublished).

8. R G. Collins, et al., Extreme Oxygen Sensitivity of Electronic Properties of Carbon Nanotubes. Science

vol. 287, Mar. 2000, 1801.

9. G.U. Sumanasekera, et al., Effects of Gas Adsorption and Collisions on Electrical Transport in SingleWalled Carbon Nanotunes. Phys. Rev. Lett. vol. 85. no 5, Jul. 2000, 1096.

10. Y. C. Choi, et al., Controlling the diameter, growth rate, and density of vertically aligned carbon

nanotubes synthesized by microwave plasma-enhanced chemical vapor deposition. Appl. Phys. Lett. vol.

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11. X. Y. Zhu, et al., Adsorption and Desorption of an 02 Molecule on Carbon Nanotubes. Phys. Rev. Lett. vol.

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12. H. Ago, et al., Work Functions and Surface Functional Groups of Multiwall Carbon Nanotunes. J. Phys.

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NASA/CP--2001-210948 763 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

ELECTRONIC STRUCTURE MODIFICATION OF MULTIWALLED CARBON

NANOTUBES BY ION BOMBARDMENT AND THERMAL TREATMENT

Jaewu Choi*

Center for Advanced Microstructures and Devices, Louisiana State University, 6980 Jefferson Hwy, Baton Rouge,

LA 70806, (Ph) 225-578-9356, (fax) 225-578-6954, (e-mail)jchoil@lsu.edu

Young Chul Choi

Material Tech. Laboratory, Corporate Research and Develop Center, Samsung SDI, Suwon, 442-391, Korea

and Young Hee Lee

Department of Physics, Sungkyunkwan University

Suwon, 447-746, Korea

ABSTRACT

The electronic structure of the vertically aligned multiwalled carbon nanotubes has been modified by Ar and Ne

ion bombardment and thermal treatment, and probed using ultraviolet and X-ray photoemission spectroscopy. With

increasing ion bombardment dose, the density of states near the Fermi level is largely increased. The carbon ls core

level spectra indicate that the sp 3 hybridization character is significantly developed at the initial exposure of the ion

bombardment. Thermal-treated carbon nanotubes after the ion bombardment largely reduce the density of states near

the Fermi level. This letter shows that the ion bombardment on carbon nanotubes can be utilized to get more metallic

property while the thermal treatment after ion bombardment can be applied for obtaining more semiconducting

carbon nanostructures.

Keywords: carbon nanotube, electronic structure, photoemission, modification

INTRODUCTION

The elementary particle irradiation can be utilized for modification of materials properties. For example,

Banhart group studied the structural phase transition of carbon nanostructures from graphite to diamond using highenergy electron irradiation at high temperature [ 1,2]. Defect-induced oxidative etching of graphite has produced

some unusual decorative etching patterns on the graphite depending on the ion-bombardment energy [3]. One of the

interesting nanostructures is carbon nanotubes (CNTs). Carbon nanotubes have unique properties, and especially the

metallic property of carbon nanotubes depends on chirality (n,m)) (metal when n - m = 3k, and semiconductor

when rt-m ¢ k where k = integer) [4-7]. The chirality-dependent electronic density of states (DOSs) near the

Fermi level from a single carbon nanotube has been observed using scanning tunneling spectroscopy (STS) [8]. The

electronic structure of carbon nanotubes has been studied using photoemission spectroscopy for single-walled CNT

films [9] and for mutliwalled CNT film [ 10]. Furthermore, the electronic band dispersion along the CNT axis from

vertically aligned multiwalled-carbon nanotubes was observed using angle-resolved photoemission spectroscopy

[ 11]. Those experimental results are obtained from CNT films with random chiarlity. However, identifying the

metallicity of the carbon nanotubes is of both fundamental and technological importance in applications of CNTs,

such as transistors [12,13], heterojunction device [14] and especially nanoscale electronic devices [15]. However,

controlling the chirality is not trivial with the current sample-preparation technology.

In this study, we focused on the modification of the electronic structure by the ion bombardment and

thermal treatment to get the desired electronic structure of carbon nanotubes, which is less related to individual

electronic structure of carbon nanotube. This could help open doors for engineers to fabricate an intramolecular

heterostructure device, which can be achieved with an anisotropic electronic structure modification of the CNT

without requiring the chirality information a priori. Therefore, carbon nanotubes could be used as a backbone to

make a heterostructure device. Ultimately, the band gap, the density of states, and the work function will be

engineered by modification of the carbon nanotubes, since these are the key physical parameters to determine the

NASA/CP--2001-210948 764 August 6, 2001
performance of thenanodevice. Asanexample, theformationoftheseamless intramolecular heterostructure will

resultinap-njunctiondiode,and/oratransistor withhightemperature operationcapabilitywithanadvantage ofits

engineered bandgap.AselectivemodificationftheCNTwillbeachievedbytheirradiationoftheionsandthermal

treatments, asshowninthisstudy.Especially, theverticallyalignedmultiwalledcarbonnanotubes havebeenusedas

afieldemissionelectronsource[16].Thefieldemissionpropertylargelydependstheelectronicstructures [17].

WeshowthathedensityofstatesneartheFermilevelincreaseswithincreasingion-bombardment dose,

andthisis relatedto thedisorderin crystalstructure, defectformation, andmixingof thesp2,sp3andsp1

hybridization. Thethermal treatment of theion-bombarded carbonnanotubes inducesreductionof thedensityof

statesneartheFermilevel,andthestructural reconstruction withlargersp3hybridizationandorder.Thisallowsthe

electronicstructuremanipulationneartheFermilevelto getthedesiredmetallicandsemiconducting carbon

nanostructures bycombination ofionbombardment andhermal treatment.

EXPERIMENTS

Thealignedmultiwalledcarbonnanotubes havebeengrownonNi-coatedSisubstrates usingthemicrowave

plasma-enhanced chmical vapordeposition(MPECVD). Thediameter, length,growthrate,thenumberofwalls,

andthedensityoftheCNTscouldbecontrolledbythegrainsizeandmorphology ofthenickelthinfilmsthatare

alteredbytheradiofrequency(RF)powerdensityduringtheRF-magnetron sputteringprocess[18].Theangleresolvedphotoemission studyhasbeenconductedin a synchrotronradiationcenter,Centerfor Advanced

Microstructures andDevices(CAMD)atLouisianaStateUniversity. A monochromatic photonhasbeenobtained

fromthedispersionof 3 m toroidalgrating.Theresultantcombinational resolutionof thephotoemission

spectroscopy is100meV.Thealignedsamplesareannealedforseveral stagesupto600°Cforthephotoemission

measurements toremovealargeamount ofhydrogen, nitrogen, water,andcarbonaceous particles. Thebasepressure

was2 x 10 -10 Torr [11]. The argon ion bombardments are conducted with 2 kV acceleration voltages. The angleresolved x-ray photoemission spectroscopic study has been done with Mg Kc_ - line (1253.6eV) with a 1 x 10 .9 Torr

base pressure. The neon bombardment has been conducted with 5 kV acceleration voltages.

RESULTS AND DISCUSSION

The structure of the thermally treated carbon nanotubes

after Ar-ion bombardment is compared to the pristine carbon

nanotube using scanning electron microscope (SEM) and highresolution transmission electron microscope (HRTEM), as shown

in Figure 1. The pristine carbon nanotubes were bombarded with

1200 btA • min Ar-ion flux at 2 kV. After Ar-ion bombardments,

the sample is annealed up to 500 °C for 5 minutes. As a result of

the heat treatment, the multiwalled carbon nanotube is changed to

a double-walled nanotube, and the relatively ordered carbon

particles are filled in-between the two graphitic layers. This

ordered structural change is not surprising, since a structural

change from the graphite to the diamond induced by the electron

irradiation at high temperature has been probed using HRTEM [ 1].

This indicates that the multiwalled carbon nanotube is transformed

to a mixture of graphite and diamond-like structure by the thermal

treatment after Ar-ion bombardments. The detailed structural

change near the nickel particles was not detected due to the

employed HRTEM resolution limit, even though Banhart group

particles [2].

::::::::::::::::::::::::::::::::::: iiiii_!:_iii:i:i:i:i:i:i:i:';'i :i:?i:?!:i:i:!:i:;!:i:!:i:!:i:?i:?!:i ii:i:::::::::::'.":::i:i:i:i:i:i:!.:,:

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

??!:!:;!:?:;??!:i:i:i.!:i:!:i:!:i:i:7:!:i:i:;

::==========================================

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_" ,.'_::: .,., ,,.. ,,....

,... • • ° -_,%%-..o_ !i:?!:i_i!_!!!:!!!!?!!i:!:i_![_i?:.::?:.::.,N.' _:::::::::::::.'_

Figure 1. a) Scanning electron microscopic

image of pristine carbon nanotube films,

and high resolution scanning tunneling

microscopic images from b) pristine, and c)

argon-ion bombarded carbon nanotubes.

showed the structural change near the nickel

The detailed effect of the ion-bombardment and the annealing effect on the electronic structure are studied using

photoemission spectroscopy, as shown in Fig. 2 and 3. As a function of the Ar-ion bombardment, the valence band

electronic structures are measured with 65 eV photons at normal emission of photoelectrons, as shown in Fig. 2(a).

The unique electronic structures of the pristine carbon nanotubes show the finite density of states near the Fermi

level, the largely overlapped states near 3 eV, and 7.5 eV binding energies [9]. These degenerate states are shown as

sharp peaks in Fig. 2(a).

NASA/CP--2001-210948 765 August 6, 2001
AstheAr-ionbombardment doseincreases, the

valencebandelectronicstructureislargelychanged. Thes-p

bandnear20eVislargelydevelopedattheinitialiondose.

Thehighlydegenerated fatureat7.5eVof thepristine

carbonnanotubes significantlybroadenswithincreasingthe

ion bombardment dueto the structural disorderand

formationof defects.Thepeakat3 eVbindingenergy

becomesunresolveddueto theincrement of the1 eV

featureandbroadeningof features. Especially, thefeature

intensitynear1 eV belowtheFermilevelis largely

increased, whilethereisnosignificant changeofthedensity

ofstatesattheFermilevel.Thenormalizedpeakintensityat

1 eVbindingenergyto theintensityof pristineCNTsis

largelyincreasedattheinitialexposureandthenitisslowly

increasedwithincreasingoftheAr-ionbombardment dose

asshownintheinsetofFig2(a).Thisisanindicationofthe

_-bandformation. Thesechangeshavebeenalsoobserved

in thepolycrystallinegraphite(sp2)andpolycrystalline

diamond(sp3)[19].Especially, thefeatureat1eVbinding

energyhasbeendetectedusingthedoublederivativeofthe

carbonKVVAugerlines,andassignedastheg-satesin

amorphous carbon[20].

Thecarbonls corelevelspectrafromtheCNTs

aretakenasafunctionof theneonbombardment dose,as

L.,_'_--_, , .. _

o 500 1 ooo

..... 7 ..... n ...... _...... _......

20 15 10 _ 0

...............................

4_. ----_ a', \ 7"50

"_ ', ' "_ 460 _o

292 2_31B 28,_ 2_30

_31ndltlq _er_v {eV)

Figure 2. a) Ultraviolet photoemission spectra,

which are taken at normal emission using 65 eV

photons as a function of Ar-ion bombardment

dose. b) Carbon 1 s core level spectra as a

function of neon-ion bombardment dose. The

inset of Figure 2a is the integrated intensity from

the Fermi level to 1.5 eV binding energy as a

function of Ar-ion bombardment.

shown in Fig. 2(b). These carbon is core levels are deconvoluted with three peaks at 284.6 eV, 285.5 eV, and 286.7

eV, respectively [20]. In the case of the pristine carbon nanotubes, the carbon ls core level spectra mainly consisted

of a large peak at 284.6 eV with the sp 2 hybridization character. There still exists the sp 3hybridization peak at 285.5

eV with 20 % of the sp 2 peak intensity. The origin of the feature at 286.7 eV is not clear, even though it has been

assigned as a satellite of the sp 2 hybridization peak [21]. The pristine carbon nanotubes also show a shake-up state

near 6.5 eV away from the carbon ls peak, as marked with an

arrow in Fig. 2(b). This feature originates from a g-plasmon

state, which shows sp 2 character [21]. With increasing neon

bombardment dose, the main carbon ls core level peak is

broadened and the shake-up states disappeared, as shown in

Fig. 2(b). This is related to the formation of the sp 3 and sp 1

hybridization and disorder in crystal structure. At the initial

exposure of the neon-ion bombardments, the peak intensity of

the sp 3 hybridization character is largely increased with

expensing the sp 2 hybridization peak intensity. With further

increasing the neon doses, the peak intensities corresponding

to the sp 2 and sp 3 hybridizations are slightly changed. This

trend is also observed in the broadening of the peaks at 7.5

eV and 3 eV in the valence band spectra, as shown in Fig.

2(a). The feature intensity at 1 eV binding energy is also

slowly increased with increasing of the ion bombardment

dose. These imply that at the initial exposures of ion

bombardment dose, the formation of defects and the crosslinking between carbon layers largely occurs, and then

disorder in the crystal structure is followed with further

exposures.

The valence band spectra of the ion-bombarded CNTs

are taken as a function of annealing temperature, as shown in

Fig. 3. With increasing annealing temperatures, the density of

1400 _-t2oo ; i 3

3 l°°° : "_

800 , 500

_" BOO: 400

400 320

__ 300 "_

280

200 230

0

2'o _'5 _ ; o

-- Binding Energy (eV)

_5

o

-- 0 100 200 300 400 500

Annealing Temperature (°C)

Figure 3. a) Ultraviolet photoemission spectra,

which are taken at normal emission after Ar-ion

bombardment with 1200 NA • min dose as a

function of the sample annealing temperature,

and b) the intensity ratio of the spectra near

Fermi level with respect to the pristine carbon

nanotubes as a function of annealing

temperature.

states near 1 eV binding energy is dramatically reduced as shown in Figure 3(a) while there is no significant change

NASA/CP--2001-210948 766 August 6, 2001
inthehigherbindingenergyfeatures(nobindingenergyshiftandnopeaksharpening). Thenormalizedintensity

ratioofthefeatureneartheFermileveltotheintensityofthepristinecarbonnanotubes dramatically decreases with

increasingoftheannealingtemperature. Above200°C,theratioisalmost linearlydecreased whilethereisahuge

reductionaround300°C.Theseclearlyindicatethati) thedefectsformedin thetubewallsduringtheion

bombardment maybereconstructed intosomeotherformsorthecapofthecarbonnanotubes isreconstructed anii)

moresp3hybridizationthansp2hybridizationis expectedin theamorphous carbonwithincreasingannealing

temperatures. Especiallyafterannealing500°C,thedensityofstatesneartheFermilevelislessthantheoneofthe

pristineasshowninFigure3(b)(ratio= 0.9 at 500 °C). It clearly indicates that the metallicity of the pristine carbon

nanotubes is changed to semiconductor with increasing annealing temperature after ion-bombardment.

CONCLUSION

In summary, the electronic structure of the carbon nanotube is modified with ion bombardment and thermal

treatments. The ion bombardment on carbon nanotubes induces i) the large density of states at 1 eV binding energy,

ii) broadening of the degenerate states, and iii) the formation of the sp 3 hybridization, disorder, and defect in the

crystal structure. Thermal treatment of the ion-bombarded carbon nanotubes reduces the density of states near Fermi

level related to the reconstruction and the increment of the sp 3 hybridization. Thus the application of ionbombardment and thermal treatment to carbon nanotubes gives the desired electronic structure modification

accompanied with structural transformation.

Acknowledgements

Authors thank to the help of H.S. Woo, J. Jiang, P.A. Dowben, and Josef Hormes. This project is conducted

in part by the support from CAMD/LSU and in part by the MOST through NRL program and New Frontier program

in Korea.

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7. J.W. Mintmire and C.T. White: Universal Density of States for Carbon Nanotubes, Phys. Rev. Lett. 81, 12,

Sep. 1998, 2506-2509.

8. C. Dekker, Carbon Nanotubes as Molecular Quantum Wires, Physics Today 22 ,May, 1999, 22-28.

9. S. Suzuki, C. Bower, Y. Watanable, and O.Zhou,: Work functions and valence band states of pristine and

Cs-intercalated single-walled carbon nanotubes. Appl. Phys. Lett. 76, 26, June 2000, 4007-4009.

10. H. Ago, et al: Work Functions and Surface Functional Groups of Multiwall Carbon Nanotubes. J. Phys.

Chem. B103, 1999, 8116-8121.

11. Jaewu Choi, S.M. Lee, Y. C. Choi, Y.H. Lee, and J.C. Jiang, Electronic band dispersion of vertically aligned

multiwall-carbon nanotubes, submitted for publication.

12. S.J. Tans, M.H. et al: Individual single-wall carbon nanotubes as quantum wires. Nature 386, Apr. 1997,

474-477.

13. S.J. Tans, A.R.M. Verschueren, and C. Dekker: Room-temperature transistor based on a single carbon

nanotube, Nature 393, May 1998, 49-52.

14. L. Chico, et al: Pure Carbon Nanoscale Devices: Nanoheterojunctions. Phys. Rev. Lett. 76, 6, 1996, 971-

974.

NASA/CP--2001-210948 767 August 6, 2001
15.J.Kong,H.T.,etal:Synthesisofindividual single-walled carbonnanotubes onpatternedsiliconwafers:

Nature, 395,Oct.1998,878-881.

16.H.Murakami, etal:Fieldemissionfromwell-aligned, patterned, carbonnanotubeemitters. Appl.Phys.Lett.

76,13,Mar.2000,1776-1778.

17.RR.Schwoebel, andI. Brodie:Surface-science aspectsofvacuumicroelectronics. J.Vac.Sci.Technol.

B13,4,Jul/Aug1995,1391-1410.

18.Y.C.Choi,etal:Controllingthediameter, growthrate,anddensityofverticallyalignedcarbonnanotubes

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20.J.C.Lascovish, V.Rosato, andA.Santoni: Thestudyofthemicroscopicscalestructureofcarbon-based

materials: a comparisonbetweenAuger-derived densityof statestructuresandmoleculardynamics

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NASA/CP--2001-210948 768 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

WORKING TOWARD NANOTUBE COMPOSITES

Sivaram Arepalli, Pavel Nikolaev, and Olga Gorelik

GB Tech/Lockheed Martin, Houston, TX

Victor G. Hadjiev

University of Houston, Houston, TX

Carl D. Scott and Bradley S. Files

NASA Johnson Space Center, Houston, TX

One of the most attractive applications of single-wall carbon nanotubes (SWNT) is found in the area of

structural materials. Nanotubes have a unique combination of high strength, modulus and elongation to

failure, and therefore have potential to significantly enhance the mechanical properties of today's

composites. This is especially attractive for the aerospace industry looking for any chance to save weight.

This is why NASA has chosen to tackle this difficult application of single wall carbon nanotubes. Nanotube

properties differ significantly from that of conventional carbon fibers, and a whole new set of problems,

including adhesion and dispersion in the adhesive polymer matrix, must be resolved in order to engineer

superior composite materials. From recent work on a variety of applications it is obvious that the wide

range of research in nanotubes will lead to advances in physics, chemistry, and engineering. However, the

possibility of ultralightweight structures is what causes dreamers to really get excited.

One of the important issues in composite engineering is aspect ratio of the fibers, since it affects load

transfer in composites. Nanotube length was a gray area for years, since they are formed in bundles,

making it impossible to monitor individual nanotube length. Even though bundles are observed to be tens

and hundreds of microns long, they can be built of relatively short tubes weakly bound by Van der Waals

forces. Nanotube length can be affected by subsequent purification and ultrasound processing, which has

been necessary in order to disperse nanotubes and introduce them into a polymer matrix. Some calculations

show that nanotubes with -105 aspect ratio may be necessary to achieve good load transfer (ref. 1). We

show here that nanotubes produced in our laser system are as much as tens of microns long and get cut into

lengths of hundreds of nanometers during ultrasound processing (ref. 2).

Nanotube length was measured by AFM on pristine nanotube specimens as well, as after sonication. In

each case great care was taken to measure individual nanotubes, rather than bundles. Pristine nanotubes

were collected on quartz substrates placed directly in the laser oven and exposed for 0.5 s. This results in an

equal mix of bundles and individual nanotubes (Figure 1). Nanotube length measurements were limited

practically by the lateral span of the AFM scanner, but nanotube length is certainly in excess of 20 gm.

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

:::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::: 96.64 1._ 5

l,- J \ ]

ii_i!!!_i!!ii_!_i_!_i_!iiiiIii_ii,iiiiiii!_:_.;.i_:!!!_iiiii_:i:!!i_iiiiii_i!_!:!: 71.1] --- C Im 0 21

\,,,,, /

C Im 0,16

FIGURE 1. AFM image and corresponding height profiles of one of the substrates exposed for 0.5

seconds, without inner tube and argon flowing at 100 sccm. a) Individual tube, 1.25 nm diameter and

>22 pm long. b) Individual tube, 0.86 nm diameter, >18 pm long. c) Short tube, 0.8 nm diameter, 0.38

pm long. d) Tapered bundle.

NASA/CP--2001-210948 769 August 6, 2001
-.-.%-.-.-o-I%_,-a-.-,-o-,-0-.-,-o-.-.-,-,-.-.-.-, -,-,-,-,-.-, ,-.-,-. -.-.-_v-e'-i-o-

==================================================

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

_?!_'_":!:_:i:i:i:!:!:i:i:!:!:_:?i:i:i:i:!'i'!:.'i:i:i:i_

:: :;iii i'.iii iiiiiiiiiiiii!iiii!i

i:i:i'i:i_iii_iiil !:i:i:i:??!:!:i:_:i:??!'????'?_'_'_'??,_*:i

60

50

40

Z 30

20

10

0

JSC laser tubes, with inner tube

, ,._ I mean = 0 3856 um

!!!

1 2 3 4

I, um

FIGURE 2. AFM image and corresponding length distribution of laser material, produced at Johnson

Space Center, dispersed in SDS / PVP water solution after 10 hours of sonication. Nanotube mean

length is -380 nm.

Location and size distribution allowed us to conclude that nanotubes form very close to the target and

continue to grow and bundle up in collisions as they travel away from it.

In order to prove that sonication shortens tubes, several specimens from various sources were

dispersed by 10 hours of sonication in water solution of SDS (sodium dodecyl sulfate) followed by

incubating with PVP (polyvinyl pyrrolidone). This process separates most bundles into individual

nanotubes, resulting in an even mixture of nanotubes and bundles (ref. 3). Nanotubes were subsequently

deposited on silicon substrates for AFM imaging. Lengths and diameters were measured for over 300 tubes.

Nanotubes vary in length from 50 to 2000 nm, which confirms that sonication cuts them (Figure 2).

Nanotube chemistry is also important for composites, especially in an attempt to separate bundles into

individual tubes while optimizing the interaction between tubes and a matrix polymer. Recent work

includes in-situ studies of the load transfer in composites using Raman spectroscopy (ref. 4). Nanotubes

exhibit a number of active Raman modes, some of which were shown previously to be strain-sensitive in

hydrostatic compression tests (ref. 5). For our experiments, nanotube bundles were embedded in epoxy

resin matrix. We measured the strain4nduced frequency shift of tangential Raman active modes of SWNTs.

Epoxy curing and sample extension in the tensile strength test were found to create residual strains on the

SWNT bundles. We demonstrate that specimen compression in combination with the Raman microprobe

technique provide a means for determination of these strains and hence load transfer effectiveness.

References:

1. Yakobson, B. I., Samsonidze, G. and Samsonidze, G. G., Carbon 38, 1675-1680 (2000).

2. Arepalli, S., Nikolaev, P., Holmes, W. and Files, B., Applied Physics Letters 78 (12), 1739-1741

(2000)

3. Michael J. O'Connell, Peter Boul, Lars M. Ericson, Chad Huffman, Yuhuang Wang, Erik Haroz,

Kevin D. Ausman, and Richard E. Smalley, to appear in Chemical Physics Letters;

4. V.G. Hadjiev, S. Arepalli, P. Nikolaev, B. S. Files and C. D. Scott, to appear in Applied Physics

Letters

5. U.D. Venkateswaram, A. M. Rao, E. Richter, M. Menon, A. Rinzler, R. E. Smalley, and P. C.

Eklund, Physical Review B, 59 10928 (1999).

Keywords: nanotubes, composites, materials, aerospace

NASA/CP--2001-210948 770 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Composite Materials from Modified Carbon Nanotubes

Rodney Andrews, Terry Rantell

Center for Advanced Energy Research and the Advanced Carbon Materials Center, University of Kentucky,

Lexington, KY 40511

Robert C. I-Iaddon

Department of Chemistry, University of California, Riverside, CA 92521

Elizabeth C. Dickey

University of Kentucky Department of Chemical and Materials Engineering, Lexington, KY 40506-0046

Robert G. Bergosh, Hui Hu, Chad Landis, and Mark S. Meier*

University of Kentucky Department of Chemistry and the Advanced Carbon Materials Center, Lexington, KY

40506-0055

ABSTRACT

We have been involved in a program directed toward development of composite materials that manifest

the extremely high tensile strength of carbon nanotubes. In order for such materials to be prepared, it is necessary

to provide improved adhesion between the nanotube and the polymer matrix. Sidewalls of (idealized) nanotubes

are essentially graphite, and therefore derivatization will require extremely reactive reagents.

We have demonstrated that dichlorocarbene and benzyne can attack the "graphitic" equatorial region of

the C7o fullerene. These reagents are therefore prime candidates for attack on the similar sidewalls of nanotubes.

We have treated a variety of nanotubes with these reagents and have characterized the resulting

derivatives. Solubilized single wall carbon nanotubes were treated with dichlorocarbene and multiwall nanotubes

with carbenes and benzynes. Spectra data from these samples has been compared with data from the

corresponding fullerene derivatives, and these results the conclusion that we have modified the nanotubes.

Composites have been made using these modified nanotubes and a variety of polymers. Results on

polystyrene/nanotube composites have shown that processibility significantly increases with modification of the

nanotubes with benzyne. Composites formed with polystyrene and benzyne-modified MWNT's are significantly

more flexible and less brittle than polystyrene or polystyrene composites made with raw, unmodified MWNT's.

The modulus increases as well.

Corresponding Author:

Mark S. Meier

Department of Chemistry

University of Kentucky

Lexington, KY 40506-0055 USA

Tel. (859) 257-3837

Fax (859) 323-1069

meier@pop.uky.edu

Keywords: Nanotubes, composites, nanocomposites, functionalization, fullerenes.

NASA/CP--2001-210948 771 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 3001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Mechanics of Carbon Nanotubes

Prof. Rod Ruoff

Dept of Mechanical Engineering

2145 Sheridan Road

Northwestern University

Evanston, IL 60208-3111

phone 847 467 6596

fax 847 491 3915

r-ruoff@northwestern.edu

We have built and used new tools, as well as employed

conventional methods like tapping mode AFM in new ways, to

measure

mechanical properties of carbon nanotubes, including: (a)

tensile response of MWCNTs and SWCNT ropes (b) response to

tensile load of nested, neighboring shells (shell sliding or

"nanobearing"), thus mapping out interesting tribology (c)

collapsed or partially collapsed MWCNTs on surfaces and on

TEM grids as examined by TEM. In this talk, I present the

new tools and methods, and results achieved.

NASA/CP--2001-210948 772 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FRACTURE STRENGTH OF CARBON NANOTUBES

IN HYDROGEN ENVIRONMENT

Long Guang Zhou

Institute for Metal Research, Chinese Academy of Sciences, Shenyang, 110006, China

San-Qiang Shi

Department of Mechanical Engineering, Hong Kong Polytechnic University

Hung Hom, Kowloon, Hong Kong, China

Mmsqshi (o2polvu.edu.hk, phone (852) 2766-7821, fax (852) 2365-4703

ABSTRACT

Molecular dynamics simulations have been performed to study the mechanical properties of two types of singlewalled carbon nanotubes under tensile loading with and without hydrogen storage. Advanced bond order potentials

were used in the simulations. (10,10) armchair and (17,0) zigzag carbon nanotubes have been studied. Three

deformation stages of armchair carbon nanotube are found. In the first stage, the elongation of nanotube was

primarily due to the altering of angles between two neighbor carbon bonds. The Young's Modulus observed in this

stage was comparable with experiments. In the second stage, the lengths of carbon bonds are extended, while stage

3 was corresponding to the break up of the carbon-carbon bonds. The tensile strength in this stage was higher than

that observed in the first stage. Similar result was also found for the zigzag carbon nanotube but at lower level.

Hydrogen molecules stored inside or outside of nanotubes reduced the fracture strength of both types of carbon

nanotubes. The competition in formation of the hydrogen-carbon and carbon-carbon bonds was found to be

responsible for the detrimental effect. During the deformation, some carbon-carbon bonds were broken and

reconstructed. If hydrogen molecules were around, H atoms would compete with the carbon atoms, to form the

hydrogen-carbon bonds.

Keywords: carbon nanotubes, hydrogen, fracture, computer simulation

INTRODUCTION

Due to the nano-size of single-walled carbon nanotube (SWNT), a direct measurement of its mechanical

properties is rather difficult. However, this extreme size is very suitable for performing atomistic simulations.

Numerous theoretical studies have been carried out in recent years. Yakobson et al. performed a molecular dynamic

(MD) simulation to study the high strain rate fracture in nanotubes (ref. 1). They proposed that nanotubes have an

extremely large breaking strain. More recently, tight-binding electronic calculations on the mechanical properties of

SWNT were reported by Ozaki, et al.(ref.2). Their results revealed that under large strain (-30%), a zigzag nanotube

and an armchair nanotube are the stiffest under elongation and compression regimes, respectively. Currently, the

atomistic simulations on hydrogen in SWNTs are mainly focused on two issues: how much hydrogen can be stored

and where the atoms are stored. It is known that in the metal-based materials, the absorption of hydrogen often

results in embrittlement. A question is, since carbon nanotubes are both promising for use in structure and hydrogen

storage, what is the effect of hydrogen on mechanical properties of carbon nanotubes? This report is an attempt in

answering this question.

METHODOLOGY

We have adopted Brenner's hydrocarbon potential, a bond order potential, in our simulation (refs.3 to 5).

Nanotubes are (10,10)xl00u armchair type and (17,0) X58u zigzag type, respectively. The subscript u denotes a

NASA/CP--2001-210948 773 August 6, 2001
repeatunitofSWNTalongtheaxialdirection.Theseunitswerechosensothathetwotypesofnanotubes havea

similardiameter andtubelength.A simulationcellisshowninFigure1 schematically. A periodicboundary

conditionhasbeenusedalongtheaxialandhorizontal directions. Tosimulatehetubesundertensileloading,we

followedtwosteps.First,thetubeswereannealedatsimulationtemperature for5,000MDsteps.Thetimeinterval

betweentwoMDstepsis0.5fs.Then,thetubewaspulledinaxialdirection(i.e.,z-direction) withastrainof 5×10 -4.

Following each step of pulling, some additional MD steps were used to relax the structure. The simulations were

performed at 300 K and 600 K, respectively.

RESULTS AND DISCUSSION

Unlike the continuum shell model, SWNT is constructed by hexangular carbon rings so that its mechanical prop

erties are strongly depended on its chiral directions. Bond angle and bond length are the two crucial factors that cont

rol the deformation. For armchair SWNT, the elongation of tube due to the altering of bond angles can be up to 15

% if the C-C bonds are assumed to be rigid. While for zigzag SWNT, one-third of the C-C bonds are parallel to the

axis. The loading force is then directly acting on these bonds so that they are easy to break. Figure 2 shows the tensil

e force (Fz) along the axial direction of the armchair (10,10) and zigzag (17,0) SWNTs as a function of strain (Fz-e c

urve) without hydrogen. The tensile force Fz is deduced from an early work (ref.6) by changing the volume to tube 1

ength. Young's modulus of the tubes can be directly evaluated from the figure from liner region in Stage 1. If we co

nsider the wall thickness of tube to be 3.44 A (ref.7), which is the interlayer separation between stacked graphene la

yers, the Young's modulus is then about 753GPa. This is comparable with the recently experimental result (ref.5).

A detailed examination of nanotube structures during deformation revealed that the elongation of the (10,10)

nanotube is initially due to the altering of bond angles (Stage 1). Under further pulling, the contribution from the

elongation of the C-C bonds becomes significant and plays the main role (Stage 2). When the strain is up to a critical

level, some groups of the C-C bonds are broken. Then, the tube starts necking and the force Fz decreases

dramatically (Stage 3).

Compared with the (10,10) SWNT, the (17,0) zigzag SWNT has significantly smaller maximum strain and

maximum tensile force. Here, we define the maximum strain (MS) and the maximum tensile force Fz (MTF) at the

turning point on Fz-strain curve that has the highest value of Fz. Due to the nature of the hexagonal carbon ring,

pulling the zigzag tube along its axial direction would cause some second nearest neighbor C-C atoms to become

closer and to form new C-C bonds. In the local regions of these newly formed bonds, some old carbon-carbon

bonds have to break due to the saturation of covalent bonds. This would lead to the necking and breakage of the

zigzag SWNT.

A range of 4-10wt% of hydrogen stored in SWNT were reported (ref.8). In the following simulations, the

armchair (10,10) and zigzag (17,0) SWNTs were used again for comparison. In our simulation, these SWNTs were

pre-stored with 4.17 or 8.34 wt% of hydrogen gas (H2), both are in the reported range. The absorption of hydrogen

in carbon nanotubes has been theoretically studied (refs. 9 and 10). In the current simulation, however, the details of

the absorption procedure are not of primary concern. We place H2 molecules into tubes directly. The initial

positions and the orientations of H2 molecules in the tubes are chosen randomly. To avoid the overlap of atom

positions, atomistic relaxation was performed. When this is done, 5000 MD steps were used to anneal the structures

of carbon and hydrogen atoms at simulation temperatures and the tensile loading were then carried out. The results

show that the maximum tensile force and the maximum strain both decreased due to the storage of H2, see Figure

3(a) and (b). The effect of H2 on the zigzag (17,0) SWNT seems not as significant as on the armchair (10,10) one.

The effect of stored molecular hydrogen on mechanical properties of SWNT was found to strongly depend on

temperature. It can be seen from Figure 3(c) and (d) that the reductions of MS and MTF caused by hydrogen

storage at 600K are much larger than that at 300K.

The interstitial channels between adjacent nanotubes in a rope of SWNTs may also be the possible sites to

accommodate H2, although there are still some debates at present on this issue. Motivated by this possibility, we

also studied the case in which H2 is stored in outside surface of the tubes, i.e., between the tubes. The simulation

procedure is as follows. A (10,10)xl00u SWNT was placed in the center of a box of 3nmx3nm in length and width.

The height of the box was set equal to the length of the SWNT. H2 molecules were stored in the box, while the

space inside the tube was kept empty. In order to compare to those tubes with H2 stored inside, the molar fractions

NASA/CP--2001-210948 774 August 6, 2001
ofH2werechosentosatisfythathepressureofthehydrogengasoutsidethetubesisequaltothatofthosetubes

withH2storedinside(4.17and8.34wt%).PeriodicboundaryconditionswereusedintheX,YandZdirectionsof

thebox.Thesimulationtemperature was300K.Comparedtothehydrogen-free SWNT,thereductionsinthe

maximumforceare37.7eV/Aand82.7eV/Afor4.17and8.34wt%hydrogenstorage, respectively. Theincreasein

reductionofMTFmightbeduetotheincreaseineffectivecontactareabetweenH2andthenanotubesurface,

becausetheoutersurfaceareaofthenanotubeislargerthantheinnersurfacearea.

Theremaybeacoupleofreasonsthathadcausedthereductionofstrengthofnanotubes byhydrogenstorage.

OneofreasonsmightbethehighpressureofH2actingonthewallofthetubes.Takethe(10,10)armchair tubewith

4.17wt%storedhydrogenasanexample.TheatomicratioofH/Cis1:2.If weassumethatallhydrogenatoms

storedinsidethetubeareinthegasphase, thepressureofgaswouldbe95.2MPaat300K.Suchighpressureof

hydrogengasactingonthewalloftubeswouldchangetheloadingmodeofthetube.In addition, underthehighstraintensileloading, somebondsinthecarbonringswerebrokenandthiscreatedlocaldefectregionsonthewall.

Thepressureofhydrogengaswoulddrivemolecular hydrogenpassingthroughthesedefectregionsandcausethe

regionstoextendintoholes.Anothermoreimportant reasoncomesfromthecompetitioninformationsof the

hydrogen-carbon andthecarbon-carbonds.Oursimulationsshowthatthetotalnumberof theC-Hbonds

increasedsharplyattheStage3ofthetensiledeformationwhentheC-Cbondstarttobreakup.SincesomeC-C

bondswereelongatedandbrokenatthestage3ofthedeformation, hydrogenatomswerelikelyto"catch"someof

thefreecarbonbondstogeneratetheC-Hbonds(seeFigure4),sothathefractureofSWNTwasthenaccelerated.

CONCLUSIONS

MDsimulations basedonanewbondorderpotential havebeenperformedtostudythemechanical propertiesof

SWNTundertensileloadingwithandwithouthydrogenstorage. Theresultsshowthat:

(1)ThetensiledeformationofSWNTsuptothepointofneckingexperiences twostages, controlledbyalteringof

theC-CbondangleandtheC-Cbondlength,respectively;

(2)HydrogenstorageinSWNTreducesthemaximumtensilestrengthofthetubeaswellasthemaximumtensile

strain;and

(3)ThecompetitionbetweentheformationsoftheH-CandtheC-Cbondsmayberesponsible forthereductionin

themechanical strengthoftheSWNTwithH2storage.

ACKNOWLEDGEMENT

TheMDprogramusedincurrentsimulations wasmodifiedfromthatofBrenner's. Theauthorsaregrateful to

Prof.DonaldW.Brennerforallowingustosharehisprogram.Thisworkwassupportedbyacentral researchgrant

fromtheHongKongPolytechnic University(A-PC04).

References

1. Yakobson, B.K.;Campbell, M.P.;Brabec, C.J.;Bernholc, J.,Comp. Mater.Sci.,8(1997)341.

2. Ozaki,T.;Iwasa, Y.;Mitani,T.,Phys.Rev.Lett.,84(2000)1712.

3. Mowrey,R.C.;Brenner, D.W.;Dunlap, B.I.;Mintmire, J.W.;White,C.T.J.,Phys.Chem. 95(1991)7138.

4. Robertson, D.H.;Brenner, D.W.;White,C.T.J.,Phys.Chem., 96(1992)6133.

5. Dunlap,B.I.;Brenner, D.W.;Schriver, G.W.J., Phys.Chem., 98(1994)1756.

6. Haile,J.M.,Molecular dynamicssimulation: elementary methods, AWilry-Interscience Publication, 1992.

7. Saito,R.;Dresselhaus, G.;Dresselhaus, M.S.,Physical Properties ofCarbonNanotubes, Impreial CollegePress,

London, 1998.

8. Simonyan, V.V.;Diep,P.;Johnson, J.K.,J.Chem.Phys.,111(1999)9778.

9. Dresselhuas, M.S.;Williams,K.A.;Eklund, P.C.,MRSBulletin,November (1999)45.

10.Darkrim,F.;Levesque, D.J.Chem.Phys.,109(1998)4981.

NASA/CP--2001-210948 775 August6,2001
I

, .......... J--- l

S

_S _ j

S S

i"41-................................ ID

lOnm

z

Y

Figure 1. Schematic plot of simulation cell.

E

O

co

O_

cv

Lj_ _

250

2OO

15O

lO0

5O

0

0.0

-/i oorm

_. ,_zz'" . ,'", . ---- (17,0) zigzag SWNT....

i I i I i I i I i

0.2 0.4 0.6 0.8

Strain

1.0

Figure 2. Tensile force along axes of the armchair (10,10) and zigzag (17,0) SWNT as a function of strain. The

numbers 1, 2 and 3 denote the three deformation stages of armchair SWNT under tensile loading.

NASA/CP--2001-210948 776 August 6, 2001
23.4

23.2

23.0

co

22.8

22.6

0 2 4 6 8

, , i , i

o_....... (a) • Ms

__

I , I , I , I , I

0 2 4 6 8

wt%H 2

10

112

111

c

11o _<

109

' 108

10

40.0

39.5

_ 39.0

38.5

38.0

37.5

0 2 4 6 8

, , , , ,

o%.. (b) --,-- Ms

"_. --o--- MTF

%

i , i , i , i , i

0 2 4 6 8

wt%H 2

10

22O

215 £E

210 _

205 __

200

195

,

10

40

30

(c) _ 330 K

_600K

HydrcgenFree 4.12wt%_ Storage

180

"_ 160

E

_ 140

(d) _ 300 K

_600K

Hydrogen Free 4.1k:_/o H2Storage

Figure 3. The effect of hydrogen on the strength of (a) armchair and (b) zigzag nanotube8 at 300 K. The effect of

temperature and hydrogen on (c) maximum tensile strain (MS) and (d) maximum tensile forces (MTF) of an

armchair nanotube.

NASA/CP--2001-210948 777 August 6, 2001
Figure4.Snapshotsfthe(10,10)SWNTwith4.17wt%hydrogenstorage: (a)acrossectionat4.7ps(stage2);(b)

thesamecrossectionat7.6ps(stage3);and(c)aviewatthepointoffracture.Smallwhiteballsarehydrogen

moleculesoratoms.

NASA/CP--2001-210948 778 August6,2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

FRACTURE STRESS ENHANCEMENT BY

EB TREATMENT OF CARBON FIBER

Atsushi KIMURA, Akihiro MIZUTANI, Takashi TORIYAMA

Kazuya OGURI, Akira TONEGAWA and Yoshitake NISHI

Department of Materials Science, Tokai University,

1117, Kitakaname, Hiratsuka, Kanagawa, 259-1292 JAPAN

TEL: +81-463-58-1211(Ext. 4206), FAX: +81-463-50-2096

e-mail: '_a_m_ _026_@9_@]__y'_2__:_ ._u_-_t_o_k_@'_a__i£

ABSTRACT

A recovery treatment of fracture origins has been developed for carbon fiber by

using sheet electron beam irradiation. EB irradiation enhanced Weibull modulus and

also increased the tensile strength of fracture. Namely, it enhanced design tensile strength

for aircraft materials.

Key Words: tensile strength, fracture, carbon fiber, EB, Weibull modulus

I.INTRODUCTION

Carbon fiber composites have been applied as high-strengthened light

structural materials in the fields of aerospace and rapid transit engineering. 1'2

Furthermore, in first step in producing, it is often difficult to handle the fractured

fibers in the bundling machine, because carbon fiber often fractured. Since carbon

atoms diffuse at high temperatures, heat treatment is typically used to enhance

the fracture stress of carbon fiber [Ref.3]. Such carbon atom migration probably

reconnected the separated crystal planes near the crack tips and might occupy

free volume [Ref.4] in glassy structure. Thus, the crack tips became dull and the

inter-atomic bonding density became high in carbon fiber. However the heat

treatment takes long time for heat treatment. In order to modify the fracture

strain for short time at lower temperature, carbon fiber has been often treated

with electron beam (EB) irradiation. 5 Electron beam (EB) irradiation of high

electrical potential (170 kV) is used to migrate carbon atoms in a carbon fiber.

Namely, the EB irradiation not only enhanced the fracture strain, but also may

enlarge the tensile strength on fracture. Therefore, the purpose of the present

work is to study effects of EB treatment on tensile strength of fracture of carbon

fiber.

NASA/CP--2001-210948 779 August 6, 2001
II.EXPERIMENTAL

The carbon fiber (6 x 10 -6 m in diameter, Filament 12000f ASAHI-NIPPON CARBON

FIBER Co. Ltd. TOKYO JAPAN) was cut into pieces 50 mm in length. The sheet electron beam

irradiation was homogeneously performed using an electron-curtain processor (Type

CB 175/15/180L, Energy Science Inc., Woburn, MA) [5-8 ].

The acceleration potential and the irradiating current densities were 170 kV and 0.89 x

10-2mA/cm 2, respectively. The EB treatment was applied intermittently (i.e., not continuously).

The conveyer speed was 10 m/rain. Irradiation time was kept constant at 0.23 s in order to

control the temperature in each of the four samples. The temperature of the sample was below

323 K just after the EB irradiation. The irradiation dose was controlled by the integrated

irradiation time in each of the samples. Here, the total amount of absorbed dose value was

converted by the absorbed dose of the distillation water. Although the EB generation was in

vacuum, the irradiated specimen was kept under protective nitrogen at atmospheric pressure in

the apparatus. Namely, specimen was irradiated by electron beam through the titanium thin film

window attached to vacuum chamber (240 mm in diameter). The distance between sample and

window is 35 ram. The oxygen concentration was less than 400 PPM in this atmosphere.

To evaluate tensile strength on fracture, a tensile test was performed. The stress rate was 0.3

N/s. Since d_/de is constant without small strain, the strain rate is approximately constant.

RESULTS AND DISCUSSION

1.Fracture ratio

To evaluate fracture resistance, the

fracture ratio (RO was defined and

expressed by a following equation.

Rf = Nf / NO (1)

Here Nf and Nowere numbers of

fractured and total samples examined

Multi Electro_ Vacuum

Ti Foil window(_ y Chamber

., Lo.e

_rIn N2 Ras _'_gnmp_ ,_

_0) Convcvc r "(0)'_

Figure 1. Schematic drawing

of the apparatus.

for a given fracture stress (6f), respectively. The fracture ratio (Rf) increased with

increasing fracture stress (cyf) (see Figure 2). Midpoint (Rf--0.5) of fracture stress was

4300 MPa for carbon fiber before EB treatment.

To enhance the fracture stress, a recovery treatment of fracture origins has been developed for

carbon fiber by using sheet electron beam irradiation. Its high electrical potential (170 kV) was

probably used to dull the edge of sharp crack tips and to relax the stress concentration. Figm, _ 2

NASA/CP--2001-210948 780 August 6, 2001
showschangeintheRfagainst fracturestressonthetensiletestofthecarbonfibersampleswith

andwithoutEBirradiation.

TheEBirradiationenhanced thefracturestress(_rf)atdifferentRfvalues. Namely, itenlarges

thefracturetoughness forcarbonfiber. WhentheEBtreatment of42.3Mradisperformed,

thetensilestrengthofcarbonfiberwasimproved. Namely, EBirradiationincreasedthetensile

strengthof fracture.TheEBirradiationof 17.3Mradenhancedthemidpoint(Rf--0.5)of

fracturestressupto4800MPa.Midpoint(Rf--0.5)offracturestresswas5500MPaforthe

sampletreatedbyEBirradiationof42.3Mrad.TheEBtreatment of42.3Mradenhances the

fracturestress, whichisabout2000MPalargerthanthatbeforeEBtreatment.

2.Fractureprobabilit wWeibullcoefficient anddesignstressforairplane

The fracture probability(P)dependson risk of rupture (_f/_o)and is

expressedbya followingequation[9].

P= 1- exp [-{_f/{_o In] (2)

(Jfis experimentally obtained fracture stress. When P is equal to 1.0, an ideal fracture stress

((3o) can be obtained. By using Median Rank method broadly applied (Ref.10), integrated

fracture probability (P) is expressed by a following equation.

P = (I- 0.3)/(n + 0.4) (3)

Here, n and I are total sample number and fractured order of each sample, respectively. A

Weibull coefficient (m) indicated statistic distribution of fracture stress. Linear regression curve

as shown in Ggi,_ce _ can be obtained of carbon fiber samples before and after EB irradiation.

Fracture stress of the carbon fiber samples were distributed from 2000 MPa to 6000 MPa. Thus,

Weibull modulus of the carbon fiber was about 3.2. The irradiation enhanced the Weibull

modulus, as shown in l_igt_re 5. The EB irradiation of 17.3 Mrad largely enhanced the modulus.

Since the EB irradiation of high electrical potential (170 kV) probably migrate carbon atoms in

a carbon fiber, such carbon atom migration probably reconnected the separated crystal planes

near the crack tips [5]. Therefore, we concluded that the EB irradiation decreased the statistics

distribution of fracture toughness.

In order to design the aircraft, a design tensile strength at 105 of integrated fracture

probability (P) is important factor. When integrated fracture probability (P) was 105 and 103 for

carbon fiber before EB treatment, the design stresses were 1000 MPa and 1900 MPa (see Fig.5)

respectively. The EB irradiation of 17.3 Mrad enhanced the design tensile strength (2000 MPa

at P=105, 3000 MPa at P=103). The design tensile strength were 3000 MPa at 105 of P and 4000

MPa at 103 of P for samples treated by EB irradiation of 42.3 Mrad. The EB treatment of 42.3

Mrad enhances the design tensile strength, which is approximately 2000 MPa larger than that

before EB treatment.

NASA/CP--2001-210948 781 August 6, 2001
1.0

0.8

('_ 0.6

_a_ 0.4

¢.1

_'_ 0.2

©

.=_

0.0 2C

-I-'1- o Mr_d I L-_ k,/

-L_-17.28_-_dI J_ 0 /

-0- 43'2Mrad _ -_ 6

I i DA o

] A" 0

 o'oo

fracture stress (Mpa)

8000

1.5 • • , • • •

1.0 l/

_-_ 0.5

_0.0 ]I_ D

_-o5 ,,4_F R

,,_-1.5 r.1/AT (/

-2.0 / /

-2.5 V./.s I

"3"_.6 718 8.0 812 814 816 818 9.0

lll_Sf (_Sf :MPa)

Figure 2. Change in the Rf against

fracture stress on the tensile test of the

carbon fiber samples with and without

EB irradiation.

Figure3. Change in probability of

fractured samples against fracture

stress on the tensile test.

_6

0

EB irradiation dose (Mrad)

4500

"i 2500

2 200 A_/j,k / ¢

1500

5

100

5 1015202530354045 0 5 1015202530354045

EB irradiation dose (Mrad)

Figure4. Change in Weibull modulus

auainst EB irradiation dose.

Figure5. Change in design tensile strength

a_ainst EB irradiation dose.

NASA/CP--2001-210948 782 August 6, 2001
3.Discussion for fracture stress enhancement by EB treatment of carbon fiber.

Figure 6 shows X-ray diffraction patterns before and after EB irradiation. The amorphous

structure can be confirmed before and after EB irradiation.

The large structure change cannot be observed. However large shape change of stress-strain

(o-o) curves for carbon fiber samples before and after EB irradiation had been observed, as

shown in Figure 7. The do/do value was above 1.0 of carbon fiber sample before EB irradiation.

The curve shape was looks like metallic amorphous. On he other hand, the do/do value was

below 1.0 of carbon fiber sample after EB irradiation. The curve shape was looks like metallic

crystal. Amorphous materials were used to involve free volume [4]. If the free volume

decreased in carbon fiber, the density of dangling bonds should decrease.

10 20 30

, QfiE(deg)

2500 _, ,_ ,_ a . _,

• befo_e_a_a_on I g"

• 4.36Mrad

2000 -- _//_

/.. /

 000 f

50

. ,

0.0 0.2 0.4 0.6 0.8 1.0 1.2 1.4

Strain a

Figure 6. X ray diffraction patterns of

carbon fiber samples.

Figure 7. Stress - strain curve

of carbon fiber samples with

and without EB irradiation.

If the EB treatment forcibly diffuses carbon atoms into free volume and incoherent interface

because of high electrical potential (170 kV), Such carbon atom migration probably decreased

the dangling bonds, dulled the edge of sharp crack tips and probably reconnected the separated

crystal planes near the crack tips. In order to evaluate the inter-atomic bonding density, density

of dangling bonds (DDB:mol 1) is obtained by dangling bond signals of graphite, which can be

detected by means of electron spin resonance (ESR) spectrometer [5]. Here the microwave

frequency range used in the ESR analysis was the X-band. The field modulation was 100 kHz.

Spin density was calculated using a Mn 2+ standard sample. The EB irradiation decreased the

density of dangling bond (spin) [5]. Namely, the EB treatment forcibly diffused carbon atoms

to free volume and incoherent interface. Since the inter-atomic bonding density increased in

carbon fiber, the EB treatment improved the tensile strength on fracture.

NASA/CP--2001-210948 783 August 6, 2001
4. EB treatment application for C/C composite

15000

In order to apply carbon-carbon

composite, C/C composite coil spring

L

(50 mm in diameter, 30 mm in height) "_10oo0

was treated by the sheet electron

beam treatment. When the compressive _, 5ooo

test was performed below 1500 N/m 2, "_

the 40Mrad EB treatment enhances the _,

spring elasticity, as shown in Figure 8.

The spring compressive elasticity for

the treated sample is 170% larger than

that for untreated sample.

C/C composite coil spring

--l-- 0Mrad T /

-0-- 40Mradl/[ /

• lO /

0.00 0.06 0.12

compressive strain

Figure 8. Compressive stress -

strain curve of carbon-carbon

composite spring samples with

and without EB irradiation.

CONCLUSION

A recovery treatment of fracture origins has been developed for carbon fiber by using sheet

electron beam irradiation. When the EB treatment is performed, the tensile strength of carbon

fiber was improved. The EB irradiation of 17.3 Mrad largely enhanced Weibull modulus. It

shows that the EB irradiation decreased the statistics distribution of fracture toughness. The EB

treatment of 42.3 Mrad enhances the design tensile strength, which is approximately 2000 MPa

larger than that before EB treatment.

References

[1] J. D. Brooks and G. H. Taylor, Carbon, 3, 185(1965).

[2] T. Chang and A. Okura, Transactions ISIJ,2_7 229(1987)

[3] C. R. Rowe, D. L. Lowe, Bienn. Conf. Carbon, 13,170(1977).

[4] M.H.Cohen and D. Turnbull, J. Chem. Physics., 43139 (1965).

[5] Y. Nishi, T. Toriyama, K. Oguri, A. Tonegawa & K. Takayama, J. Maerials Research

2001, in press.

[6] Y. Nishi, S. Takagi, K. Yasuda & K. Itoh, J. Applied physics, 70,367(1991).

[7] Y. Nishi, H. Izumi, J. Kawano, K. Oguri, Y. Kawaguchi, M. Ogata, A. Tonegawa, K.

Takayama, T. Kawai, M. Ochi, J. Materials Science, 32, 3637.(1997).

[8] K. Oguri, K. Fujita, M. Takahashi, Y. Omori, A. Tonegawa, N. Honda, M. Ochi, K.

Takayama and Y. Nishi., J. Materials Research, 13, 3368(1998).

NASA/CP--2001-210948 784 August 6, 2001
9

GROWTH AND MODELING OF CARBON NANOTUBES

AND NANOSTRUCTURES

Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

COMPUTATIONAL STUDIES ON FORMATION AND PROPERTIES OF CARBON

NANOTUBES

Tahir t_agin, Jianwei Che, Weiqiao Deng, Xin Xu, and William A Goddard, lII

Materials and Process Simulation Center, Beckman Institute, 139-74, California Institute of Technology, Pasadena,

California, USA

ABSTRACT

The discovery of lower dimensional forms of Carbon with unique mechanical and electronic properties has

generated new possibilities in many areas of technology especially in nanotechnology. Recent emergence of some

nanoscale device applications show how this potential is turning into a reality. Over the years, we have employed

various levels of theory to study the structure and properties of carbon based materials for nanoscale applications.

In this paper, we present two new theoretical studies. We present a study on the transition metal catalyzed growth of

single wall carbon nanotubes. In the second study, we investigate the relation between mechanical deformation and

excess charge in order to understand how introducing and controlling the charge at various locations might modify

the mechanical and acoustical properties of carbon nanotubes. We demonstrate that introducing excess charges into

single wall carbon nanotubes can lead mechanical deformations that do mechanical work. The results suggest a wide

range for practical applications, such as NEMS, acoustic sensors and nanoactuators.

INTRODUCTION

The peculiar chemistry of carbon results in diverse forms of structure: the 3-dimensional network of diamond and

the 2- dimensional sheets of graphite have been known through the ages. The discovery of lower dimensional forms

of Carbon with unique mechanical and electronic properties has generated new possibilities in many areas of

technology especially in nanotechnology. Recent emergence of some nanoscale device applications show how this

potential is turning into a reality.

Over the years, we have been using ab initio quantum chemistry, density functional theory and molecular

dynamics methods to study the structure and properties of carbon based materials for nanoscale applications. The

applications such as; a) Structural and mechanical properties of nanotubes, (refs 1-3); b) Tribological properties of

carbon for NEMS applications (ref 4); c) Thermal transport properties of nanotubes and carbon based heterostructures (refs 5-6), and d) Formation of fullerenes, (ref7) have appeared elsewhere.

In this paper, we will present two new theoretical studies. Over the past several years a large number of synthetic

procedures for the production of Carbon nanotubes have been developed. Since the electronic properties of carbon

nanotubes depend on structure the control of growth is essential. Theoretical studies of growth mechanisms may

shed some light on how to control the growth. We first present a study on the transition metal catalyzed growth of

single wall carbon nanotubes. In the second application, we investigate the relation between mechanical deformation

and excess charge in order to understand how introducing and controlling the charge at various locations might

modify the mechanical and acoustical properties of carbon nanotubes. We demonstrate that introducing excess

charges into single wall carbon nanotubes can lead mechanical deformations that do mechanical work. The results

suggest a wide range for practical applications, such as NEMS, acoustic sensors and nanoactuators.

MECHANISM OF TRANSITION METAL CATALYZED GROWTH OF SINGLE-WALL CARBON NANOTUBE

Introduction

Single-wall carbon nanotubes have been produced in the outflow of a carbon arc discharged method (ref 8) and

in even with higher yield by the laser vaporization technique (refs 9-12) and more recently through chemical vapor

deposition (ref 13). The transition metal catalysts such as Ni and Co are known to play important role in all of these

methods. In the laser vaporization assisted growth, the concentration of metal catalyst in the graphite is very low (<

1%). The high temperatures in the experiment suggest that the catalyst may affect the growth as single atoms.

NASA/CP--2001-210948 787 August 6, 2001
In1997,Leeetalproposedagrowthmechanism(ref14).Inthismechanism, themobileNicatalyst aomsare

absorbedatthegrowthedgeofSWCNwheretheyprevent theformationofcarbonpentagons andcatalysis

continuesthroughformationofcarbonhexagons. However, Froudakis' groupprovidedanother mechanism(ref.15):

Niatomsdon'tattachtotheedgeofthegrowthfont.It actuallyfirstcreatesandstabilizesthedefectsinnanotube

andthentheincomingcarbonatomsanneal theNi-stabilizeddefectsbyfreeingtheNiatombacktothecatalytic

cycle.Thesetwomechanisms arecontradicting. Besidesthisapparent contradiction, therestillaresomeunsolved

problemsthatneedtobeaddressed. First,inthegrowth,whatisthedifferenceinthegrowthmechanisms withor

withoutmetal?Second, cantheexistingmechanisms explainwhyametalworks,whileanother metaldoesnotwork?

Third,whatisthedeterminingstepinametalcatalyzedgrowthofcarbonnanotube? Fourth,where,attheedgeoron

thewall,doesthemetalatomlocateduringthegrowth?

Theoretical model

Here,wepresent resultsofadetailedtheoretical studyofthedynamical interactionbetweentheNi,Co,Ptand

CucatalystsandSWCNwithaviewtowardsanunderstanding ofthenanotubegrowthmechanismandadetailed

discussiontheunresolvedproblemabove.Ourcalculations arebasedonfirstprinciple- densityfunctional

theory,asimplemented intheJaguar code,foratubefragment withametalatomattached. TheB3LYPareusedto

describeexchangeandcorrelation. Westartwitha(10,10)nanotube. Thestudiedfragment iscutfroma(10,10)

nanotubestructureandfixingtheedgeatoms(whichareneutralizedbyH atom)to keepthecurvatureofthe

structure.

Resultanddiscussion

Firstquestiontoaddressiswhethertheroleofthemetalcatalyst isanecessary prerequisitefortheformationof

single-wall nanotubes. Inordertounderstand thisissue,themechanismofterminationofnanotubeinabsenceof

metalatomsisundertaken. InFigure1,wecanseethetwopathwayswhilecarbonatomsgrowontheedgeof

nanotube. Onepathwayistoformadefectwithtwohexagons andtwopentagons called6,6,5,5defect, consequently,

aftermoreandmoredefectsareformed, thetipofnanotubeisclosed. Anotherpathwayisthathesetwopentagons

rearrangetoformtwohexagonsandthenthegrowthwillproceed. Theresultsshowthatforpathway1theenergy

willdropdownby81.9kcal/moleandforpathway2theenergywillgodownby51.7kcal/mole. Thelargeenergy

differencebetweenthesetwopathwaysallowsustobelievethatintheabsenceofmetalatomsthegrowthwillbe

terminatedimmediately. ThisisinfavorofFroudakis' mechanism: Themetalmustattachontheedgeofthenanotube

growthfrontandpreventsthedefectformation.

Second, wediscussthedetailedcatalyticmechanismwithcatalyst atomstoavoidthedefectformation. First,we

explorewhether theNiatomaffectsthegrowth.Ourcalculations showthatNiatomneitherdestroysthepentagonto

formhexagonnorstabilizesthehexagonring,becausebothareenergetically unfavorable. Fortheformercase,we

startedthesimulationbyinsertingtheNi atomintoapentagontoformhexagon. Thisprocedureisenergetically

unfavorable, whichrequiredenergyof12.4kcal/mole. ThismeansthatNiatompreferstositabovethepentagonrig

withoutbreakit. Forthelatercase,thecalculations idicatethearrangement tobeenergetically favorablewithan

energygainof27.2kcal/molewithoutNi catalyst andwithanenergygainof24.6kcal/molewithNi catalyst. The

outcomeofthiscalculationshowsthathearrangement isotbenefitedfromNimetalatomsinvolveattheadjacent

site.TheseresultsareinlinewithLee'smechanismbyconcludingthatNi atomsdon'tpreventthepentagon

formationorassisttheassemblyofcarbonhexagons atthegrowingedge.

Basedonthis,weprovideanewmechanismhowninFigure2.It saysthatNi atomsblocktheadjacent sitesof

pentagontoprevent thedefectformation. Niatomcananneal theexisteddefects. Figure2ashowsthatwhencarbon

atomscomein,Ni atomsblockbothadjacent sitesoftheothersitessothatpentagoncarbonringjustisableto

rearrangeashexagonringthatisenergetically favorable, i.e,it reduces24.60kcal/mole. Figure2bindicatehowaNi

atomannealsthedefect.AfterNiatomattachestheedgeofthenanotubegrowthfont,it willrearrangewithcarbon

ringsintoseveral possiblestructures: structure1-5inFigure2b.Forgrowth,thestructure(1)and(3)(4)arenot

helpfulforannealingdefectbecausethepentagons inthestructurearenotdestroyed. Attheotherhand,structure

(2)and(5)aregoodforanneal defects. Wecomparetheenergiesofthesestructuresanddeterminewhichstructure

thereactionpathistrapped. ForNi metalatoms,structure(2)isthemoststableonethatwillleadtoanannealing

pathway. Therefore, ourmechanismuccessfully explainstheNicatalyst effectofnanotubegrowth.

WealsostudythecatalysiseffectofothermetalssuchasCo,PtandCu.Inthetable1,wegiveasummaryof

thesemetals. Basedonourcalculationresults, it showsthatCodoeswork,PtandCudon'twork,whichisingood

agreement withtheexperiments.

NASA/CP--2001-210948 788 August6,2001
Conclusion

Basedonfirstprinciplecalculation, westudythemicroscopic mechanismofsingle-wall carbonnanotubes

growthbyusingthelaservapordepositiontechnology. Byarguingthepreviousmechanism, weprovideanewone

here.Itsaysthatmetalcatalystsatomabsorbedatthegrowthedgewillblocktheadjacent growthsiteofpentagon

andthusavoidtheformationofdefect. Metalcatalystscanalsoanneal theexisteddefects. Additionally, our

resultsshowthathenanotubegrowthwillterminateintheabsenceofmetalcatalystsandalsoexplainwhyPtand

Cuarenotgoodcatalysts.

Table 1 Relative energies (kcal/mole) of various metal cluster structures in figure 2b

(1) bad for annealin_

(2) Rood for annealin_

(3) bad for annealin_

(4) bad for annealin_

(5) good for annealing

Ni Co Pt Cu

-19.65 -1.11 -- -27.34

-24.24 -14.34 22.82 -17.00

-9.55 2.25 41.11 -15.30

0.0 0.0 0.0 0.0

-5.66 9.8 2.13 -7.26

Figure 1. The nanotube grows at the absence of metal catalysts. Pathway (a) is closure pathway; Pathway (b) is

continue growth pathway. The colorful carbon atoms are our cluster model.

AE -24.60 kcal/mole

Figure 2a. The Ni atoms block adj acent site of pentagon to avoid the defect formation. The darker atoms are Ni

atoms.

NASA/CP--2001-210948 789 August 6, 2001
Figure2b. Metal Ni catalysts can anneal the formed defect. The darker atoms are Ni atoms.

CHARGE EFFECTS ON MECHANICAL PROPERTIES OF CARBON NANOSTRUCTURES

Introduction

Over the past decade both theory and experiments have shown that carbon nanotubes have both unique

electronic properties (e.g. it may be semiconducting or metalic depending on the charality) and extreme mechanical

strength (e.g. tensile modulus ~1000 GPa). Consequently, there is considerable interest in designing and

manufacturing functional devices and novel composite materials based on carbon nanotubes. In particular we

consider here how introducing and controlling charges might modify the mechanical properties of carbon structures.

The smaller nanostructures are moderate size molecules. To understand the systematic behavior of carbon

nanotube's electronic and mechanical coupling, we started with the smallest molecule that resembled carbon

nanotube structure, i.e. benzene. While the expansion of benzene molecule approaches the limit of graphene

structure, we believed that the electro-mechanical coupling in graphene is in close resemblance to that in carbon

nanotube. We will see this behavior later in our calculations.

Results and Discussion

Staring from benzene, we investigated naphthelene, pyrene, coronene, intercalated graphite, and single walled

carbon nanotubes. Based on group theory, we know that the two HOMOs and two LUMOs of benzene are

degenerate. In each PI MO, we can characterize the bonding between atoms by looking at the phase of their _p

orbital. When the two adjacent orbitals have opposite phase, an anti-bonding dominates the interaction between

them in that specific MO. On the contrary, the same phase orbital generates bonding forces. Consequently, an

electron in anti-bonding form will elongate the distance between two adjacent atoms, while an electron in bonding

form will shrink the distance. The picture in the presentation clearly shows the structural change with respect to

charging into different MOs. In larger molecules such as pyrene and coronene, we also observed the similar behavior.

A similar system at large spatial scale is intercalated graphite. The quantum calculations for periodic systems were

carried out using CASTEP. When positive charges are injected into the graphite crystal, the graphite sheets tend to

shrink. On the other hand, negative charges make the graphite to expand. The structural deformation is mainly caused

by the change in electronic structure rather than Coulomb interaction. In other words due to the filling of electrons in

anti-bonding/conduction band or holes in bonding/valence band. Coulomb force is independent of charge signs. In

addition, Coulomb interaction depends on charges in the quadratic order. For the intercalated graphite, we also found

that AA stacking structure is more stable than AB stacking in agreement with experimental observations.

Similar to intercalated graphite, carbon nanotubes also exhibit charge induced structural deformations. For single

walled carbon nanotube, we chose (5,5) tube as an example. Although it has very small size, we believe that the

results can demonstrate the essence of electro-mechanical coupling. The simulation cell consisted of two layers of

NASA/CP--2001-210948 790 August 6, 2001
(5,5)unitcell.Eitherelectronsorholesarefilledintocorresponding bands, andauniformbackground chargeisused

toneutralizethetotalsimulationbox.Fromourcalculations, wealsofoundthatpositivechargetndstoshrink

._o

"o

o

OoE

_g

_D3

<

2.49

2.48

2.47

2.46

2.45

2.44

2.43

2.42

2.41

.,3

\,\

"%.,..

i i i i i i

-2 -1 0 1 2 3

Charge Transfer (e)

"E" 14

p

13

< 12

E

L3 9

_'__,e Cell constant a

. ,:, Cell constant b

-4 -3 -2 -1 0 1 2 3 4

4

Charge Transfer (e)

Figure 3. Deformation as a function of charge transfer: Left deformation in the tube direction, Right

deformation in the lateral directions.

the tube and the tube tends to expand under negative charging. The reason is very similar to intercalated graphite.

The charge injected into valence or conduction band caused the electronic structure to shift, and this shift is charge

sign dependent. Intuitively, it can be viewed as new electronic structure under screened nuclear cores. In addition to

the deformation along tube axis, we also saw the change in cell length in nanotube bundles. Although LDA usually

does not give accurate results for nonbond interactions such as van de Waals forces, we think that the change in the

lateral directions is mainly caused by the Coulomb forces. As it can be seen in electrostatic potential map, the inter

tube space is mainly electronegative. When a charge variation is introduced, the inter-tube repulsion due to charge

variations in positive core is on the second order, and the first order interaction is between the charge variation and

the electrostatic potential. Therefore, the bundle size also shrinks or expands due to different type of charges.

Figure 4. Electron density map

In summary, we find that introducing excess charges into nanotubes can lead to mechanical deformations that do

mechanical work in agreement with experiments by Baughman et. al. (ref 16). These results suggest a wide range of

practical implications, including the design ofnano-electronic-mechanical systems (NEMS) and nano-actuators.

NASA/CP--2001-210948 791 August 6, 2001
REFERENCES

1. GGao,T.Cagin,W.A.Goddard, III:Energetics, structure, thermodynamic andmechanical propertiesof

nanotubes, Nanotechnology, vol.9,no.3,(1998)183-191.

2. G.Gao,T.Cagin, W.A.Goddard, III:PositionofKatomsindopedsinglewalledcarbonnanotubecrystals,

Phys.Rev.Lett.,80,(1998) 5556-5559.

3. J.Che,T.Cagin, W.A.Goddard, III:StudiesofFullerenes andCarbonNanotubes byanExtendedBond

OrderPotential, Nanotechnology vl.10,(1999)263-268.

4. T.Cagin, J.Che,M.N.Gardos, A.Fijany,W.A.Goddard, III:SimulationadAnalysisofExperiments on

FrictionandWearofDiamond: Amaterial forMEMSandNEMSapplications, Nanotechnology vol.10,

(1999) 278-284.

5. J.Che,T.Cagin, W.Deng,W.A.Goddard, III:Thermal Conductivity StudiesbyMolecular Dynamics

Simulations," J.Chem. Phys. Vol.113,(2000) 6888-6900.

6. J.Che,T.Cagin, W.A.Goddard, III:Thermal Conductivity ofCarbonNanotubes, Nanotechnology vol.11,

(2000) 65-69.

7. X.Hua,T.Cagin, J.Che,W.A.GoddardIII:QM(DFT)andMDstudiesonformationmechanisms ofC60

fullerenes, Nanotechnology vol.11,(2000)85-88.

8. S.Iijima,T.Ichihashi: sngle-shell carbonnanotubes of1-nmdiameter, Nature, vol.363,no.6430,Jun1993,

603-605.

9. T.Guo,PNikolaev, AThess, etc.:Catalyticgrowthofsingle-walled nanotubes bylaservaporization, Chem.

Phys.Lett., vol. 243, no.l-2, Sep. 1995, 49-54.

10. J.H. Hafner, M.J. Bronikowski, B.R. Azamian, et al.: Catalytic growth of single-wall carbon nanotubes from

metal particles, Chem. Phys. Lett. vol. 296, no. 1-2, Oct. 1998, 195-202.

11. K.B. Shelimov, R.O. Esenaliev, A.G. Rinzler, et al.: Purification of single-wall carbon nanotubes by

ultrasonically assisted filtration. Chem. Phys. Lett. vol. 282, no. 5-6, Jan. 1998,429-434.

12. A.A. Puretzky, D.B. Geohegan, X. Fan, et al.: Dynamics of single-wall carbon nanotube synthesis by laser

vaporization. Appl. Phys. A, vol. 70, Jan. 2000, 153-160.

13. H.J. Dai, J. Kong, C.W. Zhou, etc.: Controlled chemical routes to nanotube architectures, physics and

devices, J. Phys. Chem. B, vol. 103, no. 51, Dec. 1999, 11246-11255.

14. Y.H. Lee, S.G. Kim and D. Tomanek: Catalytic growth of single-wall carbon nanotubes: an ab initio study,

Phys. Rev. Lett., vol.78, no. 12, Max. 1997, 2393-2396.

15. A.N. Andriotis, M. Menon and G. Froudakis: Catalytic action of Ni atoms in the formation of carbon

nanotubes: A molecular dynamics study, Phys. Rev. Lett., vol. 85, no. 15, Oct. 2000, 3193-3196.

16. R.H. Baughman, C.X. Cui, A.A. Zakhidov, et al. Science vol. 284:no.5418 (1999) 1340-1344.

NASA/CP--2001-210948 792 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

NANOMACHINED CARBON: NANOTUBES, FULLERENES, PEAPODS

David Tomfnek

Department of Physics and Aslronomy

Michigan State University

East Lansing, Micldgan 48824-1116

tomanek@pa: msu.edu

h_ap:llww, w pa msu.edtd-tomanek/

517-355-9702(phone) 517-353-4500 (fax)

ABSTRACT

In recent years, following the discovery of the C_0 "buckybalt", a plethora of nanostructures attracted rite

attention o[ carbon scientists: Examples of these intriguing systems are other fullerenes, onions, nanotubes, and

hybrid structures such as peapods. Whereas most of these sp 2bonded structures are known to grow from carbon

vapor under extreme conditions, structures like multi-wall nanocapsules may form by activated bond rearrangement

from ultra-disperse diamond ill or higttly defective graphitic structures.

I will show how macroscopic concepts from elasticity theo_ translate down to the nanometer scale, where they

provide quantitative predictions for the relative stability, of nanotubes, scrotIs, fulIerenes, and peapods. With the help

of ab in,tie calculations, it is now possible to study the formation and inter-conversion mechanism ef other carbon

aanostructures. Molecular dynamics simulations (see _vw___e&i/cmp_sc/mid the figure below) provide

intriguing insights into the formation mechanism of pcapods [2] mtd rite tmexpected structural stabilily of entangled

nanotube "hooks", reminiscent of the velcr_ bond.

[1] Yot,ng-K)_an Kwon, David ToIr&nek, and Sumio lijima, "Bucky-Shuttte" Memot3" Device: Synthetic Approach

and Molecular Dynamics Simulations, Phys. Rev. Lett. 82, t470 (1999).

[2] Savas Berber, Young-Kyun Kwon, and David Tom_iek, Nano-gulf: Putting a C6o imo a Nanotnbe (submitted tbr

publication).

c-.,./:<:.;-:-:i_:..:-:; _;_-: _

y .,..,,¢ .:y,?;_:_

Figure: Visualization of a possible end-on encapsulaaon process qJ'a Ceofid!erer_e m a (10,t0) carbo__ ncmotube

(see also _.t_/_._._>v. .2a_e_._s.u..:_e...d.._/._._.n._p'!c._/s. i..m._p2d _1e#/_.

Keywvrds: nzmetubes, fullerenes, energy, tbrmation

NASA/CP--2001-210948 793 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Growth of Carbon Nanotubes on Patterned Catalysts: Towards a Carbon

Nanotechnology

G A. J. Amaratunga, M ChhowaUa, K. Teo, C. Ducati, I. Alexandrou, N. L.

Rupesinghe, A. Ferrari, W. I. Milne and J. Roberstson, Engineering Dept.,

D. Roy and T. W. Clyne, Dept. of Materials Science,

D. Hasko, Microelectronics Research Centre, Cavendish Laboratory,

Cambridge University, Cambridge CB2 1PZ, UK.

A. J. Papworth and C. J Kiely, Engineering Dept., Materials Science and Engineering,

University of Liverpool, Liverpool L69 3BX, UK.

It is now 10 years since the TEM observations of Ijima which gave rise to the radical

re-examination of graphitic carbon behaviour on the nanometer scale. The scientific

understanding of the expected ( and observed) mechanical and electronic behaviour of

carbon nanombes gained over the last decade points to it being a near ideal system for

nanoscale mechanical and electronic devices. This in turn has inspired may groups

world wide to attempt to develop a technology which will allow the properties of

carbon nanotubes to be exploited.

The highest quality nanotubes available at present are the single wall type produced

through laser ablation. From a technological perspective they are , however, not

optimal as it is necessary to purify them in solution, and then assemble them into

desired structures. The latter has proved to be a non-trivial task.

The alternative approach is to use a chemical vapour deposition (CVD) process to

grow where they are required for the fmal device structure. This process has been

approached on the basis of the controlled and defined placement of the catalyst on a

substrate. However, the quality of CNTs and their orientation have been difficult to

control in the CVD process.

Here we report our latest results which show that the CVD process with plasma

enhancement can be used to grow oriented nanotubes with very high definition and

selectivity over large areas. The methods of catalyst control required to achieve this

are discussed. The quality of the carbon nanotubes can also be controlled to some

extent. In fact we show evidence for the possibility of synthesising single wall carbon

nanotubes by the CVD method. We also show the differences in the catalytic process

under plasma conditions which lead to the alignment of CNTs. Taken together, the

results show that the initial steps required for the development of a technology for

electronic and mechanical devices exploiting the properties of CNTs are now in place.

NASA/CP--2001-210948 794 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Surface decomposition mechanisms on SiC (0001)

M. Kusunoki °, T. Suzuki, T. Hirayama and N. Shibata

Japan Fine Ceramics Center FCT Central Research Department, Atsuta, Nagoya, Japan

E-mail:kusunoki@j fcc.or.jp

ABSTRACT

Decomposed structures of SiC will be shown between the Si(0001)-face and the

C(000-1)-face after heating at 1700°C for a half hour in a vacuum. On the C-fac e an

aligned carbon nanotube film was self-organized perpendicular to the surface. On the

contrary, a layer of very thin graphite sheets parallel to the surface was formed on the

Si-face. Decomposition mechanisms on both faces are proposed from high-resolution

electron microscopy (HREM) results.

Keywords: decomposition, SiC, carbon nanotube, alignment, TEM

INTRODUCTION

Silicon carbide (SIC) is one of the most promising materials for engineering applications because of its

attractive properties. It has wide band gaps, excellent mechanical properties and chemical stability, and

exhibits poly typism (ref.1) Therefore, it is expected to be used for high-power and high-temperature

semiconductor devices. The surface structures or the surface reconstruction have been investigated

intensively by scanning tunneling microscopy (STM) (ref.2 to 4) and low-energy electron diffraction (LEED)

(ref.2 and 5). These experiments have shown that by increasing the heating temperature above 1000°C,

surfaces of SiC are covered with carbon-rich structures such as 6_/3 x 6_/3, 30°(ref.2,3 and 5) or 6 x 6 (ref.4

and 5). This phenomenon has been confirmed by the selective desorption of Si atoms by annealing SiC in a

vacuum.

Hexagonal (x-SiC is a zinc blend type material with a polar [0001] axis, in which the ideal (0001) and

(000-1) surfaces are terminated by a layer of Si and that of C atoms respectively. Van Bommel et al.(ref. 2)

investigated the difference of surface decomposed structures on both Si- and C-faces by heating up to 1500°C.

They examined LEED patterns and Auger electron spectra (AES) obtained from each face, then concluded

that a monocrystalline graphite layer and a polycrystalline graphite layer formed on the Si-face and C-face,

respectively.

Recently, the authors discovered an aligned carbon nanotube film being self-organized by surface

decomposition of SiC by TEM in-situ observations (ref.6). Additionally, it was shown by using HREM and

electron energy-loss spectroscopy (EELS)that the surface decomposition progressed as residual oxygen gas

produced an oxidation reaction (ref.7 and 8).

In this paper, it will be clarified that the decomposed structures on the Si-face and the C-face are different,

and they will be explained by the different oxidation mechanisms on the respective faces from cross-sectional

HREM results.

NASA/CP--2001-210948 795 August 6, 2001
EXPERIMENTAL PROCEDURE

Commercial single crystal wafers of 6H-SiC with the polished Si-face and those with the polished C-face

(CREE Research, Inc.) were cut to the size of 1.0 x 4.0 x 0.2 mm. The wafers with the C-face were heated at

1200, 1250, 1300 and 1700°C for 0.5 h, and those with the Si-face at 1350 and 1700°C for 0.5 h in a vacuum

furnace (lxl0 -4 Torr) with an electric resistance carbon heater. These specimens were carefully thinned by

dimpling and ion bombardment for HREM observation of the cross section using a Topcon 002B TEM

operated at 200 kV.

RESULTS AND DISCUSSION

Figures l(a) and (b) show images and the diffraction patterns obtained from the Si-face and the C-face of

SiC single-crystal wafers heated to 1700°C for 0.5 h. On the C-face, aligned carbon nanotubes 3-5 nm in

Figure 1. Micrographs and the diffraction patterns of the surface of SiC single-crystal

wafers heated to 1700°C for 0.5 h, on (a) the C(0001)-face and (b)the Si(0001)-face.

NASA/CP--2001-210948 796 August 6, 2001
diameter and0.25_minlengthwereformedperpendicular tothesurface, asshowniFig.l(a) (ref.7and

8).Onthecontrary, inthecaseof theSi-face,seventeen graphitesheets, 5 nminthickness, wereformed

parallel totheSi-faceasshowninFig.l(b).The0002latticedistances ofthegraphitewere0.353nmand

0.343nmontheC-faceandontheSi-facerespectively• ThecarbonlayerformedontheC-faceisalmost50

timesthickerthantheoneontheSi-face. Thedifferences intherespectivefacesregardingthicknessandin

thestructuresofcarbonlayersareremarkable, nvertheless bothbulkstructuresarecommon• Thedifference

inthethicknessofdecomposed layersontherespectivefaceshasalreadybeenreportedfromtheresultsof

ellipsometric measurements (ref.9)andAESresults(ref.2).Anumberofresearchers dealingwithSiCsingle

crystalshaveshownthatheatingSiCgavearoughersurfaceontheC-facethanontheSi-face. However, the

causeofit hasnotbeenexplained•

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_°°.oO%,,%. %%%_o,.°.•,.._/,.°,.%,°... _-°,%-, %,.,%. i_ .._.% ° %.._ .,%, %o • ,o ,...o%.. _4%,, _ %. o.,°.. • ,¢.%%%,.,

_.; 2.'-:-:.i.:.12.::.,:.. : :.'-_,:..-_.'..1-'.'...R.', ."5R-...,_ ..,G'.I _,J.'.'.'..',:-:.,.R,:.: '.7.'.:,::i.:-Z,','.°.2 .'_'. "'""''L'i2''_' "':':',';° '1

Figure 2. HREM images showing surfaces of the C-face heated at (a) 1250°C and

(b)1300°C for 0.5 h.

Figure 3. A HREM image showing the surface of the Si-face heated at 1350°C for 0.5 h.

NASA/CP--2001-210948 797 August 6, 2001
Cross-sectional HREM of decomposed surfaces on the initial stage was carried out to investigate the

mechanism of the surface decomposition on both the Si-face and the C-face. Figures 2(a) and (b) show

micrographs of surfaces of the C-face heated at 1250°C and 1300°C for 0.5 h, respectively. The amorphous

contrasts on the upper side of each micrograph in Fig. 2 and in Fig. 3 are due to the glue used during the

specimen preparation. The glue played an important role in protecting the carbon structures from the ion

bombardment, though the amorphous contrast strongly influences the weak contrasts of carbon structures at

the initial stage.

At 1250°C, 4 layers of slightly arced graphite sheets of 5 nm in diameter and 1-2 nm in height could be

found dispersed on the C-face as shown by arrows in Fig. 2(a). This corresponds to the initial stage of the

decomposition of the C-face. Furthermore, when it was heated to 1300°C, 2-3 layered carbon nanocaps

3-5nm in diameter and 3-5 nm in height were generated densely over the surface as shown in Fig. 2(b). Here,

cross-sectional HREM was used to clarify an atomically uneven-surface structure of the C-face, which has

not been easy to detect by STM.

When the Si-face was heated to 1350°C for 0.5 h, only two graphite sheets were observed on the surface, as

shown by arrows in Fig. 3. The graphite sheets were observed flat and parallel to the surface for every

location.

From these experimental results, the decomposition processes of the Si- and C-faces can be expressed as

follows: On the Si-face, a mono-atomic layer of graphite formed epitaxially, and the layers grow as graphite

(0001)//6H-SiC(0001). Two- and three-layer collapse mechanisms (ref. 2 and 10) were proposed to explain

the orientation and the carbon density of the graphite layers formed by the surface decomposition of SiC. The

carbon density of the three condensed layers, 3.67x1015 atoms cm -2, is also close to the carbon density of a

graphite layer 3.80x1015 atoms cm -2 (ref. 2). This mechanism somehow explains the formation of the

monocrystalline graphite layer parallel to the surface seen on the Si-face. However, it is difficult to explain

the distinctive differences between the faces.

On the C-face, carbon islands would be nucleated at impurity atoms or some faults on the surface, and

grow up into the particles of about 5 nm in size. Here, two types of carbon islands can be considered, such as

spherical particles of amorphous carbon and graphite fragments. The spherical carbon particles would be

graphitized from the surface. Whereas, the two to five layered topotaxial graphite fragments would be

generated at the steps of the SiC surface. Since g-electronic state at the edges of the nano-scale graphite

sheets is considered to be active, the edge of the graphite tends to curve to connect with the carbon atoms on

the surface. Furthermore, SiO gas molecules would be continuously generated at the boundary between SiC

and the graphite clusters. As a result, carbon nanocaps (ref. 8) should be finally formed. The diameter of the

nanocaps would depend on the SiC decomposing rate, the surface diffusion coefficient of carbon, the

cohesive energy of carbon at the decomposing temperature, and the flatness of the SiC surface.

The difference in thickness of the carbon layers on the respective faces is also noteworthy, nevertheless

both of the bulk structures are identical. The difference in thickness of the decomposed layers has also been

reported from the results of ellipsometric measurements (ref. 9) and AES results (ref. 2). However, its cause

has not yet been elucidated. As explained in the present model, the dense graphite sheet grows one by one

parallel to the surface in the case of the Si-face, which would protect the inner side of the Si layer from the

oxidation process. In the case of the C-face, carbon nanocaps form from the initial stage of the decomposition

NASA/CP--2001-210948 798 August 6, 2001
andthealignedCNTsgrowperpendicular tothesurface, whichallowstheoxygenatomstomigrateeasilyto

theSiCsurfacethroughthedecomposed carbonlayer.Thesestructural differencesfromtheinitialstage

mightbeoneoftheoriginsforthedifferenceofcarbon-layer thicknessbetweenthebothfaces.

Insummary, it wasclarifiedthataCNTfilmwillorganizeonlyontheC(000-1) faceofSiC,andnotonthe

Si(0001)-face, bythesurfacedecomposition oftheSiCcrystal.Additionally, weshowedthatCNTscouldnot

growwithouttheformationofnanocapsduringtheinitialstageof thegrowthprocessbycomparingthe

differenceofthestartingdecomposition mechanisms betweentheSi-andC-facesofSiC.Thisinitial-nanocap

mechanismmightbesignificant forelucidatingoneofthepreviouslyreportedCNTsynthesismethods.

Thissurfacedecomposition methodhaspotential foruseinnewcarbondesigning, utilizingeachsurface

structureofmetalcarbides.

ACKNOWLEDGEMENT

TheauthorsthankDr.S.Harafor valuablesuggestions. ThisworkwassupportedbyFCTproject,

whichwascosignedtoJFCCbyNEDO.

REFERENCES

1. W. F. Knippenberg, "Growth Phenomena in Silicon Carbide," Philips Research Reports, 18, 161-274

(1963).

2. A. J. Van Bommel, J.E. Crombeen, and A. Van Tooren, "LEED and Auger Electron Observations of the

SIC(0001) Surface," Surf. Sci., 48, 463-72 (1975).

3. F. Owman and P. Martensson, "The SiC(0001)6.3x6.3 reconstruction studied with STM and LEED,"

Surf. Sci., 369,126-36 (1996).

4. M. A. Kulakov, P. Heuell, V. F. Tsvetkov, and B. Bullemer, "Scanning Tunnelling Microscopy on the 6H

SIC(0001) Surface," Surf. Sci., 351,248-54 (1994).

5. C. S. Chang and I.S.T. Tsong, " Scanning Tunneling Microscopy and Spectroscopy of Cubic _3-SiC(111)

Surfaces," Surf. Sci., 256, 354-60 (1991).

6. M. Kusunoki, M. Rokkaku, and T. Suzuki, "Epitaxial Carbon Nanotube Film Self-Organized by

Sublimation Decomposition of Silicon Carbide," Appl. Phys. Lett., 71, 2620-22 (1997).

7. M. Kusunoki, M. Rokkaku, and T. Hirayama, Aligned Carbon Nanotube Film Self-Organized on a SiC

Wafer," Jpn. J. Appl. Phys., 37, L605-L606 (1998).

8. M. Kusunoki, T. Suzuki, K. Kaneko, and M. Ito, "Formation of Self-Aligned Carbon Nanotube Films by

Surface Decomposition of Silicon Carbide," Phil. Mag. Lett., 79, 153-61 (1999).

9. F. Meyer and G. J. Loyen, " Ellipsometry Applied to Surface Problems," Acta Electronica, 18, 33-38

(1975).

10. D.V. Badami, "X-ray Studies of Graphite Formed by Decomposing Silicon Carbide," Carbon, 3, 53-57

(1965).

NASA/CP--2001-210948 799 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

CARBON NANOTUBES AND CARBON NANOFIBERS SYNTHESIZED BY

SUBLIMATING DECOMPOSITION OF SILICON CARBIDE WITH CATALYSTS

Kazuo Tsugawa and Akiko Goto

Joint Research Consortium of Frontier Carbon Technology (FCT), Japan Fine Ceramics Center (JFCC), c/o Research

Center for Advanced Carbon Materials, National Institute of Advanced Industrial Science and Technology (AIST)

Tsuknba Central 5, 1-1-1 Higashi, Tsuknba, Ibaraki 305-8565, Japan

Yoshinori Koga

Research Center for Advanced Carbon Materials, AIST Tsuknba Central 5, 1-1-1 Higashi, Tsuknba, Ibaraki 305-

8565, Japan

ABSTRACT

Carbon nanotubes (CNTs) and carbon nanofibers (CNFs) have been synthesized by heating a polycrystalline SiC

substrate in a vacuum (10 2Torr) at 1000 - 1200 °C for 1 hr. The present work utilizes metal catalysts, such as Au, Pt,

Fe, Ni, Co, Pd, Rh, and PtPd, in order to lower the heating temperature. As a result, CNTs and CNFs have grown at

1200 °C in the case of Pt, Ni, Co, Rh, PtPd, while CNT growth requires the heating temperature of over 1500 °C

without a catalyst. Especially in the case of the PtPd catalyst, CNTs and CNFs have grown at 1100 °C. The

synthesized CNTs or CNFs indicate a diameter from 2 to 300 nm and a length of 100 nm - 10 gin. In addition, most

of CNTs and CNFs include, at the tip of them, a particle with a similar diameter to those of them, which is thought to

be the metal catalyst or its carbide.

Keywords: carbon nanotube, carbon nanofiber, sublimating decomposition, silicon carbide, catalysts

INTRODUCTION

Since the first observation of carbon nanotubes (CNTs) in deposition on the carbon cathode electrode used in an

arc-discharge evaporation method (ref. 1), various CNT growth techniques have been developed, such as carbon arcdischarge (ref. 2), laser ablation (ref. 3), chemical vapor deposition (CVD) (ref. 4), and so forth. Arc-discharge and

laser ablation methods involve the condensation of carbon atoms generated fi'om evaporation of solid carbon sources.

CVD method involves the dissociation of hydrocarbon gas catalyzed by transition metal, and dissolution and

saturation of carbon atoms in the metal particle. It is considered that CNTs grow fi'om the carbon saturated metal

catalyst.

As another approach to CNT formation different fi'om these methods in growth mechanism, Kusunoki et al. (refs.

5 to 7) have synthesized CNTs by sublimating decomposition of SiC. They have used single crystalline c_-SiC (0001)

wafers for the substrates and heated them by laser (ref. 5) or by an electric furnace (ref. 6) in a vacuum (10 4 Torr) at

1500 - 1700 °C without a catalyst. In this case, Si atoms are removed fi'om the SiC substrate by active oxidation, and

CNTs grow perpendicularly to the SiC (0001) surface, eroding the subsU'ate downward. In this manner, a selforganized aligned CNT film is formed (ref. 7).

SiC is the only known, naturally stable, IV-IV compound semiconductor. SiC has been of particular interest due

to the fact that it is an important material for devices designed to operate at high temperature, high power, high

fi'equency, and in harsh environment (ref. 8). Actually, various prototype SiC devices have been fabricated using

planar process technology. On the other hand, CNTs are unique nanostructures with remarkable electronic and

mechanical properties, which are promises for practical applications (ref. 9). With this respect, the combination of

SiC and CNTs may bear unique applications. Thus the further investigations of CNTs on a SiC substrate are required.

We have utilized metal catalysts, such as Au, Pt, Fe, Ni, Co, Pd, Rh, and PtPd, to lower the heating temperature.

Polycrystalline c_-and 13-SICdisks have been used for the substrates since the aim of this study is to lower the

heating temperature, not to form aligned CNT films. As a result, the heating temperature is reduced to 1100 °C using

the PtPd catalyst.

NASA/CP--2001-210948 800 August 6, 2001
EXPERIMENTAL

Thestartingsubstrateispolycrystalline c_-and13-SICdisks.Commerciallyvailablec_-and13-SICpowders,

rangingfrom0.1to2gminsize,werecompressed underapressureof500 kgw/cm 2 to form a bulk specimen. Then

it was sintered at 1650 °C for 20 min in a N2 atmosphere of 300 Torr to obtain a reasonable hard sample. The

sintered sample was cut into small disks of 6 x 3 x 1 mm in size. The SiC disk was coated with a metal catalyst using

vacuum evaporation or sputtering. Au, Pt, Fe, Ni, Co, Pd, Rh, and PtPd (Pd:20 wt. %) were used as the catalyst. The

nominal coating thickness was 100 nm. The metal-coated substrates were annealed in a vacuum electric furnace

illustrated in Fig. 1 at 1000 - 1200 °C for 1 hr with a base pressure of 10 2 Torr. The heating rate from room

temperature until the setting temperature (1000 - 1200 °C) is 6 - 8 °C/min. The heating (substrate) temperature was

monitored with a thermocouple contacted with the dummy SiC substrate. After the annealing, the surface of the

substrate was observed by a Hitachi S-5000 scanning electron microscope (SEM) and a JOEL JEM-2000SX II

transmission electron microscope (TEM).

Quartz tube Diaphragm manometer

Electric furnace

vaouum, pump

Quartz boat Thermocouple i controller eater I

Fig. 1 Experimental setup.

RESULTS AND DISCUSSION

Figure 2 shows SEM images of a polycrystalline SiC substrate before the metal coating (Fig. 2 (a)) and after

the PtPd coating by sputtering (Fig. 2 (b)). The SiC grain size ranges from 0.1 to 2 gin, and each SiC particle is

covered with the PtPd catalyst after the coating. The situation is almost the same in the case of other metal catalysts.

":i.. i_... :

: "ii!_ " .::_

,,_

•. .._

(a) (b)

Fig. 2 SEM images of a SiC substrate. (a) Before metal coating. (b) After metal coating.

NASA/CP--2001-210948 801 August 6, 2001
Themetal-coated SiCsubstrates wereannealedat1000- 1200 °C in a vacuum electric furnace shown in

Fig. 1 for 1 hr. As a result, tubular nanostructures grew in the case ofPt, Ni, Co, Rh, and PtPd at 1200 °C. In the case

of the PtPd catalyst, the same nanostructures were formed at 1100 °C. Figure 3 shows the tubular structures, ranging

from 2 to 50 nm in diameter and from 500 nm to 10 gm in length, formed at 1200 °C using the PtPd catalyst. In other

cases of catalysts (Pt, Ni, Co, Rh), the diameter of the tubular structures is 20 - 50 nm for Pt, and 50 - 300 nm for

the others. Most of the tubular structures include a particle at the tip with the same diameter as those of them, which

thought to be the metal catalyst or its carbide. Figure 4 (b) indicates a reflection electron image of the tubular

structures shown in Fig. 3. The secondary electron image of the same area as the reflection electron image is shown

in Fig. 4 (a). The high contrast of the particles in Fig. 4 (b) exhibits different material from the tubular structures

including them. Thus the particles are thought to be the metal catalysts or their carbide. Using the PtPd catalyst, the

same structures have grown at 1100 °C. However, no tubular structures have grown at 1000 °C with any catalyst. In

addition, the results are not different between c_-and B-SiC substrates.

Fig. 3 SEM photograph of the synthesized CNTs and CNFs (heating temperature = 1200 °C, PtPd catalyst).

(a)

:::::::::::::::::::::::::: ..

.. ._._.. ..-.....

_.i'.. ..:::._::. """

(b)

Fig. 4 SEM photographs of the synthesized CNTs and CNFs (heating temperature = 1200 °C, PtPd catalyst).

(a) Secondary electron image. (b) Reflection electron image.

The TEM observation was carried out on the sample heated at 1200 °C with PtPd catalyst. To prepare the

TEM sample, the surface of SiC substrate with the tubular structure on it was scratched. The powders including the

tubular structures were suspended in ethanol, sonicated, and then a drop was put on a Cu microgrid covered with a

carbon thin film. The observation under TEM at an acceleration voltage of 200 keV has revealed that the tubular

structures consists of CNTs and carbon nanofibers (CNFs) with amorphous-like random structure. Figure 5 shows

NASA/CP--2001-210948 802 August 6, 2001
typical TEM images of the synthesized tubular structures. Figure 5 (a) shows a TEM photograph of the synthesized

CNT. This image indicates a multi-walled CNT (MWNT), which consists of 4 graphitic layers. The inner diameter is

approximately 0.7 nm and the outer diameter is 2.5 nm. Figure 5 (b) indicates a TEM photograph of the synthesized

CNFs with a diameter of 15 - 20 nm. The CNFs in Fig. 5 (b) exhibit amorphous structure without a hollow structure.

Generally speaking, tubular structures with a diameter of over 10 nm are found to be CNFs and ones with a diameter

of less than 10 nm are found to be CNTs.

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::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::::

..................... • o*o**o*o*

:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:i:iii:i:i:iii_

i!!!iii!iiiiiii;!iii!!ii

(b)

Fig. 5 TEM images of the synthesized CNTs and CNFs. (a) CNTs. (b) CNFs.

In this experiment, a metal particle exists at the tip of the CNTs and CNFs. This situation is similar to the tip

growth mode of CVD method (ref. 9). During CVD growth, carbon atoms to construct a CNT are provided fi'om the

hydrocarbon atmosphere into the metal particle. However no carbon source exist around the metal particle at the tip

of CNTs and CNFs here. Carbon atoms to form CNTs and CNFs are provided only fi'om their bottom, i. e., the SiC

substrate. By SEM and TEM analysis, the metal catalyst is not always observed around the bottom of the CNTs and

CNFs. However the melted metal catalyst is thought to exist around the bottom during the heating process, since

CNTs do not grow at a temperature of less than 1500 °C without any catalyst (ref. 7). During the cooling process to

room temperature, the metal catalyst condenses into solid and forms islands. Actually discrete metal islands have

been observed by SEM after the heating, while the metal covers all over the substrate surface before the heating. In

addition, the lower the metal coverage on the SiC substrate becomes, the lower density of CNTs and CNFs yields.

This implies the growth mechanism is similar to the base growth mode of CVD method (ref. 9), and the metal

particles at the tip contribute only the first stage of the growth. That is to say, a CNT or a CNF starts to grow around

a metal particle, and the diameter of the particle defines the diameter of the CNT or the CNE

In summary, we assume the following growth mechanism. The SiC substrate starts to decompose at over

1000 °C by the oxidation of Si by the residual oxygen in the vacuum chamber, and the oxidized Si is released fi'om

the substrate as a SiO gas (ref. 7). The residual C atoms dissolve and saturate in the metal particle. The precipitation

of carbon fi'om the saturated metal particle leads to the formation of tubular carbon solid. Once a CNT or a CNF

starts to grow, it continues to grow acquiring C atoms fi'om the saturated metal around its bottom. At this stage, the

metal catalyst accelerates the decomposition of SiC and growth of a CNT or a CNF. In fact, CNTs and CNFs

synthesized in this study are much longer than those of CNTs synthesized without a catalyst. For example, the

longest one by 1 hr heating with the PtPd catalyst exhibits approximately 10 gm in length, while CNTs by 6 hr

heating without a catalyst exhibit only 180 nm in length. In other words, the growth mechanism is thought to be

similar to that of CVD method, namely solid-source CVD, rather than the sublimating decomposition method

without a catalyst.

NASA/CP--2001-210948 803 August 6, 2001
SUMMARY

CNTs and CNFs have been synthesized by heating metal-coated polycrystalline SiC substrates at 1100 -

1200 °C in a vacuum (10 2 Torr). Au, Pt, Fe, Ni, Co, Pd, Rh, and PtPd have been used for the metal coating. As a

result, CNTs and CNFs have grown at 1200 °C using Pt, Ni, Co, Rh, and PtPd, and at 1100 °C using PtPd, while a

heating temperature of over 1500 °C is required without a catalys. SEM observations reveal that the synthesized

CNTs and CNFs range from 100 nm to 10 gm in length. TEM observations reveal that the synthesized CNTs are

MWNTs and range from 2 to 10 nm in diameter, and CNFs range from 10 to 300 nm in diameter. The growth

mechanism is thought to be similar to a solid-source CVD method.

ACKNOWLEDGEMENTS

This work is supported by the Frontier Carbon Technology (FCT) Project, which is consigned to Japan Fine

Ceramics Center (JFCC) by New Energy and Industrial Technology Development Organization (NEDO) of Japan.

REFERENCES

1.

2.

3.

4.

5.

6.

7.

8.

9.

Iijima, S.: Helical Microtubules of Graphitic Carbon. Nature, vol. 354, Nov. 1991, 56.

Ebbesen, T. W.; and Ajayan, P. M.: Large Scale Synthesis of Carbon nanotubes. Nature, vol. 358, Jul. 1992,

220.

Thess, A., et al.: Crystalline Ropes of Metallic Carbon Nanotubes. Science, vol. 273, Jul. 1996, 483.

Kong, J.; Cassell, A. M.; and Dai, H.: Chemical Vapor Deposition of Methane for Single-Walled Carbon

Nanotubes. Chem. Phys. Lett., vol. 292, Aug. 1998, 567.

Kusunoki, M.; Rokkaku, M.; and Suzuki T.: Epitaxial Carbon Nanotube Film Self-Organized by

Sublimation Decomposition of Silicon Carbide. Appl. Phys. Lett., vol. 71, no. 18, Nov. 1997, 2620.

Kusunoki, M., et al.: Aligned Carbon Nanotube Film Self-Organized on a SiC Wafer. Jpn. J. Appl. Phys.,

vol. 37, part 2, no. 5B, May 1998, L605.

Kusunoki, M., et al.: Formation of Self-Aligned Carbon Nanotube Films by Surface Decomposition of

Silicon Carbide. Phil. Mag. Lett., vol. 79. no. 4, Apr. 1999, 153.

Camassel, J.; Contreras, S.; and Robert, J.-L.: SiC Materials: A Semiconductor Family for The Next

Century. C. R. Acad. Sci. S6rie IV, vol. 1, no. 1, Jan. - Feb. 2000, 5.

See, for instance, Dresselhaus, M. S.; Dresselhaus, G.; and Avouris, Ph. Eds.: Carbon Nanotubes: Synthesis,

Structure, Properties, and Applications. Springer, Berlin, 2001.

NASA/CP--2001-210948 804 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

SYNTHESIS OF CARBON NANOTUBES BY MICROWAVE HEATING

Eun Hwa Hong, Kun-Hong Lee*

Department of Chemical Engineering, Electrical and Computer Engineering Division, Pohang University of Science

and Engineering, San 31, Hyoja-Dong, Nam-Ku, Pohang, Kyungbuk, 790-784, South Korea

Sang Ho Oh, Chan-Gyung Park

Department of Materials Science and Engineering, Pohang University of Science and Engineering, San 31, HyojaDong, Nam-Ku, Pohang, Kyungbuk, 790-784, South Korea

Chang-Mo Ryu

Department of Physics, Pohang University of Science and Engineering, San 31, Hyoja-Dong, Nam-Ku, Pohang,

Kyungbuk, 790-784, South Korea

Jong Hun Han, Jae Eun Ryu

Iljin Nanotech Co., Ltd., Nano Technology Center(R-502), Kayang Techno Town 1487, Kayang-Dong, Kangseo-Ku,

Seoul 157-810, South Korea

Corresponding Author : Kun-Hong Lee*

Department of Chemical Engineering, Pohang University of Science and Engineering, San 31, Hyoja-Dong, NamKu, Pohang, Kyungbuk, 790-784, South Korea

e-mail : ce20047(_posle, cl-.ac.kr

Fax : (82)54-279-8298

Tel : (82)54-279-2271

ABSTRACT

We report a novel method for the synthesis of carbon nanotubes (CNTs) by microwave irradiation. CNTs were

successfully synthesized on various supports such as carbon black, silica powder and organic polymer substrates. In

our method, microwave was directly irradiated on the catalyst particles which were impregnated or painted on the

substrates. Since the microwave energy is selectively absorbed by the catalyst particles, not by the substrates, local

heating occurs, resulting in the CNTs synthesis even on the organic polymer substrates with minimum degradation of

the substrate materials.

Cobalt or cobalt sulfide was loaded on the substrates as catalysts. Acetylene was used as a hydrocarbon source

under microwave irradiation. We have observed temporary arcing phenomenon during microwave irradiation and the

substrate was covered with CNTs without any damage in its appearance. Various shapes of carbons were observed

including CNTs, nanofibers, nanoparicles and amorphous carbons. It is an important merit of our method that we

can grow CNTs virtually on any substrate provided its absorption of microwave energy is small.

Keywords: carbon nanotubes, synthesis, microwave

INTRODUCTION

CNTs have drawn a great deal of attention since Iijima reported the existence of them in 1991 (ref. 1). Various

methods have been reported for the synthesis of CNTs : arc-discharge (ref. 1 to 4), laser ablation (ref. 5), plasma

synthesis (ref. 6 to 7), chemical vapor deposition (CVD) (ref. 8 to 12), flame synthesis and Smalley's recent

invention of high pressure CO (HIPCO) process (ref. 13 to 14). None of them is able to synthesize CNTs at low

NASA/CP--2001-210948 805 August 6, 2001
temperature (say,lessthan200°C), sothatheincorporation ofCNTswithlowmeltingpointmaterialssuchas

organicpolymershasbeenseverelylimited.

Plasmasynthesis orCVDtechniques havebeenusedforthelowtemperature synthesis ofCNTs,(ref.15to18)

thoughexperiments werecarriedoutaround500°Cevenatreducedpressures. Therefore, it'sdifficulttoadapt

organicpolymersforvariousapplications ofCNTs,especially forthesubstrateoffieldemissiondisplays(FEDs).

Inthispaper,weintroducemicrowavehatingasasubstitutewithconventional heatingmethodtoproduce

CNTsonlowmeltingtemperature materials. Intheconventional heatingmethods, wholebodyofthereactorwas

heatedbyaheatingelement. Ontheotherhand,localheatingispossiblebyusingsuitablematerialsthatstrongly

absorbmicrowaveradiationandthenefficientlytransferthisenergytoreactants. Sincemostmaterialsdonotabsorb

electromagnetic energyofthemicrowave(ref.19to20),onlycatalyst particlescanbeselectivelyheatedtothe

temperature ofCNTsynthesis. Thesubstrateisnotheatedbythemicrowave, sothatit ispossibletoincorporate

CNTswithorganicpolymersinvariousapplications.

Inthisstudy,wereporttheCNTsynthesis bymicrowaveirradiationonthecatalyst particlesunderatmospheric

pressure. Weinvestigatedhemicrowave-induced reactionusingcobaltcatalyst loadedonvarioussupportsanda

substrate, whichisoneofthecommontransitionmetalcatalystsforCNTgrowth.

EXPERIMENTALDETAILS

Preparation of the catalysts

Cobalt (Co) was chosen as a catalyst and cobalt nitrate was used as its precursor material. Aqueous solution of

cobalt nitrate was loaded on two types of supports by impregnation : a microwave receptor material (carbon black,

CABOT, VULCAN) and an insulator material (SiO2 powder, Devisil). Their characteristics and metal contents are

shown in Table 1. The impregnated catalysts were dried overnight at ll0°C and then reduced in the flowing H2 at

500°C for 3hrs. The reduced cobalt catalyst loaded on SiO2 powder was additionally sulfided using 10 vol% H2S/H2

gas mixture at 400°C for 2hrs. Hereafter, it is denoted as Co(10)/SiO2. The phase of the catalyst was identified by

XRD (x-ray diffraction, MAC Science Co.) which utilizes Ni-filterd CuKc_ radiation ()v=l.5405A). It was observed

that cobalt existed as a metal phase after reduction and cobalt sulfide formed after sulfiding.

Table 1. Characteristics of the supports

Bet surface area Metal Content

Support (m2/g) Type (wt%)

C (carbon black) 199.6 Microwave receptor 5wt%

SiO2 (60-100mesh) 509.0 Microwave insulator 10wt%

Preparation of the substrate

To prepare cobalt sulfide loaded substrate, Teflon sheet of 2mm thick was used as a substrate. 1 gram of cobalt

black (Junsei co. Japan) were mixed with 10ml of Nation solution (5wt% aqueous solution, Aldrich) and sonicated

for 10minutes to disperse them well. The mixture was then painted on a sheet of Teflon to form a catalyst film. After

drying overnight at ll0°C, the substrate was sulfided with 10 vol% H2S/H2 mixed gas at 300°C for 10hours. Cobalt

sulfide particles, which acted as the catalyst for CNT growth, formed on the substrate.

Synthesis and Characterization

A quartz boat which contained catalyst powders was placed in a quartz reactor. Purified acetylene (C2H2),

hydrogen (H2), and hydrogen sulfide (H2S) gases were used to purge any air in the reactor before the start of the

microwave irradiation. Microwave was irradiated by a commercial microwave oven (2.45GHz, 800W, SAMSUNG)

with flowing reactant gases.

The formation of CNTs was identified by SEM and TEM. The morphology of CNTs and the catalyst particles

was observed by FE-SEM (field emission scanning electron microscope, Hitachi) and the structure by HR-TEM

(high resolution transmission electron microscope, JEOL 2010).

NASA/CP--2001-210948 806 August 6, 2001
RESULTS AND DISCUSSIONS

The reaction conditions for CNTs synthesis are summarized in Table 2. Carbon yield (%) was calculated by the

equation (1) based on the weight of catalysts.

Carbon yield (%) = (weight after reaction - weight of catalyst) x 100 / weight of catalyst ................ (1)

Table 2. Reaction conditions for CNTs Synthesis

Reactant gas Microwave CNT Carbon yield

Entry Catalyst* Weight (cm3/min) irradiation time

(g) synthesis (%)

C2H2/H2/H2S (min)

Fig. l(b) Co(5)/C 0.05 5 / 0 /0 30 yes 354

Fig. 2(a) Co(10)/SiO2 0.2 10 / 0 / 0 30 no 0

Fig. 2(b) Co(10)/SiO2 0.2 10 / 0 / 3 30 yes 7.5

Fig. 2(c) Co(10)Sx/SiO2 0.2 10 / 0 / 0 5 yes 15.5

Fig. 3 CoSx/Teflon 10 / 40 / 0 0.5 yes

* The number in the parenthesis is the weight percent.

Before the main experiments, we have performed control experiments in which microwave was irradiated with

flowing hydrocarbon gases, but without catalysts. No carbon was synthesized without catalysts. This confirmed that

microwave alone couldn't stimulate acetylene (C2H2) gas to produce carbon.

Addition of catalysts gave different results. In case of the catalysts loaded on the microwave receptor, Co(5)/C,

CNTs were grown even with a small amount of C2H2. Carbon yield went up to 354% after 30 minutes of microwave

irradiation. C (carbon black) without cobalt catalyst didn't result in the CNTs synthesis though there was a little

amount of carbon production. Both cobalt and carbon are microwave receptor materials. Therefore, they both absorb

microwave energy, so that they are easily heated to the temperature of carbon formation. However, it seems that

cobalt catalyst is responsible for the synthesis of CNTs.

Fig. 1 shows the FE-SEM images of fresh catalyst Co(5)/C and CNTs synthesized by microwave-induced

reaction of C2H2. There are no filamentous carbons before the reaction and only the carbon black supports are

noticeable. After microwave-induced reaction, long and bent shapes of CNTs are clearly observed as well as

amorphous and filamentous carbons.

Insulating support, SiO2, behaved differently. No CNTs were synthesized by just irradiating microwave even

with the cobalt catalysts as shown in Fig. 2(a). Although cobalt can absorb microwave energy, that seems to be not

sufficient for CNTs synthesis because SiO2 is an insulating material for microwave. It doesn't absorb microwave

energy, but it rather takes up the energy absorbed by the cobalt catalysts and dissipates it.

To solve this problem, we considered a candidate reaction system with a higher catalyst activity. We adopted

hydrogen sulfide (H2S) as a co-reactant gas to improve the activity of the catalyst loaded on the SiO2 support. CNTs

growth occurred with the addition of a small amount of H2S to C2H2 in the presence of Co/SiO2 catalyst after

30minutes of microwave irradiation,as shown in Fig. 2(b).

Fig. 1. Growth of CNTs on the catalyst-loaded Carbon black support.

(a) fresh catalyst, Co(5)/C.

(b) CNTs grown on the catalyst after microwave irradiation.

NASA/CP--2001-210948 807 August 6, 2001
Carbonyield,however, wasmuchsmaller(only7.5%)thanthecaseofCo/C(354%). Varioustypesofcarbon

structures wereobservedbyHR-TEM:CNTs,carbonnanofibers havingherring-bone wallstructureandamorphous

carbonparticles. H2Sitselfdidn'treactwithC2H2bymicrowaveirradiationintheabsenceofcatalysts. Wecouldnot

observeCNTsformationwithothergasesthanH2SevenwithCocatalysts.

Weobservedthathechangeofcobaltmetalphasetocobaltsulfide.Cobaltsulfidemaybeamuchefficient

catalystforthesynthesis ofCNTs.Toanswer thisquestion, cobaltsulfidewasloadedonSiO2(named, CoSx/SiO2) at

thefirstplace,thenmicrowavewasirradiatedwithflowingC2H2gas.Again,amorphous carbons, filamentous

carbonsandherringbonetypecarbonnanofibers aswellasCNTswereobserved(Fig.2c).Comparedtotheprevious

case,however, thereactiontimeforCNTsgrowthwasreducedto5minutes, whilethecarbonyieldusingCoSx/SiO2

catalystbecamealmost doubledfromthatofCo/SiO2catalyst after30minutes reaction. Therefore, Itcanbe

concludedthatheexistenceofsulfurasacomposition ofreactant gasorcatalyst isnecessary tothemicrowaveinducedCNTsynthesiswhereaninsulatingmaterial isusedasasupport, andmetalsulfideismucheffective

catalyst thanpuremetalcatalysts.

.i°-° ° ° .,..,°°

i

:N:_.:::::..;_.,' .°.%°. ° °o °°°°,*o.

:..,,..:>:.:.:.:.:.:.:.:

::::::::::::::::::::::::

tN_....:_.'._::_:!:_i_

N

_..;_ ._...,,

Fig. 2. Growth of CNTs on the catalyst-loaded SiO2 support.

(a) No CNTs without H2S°

(b) CNTs with H2S°

(c) CNTs grown on Co(10)Sx/SiO2.

As an extension of these experiments, CNTs growth on an organic polymer substrate was carried out. Cobalt

sulfide particles, which acted as the catalyst for CNTs growth, were loaded on a Teflon sheet. Microwave was

irradiated with flowing C2H2 without H2S. C2H2 was pyrolized by an extremely short period of microwave

irradiation and CNTs indeed formed on the Teflon sheet. We observed temporary arcing phenomenon during

microwave irradiation. Teflon substrate was covered with carbon nanotubes without damage in its appearance.

Fig. 3(a) shows the FE-SEM image of CNTs synthesized after 30 seconds of microwave irradiation. Carbon

nanotubes grown on cobalt sulfide catalyst particles were short and straight. Various shapes of CNTs and carbon

nanoparticles were observed as well as amorphous carbons. HR-TEM images of the multi-walled carbon

ii_k_-!_'_:!!!;_!i!!!!!_iii!:_.:.,::::.::::ii'

Fig. 3. (a)FE-SEM image of CNTs grown on CoSx/Teflon. (b)HR-TEM image of a carbon nanotube

NASA/CP--2001-210948 808 August 6, 2001
nanotubes producedinthiswaywereshowniFig.3(b). Theyaretensofnanometers thick,havehollowcoresand

closed-ends, andtheirfringestructures, whichwerecoveredwithamorphous carbon, werewellconnected.

CONCLUSIONS

Westudiedanewtechniqueforthesynthesis ofCNTsbymicrowavehating. CNTsweresuccessfully

synthesized onthecatalystsloadedonmicrowavereceptorinsulatorsupports, andevenonanorganicpolymer.

Sulfur,asH2Sgasormetalsulfidecatalyst, haveacriticalroleinCNTsynthesiswithmicrowaveirradiation.

Althoughpresent resultcontainsvariousformsofcarbonmaterials, thequalityofCNTscanbeimprovedbythe

optimizationfreactionconditionsandbytheselectionofbettercatalyst. Inconclusion, microwaveinducedreaction

madeitpossibletoincorporateCNTswithlowmeltingpointmaterialssuchasorganicpolymers. It maybecomean

important advancement tothefabricationofaflexiblefieldemissiondisplay(FED).

REFERENCES

1. Iijima, S.: Helical microtubules of graphitic carbon. Nature, 354, 1991, 56.

2. Bonard, J.-M. et al.: Field emission from single-wall carbon nanotube films. Appl. Phys. Lett., 73, 1998,

918.

3. Dillon, A.C. et al.: Storage of hydrogen in single-walled carbon nanotubes. Nature, 386, 1997, 377.

4. Journet, C. et al.: Large-scale production of single-walled carbon nanotubes by the electric arc technique.

Nature, 388, 1997, 756.

5. Thess, A. et al.: Crystalline ropes of metallic carbon nanotubes. Science, 273, 1996, 483.

6. Hatta N., Murata K.: Very long graphitic nano-tubules synthesized by plasma-deposition of benzene. Chem.

Phys. Lett., 217, 1994, 398.

7. Iijina S., Ichihashi, T.: Single-shell carbon nanotubes of 1-nm diameter. Nature, 363, 1993, 603.

8. Li W.Z. et al.: Large-scale synthesis of aligned carbon nanotubes. Science, 274, 1996, 1701.

9. Fan. S. et al.: Self-oriented regular arrays of carbon nanotubes and their field emission properties. Science,

283, 1996, 512.

10. Hernadi K. et al.: Catalytic synthesis and purification of carbon nanotubes. Synthetic Metals, 77, 1996, 31.

11. Huang Z.R et al.: Growth of highly oriented carbon nanotubes by plasma-enhanced hot filament chemical

vapor deposition. Appl. Phys. Lett., 73, 1998, 3845.

12. Qin L.C. et al.: Growing carbon nanotubes by microwave plasma-enhanced chemical vapor deposition.

Appl. Phys. Lett., 72, 1998, 3437.

13. Randall L. et al.: Diffusion flame synthesis of single-walled carbon nanotubes. Chem. Phys. Lett., 323, 2000,

217.

14. Smalley, et al.: Gas-phase nucleation and growth of single-wall carbon nanotubes from high pressure CO.

WO00/26138.

15. Ren Z.R. et al.: Synthesis of large arrays of well-aligned carbon nanotubes on glass. Science, 282, 1998,

1105.

16. Byron K.E.: Plasma catalysis of carbon nanofibers. WO99/58748.

17. Choi Y.C. et al.: Low temperature synthesis of carbon nanotubes by microwave plasma-enhanced chemical

vapor deposition. Synthetic Metals, 108, 2000, 159.

18. Lee C.J. et al.: Low-temperature growth of carbon nanotubes by thermal chemical vapor deposition using

Pd, Cr, and Pt as co-catalyst. Chem. Phys. Lett., 327, 2000, 277.

19. Thostenson E.T., Chou T.-W.: Microwave processing: fundamentals deposition and applications.

Composites: Part A, 30, 199, 1055.

20. Galema S.A.: Microwave chemistry. Chem. Soc. Rev., 26, 1997, 233.

NASA/CP--2001-210948 809 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

MULTI-WALLED CARBON NANOTUBES FROM ETHYLENE DIFFUSION FLAMES

Liming Yuan and Kozo Saito*

Department of Mechanical Engineering, University of Kentucky, Lexington, KY 40506

Wenchong Hu

Department of Chemical and Material Engineering, University of Kentucky, Lexington, KY 40506

Zhi Chen

Department of Electric Engineering, University of Kentucky, KY 40506

ABSTRACT

Multi-walled carbon nanotubes (MWNTs) were synthesized from a laminar ethylene diffusion flame with a stainless

steel grid as the substrate. The grid was first oxidized using a premixed propane flame to generate a layer of metal

oxide particles (iron oxide, chromium oxide and nickel oxide) which could act as catalyst particles for the nanotube

growth. The as-grown nanotubes were entangled and curved with diameters ranging from 10 to 60nm. Carbon

nanofibers were also found; they might grow by thickening the nanotube walls. The maximum growth rate of

nanotubes was approximately 2-5gm/min and 3 mg/min. A nitrogen-diluted ethylene flame reduced the growth rate

of carbon nanofibers, probably by lower concentrations of pyrolyzed hydrocarbons due to a lowered flame

temperature. A cobalt-electrodeposited stainless steel grid produced vertically oriented, well-aligned and wellgraphitized carbon nanotubes consisting of each nanotube diameter 20nm and length 10 gm.

Keywords: ethylene diffusion flames, multi-walled carbon nanotubes, synthesis, growth mechanism.

INTRODUCTION

Carbon nanotubes are predicted to possess exceptional mechanical, electrical and thermal properties because of their

high strength, low density, and high electrical and thermal conductivity (ref. 1). Since their discovery in 1991 (ref.

2), much progress have been made on nanotube synthesis. Carbon nanotubes can be synthesized using arc-discharge

(ref. 3), laser vaporization (ref. 4), chemical vapor deposition (CVD) (ref. 5), doped diffusion flames (ref. 6), or

electrochemical synthesis (ref. 7). Here, we report a novel method for the synthesis of Multi-walled carbon

nanotubes using a simple unseeded ethylene diffusion flame.

Flame can be used for commercial synthesis of a broad range of materials such as carbon black. Although flame

synthesis of fullerenes was reported, the flame synthesis of CNTs has met with little success. Some carbon

nanotubes were found in the low pressure premixed hydrocarbon flames and diffusion flames (refs, 8 to 9, 6). But

the nanotube yield was too low to be found under SEM analysis. We present here a unique method to synthesize

carbon nanotubes without directly seeding materials into the flame. We used our previous experimental apparatus

and established a steady and stable laminar ethylene-air co-flow diffusion flame (ref. 10). The growth mechanism of

carbon nanotubes and the effect of nitrogen addition to the fuel stream on the nanotube growth rate were studied.

EXPERIMENTS

An unseeded laminar ethylene-air coflow diffusion flame was used to synthesis MWNTs. Figure 1 shows a

schematic diagram of the experimental apparatus. Ethylene (99.5% purity) was issued from a 1.1-cm diameter

stainless steel tube, which was surrounded by a 5 cm-diameter tube through which air flowed. With the average

linear fuel flow rate of 4.7cm/s and the average linear airflow rate of 63 cm/s, a steady and stable laminar flame with

a visible flame height of 33 mm was established on the burner port. All syntheses were performed at normal

atmospheric pressure. A stainless steel grid (type 304) was used as the substrate for the nanotube growth. The

material deposited on the wire was examined directly by SEM (Hitachi $900, 3kV). The deposit material was

dispersed in ethanol using a mild sonication. Drops of the dispersion were placed on the copper TEM grids and

dried for TEM analysis (JEOL JEM-2000-FX, 200kV). TEM-EDX (energy dispersive X-ray) analysis was

performed using an Oxford INCA detector attached to the TEM. A high-resolution TEM (JEOL 2010F, 200kV) was

also employed to study the microstructure of carbon nanotubes.

NASA/CP--2001-210948 810 August 6, 2001
CATHETOMETERr.rl A_S CHIMNEY

"_1_ S,,MPL,N_PROBE,._ II

'7 I QUARTZ BURNER _ t I I

1 [1_ GLASSBEA_ _1

FUEL THREE WAY 111[ [] SWITCHING VALVE

_ _ OO ® ............

N 2 _ El

FLOW METER

Figure 1. A schematic diagram of flame-synthesis experiment to produce carbon nanotubes.

RESULTS AND DISCUSSION

When the stainless steel grid was inserted into the flame 4mm above the exit port without any pre-treatment, and

stayed there for 10 minutes, some black material deposited on the grid. There were no nanotubes found under SEM

analysis. But some MWNTs were found under TEM analysis. Most of the deposit material was amorphous soot

particles. It indicates that the metal oxide (iron oxide and nickel oxide) particles formed on the metal surface were

not suitable as the catalyst particles for carbon nanotube growth. In our previous study, a Ni-Cr wire was placed in

the laminar methane diffusion flame for 10 minutes, the formed nickel oxide particles were favorable for the carbon

nanotube growth (ref. 11). Although the iron oxide was easily formed than nickel oxide, the morphology of iron

oxide layer on the metal surface was different from that of nickel oxide. When the Ni-Cr wire was placed in the

ethylene flame for 10 minutes, some carbon fibers loaded with amorphous carbon particles were produced,

indicating the oxidation process in the ethylene flame to be somewhat different from the methane flame.

To synthesis a pure form of carbon nanotubes, we pre-oxidized the stainless steel grid using a lean premixed

propane flame for about 5 seconds. SEM and TEM analyses on the pre-oxidized grid showed that a layer of several

different types of particles whose diameter range from 10 to 200nm was formed. The TEM-EDX spectra identified

these particles to be iron, chromium, nickel, oxygen and carbon. We believe that most of the metal elements exist as

metal oxide form, some as metal carbide. When the grid was placed in the ethylene flame for 10 minutes, formed

was a layer of dark gray deposit material that was identified by SEM as a bundle of entangled and curved MWNTs

with their diameter of 10-60nm and some carbon fibers with diameter larger than 100nm (Fig. 2). The length of

nanotubes is at least several gm and the growth rate of gray material was more than 3mg/min.

The Sampling Time

Sampling time was changed between 1 and 10 minutes to see whether or not the morphology of carbon nanotube

changes. With an increase of sampling time, the production rate of carbon fibers increased, probably due to a

continuous deposition of pyrolytic carbon onto the carbon nanotubes (ref. 12). We found that the length of carbon

nanotubes reached the maximum length 2-5 btm within less than 1 minute, while with an increase of the sampling

time beyond 1 minute, the thickness of nanotube walls increased but the length remained the same. We also found

after the deposition of pyrolytic carbon on the nanotube surface, the nanotube grew mainly in thickness, but not

length, and eventually became a solid nanofiber. With a 10-minute sampling time, the nanotube diameter increased

to 20-60nm, but the thickness remained the same probably due to the formation of pyrolytic carbon on the nanotube

surface, different from our previous methane flame synthesis, where pyrolysis products are stable and the pyrolytic

carbon concentration is much lower than ethylene flames.

NASA/CP--2001-210948 811 August 6, 2001
(a) (b)

Figure 2. SEM (a) and TEM (b) images of carbon nanotubes synthesized from an ethylene diffusion-flame

using a stainless-steel sampling grid coated with catalyst particles created by a premixed propane

flame (magnification is shown in each figure).

The Effect of Nitrogen Addition

Our ethylene flame was diluted with nitrogen, [C2H4]/[N2] = 0.25, to change temperature of flame. The stainless

steel grid was pre-oxidized with a premixed propane flame, and was placed for 10 minutes in the ethylene flame at

the same position where we obtained the Figure 2 result. Figure 3 show SEM and TEM images of the carbon

nanotubes that contain few carbon fibers. The flame temperature at the sampling location in this diluted ethylene

flame is 1244°C, while the corresponding temperature for an undiluted ethylene flame is 1547°C. Thus, the

observed significant reduction of carbon fibers may be due to lower temperature and lower fuel concentration

conditions. The lower temperature that can reduce the pyrolysis rate of hydrocarbon products in the flame may be

something to do with an increase of nanotube formation but suppress the growth of carbon fibers. The nitrogen

dilution effect may also have something to do with the creation of the same diameter nanotubes, but at this point we

do not know why.

Cobalt-Coated Grid

The oxidized stainless steel grid produced entangled and curved multi-walled carbon nanotubes. To obtain straightly

aligned and uniform diameter nanotubes, cobalt particles were electrodeposited on to the stainless steel surface by a

two-probe dc method in 5 wt. % COS047H20 and 2 wt. % H3BO 3 solution for 5 minutes. This cobalt-coated

grid was placed for 5 minutes at the same location as the Figure 2 and 3 results. Figure 4 presents SEM and TEM

images of well-aligned and nearly uniform diameter carbon nanotubes obtained by the cobalt-coated grid.

Interestingly, some nanotubes bend down (Fig. 4), while others are straight (Fig. 5), although both nanotubes

NASA/CP--2001-210948 812 August 6, 2001
wereproducedbythesamecobalt-coated gridatthesameflamelocation.Wenoticedspacingbetweeneach

nanotubeisdifferent: nanotubes intheFigure4aremorepackedthanthatintheFigure5,suggesting thathe

mechanical strengthforeachnanotubetoshootupstraight maybeweakenedbyfrictionbetweeneachnanotube.

TheGrowthMechanismofMWNTs

Thegrowthmechanismwasinvestigated byexaminetheSEM,TEMimagesandTEM-EDXspectra. IntheSEM

images,omeparticleswerefoundattachedtotheoneendofthenanotubes. Theseparticleswerenotwellseparated,

andwerecloselyconnected eachother.TheTEMimagesconfirmedthatveryfewparticleswerencapsulated inside

thenanotubes. Thenanotubes wereprobablygrownfromsomesmaller particles, andentangledachother.The

nanotubes seemedtogrowmainlybyarootgrowththatwasdescribedbyBaker(ref.14).TheTEM-EDS(energy

dispersionspectrum) showedthatheseparticlesweremainlyironoxide,somearenickeloxide,confirmingour

examination ontheoxidizedmetalsurfacebeforenanotubegrowth.Theseparticlesactedascatalyst particlesfor

nanotubegrowth.Theinteractionbetweenmetalsurfaceandmetaloxideparticlesseemsveryweak.SEMimages

showedthatsomeparticleblockwasdetachedfromthesurfacebecauseofthenanotubegrowth.Buttheinteraction

betweentheseparticlesmaybestrong, becauseevenaftermildsonication, someparticlesarestillconnectedeach

otherontheTEMimages.Experimental resultsshowedthathediameters ofcarbonnanotubes werewellcorrelated

withthediameters ofcatalyst particles(ref.15).Thesmallercatalyst particlesgenerated smallernanotubes.

Becausethemetaloxideparticleswerestronglybondedeachother,nocatalyst particlescouldbepushedupto

encapsulate nanotubes.

NASA/CP--2001-210948 813 August6,2001
• %-,-°-. ,.-°% °, ,., • ,_o,o%O ,%o ,%°°°°° o%°°%O_,o%°°o°°_o°° °°,*.%o.%%Oo°o°°°°o°.°O.Oo%O,% o,y

i:i:i:i:!:i:i:!:i:i:i:i:i:!:i:i:i:i:i_:i.i:!:i:i:i:i:i:!:!:i:i:i:!:i:i:i:i:i:i:i:i:i:!:i:i:i:i:!:i:i:i:

:.-...-.:°:..+-.;.',_,._._e_v :.:.:.:.:.:.:.-.:.-.:.:.:°-.;..÷:.:+.°:.:,:::.:°:+ :.:.:.:

.: -._.>>) :.:.:._°_._°:°:c.:..%%_e,-_,','_°:e.'°'A.:°?>:.:.'. :.:.:.:.:.:°_.;o;,,'.{°'.,',;,

!!:!:!ii:!ii:!ii:?!:ii!:?.`_:i:i:i:iii:!!!!ii!:_!i:!iii!iiii:i:!:i:i.!ii!i:i:i!!!!:!ii:i:i:i:i!!!!:i:!i

iiii!!!iiii_iiiii!iiiiii!iiii!i_!!!i!!!!!iii_!!iiii!iii_!iiii!!i!iiiiii_!!iii!!iiii!iiii!iiiiii!iiiii!

!:?!i ii_:!:!:!:!:?`ii_ii:!:i:i:i:i:i:!:i:i:!:i:i:i:i:i:!:i:i:!:!:i:!:i:_:!:i:i:i:i:i:i:!:i:!:i:!_:._::i;:

_%_..°°°_ .'.-....**...., .•. -., ,..,..- • _...,o.,,%-.,-,°°

(a) (b)

Figure 4. SEM (a) and TEM (b) images of carbon nanotubes synthesized from an ethylene diffusion-flame

using a cobalt-coated stainless-steel grid (magnification is shown in each figure).

Figure 5. Vertically oriented, well-aligned

nanotubes synthesized by a cobalt-coated stainless

steel grid.

Figure 6. HRTEM image of a typical nanotube

synthesized by a cobalt-coated stainless-steel grid

(Scale bar: 5nm).

NASA/CP--2001-210948 814 August 6, 2001
Forthenanotubegrowthoutofcobaltparticles, TEM-EDXspectraconfirmtheexistenceofcobaltandoxygen.

Someofcobaltmaybeoxidizedintheflame.Figure4(a)SEMimageofcarbonnanotubesynthesized fromcobalt

particlesshowsnoobviousparticlesorblockofparticlesattachedtothetubesthatwereobservedinthepropane

oxidizedgridsampling. Figure4(b)TEMimagealsodisplaysnocatalyst particlesencapsulated insidethe

nanotubes indicatingastrongbondbetweenthecobaltcatalyst particlesandthemetalsurface. Figure6highresolutionTEMimageshowsthathenanotubes synthesized fromourethyleneflamehavebeenwellgraphitized.

TheproductionrateofMWNTisverysensitivetothesamplinglocation.Withanincreaseofthesampling

locationabovethecurrent location, theconcentration ofpyrolyzedhydrocarbon productsincreased. Someofthe

pyrolyzedhydrocarbon productscanbedecomposed bycatalyst particles, butherestwillproduceamorphous

carbons, acontaminant forcatalyst particles.

CONCLUSIONS

Throughtisstudy,wefoundthatanunseeded laminarethylenediffusionflamecanbeanexcellent meansto

synthesize MWNTs.Withastainlessteelgridasthesubstrate, entanglednanotubes havingeachtubediameter 10-

60nmwereformed.Wefoundthemetaloxideparticlesformedduringoxidizationtohaveactedascatalyst particles.

Wealsofoundthatdilutionoftheflamewithnitrogenhelpedtoreducethepyrolyticcarbondepositionthe

nanotubes. Withacobalt-electrodeposited stainlessteelgrid,verticallyorientedwell-alignednanotubes consisting

ofanearlyuniformdiameter nanotubewereproduced.

ACKNOWLEDGMENTS

Thisstudywassponsored byNational ScienceFoundation undertheMRSECprogram.

REFERENCES

1. Ebbesen, T.W.: Carbon nanotubes: preparation and properties. CRC Press, Boca Raton, FL, 1997.

2. Iijima, S.: Helical microtubules of graphitic carbon. Nature, 354, 1991, 56-58.

3. Ebbesen, T.W. & Ajayan, RM.: Large scale synthesis of carbon nanotubes. Nature, 358, 1992, 220-222.

4. Guo, T., et al.: Catalyst growth of single-walled nanotubes by laser vaporization. Chemical Physics Letters,

243, 1995, 49.

5. Li, W.Z., et al.: Large synthesis of aligned carbon nanotubes. Science, 274, 1996, 1701.

6. Vander wal, R. L., et al.: Diffusion flame synthesis of single-walled carbon nanotubes. Chemical Physics

Letters, 323, 2000, 217.

7. Matveev, A.T., et al.: Synthesis of carbon nanotubes below room temperature. Carbon, 2001,137.

8. Howard, J.B., Das Chowdhury, K. & Vander Sande, J.B.: Carbon shells in flames. Nature, 370, 1994, 603.

9. Richter, H., et al.: Formation of nanotubes in low pressure hydrocarbon flames. Carbon, 34, 1996, 427.

10. Saito, K., et al.,: A study of the early history of soot formation in various hydrocarbon diffusion flames.

Combust. Sci. and Tech., 80, 1991, 103.

11. Yuan, L., et al.: Carbon nanotubes from unseeded methane flames. Chemical Physics letters,

2001 submitted.

12. Endo, M., et al.: Pyrolytic carbon nanotubes from vapor-grown carbon fibres. Carbon, 33, 1995, 873.

13. Cassell, A.M., et al.: Large scale CVD synthesis of single-walled carbon nanotubes, J. Phys. Chem. B. 103,

1999, 6484.

14. Baker, R.T.: Catalytic growth of carbon filaments. Carbon, 27, 1989, 315.

15. Li, W.Z., et al.: Controlled growth of carbon nanotubes on graphite foil by chemical vapor deposition.

Chemical Physics Letters, 335, 2001, 141.

NASA/CP--2001-210948 815 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

OPTICAL PLASMA CONTROL DURING ARC CARBON NANOTUBE GROWTH

I. Hinkov, S. Farhat*

LIMHP Universit_ Paris 13, Avenue J. B. Clement 93430 Villetaneuse, France

M. Lamy De La Chapelle, S. S. Fan

Nanostructures laboratory, Department of physics, Tsinghua University, Beijing 100084, China

H. X. Han, G. H. Li

National Laboratory for Superlattices and Microstructures, Institute of Semiconductors, Chinese Academy of

Sciences, Beijing 100083, China

and C. D. Scott

NASA, Lyndon B. Johnson Space Center EM2, Houston, TX 77058, USA

*Corresponding author: farhat@limhp.univ-parisl 3.fr - tel.: (33) 1 49 40 34 18 - fax: (33) 1 49 40 34 14

ABSTRACT

To improve nanotube production, we developed a novel optical control technique, based of the shape of

the visible plasma zone created between the anode and the cathode in the direct current (DC) arc process. For a

given inert gas, we adjust the anode to cathode distance (ACD) in order to obtain strong visible vortices around

the cathode. This enhance anode vaporization, which improve nanotubes formation. In light of our experimental

results, we focus our discussion on the relationship between plasma parameters and nanotube growth. Plasma

temperature control during arc process is achieved using argon, helium and their mixtures as a buffer gases. The

variation of the gas mixture from pure argon to pure helium changes plasma temperature. As a consequence, the

microscopic characteristics of nanotubes as diameter distribution is changed moving from smaller values for

argon to higher diameters for helium. We also observe a dependence of the macroscopic characteristics of the

final products as Brunauer-Emmett-Teller (BET) surface area.

Keywords: carbon, nanotubes, diameter, BET, control

INTRODUCTION

The interest on single-walled carbon nanotubes (SWCNTs) has been stimulated by the properties and

potential applications of these novel one-dimensional objects (ref. 1). Currently, there is an important challenge

to understand the growth mechanism and to control the structure of the produced nanotubes. For the arc process,

the plasma temperature distribution is an intimate component of SWCNTs synthesis (ref. 2). This parameter

controls species production and condensation, hence nanotube growth and morphology. To control this parameter

the majority of laboratory scale arc reactors use a water cooling system. Plasma temperature control is

accomplished by one of the following methods: 1) changing the plasma conditions by varying the electric current

and/or 2) varying the anode to cathode distance (ACD) to maintain a constant voltage. These methods do not

allow active control on the plasma temperature, independently of plasma parameters. An alternative approach

was developed in our recent work (ref. 3), which consists of varying the conductance of the gap between the

electrodes, using an argon-helium mixture as the inert background gas. Unfortunately, the yield of nanotubes was

found to be much higher with helium than with argon. To improve nanotube production independently of the

used inert gas, we developed an optical control technique based on the shape of the visible plasma zone. For a

given inert gas composition, we adjust the gap distance in order to obtain a strong visible vortices at the edge of

the cathode. These vortices are suspected to ensure a maximal mass flux of carbon and catalyst species from the

plasma zone to the cold cathode region, and to enhance the nanotube growth. In the present work we have

achieved a set of experiments with different argon-helium mixtures as inert gases. In addition, the morphology

and the structure of SWCNTs were analyzed using Raman spectroscopy to determine diameter distribution, and

Brunauer-Emmett-Teller (BET) surface analysis to estimate the surface area of as produced nanotubes.

NASA/CP--2001-210948 816 August 6, 2001
EXPERIMENTAL

SWCNTswerepreparedbyconventional electricarcdischargemethod.Thereactorconsistsofawatercooledreactionchamber, inwhichhasbeenintroducedananode, madeofpuregraphite: AlfaUltracarbonrod,6

mmin diameter and150mmlength.Thecathodeisalsomadeinhighpolishedgraphite. Nickelandyttrium,

wereintroducedascatalystsinadrilledholeinsidetheanodeof3mmdiameter and100mmlength.Finallythe

anode,includingtheoutergraphiteshell,consistsof94.8%atom.carbon,4.2%atom.nickeland1%atom.

yttrium.Inthepresent workwehaveachievedasetofsixexperiments withdifferent argon-heliummixturesused

asinertgasesundercontrolledpressure(ref.3).Optimumnanotubeyieldwasfoundatapressureof660mbarfor

pureheliumand100mbarforpureargon.Varyingthepercentage ofheliumfrom0to 100withastepof20%,

thepressurewasarbitrarilyfittedusingalinearfitbetween100mbarand660mbar.Thethermal conductivityof

themixturedependsontherelativeconcentration ofthetwogases.Thethermal conductivityofargonisabout

eighttimessmallerthanof thehelium.Varyingtherelativeconcentration of thetwogasesinfluencesthe

temperaturegradient betweentheplasmaandthecathode. Thisaffectsthecondensation ofatomicarbonand

metalcatalysts, hencethenanotubediameter distribution.

Theelectricparameters ofalltheexperiments weremaintainedatadirectcurrentof100Atocreatethe

arcdischarge. Inthebeginningtheplasmais ignitedbycontactbetweentheelectrodes, whichincreasesthe

temperature ofthecontact pointuntilevaporationftheanodematerial. Whenthedischargeruns,theplasmais

observedthroughthequartzwindowofthereactionchamber, andamagnifiedimageoftheluminescent plasma

isprojectedonascreen, usinglens.Theanodeismovedbackuntilvorticesappearbetweenthegrowingdeposit

onthecathodeandtheburninganode.Thegapdistancebetweentheelectrodesismaintainedbymovingthe

anodetowardthecathode. Thehightemperatureneartheanodeandthehighenergydensitiesin theplasma

ensureatotalvaporizationftheanodematerial, whilethewatercooledcathodeladstohighquenchratesand

tonanotubeformation. Theexperimental resultsindicatehathecarbonnanotubes aregeneratedonthefront-end

surfaceofthedeposit. Thequenchprocessicomplexanduncontrolledbutweusuallyobtainvariousproducts:

sootonthereactorwalls,web-likestructuresbetweenthecathodeandthechamberwalls,a depositonthe

cathode'sedge,andarubber-likecollaretaroundthisdeposit.

Pictures of the projected plasma have been taken with a video camera and using the recorded film we

can determine the distance between the electrodes in every moment of the experience. A data acquisition system,

composed of a power and harmonics clamp and a multimeter connected to a PC, allows recording the variations

respectively of the intensity of the current, and of the difference of the potential. Nanotube diameter was

determined using Raman spectroscopy using spectrophotometer Jobin Yvon T64000. Spectra have been recorded

at room temperature in ambient air with an argon laser at 514.5 nm. Surface analyses was accomplished with a

BET COULTER SA 3100 instrument using nitrogen gas in order to determine the BET surface areas as well as

the pore size distributions. The method consists of one-hour gas evacuation at 393 K and of determination of the

pore volume introducing helium gas, followed by adsorption of nitrogen at temperature of liquid nitrogen (77 K).

RESULTS AND DISCUSSION

Plasma characteristics

In the present work, we have achieved the set of six experiments with different argon-helium mixtures as

discussed above. The plasma characteristics are reported here only for pure argon plasma but are reproducible for

all the experiments. The conditions of the synthesis of nanotubes namely ACD were modified in order to improve

the yield of nanotubes. This parameter was not maintained constant as suggested in refs. 4,5, but varied to

establish a visible vortices around the cathode. These vortices are visible probably because of the presence of

small particles of carbon acting as markers in the stream.

During all the experiments, we distinguished qualitatively two operative regimes concerning the shape

and the color of the plasma.

1) The first is the stable regime, during which the arc flame is uniform, in a blue-green color, and does

not change its shape and brightness (figure la, lb). We adjust manually the distance between

electrodes, until the appearance of the vortices (figure lc, 1d). We found in our geometry that when

we use argon as inert gas, the optimal distance between the electrodes, encouraging the formation of

vortices is between 15 to 17 ram.

NASA/CP--2001-210948 817 August 6, 2001
2) The second regime is not desirable, the vortices do not appear; the color of the arc flame is bluepurple and the brightness is weaker (figure 2a, 2b). This regime is unfavorable for the synthesis of

nanotubes. It appears when the anode is very near (ACD 1 to 7 mm) or when it is very far from the

cathode, (ACD greater than 18 mm). If we continue to move away the anode from the cathode the

arc extinct.

Figure 1. Development of vortices in argon plasma (Pressure = 100 mbar, Current = 100 A). The

streamlines are curved outside the plane of symmetry.

Figure 2. Pictures of the non-desirable regime. The vortices do not appear

Finally, the effect of the control of the ACD on the yield of nanotubes is reported on figure 3. From this

figure we can conclude that optimal ACD giving vortices improve significantly the yield of collaret, compared

with ACD giving a constant voltage (U 40 V). For some mixtures namely 0 % and 40 % argon, the deposition

of webs on the quartz window makes ACD control difficult and explain the decrease of the mass of collaret. For

pure argon the yield is improved by a factor 3.5.

NASA/CP--2001-210948 818 August 6, 2001
Ourfirstexperimental resultsdemonstrate clearlytheefficiencyofACDcontroltechniquefordifferent

inertatmospheres. Thisnewdesignprovidesameansto improvenanotubeyieldbutalsoto controlgradient

temperaturein thearcindependently of plasmaconditionsi.e. current and voltage. This increases the

repeatability of experiments for various process conditions. It also provides the opportunity to control the

morphology of nanotubes. Indeed, since the thermal conductivity of the argon-helium mixture is affected by the

relative concentration of the two gases, we expect that control of thermal transfer and hence condensation of

atomic carbon and metals between the plasma and the vicinity of the cathode can control nanotube diameter in

arc process. This result implies that single-layer tubules nucleate and grow on a metal particle with different sizes

depending on the quenching rate in the plasma and suggests that temperature and carbon and metal densities

affect the diameter distribution.

1000

800

g

600

O

o 400

200

nConstant ACD (3 mm) mOpt ma (ACD)

0

0 20 40 60 80 100

% Ar

Figure 3. Effect of the anode to cathode distance on collaret mass collected in the cathode using following

parameters: burned anode length 100 mm and current I=100 A. [] ACD was maintained constant at 3 mm.

• ACD was varied to obtain vortices.

Microscopic characteristics: nanotube diameter

For each sample, we have recorded spectra on at least ten different points to get the best view of all the

sample. Two main regions of the Raman spectrum can be attributed to SWCNTs: 1) the high frequency (1500-

1600 cm -1) and 2) the low frequency (100-250 cm -1) ranges. In figure 4a, we can see the overall spectrum and in

figure 4b, the low frequency range has been scaled up to better observe the breathing modes ofnanotubes.

v __.A

¢-

¢g

E

100 %

8O %

60 %

4O %

2O %

)' 0 %

'2(?0 '4(?0 ' 6(?0' 8(?0 '10'00'12'00'14'00'1600'18'00 '

Raman shift (cm -1)

b

7 II

n1;o 1;o

_/o

_______8o %

<____%

20 %

_%

2;0 2;0

Raman shift (cm -1)

Figure 4. (a) Overall Raman spectrum and (b) Radial breathing mode of the SWNT produced with

various concentration of Ar in the inert atmosphere. The percentage of Ar is indicated on the rigth side.

NASA/CP--2001-210948 819 August 6, 2001
Therelativeintensityofthesebandschangesatheinertgasvariesfrompureheliumtopureargon.To

determineexactlythevariationinthediameter distribution, wehavefittedseveral lorentzianshapes(oneshape

perband)tothewholegroupofpeaksassignedtothebreathingmode.Inthisway,wehavebeenabletogetthe

proportionofeachdiameter forallrecordedspectra. Wecanthencalculateanaveragediameter distributionand

anaveragediameter fromallspectraobtainedonthesamesample. Weassumethathediameter distributionand

averagediameter dependontheinertatmosphere condition. Thisdependence isgiveninthefigure5.Varyingthe

gasfromheliumtoargonchangesaveragediameter tosmallervalues.

1,4

1,35

E

o

1,3

c

c

o

1,25

Figure 5. Variation of the average nanotube diameter as determined by Raman spectroscopy with the

argon percentage in the arc chamber obtained with I=100 A and optimal ACD giving vortices around the

cathode.

Macroscopic characteristics: BET surface area

The BET analysis of the as-produced nanotubes was achieved during this work and permits to measure

the BET surface area. On the figure 6, we report the variation of the BET surface area of the collected collaret

with the argon mole fraction in the inert gas obtained with the optimal ACD discussed above. On the figure 7 we

report the corresponding isotherms of nitrogen adsorption on nanotubes. BET surface areas varies in the range of

[200 450 m2/g] and are less than the outer surface area of single nanotube (1300 m2/g) as given by Ye ef al.

(ref. 6). This is a proof that the measured BET surface area is attributed to the outer surface of a nanotubes

organized in bundles, or "ropes" of SWCNTs as discussed in ref. 6. Hence, the parameter controlling BET

surface area must be the rope diameter. In spite the fact that the relationship between individual nanotube

diameter and rope diameter is not yet fully understood, we can conclude that the macroscopic surface area of

nanotubes is controlled by the argon-helium mixture gases. In our knowledge, the maximal value of BET surface

area published in the literature never exceed 350 m2/g. Further High Resolution Transmission Electron

Microscopy (HRTEM) observations of our samples are needed to determine the number of individual nanotubes

per rope and to better understand this result. In the present work, the maximal value of BET surface area is found

with mixture of 80 % helium and 20 % argon, and it decrease increasing the argon percentage. In the most

favorable case the BET surface is 440 mTg, which is bigger than the values reported in the ref. 6 (285 m2/g).

500

450

400

g

350

300

250

200

150

100

0 20 40 60 80 100

% Argon

Figure 6. Variation of the BET surface area of the collaret with the argon mole fraction.

NASA/CP--2001-210948 820 August 6, 2001
_0

x_

.<

=

400 .........................-_--.....................r.........................T.........................1..........................

100%A]

350 _ 80%Ar

300 _ 60%Ar j

40 %Ar

250 _ 20%Ar

0%Ar

200 t i t

0

0 0,2 0,4 0,6 0,8

Relative Pressure Ps/Po

Figure 7. Adsorption isotherms of nitrogen at 77 K on as prepared nanotubes at different argon

mole fractions.

CONCLUSION

The mechanism of plasma control with a variable ACD and a variable thermal conductivity of the

plasma is demonstrated in this paper and validated in DC arc process for specific SWCNTs synthesis. More

precisely, it was demonstrated that varying the arc plasma parameters and basically the distance between the

electrodes permits to work in optimal regime, with a strong vortices around the cathode, which increase

significantly the nanotube yield. This suggests that controlling the argon-helium mixture and/or the ACD provide

the opportunity to control simultaneously:

1) microscopic nanotube parameters as nanotube diameter in the range of 1.27 to 1.37 nm

2) macroscopic nanotube parameters as BET surface area in the range of 200 to 450 m2/g.

This opens new horizons to investigate hydrogen adsorption in SWCNTs of different diameters.

REFERENCES

1. Iijima, S.: "Helical microtubules ofgraphitic carbon", Nature, vol. 354, Nov. 1991, p. 56-58.

2. Takizawa, M., et al.: "Effect of temperature for Synthesizing Single-Wall Carbon Nanotubes by Arc

Vaporization Method", Chem. Phys. Lett., vol. 302, no.l-2, March 1999, p.146-150.

3. Farhat, S., et al.: "Diameter Control of Single-Walled Carbon Nanotubes using Argon-Helium Mixture

Gases", submitted in J. Chem. Phys. 2001.

4. Gamaly, E. G.; and Ebbesen, T. W.: "Mechanism of carbon nanotube formation in the arc discharge", Phys.

Rev. B vol. 52, no.3, Jul. 1995, p. 2083-2089.

5. Shi, Z., et al.: "Mass-production of single-wall carbon nanotubes by arc discharge method", Carbon vol. 37,

1999, p. 1449-1453

6. Ye, Y., et al.: "Hydrogen adsorption and cohesive energy of single-walled carbon nanotubes", Appl. Phys.

Lett. Vol. 74, no. 16, APR. 1999, p. 2307-2309.

NASA/CP--2001-210948 821 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

Boundary-layer control during carbon nanotube growth by arc method

Single-walled carbon nanotubes (SWCNTs) have interesting properties and potential applications [ 1].

Many experimental and theoretical works have been done to provide an understanding of the growth

mechanism and to optimize the process. Several laboratory-scale methods have been proposed including arc

[2,3] and laser [4,5] condensation of vaporized carbon in the presence of catalyst. In the present work,

nanotubes were produced using are-evaporation method in helium gas ambient under controlled pressure. The

reactor consists of a water-e0oled reaction chamber with two cooled graphite rods. For the anode, a mixture of

graphite with nickel and yttrium acting as catalyst is used. The cathode is made in pure graphite. A direct

current of 100 A driven by 40 volts passes through the electrodes and a plasma is created in the inter-electrode

region. The gap between the electrodes is controlled manually by moving the anode toward the cathode and by

projecting a magnified image of the luminescent plasma on a screen using a lens. This image is used to in situ

characterize and control flow by moving the anode as nanotubes grow. The light can be also spectroscopicallyanalyzed to determine populations of species and temperature.

The plasma is firstly ignited by a contact between the anode and the cathode, which increases the

temperature in the contact point until evaporation of the anode material. Then, the anode is moved back to

maintain a desired gap between growing deposit on the cathode and burning anode. An active plasma zone

bounded by the deposit and the anode is created. The function of this zone is to produce optimally carbon and

catalyst vapors which diffuse then to the cooled reactor regions. The carbon species are deposited as cellaret in

the vicinity of the cathode deposit and soot on the reactor walls. The high temperature near the anode and the

high energy densities in the plasma ensure a total vaporization of the anode material, while"the water cooled

cathode leads to high quench rates and nanotubes formation. The quench process is complex and uncontrolled

but we usually obtain various products: soot on the reactor walls, web-like structures between the cathode and

the chamber walls, a deposit on the cathode's edge, and a rubber-like cellaret around this deposit. We only

focused this study on the cellaret which is known to contain the highest density of SWCNTs [3] and we need to

maximize the amount of this cellaret.

First, we experimentally observe a macroscopic dependency of the amount of cellaret collected and the

anode to cathode distance. Then, we analyze the flow under the exact conditions of nanotube growth. At high

gap of about 6 mm corresponding to the anode diameter, visible parallel streamlines traced by blackbody

radiation emitted from hot carbon particles are present in the plasma. When decreasing this gap to about 3 ram,

streamlines are curved and vortices appears in the cathode region. At smaller gap values, deposit temperature is

so high that cathode deposit material is evaporated. We found that the optimal regime for nanotube production

in the arc process is not dictated by the maintenance of a constant voltage as reported in [3], but the

maintenance of a vortices close to the cathode. This is achieved by adjusting anode to cathode distance during

nanotube production. A suitable shape of the pIasma corresponds to the creation era strong visible vortices

which ensure a maximal time of stay of carbon species and catalyst particles in cold region and then enhance

nanotube production.

REFERENCES

[1] Iijima, S., "Helical microtubules of graphitic carbon", Nature, 354 p. 56-58, (1991).

[2] Bethune, D.S., C. H. Kiang, M. S. DeVries, G. German, R. Savoy, J. Vazquez and R. Beyers, 1993.

"Cobalt-catalysed growth of carbon nanotubes with single-atomic-layer walls" Nature, 363 (6430) p. 605-7,

(1993).

[3] Journet, C., Maser, W.K., P. Bernier, A. Loiseau, M. L. Delachapelle, S. Lefrant, P. Deniard, R. Lee, and J.

E. Fischer, "Large-Scale Production of Single-Walled Carbon Nanotubes By the Electric-Are Technique",

Nature, 388 (6644)p. 756-758,(1997).

[4] Guo, T. and R. E. Smalley, "Production of single-walled carbon nanotubes via laser vaporization

technique", in Recent Advances in the Chemistry and Physics of Fullerenes and Related Materials.

Electrochem. Soc, Pennington, N J, USA. Reno, Nevada, (1995).

[5] Thess, A., R. Lee, P. Nikolaev, H. J. Dai, P. Petit, J. Robert, C. H. Xu, T. H. Lee, S. G. Kim, A. G. Rinzler,

D. T. Colbert, G. E. Scuseria, D. Tomanek, J. E. Fischer, and R. E. Smalley, "Crystalline Ropes of Metallic

Carbon Nanotubes", Science, 273 (5274) p. 483-487, (1996).

e-mail: farhat@limhp.univ-parisl 3.fr

NASA/CP--2001-210948 822 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

GROWTH OF ALIGNED CARBON NANOTUBES ON SILICON BY RF-PLASMA

ASSISTED PULSED-LASER DEPOSITION

Y. K. Yap*, M. Yamaoka, M. Yoshimura, Y. Mori, T. Sasaki

Department of Electrical Engineering

Osaka University, 2-1 Yamadaoka, Suita, Osaka 565-0871, Japan.

Tel: +816-6879-7707 Fax: +816-6879-7708

Email: yap @ ssk.pwr.eng.osaka-u.ac.jp

To Hanada

Institute of Science and Industrial Research

Osaka University, 8-1 Mihogaoka, Ibaraki, Osaka 565-0847, Japan

H. Furuta, To Hirao

Department of Electrical Engineering

Osaka University, 2-1 Yamadaoka, Suita, Osaka 565-0871, Japan.

S. Honda and K. Oura

Department of Electronics Engineering

Osaka University, 2-1 Yamadaoka, Suita, Osaka 565-0871, Japan.

ABSTRACT

We present the first report on growing vertically aligned carbon nanotubes (CNTs) on Si substrates by means of

a physical vapor deposition technique: RF-plasma assisted pulsed-laser deposition (PLD). Our CNTs are grown at

800o by using Fe nano-particles as the catalysts. Two RF-plasmas are used with the first one applied on a ring

electrode placed between the graphite target and substrates. The second RF-plasma coupled on the substrate holder

initiates negative dc self-bias voltage on the substrates. Hydrogen is used as the working gas. These CNTs are multiwalled, 200-400 nm in length and about 40 to 80 nm in diameter. The density of such CNTs is estimated as >1 x 108

cm -2. Some CNTs are as long as 1 gm in length. According to high-resolution transmission electron microscope

(HRTEM), these tubes are constructed of high-quality graphitic structure. We speculate that CNTs with diameter of

4 • are confined inside our multi-walled CNTs.

Keywords: carbon nanotubes; catalyst; Fe, RF-plasma; pulsed-laser deposition

INTRODUCTION

Large scale, vertically aligned multi-walled carbon nanotubes (CNTs) can be grown on various substrates by

chemical vapor deposition (CVD) techniques like thermal CVD (ref. 1), plasma enhanced hot filament CVD (refs. 2

to 3) and plasma CVD (ref. 4). These CNTs are promising for the application of electron filed emission like coldcathode flat panel display. However, most of these CNTs are more properly called carbon nanofibers (refs. 2 to 3 and

5) because of the absent of high order graphitic crystalline. Distorted of graphitic structures can happened on the

whole tube or at least near the outermost walls of the tubular structures. Likewise, both single wall and multi-walled

CNTs can be grown by physical evaporation techniques like arc discharge (refs. 6 to 8) and laser ablation (ref. 9).

These CNTs are constructed from high-quality graphitic sheets, free standing but is mixed of carbon impurities. It is

desired to understand way to growth high quality, vertically aligned multi-walled CNTs on substrates. On the other

hand, techniques to growth large scale, single wall CNTs on substrates are important if single wall CNTs are to be

used for the purpose of nano-devices. So, new synthesis routes of CNTs that combined abilities of CVD and physical

evaporation techniques are desired.

We attempt to grow vertically aligned multi-walled CNTs by a physical vapor deposition (PVD) technique. The

foremost step to begin such an attempt is trying to simulate the synthesis condition of PVD to that of CVD for

growing aligned multi-walled CNTs. In this work, we report on the growth of aligned CNTs on Si(100) substrates by

NASA/CP--2001-210948 823 August 6, 2001
RF-plasma assisted pulsed laser deposition (PLD). Details on the synthesis procedure and the structure of these

CNTs are discussed.

EXPERIMENTAL DETAILS

Deposition system

Our deposition system consist of a solid-state UV laser (5c0, fifth harmonic generation of Nd:YAG lasers at a

wavelenghth of 213 nm) with pulse duration of 3 ns (ref. 10). Because of the short pulse duration, this UV radiation

can provide 5-10 times higher intensity than excimer lasers (KrF, ArF) when identical energy density is applied. The

schematic diagram of RF-plasma assisted PLD system is shown in Figure 1. As shown, two RF-plasmas are

applicable to assist deposition. The RF plasma coupled to a Mo substrate holder induced negative self-bias voltage

on the substrate (Vb) and initiated ion bombardment on the growth surface. We choose to use slow deposition rate for

the growth of CNTs by focusing the UV light on the graphite target (99.99 at. %) with an intensity of 0.27 GWcm -2.

Similar approach was used to prepare carbon nitride films with predominantly sp3C-N bonds (ref. 11). The carbon

plume generated from such ablation is the only carbon source for the deposition on the treated Si substrates attached

on a heater 4 cm away from the target.

For the synthesis of CNTs, we installed another RF-plasma on a ring shape Mo electrode located in between the

target and substrates. This plasma induced a negative self-bias voltage on the electrode and is referred as V_ing

hereafter. Our CNTs were grown at 800 • with a fixed bias voltage on the electrode, V_ing = -500V. The ambience gas

pressure was fixed at 5x10 -1 mbar. Because of the low ablation rate, all samples were deposited for about 8 hours.

Nonlinear optical

CsLi_010 crystals

 PellinBroom Prism

_2_321!!!i_or / _Beam dump

. _to ist RF generat

(04g radiat i_ _

ubstrates

m5generWcor

\' fl!!! - Ill

<ggN<

to 2nd R:

%o

generato:

Figure 1. Experimental setup of the RF-plasma assisted PLD system.

NASA/CP--2001-210948 824 August 6, 2001
Substrate pretreatments

Thegrowthofverticallyalignedmulti-walled CNTsrequirednano-size transitionmetalcatalysts. Weuse

commerciallyvailableFenanopowder (20nm.). This powder was coated on Si substrate by alcohol suspension

method and dry in air at RT. The treated substrates were than mounted into the vacuum chamber and treated with H2

RF-plasma and Vb = - 0 V, V_ing= -500V for about 15 minutes at 800 .. After the plasma treatment, laser ablation

was started at 800 .under various Vb. During the plasma treatment, the Fe nano-particles will rearranged themselves

uniformly into particles of < 100 nm• as confirmed by field emission scanning electron microscope (FESEM) under

the controlled experiments.

RESULTS AND DISCUSSION

The formation of CNTs depends strongly on the laser intensity, configuration of RF-plasma, substrate bias

voltage and the synthesis gas ambience. Details of these results are to be discussed elsewhere. In short, substrate bias

voltage is indispensable. Next, the deposition condition must not allow formation of amorphous carbon film. Also,

H2 plasmas are necessary for the formation of high-quality graphitic structure. All these are crucial for the formation

of CNTs instead of carbon nanofibers with poor crystalline structure.

For example, at Vb = - 100 V and V_ing = -500V, formation of CNTs is detected by FESEM as shown in Figure 2.

From such a magnified view tilted at 45., these CNTs are 200-400 nm in length and about 40-80 nm in diameter. As

shown, within an area of four square micrometers, more than four CNTs can be observed through out the whole

surface of the sample. The density of such CNTs is thus estimated as > lxl08 cm -2. Such a density is at similar order

of magnitude to those typically obtained by CVD techniques. Some CNTs on our sample are as long as 1 btm.

Figure 2. Vertically aligned carbon nanotubes on Si substrates.

NASA/CP--2001-210948 825 August 6, 2001
Structural property of our CNTs was examined by High-resolution transmission electron microscopy (HRTEM)

using a FEG-TEM (JEOL 3000F) operating at 300 keV. CNTs sample was prepared by scratching CNTs from the Si

substrate on to a microgrid. HRTEM indicates that we have grown multi-walled CNTs with high-quality graphitic

sheets. Figure 3 shows a magnified view at the outermost walls of the tubes. As shown, structural distortions are

rarely observed. We think that this might attributed to the high density of atomic hydrogen in our synthesis ambience.

Low ablation rate of graphite and highly reactive H2 RF-plasma as generated by the ring electrodes are crucial. Note

that no CNTs can be grown when substrate bias voltage is not applied (Vbias = 0).

Figure 3. HRTEM image of carbon nanotubes at the outermost graphitic walls.

According to our result, the whole volume of our CNTs are filled of highly-parallel graphitic planes with regular

inter-plan spacing of 3.4.. The possible explanation is that a CNT with diameter of 4. is confined at the center of the

multi-walled CNTs. Further analysis is required to confirm our speculation. Note that CNT with diameter of 4. was

recently produced by arc discharge of graphite in H2 atmosphere (ref. 12).

CONCLUSION

As a conclusion, we have presented the first report on growing vertically aligned carbon nanotubes (CNTs) on

Si substrates by means of a physical vapor deposition technique: RF-plasma assisted pulsed-laser deposition (PLD).

These CNTs are multi-walled, 200-400 nm in length and about 40 to 80 nm in diameter. The density of such CNTs is

estimated as >1 x 108 cm -2. Some CNTs are as long as 1 gm in length. According to high-resolution transmission

electron microscope (HRTEM), these tubes are constructed of high-quality graphite structure. We suspect that CNTs

with diameter of 4 • are confined inside our multi-walled CNTs. Reactive H2 ambience and substrate negative bias

voltage are important for the formation of CNTs with good crystalline property or else, carbon nanofibers are formed.

NASA/CP--2001-210948 826 August 6, 2001
REFERENCES

1. Li W. Z., et al.: Large-scale synthesis of aligned carbon nanotubes. Science 274, Dec. 1996, 1701-1703.

2. Ren. Z. F., et al.: Synthesis of large arrays of well aligned carbon nanotubes on glass. Science 281, Nov.

1998, 1105-1107.

3. Huang Z. R, et al.: Growth of highly oriented carbon nanotubes by plasma-enhanced hot filament chemical

vapor deposition. Appl. Phys. Letts. 73, Dec. 1998, 3845-3847.

4. Tsai. S. H., et al.: Bias-enhanced nucleation and growth of the aligned carbon nanotubes with open ends

under microwave plasma synthesis. Appl. Phys. Letts. 74, June 1999, 3462-3464.

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NASA/CP--2001-210948 827 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

DIRECT SYNTHESIS OF CARBON NANOTUBES ON PLASTIC SUBSTRATES

Eun Hwa Hong, Kun-Hong Lee"

Department of Chemical Engineering, Electrical and Computer Engineering Division, Pohang University of Science

and Engineering, San 31, Hyoja-Dong, Nam-Ku, Pohang, Kyungbuk, 790-784, South Korea

r Sang Ho Oh, Chan-Gyung Park

Department of Materials Science and Engineering, Pohang University of Science and Engineering, San 31, HyojaDong, Nam-Ku, Pohang, Kyungbuk, 790-784, South Korea

Chang-Mo Ryu

Department of Physics, Pohang University of Science and Engineering, San 31, Hyoja-Dong, Nam-Ku, Pohang,

Kyungbuk, 790-784, South Korea

Jong Hun Hart, Jae Eun Ryu

Iljin Nanotech Co., Ltd., Nano Technology Center(R-502), Kayang Techno Town 1487, Kayang-Dong, Kangseo-Ku,

Seoul 157-810 (South Korea)

ABSTRACT

Since the discovery of carbon nanotubes(CNTs), various methods of the synthesis of CNTs have been reported :

arc-discharge, laser ablation, chemical vapor deposition, flame synthesis and Smalley's recent invention of high

pressure CO(HIPCO) process. None of them is able to synthesize CNTs at low temperature (say, less than 200°C), so

that the incorporation of CNTs with low melting point materials such as organic polymers has been severely limited.

Here, we report a novel method for CNT synthesis by microwave irradiation. Carbon nanotubes were

successfully synthesized on various supports such as carbon black, silica powder, a sheet of sodalime glass, and

finally organic pot)aner substratesO_eflon and polycarbonates) [_1.In our method, microwave was directly irradiated

on the catalyst particles which were impregnated or painted on various substrates. Since the microwave energy is

selectively absorbed by the catalysts, not by the substrates, local heating occurs, resulting in the CNT synthesis even

on the organic polymer substrates, and degradation of the substrate material is minimum.

Microwave(2.45GHz, 800W) was irradiated using acetylene as a hydrocarbon source and cobalt naphtenate and

cobalt sulfide as the catalysts loaded on the substrates. We have observed temporary arcing phenomenon during

microwave irradiation and the substrate was covered with CNTs without any damage to its appearance. The

morphology and structure of carbon nanotubes were observed by using FE-SEM (field emission scanning electron

microscope) and HR-TEM (high resolution transmission electron microscope). Various shapes of carbon nanotubes

and nanoparticles were observed as well as with amorphous carbons, and CNTs produced in this way were short,

straight, and multi-walled. They varied depending upon the catalyst phases and reaction conditions.

It is an important merit of our method which make it possible to grow CNTs virtually any substrate provided its

absorption of microwave energy is small. This novel synthesis technique is expected to be applicable to the

fabrication of the flexible field emission displays (FEDs).

References

[1] Eun Hwa Hong, Kun-Hong Lee, Chang-Mo Ryu, Jong Hun Han and Iae Eun Ryu, "Method of synthesizing

Carbon Nanotubes and apparatus being used therein", Korea Patent Application No. 10-2000-55829 (2000. 9. 22.)

Keywords: carbon nanotubes, synthesis, substmte, microwave

Corresponding Author : Kun-Hong Lee"

Department of Chemical Engineering, Pohang University of Science and Engineering, San 31, Hyoja-Dong,

Nam-Ku, Pohang, Kyungbuk, 790-784, South Korea

e-mail : .c.e20.O47._t_s.te.c...h..ac._

Fax: (82)54-279-8298

Tel: (82)54-279-2271

NASA/CP--2001-210948 828 August 6, 2001
A NOVEL APPROACH TO GROWTH OF LARGE SCALE ALIGNED CARBON

NANOTUBES OR CARBON NANOSHEETS ON SI SUBSTRATE BY ECR CHEMICAL

VAPOR DEPOSITION

Chao Hsun Lin, FIui Lin Chang, Ming Her Tsai and Cheng Tzu Kuo

Department of Materials Science and Engineering, National Chiao Tung University, Hsinchu 300, Taiwan

E-mail; ctkuo@ee.nctu,edu.tw

ABSTRACT

A novel microwave ECR CVD process to grow large scale (4 x 4 cm 2) aligned carbon nanotubes or seaweedlike "carbon nano-sheets" on p-type Si wafer was developed with CI-h and/or H2 as source gases. Before nanotube

deposition, the Si substrate was etched by HF solution and then dispersed with a layer of catalyst, including Fe, Co,

Ni or Pd. The deposited surfaces at different deposition stages were characterized by SEM, TEM, Ramaa and CL

spectroscopy. The field emission properties were determined by I-V measurements. Effects of processing parameters

including various pretreatments on growth rate, nanotube morphology, crystal structure, formation mechanism and

field emission properties were examined. The preliminary results show that the well-aligned nanotubes or

nanosheets can be obtained. A unique feature of the seaweed-like nanosheets of few microns in length has not been

reported in the literature. Effect of ECR on nanotubes and nanosheets formation will be discussed.

Keywords: Carbon nanotubes, Carbon nanosheets, ECRCVD, Catalyst, Field emission

Corresponding author information

Full name: Chao Hsun Lin

Affiliation: Department of Materials Science and Engineering

National Chiao Tung University

Hsinchu 300, Taiwan

Address: 195-5 Chung Hsing Rd., Section 4

Chutung, Hsinchu 310, Taiwan

Phone Number (country code +): +886-3-5914 I74

Fax Number (country code +): +886-3-5820207

E-mail address: 800589@itri.org.tw

NASA/CP--2001-210948 829 August 6, 2001
Proceedings of the Sixth Applied Diamond Conference/Second

Frontier Carbon Technology Joint Conference (ADC/FCT 2001)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga (Editors)

RADIAL AND VERTICAL VAPOR GROWN CARBON FIBERS (VGCF)

Maheshwar Sharon, Debabrata Pradhan

Department of Chemistry, Indian Institute of Technology, Bombay 400 076, India.

Different type of carbon fibers are grown over the alumina, quartz and

oxidized silicon substrate taking camphor as a precursor material. The

vaporized camphoric gas is pyrolysed in wide range of temperature from

700°C to 1100°C in different catalytic medium and different atmosphere. In

some cases fibers are found to grow horizontally starting from a point to

outward like spoke in the bicycle wheel but in other cases it protrudes

towards image analyzer. This paper will present a very simple experimental

method of production and a detail comparative study of the effect of

substrate, temperature, catalyst and medium of reaction along with a

proposed mechanism. All analysis has been done by SEM, TEM and EDAX.

This is the first time we obtained Vapor Grown Carbon Fibers using a new

source camphor. Camphor, a volatile material found much easier to vaporize

and get deposition.

NASA/CP--2001-210948 830 August 6, 2001
AUTHOR INDEX

Abraham, T. 44

Adamschik, M. 90

Ago, H. 651

Ahn, J. 473, 621

Akagi, K. 523,528

Akhvlediani, R. 120

Aksenov, I.I. 509

Albergo, S. 78

Albin, S. 170, 275, 718,726

Alexenko, A.E. 284

Alexenskii, A.E. 264

Amaratunga, G.A.J. 794

An, K.H. 33, 669, 675, 695,758,759

Anderson, A.B. 164

Ando, Y. 110, 323

Andrews, R. 701,771

Angelone, M. 80

Angus, J.C. 164

Aoi, Y. 609

Arepalli, S. 769

Arnault, J.C. 263,551

Arumugam, R 210

Ashfold, M.N.R. 291

Au, F.C.K. 717

Bae, D.J. 675,695, 758,759

Bae, E.-J. 267

Bagayoko, D. 544

Baidakova, M.V. 264

Baik, Y.-J. 304

Baker, RA. 452

Banhart, F. 551

Baraty, N.J. 61

Batzer, S. 210

Beghi, M.G. 521

Bellini, V. 78

Bello, I. 262, 552, 565, 717

Belous, V.A. 327

Bergosh, R.G. 771

Billups, W.E. 687

Blumenthal, R. 349

Bobrovnitchii, G. 245,590

Bochko, A. 590

Booth, R.A. 180

Borzdov, Y.M. 171

Bottani, C.E. 521

Breidt, D. 198

Brown, B. 126, 132, 133

Brown, W.D. 589

Bulusheva, L.G. 550, 694

Butenko, Y.V. 550

Butler, J.E. 368

_agin, T. 787

Catledge, S.A. 452

Chalamala, B.R. 708

Chanana, R.K. 618

Chang, H.L. 829

Chang, S.-K. 214, 343

Che, J. 787

Chen, C.-F. 109

Chen, G. 642

Chen, G.Z. 661

Chen, J.C. 366

Chen, Y.-T. 305

Chen, Z. 810

Cheng, H.-F. 724

Chew, K. 621

Chiang, I.W. 687

Chien-Min Sung, J. 32

Cho, W.-K. 716

Choi, J. 759, 764

Choi, Y.C. 764

Chung, G.Y. 618

Chung, H.-K. 254, 305, 334, 364

Chuvilin, A.L. 694

Cizek, J. 595

Clarke, R. 577

Coll, B.F. 708

Collins, A.T. 223

Comins, J.D. 376

Cooke, R. 424

Cooper, J.B. 275, 170

Cui, J.B. 114

Cutshaw, C. 736, 742

Dai, C.-W. 453

Dai, Y.B. 386

Daniel, A. 577

Das, M.K. 618

Davidenko, V.M. 237

Davidson, J.L. 82, 104, 144

DeSio, A. 179

De Weerdt, F. 232

Dean, K.A. 708

DeJoseph, C.A. 597

Dementjev, A.P. 257, 554

Deng, W. 787

Deng, X. 615

Desai, V. 424

Dickey, E.C. 466, 771

Dideikin, A.T. 264

Doh, C.-H. 33

Doll, G.L. 460

Duan, X.F. 552

NASA/CP--2001-210948 831 August 6, 2001
Dzbanovsky, N.N. 68

Eaton,S.C. 164

Ehret,G. 551

Eibisch, H. 333

Encarnacidn, RA. 577

Erasmus, R.M. 376

Erdemir, A. 437

Ertl,S. 90

Espinosa, R. 399

Eun,K.Y. 304,454,522

Evans, R.D. 460

Evstefeeva, Y.E. 164,171

Fan,S.S.816

Farhat, S. 816,822

Feldermann, H. 571

Feldman, L.C. 618

Feoktistov, N.A. 264

Ferrari, A.C. 521

Files,B.S. 769

Fink,R.L. 399

Fisher, T.S. 144

F1/Ster, A. 90

Fogarassy, E. 616

Fonseca, A. 550

Frank,M. 198

Fray,D.J. 661

Freeman, M. 395

Friedbacher, G. 340,344

Fries-Carr, S. 395

Fryda,M. 158

Fu,W. 275,718,726

Fuchigami, K. 486

Fuchs, C. 616

Fuji,K. 348

Fujishima, A. 150

Fujiwara, S. 505,632, 657

Furuta, H. 823

Gao,R.R 752

Gao,X. 197

Garguilo, J.M. 126,132,133

Gaudin, O. 178

Gavrilov, S.A. 68

Glass, J.T. 460

GoddardIII,W.A. 787

Golubev, V.G. 264

Gorelik,O. 769

Goto,A. 458,523,528, 800

Grill,A. 397

Grudinkin, S. 264

Grulke, E. 701

Gu,C. 266,273

Gullickson, E. 522

Gunda, N. 424

Guo,W.T. 197

Haddon, R.C. 771

Hadjiev, V.G. 769

Haidar, J. 498

Hajek,V. 595

Hamada, Y. 583

Han,H.X. 816

Han,J. 700

Han,J.H. 805,828

Hanada, T. 823

Handa, Y. 492

H_inni, W. 158

Hanrath, T. 604

Hartman, J. 604

Hasan, S. 544

Hasegawa, M. 121

Hatta,A. 348

Hauge, R.H. 687

Haung, J.-L. 509

Hayashi, K. 24

He,X.C. 386

Heggie, M.I. 550

Heidger, S.L. 61

Heo,J. 747

Herrmann, D. 158

Hinkov,I. 816

Hirahara, T. 609

Hirakuri,K.K. 340,344

Hirao,T. 110,323,823

Hirata,A. 538

Hirayama, T. 795

Hoch,B. 652

Hoffman, 1st. 120

Hoffmann, V. 400

Hofs_iss, H. 400,571

Holland,O.W.618

Holmes, K. 82

Hommet, J. 616

Honda, S. 823

Hong,E.H. 805,828

Horiuchi, M. 486

Hosomi, T. 267,372

Howell,M. 82

Hsu,K.-H. 305

Hsu,T. 410

Hsu,W.K. 661

Hu,C.G. 139,406, 579

Hu,H. 771

Hu,W. 810

Hu,X.J. 386

Huang, Q.E 473

Huang, T.B. 285

Hung,J.-L. 343

NASA/CP--2001-210948 832 August6,2001
Hwang, H.Y. 716

Hyodo,H. 412

Igarashi, M. 538

Ii'ichev,E.A. 68

Imai,T. 31,110

Ingram, D.C. 597

Inoue,H. 651

Inoue,K. 24

Ishihara, M. 51,636

Ivahknenko, S.A. 355

Iwataka, N. 369

Jackman, R.B. 178

Jagtoyen, M. 701

Jaskie, J.E. 708

Jentsch, H.-G. 333

Jeon,K.K. 33,669,675

Jeong, H.J. 669,675,695,747

Jeong, S.-H.716

Jeong, S.Y. 675

Jeong, T. 747

Jha,S. 424

Jin,M. 204

Jin,S. 747

Jin,Z. 266,273

Jo,C. 695,753

Jung,H.J. 759

Kaito,T. 538

Kakudate, Y. 632,657

Kalish,R. 120

Kamijo,E. 609

Kan,M.-C. 516,627

Kanda, H. 279

Kang,W.R 82,104,144

Kassabian, K. 431

Kawakami, N. 24

Kehrli,M. 370

Kempinski, W. 264

Kerns,D.V. 104

Khmelnitskiy, R.A. 370

Kidalov, S.V. 237

Kiely,C.J. 794

Kim,J.M. 747

Kim,K.-Y. 716

Kim,S.-H.267

Kim,T.-G.267

Kim,W.S. 33,669,675,747

Kim,Y.-H. 267

Kimura,A. 779

Kirkici,H. 186

Kitamura, J. 632

Kiyohara, S. 317

Kobashi, K. 16,24,110,323

Kobayashi, T. 267,344,372

K/Sck, F.A.M.126,132,133

Koenigsfeld, N. 120

Koga,Y. 9,51,447,458,485,505,523,528,636,

681,800

Kohler-Redlich, R 551

Kohn,E. 90

Koidl,R 292

Koizumi, S. 279

Kokai,F. 51,505,636

Kolupaeva, Z.I. 327

Komatsu, T. 651

Kononenko, V.V. 370

Konov, V.I. 370

Koretz, A. 255

Kormunda, M. 595

Kosai,H. 395

Kosel,RB. 62,395

Krauser, J. 400

Kroto,H.W. 661

Krotova, M.D. 171

Krt_ger, A. 264

Krzanowski, J.E. 598

Kuo,C.T. 829

Kuo,J.-K. 453

Kuo,T.-F. 724

Kusunoki, M. 795

Kuznetsov, V.L. 550

Kyotani, M. 523,528, 681

Lai,T.-H. 724

LamyDeLaChapelle, M. 816

Landis,C. 771

Lansley, S.R 178

Lanter,W.C.597

Laptev, V.A. 171

Larsen-Basse, J. 3

LeNormand, F. 263,551,616

Lee,C.S. 262,522,552,717

LeeJ. 747

LeeJ.-K. 304

LeeK.-H. 701,716, 805,828

LeeK.-R.454,522

LeeN.S. 707

LeeS.-Y.343

LeeS.L. 648

LeeS.M. 33

LeeS.T. 262,552,565,717, 730

LeeT.-C. 667

LeeY.H. 33,669,675,695,747, 753,758, 759,764

Lemmer, O. 198

Lewis,K.L. 177

Ley,L. 114

Leyendecker, T. 198

Li,G.H. 816

NASA/CP--2001-210948 833 August 6,2001
Li,J.Q. 197

Li,Q. 552,565

Li,Z. 139

Lian,G. 466

Liao,K.J. 139,406, 579

LiBassi, A. 521

Lifshitz,Y. 552

Ligatchev, V. 621

Lim,J.-H. 33

Lim,S.C. 675,695, 758,759

Lin,C.-L. 109

Lin,C.H. 829

Lin,G.-M. 410

Lin,I.-N. 410,724

Liu,C. 736,742

Liu,Y. 254

Lu,F.X. 285

Luo,S.Y.453

Lt_thy, W. 370

Ma,T.C. 491,615

Ma,Z.Q. 638

Machi,Y. 340,344

Maki,T. 267

Malogolovets, V.G. 355

Malshe, A.R 210,589

Malyi,V. 255

Mao,Z. 701

Margrave, J.L. 687

Marinelli,M. 78,80

Martin,Z. 376

Maslakov, K.I. 257,554

Matsui,H. 523

Matsumoto, S. 561,565

May,RW. 291

McCulloch, D.G. 457

McDonald, K. 618

McKeag, R.D. 178

McKenzie, D.R. 457

Meguro, K. 110

Mei,X.X. 491,615

Meier,M.S. 771

Melehin, V.G. 264

Meng,F.Y. 552

Merchant, A. 457

Mermoux, M. 293

Mickelson, E.T. 687

Milani,E. 78,80

Minakov, R 68

Mironov,E. 255

Mishuk,V.Y. 171

Misra,D.S. 360

Mitra,S. 604

Miyake,S. 599

Mizutani, A. 779

Mofokeng, V. 376

Mollart,T.R 177

Momose, H. 479

Moon,J.-M. 33,669,675,695, 758

Moon,S.-I. 33

Mori,K. 317

Mori,Y. 533,823

Morrison, Jr.,RW. 460

Motta,N. 263

Moulton,J.A. 170,275

Mtiller-Sebert, W. 292

Murakawa, M. 204,382,599

Musumarra, A. 78

Mutsukura, N. 340,344,492

Nachalnay, T.A. 355

Nakahigashi, T. 423

Nakajima, H. 583

Nakamura, E. 340,344

Nakamura, T. 51,458

Naumkin, A.V. 554

Nemanich, R.J. 126,132,133,681

Nikolaev, R 769

Nishi,Y. 369,779

Nishibayashi, Y. 110,323

Niu,C. 652

Noguchi, H. 204

Nolting,F. 454

Novoselova, I.A. 355

Oguri,K. 369,779

Oh,S.H. 805,828

Ohana, T. 458

Ohashi, T. 583

Ohshima, S. 651

Ohtake, N. 479

Ohyama, M. 340

Okotrub, A.V. 550,694

Okushi, H. 121

Okuzumi, F. 55

Omino,K. 583

Onishi,F. 335

Ono,T. 121

Opalev, O.A. 327

Ordanyan, S.S. 237

Orr-Ewing, A.J. 291

Osaki,H. 413

Osawa, E. 264

Osipov, V.Y. 264

Osipowicz, T. 473,621

Oura,K. 110,323,823

Ozawa, M. 264

Pace,E. 179

Padmore, H. 454,522

Palmour, J.W.618

Paltnikar, U. 360

NASA/CP--2001-210948 834 August6,2001
Palyanov, Y.N. 171

Pan,Z.W. 730,752

Panich, A.M. 264

Panosyan, Z. 431

Pantelides, S.T. 618

Paoletti, A. 78,80

Papworth, J.J. 794

Park,C.-G.805,828

Park,G.-S.695

Park,S. 454

Park,Y.S. 33,669,675,695

Pashnev, V.K. 327

Paulmier, D. 293

Pavlovsky, I. 399

Payandeh, B. 667

Pazik,J.4

Pecoraro, S. 263

Perret, A. 158

PersNov, E. 245

Petrov, E. 255

Phani, A.R. 598

Piao,G. 528

Pillon,M. 80

Pimenov, S.M. 370

Pini,A. 179

Pleskov, Y.V. 164,171

Pleuler, E. 292

Podzyarey, G.A. 355

Poliakov, V.R 239

Polini,R. 263

Poltoratsky, E.A. 68

Popov, M. 681

Potenza, R. 78

Pradhan, D. 830

Prater, J.T. 55

Pucella, G. 78

Ramamurti, R. 62

Randieri, C. 78

Rantell, T. 771

Reinhard, D.K. 180

Richkov, G.S. 68

Ristein, J. 114

Rizvi,N. 178

Robert, M. 667

Robertson, J. 521

Romanenko, A.I. 694

Romano, V. 370

Romeo, M. 551

Ronning, C. 400,571

Rosiwal, S.M. 333

Rosser, K.N. 291

Ruoft,R. 772

Rusli,473,621

Russell, W. 589

Ryu,C.-M. 805,828

Ryu,J.E. 805,828

Saini,R.K. 687

Saito,D. 71

Saito,K. 810

Sakai,T. 121

Sakuma, N. 121

Samiee, M. 62

Sasa, T. 486

Sasaki, K. 523

Sasaki, T. 533,823

Sastri,S.A. 424

Sch_ifer, L. 158

Schlesser, R. 55

Schmid, R 90

Schmitt, M. 293

Scholl,A. 454

Schultrich, B. 400

Scott,C.D. 769,816

Scuderi, S. 179

Shames, I. 264

Shang, N.G. 552

Shanov, V. 62

Sharon, M. 830

Shen,H.S. 386

Sheng, Y. 577

Shi,S.-Q.773

Shibata, N. 795

Shibata, T. 98

Shin,Y.M. 675,747,759

Shmal'ko, Y.F. 327

Shvets, D. 553

Sikder, A.K. 360

Siklitski,V.I. 264

Singer, R.F. 333

Singh,M.K. 360

Singh,R.N. 62

Sinnott, S.B. 701

Sitar,Z. 55

Smith,J.A. 291

Song, J.H. 285

Song,X. 642

Sperduto, M.L. 78

Spitsyn, B.V. 284

Stammler, M. 114

Stelitano, A. 78

Stepanyan, A. 431

Stoner, B.R. 55

Strel'nitskij, V.E. 327,509

Suetin, N.V. 68

Suh,E.K. 759

Summitomo, T. 348

Sun,Yue 266

Sun,Ye 266

NASA/CP--2001-210948 835 August 6,2001
Sun,Y.N. 197

Sung,J.C. 211,213,214, 253,254, 305,334, 343,364,

365,366,436, 453,504,596,617,627,648

Sunkara, M.K. 466

Sussmann, R.S.43

Suzuki,H. 657

Suzuki,K. 486

Suzuki, T. 795

Szorenyi, T. 616

Tachibana, T. 24,323

Tai,M.F. 648

Takagi, Y. 335

Takeuchi, S. 382

Talin,A.A. 708

Tanaka, A. 11,447,458,485

Tang,W.Z. 285

Taniwaki, M. 636

Tanner, B.K. 521

Tashiro, T. 71

Tennent, H. 652

Teo,E.J. 473,621

Teraji,T. 279

Thuesen, L.F. 399

Tin,C.C. 618

Titus,E. 360

Tochimoto, T. 609

Tomfinek, D. 793

Tonegawa, A. 369,779

Tong,Y.M. 285

Toriyama, T. 779

Torosyan, G. 431

Toyoguchi, T. 412

Trautmann, C. 400

Tricomi, C.M. SuteraeA. 78

Tr/Sster, I. 158

Trujillo,J. 708

Tsai,C.-H. 724

Tsai,C.-L. 109

Tsai,M.H. 829

Tsai,T.J. 453

Tsugawa, K. 528,800

Tsuji,K. 372

Tucciarone, A. 78,80

Tzeng, Y. 736,742

Uchida, K. 651

Uematsu, K. 486

Uemura, M. 215

Ueno,M. 466

Usuba, S. 632,657

Vaseashta, A. 72

Verona-Rinati, G. 78,80

Vlcek,J. 595

Vo,T. 667

Vohra,Y.K. 274,452

Vul',A.Y. 237,264

Vyrovets, I.I. 327

Waiblinger, M. 400

Walker, D.G. 144

Walton, D.R.M.661

Wan,Y.Z. 386

Wang,J.T. 367

Wang,S.-F.504,627

Wang, W.L. 139,406, 579

Wang,Z.L. 730,752

Watanabe, A. 323

Watanabe, K. 279

Watanabe, S. 599

Watanabe, T. 485,505

Watt,F. 473,621

Watts,RC.R 661

Wazumi, K. 447,486,505

Weber, H.R 370

Wei,Y. 708

Weidinger, A. 400

Weimer, J.A. 61,395

Weller,R.A. 618

Werckmann, J. 263

Whitfield, M.D. 178

Wild,C. 292

Williams, J.R. 618

Wisitsora-at, A. 104

Wolter,S.D. 55

Wong,M.S. 254

Wong,W.K. 717

Wrbanek, J. 395

Wu,B.R. 648

Wu,R.L.C.395,597

Xia,B. 45

Xiao,B. 170,275,718, 726

Xie,S.S.730

Xu,J. 406,579,615

Xu,X. 787

Yagi,H. 291

Yamaguchi, C. 523

Yamamoto, K. 485,505,583, 632

Yamamoto, T. 412

Yamaoka, M. 533,823

Yang,X.D. 479

Yaniv,Z. 399

Yao,A. 45

Yap,Y.K. 533,823

Yashin, V.A. 237

Yedave, S.N. 589

Yeh,B. 453

Yeh,W.T. 504,509,627

Yengibaryan, Y. 431

Yi,W. 747

Yokoi,H. 632,657

NASA/CP--2001-210948 836 August6,2001
Yokota, Y. 24,323

Yoon,K.-H. 522

Yoon,S.F.473,621

Yoshida, H. 121

Yoshikawa, M. 5

Yoshimura, M. 533,823

Yu,J. 561

Yu,S. 747

Yuan,L. 810

Yudanov, N.F. 694

Yue,J. 642

Yugo,S. 71

Yumura, M. 651

Yushin, G.N. 55

Zabusov, O.O.257,554

Zaleskij, D.U. 509

Zanevskiy, O.A. 355

Zavedeev, E.V. 370

Zhan,Z. 45

Zhang, J. 615

Zhang, J.H. 197

Zhang, L. 121

Zhang,Q. 473,621

Zhang,S. 642

Zhang,W.J. 561,565

Zhang,Z.M. 386

Zhao,G.L. 544

Zhou,L.G. 773

Zhou,O. 668

NASA/CP--2001-210948 837 August 6,2001
REPORT DOCUMENTATION PAGE Form Approved

OMB No. 0704-0188

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1. AGENCY USE ONLY (Leave blank) 2. REPORT DATE 3. REPORT TYPE AND DATES COVERED

July 2001 Conference Publication

4. TITLE AND SUBTITLE 5. FUNDING NUMBERS

Proceedings of the Sixth Applied Diamond Conference/Second Frontier

Carbon Technology Joint Conference (ADC/FCT 2001)

6. AUTHOR(S)

Y. Tzeng, K. Miyoshi, M. Yoshikawa, M. Murakawa, Y. Koga,

K. Kobashi, G.A.J. Amaratunga, editors

7. PERFORMING ORGANIZATION NAME(S) AND ADDRESS(ES)

National Aeronautics and Space Administration

John H. Glenn Research Center at Lewis Field

Cleveland, Ohio 44135-3191

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National Aeronautics and Space Administration

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E-12802

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NASA CP--2001-210948

11. SUPPLEMENTARY NOTES

Proceedings of a conference held at Auburn University cosponsored by Auburn University, Alabama Microelectronics

Science and Technology Center, Space Power Institute, Alabama Space Grant Consortium, NASA Glenn Research

Center, Japan New Diamond Forum, and Japan Fine Ceramic Center, Auburn, Alabama, August 6-10, 2001. Responsible

person, Kazuhisa Miyoshi, organization code 5160, 216-433-6078.

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Available electronically at ]£_p_'._//_g_r_._gr_.c_n_a__a_..g,_a_vJ_G_L_T_R__S.

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13. ABSTRACT (Maximum200 words)

These are the Proceedings of the Sixth Applied Diamond Conference/Second Frontier Carbon Technology Joint Conference hosted by Auburn University fromAugust 6 to 10, 2001. The diamond CVD process was first reported by Dr.

Spitsyn in 1981 and Prof. S. Iijima reported his discovery of carbon nanotubes in 1991. In the past years, both diamondrelated materials and novel carbon materials have attracted considerable interest by the scientific, technological, and

industrial community. Many practical and commercial products of diamond materials are reported in these proceedings.

A broad variety of applications of carbon nanotubes and novel carbons have also been explored and demonstrated.

Having more than 200 invited and contributing papers by authors from over 20 countries for presentations at ADC/FCT

2001 clearly demonstrates that these materials, due to the combination of their superior properties, are both scientifically

amazing and economically significant.

14. SUBJECT TERMS

Coating; Nanotechnology

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ISBN 0-9710327-0-X