United States Department of Commerce

National Institute of Standards and Technology

NIST

PUBLICATIONS

A111D3 M7flMSb

MSr Special Publication 804

Chemistry of Electronic Ceramic Materials

Peter K. Davies and Robert S. Roth, Editors

100

.U57

1991

C.2

NIST Special Publication 804

Chemistry of Electronic Ceramic Materials

Proceedings of the International Conference on

the Chemistry of Electronic Ceramic Materials,

August 17-22, 1990 held at Jackson, WY

Edited by

Peter K. Davies

University of Pennsylvania

Robert S. Roth

National Institute of Standards and Technology

Editorial Assistant

Mary A. Clevinger

National Institute of Standards and Technology

Sponsored by

NIST

NASA

ONR

DuPont

Idaho National Engineering Laboratory

University of Pennsylv£inia

January 1991

U.S. Department of Commerce

Robert A. Mosbacher, Secretary

National Institute of Standards and Technology

John W. Lyons, Director
National Institute of Standards U.S. Government Printing Office

and Technology Washington: 1991

Special Publication 804

Natl. Inst. Stand. Technol.

Spec. Publ. 804

513 pages (Jan. 1991)

CODEN: NSPUE2

For sale by the Superintendent

of Documents

U.S. Government Printing Office

Washington, DC 20402
PREFACE

The International Conference on the Chemistry of Electronic Ceramic Materials was

held at The Sojourner Inn, in the Teton Village, Jackson Hole, Wyoming from August

17-22, 1990. In an attempt to maximize the development of this rapidly moving,

multidisciplinary field, this conference brought together major national and international

researchers to bridge the gap between those primarily interested in the pure Chemistry

of Inorganic Solids and those interested in the Physical and Electronic Properties of

Ceramics. With the many major discoveries that have occurred over the last decade, one

of the goals of this meeting was to evaluate our current understanding of the Chemistry

of Electronic Ceramic Materials, and to assess the state of a field that has become one of

the most important areas of Advanced Materials Research.

The conference consisted of 18 extended invited lectures, 10 contributed oral

presentations, and 37 poster presentations. The meeting was attended by scientists from

a broad spectrum of fields based at universities, industries, and national laboratories.

Approximately one-third of the attendees were from outside the United States and

represented 10 different countries. In addition to the quality of the science, the meeting

was also memorable for the large number of family members who were able to attend,

resulting in many new friendships.

The scientific presentations covered many topics including new methods for the

synthesis of ceramics, the structures and properties of dielectric and ferroelectric

materials, crystal chemistry, surface chemistry, high-T^, superconductors, and theory and

modeling. The papers appearing in these proceedings provide an up-to-date review of

many aspects of these areas. For most of the oral presentations the discussions following

the paper were successfully transcribed. Edited versions appear in these proceedings after

each paper. In addtion to clarifying certain aspects of the material covered in the lecture,

the unedited comments, transcribed by a local court reporting service, also provided

considerable amusement for the editors and speakers.

The success of a meeting depends on many parameters ranging from the science to the weather. One of the most important is sponsorship and support. For this we are

grateful to NIST, NASA, ONR, DuPont, the Idaho National Engineering Laboratory, and

the University of Pennsylvania. The smooth running of a meeting also depends upon

considerable work by those who often receive little recognition. In this respect we were

truly fortunate to have Mary Clevinger as our conference coordinator who has also

ensured that these proceedings appeared in an accurate form in a timely manner. We

would also like to thank Audrey Roth and Kathy Davies for organizing an outstanding

activities program for the many family members who attended the meeting. Thanks are

also due to Denice Gilbert, Nancy Dickinson, and Tony Kostick in the Materials Science

Department at Penn, who assisted in many typing duties and several tedious budgetary

matters.

iii
We hope that these proceedings will contribute to the development of the Chemistry

of Electronic Ceramic Materials and, judging from the enthusiastic response of all who

attended this meeting, we look forward to the Second International Conference sometime

in the not too distant future.

October 1990 Peter K. Davies (Univ. of Pennsylvania)

Robert S. Roth (NIST)

iv
TABLE OF CONTENTS

PREFACE

I. DIELECTRIC CERAMICS

CHEMICAL REACTION CONTROLLED MICROSTRUCTURES 3

AND PROPERTIES OF FERROELECTRIC CERAMICS (Invited)

D. Kolar

CHEMISTRY AND PROPERTIES OF TEMPERATURE 21

COMPENSATED MICROWAVE DIELECTRICS (Invited)

T. Negas, G. Yeager, S. Bell, and R. Amren

TUNABLE TRANSDUCERS: NONLINEAR PHENOMENA IN 39

ELECTROCERAMICS (Invited)

R. E. Newnham

STRUCTURAL AND ELECTRONIC PROPERTIES OF SOME 53

PEROVSKITES (Invited)

P. D. Battle

FERROELECTRIC AND FERROELASTIC DOMAIN STRUCTURES 67

IN PIEZOELECTRIC CERAMICS (Invited)

L. A. Bursill and P. JuLin

CATION SUBSTITUTION IN BARIUM TITANATE FOR 77

DIELECTRIC CERAMIC APPLICATIONS

F. Batllo, A. Beauger, B. Jannot, J. C. Jules, M. Laurent, M.

Maglione, and J. C. Niepce

LIQUID PHASE SINTERING OF LiF-FLUXED BaTiOa 85

S. F. Wang, W. Huebner, and C. Randall

A REAL SPACE ANALYSIS OF SHORT RANGE ORDER IN 93

FERROELECTRIC Pb(Mgi/3Nb2/3)03

H. D. Rosenfeld, T. Egami, and A. Bhalla

V
MECHANISM OF PNN BASED PEROVSKITE CERAMICS

FORMATION

Y. Sasaki, A. Nagai, and T. Yoshimoto

CONVENTIONALLY PREPARED SUBMICRON ELECTROCERAMIC POWDERS BY REACTIVE CALCINATION

T. R. Shrout

LOW-TEMPERATURE SYNTHESIS AND CHARACTERIZATION

PREPARATION OF YBa2Cu307.5 FROM HOMOGENEOUS METAL

ALKOXIDE SOLUTION: SYNTHESES OF MIXED-LIGAND

COPPER (II) ALKOXIDES

D. M. Millar and D, A. Payne

MOLECULAR CHEMISTRY AND THE SYNTHESIS OF

PRECURSORS TO ELECTRONIC CERAMIC MATERIALS (Invited)

J. J. Ritter, J. F. Kelly, D. E. Newbury, and D. B. Minor

ION EXCHANGE REACTIONS OF LAYERED

PHOSPHATOANTIMONIC ACIDS: A ROUTE FOR NEW

CATALYSTS AND LUMINESCENT MATERIALS (Invited)

Y. Piffard

MIXED METAL OXIDES WITH THE PYROCHLORE STRUCTURE

AS CATALYSTS FOR METHANE OXIDATIVE COUPLING

(Invited)

A. J. Jacobson, L T. Lewandowski, C. A. Mims, R. B. Hall,

and G. R. Myers

SOFT CHEMICAL SYNTHESIS OF METASTABLE TITANIUM,

VANADIUM, AND MOLYBDENUM OXIDES

T. P. Feist and P. K. Davies

PRELIMINARY SOLID-STATE MULTINUCLEAR NMR OF

TITANIUM AND ZIRCONIUM OXIDE CERAMICS AND

PRECURSORS

J. J. Fitzgerald, S. S. Han, S. F. Dec, M. F. Davis, C. E.

Bronnimann, and G. E. Maciel

vi
POLYMERIC SYNTHESIS OF PEROVSKITE POWDERS AND 179

FILMS

H. U. Anderson, M. M. Nasrallah, F. D. Blum, and M. S. Smith

ELECTRODEPOSITION OF NANOMODULATED ELECTRONIC 185

CERAMIC THIN FILMS

J. A. Switzer

III. SOLID-STATE SYNTHESIS AND CHARACTERIZATION

PROPERTIES OF SOME MIXED URANIUM OXIDES 195

P. G. Dickens, A. V. Powell, and G. P. Stuttard

INVESTIGATIONS ON THE STRUCTURAL, ELECTRICAL AND 207

MAGNETIC PROPERTIES OF Nd2.xM^Ni04+5, M = Ca^+ and Ba^+

B. W, Arbuckle, Z. Zhang, and M. Greenblatt

SYNTHESIS AND CHARACTERIZATION OF La^^SfJiOa

(0<x<0.05) 217

S. M. Kauzlarich, J. E. Sunstrom, IV, and P. Klavins

SYNTHESIS AND CHARACTERIZATION OF PHASES IN THE

SYSTEM Ba-Au-Ag-02 225

R. S. Roth, C. J. Rawn, and M. D. Hill

STRUCTURE AND PROPERTIES OF REDUCED EARLY

TRANSITION METAL OXIDE SINGLE CRYSTALS GROWN 237

FROM BORATE FLUXES

S. A. Sunshine, B. Hessen, T. Siegrist, A. T. Fiory, and J. V.

Waszczak

THE KINETICS AND MECHANISM OF THE CRYSTALLIZATION

OF Mg2Al4Si50i8 FROM MgAl204 AND Si02 IN THE PRESENCE 245

OF A BISMUTH OXIDE FLUX

R, W. Dupon, A. C. Tanous, and M. S. Thompson

vii
IV. CRYSTAL CHEMISTRY

SOLID STATE CHEMISTRY OF PEROVSKITE AND RELATED

OXIDES CONTAINING Fe^^, Ni^+, and Cu^^ (Invited)

M. Takano

IMPURITY INCORPORATION MECHANISMS IN LiNbOj

D. P. Birnie, III

ALKALINE EARTH NITRIDES AND HYDRIDES

N. E. Brese

OXYGEN DIFFUSION IN Y2O3-CONTAINING TETRAGONAL

ZIRCONIA POLYCRYSTALS (Y-TZP)

Y. Ikuma

A STRUCTURAL BASIS FOR THE DIFFERENT DIRECTIONS OF

OXYGEN NONSTOICHIOMETRY IN La2Cu04 AND Nd2Cu04

D. J. L. Hong and D. M. Smyth

RUDDLESDEN-POPPER PHASES A„+iM„03„+i: STRUCTURES

AND PROPERTIES

J. S. Kim, J. Y. Lee, J. S. Swinnea, H. Steinfmk, W. M. Reiff,

P. Lightfoot, S. Pei, and J. D, Jorgensen

V. SUPERCONDUCTORS

MICROSTRUCTURES IN HIGH TEMPERATURE

SUPERCONDUCTORS (Invited)

H. W. Zandbergen

CRYSTAL CHEMISTRY OF OXIDE SUPERCONDUCTORS

(Invited)

A. W. Sleight

OUT OF PLANE DISPLACEMENTS OF OXYGEN FROM THE

CUO2 SHEETS IN Ca ^^St 15CUO2 BY ATOM-PAIR DISTRIBUTION

FUNCTION ANALYSIS

S. J. L. Billinge, P. K. Davies, and T. Egami

viii
CRYSTAL CHEMISTRY AND OXYGEN ACTIVITY EFFECTS ON 343

T, T', AND T* PHASE STABILITIES (La,RE)2Cu04 SYSTEMS

J. F. Bringley, B. A. Scott, and S. S. Trail

ROOM TEMPERATURE MOBILITY AND DIFFUSION 351

COEFFICIENT OF OXYGEN IN POLYCRYSTALLINE YBa2Cu307

Y. Scolnik, M. Schwartz, E. Sabatani, and D. Cahen

PHASE SEPARATION IN Ndj ,,Ce,Cu04

B. Dabrowski, P. Lightfoot, D. R. Richards, D. G. Hinks, J. ^57

D. Jorgensen, S. Pei, Y. Zheng, D. T. Marx, and A. W.

Mitchell

NEW RARE-EARTH CUPRATES WITH THE NaCuOj

STRUCTURE 365

P. K. Davies, E. Caignol, and T. King

THERMODYNAMIC MEASUREMENTS IN THE Y-Ba-Cu-O

SYSTEM 373

F. H. Garzon and I. D. Raistrick

CALORIMETRIC STUDIES OF CERAMICS

A. Navrotsky

SYNTHESIS AND PROPERTIES OF Ba-FREE

SUPERCONDUCTIVE (Eu,Ce)4(Eu,Sr)4Cu6.^M^O^ (M:Fe,Co,Al) 393

A. Ichinose, T. Wada, Y, Yaegashi, A. Nara, H. Yamauchi,

and S. Tanaka

EXCESS OXYGEN DEFECTS IN LAYERED CUPRATES

P. Lightfoot, S. Pei, J. D. Jorgensen, X.-X. Tang, A. 399

Manthiram, and J. B. Goodenough

CAN Co(Fe) SUBSTITUENT IN YBa2Cu307.5 MIGRATE BACK

AND FORTH BETWEEN Cu(l) AND Cu(2) SITES? 407

A. Nath, Z. Homonnay, G. W. Jang, S. 1. Nagy, Y. Wei, and

C. C. Chan

ELECTROCRYSTALLIZATION IN THE Ba-K-Bi-O SYSTEM

M. L. Norton, H. Y. Tang, and J. Eglin 413

ix
STRUCTURAL DIVERSITY IN OXYGEN-DEFICIENT

PEROVSKITES

J. T. Vaughey and K. R. Poeppelmeier

NEW LaCu03.j PEROVSKITES PREPARED AT HIGH OXYGEN

PRESSURE

J. F. Bringley, B. A. Scott, S. J. LaPlaca, R. F. Boehme, M.

W. McElfresh, T. M. Shaw, S. S. Trail, and D. E. Cox

VI. SURFACE CHEMISTRY

SURFACE ENERGY BARRIER FORMED BY ADSORBED

OXYGEN IN POROUS ZnO

S. Fujitsu, K. Koumoto, and H. Yanagida

SURFACE PROPERTIES OF BaTiOj AT ELEVATED

TEMPERATURES

J. Nowotny and M. Sloma

A SCANNING TUNNELING MICROSCOPY STUDY OF SINGLE

CRYSTAL ZnO AND TiOj SURFACES

G. S. Rohrer and D. A. Bonnell

VIL THEORY AND MODELING

FACTORS DETERMINING THE DIELECTRIC CONSTANTS OF

OXIDES AND FLUORIDES (Invited)

R. D. Shannon

INTERNAL STRAIN IN PEROVSKITE RELATED MATERIALS

(Invited)

I. D. Brown

EMPIRICAL METHODS IN OXIDE CRYSTAL CHEMISTRY

(Invited)

M. O'Keeffe

COMPUTER SIMULATION STUDIES OF ELECTRONIC

CERAMICS

A. N. Cormack
COMPUTER SIMULATION OF DOPANT SUBSTITUTION IN 507

YBa2Cu307

M. S. Islam and C. Ananthamohan

MOLECULAR DYNAMICS SIMULATIONS OF ION MOTION IN 513

DIVALENT AND MIXED MONOVALENT-DIVALENT BETA"-

ALUMINA

C. A. Lane and G. C. Farrington

AUTHOR INDEX 519

ATTENDEES 521

xi

DIELECTRIC CERAMICS

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

CHEMICAL REACnON CONTROLLED MICROSTRUCTURES AND

PROPERTIES OF FERROELECTRIC CERAMICS

D. Kolar

"J. Stefan" Institute, University of Ljubljana, Ljubljana, Yugoslavia

ABSTRACT

Electronic ceramic components are usually manufactured from complex mixtures of oxides or partially prereacted compounds. During sintering at high

temperatures, chemical reactions take place which influence the kinetics and

mechanism of densification and development of ceramic microstructure. The

densification process is also critically dependent on minor impurities. It is recognised that the inadequate control of chemical reactions frequently does

not allow generaHsation of research results and is the main reason for poor

reproducibility in manufacture of electronic ceramics. The present review is concentrated on interdiffusion phenomena and growth in presence of minor

amounts of reactive Hquid.

1. INTRODUCTION

Ceramic ferroelectrics are extensively used for manufacture of important

electronic devices, notably ceramic capacitors, nonlinear resistors and piezo

elements. More than 40 years of extensive investigation of ferroelectric

phenomena resulted in a good understanding of the fundamental physical

principles underlying the function and performance of various devices.

The principles and latest achievements in the field of ceramic capacitors were

thoroughly discussed in recent reviews by Payne (1), Goodman (2), Hennings

(3) and Kahn et al. (4).

At present, two families of materials are used in manufacture of multilayer

capacitors: barium titanate based compositions and lead niobate based com3
positions, known as relaxors. The barium titanate family remains the major

component for manufacture of MLC dielectrics with permittivity above 1000 to

15000, and above.

In spite of extensive development, several problems remain to be solved. The

present review stresses the importance of chemical reactions during sintering

at high temperatures on the microstructure development and electrical properties of BaTiOs based dielectric ceramics.

2. THE IMPORTANCE OF CHEMICAL REACTIONS AT HIGH

TEMPERATURES

High resistivity ceramic capacitors with low temperature coefficients of

capacitance are chemically inhomogeneous materials in metastable thermodynamic equilibrium.

Due to this thermodynamic instability, the technology is rather demanding and

requires careful control of manufacturing parameters, such as chemical composition, initial homogeneity, powder grain size and morphology, homogeneity

of the formed body, heating rate, sintering temperature and time.

In general, maintainance of chemical heterogeneity demands a low sintering

temperature and short sintering times to limit sohd state diffusion, which causes homogenisation, A low sintering temperature is also required to allow the use

of less expensive electrodes in multilayer devices.

On the other hand, sintering time and temperature must assure densification to

low and closed porosity. Porous bodies exhibit a low dielectric constant, low

breakdown strength and high loss due to water permeability.

High density at lower sintering temperatures is usually achieved with additives

which promote densification by forming liquid phase at sintering temperature,

as will be discussed later. Sintering additives dilute the high dielectric constant

phase, therefore the amount should be as small as possible.

The third deciding parameter which determines electrical properties is ceramic

microstructure. Small grain size with narrow size distribution is normally re- quired for most applications. Liquid phase sintering with small amounts of liquid

phase frequently results in coarse microstructure with broad size distribution.

It is clear that carefull optimization of chemical composition and processing

parameters is needed to achieve the best compromise, i.e. dense, fine grained

and chemically heterogeneous ceramics.

4
3. LIQUID PHASE SINTERING OF BARIUM TTTANATE BASED

COMPOSITIONS

The basic approach to preparing high capacitance material is to shift the BaTiOs

Curie peak close to room temperature and to broaden it. Shifting is ac- complished with solid solution additives, for instance alkaline-earth titanates,

lead titanate, various stannates and zirconates (2).

Broadening of the Curie temperature maximum is achieved by compositional

inhomogeneity. Inhomogeneity causes a diffuse phase transition (DPT). A

simple model of the inhomogeneous ceramic is that it consists of a symmetric

compositional distribution of micro regions, each having a Curie point temperature characteristic of its composition (5).

In accordance with the model, the width of the distribution of Curie temperature

maxima may serve as an indication of the homogeneity of the dielectric (6).

Barium titanate with solid solution additives are refractory compounds which

demand high sintering temperatures, 1350°C and above.To decrease the high

sintering temperature, which favours homogenisation and yet to enable densification to closed porosity, barium titanate compositions contain sintering

aids,which form a liquid phase at sintering temperatures. The liquid phase

formed by the presence of small amounts of sintering additives in the basic

composition accelerates densification and grain growth by rearrangement and

solution-precipitation processes (7).

Acommon liquid phase forming aid in BaTiOs based compositions is Ti02. The

well known phase diagrams by Negas et al (8) and O'Bryan and Thomson (9)

indicate a eutectic in BaTiOs-excess Ti02 compositions at 1320 and 1312°C,

respectively. This temperature is further lowered to 1240°C by the presence of

impurities such as Si02 and AI2O3 (10). The Hquid phase triggers discontinuous

grain growth which results in a coarse structure. Such structures can lead to low

breakdown strength and mechanical weakness (11).

Discontinuous grain growth is characterised by rapid growth of a limited amount

of grains to sizes much larger than those of the average grain population (7). It may originate from many possible causes such as wide initial particle size

distribution, a small amount of liquid phase, chemical heterogeneity,

heterogeneous density distribution, intrinsic boundary properties (preferential

grain boundary mobihty) and others.

The discontinuous grain growth observed to occur in the presence of a small

amount of liquid phase is beleived to be caused by the combined effects of the

heterogeneous distribution of the Hmited amount of hquid, and the presence of

5
grains much larger than the average population. The growth mechanism is described as Ostwald ripening caused by the difference in solubility for particles

of different size, which causes matter to be transported from the smaller to the

larger particles. Discontinuously growing grains grow at a constant rate, until

they impinge on other rapidly growing grains and the growth rate abruptly slows

down to a value encountered in normal grain growth. Thus, the resulting grain

size after prolonged sintering is controlled only by the number of secondary

grains that are nucleated and not by the time of heating (12).

Barium titanate is one of the first ceramic materials for which discontinuous

grain growth was recorded (12). When triggered, discontinuous growth is very

rapid. Growth rates on the order of 10 /im per minute were observed directly

under a heating microscope (13). Since the final coarse grain size may be quite

homogeneous, discontinuous grain growth may go unnoticed when only final

microstnictures were examined. Fig. 1. illustrates rapid growth and shows

typical microstnictures developed during sintering of BaTiOs with 2 mol %

excess Ti02.

Extensive grain growth by a solution-preciptation process helps in rapid chemical homogenisation, decreases permittivity below the Curie point maximum and

increases temperature dependence of the maximum. Therefore, special additives are usually incorporated in dielectric compositions to prevent grain growth.

However, hquid phase sintering is in general difficult to control. The effect of

additives may depend on processing conditions, as demonstrated in BaTi03-

SrTiOs and BaTiOs-CaTiOs systems (14).

Figs 2a-2c show the influence of 8 mol % CaTiOs and SrTiOs additions,

respectvely, on the microstructure of BaTiOa - 2 mol % Ba6Tii7O40 ceramics

sintered 24 hours at 1350°C. The addition of CaTiOs strongly increases the final

grain size, and the addition of SrTiOz reduces it. Whereas reduction of grain

size due to additives is a common phenomenon in ceramics, the coarsening

influence of CaTiOs is unusual. In contrast, it was reported (15) that CaTiOs

addition to PTCR ceramics based on a BaTiOs-PbTiOs solid solution containing, besides Y2O3 dopant, 2 mol % Si02 and 0.5 mol % TiOi inhibits the grain

growth. It is obvious that the influence of CaTiOs on microstructure development depends on the grain growth mechanism, firing conditions and composition.

Examination of microstructures after short sintering times revealed that the

addition of 8 mol % CaTiOs to the BaTiO3-Ba6Tii7O40 composition greately

reduced the nucleation rate of discontinuously growing grains. These grains

could grow unhindered to comparatively large size, the maximum size of the

grains exceeding 300// m. On the other hand, addition of 8 mol % SrTiOs caused

a high nucleation rate with consequent early impingement of discontinuously

6
5 10 20 60

TIME (min )

Fig 1: Average intercept length vs. sintering time and corresponding

microstmctures for BaTiOs - Ti02 excess ceramics, smtered at 1350 C

7
Fig. 2: Microstructures of BaTiOs ceramics after heat treatment 24 hours at 1350°C (a) BaTi03-2 mol % Ba6Tii7O40, (b) BaTiOs - 2 mol % Ba6Tii7O40 - 8 mol % CaTiOs , (c) BaTiOs - 2 mol % Ba6Tii7O40 - 8 mol % SrTiOs

8
growing grains. The microstructure remained fine grained even after prolonged

sintering.

The final grain size is influenced by the relative values of the rate of nucleation

N and the rate of growth of nuclei G. If the ratio of N to G is large, the final

grain size will be small. Conversely, if the ratio of N to G is small, the final grain

size will be large. Both rates strongly depend on processing parameters such as the composition, the amount of liquid phase and temperature.

Analysis of microstructure evolution in undoped and doped BaTiOs samples at various temperatures revealed that the grain growth rate was not particularly

sensitive towards increase in temperature, whereas the nucleation rate strongly

increased, as already noted (16). Consequently, the ratio N/G increases with

temperature resulting in finer final microstructures. It was demonstrated that a

finer grain size may be achieved by avoiding the critical temperature region in

which ratio N/G is low. Similar temperature dependence of grain growth was

reported by Matsuo and Sazaki (10), who observed exaggerated grain growth in

BaTiOs doped with TiOz, Si02 and AI2O3 in the temperature region of 1240-

1250°C. In their case, the eutectic temperature was estimated to be at 1240°C.

The reason for contrasting influences of SrTiOs and CaTiOs addition on the

nucleation rate of discontinuously growing grains in Ti02-excess BaTiOs

ceramics seems to be the different chemistry in the BaTiOs-TiOi-SrTiOs and

BaTi03-Ti02-CaTi03 systems.

The systems differ from each other in several aspects. Relevant for this discussion are the facts that whereas BaTiOs and SrTiOs exhibit soHd solubility at all compositions, BaTiOs and CaTiOs are soluble only to a Hmited extent. The

second major difference is the existence of a Ca-stabilized Ba2Ti50i2 phase in

the BaO-CaO-Ti02 system, which was not found in the BaO-SrO-Ti02 system

(17).

TEM analysis of sintered BaTiOs - 2 mol % Ba6Tii7O40 - 8 mol % SrTiOs and

BaTiOs - 2 mol % Ba6Tii7O40 - 8 mol % CaTiOs revealed the presence of a Ca-stablized Ba2Ti50i2 phase in the Ca-containing ceramics. Such a phase

could not be identified in Sr- containig ceramics.

It was proposed that the close similarity of BaTiOs and Ca- stabilised Ba2Ti50i2

causes absorption of the Ca-containing phase on the surface of BaTiOs grains

thus diminishing the nucleation rate. In the concentration range studied, the

effect is the same as that caused by the decrease in the amount of Ba6Tii7O40-

BaTiOs eutectic, which accounts for the coarse grained structure after

prolonged sintering. On the other hand, addition of SrTiOs does not destabilize

the Ba6Tii7O40 phase; it causes a monotonic decrease in the grain size of the

(Ba,Sr)Ti03 ceramic by accelerating the nucleation rate.

9
Besides Ti02 a large number of other additives which form a liquid phase at lower temperatures have been investigated. In early work various low melting

glasses were found very effective in decreasing the sintering temperatures at the

expense, however, of considerable lowering of permittivity. Later on, glass

compositions were adjusted to titanates to permit some grain growth during

sintering. During grain growth, partial homogenisation of the constituents takes

place by dissolution of some ions from the glass into the titanate lattice, thereby

lowering the volume fraction of the low dielectric constant phase. By this means,

several low fired MLC dielectrics with permittivity around 4000 were prepared

(18).

4. EFFECTS OF CHEMICAL HETEROGENEITY ON SINTERING AND

GRAIN GROWTH IN FERROELECTRIC CERAMICS

In preparation of dielectric ceramics from chemically heterogeneous mixtures

care should be taken to minimise the effects of high temperature reactions on

sintering and grain growth. In general, volume changes and microstructure

development during sintering of chemically heterogeneous compacts are mainly

controlled by interdiffusion. Analysis of diffusion phenomena which lead to

solid solution and compound formation in solid state sintering of initially

heterogeneous mixtures revealed that counterdiffusion with equal diffusion

rates such as described by Wagner (19) is merely a special case. In general,

diffusion coefficients of interdiffusing species in chemically heterogeneous

systems are not equal. Differences in diffusion fluxes cause formation of vacancies on the side of the atoms whose diffusivity is larger. Excess vacancies

precipitate at contacts between the particles and form additional porosity in the

powder compact during sintering. The effect is more pronounced in sintering

of large particles. Formation of additional porosity may cause swelling of the

compact instead of shrinkage during sintering.

In perovskite systems, swelling was observed during sintering of compacts made

of PbTiOs-PbZrOs powder nuxtures, which form a complete solid solution (20).

Expansion was more extensive when coarse powders of PbTiOs were used.

Many small cavities were seen to exist in the structure. It was suggested that

preferential diffusion of Ti occured during firing and the expansion was qualitatively explained as a Kirkendall effect.

EPMA analysis of PbTiOs-PbZrOs diffusion couples after firing at 900°C

confirmed preferential diffusion of Ti ions into PbZr03 (21). Since line analysis

did not show an appreciably higher PbO concentration of Ti02 depleted

PbTiOs, it was concluded that PbO is transported with Ti ions, most probably

10
through the gas phase. This is supported by the fact that a Hquid phase could not

be detected in sintered samples. If present in excess, PbO should be detected

locally.

In sintering of BaTiOs-SrTiOs mixtures (22), preferential diffusion of Ba ions

caused formation of the nonequilibrium polytitanate phase Ba6Tii7O40 which

in turn gave rise to a low temperature eutectic liquid with BaTiOs. Transient

eutectic liquid was proposed as a reason for discontinuous grain growth, which

was observed at contacts between sintered SrTiOs-BaTiOs powder layers at temperatures well below the hquid phase formation temperature in the BaTiOsSrTiOs system. Large Sn-xBaxTiOs soHd solution grains such as shown on

Fig. 3. are believed to be flux grown. In spite of liquid phase formation,

BaTiOs-SrTiOs compacts sintered to lower final density than pure compounds.

Expansion instead of shrinkage was also observed in sintering of a BaTiOsBaZrOs mixture. Analysis of neck areas formed during sintering of BaTiOs

spheres on a BaZrOs plate by EPMA revelaed preferential diffusion of Ti"^^

into BaZrOs (Fig. 4).

As in the case of sintering of BaTiOs-SrTiOs mixtures, nonequilibrium phases

and additional porosity are formed because of unequal diffusion rates. It is interesting to note that X-ray powder examination of the sintered mixture

indicated preferential diffusion in the opposite direction, i.e., Zr"^"^ into

BaTiOs. After heating 1:1 mole BaTiOs-BaZrOs mixtures for 15 minutes at 1100°C, X-ray reflections belonging to BaTiOs almost disappeared, while the

intensity of pure BaZrOs reflections remained strong.

High temperature firing resulted in complete solid solution. It may be easily

shown that the type of unreacted material depends on the ratio of the grain sizes

of BaTiOs and BaZrOs. In drawing up such conclusions, it must be remembered

that average grain size may be misleading. Even a small amount of coarse BaZrOs particles, say 10 %, when present, may remain unreacted after a prolonged sintering time, giving rise to a false impression of predominant

diffusion of Zr ions (23).

5. REACTION SINTERING AND MICROSTRUCTURE DEVELOPMENT

IN BaTiOs-CaZrOs MIXTURES

Among the most widely used compositions for achieving a permittivity of 6000

and above are those based on the system BaTiOs-CaZrOs (2,24-27). Both

compounds are not compatible. On heating, BaTiOs and CaZrOs react to form

a barium zirconate phase with some Ca and Ti in solid solution and a calcium

11
Fig. 3: Discontinuous grain growth in 95 w/o SrTiOs - 5 w/o BaTiOs ceramic

sintered at 1450°C, 3 hours

titanate phase with some dissolved BaTiOs (29). BaTi03-BaZr03 form a sohd

solution Ba(Tii-yZry)03 over the whole composition range. The solubility of Ca

in the Bai-xCax Tii-yZryOs solid solution was reported to be 17,5 mol % at 1300''C (25) and increases to 28 mol % at 1600°C (28). Zr"*"^ ions shift the Curie

temperature maximum to the vicinity of room temperature, whereas Ca ions are effective in broadening the maximum (27).

In our experiments (29), reaction between BaTiOs and CaZrOs was detectable

by X-ray analysis already after 1 hour's heating at 1100°C. At this temperature,

reaction is slow and no heat effect could be detected on DTA curves. It was

demonstrated by dilatometric analysis that chemical reaction interferes with the

sintering process. Shrinkage curves presented in Fig. 5 show that BaTiOs started

sintering at around 1100 C and sintered to high density in the temperature

region 1300-1350°C; CaZrOs, being more refractory needed higher sintering

temperatures. It began to shrink at around 1250°C and sintered to high density

in the temperature region 1450-1500°C. BaTiOs - 16 mol % CaZrOs started to

shrink at approximately the same temperature as BaTiOs; however, after a few

% shrinkage the kinetics slowed down and temperatures of 1550-1600°C were

needed for densification. Microstructural examinations showed high porosity

which was ascribed to the Kirkendal effect due to differences in the diffusion

coefficients of various constituents in the BaTiOs- CaZrOs ternary mixture.

12
Fig. 4: Microprobe analysis of BaTiOs spheres sintered to BaZrOs plate at 1400°C, 3 hours. ES electron image, Ti, Ba, Zr X-ray image

Sinterabiltiy was improved by avoiding solid state diffusion. As shown on Fig.5,

sinterability of BaTiOs - 16 mol % CaZrOs - 2 mol % Ti02 was improved by

reactive liquid formation so that sintering temperatures of 1350-1380°C were

sufficient to achieve high density - above 95 % of theoretical value in 1 hour

firing time. Common additives which lowered the eutectic temperatures (i.e.

Si02 and AI2O3), decreased sintering temperature to below 1300°C. 16 mol %

ZrTiOs effectively reduced discontinuous grain growth to an average intercept

length 3-4 pim (Fig.6), The solution- precipitation process resulted in a coreshell type structure (1,3,30,31), as shown on Fig. 7. It is beheved that the

compositional inhomogeneities within the grains is related to the known flat- tening of the dielectric constant vs. temperature characteristic of BaTiOsCaZrOs dielectric ceramics. Compositions are sensitive to firing conditions.

Small increases in sintering temperature and time increase the permittivity and

its temperature dependence due to diffusional homogenisation.

13
800 900 1000 1100 1200 1300 1400 1500

Temperature 1°C|

Fig. 5: Dilatometric curves for (1) BaTiOs - 2 mol % Ti02 (2) BaTiOs - 2 mol % Ti02 - 8 mol % CaZrOs (3) BaTiOs - 8 mol % CaZrOs (4) CaZrOs.

Heating rates lO^C/min

I . I I I I i 1 '

' ^

0 2 4 6 8 10 12 14 16 18 20

molVo CaZr03

Fig. 6: Average grain size vs. composition in BaTiOs - 2 mol % Ti02-CaZr03

ceramics sintered 1 hour at 1360°C

14
Ba,Ti Baji

Zr Ca , Cu Zr Cu

-F:i=255 •FS= ch 51

iy.5 >> IS cx-s;

Fig. 7: TEM photomicrograph and corresponding EDS spectra of 90 mol %

BaTi03-8 mol % CaZrOi - 2 mol % Ti02 ceramic, fired 2 hom-s at 1260 C.

(A) SoUdified TiOi-rich phase at grain comers and along grain boundaries,

(B) Ca and Zr modified domain free region and (C) ferroelectric grain core.

15
REFERENCES

1. D.A.Payne, pp 413-431 in Processing of Heterogeneous Ceramics for

Dielectric Applications, Mat.Sci.Res. 2Q Plenum New York 1986

2. G.Goodman, Ceramic Capacitor Materials, pp 79-138 in Ceramic Materials

for Electronics, ed. by R.C. Buchanan, Marcel Dekker, New York 1986

3. D.Hennings, Barium Titanate Based Ceramic Materials for Dielectric Use,

IntJ.High Technology Ceramics 1 91-111 (1987)

4. M.Kahn, D.P. Burks, LBurn and W.A.Schulze, Ceramic Capacitor Technology, pp 191-274 in Electronic Ceramic, ed. by L.M.Levinson, Marcel Dekker, New York 1988

5. G.A.Smolenskii, J.Phys.SocJpn.Suppl. ZS, 26 (1970)

6. S.M.Neirman, The Curie Point Temperature of Ba(Tii-xZrx)03 SoHd Solutions, J.Mat.Sci. 22 3973-3980 (1988)

7. W.D.Kingery, H.K.Bowen, D.R.Uhlmann, Introduction to Ceramics, Wiley,

New York 1976 498-

8. T.Negas, R.S.Roth, H.S.Parker and D.Minor, Subsolidus Phase Relations in

the BaTi03-Ti02 System, J.Solid State Chem. ^ 297 (1974)

9. H.M.O'Bryan Jr. and J.Thomson Jr., Phase Equilibria in the Ti02-Rich

Region of the System BaO-TiOi, J.Am.Ceram.Soc. 51 (12) 522-526

(1974)

10. Y.Matsuo, H.Sasaki, Exaggerated Grain Growth in Liquid Phase Sintering

of BaTiOs, J.Am.Ceram.Soc 54 (9) 471 (1971)

11. J.M.Herbert, Ceramic Dieletrics and Capacitors, Gordon and Breach

Publ., New York 1985, p 86

12. J.E.Burke, "Grain Growth in Ceramics", pp 109-116 in Kinetics of High

Temperature Processes, ed. by W.D.Kingery, Wiley, New York 1959

13. H.Heydrich, quoted by H. Schmelz and H. Thomman, cfi/Ber. DKG

No. 4/5 199-205 (1984)

14. D.Kolar, V.KraSevec, "Influence of Additives on Discontinuous Grain

Growth During Liquid Phase Sintering of Barium Titanate", pp 375-380

in Sintering '89, ed. by S.Somiya, M.Shimada, M.Yoshimura and

R.Watanabe, Elsevier, Amsterdam 1988

15.W.Y.Howng and C.McCutcheon, "Electrical Properties of Semiconducting

BaTi03 by Liquid Phase Sintering", Cer. Bull ^, 231-243 (1983)

16. D.Hennings, "Recrystallization of Barium Titanate Ceramics", Sci. Ceram.,

12405-409(1984)

17. W.Kwestroo, H.A.M. Paping, 'The Systems BaO-SrO-Ti02, BaO- CaOTi02 and SrO-CaO-TiOz", J.Am.Cer.Soc. 42 (6) 292-299 (1959)

18. 1.Burn, Flux-Sintered BaTiOs Dielectrics, J.Mat.Sci. 12 1398-1408 (1982)

16
19. C.Wagner, J.Phys.Chem. B_M209 (1936)

20. Y.Nakamura, S.Chandratreya and R.M.Fulrath, "Expansion During the

Reaction Sintering of PZT', Ceramurgia Int. ^57-60 (1980)

21. M.Kosec and D.Kolar, "PZT Solid Solution Formation from PbZrOs and

PbTiOs", pp 421-427 in "Ceramic Powders", ed. by P.Vincenzini, Elsevier,

Amsterdam 1983

22. D.Kolar, M.Trontelj and Z.Stadler, "Influence of Interdiffusion on Solid

Solution Formation and Sintering in the System BaTiOs-SrTiOs",

J.Am.Cer.Soc. ^470 (1982)

23. D.Kolar, "Sintering in Chemically Heterogeneous Systems", Ceramic

Transactions, in print

24. M.McQuarrie and F.W.Behncke, "Structural and Dielectric Studies in the

System (Ba,Ca)(Ti Zr)03, J.Am.Cer.Soc. 22 539-543 (1954)

25. G.H.Jonker, Kondensatormaterialien mit Hoher Dielektrizitatskonstante,

Philips Tech. Rundschau 12 127-135 (1955)

26. O.Drexler, B.R.Schatt, "Development of Ceramic Materials With a

Dielectric Constant of 10000 at Room Temperature", Science of

Ceramics 1, 239-254 (1962)

27. H.J.Hagemann, D.Hennings and R.Wernicke, "Ceramic Multilayer

Capacitors", Philips Tech.Rev. 41, No 3, 89-98 (1983/84)

28. D.Heimings and H.Schreinemacher, 'Temperature Dependence of the

Segregation of Calcium Titanate from Solid Solutions of (Ba,Ca)

(Ti Zr)03 and its Effect on the Dielectrics Properties", Mat.Res.Bull. 22,

1221-1226 (1977)

29. D.Suvorov and D.Kolar, "Sintering and Dielectric Properties on the System BaTiOs-CaZrOa", pp 2.314-2.318 in Euro-Ceramics 2, ed. by G. de

With, R.A.Terpstra and R.Metselaar, Elsevier 1989

30. B.S.Rawal, M.Kahn, and W.R.Buessem, "Grain Core - Grain Shell Structures in Barium-Titanate Based Dielectrics", pp 172-88 in Advances in

Ceramics, Vol. 1, Am.Cer.Soc, Columbus, Ohio 1981

31. D.Hennings and G.Rosenstein, Temperature-Stable Dielectrics Based on

Chemically Inhomogeneous BaTiOs, J.Am.Cer.Soc. §1, 4, 249-254 (1984)

17
DISCUSSION

R. J. Cava: Did you imply we do not understand the reasons for the effect of grain sizes

upon the dielectric constant?

D. Kolar: I did not imply this. It is commonly observed that with decreasing grain size

in BaTi03 and BaTi03-based dielectric ceramics, the dielectric constant below the Curie

temperature increases whereas the maximum at the Curie temperature decreases and

broadens. The dielectric constant increase below the Curie temperature with decreasing

grain size is attributed to increasing internal stress due to a smaller number of domains

per grain. A decrease in the Curie point dielectric maximum and a broadening of the

peak width in small grain size ceramics is attributed to lower degree of crystallinity,

amorphous outer layers and local chemical inhomogeneity of the grains.

What I commented was that the experimental data vary. Whereas most of the authors

find very shallow dielectric maximum in small grain size ceramics (about 1 mm), a high

dielectric maximum of ~ 1 1000 was also reported (K. Kinoshita and A. Yamaji, J. Appl.

Phys. 47, 371 (1976)).

J. Nowotny: The effect of the grain size can also be discussed in terms of the different

stoichiometry of the grain boundary layers, which in the case of oxides has a certain

thickness. Now, as the grain size decreases the ratio of the boundary layer to the bulk

increases. When we speak about the different grain sizes we should also consider the

segregation of certain ions to the boundary, the grain boundaries.

For example, titanium has a tendency to segregate. But that's not the end of the story

because there is also a structural rearrangement of the boundary layer, and, for example,

it is possible that the bulk grains remain tetragonal whereas the boundary layers may be

stabilized with a cubic structure.

D. Kolar: It was also reported (D. Hennings and A. Schnell, J. Am. Ceram. Soc., 65,

539-544 (1982)) that in small grain size BaTi03 based ceramics, the structure may remain

cubic below the Curie temperature or that the cubic, tetragonal and orthorhombic phases

coexist.

A. Sleight: I want to be clear about the explanation for the liquid phase that you gave.

Were you just postulating this disproportionation from 1:1 to 1:2 and 2:1 phases or have

you some observation of the real cause of the liquid phase.

D. Kolar: Would you please repeat that? I'm not quite sure I understand the question.

A. Sleight : My question is do you have evidence that the liquid phase is caused by going

from the 1:1 compound to 1:2 plus 2: 1 compounds? Is there experimental evidence, and

can you quench these phases and see them in the specimens?

18
R. Roth: Did you actually see the polytitanate?

D. Kolar: In some cases we detected in BaTi03-SrTi03 compositions after firing faint

x-ray diffraction lines corresponding to polytitanate (Ba.,ST)^Ti^-jO^Q solid solution. Other

evidence of polytitanate was grain morphology and size of (Ba,Sr)Ti03 solid solution

grains, which were typical for flux growth.

R. Roth : Did you also observe lines from the 2: 1 compound?

D. Kolar: Faint diffraction lines of (Ba,Sr)2Ti04 were also detected. However, the main

evidence for the existence of 2:1 compound was obtained by microprobe analysis which

showed Sr-rich layer in the boundary region between BaTi03 and SrTi03 in hot pressed

BaTi03-SrTi03 diffraction couple according to the reference (D. Kolar, M. Trontelj, and

Z. Stadler, J. Am. Ceram. Soc., 65 [10] 470-474 (1982).

R. Newnham: The key point you seem to be making at the end was that barium has a

higher diffusion coefficient.

D. Kolar: Yes.

R. Newnham: Higher than strontium, it seems counter-intuitive since barium is a larger

ion than strontium.

D. Kolar: Yes.

R. Newnham: At least it seems unusual if we think of the diffusion expected in a pure

crystal. Do you believe that this is related to the defect structure? It's known, for

instance, that in barium titanate sometimes the surface of barium titanate is barium

deficient depending on the preparation technique of the material. This could give a very

high surface diffusion rate for barium. Or do you think it's through the bulk that this is happening?

D. Kolar: The pictures which I presented demonstrate surface diffusion of barium over

the SrTi03 sphere. They also demonstrate bulk diffusion of Ba into SrTi03. SrTi03 was

coarse grained and Ba diffused along grain boundaries and into the bulk of the grains.

I agree that one would expect faster diffusion of Sr ions, however, the experiments

showed the opposite.

19

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

CHEMISTRY AND PROPERTIES OF TEMPERATURE

COMPENSATED MICROWAVE DIELECTRICS

T. Negas, G. Yeager, S. Bell, and R. Amren

Trans-Tech, Inc. Adamstown, Maryland

ABSTRACT

Modem microwave detection and communication devices rely on high dielectric

constant (E'), low loss (high Q), temperature stable oxide ceramics that can be grouped

into four major classes. These are (Zr,Sn)(Ti,Sn)04, Ba2Ti9O20, BaLnTi-oxides

(Ln = Nd, Sm), and Ba3Ta2Me09 (Me = Zn or Mg). This report focuses on relevant

chemistry/structure-property-processing features for ceramics in or derived from the last three classes.

INTRODUCTION

The past decade has seen the emergence of diverse microwave technologies that

operate in the 0.4-30GHz band. Typical systems presently encompass cellular telephones,

satellite communications of many types, and frequency sensors. Global positioning

systems for personal use and high resolution/digital communication devices loom in the near future. Oxide ceramics, developed and improved over the past 30 years, are critical

elements in these microwave circuits. By simple coupling and frequency tuning

techniques, they store, filter, and /or transfer electromagnetic energy in a circuit with

minimal loss because of their high Q. They diminish the size of a device because of their

high E'. They minimize frequency drift when temperature fluctuates due to the

environment and/or to circuit heating because of their small temperature coefficients of

dielectric constant (Te'), resonant frequency (Tf), and capacitance (TJ. The industry

requires that these chemically complex ceramics be held to stringent electrical (and

dimensional) tolerances, typically ±0.5-lppm/°C and ±0.25% for specified Tf and E',

respectively. Yet, the market demands low cost for parts ranging from a few samples to millions per year. Thus, the message to the ceramist, an implicit theme of this report,

is old but worth repeating. Make processing more efficient, minimize technical problems,

and, thereby, increase product yield and lower expenses. Therefore, know your

crystallochemical system, understand how it is influenced by processing, and design

processes accordingly within technical/economic constraints. Comprehend the sampling

and electrical measurement methodologies but be aware that subtle microstructural details,

often process or raw material dependent, can yield unexpected property results. Routine

microwave measurements are especially sensitive to microstructural/chemical details but

the evidence gained is indirect and, therefore, must be linked either to empirical models

of a chemical system that can be tested experimentally or, if time/resources permit, to direct observations at the submicron level.

21
Three well established classes of ceramics are discussed. Our objective is to help

consolidate the technical base for each by detailing some simple but pertinent features that

either are alluded to or not recognized in the open literature.

Ba.Tin0.o (B.Tn) and BaTi^On (BT^): O. Tf

ZrTi04:Sn02 ceramics, developed in the late 1970's to early 1980's, have gained

acceptance by the industry because of their high Q (> lOK, 4.5GHz), high E' (35-40),

and low, adjustable Tf. The earlier candidates, B2T9 and BT4, forwarded by work in (1)

and the pioneering efforts in (2-6), lost favor primarily because Q was "too low" and,

secondarily, because Tf could not be adjusted easily. Typical literature Q's range from

5K-10K (4.5 GHz), while 2-5 and 14-20ppm/°C are approximate Tf's reported for B2T9

and BT4, respectively. Periodically, the literature informed us that Q for B2T9 can be

improved by thermal annealing, by "chemical treatment" of powders (2-5), and/or by

additions of Mn-oxide (7). Most recently (8), we learn that for unclear reasons WO3

additions increase Q for B2T9 and BT4. We have concluded that two simple chemical

factors, apart from ill-defmed intrinsic loss mechanisms related to crystal structure, really

control Q. Figures 1 and 2 demonstrate that the Q of BT4 and B2T9 is extremely high

when certain raw materials are avoided. Modest Q ceramics are obtained through the use of titanias of 99.5-1-% purity (LP), typically anatase that contains silica or phosphate at levels near 0.1%. When 99.9 + % rutile (HP) is used, Q (4.5 GHz) exceeds 13K and

12K for BT4 and B2T9, respectively, even in routine production. Well controlled

laboratory specimens yield Q as high as 14. 8K. Indeed, the data show that BT4 has a Q

somewhat higher than B2T9 and that both rival, if not exceed, the best Q's reported for

ZrTi04:Sn02. We speculate that for nominally "pure" ceramics, silicates or phosphates

of very low Q develop during densification and these reside primarily in the grain

boundary region. O'Brycin (9) broached this very issue and suggested for the case of

phosphorus that crystalline Ba-Ti-P or glassy Ti-P oxides are formed. They diminish bulk

Q disproportionately to their small apparent volume fraction and are easily detected by

the microwave measurement. If this conclusion is nearly correct some deliberate

additives (e.g. Mn-oxide, WO3) may improve a normally low Q ceramic simply by

scavenging the detrimental impurities and "improving" the grain boundary region. Our

work indicates that when the system is "clean", Q is high and additives only diminish it. Process contaminants such as milling media [e.g. AI2O3 (10)] and sodium impurities in

many cellulose binders also deteriorate Q. We return to this argument below.

Oxygen loss, particularly for B2T9, at densification temperatures, 1325-1 375 °C,

is the second major factor that influences Q. The process, operative for almost all titanate-based ceramics, is manifested by a darkened, semiconducting interior or "core",

particularly within ceramics that are of high density and large cross-section. When coring

is present, Q is low. Figure 3 (top) shows slices of B2T9 ceramics densified at 1330°C,

air. High purity (HP) and lower purity (LP) end-members (see Fig. 2) were air-quenched

from 1330°C. The HP ceramic is uniformly blue-grey, while the LP material is not as dark near exposed surfaces. The lower set of B2T9 parts in Fig. 3 have compositions

given in Fig. 2. They were densified at 1330°C in air, then cooled 50°C/hr to RT.

Note the progressive increase in "cored" volume from LP to HP ceramics and refer to

22
Fig. 1. Q of BaTi^Og ceramics versus Ti02 raw Fig. 2. Q of Ba2Ti902Q ceramics versus Ti02 raw

material content. Temperature in °C. material content and processing in 02/air.

99.9+ %Ti02 = HP and 99.5%Ti02 = LP in text.

Fig. 2 for corresponding Q in air. Bulk density is not a variable in the coring since it remains high and constant (Fig. 2) except for the most extensively cored HP end-member

which, interestingly, has a much lower density. When ceramics are processed in oxygen,

coring is eliminated and Q increases accordingly (Fig. 2).

Speculation regarding the influence of the grain boundary region on Q is given

greater validity when considered within the context of the coring observations. First,

B2T9 reduces in air during densification, > 1325°C. Second, reoxidation occurs during

cooldown but the penetration depth of

oxygen (Fig. 3) clearly is a function of

impurity content. LP end-member parts

completely reoxidize, while coring volume

increases as the HP content increases.

Given constant bulk density and grain size

(3-6 /im) for the ceramics, the

physical/chemical nature of the grain

boundaries must control the reoxidation

kinetics. Wider, "dirty" boundaries, perhaps

dominated by defects, provide avenues for

faster oxygen transport that eliminate coring,

but the penalty is low Q. Tighter, "cleaner"

boundary regions decrease the reoxidation

rate and parts core. Therefore, more

stringent processing is required to avoid

Fig. 3. Cored Ba2Ti9O20 ceramics (1330°C, air). Slow cooled (bottom); air quenched (top);

compositions from left to right correspond to materials in Fig. 2, left to right.

23
reduction and attain high Q. In a pertinent HRTEM study (11), unsuccessful attempts were made to "explain the poorly understood variations in dielectric loss" of Ba2Ti9O20

materials including a low Q ceramic from Trans-Tech, An argument was made that loss seems to "result from different degrees of polytypic intergrowth" (within B2T9 grains).

This may be correct for overall intrinsic loss but the Q of real ceramics also is influenced

by extrinsic contributions such as from grain boundaries. Indeed in (11), one HR-image

is shown for a grain boundary region and the writers remark that it is characterized by

a Ti-O framework and extended rows of vacancies (i.e. "missing" Ba). It is this

observation that meshes with our data. Impurities probably help generate the defects and

provide for localized charge compensation. In turn, the overall concentration of

defects/impurities controls Q and reoxidation kinetics. HRTEM images for low or high

Q BT4 do not show intracrystalline extended defects but they do exist in ceramics of high

and low Q B2T9 (12). Such defects, therefore, are of marginal importance in Q

considerations. However, Q for BT4 is a function of impurity content (Fig. 1), again

suggesting that the grain boundary region is critical. Reoxidation kinetics are not relevant

since BT4 does not reduce in the practical densification range, 1275-1350°C, air.

The ability to chemically tune the Tf

of dielectrics to be device specific without

diminishing Q or changing E' by more

than several units is a practical

requirement. BT4 would be the perfect

candidate to tune because of processing

ease. Unfortunately, the most obvious

substitutions for Ba and/or Ti (e.g.

Sr,Ca,Zr) do not budge Tf from a plateau

in the 13-16ppm/°C range. Solubility

appears too limited to produce the desired

effect. Even Sn solubility (see below) is less than given in (1) and is not useful.

Thus, the next most attractive option is to drive Tf toward 4ppm/°C by making

composites of BT4 with B2T9. Figure 4

summarizes the results. Unexpectedly, a Tf minimum, ~2ppm/°C, near 30m. %BT4

is observed but its significance is not clear.

We suspect that the electrical data are

reflecting a limited homogeneity region for pig. 4. Tf (25-60°C) for the system 6311409-

B2T9 not unlike that proposed in (4). A Ba2Ti9O20-TiO2.

second practical aspect of Fig. 4 is amplified in Fig. 5. For many applications, the designer often requires a Tf

approximately "linearized" over the range -30 to 60°C. Conventional ZrTi04:Sn02

ceramics show significant departure from linearity. For example, a Tf that holds at 6ppm/°C from 25-60°C changes to 9-10ppm/°C from -30 to 25°C. Figure 5 shows that

BT4-B2T9 composites retain good linearity and, indeed, traverse near perfect linearity in

the 75-60m.% BT4 range.

24
BT4 ceramics can be forced to assume Tf values 10-15ppm/°C higher than

the normal 13-16ppm/°C without chemical

adjustments. The most obvious

microstructural feature that directly

correlates is the presence of

anomalous/discontinuous grain growth

introduced into the ceramics by deliberate

bad processing practices. When random

grain growth is extensive and quite visible

(1-6 mm long, anisotropic crystals), Tf is highly positive. Details of this work are reserved for the future but are mentioned

only to target a potential pitfall for

researchers. BT4/B2T9 composites,

however, never show massive

discontinuous grain growth (B2T9 > 15

mol.%) and unusually large Tf indicating

that B2T9 is an effective inhibitor.

Fig. 5. Af/f versus temperature (Tf = Af/fAT) for compositions in Fig. 4; 25-60, -30-25 in °C. The last need is to produce BT4/B2T9 composites with Tf less positive than

2ppm/°C. Work in (5) shows that Sn02 substitutions in/near B2T9 diminish Tf. Figure

6 illustrates Tf isopleths for solid solutions in the system Ba2Ti9O20-BaTi409-

"BaTi3 5Sno 509"-"Ba2TigSn02o". The latter two are not single phase. Our results for

B2T9 are consistent with the Tf data in (5). Sn02 enters into solid solution primarily with

B2T9 as indirectly reflected by the electrical data. Microstructural evidence (Fig. 7) from

a ceramic with near zero Tf combined with qualitative EDS (S. G. Mhaisalkar, Ohio State

Univ.) confirms that B2T9 grains contain most of the Sn02. The composition shown by

the data point at "BaTij ySno 3O9" consists of B4T13, BT4, and Ti02 with Sn02

concentrated primarily in the latter. Q is very low, — 600. Thus, at high Sn02 content,

two-phase solid solution tie-lines (not shown) must traverse compositional space from

northwest to southeast in Fig. 6. Note particularly in Fig. 6 the broad trough of low Tf

that appears to be related to the minimum on the BT4/B2T9 join. As Sn02 content is increased for any starting Ti/Ba ratio, Q progressively decreases from a > 12.5K plateau

towards about 10.5K (4.5 GHz). Large pucks (2.5" OD x 1.25" t) of several Sn02-

containing ceramics were dissected to obtain numerous randomly oriented test parts (0.5"

OD X 0.25" t) that sampled surface and interior volumes. Uniform electrical properties were measured indicating that the ceramics are not anisotropic and that vaporization of

Sn02 is not a significant problem.

Figure 8 shows that good Tf linearity is maintained for B2T9 materials with Sn02

and other substitutions. Zr02 substitutions are not very effective in decreasing Tfi

coupled substitutions, for example, of Ni/Ta for Ti, effectively increase Tf of B2T9. The

coupled substitution approach appears attractive but requires further work (e.g. Me^^/Ta,

Nb; Me = Mg, Zn, Co, etc. combinations).

<]

+ 500

+ 500

+ 400

+ 300

+ 200

+ 100

0

-100

-200

-300

-400

-500

-500

-700

Composition Tf ppr n/ 1^ "m.%BT /B T„

4' 2 9 25-60 -30- 25

- © 89 m.%BT^ 10.8 12.6 © 1 1 " @ 73 ^\ 1.2 7.8 /

/ /''\,<:;?Y5^

_

' ® ^

f2)/^ Tf ppm/°C

25-50 -30-25

© 58.5 5.4 5.2 - (TV © 33 BT 2.8 1.4 © 0 8T

4 5 3.3 100 ^2^9

' 1 -1 : 1 -30 -20 - 10 0 10 20 30 40 50 50

Reference T=25°C

Reference f=4.5 GHz T(°C)

25
Ti+Sn/Ba=4.0

MOI. Ti+Sn/Ba=4.5

Fig. 6. Tf isopleths (approximate) for the system BaTi4O9-Ba2Ti9O20 andSn02-containing

solid solutions (1330-1 350°C,O2). Hatching depicts approximate limits of solid solubility

after (1) and this work (, not single phase). Open circles are after (5).

+ 50

0

-50

-100

-150

-200

-250

-300

-350

1 Tf ppm/°C

- 25-60 -30-25 ®

® 10 SaTi^Og 0.3 -0.4

-

3.3 Ba jTi g02Q

6.7 "8a2TigSn02o / ®

"(D Boj ^'8.825 ^"0.175 "^20 2.2 0.5 / /V® _(3) eojTi 8.65 Zro.jjOjo 4 2.3 / //y®

®

(D

"(^

Tf ppm/°C

25-60 -30-25 ® "Ba'T' 4.162 Zr 0.24309.8 3.4 2.6 (BT^+BjTgSolid Solutions) ® Bai.gsSroosTigssZrojsOjo 4.8 3.2

1 1 © Ba2Tlg75Nio(jaTao 16O20 ^.3

1 1 1 1 1 1

5.3

1

1 -30 -20 -10

Reference T=25°C

Reference f=4.5GHz

0 10 20

T(°C)

30 40 50 60

Fig. 7. Microstructure of BaTi4O9/Ba2Ti9O20 Fig. 8. Af/f versus temperature (Tf = Af/fAT) for composite; each phase contains Sn02; Tf = ceramics based on Ba2Ti9O20 that contains various

-I- lppm/°C (Fig. 6). Defective grains are B2T9. cation substitutions; 25-60, -30-25 in °C.

26
Ba.Ln.Ti-OXIDES: Tf. O

High E' (smaller size) ceramics of modest Q are used in memy applications in the

0.4 to ~3GHz band. Development of products lay dormant until spurred by market

needs in the early 1980' s. The technical foundation for renewed efforts resides in work

by Bolton (13), but subsequent literature, including a plethora of patents, never acknowledges the study. Bolton examined the system, BaTiOj-LnjOj-TiOj,

Ln = didydium. A composite with low loss (IMHz), small temperature coefficient of

capacitance (TJ, and high E' -62 was found when "pure" NdjOj was used. More

importantly, a broad region of higher E' (> 80) and very negative T,. was delimited and

shown to result from a phase X plus a phase Y. Figure 9 summarizes developments to

Fig. 9. Present knowledge of the BaTi03-Ti02-Ln203 system. Solid solution (hatched)

is after (20). O, low Tc, high E' material after (13). , 1:1:4 composition (Nd203) of

special focus in (14,15,16). ®, 1:1:5 composition to 1:1:4 (Sm203) of special emphasis

in (17) and this work. Tie-lines from Ln2T9 to 1:1:4 solid solution do not exist for the Sm203 system [2:9(Sm) not stable].

date. Kolar et al. (14) completely characterized the electrical properties (IMHz) of the

system containing Nd203 and later (15) showed that additions of Bi4Ti30,2 to the

compositions 1:1:5 (BaO:Nd203:Ti02) and 1:1:4 yield high E' (88-105) ceramics with

low Te'. Wakino et al. (16) next demonstrated that in the vicinity of the 1:1:4

composition (Nd), ceramics with high E' (85-88), modest Q ( ~ 2K, 3GHz), and near-zero Tf can be obtained by substituting some Pb for Ba. Parallel work in (17), and at TransTech focused on Sm-variants which yield ceramics with negative Tf, E«80, and higher

Q«3K-4K (3GHz). The former tuned Tf through zero by substituting some Sr for Ba.

We tune Tf simply by adjusting Ti02 content along the join Ti02 - 1:1:4. Meanwhile,

crystal chemists were attempting to define phase equilibrium and structural details for

these materials. Kolar et al. (18) reported on the Nd compounds 1:1:3 and "1:1:5".

Matveeva et al. (19) determined a preliminary crystal structure for a Pr-analogue near this

compositional region. Finally, in (20) we learn that the high E' compound(s) are members of a solid solution region that extends from about 3:2:9 through 1: 1:4 to about

2:2 1/3:9 (Fig. 9). The 3:2:9 end-member is Bolton's phase X, while his phase Y is close to the 1:1:3 compound. Today, the chemistry-property cycle has been nearly

Ti0

2

1:1

mol.%Ln203,Ln = Nd(Sm)

27
closed, but one major factor has been overlooked and it casts doubt on the ability of

researchers to reproduce the available property data. In every study and patent

promulgated to date, samples were made by uniaxial pressing of prereacted powders. It will be demonstrated below that this methodology complicates the electrical data.

In (19), l:l:4-type materials are shown to have a structure that contains

perovskite-like building blocks and that

crystals with orthorhombic symmetry grow

with acicular habit. During our work with

1:1:4 (Sm), we observed that a given batch

of powder, 12(X)°C calcine, can yield

highly variable Tf depending on the

forming process used and on the method of

sampling the ceramic for electrical testing.

We conjectured that, since 1:1:4

crystals/grains are anisotropic, electrical

properties should be anisotropic. If so, the

observed Tf variation must somehow be

related to dimensional crystalline anisotropy. However, this requires at least partial

preferred orientation of crystallites in the finished ceramics. Qualitative microstructural

examination and quantitative XRD analysis (J. Simmins, Alfred Univ.) of typical ceramics

do not reveal obvious preferred orientation, albeit lath/acicular crystals are present (Fig.

10). Consequently, a series of experiments was designed to address the issue.

Fig. 10. Microstructure of ceramics in the vicinity of

1:1:4-1:1:5 (SmjOj).

Five large, dimensionally precise

pucks were formed [~ 1:1:4. 5(Sm)

composition] each at a different uniaxial

pressure. Pucks were cut in half, and one

set of halves was isostatically overpressed,

> lOOMPa. All ten pieces were densified

(>5.70 g/cc, <2% porosity) at 1330°C.

Figure 11 gives initial green densities and

shows anisotropy in fired shrinkage which

we will not discuss other than to indicate

that uniaxial pressing imposes some type of

"memory" to the powder that is not erased

by isostatic overpressing cmd is carried into

densification. Figure 12 illustrates

subsequent methodology and Tf data for the

fired parts. Every half-puck was machined

to obtain a minimum of two orthogonal sets of test parts for Tf measurement using cavity

methods and the TEqis mode (21).

Uniaxial Uniaxiol +

Pressure Isostatic

16 -

10 30 50 70 90 Pressure, MPa

Fig. 11. Pressing methodology of ~l:l:4.5(Sm)

powder (plus green densities) showing anisotropy in

fired shrinkage.

28
+3

+ 2

CJ +1

o\

CL

P= -1

-2

-3

• . Uniaxial Only o , Uniaxial + Isostatic (>100Mpa)

Orientation

Orientation (D

Fired f>5.70 g/cc

Test Ports: 12.7mm OD 5.1mm Thick

10 30 50 70

Pressure, MPa

90

STRIPLINE METHOD son -LINE

Fired Port

Network Analyzer Resonont Frequency

Uniaxial Only And

Uniaxial + Isostatic

Fig. 12. Tf of fired ceramics versus methodology Fig. 13. Method to detect E' anisotropy in of pressing (Fig. 1 1) and extraction of oriented test ceramics of Fig. 12. parts for electrical measurement.

Alternatively, orthogonal TE and TM modes for only one set of identically oriented parts

may be used to test for Tf, but this complicates procedural logistics and increases possible

error. Figure 12 demonstrates that ceramics show anisotropy in Tf that widens in

magnitude as uniaxial pressure increases. Isostatically overpressed parts retain the

anisotropy imposed by the initial uniaxial pressing. Ceramics that are isostatically pressed

only do not display anisotropy and yield Tf « 0 ± lppm/°C. Anisotropy in E' was

examined using the same test pieces and the parallel plate method (22). Small variations near the limit of error were observed, so a more sensitive but indirect method was

developed. Fig. 13. Cubes containing the uniaxial pressing direction as a unique 4-fold

axis were precisely extracted and machined from all half-puck ceramics. These were

coupled to a 50 n stripline, and resonant frequency (F^) was measured at six different

cube orientations, i.e. each 4-fold axis, (+)

and (-) direction, brought to the vertical

position. If E' is isotropic, F^ should remain

constant for all cube orientations. Figure 14

shows an E' anisotropy of —0.4 units, and

the unique uniaxial pressing direction is easily discemable. The extensive data

collected show that Sm-ceramics display

anisotropy in Tf of about 6ppm/°C and in E'

of 0.4-0.7 units. The magnitude of the

anisotropy within the 1:1:4-1:1:5-1:1:5 (+2

wt% Ti02) compositional region examined

Axis JL Substrate/Slripline

V Fr(MHz)

AA' 4635.1

A'A 4636.5

BB' 4646.5

B'B 4646.5

CC 4646.5

C'C 4647.4

» Unique Axis

c

B 7B' 500

A'

AA' = Pressure Axis

AFr=11 MHz

AE'=0.4 Units

remains constant provided uniaxial pressure Fig. 14. Methodology and example measurements

also is constant. This indicates that to determine E' anisotropy by the "cube technique",

secondary phases (e.g. Ti02) contribute little ^^s29
to the anisotropy. The largest E' and most

positive Tf are observed when the TEqi^

field during measurement is normal to the

pressing direction. Almost within

measurement error (±5% maximum), Q

seems to remain constant but systematic

trends in the data suggest that slightly lower

Q is associated with the more positive Tf

orientation. The ceramic in Fig. 15 contains

0.2 wt% less Ti02 than the materials

illustrated above. Uniform but directional

properties are clearly evident and note

sample 8 which was extracted inclined to the

pressing direction.

AA—64mm

A8=-20mm(PRESSING AXIS) PReSSURe-70MPa

SECTION A-A

SAMPLE: I3mm(O0)x5.1mm(THK) SAMPLE Q(2.9GHz) Tf £• OENSITY(g/cc)

1 3480 -0.7 80.1 5.72

2 3525 -0.5 79.6 5.71 Z 3410 -0.5 79.7 5.72

4 3710 -6.5 79.2 5.71

s 37)0 -6.3 79.2 5.71

s 3710 -6.3 79.4 5.72

7 3670 -6.1 79.4 5.72 8 3700 -2.0 79.5 5.72

Fig. 15. Uniaxially pressed and fired ceramic

[~l:l:4(Sm) composition], dissected to demonstrate anisotropic electrical properties but

good overall uniformity.

Several samples ofdense, commercial

l:l:4(Nd):PbO pucks were made available

and these were tested as above. The pucks

must have been pressed uniaxially because anisotropy in Tf of 14ppm/°C and in E' of

^0.7 units are evident (Q«2300, 3GHz). BijOj-containing materials of the l:l:4(Nd)-

type (15) were fabricated by isostatic and by uniaxial pressing, then densified. The latter show anisotropy of 22ppm/°C and ~ 3 units in Tf and E', respectively. These data might

be further influenced by a Bi-ferroelectric phase that exists as domains and/or at grain

boundaries. Solubility of 31263 m 1:1:4 also requires definition. Isopressed parts are

relatively uniform, Tf « 10ppm/°C and E' « 91 (Q « 1800, 3GHz).

Electrical properties and processing for the high E' materials can be considered

within the context of an oversimplified empirical model. First, it is assumed that powder,

calcined at 1200°C and milled to an average particle size near 2 j^m, contains nuclei of

1:1:4 that have lath/acicular habit. XRD reveals that the powder consists of B2T9, Ti02,

Sm2Ti207, and 1 : 1 :4 but definitive identification of nuclei could not be obtained by SEM.

Second, we speculate that some fraction of the nuclei is aligned during uniaxial forming

with their long axis (or axes) perpendicular to the pressing direction. Most of the

alignment is accomplished at the lowest pressures. Isostatic pressing (only) does not cause preferred orientation. When it is applied to uniaxially pressed parts of low green

density, the alignment of nuclei is preserved with only minor disruption (Figs. 11,12).

During densification, preoriented nuclei grow in the plane normal to the initial pressing

direction at the expense of smaller 1:1:4 particles and of adjacent secondary phases as chemical equilibrium is attained. The rest of the ceramic also undergoes similar changes,

but they are directionally random. The resultant product consists of a matrix of random

crystallites plus a small volume fraction of oriented grains that are not easily resolved by

routine methods. However, the sensitive microwave measurements detect the anisotropy.

It is useful to reflect on one final experiment. The model assumes that nuclei of 1:1:4

exist in processed powder. Therefore, their availability and concentration should be

dictated by the magnitude of the initial calcination temperature. Consider the extreme, no calcination and no nuclei. BaTiOj, Sm203, and Ti02 (rutile) raw materials (see bulk

30
composition pertinent to Fig. 15) were blended in alcohol, dried, and reblended. Samples

of this powder were pressed uniaxially and isostatically (90 MPa) and these were

densified, 1330°C. Fired densities were 5.60g/cc (<4% porosity) sufficient to yield

E' « 77 and Q = 3.2K (3GHz). Both ceramics gave Tf = -3.1ppm/°C and the cube

test showed no frequency variation indicating complete lack of anisotropy. Although we

grant that for these powders microstructural evolution is controlled by a completely

different set of densification/reaction mechanisms and kinetics, the assumption for the

need of 1:1:4 precursors appears to have some validity.

Exact details for the model require further work and it would be useful to have

electrical data from 1:1:4 single crystals. We recognize that electrical and dimensional

anisotropy are discussed using the simple case of two primary axes. Since 1:1:4 is orthorhombic, a third axis must influence the observed properties and add complexity to the model. Many other experiments need to be performed to address questions involving

numerous variables such as calcination and densification temperature/time, extent of grain

growth, the influence of pore size/shape, etc. versus electrical anisotropy. We note,

however, that for l:l:4(Sm) ceramics the experimental options are limited by a narrow

effective densification range, 1330±15°C, without introducing other problems.

Nevertheless, some practical processing approaches can be derived from the limited data.

Anisotropy initiated by forming methods such as uniaxial pressing, extrusion, etc., can be managed by imposing strict controls in powder calcining through densification

operations. Alternatively, parts can be formed exclusively by isostatic pressing but this

is not always cost effective nor applicable to shapes such as substrates. Finally, the

nuclei in powder can be eliminated. Reactive sintering of well blended raw materials (see

above) is appropriate provided high density, target E', and control of large shrinkage are

obtained. Calcining the blended raw materials at a temperature slightly lower than that

required to form 1:1:4 crystallites is particularly attractive since powder homogeneity is enhanced by preliminary reaction and subsequent milling. The 1050-1 100°C range is optimal.

High E' materials mildly reduce during densification in air. High density, Smceramics show distinct cores, while Nd variants have cores that are masked by their

intrinsic bluish color. Coring diminishes Q to almost unmeasureable values but

processing in oxygen, 1330± 15°C, avoids the problem when using conventional powders.

MICROWAVE PEROVSKITES

As applications entered the 10-30GHz band, ceramics with super Q (> > lOK,

lOGHz) and small E' (<30) were demanded. During the 1980's, it was demonstrated

that the well-known perovskite family, Ba3Ta2Me09 (Me = Zn or Mg), has these

requisite properties (8,23,24) when octahedral cations are ordered (1:2 type) to near

perfection. However, this seemingly simple need has important practical engineering

consequences. Multiple octahedral substitutions (Ni,Co,Mn,Zr,Sn, etc.) often are used

to tune Tf. As an example. Fig. 16 illustrates ordered perovskite solid solutions in the

system Ba3Ta2.2/3xZni.i/3x.yNiyZr^09. Since ordering is diffusion controlled, many obvious

31
Fig. 16. Tf and E' isopleths (approximate) for ordered perovskite ceramics in the system

Ba3Ta2Zn09-Ba3Ta2Ni09-Ba3Zr309. Ten Kg

batches were made; fired densities are >7.70 g/cc.

Ba^Ta^ZnOg 5 lO

Ba^Zr^Og —*-mol.%(E' = 35)

factors limit the kinetics. Of these, the "mundane" blending of raw materials is the least

appreciated and, in our judgment, the most crucial yet most difficult to quantify. For the

example given, BaC03, Ta205, ZnO, NiO, and Zr02 powders having diverse physical

properties are blended by high shear wet mixing. To test "homogeneity", samples are reacted at 1250°C. Dark agglomerates are observed when homogeneity is "poor". XRD

reveals that these products are "BaNi03"-type oxides and BaNi-oxycarbonates (25); when

present, even after rehomogenization by milling, pressed ceramics rarely reach target E'

or.Q after firing. When the calcined blend lacks these and appears "uniform", there is considerable batch to batch variation in the time required for pressed parts to attain the

high Q ordered state. Typically, 12 to 100 hr at 1450-1500°C are necessary indicating

that perfect blending is elusive. Much better uniformity is achieved when the Zn, Ni, and

Zr are added to a BaC03/Ta205 slurry as nitrate and/or acetate solutions. In this case, the high Q state is obtained within 25 hours.

The association of high Q with insulating, non-ferroelectric, perovskites prompted

us to explore other systems. Ba8Ta6Li2024 ceramics were fabricated and without special

efforts to optimize, a material was obtained with E' = 28, Tf = 26ppm/°C, and

Q = 14. 5K (5.5GHz). This perovskite-like phase features ordered Ta/Li octahedra that

share comers and faces (26). Complex perovskites also were explored in work for a

radar-related project that required a dielectric with E' = 50±5, Tf = 0±20ppm/°C, and

Q > 5K (4GHz). Unpublished records at Trans-Tech describe a ceramic, nominally

La2/3Ti03, which the literature shows (27) should consist of L2L{Th^7 (^^i) P^'JS La4Ti9024 (L2T9). Properties, reproduced in Fig. 17, reveal some peculiar correlations

with processing. At constant fired density, E' increases and Q improves each time the

powder is calcined/milled. Figure 18 shows our XRD traces for the actual products.

When E' = 45 and Q is low, the ceramic is a composite of LTj and L2T9 plus a third

32
phase that dominates when E' = 68 and Q

is high. This product is an ordered

orthorhombic perovskite ( — 3.8 x 3.8 x

7.6a) similar to La2/3Ti03.x reported in

(28). Additional details indicate that the

powder was milled with AI2O3 media and

up to ~2 wt. % "pickup" was experienced.

Furthermore, the perovskite ceramics were

"cored" and of low Q unless annealed at 1100°C. Thus, it appeared that perovskite

formation could be favored by

incorporation of AI2O3 according to Two compositions,

E'

^^2/3 +X/3T1 1 -x^^lx^S

La2/3Ti03 + 2

70

65

60

55

50

45

40

Unpublished

L.P. Domingues

R. West Trans-Tech

1 976

5.22 , Q =2.500 /

T^, =-170 /

(5-6 GHz) /• 5 26 / Q= 775

r ^=5.25 Q= 300

1

g/cc

1

Wt.%Al203 (A) and Number Of Calcines/Millings

Fig. 17. "Processing'

comp>osition.

versus E' for "1^^37103'

X = 0.10 (B) were formulated (lOg) from

pure (>99.9%) oxides which were heated

in AI2O3 crucibles at 1200% at 1350°, and

finally at 1400°C, air. After each heating, samples were cooled 200°C/hr to RT and

remixed. XRD showed that (A) consists of L2T9 plus perovskite and (B) is single-phase,

ordered perovskite. Although the evidence suggests that AI2O3 helps "stabilize" the

structure, it is not sufficient to clarify the importance of reduction. Since ceramics core, mild reduction must occur but the perovskite can be reoxidized without causing

decomposition. More work is necessary to elucidate equilibrium details including solid

solution toward LaA103. Subsequently, the system La203'Ti02'^gO was examined in

the hope of finding a similar perovskite, 1^/3 +xj3^h-i/2x^^ 1/2x^2^ ^ith smaller E' and Tf. In the vicinity of 1^2/3^103 ±MgO, the phase was obtained again but ceramics were not

'Inc

0)

Fig. 18. XRD tracings for ceramics in Fig. 17 (A,B,C correspond to E' = 45, 55, 68,

respectively). A LT2 + L2T9 + trace ordered

perovskite (P, shaded); B — LT2 + L2T9 + P; C = L2T9 + P (trace LT^. LT2 = La2Ti207;

L2T9 = La4Ti9024.

11 12 26 28 30 32

26, CuKa

33
fabricated since the XRD data from powders were plagued by many uncertainties.

Fortunately, at x = 1, LaTi,/2Mg,/203 (LTM) was found to be an orthorhombic perovskite

similar to GdFeOj. Ceramics yield E' = 33, Tf « -lOOppm/^C, and Q ~ 8K (5GHz).

Possible LTM-CaTiOj (E' « 150, Tf « 800ppm/°C) solid solutions thus became obvious

candidates and at 44 mol. % LTM an orthorhombic perovskite was obtained with E' =

47, Tf = 0, and Q = 7.4K (3.8GHz). All properties met project requirements.

ACKNOWLEDGMENT

We thank all the employees of Trans-Tech that made this report possible. D. Cronin, L. P. Domingues, R. West, and R. Huntt contributed valuable discussions. A. Arnold, D. Huntt, and K. Anderson (with his swamp group) provided their machining, forming, and powder preparation expertise.

Appreciation is extended to M. Clevinger of NIST for organizing and preparing this manuscript.

REFERENCES

1. G. H. Jonker and W. Kwestroo, J. Am. Ceram. Soc. 41, 390 (1958).

2. H. M. O'Bryan, Jr., J. K. Pluorde, and J. Thomson, U.S. Patent 3,938,064 (1976).

3. H. M. O'Bryan, Jr. and J. Thomson, U.S. Patent 4,337,446 (1982).

4. H. M. O'Bryan, Jr. and J. Thomson, J. Am. Ceram. Soc. 66, 66 (1983).

5. H. M. O'Bryan, Jr., J. K. Pluorde, and J. Thomson, U.S. Patent 4,563,661 (1986).

6. D. W. Readey, E. A. Maguire, C. P. Harting, and D. J. Masse', Raytheon Co., Contract DAAB07-

69-C-0455 (1971).

7. S. Nomura, Ferroelectrics 49, 61 (1983).

8. S. Nishigaki, S. Yano, H. Kato, T. Hirai, and T. Nonomura, J. Am. Ceram. Soc. 71, C-11 (1988).

9. H. M. O'Bryan, Jr., Am. Ceram. Soc. Bull. 66, 677 (1987).

10. S. G. Mhaisalker, W. E. Lee, and D. W. Readey, J. Am. Ceram. Soc. 72, 2154 (1989).

11. P. K. Davies and R. S. Roth, J. Solid State Chem. 71, 490 (1987).

12. S. G. Mhaisalker, Masters Thesis, Ohio State Univ. (1987) and personal communication.

13. R. L. Bolton, Dissertation, Univ. 111. (1968).

14. D. Kolar, Z. Stadler, S. Gaberscek, and D. Suvorov, Ber. Dt. Keram. Ges. 55, 346 (1978).

15. D. Kolar, S. Gaberscek, Z. Stadler, and D. Suvorov, Ferroelectrics 27, 269 (1980).

16. K. Wakino, K. Minai, and H. Tamura, J. Am. Ceram. Soc. 67, 278 (1984).

17. S. Nishigaki, H. Kato, S. Yano, and R. Kamimura, Am. Ceram. Soc. Bull. 66, 1405 (1987).

18. D. Kolar, S. Gaberscek, B. Volavsek, H. S. Parker, and R. S. Roth, J. Solid State Chem. 38. 158

(1981).

19. R. G. Matveeva, M. B. Varfolomeev, and L. S. Il'yushchenko, Russ. J. Inorg. Chem. 29, 17 (1984).

20. F. Beech, K. Davis, A. Santoro, R. S. Roth, J. L. Soubeyroux, and M. Zucchi, Abstract 251-B-87,

Aim. Meeting Am. Ceram. Soc. (April 1987).

21. D. Kajfez and P. Guillon, Dielectric Resonators, Artech House, Norwood, MA, 1987 (pp. 539).

22. E. Courtney, IEEE Trans. Mic. Th. and Tech. MTT-18. 476 (1970).

23. S. Kawashima, M. Nishida, I. Ueda, and H. Ouchi, J. Am. Ceram. Soc. 66, 421 (1983).

24. K. Matsumoto, T. Hinga, K. Takada, and H. Ichimura, Proc. 6th IEEE Int'l. Symp. App. of

Ferroelectrics, p. 118 (June 1986).

25. T. Negas and R. S. Roth, NBS Spec. Publ. 364. 233 (1972).

26. T. Negas, R. S. Roth, H. S. Parker, and W. S. Brower, J. Solid State Chem. 8, 1 (1973).

27. J. B. McChesney and H. A. Sauer, J. Am. Ceram. Soc. 45, 416 (1962).

28. M. Abe and J. Uchino, Mat. Res. Bull._9, 147 (1974).

34
DISCUSSION

D. Smyth: I'd like to suggest another effect of purity on your oxidation problem. Most

of the naturally occurring impurities have a lower valence, impurities like iron,

aluminum, and these are acceptor impurities. As the acceptor impurity content increases,

the transition from oxygen excess insulating material to oxygen deficient semiconductor

material is shifted to lower oxygen activities. So, you essentially have shifted your

reduction problem to a lower oxygen activity.

T. Negas: I'll buy that.

D. Smyth: And there is another bonus to that. Because acceptor impurities are compensated by oxygen vacancies and these increase your reoxidation kinetics you obtain a double advantage. That would be a bulk effect as opposed to your grain boundary

effect.

T. Negas: I indicated that I used 99.9 plus percent rutile which contains very low levels

of aluminum impurities. It actually is a very good, clean titania that results in very high

Q ceramics. The ceramics with less pure titania show faster reoxidation kinetics but are of low Q.

D. Smyth: That's what I'm suggesting. The less pure material behaves better because

it contains more acceptor impurities.

T. Negas: I understand what you're saying.

D. Kolar: Is it possible that your impure material was simply not reduced as much as the

pure?

T. Negas: Okay, that's a good point. If I can have the overhead. You wish to know

what X is in the Ba2Ti9O20.x ceramics. There it is (overhead shows X < 0.001) and it

is difficult to measure either the pure or impure material with conventional TGA. One

has to fill a crucible with a hundred grams of powder, then quench it and reoxidize the

powder. Both show low levels of reduction ~ the x value is very small even though you

see highly darkened material, very small.

D. Kolar: The simple way to show this would be to cool down your darkened materials

in a neutral atmosphere.

T. Negas: We have also done that.

D. Kolar: Yes? It's dark, this impure material?

T. Negas: Of course.

35
D. Kolar: Okay.

T. Negas: Ba2Ti902o~you cool it down in nitrogen or argon and it is dark. When you

pull it out of the kiln-straight out—it's blue. The less pure material starts to reoxidize

and the edges, the leading edges of the ceramic start turning a lighter color but the

interior is blue. However, the high purity material is all blue. Again, it's the reoxidation

kinetics as related to the grain boundary region that control the reoxidation process which

in turn relates to the overall Q of the material. The two are related, the impurities, the

oxidation phenomenon, I think they're really tightly correlated.

P. Davies: On the dirty materials that people use, bismuth or lead are added to affect the

temperature coefficient of resonant frequency. Does it generally drop, and if so, can you

give me a crystal chemical explanation as to why that happens to go in that direction?

Is there a good chemical reason why they have that particular effect?

T. Negas: You mean in tuning the temperature coefficient of resonant frequency?

P. Davies : Yes.

T. Negas: In these sort of things, theoretical approaches are practically useless as far as I'm concerned. There aren't that many but we have a lot of empirical data;

measurements, a lot of measurements for many materials. For example, if you look at lead titanate, the temperature coefficient of dielectric constant is something like plus

5600 ppm/°C. It stands to reason that if you have a highly negative material, I could put

in lead titanate, for example, in solid solution using simple mixing rules. You can adjust

chemical composition such that the temperature coefficient of dielectric constant and

resonant frequency go to zero. So you pick and choose materials, pick and choose based on what you know about crystal chemistry and v/hat you know about available properties

because there isn't much theory around. There is some, but it's not theory that guides

you. Unfortunately, it always comes afterward.

You can talk about Q and theoretial views on Q looking at optical frequency resonances and gamma factors which are related to lattice vibrations and crystal structure. We can

talk about that, but these are summation kind of things—a summation over all atoms, all molecules, all of the grain boundary regions, etc. It's tough and you get some guidance—

but it's the measurements, I feel that count, at least in the short term.

R. Roth : I have one question. Back in about 1980 we obtained specimens of the

tetratitanate from Trans-Tech. When looked at with high resolution electron microscopy

the material was highly defective and showed a lot of polytypic defects. The materials

you showed did not show the defects. Are you implying that the difference between what

we saw in Trans-Tech material in 1980 and what you see in 1988 or 1989 may be due

to differences in purity of starting materials?

T. Negas: You are not looking at a difference in impurity content because the

tetratitanate grains, despite what raw materials you use, remain clear. We do not see, at

36
least a group at Ohio State does not see, this gross defectiveness in the tetratitanate. On

the other hand, I should caution you about samples that you used.

R. Roth : But these came from Trans-Tech.

T. Negas : Right. You have to remember that if you really look into what you got, it is the standard Trans-Tech low Q tetratitanate material. It is not pure barium tetratitanate.

It's about 90 mole percent barium tetratitanate, 10 mole percent barium nonatitanate.

Going through and looking at the high resolution TEM you have to make sure you're

looking at the right grains because there is 10 mole percent barium nonatitanate. I know

you showed me pretty pictures of polytypism and defects, but until you guys publish that

and convince me, I'll stand pat.

R. Roth : I noticed that Peter Davies published the paper on the 1:5 and 2:9 compounds

but never wrote up the 1:4.

P. Davies : It was definitely 1:4. There is no doubt about it.

T. Negas: I believe you also have had similar discussions with the group from Ohio State

and I guess you guys seem to be in some disagreement. I'm stuck in the middle and I've seen the microstructures that have been taken with the TEM, and the evidence there

seemed to be that, at least the limited number of pictures that I've seen, the tetratitanate seems to be fairly defect free. Perhaps the resolution isn't that great. I don't know.

R. Roth : Maybe what we should do is look at your materials.

T. Negas: I've been wanting you guys to do that but I did go to the Ohio State group.

Here, I'll give you a sample now. Here's my high Q composite, and there you should

see tetratitanate and nonatitanate grains-both types. But let me tell you. Peter and Bob

had a beautiful paper on lattice imaging or high resolution work on nonatitanate. I read

the paper in 1987 and I was looking at beautiful pictures. I went from figures 1 through

13 and I kept saying, golly, get me to the grain boundary region. Finally in figure 14

in that publication they show this region and then say a few words about it, then drop it. You had it right there, I think. They showed in the grain boundary region strings of

titanium, oxygen and a lot of missing barium. More than likely this area is charge

compensated by impurities but they dropped the discussion after that. This is precisely

what I'm talking about, the reoxidation kinetics being controlled by the grain boundary

region. I think your data are in line with, at least, our very simplistic observations.

37

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

TUNABLE TRANSDUCERS:

NONLINEAR PHENOMENA IN ELECTROCERAMICS

Robert E. Newnham

Materials Research Laboratory

Pennsylvania State University

University Park, PA 16802

ABSTRACT

The concept of a fully tunable transducer is developed as an example of a very smart material whose sensing and actuating functions can be tuned to optimize behavior. By constructing the composite transducer from an elastically

nonlinear material (rubber) and an electrically nonlinear material (relaxor

ferroelectric), most of its key properties can be adjusted over wide ranges by

applying DC bias fields or mechanical prestress. These properties include

resonant frequency, acoustic and electric impedance, damping factors, and

electromechanical coupling coefficients.

The origins of nonlinear properties are considered briefly using

electrostriction as an example. The nonlinearities usually involve phase

transformations and size-dependent phenomena on the nanometer scale.

INTRODUCTION

Much of the recent interest in nonlinear phenomena stems from the desire

to build "smart materials" for intelligent systems. The words "smart materials"

have different meanings for different people, and can be smart in either a passive sense or an active sense. Passively smart materials incorporate self-repair

mechanisms or stand-by phenomena which enabl e the material to withstand

sudden changes in the surroundings.

Ceramic varistors and PTC thermistors are passively smart materials.

When struck by lightning or otherwise subjected to high voltage, a zinc oxide

varistor loses its electrical resistance and the current is bypassed to ground. The

resistance change is reversible and acts as a stand-by protection phenomenon.

Barium titanate PTC thermistors show a large increase in electrical resistance at the ferroelectric phase transformation near 130°C. The jump in resistance enables

the thermistor to arrest current surges, again acting as a protection element. The

39
R(V) behavior of the varistor and the R(T) behavior of the PTC thermistor are both highly nonlinear effects which act as standby protection phenomena, and

make the ceramics smart in a passive mode.

A smart ceramic can also be defined with reference to sensing and

actuating functions, in analogy to the human body. A smart ceramic senses a change in the environment, and using a feedback system, makes a useful

response. It is both a sensor and an acutator. Examples include vibration

damping systems for space structures and energy-saving windows for homes and

factories. The new electronically-controlled automobile suspension systems

using piezoelectric ceramic sensors and actuators consdtutes an actively smart

material.

By building in a leaming function, the defmition can be extended to a higher level of intelligence: A very smart cerat iic senses a change in the

environment and responds by changing one or more of its property coefficients. Such a material can tune its sensor and actuator functions in time and space to optimize behavior. The distinction between smart and very smart materials is essentially one between linear and nonlinear properties. The physical properties

of nonhnear materials can be adjusted by bias fields or forces to control response.

TUNABLE TRANSDUCER

To illustrate the concept of a very smart material, we describe the tunable

transducer recendy developed in our laboratory. Electromechanical transducers

are used as fish finders, gas igniters, ink jets, micropositioners, biomedical

scanners, piezoelectric transformers and filters, accelerometers, and motors.

Four important properties of a transducer are the resonant frequency f, the

acoustic impedance Za, the mechanical damping coefficient Q, the

electromechanical coupling factor k, and the electrical impedance Zg. The

resonant frequency and acoustic impedance are controlled by the elastic constants

and density, as discussed in the next section. The mechanical Q is governed by

the damping coefficient (tan 5) and is important because it controls "ringing" in

the transducer. Electromechanical coupling coefficients are controlled by the

piezoelectric coefficient which, in tum, can be controlled and fine-tuned using

relaxor ferroelectrics with large electrostrictive effects. The dielectric "constant"

of relaxor ferroelectrics depends markedly on DC bias fields, allowing the

electrical impedance to be tuned over a wide range as well. In the following

sections we describe the nature of nonlinearity and how it controls the properties

of a tunable transducer.

ELASTIC NONLINEARITY: TUNING THE RESONANT FREQUENCY

Information is transmitted on electromagnetic waves in two ways:

amplitude modulation (AM) and fi-equency modulation (FM). There are a number

of advantages to FM signal processing, especially where lower noise levels are important.

40
Signal-to-noise ratios are also important in the ultrasonic systems used in

biomedical and nondestructive testing systems, bui FM-modulation is difficult

because resonanr frequencies are controlled by stiffness (c) and sample

dimensions (t):

Neither c, t, or the density p can be tuned significantly in most materials, but

rubber is an exception. To tune the resonant frequency of a piezoelectric

transducer, we have designed and built a composite transducer incorporating thin

rubber layers exhibiting nonlinear elasticity.

Rubber is a highly nonlinear elastic medium. In the unstressed

compliant state, the molecules are coiled and tangled, but under stress the

molecules align and the material stiffens noticeably. Experiments carried out on

rubber-metal laminates demonstrate the size of the nonlinearity. Young's

modulus (E = 1/siiii) was measured for a multilayer laminate consisting of

alternating steel shim and soft rubber layers each 0.1 mm thick. Under

compressive stresses of 200 MN/m^, the stiffness is quadrupled from about 600

to 2400 MN/m^. The resonant frequency f is therefore double, and can be

modulated by appUed stress.

Rubber, like most elastomers, is not piezoelectric. To take advantage of

the elastic nonlinearity, it is therefore necessary to construct a composite

transducer consisting of a piezoelectric ceramic (PZT) transducer, thin rubber

layers, and metal head and tail masses, all held together by a stress bolt.

The resonant frequency and mechanical Q of such a triple sandwich

structure was measured as a function of stress bias. Stresses ranged from 20 to 100 MPa in the experiments. Under these conditions the radial resonant

frequency changed from 19 to 37 kHz, approximately doubling in frequency as predicted from the elastic nonlinearity. At the same time the mechanical Q

increases from about 1 1 to 34 as the rubber stiffens under stress.

The changes in resonance and Q can be modeled with an equivalent

circuit in which the compliance of the thin, rubber layers are represented as capacitors coupling together the larger masses (represented as inductors) of the

PZT transducer and the metal head and tail masses. Under low stress bias, the

rubber is very compliant and effectively isolates the ^T transducer from the head

and tail masses. At very high stress, the rubber stiffens and tightly couples the

metal end pieces to the resonating PZT ceramic. For intermediate stresses the

rubber acts as an impedance transformer giving parallel resonance of the PZT - rubber - metal -radiation load.

It is interesting to compare the change in frequency of the tunable

transducer with the transceiver systems used in the biological world. The

biosonar system of the flying bat is similar in frequency and tunability to our

41
tunable transducer. The bat emits chirps at 30 kHz and listens for the return

signal to locate flying insects. To help it differentiate the return signal from the

outgoing chirp, and to help in timing the echo, the bat puts an FM signature on

the pulse. This causes the resonant fi^equency to decrease from 30 to 20 kHz near

the end of each chirp. Return signals from the insect target are detected in the ears of the bat where neural cavities tuned to this frequency range measure the time

delay and flutter needed to locate and identify its prey. Extension of the bat

biosonar principle to automotive, industrial, medical and entertainment systems is obvious.

PffiZOELECTRIC NONLINEARITY: TUNING THE

ELECTROMECHANICAL COUPLING COEFFICIENT

The difference between a smart and a very smart material can be

illustrated with piezoelectric and electrostrictive ceramics. PZT (lead zirconate

titanate) is a piezoelectric ceramic in which the ferroelectric domains are aligned in a poling field. Strain is linearly proportional to electric field in a piezoelectric

material which means that the piezoelectric coefficient is a constant and cannot be

electrically tuned with a bias field. Nevertheless it is a smart material because it can be used both as a sensor and an actuator. PMN (lead magnesium niobate) is not piezoelectric at room temperature

because its Curie temperature lies near 0°C. Because of the proximity of the

ferroelectric phase transformation, and because of its diffuse nature, PMN

ceramics exhibit very large electrostrictive effects. The nature of this large

nonlinear relationship between strain and electric field, and of its underlying

atomistic origin, will be described later.

Electromechanical strains comparable to PZT can be obtained with

electrostrictive ceramics like PMN, and without the troubling hysteretic behavior

shown by PZT under high fields. The nonlinear relation between strain and

electric field in electrostrictive transducers can be used to tune the piezoelectric

coefficient and the dielectric constant.

The piezoelectric ^33 coefficient is the slope of the strain-electric field curve when strain is measured in the same direction as the applied field. Its value

for Pb (Mgo.3Nbo.6Tio.i) O3 ceramics is zero at zero field and increases to a maximum value of 1300 pC/N (about three times hrger than PZT) under a bias

field of 3.7 kV/cm.

This means that the electromechanical coupling coefficient can be tuned over a very wide range, changing the transducer from inactive to extremely active.

The dielectric constant also depends on DC bias. The polarization saturates under

high fields causing decreases of 100% or more in the capacitance. In this way the

electrical impedance can be controlled as well.

42
Electrostxictive transducers have already been used in a number of

applications including adaptive optic systems, scanning tunneling microscopes,

and precision micropositioners.

To summarize, two types of nonlinearity are utilized in the fully tunable

transducer: elastic nonlinearity and piezoelectric nonlinearity. By incorporating

thin rubber layers in an electrostrictive transducer several important properties can be optimized with bias fields and bias stresses. Electromechanical coupling

coefficients and electric impedance are tuned with electric field, and mechanical

damping, resonant frequency, and acoustic impedance with stress bias.

RELAXOR FERROELECTRICS

In the three remaining sections we consider the atomistic origin of

nonlinear behavior in ceramics. The magnesium - niobium distribution in PMN

ceramics have been studied by transmission electron microscopy. Images formed

from superlattice reflections show that the size of the regions with 1:1 ordering is approximately 3 nm (=30A). The ordered regions are small islands separated by narrow walls of niobium - rich PMN.

A simple energy argument explains the scale of the ordered regions.

The chemical formula of PMN can be divided into 1:1 ordered regions with

niobium-rich coatings as follows:

Pb3MgNb209 = (Pb2MgNb06)- + (PbNb03)+

assuming all ions have their usual valence states, the ordered regions will be

negatively charged and the coatings positive. The ordering must therefore be on a very fine scale in order to minimize coulomb energy.

A cubes model is adopted to estimate the size of ordered islands. Let

each island be a cube n unit cells on edge. Then there will be v?l2 Mg atoms and

n3/2 Nb atoms within the island. This means that n^ll Nb atoms are excluded

from the island because the Mg:Nb ration is 1:1 in the ordered island and 1:2 in

the chemical formula. We assume the excluded niobiums form a thin monolayer

coating on the cubes then n^l2 = 3n^, where 3n^ is the number of surface atoms

per cube, neglecting edges and corners. By keeping the coating as thin as possible, charge is neutralized on a local scale, and electrostatic energy is minimized.

Solving the equation gives n=6, or 24A, surprisingly close to the 30A

observed by TEM. It is also interesting to compare this result with the behavior

observed of small ferroelectric pardcles. X-ray studies of very fine-grained

PbTi03 show that the polar tetragonal phase becomes unstable below about

200A, the so-called ferroelectric - superparaelectric transformation. Relaxor

ferroelectrics like PMN exhibit many of the characteristics of superparaelectric

solids where the dipole moments are strongly coupled to one another, but not to a

43
crystallographic axis. The coupled electric dipoles oscillate in orientation and

respond readily to applied fields giving rise to large dielectric constants and

massive electrostriction coefficients.

SEVEN MISCONCEPTIONS ABOUT ELECTROSTRICTION

Perhaps the best way to describe the basic features of nonlinear physical

properties is to first dispel some of the common misconceptions. Some of the

key ideas are described in this section using electrostriction as an example.

Misconception #1: Electrostrictive strain is proportional to the square of the

electric field. When an electric field E is applied to an insulator it develops a polarization

P and a strain e. In tensor notation the strain can be written as a power series in E

or in P.

(1) Eij = dijkEk + Mijki EkEi +

(2) eij = dijkPk + Qijkl PkPl + . . . .

The first term in both equations represents piezoelectricity, the second

electrostriction. Piezoelectricity is a third rank tensor property found only in

noncentrosymmetric materials. It is absent in most ceramic materials, but

electrostriction is not. The electrostriction coefficients Mjjkl or Qijki constitute a fourth rank tensor which, like the elastic constants, are found in all materials,

regardless of symmetry. In the discussion which follows, we deal with

centrosymmetric media for which the piezoelectric coefficients are zero.

Returning to equations (1) and (2), which of these equations is the correct

way to describe electrostriction, or can both be used? For normal low permittivity

materials, the polarization P is proportional to the applied electric field E, and

therefore both expressions are correct, but for high permittivity solids only one is correct. Under high fields, the polarization saturates in high permittivity

materials, especially in ferroelectric ceramics just above Tc, the Curie temperature.

In capacitor dielectrics, it is not unusual to observe a decrease in dielectric

constant of 100% under high electric fields. In this situation P is not proportional

to E, and therefore electrostrictive strain cannot be proportional to both E^ and p2.

When strain is plotted as a function of electric field for PMN-PT, a

typical electrostrictive transducer ceramic, it is not proportional to E^ except for

small fields where the shape is parabolic. When plotted as function p2, a straight

line is obtained for the e(P2) relationship. Therefore the correct way to formulate

electrostriction is

44
(3) Eij = QijklPkPl (i,j, k, 1 = 1,2,3)

The Mijki coefficients in eq. (1) are not good constants except in linear dielectrics.

It is instructive to convert the fourth rank tensor coefficients to the more

manageable matrix form

(4) ei=QijPj2 (i,j=l-6)

When written out in full for cubic point group m3m, this becomes

(5)

ei Qll Ql2 Ql2 o 0 0 Pj2

£2 Ql2 Qll Ql2 o 0 0 P22

£3 = Ql2 Ql2 Qll o 0 0 P32

£4 O O 0 Q44 0 0 P2P3

£5 O o 0 0 Q44 0 P3P1

£6 O 0 o 0 0 Q44 P1P2

In this expression ei, £2 and £3 represent tensile strains along the [100], [010],

and [001] axes, respectively. £4, £5 and £5 are shear strains about the same three axes. To understand the structure-property relationship underlying

electrostriction, it is helpful to visualize the atomistic meaning of coefficients Qn,

Q12, and Q44. The high temperature structure of PbTi03 is pictured in Fig. 1 a. Lead ions are located at the comers of the cubic unit cell, titanium ions at the

body-centered position, and oxygens at the face-center positions.

When an electric field is applied along [100], polarization component Pi

develops along the same axis. This in turn causes the unit cell to elongate by an amount Aa, and the resulting tensile strain is £1 = Aa/a. Electrostrictive

coefficient Qn is equal to eifPi^. The drawing ir» Fig. lb makes it clear why

electrostrictive strain is proportional to electric polarization rather than electric

field. Polarization and strain are both controlled by atomic displacement.

The atomistic meaning of Q12 and Q44 are pictured in figs. 1 c and 1 d,

respectively.

45
(sT Oft

1

Pb ,0 - O If AO.

^^ x ! 6 6 6

I?

CUBIC Qr42:=Q„P,'^

PEROVSKITE ^« CL

(a) (b)

(0 ' ^

(d)

Fig. 1 Electrostriction in cubic pervoskite showing the physical origin of

coefficients Qn, Q12, and Q44-

46
Misconception #2. Large voltages are required to observe electrostriction.

This misconception arises from the fact that historically, the first electroceramics were high voltage insulators made from low permittivity oxides

such as porcelain, glass, steatite, and alumina. To prevent breakdown, the

insulators were fabricated in large sizes, typically with 1-10 cm separation

between electrodes. Under these conditions the voltages required to induce

measurable electrostriction are huge, but the picture has changed in the last decade.

Two factors have made the difference: smaller electrode separation and

higher permittivity. The introduction of thin and thick film technology, together

with tape-casting, has reduced the thickness of the dielectric constants to less than

50 |im with corresponding increases in the electric field levels. Fields of

megavolts/meter are common under these circumstances. At the same time, new

ferroelectric ceramics with dielectric constants in excess of 10,000 have raised the

polarization levels to new heights, further contributing to electrostriction strain.

As a result, strains of lO^^ to 10-2 are observed with relatively modest voltages.

Misconception #3 Electrostriction coefficients are about the same size for all materials.

Misconception #4 Electrostriction coefficients are largest in ferroelectric solids

with high dielectric constants.

Qll coefficients measured for normal oxides are several orders of

magnitude or more larger than those of ferroelectrics.

Qll(m4/C2)

Silica +12.8

Barium Titanate +0.11

Lead Magnesium Niobate +0.009

Misconception #5 Materials with large electrostrictive coefficients produce the

biggest strains.

The Q coefficients are largest for non-ferroelectrics but relaxor

ferroelectrics like PMN have the largest strains. The following table lists typical

values of the electrostriction coefficients Q, the dielectric constant K, and

electrostrictive strain e computed for a field of 1 MV/m.

Q K

non-ferroelectric -10 -10 -10-7

normal ferroelectric -10-1 -103 -10-5

relaxor ferroelectric -10-2 -104 -10^

47
Note that the product QK~100 m^/C^ is approximately a constant for all three classes. To see why this is so, consider the following argument based on

the atomistic models in Fig 1.

Under the action of an applied electric field, the cations and anions in a

crystal structure are displaced in opposite directions by an amount Ar. It is this

displacement which is responsible for electric polarization (P), the dielectric

constant K, and the electrostrictive strain e. To a first approximation, all three are proportional to Ar. The electrostriction coefficient O is therefore proportional to 1/K:

n E . (Ar) 1 , 1

This means then, that QK is approximately constant, as indicated earlier.

Materials with high permittivity have small electrostriction coefficients but large

electrostrictive strain.

Misconception #6. Unlike piezoelectricity there is no converse electrostriction

effect.

In the direct piezoelectric effect an applied mechanical stress produces an electric polarization Pj = dijk Oki- The converse effect relates mechanical

strain to applied electric field, = dyk Ek. The piezoelectric coefficients dijk are

identical in the direct and converse effect because both originate from the same

term in the fill energy function:

F = + dijk Pi cJjk +

Electrostriction and its converse effects arise from the free energy term

F = + Qijkl Pi Pj Oki +

Taking the partial derivatives in different order leads to three equivalent

effects. The first is the normal electrostriction effect, the variation of strain with

polarization:

dP. dP. do, , V kl J)

kl dP dP.

I J

= QIjkl

The second is the stress dependence of the dielectric stiffness (3ij, better known as

the reciprocal electric susceptibility:

48
d ( d f dF \\

kl = Q.

ijkl

The third effect is the polarization dependence of the piezoelectric voltage

coefficient gjki:

Thus there are two converse effects for electrostriction, and there are three ways

of evaluating the electrostrictive coefficients.

The three effects are used in three different applications: (1) the

electrostrictive micropositioner, (2) a capacitive stress gauge, and (3) a fieldtunable piezoelectric transducer.

Misconception #7 Thermal expansion effects make electrostrictive

micropositoners and stress gauges impractical.

This is true for normal oxide insulators but not for relaxor ferroelectrics.

For a normal oxide such as silica, the electrostrictive coefficient Q-lOm^/C^, the

dielectric constant K~10, and the thermal expansion coefficient a-lO'^K"!. For

such a material the strain produced by an electric field of 1 MV/m is about 10"^,

which is equivalent to a temperature change of only 10'^ degrees. This would

pose severe problems in an actuator or micropositioner, but the situation is quite

different for a ferroelectric.

For PMN, Q ~ 10-2, K ~ 10^ and a ~ lO-^K-l. Therefore the

electrostrictive strain for a field of 1 MN/m is about lO^^ which is equivalent to a temperature rise of 1000°! Relaxor ferroelectrics sometimes have an abnormally

low thermal expansion coefficient near the diffuse phase transformation where the

dielectric constant and the electrostrictive strain are unusually large.

Micropositioners made from ferroelectric ceramics are not troubled by small

variations in temperature.

ORIGINS OF NONLINEARITY

What do nonlinear materials have in common? The passively- smart PTC

thermistor and ZnO varistor have grain boundaries a few nm thick, insulating

barriers that can be obliterated by the polarization charge accompanying a

ferroelectric phase transformation, or, in the case of the varistor, insulating

boundaries so thin they can be penetrated by quantum mechanical tunneling.

= Q.

Ijkl

49
Small size is also key factor in the nonlinear behavior of semiconductors.

The thin gate region in a transistor allows charge carriers to diffuse through

unimpeded. The p-region in an n-p-n transistor is thin compared to the diffusion

length in single crystal silicon. Similar size-related phenomena are observed in

quantum well structures made from GaAs and Gai-xAlxAs. Planar structures with nm-thick layers show channeling behavior of hot electrons in clear violation

of Ohm's law. Current - voltage relationships are highly nonlinear in many

submicron semiconductor structures. Ohm's law is a statistical law which relies

upon the assumption that the charge carriers make a sufficiently large number of

collisions to enable them to reach a terminal velocity characteristic of the material.

When the size of the conduction is sufficiently small, compared to the mean free

path between collisions, the statistical assumption Ireaks down, and Ohm's law

is violated.

Nonlinear behavior is also observed in thin film insulators where even a modest voltage of 1-10 volts can result in huge electric fields of 100 MV/m or more. This means that thin film dielectrics experience a far larger field than do

normal insulators, and causing the polarization to saturate and electric permitivity

to decline. Because of the high fields, electric breakdown becomes a greater

hazard, but this is partly counteracted by an increase in breakdown strength with

decreasing thickness. This comes about because the electrode equipotential

surfaces on a thin film dielectric are extremely close together, thereby eliminating

the asperities that lead to field concentration and breakdown.

The influence of nanometer-scale domains on the properties of relaxor

ferroelectrics has already been made plain. Here the critical size parameter is the

size of the polarization fluctuations arising from thermal motions near the brood

ferroelectric phase transformation in PMN and similar oxides. The ordering of

Mg and Nb ions in the octahedral site of the PMN structure results in a chemical

inhomogeneous structure on a nm-scale, and this, in turn, influences the size of

the polarization fluctuations. Tightly coupled dipole • within each Nb-rich portion

of this self-assembling nanocomposite behave like a superparaelectric solid. The

dipoles are strongly coupled to one another but not to the crystal lattice, and thus

they reorient together under the influence of temperature or electric field, this in turn causes the large electric permittivities and large electrostrictive effects found

in relaxor ferroelectrics.

The importance of nanometer- scale fluctuations and the instabilities

associated with phase transformations is also apparent in the nonlinear elasticity

of rubber and other polymeric materials. The thermally-assisted movement of the

randomly oriented polymer chains under tensional stress results in large

compliance coefficients, but rubber gradually stiffens as the chains align with the

stress into pseudo-crystalline regions. The increase in stiffness with stress gives

rise to sizeable third order elastic constants in many amorphous polymers. The

effect depends markedly on temperature. On cooling, to lower temperatures

rubber and other amorphous polymers transfoixn from a compliant rubber-like

material to a brittle glass-like phase which is of httie use in nonUnear devices.

50
Nonlinear behavior is also observed in magnetic and optical systems.

Superparamagnetic behavior, analogous to the sunerparaelectric behavior of

relaxor ferroelectrics, is found in spin glasses, fine powder magnets, and

magnetic cluster materials. As in PMN, the magnetic dipoles are strongly coupled

to one another in nanometer- size complexes, but are not strongly coupled to the

lattice. Superparamagnetic solids display nonlinear magnetic susceptibilities and

unusual "AE" effects in which Young's modulus E can be controlled by magnetic

field. The effect is especially large in metallic glasses made from Fe-Si-B-C

alloys. The cluster size in spin glasses is in the nanometer range like those in

PMN.

Lead lanthanum zirconate titanate (PLZT) perovskites can be prepared as transparent ceramics for electrooptic modulators. Quadratic nonlinear optic

behavior are observed in pseudocubic regions of the phase diagram which show

relaxor-like properties.

In summary, the nonlinear properties of electroceramics are often

associated with nanometer-scale structure and diffuse phase transformations.

Under these circumstances the structure is poised i n the verge of an instability

and responds readily to extemal influences such as electric or magnetic fields, or mechanical stress.

The ready response of nonlinear ceramics allows the properties to be

tuned in space or time to optimize the behavior of the sensor-actuator systems

referred to by some as "very smart ceramics."

ACKNOWLEDGEMENTS

Numerous discussions with Q.C. Xu, M. Blaskewiecz, L.E. Cross, K.

Uchino, S.J. Jang and S. Yoshikawa are gratefully acknowledged. Carol Foster

typed the manuscript with great skill.

51
DISCUSSION

P. Davies: Concerning the relaxor ferroelectric with cation ordering in the nanodomains,

the charge imbalance between the ordered and disordered regions bothers me a little bit. Are we sure there are no oxygen deficiencies in these domains? Why are the ordered

domains stable?

R. Newnham: Well, that's a good question. If you examine various ordered

configurations, say by model-building, you fmd the one-to-one pattern goes together very

nicely because there are six unit cells around each perovskite unit cell, making it easy to pair things off. If you try, two-to-one ordering corresponding to the overall composition,

you fmd it is necessary to position niobium ions in neighboring octahedral sites. This

violates Paulings electrostatic valence rule and leads to charge imbalance on the atomic

scale.

What I'm pointing out is that if you look at other possible ordering schemes there are problems with them, too, so it comes down to which is least unfavorable.

P. Davies: Are we sure there are no oxygen deficiencies in the domains?

R. Newnham: Well, I'll let Prof. Smyth comment on that, but I can tell you that these

relaxors are excellent insulators which will withstand very high voltages, and I think if there were a lot of oxygen vacancies which were present we would be seeing conduction

and degradation phenomena which are not observed.

D. Smyth : Well, that's a question of continuity. You might have some motion within

the domain, but whether it would transport throughout the entire system, I don't know.

I don't think we have any direct evidence on that because it is difficult to probe these

very small areas.

D. Kolar: What is the effect of grain size?

R. Newnham: The samples we have prepared have normal grain sizes on the micron

scale. There have been studies of small grain sizes in ferroelectric ceramics and studies

of nanocomposite materials showing transitions from the multidomain regime, to single

domain, and to the superparaelectric regime. You have to ask in what temperature range

and in which size regime are you operating. In the high temperature regime, the

electrostrictive material behaves like a normal paraelectric, then on cooling, there is a broad diffuse phase region with relaxor phenomena, and then at low temperatures it is a normal ferroelectric with normal domain structure. Then PMN behaves pretty much like

other ferroelectrics, but it is this broad diffuse range which comes about because of

nanometer scale structure, and which we think is analogous to similar phenomena

observed when ferroelectrics are prepared in very small particle sizes. This is what is often called superparaelectricity with very large dielectric constants and very large

electrostrictive effects.

52
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

STRUCTURAL AND ELECTRONIC PROPERTIES

OF SOME PEROVSKITES

P.D. BATTLE

INORGANIC CHEMISTRY LABORATORY

SOUTH PARKS ROAD

OXFORD OXl 3QR

ENGLAND

INTRODUCTION

The perovskite structure is one of the best known structures in oxide

chemistry. It is adopted by compounds having the general formula ABO3

where A is usually (but not always) an alkaline earth or one of the larger

lanthanides and B is usually a transition metal or one of the smaller

lanthanides. It is possible to have more than one element occupying the A

or B sites in a compound; for example, in BaLaNiRuOg the A sites and B

sites are respectively occupied by random distributions of Ba/La and Ni/Ru.

These compounds can be represented by the general formula AA'BB'Og. In

this paper we shall focus on the coordination geometry around the B site

cations. In an ideal cubic perovskite (Figure 1) these cations occupy an

octahedral site, 6-coordinated by oxygen. (The larger A-site cations are 12

coordinated by oxygen.) I shall regard the BOg octahedron as the basic

building block of these strucmres. The cubic perovskite structure can then

be built up by the corner-sharing of these units. There are other

compounds, having hexagonal symmetry, which also have the basic

formula ABO3. The B site cation is again 6-coordinated and these

compounds are usually referred to as hexagonal perovskites. However, in

these compounds the octahedra are linked together by both comer-sharing

and face-sharing. Figure 2 shows the structure of SrMn03 which contains

both types of octahedral linkage. The B-O bond length in diese compounds

is of the order of 2A; the B-B distance is thus ~4A when the octahedra are linked by comer-sharing and ~2.5A when they are linked by face-sharing.

The extent of face sharing increases as the ratio of the ionic radii r^/r^

increases.

The electronic properties of perovskites are very variable. For

example, SrFe03 is a metallic antiferromagnet, SrCo03 is a metallic

ferromagnet, LaFe03 is an insulating antiferromagnet and LaNi03 is a metallic Pauli paramagnet. The transition metals in perovskites can all take more than one oxidation state. This leads to the possibility of nonstoichiometry on the oxygen sublattice, i.e. the formation of AB03^. The

oxygen content of the Sr-Co-O system, for example, can vary from SrCo03

(Co ) to SrCo02.5 (Co^"*"). The variation in the number of d electrons on

the transition metal cation can have a profound effect on the electronic

properties of the system, as demonstrated by the change from metallic

53
54
ferromagnelism in SrCoOg to insulating antiferromagnetism in SrCo02.5.

The variation is governed by the oxygen content, which is itself governed

by the method of preparation. In other words, the electronic properties of

perovskites are controlled by chemistry.

SrCoOg.y

The points made above are well illustrated by a more detailed

discussion of the phases in the system SrCoOg.y. These are all prepared

from the reaction

3SrC03 + C03O4 + n02 SSrCoOg.y + 3CO2

but under a variety of conditions as described below:

(i) heat at 1200°C in air and quench to room temperature gives

SrCo02.5 (y=0.5) i.e. Co^^ only

(ii) heat at 1200*^C and slow cool in air gives SrCo02 63 i-^- Co^"^

and Co^-^

(iii) heat as in (i), but then heat at 800*^C under argon gives

SrCo02.5 still The crystal structiu-e of the product from (i) is that of a brownmillerite

- essentially an oxygen deficient perovskite with the anion vacancies

ordered in such a way as to give alternate layers of 6- and 4-coordinate

cations. The crystal structure of the product from (ii) and (iii) is a 'rhombohedral' phase which will be discussed in more detail later in this

paper. The magnetic properties of the three products are as described

below:

brownmillerite SrCo02.5 - antiferromagnetic at room temperature

'rhombohedral' SrCo02.5 - antiferromagnetic below 25K

'rhombohedral' SrCo02.63 - paramagnetic at 4.2K

The dependence of the magnetic properties on the crystal structure and the

oxygen content, both determined by the conditions of synthesis, is thus

clear.

I write inverted commas aroimd "rhombohedral" because although

many groups have made this phase of SrCo02.5, nobody has ever got a

satisfactory unit cell for it, never mind a crystal structure. The X-ray

powder diffraction pattern strongly suggests rhombohedral symmetry, but

the fit between observed and calculated line positions is less than

satisfactory. EXAFS measurements have been able to shed some light on

this structural problem. Figure 3 shows the orthorhombic brownmillerite

structure of high temperature SrCo02.5. Although anion vacancies are present, the structure is still very similar to that of a perovskite with a B-O distance of ~1.9-2.0A, B-A at ~3.4A and B-B distance of ~4A. The

observed and calculated Co K-edge EXAFS of this phase are shown in

55
Figure 3: The brownmillerite structure

B-SfzCojOs

(Quenched from 1200°C)

Fourier filtered EXAFS

1InA 111 '

' B w

\ / \ '/ \ ^

A I \ J \ P

HI

F.T

Ai

1 1 t 1 1 1

the Fourier transform

6 8 10 12

ptiotoelectron wave number, k(A ') 12 3 4 5

distance, r(A)

Figure 4: The Co-K EXAFS spectra of brownmillerite SrCo02.5

56
Figure 4, along with their Fourier transforms, and it is clear that the

coordination geometry described above is consistent with the observed

data. The EXAFS data on 'rhombohedral' SrCo02.5 are shown in Figure

5. The most striking difference between the environments of the Co atoms

in the two phases is the presence of a coordination shell at a radius of

2. 46A in the rhombohedral phase. Figure 6 shows the observed and

calculated Fourier-filtered EXAFS from this shell, modelled with the

assumption that the back-scattering atom is Co. These data thus show a CoCo distance of 2.46A in the low temperature phase of SrCo02.5,

suggesting that, in contrast to the high temperature brownmillerite form,

face sharing octahedra are present in this phase. The EXAFS of the

partially oxidised composition SrCo02.56 show an additional coordination

shell at ~2.8A, consistent with the presence of edge-sharing octahedra,

although this is not well-established (1). The structural complexities of

SrCo03.y are still not fully resolved, but it is clear that EXAFS can make

an important contribution in some cases when diffraction methods fail. Sr2LaFe308^y

Sr2LaFe308^.y is another perovskite-like system showing a wide range of

electronic behaviour. The end member (y=0) Sr2LaFe308 is an Fe"^"*" compound which is antiferromagnetic at room temperature. It can be

prepared by firing stoichiometric quantities of SrC03, La203 and Fe203

under vacuum at 1200''C (2). The crystal structure (Figure 7) is

intermediate between those of brownmillerite and perovskite, consisting of

double octahedral layers separated by single tetrahedral layers. When the

reaction mixture is heated in air and slow cooled to room temperature the

product has the composition Sr2LaFe308 94 (3). The average oxidation

state of the Fe atoms in this compoimd is 3.63. The Mossbauer spectrum

(Figure 8) suggests that above 230K all the iron atoms are electronically

equivalent. This is consistent with the crystal structure refined from

neutron powder diffraction data collected at room temperature (4), which

contains only one crystallographic site for Fe in a perovskite which has

undergone a small rhombohedral distortion. The observed and calculated

diffraction patterns are shown in Figure 9. The Mossbauer spectra

collected below 230K suggest that the material orders magnetically at 220K

and that the Fe simultaneously di^roportionates into two oxidation states,

ideally described as Fe"^"*" and Fe^ .

i.e. 3Fe^-^^ 2Fe^-^ + Fe^"^

The areas of the two sextets in the low temperature spectrum are consistent

with this interpretation. We expected this disproportionation to produce a

structural distortion and we therefore repeated our neutron diffraction

experiments at 50K. The observed and calculated diffraction patterns are shown in Figure 10. Additional Bragg scattering is apparent, particularly

at low angle, because of the presence of long-range antiferromagnetic order

at 50K. The crystal structure could still be defined in the same

rhombohedral space group with only one Fe site, suggesting that, perhaps

surprisingly, there was no ordering of the Fe^^ and Fe^"*" cations.

However, consideration of the magnetic Bragg scattering shows this to be a

false conclusion. Perovskites are known to order magnetically in a number

of different ways, each of which gives rise to a characteristic neutron

57
H-SrzCozOs (Slow cooled from 800°C)

Fourier filtered EXAFS the Fourier transform

6 8 10 12

photoelectron wave number, k(A ')

1 2 3 4 5

distance. r(A)

Figure 5: The Co-K EXAFS spectra of 'rhombohedral' SrCo02.5

The fit to the Fourier-filtered second shell EXAFS shows that the backscattering atom is Co

6 8 10 12

photoelectron wavenumber, k(A ')

Figure 6: The Fourier-filtered EXAFS of the second

coordination shell in 'rhombohedral' SrCo02.5

58
Figure 7: The crystal structure of Sr2LaFe308

r 230 K

220 K

78 K

-8 -4 0 4 8

velocity (mm s"')

Figure 8: Mossbauer spectra of Sr2LaFe308 94 as a function

of temperature

59
*500

90 100 no 120 TWO THETA

Figure 9: Neutron diffraction pattern of Sr2LaFe308 94

at room temperature

CD

UJ

S500

5000

'iSOO

<iOOO

3500

3000

2500

2000

1500

1000

500

0 jJjLjUi

50(C P-3C1 P-Volgt

II I II I II I II II I I I 1 I I II I I II II I II HI I III I I I 10 II II

-~\ A -V -f f j f -H^ -J I I I L_

I III II n

J I I L

20 JO 40 50 60 70 60 90 100 110 120 130 KO TWO THETA

Figure 10: Neutron diffraction pattern of Sr2LaFe308 94 at 50K

60
diffraction pattern. The distribution of magnetic scattering in the

diffraction pattern of Sr2LaFe308 94 differed from these standard forms,

and was eventually accounted for using the magnetic structure shown in

Figure 1 1 . The diagram shows a network of comer sharing FeO^ octahedra

with the [111] axis of the rhombohedral unit cell running from top to

bottom. The magnetic ordering involves the formation of ferromagnetic

sheets perpendicular to this axis. The magnetic moment per cation is not

the same in all sheets, and those with the larger moment (3.61 1X3) are assumed to contain Fe"^"*" cations whilst those with the smaller moment

(2.72 (Xg), which are only half as numerous, are assvuned to contain Fe^"*". The relative orientations of the magnetic moments in neighbouring sheets is such as to ensure anti ferromagnetic coupling between pairs of Fe"^"*" ions,

but ferromagnetic coupling between Fe^"*" ions and Fe^"*" ions. This is

consistent with the different electronic configurations of the two ions

(Fe^-^:3d^, Fe^-^:3d^). The surprising conclusion to emerge from these

results is that the magnetic diffraction data indicate that the Fe"^"*" and

Fe^"^ ions are ordered (into sheets perpendicular to [111]) despite the fact

that there is no crystallographic distortion. It is likely that the

description of this compound in terms of integral oxidation states is an

oversimplification, and we are really dealing with Fe^'*"'**'-^"*' and Fe^'*"*"'^-^'*'. The electronic structure of the transition metal sublattice can then be

described in terms of spin- and charge-density waves, but it is still

surprising not to find an accompanying periodic structural distortion. It

is possible that such a distortion is quenched by the presence of disorder

among the Sr and La atoms on the A site of the perovskite structure. The behaviour of BaLaFe20^ 91 in some ways resembles that of

Sr2LaFe308 94, but in other ways it differs (5). Mossbauer spectroscopy

indicates that a high temperature, average valence state does undergo a charge disprqportionation into oxidation states that can be approximated by

Fe^ and Fe in a 3:1 ratio, and the crystal structure is always that of a primitive cubic perovskite, indicating an absence of ordering among the

transition metal cations. However, the low temperature magnetic structure

is that of a G-type antiferromagnet, perhaps the most common type of

magnetic ordering found in simple perovskites and consistent with the

presence of a disordered arrangement of Fe^"*" and Fe^"^ in this compound,

unlike the behaviour of Sr2LaFe308 91 where an ordering of the spin

density was found.

SrMni.^Co^03.y

The existence of both hexagonal and cubic perovskites leads to a number of interesting questions, for example if SrCo03 is a cubic, metallic

ferromagnet and SrMn03 is a hexagonal antiferromagnetic insulator, what

are the properties of the solid solution SrMnj_xC0j^03_y? Our attempts to

prepare this compound in air (6) led to the discovery of an unusual

incommensurate hexagonal phase for 0.4<x<0.55. The lines in the X-ray

powder diffraction pattern could be indexed using the relationship

1 4 (l+j5)^ = (h^+k^+hk) +

d^ 3a2 c^

61
Figure 11: The magnetic structure of Sr2LaFe308.94

62
Figure 1 3: The Co K-edge EXAFS of SrMno.55Coo.45O2.

where 6 = 6(x): 6o.4~0.194, 5p.55~0.135. The value of y is -0.45 for all values of x. The origin of this incommensuration is unclear, but again

EXAFS has thrown some light on the problem. The Mn and Co EXAFS

from a sample having x = 0.45 are shown in Figures 12 and 13 respectively.

The Mn data show an oxygen shell at a radial distance of 1.88A from the

central Mn atom, a transition metal shell at a distance of 2.53A, and a Sr

shell at 3.21 A. The Co data are somewhat more complex, the oxygen shell

is again followed by a transition metal shell, but there is then an

additional Co shell at 2.87A before the Sr shell at 3.2lA. The presence of

a shell at ~2.5A in both data sets indicates that both the Co and Mn cations

occupy, at least in part, face sharing octahedra. The shell at 2.87A,

observed only in the Co data, suggests the presence of edge sharing Co

containing octahedra, and is reminiscent of our results, described above, on

SrCo02.56- It ni^y be that there is a link between the structure of this

incommensurate phase and that of 'rhombohedral' SrCo03_y.

ACKNOWLEDGEMENTS

Dr. T.C. Gibb (Leeds University) played a major role in this work,

particularly in the collection and analysis of Mossbauer data. Drs. A.T.

Steel and R. Strange played an equally important part in the EXAFS work.

The work has been funded by the SERC.

63
REFERENCES

1 . P.D. Battle, T.C. Gibb and A.T. Steel

J.Chem.Soc.Dalton Trans., 1987, 667

2. P.D. Battle, T.C. Gibb and P. Lightfoot

J. Solid State Chem., 84, 237 (1990)

3. P.D. Batde, T.C. Gibb and S. Nixon

J. Solid State Chem., 77, 14 (1988)

4. P.D. Batde, T.C. Gibb and P. Lightfoot

J. Solid State Chem., 84, 271 (1990)

5. P.D. Batde, T.C. Gibb, P. Lightfoot and M. Matsuo

J. Solid State Chem., 85, 38 (1990)

6. P.D. Batde, T.C. Gibb and R. Strange

J. Solid State Chem., 81, 217 (1989)

64
DISCUSSION

M. Takano: Did you make EXAFS measurements at low temperatures?

P. Battle : On the Co system, no, we didn't. We haven't done any EXAFS

measurements on those systems at all.

R. Newnham: You have spin glass behavior in that 3''"/5'^ phase, don't you?

P. Battle: In the Ba-La one it's what physicists call a Mattis spin glass where you have

a random arrangement of Fe^"*" and Fe in the 3:1 ratio which is governed by the

oxygen content. And, the value of the magnetic moment you get from the neutron data

is again consistent with antiferromagnetic coupling between Fe-' pairs, antiferromagnetic

coupling between Fe^"*" pairs, but ferromagnetic coupling between 3+ and 5 + .

M. Takano: Did you say it's a spin glass?

P. Battle : It's what physicists, I believe, call a Mattis spin glass. Physics Letters, 56A . 421 (1976).

M. Takano: Did you make magnetization measurements?

P. Battle: No, we haven't done that measurement yet.

A. Navrotsky : Do you need Fe^"*" to stabilize Fe^"*"? That is, can you replace Fe^^ with

something trivalent of similar size such as, for example, aluminum or gallium, and still maintain the Fe^"^?

P. Battle: There is that compound LdL2^^^^^6 which is an Fe^"*" perovskite made by

Hagenmuller's group which doesn't have any other oxidation states of iron present. That

was done in about 1982.

K. Poppelmeier: Getting back to this explanation of no atomic displacement. Is it associated with the fact that you have two different A cations,

P. Battle : I think so, yes. I think the sort of variation in the strain field caused by

having disordered arrangements of strontium and lanthanum quenches the structural

distortion that you would otherwise expect to see.

A. Sleight: Just to follow up on that. That might be a good explanation for this one but

it doesn't solve the problem of the calcium iron oxide.

P. Battle: That's the one that shows the charge disproportion. It's not clear. Has

anybody actually done the structure of CaFeOg at low temperature?

65
M. Takano: No, just simple powder X-ray diffraction.

A. Sleight : I want to go back to the strontium cobalt system. This pseudo-rhombohedral

phase you were talking about. What I was curious about and you didn't mention, is it single phase?

P. Battle : We tried to index the extra peaks as other likely impurities or second phases

and got nowhere. We did think about the possibility that there was something else in

there. I suspect it may be an incommensurate phase.

D. Smyth: Are these powder samples?

P. Battle: Yes.

D. Smyth : Are you sure that when you go through these various quenchings and slow

coolings, and so forth that you have really homogeneous oxygen stoichiometrics?

P. Battle : The way we made the Og 94 was to cool it very slowly and then leave it at 200°C for 24 hours. So, I believe we came very close to homogeneity.

D. Smyth: But there is a continuous tendency to pick up oxygen as you cool down and

at some point you'll have a kinetic problem.

P. Battle: Yes, but I think we did it in what I think was a thorough way, and I'm

satisfied. From memory I can't describe everything we did. But I know at the time I believed it was thorough.

D. Smyth: I would like to see a long equilibration at a fixed condition in the quench.

P. Battle: That's essentially what we did. We cooled it slowly from high temperatures

to 200°C to give it a chance to pick up what it wanted to. We held it at 200 and then

quenched it. You get the same result from 400 as 200. We did that just to prove that

point.

66
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

FERROELECTRIC AND FERROELASTIC DOMAIN STRUCTURES

IN PIEZOELECTRIC CERAMICS

L.A. Bursill and Peng JuLin

School of Physics

The University of Melbourne

Parkville 3052 Victoria

Australia

ABSTRACT

There exist a hierarchy of domain textures in both single crystal and

polycrystalline piezoelectrics. Considerable effort has to be made to develop imaging techinques which allow the different contributions to be distinguished; e.g. ferroelectric and/or ferroelastic domains,

discommensurations versus antiphase boundaries, extended versus

localized chemical defects and multiphase versus grain boundaries.

The present paper presents a discussion of the results of conventional

and high-resolution high-voltage electron microscopic studies of two

important ferroelectrics, barium sodium niobate (BNN) and lead

zirconium titanate (PZT). The first case is relatively well-characterized

whereas the latter reveals the complexity of a commercial piezoelectric

ceramic.

BARIUM SODIUM NIOBATE (BNN; Ba2NaNb20i5)

BNN has essentially a tetragonal-tungsten bronze type framework

superstructure with the corner-shared octahedral sites fully-occupied

by Nb (Fig. la). The room temperature phase has cell parameters a =

35.18A, b = 35.24A, c = 1.99k, space group Im2a [1]. The pentagonal and

square tunnels are fully-occupied by Ba and Na respectively. A highresolution electron micrograph (Fig. lb) shows all of the cations imaged

for the [001] projection. Since the point group of the paraelectric phase

(4/m2/m2/m) has order 16 compared with the room temperature

ferroelectric phase which has point group mm2 of order 4 then we

expect two ferroelectric variants and two ferroelastic variants [2].

Physically we may say the polarization vector lies either parallel or

antiparallel to the c axis whereas the orthorhombic a axis lies either

parallel to [110], or [110]. of the tetragonal paraelectric phase.

67
Ip^-,

'^ftfe "'^li^ JHf

»

4A

(a) Ferroelectric domains.

BNN at room temperature.

# t f » II ^¥f:^

Fig. 2a shows 180° domain walls in as grown

Note the irregular nature of the walls, which

may be ribbon-like at low magnification. The fringe pattern shows the

"black-white" contrast (a -fringes) typical of inversion boundaries in

ferroelectric materials. On approaching the thin edge of the specimen

these domains become typically lenticular in shape (Fig. 2b). In this

high-resolution dark-field image asymmetrical excitation of 001

systematic reflections gave rise to a background change, whereas

interference between subcell and supercell reflections maintained a

lattice image. Thus it can be seen that the wall thickness is virtually zero. These walls are elongated in the direction of the ferroelectric c- axis. Sub-lattice steps have been visualized clearly in many areas. The

lenticular nature of these domains, with c alternately parallel and

antiparallel to the crystal surface may indicate the considerable

influence of surface depolarization mechanisms on the domain wall

topology. Crystallographic studies [3] show that all the metal cations in

BNN are displaced along c, in the same sense, away from the nearest

mean plane of oxygen atoms, which provides the origin of the

ferroelectricity and determines the macroscopic polarization vector.

Other more complex domains, including zig-zag and tightly-curved

patterns, represent alternative surface depolarization mechanisms.

Here the domain texture is strongly dependent on the orientation of the

polarization vectors with respect to the surface and edge directions of

the wedge-shaped crystals used typically for electron microscopy. Any

electric fields due to charging of the specimen serve only to complicate

the domain texture. Thus focussing the electron beam to an intense spot

may induce rapid rearrangements of the surface domains.

68
fig 2a i^avy 1 SO" domains In 8WW

fig 2b UyuUcLitcVL ISO {^2AAo 2Ze.cXAlc. domains In BUM

(b) Ferroelastic/ferroelectric domain walls. The room temperature

phase of BNN is both ferroelectric and ferroelastic, but the spontaneous

polarization and strain appear to remain uncoupled, with the

configuration of the ferroelastic domain not strongly coupled to the

180° ferroelectric domain walls. Two sets of ferroelectric-ferroelastic

domains are observed with domain walls approximately parallel to [100]^ or [010]^ directions, referred to the tetragonal subcell axes (see

Fig. 3). These 90°-ferroelastic/ferroelectric domains present distinct

contrast for appropriate bright- or dark-field imaging conditions. They

also show essentially planar walls, quite unlike the 180° domain walls,

i.e. the surface energy would appear to be much higher, due to the

piezoelectric induced strain across the walls.

(c) Commensurate superlattice domain walls (or discommensurations).

Incommensurate superlattice phases of BNN were studied by X-ray

difffraction [4]. The commensurate/incommensurate transition near 561K gives a phase having a double c axis and an incommensurate

modulation along [110]^ tetragonal axis, yielding orthorhombic

69
¥lg.3 Vu^commQ^Yi^uJtcuUovi^, i^hn-OoZoMtic domains and mvy

{^^AAooI-^ct/ilc domains In BMW

70
symmetry for the average structure. It has been suggested

theoretically [5] that the incommensurate phase is composed of

commensurate domains separated by out-of-phase domains or discommensurations. Bear in mind that the results of the present

section all refer to the commensurate phase, which is stable at room

temperature. Dark-field diffraction contrast images were obtained

using the commensurate superlattice reflections (inset in Fig. 3). Very

faint streaking parallel to [110] was also included (not visible). A

special tilt condition was required, to enhance those weak reflections to produce observable contrast. A distinct array of ribbonlike domain

contrast was then observed, accompanied usually by 180°-ferroelectric

domain walls as well as 90°-ferroelastic domains. These new features are relatively very rough (stepped) and irregular compared to the smooth

wavy 180° walls, or the essentially planar 90° walls (cf. Figs. 2). Fig. 3

refers to a crystal orientation tilted a few degrees away from the

tetragonal c-axis. A second imaging condition was obtained close to the

orthogonal [110]^ projection, which gave discommension images which were elongated essentially parallel to the c-axis. These were again

easily distinguished from the 180° and 90° domains, which could be

extinguished for this projection by appropriate choice of dark-field

imaging conditions. Thus the discommensuration domain walls are elongated along c but have an irregular, chaotic cross-section, perhaps

best described as irregular prismatic in shape. Close inspection of the

discommensuration morphology reveals a predominance of quadruple

modes when they intersect (see Fig. 4) confirming that at least four

translation variants exist in the room temperature phase.

Referring to our Fig. 1 , note the sequence of rotated octahedra C A A C

along both [100] and [010]. Thus there are four possible sequences

generated by a discommensuration which interupts the normal

sequence (Fig. 5) corresponding to four phase shifts which are multiples of nl2. The discommensuration density increases dramatically on approaching the commensurate to incommensurate phase transition

[6,7].

V-LQ 4 VouA-lotd dAJ^CLommZYiiiUJwution. VIq 5 ?hcu>(L change cut ^ouJimdd In BWM node in BNN

71
(d) Chemical defects. The isostructural phase strontium barium

niobate (SBN ; Ba^Sr j _^Nb205 ; 0.25 < x < 0.75) exhibits a further domain

texture, due to short-range ordering associated with the mixed

occupancy of Ba and Sr along the [001] tunnels [8].

(e) CQnglusions. Use of highly perfect single crystals of BNN has

enabled the hierarchy of domain textures to be sorted out using a combination of high-resolution and conventional bright- and darkfield electron optical techniques.

LEAD ZIRCONIUM TITANATE (PZT; Pb(Zrj^Ti^)03)

The phase diagram of the PbZrO^-PbTiO^ system shows a complete solid

solution at high temperature [9] with Zr and Ti distributed randomly

over the octahedral sites of the cubic perovskite structure. On cooling

there is a displacive phase transformation into a distorted perovskite

structure. Titanium-rich compositions favour a tetragonal modification

with significant elongation along [001] and a large spontaneous

polarization in the same direction. There are six equivalent polar axes

in the tetragonal state, corresponding to the <100> directions of the

cubic paraelectric state. A rhombohedral ferroelectric state is favoured

for zirconium-rich compositions. Here the distortion and polarization

are along the [111] directions, giving rise to eight possible domain

states, corresponding to the eight <111> directions of the cubic

paraelectric phase. The compositions which pole best are near the

morphotropic boundary between the rhombohedral and tetragonal

ferroelectric phases. Thus, for x 0.5 there may be up to 14 poling

directions over a wide temperature range. We note that the

piezoelectric coefficients are largest near the R-T phase boundary.

Domain walls are an important source of dielectric loss for temperature

below T^ [9]. Under applied fields, domain wall motion takes place,

dissipating energy. A number of different types of walls are expected

in PZT. Tetragonal PZT has 180° walls, and both charged and uncharged

90° walls. Charged walls are important only in conducting ceramics,

where currents can flow, neutralizing the charge. The 180° walls are generally more mobile than 90° wails because of the mechanical strain

associated with the latter. The magnitudes of the piezoelectric

coefficients depend markedly on dopants and defect structure because

of their influence on domain wall motion, due to alignment of defect

dipoles against spontaneous polarization below the Curie temperature

(~350°C) [9]. Oxygen vacancy diffusion is believed to be the

predominant relaxation mechanism contributing to dielectric losses [9].

72
The aim of the present work was to observe submicron texture for PZT

ceramics of commercial interest. Bright- and dark-field TEM images were used to characterize the ferroelectric and ferroelastic twin

domains in both tetragonal and rhombohedral phases of PZT. Attempts

have also been made to apply atomic resolution electron microscopy to investigate domain wall widths (charged walls may be diffuse, whereas

uncharged walls are sharp [10,11], as well as interactions between twin

variants and dislocations. The grain texture is also accessible to highvoltage instruments (say >.300 kV). Results are reported here for a pure

tetragonal PZT, i.e. Pb(ZrQ 53'TiQ 47) and a rhombohedral doped PZT + 6

mol.% NiNbO-^. Thin specimens were prepared by mechanical grinding

and dimpling, followed by ion-beam thinning using argon. Thin edges were carefully aligned with <001>, <011> and <111> pseudocubic vectors

parallel to the electron beam. A JEOL-4000EX instrument was used,

operated at 400 keV.

(a) Pb(Zr

Q gj . TIq

^y'lO ^ . Low magnification shows grain size about 7

(im. Complex intersecting fringe contrasts, indicative of multiple

polysynthetic twin lamellae give rise to textural detail within each

grain on the scale ^SO nm. Fig. 6 shows a whole grain, containing lath- like domains parallel to [110], as well as fringe contrasts due to extended

defects inclined to the projection axis. The overwhelming complexity of

the domain texture, and the fine scale have so far prevented a systematic identification of the many possible combinations of

ferroelectric and ferroelastic twins in this material.

Fi.g 6 Edge,- on [tdit] and lvidU,n2,d [^ght] {^QAAodizdPvic, domains

in toJi^gonat PIT

73
High-resolution images showed areas of perfect perovskite-type

structure with changes of image detail and symmetry occuring on the

scale of -10 nm. Such changes are consistent with the existence of

coherent domain walls. This specimen closely ressembled BaTiO^ and

KNbO^ in its behaviour in the electron beam. Thus the specimens often

showed charging effects.

(b) PbfZrQ

5 3 ^-^ IO j + 6 mol.% NiNbO ^. This rhombohedral

specimen was relatively very stable in the electron microscope,

showing no charging effects. Fig. 7 shows a typical grain again

showing complex domain textures. Some grain boundaries showed

precipitates. High-resolution images again showed local changes of

periodicity (doubling) and symmetry implying a fine-scale coherent

microdomain texture, with fluctuations on the scale of 5-10 nm (Fig. 8). This texture was remarkably similar to results obtained previously for

silver ion-exchanged lithium tantalate [12,13], proton- and Fe-doped

lithium niobates [14,15]. The papers [12-14] should be consulted for a discussion of the coherency and structural relationships between

various pseudocubic and rhombohedral perovskite phases.

Vlg 7 T^AKonla^tlc and {^^KA0QJi2.cX/u.c domain^ In nkombok^dAol PIT

74
38A

¥lg S HREM Image. ^hoLClngchange^ In IcuUxciQ, Imag^ on cA06^i.ng

(^QAn.odi<utlcL domains In PIT

(c) Conclusion. The present results demonstrate that knowledge of the

interaction between stoichiometry, processing routes, microstructure

and piezoelectric characteristics is rather primitive, due largely to the

complexity of PZT. It is clear however that there is a rich variety of

microstructural information accessible in PZT, using high-voltage

HREM and TEM techniques

ACKNOWLEDGEMENTS

This work was supported financially by the Australian Research

Council, the Victorian State Government and The University of

Melbourne, through the National Advanced Materials Analytical Center.

75
We are grateful for the collaboration of Dr J.R. Sellar, ICI Australian

Research Laboratories.

REFERENCES

[I] Peng JuLin and L.A. Bursill, Acta Crystallogr. B 47, 504 (1987).

[2] V. Janovec and V. Dvorak, Ferroelectrics 66., 169 (1986).

[3] P.B. Jamieson, S.C. Abrahams and J.L. Bernstein, J. Phys. Chem. Sol.

49, 4352 (1969).

[4] J. Schneck, J.C. Toledano, C. Joffrin, T. Aubree, B. Joukoff and A.

Gabelotard, Phys. Rev. B 25., 1766 (1982).

[5] W.L. McMillan, Phys. Rev. B 14, 1496 (1976).

[6] G. van Tendelloo, J. van Landuyt and S. Amelinckxz, (J.F. ScoH" and

N.A. Clark (Eds.) "Incommensurate Crystals, Liquid Crystals and

Quasicrystals", NATO A.S.L Series, Plenum Press, NY, (1987).

[7] Feng Duan, Cryst. Latt. Defects and Amorp. Mat. 12, 229 (1985).

[8] L.A. Bursill and Peng JuLin, Acta Crystallogr. B 43, 49 (1987).

[9] R. E. Newnham, Rep. Prog. Phys. 52, 123 (1989).

[10] L.A. Bursill, Peng JuLin and Feng Duan. Phil. Mag. 48, 953 (1983).

[II] Peng JuLin and L.A. Bursill, Phil. Mag. 48, 251 (1983).

[12] L.A. Bursill and Peng JuLin, Ferroelectrics 76, 175 (1987).

[13] Peng JuLin and L.A. Bursill, Ferroelectric 77, 69 (1988).

[14] Yan Yong, Feng Duan, Peng JuLin and L.A. Bursill, Ferroelectrics

77. 91 (1988).

[15] L.A. Bursill and Peng JuLin, Ferroelectrics 71, 81 (1988).

76
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

CATION SUBSTITUTION IN BARIUM TITANATE FOR

DIELECTRIC CERAMIC APPLICATIONS.

F.Ba.tlloa), A.Beaugerc), BJannotb), J.CJulesa), M.Laurentc),

M.Maglione^), J.C.Niepce^).

a) Laboratoire Reactivite des Solides - CNRS URA 23

^)Laboratoire de Physique du solide - CNRS URA 785

Universiti de Bourgogne - BP 138 - 21004 DIJON Cedex, France.

^^Laboratoire Recherche et Developpement - Thomson LCC

Avenue du Col. Prat - 21850 Saint Apollinaire, France.

A small amount of foreign element is able to modify drastically the

properties of a material but the understanding of these modifications is not always well established.

When a partial substitution on a cation site is carried out in a crystalline

material the substituting cation may differ from the substituted cation

by the mass, the volume, the electric charge, the oxydo reduction

potential,...

At the present time it is not possible to give a full explanation of the

modifications of the properties of the material resulting of the various

substitutions with reference to the above features only. Nevertheless we

would like to show, on barium titanate as an example, how some of these

features can be taken into account to influence the properties of a material.

It will be shown and explained how some of the dielectric properties of

BaTi03 are modified with various substitutions. Particularly, with

reference to some of the above features, the manufacture of BaTi03

based material able to be cosintered with a non noble metal will be

described.

INTRODUCTION

Barium titanate is the main dielectric material component of

class II ceramic capacitors. Due to particular characteristics

such as its crystalline structure, insulating properties and

phase transitions, BaTiOs exhibits a very large relative

dielectric constant £r at the ferroelectric-paraelectric transition.

For these reasons, BaTiOs, used as main dielectric material

77
component of class II ceramic capacitors, allows the highest

volumic capacitance for this kind of capacitors among all the

capacitor species.

The manufacture of ceramic capacitors under certain pressure

and temperature conditions could lead to an important drop in

BaTiOs dielectric properties. For these reasons it is always

necessary to modify the chemical composition of the material

in order to anneal the effects of such processing.

To fully understand the chemical modifications of BaTiOa prior

to its elaboration as a ceramic capacitor a review of the point

defect chemistry at high temperature of the material should be

made.

INTRINSIC POINT DEFECT IN BaTiOa

Figure 1 shows the point defect chemistry outlines of pure

BaTiOs evidenced by electrical conductivity measurements

upon a large range of oxygen pressure [1][2]. Three different

regions can be defined.

1

/ -20 -15 -10 -5 0 log POj (atm)

Figure 1 : • —• —• — Equilibrium electrical conductivity versus oxygen

pressure for undoped BaTi03 [1].

• —• —• — Equilibrium line between the stability fields of

Nickel and Nickel oxide in a P02 vs. T diagram [2].

78
- Region I is located at very low oxygen pressure and high

temperature. The point defect chemistry of BaTiOs in this

region is governed by the departure of oxygen from the

structure and the creation of oxygen vacancies associated with

the release of 2 electrons. These 2 electrons eventually remain

weakly bonded to the titanium sites and lead to a partial

reduction of Ti(IV) cations. This behavior can be described by

the reaction : Oo + 2TiTi

^

I/2O2 + Vo" + 2TiTi' (1)

- The point defect chemistry in region II and III have been

subject to controversies among the different authors. However

the most widely accepted view is that of Smyth et al.[l] who

suggest that oxygen vacancies Vq * resulting from the presence

of extrinsic point defects and in particular acceptors (i.e. cations

with lower valency than normal cations) are still preponderant

in these regions.

ELABORATION OF MULTILAYER CERAMIC CAPACITORS

One of the most used type of capacitor is the multilayer

ceramic capacitor (MLC). The MLC consists in the alternated

superposition of dielectric and metallic electrode layers as shown figure 2.

a) Disk b| Multilayer

Fi gure 2 : Ceramic Capacitors

79
This particular processing allows a larger volumic capacitance

than classical capacitors. In order to lower the MLC cost price,

non noble metal such as nickel is likely to be used as internal

electrodes.

An important stage during the manufacture of the MLC is the

cosintering of the dielectric material and the Ni electrodes at

high temperatures.

Figure 1 shows that in order to avoid an oxidation of the nickel

and the loss of its conductive properties, the MLC must be

sintered under reducing conditions (Po2 < lO'^^atm). These

pressure and temperature conditions coincide with the

reducing region of BaTiOs (region I). The sintering under low

oxygen pressure results in the departure of oxygen from the

material and the reduction of Ti(+IV) cations into Ti(+III)

cations. The ceramic is no longer an ideal dielectric but

presents the characteristics of a n type semi conductor with

large dielectric losses at room temperature (figure 3) [3].

_40 0 100 Temperature ( C)

Figure 3 : Influence of the oxygen partial pressure during the

sintering of a Z5U material with general formula

Bai-x Cax Tii.y Zry O3 [3].

To avoid the reduction of Ti(+IV) cations, one solution is to trap

the electrons on a more reducible cation than Ti(+IV). Figure 4

represents the free enthalpy variation of the reaction :

80
M + n/402 -> MOn/2 (2) for different elements [4]. This figure

shows that manganese is less oxidizable than titanium and

consequently oxidized Mn is more easily reduced than oxidized

titanium. This confirms that Mn(+IV) and Mn(+III) cations are able to trap the electrons produced during the sintering of the

MLC. Moreover, due to the fact that Mn cations are statistically

far from one another ([Mn]^0.1%), the trapped electrons are unable to jump from site to site.

The large decrease of the dielectric losses with the increase of

Mn concentration in BaTiOs shown figure 5 [2], confirms the

role of Mn as an electron trapper.

Figure 4 : Oxidation state diagram Figure 5 : Dielectric losses at Room

of the metal of the phase Temperature as a function of

transition series [4] (etablished Mn(+II) concentration in a BaTi03

with reference to the free based Z5U material [2]. enthalpy ^G^2^0 of formation of

the oxides [7], [8]).

Another unwanted effect of the sintering at low Po2 is the

appearance of oxygen vacancies (reaction (1)). In addition the

81
introduction of acceptors in the lattice like Mn(+III) in place of

Ti(+IV) is also source of oxygen vacancies through the reaction:

BaO + (l-y)Ti02 + y/2 Mn203 ^

(BaBa) (TiTi)l-y (MnTi)y (Oo)3-y/2 (Vo")y/2 (3)

Because of their mobility when T increases, the presence of

oxygen vacancies inside the lattice induces a very poor stability

of the dielectric material during accelerated life tests as shown

figure 6 [3].

To prevent the appearance of oxygen vacancies a substitution

of Ti(+IV) and/or Ba(+II) cations by donors i.e., cations with

higher electric charge, can be made. For example, the

introduction of Ce(+III) cations in place of Ba(+II) lead to the

appearance of CcBa which compensates for 1/2 Vq ".

jk log p

(n.m)

9-

Figure 6 : Resistivity of BaTi03 based Z5U material as a function of

various doping during life test (1.7V/)im and 280°C) [3].

The double substitution, in a convenient proportion, by both

donors and acceptors like Ce(+III) in Ba(+II) sites and Mn(+III)

in Ti(+IV) sites can be described by the following reaction:

(l-x)BaO + x/2Ce203 + (l-x)Ti02 + x/2Mn203 ^

(BaBa)l-x (CeBa*)x (TiTi)l-x (MnTi')x (Oo)3 (4)

82
The resulting material no longer contains vacancies. However

this substitution must be made in such a way that the oxygen

vacancies generated by the departure of oxygen during the

sintering have to be compensated prior to the sintering.

Figure 6 shows the effect of a triple substitution on the

accelerated life test evolution of a BaTiOs ceramic capacitor in

time. Mn has been added as Mn(+III) and Mn(+IV) on Ti(+IV)

sites to trap the electrons generated by the oxygen departure.

La(+III) and W(+VI) are added in substitution for Ba(+II) and

TI(+IV) in order to compensate the oxygen vacancies generated

by both the acceptor substitution and the oxygen departure

[3],[5],[6].

The resulting composition can be written:

(BaBa)l-x (LaBa')x (TiTi)l-y-z (MnTi')y (WTi")z (Oo)3

In conclusion, despite the chemical complexity of a material

resulting from different substitutions, we have shown that it is

possible to modify in a convenient way some of the dielectric

properties of a material by considering parameters such as: relative electric charges and oxydo reduction potential of the

substituting cation with reference to the substituted cation.

However, we would not try to explain completely the behavior

of a substituted material by considering only these parameters.

Nevertheless they can be a good guide.

References:

[1] D.M. SMYTH. Prog.Solid.State Chem., 15, (145-171) 1984.

[2] F.BATLLO. Thesis DUON (France), 1987.

[3] M.LAURENT. Thesis CNAM DUON (France), 1989.

[4] A.MORELL and J.C.NIEPCE. J.Mat.Education, (in press.), 1991.

[5] Patent of "Inst. Nat. Prop. Indust. Fran9ais", n^ 8817081,

(23/12/88).

[6] Patent of "Inst. Nat. Prop. Indust. Fran^ais", n° 8807404,

(03/06/88).

[7] O.KUBACHEVSKI and E.EVANS. "La Thermochimie en

Metallurgie". Gauthiers Villars Edit. Paris. 1964.

[8] L.BREWER. Chem. Rev., 52, 1,1953.

83

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

LIQUID PHASE Sintering of LIF-Fluxed BaTiOs

Sea Fue Wang, Wayne Huebner, and Clive Randall

Materials Research Laboratory

Penn State University

University Park, PA 16802

INTRODUCTION

Reduction of the firing temperature of BaTiOs-based multilayer capacitors

(MLC's) is typically achieved using fluxing agents such as CuO, ZnO, CdO, LiF, or Bi-compoundsl"6. Most of these fluxes enhance densification due to liquid

phase sintering. Liquid phase formation in these systems is either due to melting of

the flux, or by the formation of a eutectic liquid between the flux and BaTiOs.

Although the flux reduces the sintering temperature, it may also introduce

undesirable effects such as lowering the dielectric constant and mechanical strength,

and increasing the aging rate. These effects are most noticeable in systems in which

the flux remains as a continuous second phase in the grain boundary, or enters into

solid solution with the BaTi03 and hence alters the defect chemistry.

LiF is an effective sintering aid for both MgO and many perovskites^-^.

Small additions of LiF (0.5-3wt%) may yield > 98% dense BaTiOs at temperatures

as low as 850°C3-5. Changes in stoichiometry dramatically impact this behavior:

Ba-excess compositions exhibit the highest densities. Desgardin et al.^'^ proposed

that an intermediate composition of the fired ceramic is an oxyfluoride, BaTii.

xLix03-3yF3y, but this was difficult to prove due to the fact that Li and F are not

readily detectable by simple spectroscopic techniques.

The dielectric behavior of flux-sintered BaTi03 cannot be simply understood

through the application of standard mixing rules. Although these rules effectively

describe the dielectric behavior for multiple component systems, for flux-sintered

BaTi03 strong variations in the chemical homogeneity (and its impact on the phase

transition temperatures) and the defect chemistry demand a more complex analysis.

Few systematic studies on the microstructural evolution of flux-sintered BaTi03

and a direct correlation to the resultant dielectric properties have been performed.

The purpose of our investigations is to clarify the interaction of BaTi03 with

Li-compounds during Jiquid phase sintering, the chemical inhomogeneity of the

resulting microstructure, and the microstmcture-property relationships.

85
EXPERIMENTAL PROCEDURE

The effects of two flux systems are reported here: LiF+BaCOs, and BaLiF3.

The LiF and BaCOs were obtained from Aldrich. The BaLiF3 was prepared by

mixing stoichiometric amounts of BaF2 and LiF for 24 hours with a shaker mill.

Powders were milled in methyl alcohol using polyethylene jars and Zr02 media.

After drying, these powders were transferred to a Pt crucible, reacted at 850°C, and

quenched. The powder was then dry-milled and sieved through a 320 mesh screen. XRD analysis confirmed BaLiFs formation.

TAM-HPB BaTiOs was used in this study. Fluxes were mixed/milled with

the BaTiOs as described above. Milled powders had an average particle of 0.7 |j,m as determined by light scattering. To prepare the powders for pressing, they were mixed with DuPont 5200 binder using acetone as a solvent, dried, pulverized using

a motor and pestle, and then sieved through a 120 mesh screen. Disks were

prepared by uniaxial-pressing at 15000 psi. Green densities were = 64%

theoretical. These disks were calcined for 2 h at 500°C, followed by sintering in a closed crucible at temperatures ranging from 800 to 1150°C for two hours. A

6°C/min heating rate was used throughout.

Sintered densities were measured using a displacement technique with xylene

as the liquid medium. Fired samples were mounted in an epoxy, ground, polished

and etched using a 10 wt% HCl, 90 wt% H2O solution with several drops of HF.

Microstructural observations were carried out on etched surfaces using an SEM,

and on ion-milled samples with a TEM.

After sputtering on Pt electrodes, the dielectric properties of sintered samples were measured as a function of temperature and frequency using an HP 4274 RLC

bridge. Measurements were taken during cooling from 200 to -50°C, over a frequency range from 100 Hz to 100 kHz.

RESULTS AND DISCUSSION

Figure 1 shows the densification curves for 2wt% LiF+2m% BaC03, and 3

and 7w% BaLiF3. For temperatures >850°C the densities for 2wt% LiF-i-2m%

BaCOs and 7w% BaLiFs samples are all =95%. The 3w% BaLiFs flux did not

densify as well, which is probably due to an insufficient amount of liquid phase.

Since the melting temperature of LiF is 846°C, a eutectic liquid must have formed in

order to explain the onset of densification for T<800°C. Results from DTA studies on mixtures of BaTiOs and 20 wt% LiF with or without BaC03, and BaTi03 and

20 wt% BaLiF3 are shown in Figure 2. Endothermic peaks corresponding to eutectic liquid formation appear at 700°C for BaTi03 and LiF, and 610°C with the

addition of BaC03. This correlates well with studies which have shown LiF

additions to Ba-excess BaTi03 (adjusted by BaC03) lower the sintering

temperature^'5. The DTA results also show that mixture of BaTi03 and BaLiF3 did

not exhibit any liquid formation other than the incongruent melting of BaLiF3 at ~840°C. This result is consistent with the sintering behavior which shows

86
100

750 850 950 1050 1150

Temperature (°C)

Figure 1 : Densification behavior of BaTi03 sintered with LiF and BaLiF3

o

a2

UJ

V a/

U 10-00-

LD\

_1

B < r

U

c_

A - BaTi03 + LiF 11

B - BaTiOs + BaLiFa ' f

aoo- C - BaTi03 + LiF + BaCOs

330.00 4^00 510.00 oiaoo eaaoo 780.00 skoo sdaoo TEMPERATURE (C) DTA

Figure 2: DTA results for mixtures of LiF and BaLiFs with BaTi03.

87
significant densification occurs for T >850°C; i.e. the onset of densification is coincident with liquid phase formation. Some authors^ have proposed that the

mechanism for densification in the LiF system is due to intermediate formation of

BaLiF3. However, our results show BaLiF3 is only effective at higher

temperatures.

XRD analysis on identical mixtures heated and quenched at temperature from

600 to 900°C indicate : a) no BaCOs or LiF peaks are observed at temperatures

above the eutectic temperature, b) for T>800°C there are no detectable differences in

phase formation for mixtures of BaTiOs and LiF with and without BaCOs, and c)

only BaTiOs peaks exist when the BaLiFs mixture is heated at T>900°C.

Based on these DTA and XRD results, liquid formation at the eutectic

temperature of BaTi03-LiF-BaC03 at 610°C is responsible for the lower

densification temperatures observed for Ba-excess compositions. Since only

BaTiOa peaks are indicated in XRD pattern of BaTiOs and BaLiF3 mixture, this

confirms that the formation of a perovskite solid solution BaTii.xLix03.3yF3y from

BaTiOs and BaLiFs.

Figures 3a and b contain SEM micrographs of 2 wt% LiF-BaTi03

(BayTi=1.02) sintered at 900°C and 1 100°C for 2 hours. For the sample sintered at 900°C, the grain size is = 1-2 |im. For samples sintered at 1 100°C the micrograph

shows faceted grains of an average diameter >50 |im, and the presence of many

<111> twins. This microstructure is similar to that obtained when Ti-excess

BaTi03 is sintered at T> 1312°C. Figure 4 contains a representative TEM

micrograph of a sample sintered at 900°C. A grain core-shell microstructure is observed, with both 90° and 180° domain walls inside the core region, and line and

dislocation loops located at the core-shell interface. Heating and cooling the sample

in-situ and observing the disappearance and appearance of the domains confirmed

the core was pure BaTiOs. Importantly, the shell bears an epitaxial relationship to the core, and the size of the core is roughly equal to the original starting powder

size prior to the liquid phase sintering. A residual glass phase was also observed at grain boundaries and triple points. These results indicate that the shell is the

perovskite solid solution, BaTii.xLix03.3yF3y, which formed due to a solutionreprecipitation process. The interface is due to the lattice mismatch associated with

the discontinuous Li-concentration.

The permittivity response of a LiF-BaC03-BaTi03 sample sintered at 900° for

2 hours is shown in Figure 5, with the corresponding Curie-Weiss plot in Figure 6. A broad peak in the dielectric constant is observed at ~10°C, and a slight peak at = 120°C. The dielectric loss is high and dispersive. These results correlate well

with the hypothesis that the shell is a BaTii-xLix03.3yF3y solid solution. The peak

at ~10°C is associated with the decrease in the cubic-tetragonal transition, while the

broadness is associated with spatial variation in x and y. The linearity of the CurieWeiss plot for T>125°C is due to the pure BaTi03 core. The high, dispersive

dielectric loss mq^t likely indicates the shell region exhibits a high conductivity

88
0.1 im

Figure 4: TEM micrograph of a 2wt% LiF-BaTiOs (BayTi=1.02) sample

sintered at 900°C.

89
c

-pC

o

u

^000

Q 2000 -

U

Q

•50 0 50 100

Temperature (C)

-i . 4

. 3

. 2

. 1

0

150

<UM

O

Figure 5: Dielectric behavior of a 2wt% LiF-BaTiOs (Ba/ri=1.02) sample

sintered at 900° for 2 h.

1 /K

0007

r

.0006 -

.0005 -

. 0004

^ 0002 L-A J I i I I I L ''HI

-50 0 50 100 150

Temperature (C)

Figure 6: Curie-Weiss plot of the dielectric behavior exhibited in Figure 5.

90
associated with either ionic conductivity of Li or F ions, or due to an enhanced

electronic conductivity associated with an altered defect chemistry.

Summary

1. The onset of densification in the flux systems studied can be direcdy

correlated with the temperatures at which a eutectic liquid forms. The lower

sintering temperatures of Ba-excess compositions (due to BaC03) is clearly

due to the eutectic liquid which forms at 610°C, compared to 700°C when

BaC03 is not present.

2. TEM micrographs reveal BaTiOs sintered with 2 wt% LiF (Ba/Ti=1.02) at 900°C for 2 hours results in a grain core(tetragonal)-shell(cubic) structure. The shell region is most likely the perovskite solid solution BaTii.xLix03.

3yF3y, which is formed due to a solution-reprecipitation process. The

temperature dependence and dispersive behavior of the dielectric properties

confirm this hypothesis.

References

1 . I.Bum, "Hux-Sintered BaTi03 Dielectrics," /. Matr. Sci., 17[5], 1398-1408

(1982).

2. M. W. Benecke, N. E. Olson, and J. A. Pask, "Effect of LiF on the HotPressing of MgO," /. Am. Ceram. Soc, 50[7], 365-368 (1967).

3. A. Potten, J. Ravez, and J. P. Bonnet, "STEM Study of the Sintering

Mechanism of BaTi03 with LiF," High Tech Ceram. . edited by P. Vincenzini, 1525-1532 (1987).

4. J.M. Haussonne, G. Desgardin, P.H. Bajolet, and B. Raveau, "Barium

Titanate Perovskite Sintered with Lithium Fluoride," /. Am. Ceram. Soc,

66[11], 801-807 (1983).

5. D. A. Tolino and J. B. Blum, "Effect of Ba:Ti Ratio on Densification of LiFFluxed BaTi03," /. Am. Ceram. Soc, 68[111, C-292-C-294 (1985).

6. G. Desgardin, I. Mey, and B. Raveau, "BaLiF3-A New Sintering Agent for

BaTi03-Based Capacitors," Am. Ceram. Soc. Bull., 64[4] 563-70 (1985).

91

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

A REAL SPACE ANALYSIS OF SHORT RANGE ORDER IN

FERROELECTRIC Pb (Mg^ /3Nb2 /3 ) O3

H. D. Rosenfeld and T. Egami

(Department of Materials Science and. Laboratory for Research on the Structure of Matter, University of Pennsylvania, Philadelphia,

PA 19104, USA)

A. Bhalla

(Materials Research Laboratory, Pennsylvania State University,

University Park, PA 16802, USA)

ABSTRACT

The local atomic structure of Pb (Mg;;^/3Nb2/3 ) O3 (PMN) was

investigated using pulsed neutron powder diffraction.

Changes in nearest neighbor distances were studied by

pair-distribution function, or real space, analysis. We

find a decrease in the number of intermediate Mg/Nb-0

and Pb-0 pair distances below the Curie temperature due

to interatomic correlations. These correlated

displacements may be related to the onset of bulk

f erroelect ricity . We find evidence supporting the

existence of ferroelectric cells above T^ in PMN.

INTRODUCTION

Short range order in the relaxor ferroelectric lead

magnesium niobate Pb (Mg2^/3Nb2/3 ) O3 (PMN) has been a subject of considerable interest, since the

ferroelectric transition may be attributed to local

deviations of the atomic structure from that of an

ideal perovskite [1]. Bonneau et al . [2] have suggested

a disorder model in which Pb and Nb atoms are displaced

along the cubic [100], [110] and [111] directions, and

Smolenskii and Agranovskaya [1,3] have discussed the

distribution of Mg and Nb ions as a means of explaining

the diffuse nature of the ferroelectric phase

transition. Recently Husson et al . [4] have reported the

results of a Raman spectroscopy study of short range

93
order in PMN suggesting the existence of ferroelectric

cells above the -7°C Curie temperature.

We have studied short range order in PMN by Rietveld and

pair-density function (PDF) , or real space, analysis of

neutron powder diffraction data. The results of the

Rietveld refinement have been reported elsewhere [5].

PDF or real space analysis provides extensive insight

into the local atomic structure of a material, and

provides a simple and direct means of connecting trial

structures of materials not possessing long range order

to diffraction experiments. The pair-density function

p(r) is the Fourier transform of the structure factor

S (q) . Periodicity of the structure need not be assumed

in calculating the PDF and non-periodic structures may

be treated exactly. The PDF is sensitive to details of

the local structure, whereas Rietveld analysis is best

suited to studying long range order in periodic

structures. Neutron diffraction was chosen owing to the

relative uniformity of neutron scattering lengths which

makes it possible to refine the positions of Mg, Nb and

O atoms in the presence of Pb

.

EXPERIMENTAL METHODS

Neutron diffraction data were collected using the

Time-of-Flight method on the Special Environment Powder

Dif f ractometer (SEPD) at the Intense Pulsed Neutron

Source (IPNS), Argonne National Laboratory. A burst of

highly energetic neutrons is produced at time tg . A

liquid methane moderator placed close to the target

produces a neutron wavelength spectrum with a large

Maxwellian thermal component in addition to an

epithermal portion, producing a continuous spectrum of

neutrons with wave vectors ranging from 0 . 3A"^ to 30A~^.

Assuming elastic scattering, the magnitude of the wave

vector k may be determined by measuring the time

required. At, for the neutron to traverse the flight

path from target to detector, D, via

k=27CmD/hAt . (1)

The sample is located 14m from the target and is

surrounded at a distance of 2m by detector banks at

angles of 29 equal to 15°, 30°, 60°, 90° and 150°. The

magnitude of the scattering vector q is related to k by

94
q=2k sin (6) (2)

S(q) was obtained over q equal to 0.5-50A~i. Obtaining

the structure factor S (q) to such large values of q

minimizes termination errors when Fourier transformed to

produce the pair-density function. The lack of energy

resolution in this experiment implies an integration

over all neutron energies during data collection. In

interpreting the data, this integration was assumed

taken at constant scattering vector q. A correction for

this approximation, the Placzek correction [6], was

applied to the data. In addition corrections were made

for absorption, background and incident spectrum (see

for example [6]) . S(q) is related to the corrected

intensity by

S(q) = + «b»--«b^» .

10.37g of powdered PMN was placed in a vanadium can and

mounted in a Displex in SEPD. Diffraction data were

collected at three temperatures, lOK, 200K and 295K.

Detector counts were typically 32000 per bank.

Scattering from the empty vanadium can and empty Displex

were measured for absorption and background corrections.

Scattering from a vanadium standard was also measured to

determine the incident spectrum.

PDF ANALYSIS

The pair-density function p(r) is related to the

structure factor S (q) by the Fourier transform,

P^^^'Pq = (^:)3jdq[S(q)-l]e''^"' (4)

where is the average atomic number density and q is

the scattering vector [7]. It should be noted that S (q)

is the square of F(q), the structure factor

conventionally used in crystallography. If S(q) is

isotropic as in an amorphous material, or as the result

of powder averaging, the integral in (4) reduces to

P(r)-Po= Jdq[S(q)-i]q3in(qr) . (5)

95
The physical interpretation of p(r) may be understood by

considering the single and two particle density

functions . In general the two particle density function

is not equal to the product of the two single particle

density functions, due to interatomic correlations,

rather

n<2) ( r^^ r2) = n'D ( r^) g(r) (6)

r= r^- r2

where

g(r) = p(r)/p^ (7)

is the pair-distribution function. As r approaches

infinity, interatomic correlation becomes negligible and

g{r) tends to unity or p(r) tends to p^

.

The PDF's determined from neutron powder diffraction at

lOK, 200K and 295K are shown in figure 1. In the ideal

perovskite we would expect peaks at 2.02A corresponding

tp Mg(Nb)-0 pairs, 2.8 6A corresponding to Pb-0 and 0-0

pairs, 3.51A due to Mg/Nb-Pb pairs and 4.05A the unit

cell size. Rietveld analysis [5] suggests displacements

of Pb, Mg and Nb along the cubic [111] and of oxygen

along [112] . These displacements account for the

96
additional peaks near 2.41A and 3.30A observed most

clearly in the lOK data. Peak positions, integrated

intensities and contributing pairs at the temperatures

studied are given in table 1. The aforementioned

displacements, when not constrained by the requirement

of translat ional invariance, do not produce single pair

separations at the peak centers, but rather produce a range of distances about the peak positions. This

accounts for the broad nature of the peaks, a phenomenon approximated in Rietveld refinement by

anomalously large thermal factors. That these broad

peaks are composed of multiple pair distances may be

seen in the Mg(Nb)-0 peak. At lOK the maximum is at

2.06A, but the peak shape suggests the presence of a second peak near 1.9A. At 200K the maximum has shifted

to 1 . 95A, however substantial intensity is observed near

2.06A. In the 295K data only a single broad peak may be

seen at 2.02A, between the lOK and 200K peak positions.

Lead-oxygen distances are distributed over three ranges.

The first range centered near 2.45A, the second covers

2.8A-3.0A and the third centered near 3.3A. At 200K the

number of short Pb-0 distances decreases in favor of

peak position integrated intensity contributing pairs

295K 200K lOK 2 95K 200K lOK

2 . 02 1. 95 2.06 4.282 4.279 4.281 Mg/Nb-0

2.42 2.451 2 . 47 1. 825 1.742 1. 924 Pb-0

2. 82 2 . 83 2 . 85 5. 782 5. 692 5.452 Pb-0, 0-0

3.37 3.30 3.30 2 .888 2. 904 2. 978 Pb-0

3. 61 3.57 3. 65 1. 851 1.830 1. 882 Pb-Pb, Mg/Nb-Pb

1. No di stinct peak observed.

Table 1. Nearest neighbor peak positions and integrated

intensities . long and, to a lesser extent, intermediate Pb-0

distances as indicated by the integrated intensities.

At lOK the intermediate range has been depleted in favor

of both short and long Pb-0 nearest neighbor distances.

CONCLUSIONS

We find significant temperature dependent shifts in the

distribution of both Mg/Nb-0 and Pb-0 nearest neighbor

distances, the trend being toward a decrease in the

number of intermediate pair separations as the

temperature is reduced below the Curie temperature. We

97
believe the increase in number of extreme Mg/Nb-0 and

Pb-0 nearest neighbor distances below the Curie

temperature is responsible for bulk f erroelectricity in

PMN . The presence of these short and long pair

separations at all temperatures suggests that PMN is

locally ferroelectric above . Acknowledgements

The authors would like to acknowledge Prof. S. K. Kurtz

for the suggestion of applying the real space method to

relaxor f erroelect rics , Brian Toby and Prof. P. K.

Davies for useful discussions and Jiu-qi Huang for

assisting with analysis. Work at the University of

Pennsylvania was supported by the National Science

Foundation through the MRL Grant DMR-8819885. The

Intense Pulsed Neutron Source is operated as a user

facility by the U.S. Department of Energy, Division of

Materials Sciences, under contract W-31-109-Eng-38

.

REFERENCES

[1] L. E, Cross, "Relaxor Ferroelectrics,"

Ferroelectrics . Vol. 76, pp. 241-267, 1987.

[2] P. Bonneau, P. Garnier, E. Husson and A. Morrell,

"Structural Study of PMN Ceramics by X-ray

Diffraction Between 297 and 1023 k, " Mat. Res . Bull. , Vol. 24, pp. 201-206, 1989.

[3] G. A. Smolenskii and A. I, Agranovskaya, Soviet

Physics Solid State . pp. 1429, 1959.

[4] E. Husson, L. Abello and A. Morell, "Short Range

Order in PbMg;^ / 3^ ^ 2 / 3^ 3 Ceramics By Raman

Spectroscopy," Mat . Res . Bui 1 . . Vol. 25, pp.

539-545, 1990.

[5] H. D. Rosenfeld, T. Egami and A. Bhalla, "A Neutron

Powder Diffraction Study of Short Range Order in

Ferroelectric Pb (Mg^/3Nb2/3) O3, " (To Be Published)

[6] IPNS Note 19, IPNS, Argonne National Laboratory.

[7] B. E. Warren, X-ray Diffraction . Reading, Mass.:

Addison-Wesley, 1969, ch . 10, pp. 120-130.

98
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

MECHANISH OF PNN BASED PEROVSKITE CERAMICS FORMATION

Yukichi Sasaki, Akira Nagai, Tetsuo Yoshinoto

Odawara Research Center, Nippon Soda Co., Ltd. 345 Takada Odawara Kanagawa, 250-02 JAPAN

Introduction

It is well known that multiconponent lead perovskite materials are promising candidates

for applications in multilayer-capacitor. Pb(Hgi/3Nb2/3)03 (PHN), Pb(Nii/3Nb2/3)03 (PNN),

PbTiOs (PT) are representatives of lead perovskite materials. The dielectric constant

for the three component solid-solution system of PNN-PMN-PT is about 20,000 to 23,000

and the value of T.C.G(Thermal Coefficient of Capacitance) for the solid-soluton

satisfies the requirements of the Y5V specification by the E.l.A. When that solid solution is prepared by the conventional oxide nixing process, lead-based pyrochlore phases mainly occur at first. However, it is reported that

the columbite precursor method is very effective restraining the formation of the pyrochlore phase. The calcination reactions were studied in detail to confirm the relationship between the crystal structure and the calcining temperature which is much

influenced by powder processing. The XRD studies showed many differences in the calcination reaction between PNN and PMN in spite of having the same crystal structure

which has the general fomula Pb(Bli/3B22/'3)03. Therefore, for PMN there was little

pyrochlore phase produced by the columbite process, but for PNN the pyrochlore phase was formed as much as in the conventional oxide process even when the columbite process was used. And also the sintered pellets of the conventional oxide mixing process often

contained inclusions in the grains. The microstructure and the chemical compositions of the inclusions were determined by SEM-VIDS and TEH-EDS.

Experiiental Procedure

(1) Powder preparations

The material composition we used for this experiment was ternary multicomponent lead

perovskite. system, with the chemical formula 20Pb(Hgi/3Nb2/3)03-60Pb(Nii/3Nb2/3)03-

20PbTi03. Powders were prepared using three different methods and the flowcharts for those three methods are given in fig.l.

(2) Studies of the calcination reaction To investigate the calcination reaction each powder was heat-treated at 600-9001C,

using a magnesia crucible in a closed atmosphere. These calcined powders were identified by X-ray diffraction at a scanning rate of 2.4° /nin.

Furthermore, PMN and PNN, which are the main components of this solid solution, were investigated individually.

(3) Microstructure of sintered pellet

The microstructure of these sintered pellets were observed by TEM-EDS and SEM-WDS. The samples were polished and etched with 20vt% phosphoric acid solution before SEM observation.

Then for each polished sintered pellet(lOnin^disk), 18 points on the surface were analyzed by high resolution SEH-WDS and the measured values were statistically

calculated in order to investigate the homogeneity of the constituent elements.

99
Results and Discussion

(1) Calcination reaction of PNN-PMN-PT system (fig. 2) ©Conventional oxide process The calcination reaction on the conventional oxide process occurs approximately as

follows.

PbO+NiO+MsO+TiOj+NbzOs

~600t: -> Pyrochl ore ( 1

) +NiO+HgO+PbTi O3 <1>

850~700t: -» Pyrochl ore ( 1

) +Pyrochlore (2) +PbO

+(NiO+HgO+PbTi03) <2>

-> Pyrochlore (2) +PbO+ (NiO+HgO+PbTiOs) <3>

800~850T: -» Pyrochl ore ( 2) +PbO+Perovski te*+ (NiO+MgO) <4>

-> Perovskite* <5>

Pyrochlore(l):Pbi4Nbio039 1 NiO and/or HgO

Pyrochlore(2)IPb3Nb40i3 are solid-solved

Perovskite* :PNN-PHN-PT solide-solution

In this reaction, most of the PbO^NbaOs and TiOa were changed to pyrochlore(l) and PbTiOs at about 600°C. XRD peaks of the NbzOs and PbO vanished at that

temperature, and a part of pyrochlore(l) decomposed to pyrochlore(2) and PbO at 850-^

700°C. Then, the PbO was reproduced again in this system. The decomposition

reaction is Pb^NbioOag -> 5/2Pb3Nb40i3 + 13/2PbO, and it was accelerated at 700~750°C.

At SOO-^SSO'C other perovskite compounds were produced and formed a solid solution with

the PbTi03. Finally the reaction product was changed to single phase. On the other

two processes, this three component solid solution was produced under the similar

reaction mechanism, but there were some differences from the conventional oxide process.

©Colunbite precursor process

1. The amount of reproduced PbO was less in the columbite precursor process than in the conventional oxide process, and little PbO was observed at 800X1.

2. It is possible to obtain a single phase of perovskite in the columbite process at about 50°C lower than in the conventional oxide process. ®Alkoxide process

1. The same amount of PbO was reproduced during the calcining reaction in both the alkoxide process and the conventional oxide process, but the PbO was not observed

at 800T:.

2. The mechanism of the calcination reaction for the alkoxide process was the same as for the conventional oxide process up to 750°C, but in the alkoxide process the reaction proceeded quickly, and the complete monophase material was obtained at 800°C.

(2) Calcination reaction of PNN and PHN The calcination reaction was examined for PNN and PHN individually in order to investigate the influence of the powder processing in more detail (fig. 2). As the result it was confirmed that the mechanism of the calcination reaction depended

on the process for preparing the PHN. Vhen the PHN was prepared by the conventional

100
oxide process, that reaction proceeded like <1>~<5> reaction. But, by the colunbite precursor process the PMN perovskite phase vas produced directly in the calcination

reaction and little pyrochlore phase appeared. On the other hand, PNN did not depend

on the powder processing, and even though the previously prepared NiNbaOe was reacted

with PbO, the pyrochlore phase was appeared. The difference between the NiNb206 and HgNb206 in the reaction with the PbO is supposed as follows.

Thus, the PMN and PNN are the sane Pb(Bli/3B22/3)03 perovskite compounds, but there is

a difference in the calcination reaction according to the powder preparation processes between the PMN and the PNN.

(3)Micro3tructue study

Moreover, to confiri the above presunption the nicrostructure of the sintered pellets

was observed by SEM and TEM. After thermal and chemical etching, we could see small circular traces in SEM images. Since some of those traces were not falling off and

remained, we considered that these were not pores or defects but inclusions which were observed in the sintered grains by SEN (fig. 3) These inclusions were seen in both the conventional oxide process and the columbite precursor process. In the alkoxide process we could not see such inclusions as those

observed in the oxide process or colunbite process, and by TEM-EDS study, it is

understood that the chemical composition of these inclusions was different between the oxide process and the colunbite process. Namely, the inclusions in the conventional

oxide process pellets were composed of NiO and MgO but those in the colunbite process pellets were conposed of only NiO (fig. 4). The reason of this difference is considered

as follows.

l.In the conventional oxide process, the PbO reacted with NbaOs more preferably than NiO

or MgO, and the pyrochlore phase formed prior to the perovskite phase. So some of the free NiO and MgO renained as inclusions in the grains or between the grains after

perovskite formation was finished.

2. In the colunbite process, MgNbaO^ colunbite reacted with PbO to forn the perovskite

phase directly but NiNb206 colunbite reacted with PbO to decompose its structure and to form NiO and the pyrochlore phase to sone extent. So NiO only was observed for the inclusion in the sintered pellet.

Then we studied the conpositional honogeneity of the sintered pellets by SEM-VIDS analysis. The honogeneity of B-site ion constituents is not good in the oxide and is very good in the alkoxide process. For the colunbite process the nediun value between

the oxide process and the alkoxide process was obtained. Therefore, it is supposed that

this conpositional honogeneity neasured by SEM-WDS has a close relation with the presence of the inclusions in the sintered pellets.

Conclusion

(l)There is a large difference in the fornation reaction nechanisn between PMN and PNN

in spite of having the sane Pb(Bli/'5B22/3)0j perovskite structure. For PMN it is easy

PMN: 3PbO+MgNb206 -> 3Pb(Mgi/3Nb2/3)03 <6>

PNN: 3PbO+NiNb206 -> Pyrochlore(l)+NiO

Pyrochlore(2)+PbO+NiO

-» Perovskite

<7> <8>

<9>

101
to form the perovskite single-pahse compounds by using the columbite process. But for PNN it is difficult to forn the perovskite phase even if columbite process is used.

(2) The reaction nechanisn of PNN-PMN-PT ternary solid solution formation varies with the powder processing method, and this is attributed mainly to the reaction mechanisms of the PMN formation.

(3) The chemical composition of the inclusions in the sintered pellets varies with the

powder processing methods, and this inclusion formation has a strong influence on the compositional homogeneity of the sintered pellets.

l.S.L.Swartz, T.R.Shrout Mat.Res.Bull., 17 (1982) 1245-1250

2.T.R.Shrout, Arvind Halliyal Am.Ceram.Soc.Bull., 66[4] (1987) 704-711 a.H.Backhaus-Ricoult, D.Ricoult J. Materials Sience, 23 (1988) 1309-1316

4. A.Nagai, Y.Sasaki, T.Yoshimoto Perroelectrics, 102 (1990) 117-127

5. J.P.Guha, H.U.Anderson J.Am.Ceram.Soc, 70[3] (1987) C-39 - C-40

PbO„ KtO, NiO IiO„ |lb,0.

V « i { h 1 0 t

Ball NlUlrK

Calcination

Ball HilHnt

P 0 » i t r

nio, Nio. no,

«b,0.

y<l(htin(

Ball Hilllnt

Calcination

Ball Killint

Calcination

Ball Iflllinc

P 0 a d « r

PbO

Pb(OAc),. NKOAc),

Nc(OAc).

(In H,0)

Solution 1

TKOBu).

Et.NH

. Nb(OBuK

(in EtOH)

S 0 1 u t 1 0 n 2

Bydrol ysi s

Filtration

Dry U f

Calcination

Ball nillint

P 0 . dor

Conventional oxide process Coluibita precursor process Alkodde process

Figure 1 Processing flowchart for powder preparation.

102
PNN-PHN-Pf

(Conventional onide process)

PNN-PMN-PT (Colmbite precursor process)

pnV-phn-pt

(Alkoxlde process)!

-H—

inzt

zfi

• PirrMkl«r-<U AM

• rrr«U<r>(U

I I I i -i <ViA I

'^Vi ^i -i^ I I I I I 1

I I I r

iL Wan caIcImUm

'PNN' (Conventional oxide process)PMN (Conventional oxide process)

u

J A

—J—no -A-„

\ 1^ 1^ -^1 I ' I

J L

PNN (Coluiblte precursor process)

II.

PMN (Coluiblte precursor process)

Ilt/LAA_A_><-.

-A_„

T ! r I I ! ! "I" I

I I I T-T—r-i—r i i [ -r i -t- r i i r i i [

Ofar*nUU

0^^wU•r^d)

Figure 2 X-ray diffraction patterns of the calcineci

powder.

103
Figure 3 SEM photographs of the etched surfaces.

The surface was chemically etched by 20vjt^ phosphoric

acid after thermal etching.

1 nc 1 US! Oil 1 , 1

i ,

1

p. latrix —

i ; t - - —

--: -Jl - -

•

f! '

' i \

Conventional oxide process

! ncl us i on 3

i i !; !

I ; 1 ' i 1

l^nclusioii 4 -Hi

•Se -jSj- ; i J. ;

Columbite precursor process

Figure 4 TEM photographs and EDS spectrums of sintered

pellets.

104
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

CONVENTIONALLY PREPARED SUBMICRON ELECTRO-CERAMIC

POWDERS BY REACTIVE CALCINATION

Thomas R. Shrout

Pennsylvania State University

Materials Research Laboratory

University Park, PA 16802

ABSTRACT

It is the physiochemical nature of lead-based perovskites Pb(B)03 that upon

reaction of the component oxides a large volumetric expansion occurs. In addition

to perovskites, Pb- and Bi-based pyrochlores (A2B2O7) and the layered structure compound Bi4Ti30i2 also exhibit similar physiochemical behavior. At the

temperature of maximum expansion, the associated morphological development

results in a porous skeletal type structure consisting of fine particulates that can be

readily broken down further by milling. The level of expansion and ease of

comminution was shown to be strongly dependent on the starting powder size,

using perovskite Pb(ZrTi)03 as the example. Using this concept of "reactive

calcination," the state of optimum soft agglomeration and subsequent milling can

allow for fully reacted powder with submicron particle size. INTRODUCTION

With the continuing miniaturization of electronics including ceramic devices

such as multilayer capacitors, ultrasonic transducers, sensors, piezo motors, and

electrostrictive actuators, the importance of fine and phase pure starting powders

becomes ever so more evident. Such powders also allow the potential for novel

applications such as ferro-fluids which require finely dispersed submicron powders

in a liquid medium.(^) Fundamentally, submicron to nano-sized powders are of

interest in the understanding of the role of "scale" on intrinsic physical phenomena

such as ferroelectricity and related behavior.(2)

Numerous methods to produce submicron powders have been developed

with emphasis on the perovskite ABO3 family of materials. Methods include

chemical synthesis techniques such as co-precipitation of alkoxides, molten salt,

and hydrothermal. However, most chemical synthesis techniques are relatively

costly and do not lend themselves for mass production.

Recently, submicrometer powders of Pb-based perovskites including

PbTi03, PbZr03, Pb(Zr.53Ti 47)03 and Pb(Mgi/3Nb2/3)03 were prepared by a reactive calcination process.^^) Using only reagent-grade raw materials and

conventional processing techniques, highly reactive powders were produced by

reacting the materials near the temp>erature of maximum volumetric expansion. At

this point, the associated morphological development results in a skeletal-type

105
structure consisting of ultra-fine particulates than can be readily broken down by

milling (see Fig. 1). Powder sizes < 0.1 )im were produced allowing enhanced

densification and correspondingly small grain sizes.^"^*^)

The objective of this work was to further optimize the "reactive calcination"

process of lead-based perovskite through control of the starting powder

characteristics. It was a further objective to explore the potential of reactive

calcination in non-PbO perovskites and other electro-ceramic structural families

through knowledge of their physiochemical behavior.

Figure 1. Schematic representation of the perovskite PbBOs formation and

associated morphological changes.^^)

EXPERIMENTAL PROCEDURE

The Pb-based perovskite material chosen to examine the role of component

powder characteristics on reactive calcination was the solid solution

Pb(Zr.53Ti.47)03 (PZT). Non-perovskite materials investigated included the

pyrochlores Pb2Nb207 and Bi2Ti207 and Bi4Ti30i2 layer structure compound.

TTie non-PbO perovskite (Nai/2Bii/2)Ti03 was also suidied.

Reagent-grade raw materials were used to prepare the compounds as described in ref. 3. The component powders were characterized by determining

their as-received surface area and particle size. Since intermediate reactions occur in the formation of the complex

perovskite PZT, the B-site precursor method developed by Swartz and Shrout^^)

was used whereby the Zr02 and Ti02 oxide were pre-reacted prior to reaction with

PbO as follows:

0.53 Zrq + 0.47 TiO, ) Zr^Ti,,q

^^^^^^ PbO +Zr^3Ti,4702 PZT (perovskite) [2]

Variations in particle sizes of both the PbO and Zr,Ti02 oxides were

prepared by vibratory or attrition milling. The powder size schemes for reactive

calcination of PZT are schematically shown in Fig. 2.

106
Stoichiometric mixtures of the component oxides for all the compounds

studied were prepared by vibratory mixing in ethyl alcohol or de-ionized H2O and

polyelectrolyte dispersant at a pH where minimal dissolution of Pb"'"^ occurs.(7)

Upon drying, the powders were pressed into disks of similar packing densities.

The possibility and/or extent of reactive calcination was determined by the

degree of volumetric expansion and corresponding temperature reaction range.

This was performed by placing pressed disks in a furnace (open air) and heated at a

rate of ~ 100*C/hr. Above SOO'C, disks were quenched at 25 to 50°C intervals by

being pulled directly from the furnace and allowed to cool in air. Sample

geometries were used to determine the degree of volumetric expansion or

shrinkage. The disks were then crushed and powder XRD analysis performed to detennine the amount of phase reaction and subsequent particle characteristics.

SEM microstructural analysis of selected powder compacts (before and after

crushing) were also examined to examine morphological development during

reaction.

Based on these results, large quantities of (~ 1 kg) PZT packed powders

were prepared. The powders were calcined for 4 hrs approximately 50° C below the

maximum expansion (to allow for kinetics) and at higher temperatures to demonstrate the point of optimum skeletal formation and subsequent consequence

on milling efficiency. Slurries of the calcined powders were prepared as before and

attrition milled while samples were removed at various intervals. The milled

powders were then characterized by specific surface area. The primary particle size

was determined through the equation

PSA [3]

where p is the theoretical density, SA the surface area, and D the calculated primary

particle size.

Figure 2. Component powder size schemes for the formation of perovskite PZT.

107
RESULTS AND DISCUSSION

Instinctively, one expects that ceramic powder compacts should shrink on

reaction as a consequence of higher product densities, as found for non-PbO

perovskites. The volumetric expansion observed during the formation of Pb-based

perovskites is reportedly the result of molar volume differences in the component

oxides and subsequent perovskite phase. Attempts to predict relative volume

expansions for various compounds clearly suggests the expansion behavior is also

dependent on the particle packing and size characteristics as proposed in Fig. 2. The importance of particle characteristics on reactive calcination is clearly

evident in Fig. 3, showing the impact of PbO particle size on the volume expansion

during the formation of perovskite PZT. As presented, fine PbO (Case I) resulted

in shrinkage prior to a relatively small volume expansion and associated formation

of the perovskite phase. The shrinkage is believed to be the result of the highly

reactive nature of the fine PbO particle/agglomerates which sinter prior to reaction

with the coarser Zr,Ti02 particles. The impact of Zr,Ti02, particle size, clearly

shows that the finer material enhanced the formation of perovskite PZT, but the

reactive nature of the fine powder and less open skeletal structure lead to densification at relatively low temperatures (Case IV). For Case m, where PbO

and Zr,Ti02 are of similar size, resulted in the largest volume expansion. Of

particular interest is that once the skeletal structure was formed, densification did

not occur even up to temperatures of KXWC. Hence, inhomogeneous densification

(low coordination number of reaction regions) was optimized. The importance of

this behavior is clearly evident in Fig. 4, whereby the ease in which the skeletal structures can be broken down is given in terms of a milling factor. As shown, the

skeletal sructure associated with Case III was readily broken down by milling with an order of magnitude decrease in particle size in only one hour. A large volume

change found over an extensive temperature range (Case m), thus, allows for a broader calcination process window. Further implications of the morphological

development of these materials will be presented in the summary.

The morphological behavior presented above was not found to be unique to Pb-perovskites but was also observed in the Pb-based pyrochlore phase Pb2Nb207

as shown in Fig. 5. Analogous to Pb, the Bi-based perovskite Nao.sBio.sTiOa,

pyrochlore Bi2Ti207 and layer structure Bi4Ti30i2 compounds also exhibited

similar behavior, but to a lesser degree. It is important to note that at the point of

maximum volumetric expansion corresponds to the formation of the expected

phase.

CONCLUSION

The volumetric expansion and morphological development associated with

the formation of Pb-based perovskites was found to be strongly dependent on the

starting powder characteristics, using perovskite Pb(ZrTi)03 as the example. The

resulting porous skeletal structure of perovskite particulates was found to be readily

broken down into submicron particles by milling. In addition to Pb-based

perovskites, Pb- and Bi-based pyrochlores (A2B2O7), the layered structure

compound BL4Ti30i2 and (Nai/2Bii/2)Ti03 perovskite were also found to exhibit

108
similar physiochemical behavior. Using this concept of "reactive calcination" the

point of maximum expansion can be used as a characterization tool to determine the

optimum state of soft agglomeration to provide highly reactive powders.

Further investigation and modeling of particle size and packing effects on

reactive calcination are still required.

50

-20 -I 1 • 1

' 1 1 -1 500 600 700 800 gOO 1000 1100

Temperature CC)

Figure 3. Volumetric expansion-shrinkage characteristics of PZT as a function of

temperature for various component sizes.

Milling Time (Hrs)

Figure 4. Milling factor (normalized) for reactively calcined PZT (Cases III (3.1

lom) and IV (0.9 pm)) as a function of milling (attrition) time.

109
80

-40 I

1

' 1

'

1

' 1

' 1 ' —

400 500 600 700 800 900 1000

Temperature (°C)

Figure 5. Volumetric expansion-shrinkage characteristics of Pb2Nb207, Bi2Ti207

(pyrochlore), Nao.sBio.sTiOs (perovskite) and Bi4Ti30i2 (larger structure).

ACKNOWLEDGEMENTS

The author wishes to thank Beth Jones for her technical skills and Drs.

Venkataramani, G.S. Lee, and J.H. Adair for their helpful discussions.

REFERENCES

1. J.U. Muller and K. Bamer. "Polydisperse Suspensions of BaTiOa

Particles," Ferroelectrics, IM, 83-88 (1990).

2. K. Uchino, E. Sadamaga, and T. Hirose. "Dependence of the Crystal

Structure on Particle Size in Barium Titanate,"/. Am. Ceram. Soc, 22[8],

1555-1558 (1989).

3. T.R. Shrout, P. Papet, S. Kim, and G.S. Lee. "Conventionally Prepared

Submicrometer Lead-Based Perovskite Powders by Reactive Calcination,"

/. Am. Ceram. Soc, 21[3], 1862-1867 (1990).

4. P. Papet, J.P. Dougherty, and T.R. Shrout. "Particle and Grain Size

Effects on the Dielectric Behavior of the Relaxor Ferroelectric

Pb(Mgi/3Nb2/3)03," (to be published in J. Mat. Res.).

5. S. Kim, G.S. Lee, and T.R. Shrout. "Fabrication of Fine Grain

Piezoelectric Ceramics Using Reactive Calcination," (to be published in J. Mat. Sci.).

6. S.L. Swartz and T.R. Shrout. "Fabrication of Perovskite Lead Magnesium

Niobate," Mater. Res. Bull., H, 1245-1250 (1982).

7 . M. Pourbaix. Atlas of Electrochemical Equilibria in Aqueous Solutions. National Association of Corrosion Engineers (1974).

110
LOW-TEMPERATURE SYNTHESIS

AND CHARACTERIZATION
I

I
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

PREPARATION OF YBa2Cu307.6 FROM HOMOGENEOUS

METAL ALKOXIDE SOLUTION; SYNTHESES OF MIXED-LIGAND

COPPER (H) ALKOXIDES

D.M.Millar* andD. A. Payne

Department of Materials Science and Engineering,

Materials Research Laboratory, Beckman Institute,

and Science and Technology Center for Superconductivity,

University of Illinois at Urbana-Champaign, Urbana, EL 61801

ABSTRACT

Mixed-ligand copper (II) alkoxides of the type (C5H702)2Cu2(M.-

OR)2,C5H702 = 2,4-pentanedionate, R = CH2CH2OCH3 (1) and n-C4H9 (2),

have been synthesized as precursors for the preparation of superconducting

YBa2Cu307 thin layers. Crystallization from n-butanol yielded crystals of 2

suitable for x-ray structure determination. The molecule consists of two

approximately planar n-butoxy-bridged dimers (C5H702)Cu(|j,-

OC4H9)2Cu(02C5H7) linked by four apical Cu-O bonds (2.443(4)A) involving

each of the n-butoxy-oxygen atoms. The dimers are situated such that the two

CU2O2 units are rotated 90° with respect to each other. Comparison of the IR

spectra of compounds 1 and 2 suggest that in 1, dimer Unkage also occurs between

one pentanedionate oxygen atom of each unit and a copper atom of the opposing

dimer. This dual coorchnation of the pentanedionate prevents subsequent reactions

of this ligand with barium alkoxides, making compound 1 a suitable precursor for

formation of 1:2:3 Y:Ba:Cu solutions. Superconducting thin-layers have been

prepared from these solutions by spin-casting methods.

INTRODUCTION

Ever since Bednorz and Miiller^ first reported the unusually high

temperature superconductivity in (La,Ba)2Cu04, increased attention has been

focussed on a broad class of high Tc ceramic superconductors. In addition to attempts to increase Tc by new oxide compositions, much work has been directed

towards processing superconducting oxides such as YBa2Cu307.s into

technologically useful forms.^ Of particular interest has been the preparation of thin

layers due to their potential applications on semiconductor and electronic devices.

Preparation of such layers by rf sputtering,^ electron beam de-composition,^ or metal-organic chemical vapor deposition^ methods have been successfully

demonstrated. However, these techniques require sophisticated equipment and

* Currently with Central Research, The Dow Chemical Company, Midland, MI, 48674

113
extreme processing conditions. Deposition of YBa2Cu307.6 layers from solutions

of the corresponding alkoxides offers the possibility of avoiding these extreme

processing conditions, as well as the potential for preparing highly

homogeneous products.

As part of a larger effort of research on the chemical processing of electrical

ceramic materials, we have investigated the preparation of YBa2Cu307^ thin layers

and powders from homogeneous metal-alkoxide solutions. This paper describes

the synthesis and characterization of a series of unique mixed-ligand copper (II) alkoxides of the type (C5H702)2Cu2(|i-OR)2.

/

Hf

IR CH,

This system avoids the usual oUgomerization of copper (II) alkoxides, and

the subsequent insolubility by filling two copper coordination sites with a chelating

B-diketonate, 2,4-pentanedionate (C5H7O2). The remaining coordination sites

about the copper atoms are filled with doubly bridging alkoxy groups. Of particular

interest is the system where these bridging alkoxides are derived from 2-

methoxyethanol, R = CH2CH2OCH3. Reaction of this readily soluble material

with the corresponding yttrium and barium alkoxides results in homogeneous 1:2:3

Y:Ba:Cu solutions from which thin layers or powders can be obtained.

EXPERIMENTAL

REAGANTS, SOLVENTS AND GENERAL PROCEDURES

The following chemicals were purchased from commercial sources and used

without further purification: Cu(N03)2 • 2H2O (Mallinkrodt) 2,4-pentanedione

(Aldrich); NH4OH (Mallinkrodt); CH3OH (Baker); KOH (Fischer); n-butanol

(Aldrich); barium granules (Alfa); yttrium isopropoxide (Alfa). HPLC grade 2-

methoxyethanol (Aldrich) was obtained in Sure-Seal bottles and transferred and

stored under dry nitrogen.

Sapphire wafers (1.5 cm diameter, 350 |im thickness, one side polished,

0001 axis parallel to surface) (Insaco, Inc.) and platinum substrates (12 mm x 12

mm X 0.5 mm) were cleaned by sequential ultrasonication in trichlorethylene,

acetone, and 2-methoxyethanol, then air dried.

Thin layers were prepared as described below by spin-casting onto the

substrates using a Headway Research Photoresist Spinner.

114
ANALYTICAL PROCEDURES. Elemental Analyses were carried out by the

University of Illinois School of Chemical Sciences Analytical Laboratory. Mass

Spectrometry was done in the School of Chemical Sciences Mass Spectrometry

Laboratory. FTIR spectra were obtained from KI pellets using an IBM IR/32

spectrometer. Single-crystal X-ray diffraction studies were carried out by Dr. Scott

Wilson in the School of Chemical Sciences X-ray Diffraction Laboratory.

PREPARATION OF (C^H209)2Cu (ID.8 A IL flask was fitted with a pressureequalizing addition funnel andloaded with 50 g (0.224 mol) Cu(N03)2- 2H2O and

500 mL deionized water, followed by 75 mL concentrated NH4OH. To the

rapidly stirred, deep blue solution, 55 mL (0.534 mol) 2,3-pentanedione were

added dropwise over a 1 hr period. Pale blue (C5H702)2Cu was collected as a wet

precipitate by vacuum filtration 30 minutes after complete addition. This material

was dissolved in 900 mL HCCI3 and a clear colorless, aqueous layer separated

from the deep blue organic phase in a separatory funnel. The volume of the HCCI3

solution was reduced one-third by removal of the solvent by distillation. Upon

slow cooling to room temperature, dark blue crystalline (C5H702)2Cu was obtained

and collected by vacuum filtration. This procedure was repeated until no further

crystalline product was obtained. A total of 46.17 g (0.178 mol) were collected, a yield of 80% based on the original amount of Cu(II). IR (KI pellet, 1400-400 cm1) 1356(s), 1275(s), 1190(m) 1020(s), 937(s), 783 (vs), 685(m), 654(m), 613(m),

455(vs), 430(w).

PREPARATION OF (Cs H702)2Cu9(|li-OCH2)2 d^-^ A IL flask fitted with a pressure equalizing addition funnel and a reflux condensor vented to a Nujol bubbler

was loaded with 5.4 g (20.9 mmol) (C5H702)2Cu and 500 mL CH3OH. The

addition funnel was charged with a methanolic KOH solution (1.2 g KOH in 50 mL

CH3OH). This solution was added dropwise to the refluxing (0511702)2 Cu

suspension over a 2 hr period. Thirty minutes after complete addition, heating was

stopped and the mixture cooled to room temperature. Vacuum filtration of the

resulting suspension produced 3.18 g (8.2 mmol) violet (0511702)2 Cu2(|a,-

OCH3)2, a yield of 79% based on the original amount of Cu(II). Anal. Calcd for

C6Hio03Cu:C,37.21; H,5.21; Cu, 32.81. Found: 0,37.14; H,5.23; Cu,32.91.

IR (KI pellet, 1400-400cm-l, bands in addition to those for (C5H702)2 Cu) 1071

(s), 554 (vs).

PREPARATION OF (C5H702)2Cu2(^-OCH9CH90CH2)9 (2) A 100 mL flask

was loaded with 1.0 g (2.5^ mmol) of 1 and 4'0 mL 2-metEoxyethanol added by

syringe. A reflux condensor vented to a N2 inlet/Nujol bubbler was added and the

slurry heated in a 90°C oil bath. After 1.5 hr at this temperature, all volatiles were removed under vacuum. Turquoise- solid 2 was obtained as a fine powder and

1.19 gl (2.50 mmol) were collected, a yield of 97% based on the original amount of

Cu(II). Anal. Calcd for C8H14O4CU: 0,40.42; H. 5.95; Cu, 26.73. Found: C,

40.23; H, 5.99; Cu, 26.62. MS. m/z = 476 (M+). IR (KI pellet, 1400-400 cm-1,

bands in addition to those for (C5H702)2 Cu) 1240 (w), 1123 (s), 1075 (s), 966

(m), 901 (m) 839 (w), 764 (m) 588 (m) 575 (w), 513 (m), 450 (m).

115
PREPARATION OF (CsH202)2Cu2(|Li-OC4HQb (3) This compound was prepared

identically to 2 above using n-butanol as the solvent. Blue-green 3 was obtained as a fine powder with 1.21 g (2.56 mmol) collected, an essentially quantitative yield. IR

(KI pellet, 1400-400 cm-l, bands in addition to those for (C5H7O2) Cu) 1240 (w),

1117 (m), 1073 (s), 950 (w), 840 (w), 768 (m), 630 (w), 552 (m), 507 (m).

PREPARATION OF Ba(OCAHQ)7 SOLUTION. In a glove box, an oven-dried

50niL flask was loaded with 3.43 g (25.0 mmol) Ba granules, and an oven-dried

addition funnel and reflux condenser with N2 inlet added. In a fume hood, lOmL

Na2S04 dried pentane was added by syringe and the suspension stirred vigorously.

To this mixture, 25 mL (25.0 mmol) n-butanol was added over 1.5 hours. Stirring at room temperature was continued until gas evolution ceased and no solids remained,

about 18 hours. Removal of pentane under vacuum yielded a cloudy, viscous

solution. The fmal clear solution was obtained after filtration through a 0.45 |J.m filter and dilution to 25 mL volume with n-butanol. Analysis for Ba determined a final

concentradon of 0.90 M.

PREPARATION OF 1:2:3 Y:Ba:Cu SOLUTIONS. In a typical preparation, an ovendried 25 mL flask was charged with 300 mg (0.631 mmol) of 2, thoroughly purged

with dry nitrogen, and 10 mL 2-methoxyethanol added. To this clear, deep blue

solution, 631 [iL of a IM Ba(OCH2CH20CH3)2 (0.631 mmol) 2-methoxyethanol

solution were added and the homogeneous solution stirred rapidly at room

temperature. Over the next 1.5 hours, finely divided Cu(OCH2CH20CH3)2

precipitated and was removed by filtration of the solution by cannula. Addition of 0-

90 equivalents (based on yttrium) of deionized water was followed after 30 minutes

by addition of 315 ^L of a IM Y(OCH2CH20CH3)3 (0.315 mmol) 2-

methoxyethanol solution. After vigorous stirring for 2 hours at room temperature,

the clear solution was filtered by cannula and stored in a refrigerator (5°C) for 24

hours. Analysis for Y, Ba, and Cu determined a ratio usually on the order of

1:2.2:2.9.

PREPARATION OF LAYERS ON PT. Prior to deposition, each solution was

filtered through a 0.2 fim filter. Twelve sequential layers were deposited with a 10

second, 300°C heat treatment between layers and a final 30 second, 300°C heat

treatment upon completion. Layers were spun at 2000 rpm for 50 seconds.

PREPARATION OF LAYERS ON SAPPHIRE. Similarly, layers were spin-cast

from solutions filtered through a 0.2 |j,m filter for 50 seconds at 2000 rpm. Between

layers, substrates were placed in a 630°C box furnace (Lindburg) for 15 seconds,

then quickly cooled to room temperature. Ten layers per substrate were deposited in

this manner and the resulting specimens stored in a desiccator.

RESULTS AND DISCUSSION.

Although the alkoxide chemistry of copper (II) has been well explored,

such products are generally non-volatile and insoluble in common organic solvents.

These properties, a direct consequence of the oligomeric nature of copper (II)

alkoxides, can be ascribed to the strong tendency of the metal for higher coordination.

These coordination requirements can partially be satisfied by chelation of the copper

116
(II) with 6-diketonate ligands such as 2,4-pentanedionate, C5H7O2. I. In this

manner, soluble mixed-ligand copper (II) alkoxides can be prepared.

In 1965, Bertrand and Kaplan^ first reported the synthesis of a dimeric 2,4-

pentanedionate, methoxy-bridged compound formulated as (C5H702)2Cu2(}J.-

0013)2 (1) based on spectroscopic and molecular weight measurements. This basic

dimeric structure was later confirmed by Andrew and Blake^^ as the essential unit in

the solid state structure of (C5H702)2Cu2(|i-OCH2C6H5)2, prepared by alcohol

exchange of 1 with benzyl alcohol. This compound was not found to crystallize as discrete Cu2-dimers, but rather as strongly associated pairs of dimers. The authors

also reported the synthesis by alcohol exchange of the ethoxy and n-propoxy-bridged

compounds. Based on the spectroscopic similarities observed between these

systems, it was suggested that compounds of the type (C5H702)2Cu2(|i-OR)2

possess a similar associated-dimer solid state structure. In addition, each of these

compounds shows some solubility in the parent alcohol, as well as limited solubility

in other polar organic solvents.

As detailed in the experimental section, the methoxy-bridged dimer 1 readily

undergoes alcohol exchange to produce, in addition to those compounds just

described, products of the type (C5H702)2Cu2(|i-OR)2 (R = OCH2Ch20CH3 (2),

n-C4H9 (3)) according to reaction 1:

(C5H702)2Cu2(^l-OCH3)2 + 2ROH ^(C5H702)2Cu2 (OR) + 2CH3OH 1

90PC

Although compounds 1, 2, and 3 share the same dimeric formulation, they differ

markedly in their solubility in the parent alcohol, and in the case of 2 and 3, in their

subsequent reactivity with Ba(OCH2CH20CH3)2. Violet solid 1 is sparingly soluble

in methanol, whereas turquoise dimer 2 demonstrates good room temperature

solubility in 2-methoxyethanol. Blue solid 3 is soluble in room temperature n- butanol, but is an order of magnitude less soluble in this solvent than 2 is in 2-

methoxyethanol.

The range of solubilities and disparate reactivities for compounds 2 and 3

could be a simple consequence of the nature of the bridging alkoxide, or they may

reflect more subtle structural differences, such as a different degree of dimer

association. Consequentiy, attempts were made to determine the solid-state structures of compounds 1, 2, and 3. Single crystals of 3 were readily obtained from

solutions of hot n-butanol, however, no suitable solvent for the crystallization of 1

was found. Crystals obtained of 2 using a variety of crystallization procedures and

solvent systems, were consistentiy twinned and unsuited for X-ray analysis.

Figure 1 shows the results of a single-crystal X-ray determination for crystals

of dimer 3.^2 The structure consists of two approximately planar n-butoxy-bridged

dimers (C5H702)2Cu(|J.-OC4H9)2Cu(02C5H7) joined by four apical Cu-O bonds

involving each of the n-butoxy oxygen atoms. Each copper atom possesses

approximately square-planar coordination. The dimers are situated such that the two

CU2O2 units are rotated 90° with respect to each other. This unit is essentially planar

117
Figure 1 . Perspective view of compound 3 with atoms labeled. Thermal ellipsoids are drawn at the 35% probability level for the non-hydrogen atoms.

118
in each dimer with mean Cu-O bond lengths of 1.94A and are linked by 2.44A apical

Cu-O bonds. The planes of the dimers containing the 2,4-pentanedionate ligands are slighdy bent away from each other with the n-butoxy groups extending perpendicular

to these planes. Distances and angles within the 2,4-pentanedionate and n -butoxy

ligands are as expected.

The structure of 3 can be contrasted with that reported by Andrew and

Blakel2, the only other structurally characterized (C5H702)2Cu(|i-OR)2 system

reported. Both structures are similar in that they consist of apically associated

dimers. However, in the structure of the benzoxy product, the dimers are linked by

Cu-O bonds involving one benzoxy and one 2,4-pentanedionate oxygen atom of each

dimer. Consequently, the two dimers lie above and below each other without rotation

of the CU2O2 units. Within each CU2O2 unit, the two Cu-O bonds to the threecoordinant bridging oxygens are significantly shorter (1.88 A) than those to the fourcoordinant bridging oxygens (1.94 A).

In the absence of single crystal data for compounds 1 and 2, information

regarding their molecular structures can be inferred by comparison of their infrared

spectra with those of 3 and the benzoxy-bridged product. Figure 2. Absorptions

assigned to the (C5H702)2Cu unit^^ are indicated. Of particular interest is the region

of Cu-O absorptions found between 570 and 420 cm-l. In the spectrum of

compound 3, Figure 2c, three distinct Cu -O absorptions are observed and assigned

to 455 cm-i, Cu-O (2,4-pentanedionate); 507 cm-^ Cu-O (intramolecular bridged

alkoxide); 552 cm-l, cu-O (intermolecular bridged alkoxide). These Cu-O

assignments are based on the relative bond lengths (1.94A-intramolecular vs. 2.44Aintermolecular) of the corresponding units.

Examination of the Cu-O absorption region in the spectrum of

(C5H702)2Cu2(|i-OCH2C6H5)2, Figure 2d, shows the metal-oxygen bands

consistent with the molecular structure reported. As described earlier, the two dimer

units are linked through one of the two bridging benzoxy- and one of the chelating 6-

diketonate ligands. As a result, two distinct Cu-O (2,4-pentanedionate) absorptions,

451 and 448 cm-1, and two Cu-O (intramolecular bridged alkoxide) bands 580 and

548 cm-1, are observed. In addition, a weak absorption in the region assigned to the

dimer-linking Cu-O (intermolecular bridged alkoxide) bond, 490cm- 1, is observed.

No band could be assigned to the second type of dimer-linking Cu-O (2,4-

pentanedionate) band seen in the structure.

In comparison with the IR spectra of 3 and the benzoxy-bridged dimer, the

spectra obtained of dimers 1 and 2, Figures 2a and 2b respectively, show similarities

and significant differences. In the spectrum of 1, only two metal-oxygen absorptions

are observed. The lower frequency band, 455 cm-1, falls in the region assigned to the Cu-O bond to an unperturbed 2,4-pentanedionate ligand, as in (0511702)2 Cu^^

and dimer 3. By analogy with the group assignments previously made, the

absorption at 554 cm^l is assigned to the Cu-O bond of an intermolecular bridged

methoxide ligand. The absence of any bonds in the intermolecular bridged alkoxide

region suggests that compound 1 does exist as a singular, unassociated methoxybridged dimer as originally formulated.

119
I

1 r—I—I—I t—T I r~l

600 500 400 600 500 400

CM

Figure 2. FTIR spectra (650-400 cm-1) of KI pellets of A) compound 1, B)

compound 2, C) compound 3, and D) (C5H702)2Cu2(^l-OCH2C6H5)2.

120
Interestingly, the IR spectrum of the methoxyethoxy bridged dimer 2, Figure

2b, exhibits characteristics observed in both 3 and the benzoxy-compound. As seen

in these systems, bands are present which correspond to both intramolecular, 561 cm- ^ and intermolecular 512 cm'^, bridged alkoxides. It must be noted that in contrast

to the phenoxy-system, only one of each type of absorption is found. However, an

additional band at 433 cm-i is observed with the 455 cm-l Cu-O (2,4-

pentanedionate) moiety suggesting a linkage of dimers through an oxygen of the B- diketonate. These results suggest that dimer 2 may well possess a solid state structure similar to that found for (C5H702)2Cu2(|J.-OCH2C6H5)2. These studies

clearly demonstrate that generalization of the molecular structure for these types of

alkoxy-bridged dimers cannot be made based solely on observed physical behavior.

As was stated earlier, the differences between dimers 2 and 3 extend beyond

their relative solubilities and structures. Most significant is the behavior of 2 is such

that homogeneous 1:2:3 Y:Ba:Cu solutions can be prepared, whereas they cannot by

using 3. Solutions of Y:Ba:Cu with stoichiometries of 1:2:3 were obtained by

reaction of 2 with IM 2-methoxyethanol solutions of Ba(OCH2CH20CH3)2, and

Y(OCH3CH20CH3)3 in the ratio 2:2:1 (see Experimental Section). This

stoichiometry was required due to reaction of 2 with the barium alkoxide which

resulted in the precipitation of one equivalent of Cu(OCH2CH20CH3)2 Attempts to characterize this reaction and to identify the soluble barium and copper species using

NMR spectroscopy were complicated by the presence of paramagnetic copper (II). As a result, resonances for species other than the solvent employed were not

observed. A similar reaction between 3 and the corresponding barium alkoxide also

produces copper precipitation as Cu(OC4H9)2, however all of the copper and 80% of

the barium are lost from solution. Similar destructive reactions between 2 or 3 with

the yttrium alkoxide solutions were not observed.

Solutions of Y, Ba, and Cu suitable for deposidon of thin layers and

•preparation of powders were prepared according to reaction 2, where R = OCH2CH2OCH3 and n = 9-90:

Y(OR)3 + 2 Ba(OR)2 + 2(C5H702)2Cu2(^l-OR)2 + nH20 ^ 2.

Characterization of layer deposited on platinum and sapphire wafers has been

reported. 14 From such layers, superconducting films of YBa2Cu307 have been

obtained.

CONCLUSIONS

Two unique mixed-ligand copper (II) alkoxides, (C5H702)2Cu2(HOCH2CH20CH3)2. and (C5H702)2Cu2(|i-OC4H5)2.were synthesized by alcohol

exchange of (C5H702)2Cu2()i-OCH3)2. A single crystal X-ray diffraction study

determined the structure of the n-butoxy bridged dimer to be two apically associated

dimers linked by Cu-O bonds to each of the two bridging alkoxide oxygen atoms.

The CU2O2 units of each dimer are rotated 90° with respect to each other. Comparison

of the FTIR spectra of methoxy- and methoxyethoxy-bridged dimers with the n121
butoxy and benzoxy-bridged products suggested that (C5H702)2Cu2(m.-OCH3)2

exists as discrete, unassociated dimers in the solid state. Similar comparison

suggests that (C5H702)2Cu2(^-OCH2CH20CH3)2 and (C5H702)2Cu2(|lOCH2C6H5)2 possess similar solid state structures.

Homogeneous solutions of Y:Ba:Cu can be obtained with stoichiometrics of

1:2:3 by reaction of (C5H702)2Cu2(|i-OCH2CH20CH3)2 with IM solutions of

Ba(OCH2CH20CH3)2 and Y(OCH2CH20CH3)3. Hydrolysis of these solutions

with 9-90 equivalents of water per equivalent of yttrium results in homogeneous

solutions from which both powders and thin layers can be prepared.

ACKNOWLEDGEMENTS

This work was supported by the .U.S. Department of Energy, Division of

Materials Sciences, under contract DE-ACO2-76ER01 198. We gratefully

acknowledge the assistance of S. R. Wilson for X-ray diffraction studies, and with

W. G. Fahrenholtz for ceramic processing.

REFERENCES

1. Bednorz, J. G.; Miiller, K. A., J. Phys. 6-Condensed Matter, 1986. 64, 189.

2. See, for example, Advanced Ceramic Materials, 1987, 2 (3B), 267.

3. Adachi, H.; Setsune, K.; Mitsuyu, T.; Hirochi, K.; Ichikawa, Y.; Kamada, T.; Wasa, K. Jap. J. App. Phys. 1987 . 26 (5), 709.

4. Naito, N. submitted to J. Mat. Res.

5. Chandhari, P.; LeGoues, F. K.; Segmuller, A., Science, 1987. 1238 . 342.

6. Kramer, S.; Wu, Ki; Kordas, G., Mat. Res. Soc. Proc. 1988, 5L323.

7. Hirano, S.; Hayashi, T.; Miura, M.; Tomonaga, H. Presented at the 90th

Annual Meeting of the American Ceramic Society, May, 1988, Cincinnati, OH.

8. Jones, M. M., J. Am. Chem. Soc.. 1959. 81. 3188.

9. Bertrand, J. A.; Kaplan, R. I.; Inorg. Chem.. 1965. 4 (11). 1657.

10. Singh, J. v., Baranwal, B. P.; Mehrotra, R C. Z Anorg. Allg. Chem., 1981.

477. 235.

11. Andrew, J. E.; Blake, A. B.; J. Chem. Soc. Dalton, 1973. 1102.

12. Data were collected at -75°C. The structure was solved by direct methods

(SHELXS-86); correct positions for copper, oxygen and 12/18 carbon atoms were deduced from an E-map. Subsequent least squares-difference Fourier

calculations gave positions for the remaining carbon and hydrogen atoms. The

final difference Fourier map had no significant features. Final agreement factor,

R = 0.0033.

13. Nakamoto, K., Infrared Spectra of Inorganic and Coordination Compounds; J. Wiley and Sons, Inc., NY, 1978.

14. (a) Fahrenholtz, W. G.; Millar, D. M.; Payne, D. A., Ceramic Superconductors

III Research Update 1989, M. Y. Yan, Editor, 141, Am. Ceram. Soc.

(b) Fahrenholtz, W. G., M. S. Thesis, University of Illinois, 1989.

122
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

MOLECULAR CHEMISTRY AND THE SYNTHESIS OF PRECURSORS TO ELECTRONIC

CERAMIC MATERIALS

J. J. Ritter, J. F. Kelly, D. E. Newbury^, and D. B. Minor

NIST

Ceramics Division

Gaithersburg, MD 20899

ABSTRACT

The use of molecular chemical reactions to generate precursors to ceramic

materials has been the subject of considerable study for more than three

decades. This approach, in many cases, offers distinct advantages over conventional high- temperature syntheses particularly in the areas of

improved homogeneity, purity, particle size regulation and amenability to

intelligent processing control. Generally, within the broad spectrum of

chemical schemes available for ceramic precursor synthesis, the molecular

processes are poorly understood. Moreover, the reactions are often

interpreted either through "chemical intuition" or by analogy to similar

chemistries, rather than from measurements on the relevant systems. An

additional complicating factor arises from our lack of knowledge about the

mechanisms operating during the conversion of precursors to ceramic phases.

This presentation elaborates on several chemical routes which have been

used to synthesize ceramic precursors. The discussion includes systems

where control over the molecular architecture may lead to ceramic phases at

low temperatures. Other systems, which essentially deliver random,

intimately mixed assemblies of the elements, are also examined. The

various approaches are illustrated with examples from research programs at

NIST which have focused on the chemical synthesis of precursors to barium

polytitanates , metal carbides and borides, electrically conductive oxides

and nitrides, and the YBCO high- temperature superconductor.

INTRODUCTION

Molecular chemical reactions have been used for more than 30 years to

generate precursors to ceramic materials. In many cases, these approaches

offer advantages over conventional high- temperature syntheses in the areas

of homogeneity, purity, particle size regulation, and amenability to

intelligent processing control.

Synthetic methods reported in the literature include the hydrolysis of

metal alkoxides, the reductive dehalogenation of elemental halides (RDEH)

the formation and pyrolysis of organometal ion complexes, and the

coprecipitation of insoluble salts. Some of these methods e.g., alkoxide

hydrolysis and RDEH offer the potential for control of protostructural

development. The concept of protostructural development may be viewed as

Center for Analytical Chemistry, NIST

123
the placement of atoms in a precursor in nearly the same locations they

will have in the final ceramic phase. Alternatively, one can envision the

fabrication of a precursor structure whose favored decomposition pathway

gives the desired ceramic phase through a facile relocation of atoms. From

the molecular chemist's viewpoint, control for protostructural development

can be approached by altering the reaction mode of participating molecules

through changes in substituents , solvent systems, pH, etc.. Recent

publications by Phule and Risbud[l,2] describe work related to this concept

with the synthesis of barium titanates from the hydrolysis of modified Ba

and Ti alkoxides. Protostructural precursors generally will require only a little additional energy to promote their conversion. Other synthesis

methods such as coprecipitation and the decomposition of organometal

complexes usually provide only an intimately mixed, random assembly of

atoms which require relatively more additional energy to form the ceramic

material

.

The use of metal alkoxides to produce oxide ceramic powders has been

reviewed extensively in the literature [ 3 ] . NIST work on barium

polytitanates derived from alkoxide routes has also been published

previously [4 , 5 ] . This presentation will focus on RDEH as a route to

carbides and borides, organometal complex routes to a conductive ceramic

powder, and a copreciptation route to the YBCO superconductor.

SYNTHESIS OF PRECURSORS TO CARBIDES AND BORIDES BY REDUCTIVE DEHALOGENATION

Background : Electrical and electronic devices of the future will be

expected to operate reliably in a variety of increasingly adverse

environments. To meet these requirements, designers are seriously

considering nonoxide ceramic materials for certain applications. For

example, silicon carbide has been proposed as a microchip substrate, boron

carbides may have applications as high- temperature semiconductors, while

titanium nitride has been shown to perform well as an electrical contact

material. The development of molecular chemical routes to these and other

nonoxide ceramic materials continues to gain momentum. I shall describe

here the application of an inorganic redox chemistry, known to chemists for

more than ICQ years, for the synthesis of precursors to carbide and boride

powders. Schematically, for silicon carbide, the reaction may be

represented as follows:

n-heptane

SiCl^ + CCl^ + 8Na > "SiC" + 8NaCl (1)

130°C (precursor)

The precursor to boron carbide is synthesized from a n-heptane solution of

BCI3 and CCI4 , whereas precursors to TiC and TiB2 are generated from TiCl^

and CCli, or BCI3 , respectively. The reaction chemistry is general for any

combination of elemental halides which can be reduced by alkali metals and

have at least some solubility in the solvent system used.

Synthesis and Characterization : The details of the reaction procedures

have been given previously[6] . In brief, a dilute solution of the

elemental halides in n-heptane is added dropwise to a vigorously stirred

suspension of molten sodium in the same solvent, under vacuum. Potassium

124
can be substituted for Na, or if Na/K alloy is used, the reaction proceeds

at room temperature. Under these carefully controlled conditions, the

normally explosive reaction between elemental halides and alkali metals is

smooth and only mildly exothermic. However, it is recommended that only

personnel familiar with handling reactive materials attempt these

procedures and that maximum safety precautions be implemented. After

removal of the solvent by distillation, the alkali halide is separated from

the precursor by vacuum sublimation at -900 "C. Silicon carbide and boron

carbide precursors require further heating in vacuum at 1400, and 1900 °C,

respectively, to crystallize, while TiC and TiB2 crystallize at -900 "C,

during the removal of the alkali halide. The materials are identified by

x-ray powder diffraction (XRPD) and morphology determined by scanning

electron microscopy (SEM) . Surface areas are measured with N2 by the

standard B.E.T. method.

Results and Discussion : Characterization data are summarized in table I. In most cases, the precursors are pyrophoric and must be handled in vacuum

or under argon (< Ippm water and oxygen) in a glove box. The powders are recovered with moderate -to high- surface areas while some, particularly the

boron carbide, contain a few percent excess carbon. The use of Na/K with

concomitant lower reaction temperatures reduces the excess carbon problem.

The TiC produced by this process is readily converted to TiN by heating to

-900 °C in N2

While the dehalogenation reaction mechanism has not been extensively

investigated, it is thought to be of the Wurtz-Fittig nucleophilic

displacement type well known in organic chemistry [7] . It is proposed that

the initial interaction of an elemental halide, MX4 , with an alkali metal

forms a nuleophile, X3M' , and an alkali halide, e.g., NaX, as shown in

eq(2).

MX^ + 2Na > Na"^ 'MXg + NaX (2)

Due to the electronegativities of halogen substituents , the central atom,

M, in MX^-type molecules bears a partial positive charge. Thus , the

nucleophile initially forms a complex, (I), with the central atom of another

elemental halide molecule, e.g., M'X^ , and in so doing displaces a halogen

from this receptor molecule. The coupled M-M' product (II) and another

molecule of NaX are produced, as shown in eq(3).

XgM"- M'Xg > XgM-M'Xg + NaX (3)

Na"" : X

(I) (II)

Each halogen and M site on the coupled product (II) is subject to attack by

either Na or another nucleophile with the final result that all of the

halogens are removed and a three-dimensional M-M' matrix is developed.

This three-dimensional M-M'matrix is the ceramic precursor. To a first

approximation, the manner in which this precursor is built up can be

influenced by varying the substituents, X, on M and M' . Changes in the

substituents will affect the electron densities of the M and M' centers,

125
and thus, their susceptibility to nucleophilic attack. In practice,

control of protostructure development can be realized by changing halogens

using a mixed halogen system, introducing a precoupled moiety similar to

(II) as a reactant, or perhaps even substituting pseudohalogens . Another

approach would involve adding the alkali metal in carefully measured

increments which are designed to control the coupling reactions in a sequential manner.

Table I

Characterization data for boride and carbide powders

silicon carbide boron carbide titanium diboride

P C P C P C

color black grey black black black black

B.E.T Surface Area (m^/g) 65.5 26.4 215.6 29.5

morphology

(SEM) spongelike

whiskers &

crystallites

spongelike

agglomerated

irregular

particles

spongelike

irregular

particles

XRPD eunorphous cubic (8F) & hexagonal

(AH)

amorphous B^C, excess C amorphous TiB2

Residual

CI (vrt.%) 0.090 0.086

P = precursor C = crystallized powder

ORGANOMETAL COMPLEX SYNTHESIS OF A CONDUCTIVE OXIDE CERAMIC POWDER

Background : Current practice in the construction of cathodes for alkaline

fuel cells involves the coating of small polytetrafluoroethylene (PTFE)

particles with 2-5 nm gold particles at such loadings to assure Au-Au

contact for electrical conductivity. While fine gold particles are a good

electrocatalyst for the cathodic reduction reaction, they tend to sinter

when cell temperatures exceed -80 "C, with a concomitant loss in

electrocatalytic activity. A new generation of cathodes is under

consideration, wherein the PTFE will be replaced by a "conductive particle

lightly loaded with gold. The elimination of extensive Au-Au contact is

expected to greatly alleviate the loss of activity at elevated

temperatures. Alternatively, if the "conductive particles" show

electrocatalytic activity, they could be used directly without the Au

deposition. The requirements for these "conductive particles" include,

electrical conductivity in the 10'^ -10^ /ohm- cm range, 20-70 nm particle

size, resistance to chemical attack by 40% aqueous KOH at -100 "C, and

stability under cell reduction potentials in the range 0.4-1.2 V,

Reversible Hydrogen Electrode. Among the candidates considered for this

application is a lithium-doped lanthanum nickelate, La(Li)Ni03j^ .

^ Investigation suggested by Dr. Joseph Singer, NASA-Lewis Research

Center, Cleveland, OH.

126
The synthetic challenge was to effectively substitute 1 mol % Li for Ni

while maintaining the required particle size.

Synthesis and Characterization : The known solubility of many lithium

compounds in water and other common solvents suggested that coprecipitation

would be inappropriate to assure the quantitative inclusion of Li in the

product. Thus, an alternative organometal complex approach to prepare

initial samples of the nickelate was selected. Three options were

investigated, a citric acid-ethylene glycol route[8], a glycine complex

route[9] and a citric acid route[10]. In each case, aqueous nitrate

solutions of the metals with molar ratio La : Ni : Li = 1.0 : 0.99 : 0.01, were combined with 1 mol of complexing agent per mole of metal ion. Solvent

was removed thermally until polymerization occurred. The resultant

polymers were calcined up to 500 C, and the desired phase formed at 600-

650 °C in O2 . Phases were identified by XRPD, surface areas determined

with N2 on a commercial B.E.T. apparatus, Li analysis conducted by flame

emission spectroscopy (FES) , while powder morphology was examined by SEM.

Small portions of the powder were ultrasonicated in 2%(v/v) acetic acid, in an attempt to increase surface area.

Results and Discussion : All three methods generate the desired material

which contains, in all cases, a few percent of NiO. However,

experimentally, the simplest approach is the citrate route. Surface area measurements on citrate -derived powders by standard B.E.T. methods give

values typically in the 3-6 sq m/g range. Results from FES verified that

Li is present at the intended 1 mole % level, whereas examination of the

material by SEM shows a matrix of fused ~100 nm particles. After

ultrasonication in 2% acetic acid, only 20% of the powder is recovered.

However, this recovered material exhibits a surface area of ~15 sq m/g.

Examination by SEM indicates considerable erosion of the connective "necks"

between particles but general retention of the connective structure.

Lithium assay by FES indicates that more than 80% of the original Li is

retained.

Studies by Delmon and Droguest[ll] , suggest that the polymerization of

citric acid-metal complexes form a vitreous phase which contains the metal

ions in a random array. It is unlikely that the polymer formation can be

controlled to develop a "protostructure" to the ceramic phase, but its

chief advantage is in its ability to retain ions in a relatively

homogeneous dispersion. For the synthesis of the Li -doped lanthanum

nickelate, the citrate complex route has the demonstrated advantage of

incorporating Li quantitatively. The high Li content of the powder after

vigorous acetic acid treatment suggests that a major portion of this dopant

is part of the nickelate crystal structure.

SYNTHESIS OF A COPRECIPITATED HYDROXYCARBONATE PRECURSOR TO THE YBCO

SUPERCONDUCTOR

Background : Within the past few years, a large number of chemical

procedures for the synthesis of precursors to the YBCO superconductor have

been published. The coprecipitation approach to be considered here has

particular merit in that it is a water-based process, uses relatively

127
inexpensive chemicals, and is readily adapted to continuous production

methods with intelligent processing control.

In conventional practice, coprecipitation reactions are often conducted by

adding one reactant solution to another in a dropwise fashion with some attendant nominal stirring. This approach favors the formation of large

concentration gradients in the vicinity of the entering reactant solution

causing significant initial inhomogeneities in the formation of the

coprecipitate . Moreover, the ionic concentrations are subject to

continuous change as the addition proceeds, resulting in a continuous

change of the reaction rate. In addition, the ever increasing bulk of

coprecipitated material captive in the receiving vessel acts as an occlusion trap for unreacted species, thus both slowing down the

coprecipitation process, and increasing the probability for inhomogeneous

precipitation. The latter problems become particularly serious in the

final stages of addition when the viscous slurries become difficult to

stir.

Alternative methods of combining reactant solutions have been developed,

but these are infrequently used in the research laboratory. For example,

it is well known that a fluid moving under laminar flow conditions within a straight tube develops a parabolic flow velocity profile. Fluid near the

center of the tube travels at twice the average fluid velocity, while fluid

near the walls travels much more slowly. Samples of reaction fluid taken

along a laminar flow tube radius will have very different residence,

reaction and thermal histories. Thus, since a laminar flow tube provides

practically no radial mixing, it is a poor choice for a chemical reactor.

The flow profile , and thus the radial mixing can be improved somewhat by

increasing the Reynolds number of the system to >10'' so that the fluid

experiences turbulent flow.

However, within the past two decades, a technology has been developed in

the chemical process industry which provides efficient mixing even under

laminar flow conditions. This advance, often referred to as "motionless"

or "fixed-element" mixing, incorporates a series of carefully designed,

stationary, geometric elements within a tube of flow. The shape and

placement of these elements is such that they interact with the moving

fluid causing it to divide, swirl and recombine in a mathematically

predictable fashion to impart an effective radial mixing. The overall

effect is to give all of the flowing material nearly the same velocity,

residence time and degree of radial mixing. As a result, each increment of

fluid emerging from the flow tube has the same uniform mixing history. The

output of a well -designed "fixed- element" mixer very closely approximates

that of an ideal plug flow reactor (12)

The chemical flow reactor used at NIST is a commercially available research

model. It consists of a series of 52 helical elements, each of aspect

ratio 1.6, stacked within a flow tube with their respective ends oriented

at 90 "C to each other. Each element is placed so that its "twist,"

clockwise or counterclockwise, is opposite to that of its precedent. The

number of divisions experienced by an increment of fluid in its passage

through the reactor is given by 2" , where n is the number of elements in

the tube. For this reactor the value is 2^^. The two reactant solution

streams are combined in a region of confluence at the first helical element

128
in the tube. With proper adjustment of flow rates and reactant

concentrations, the efficiency of this system allows, for most

precipitation reactions, an effective blending of the solutions within the

induction period of the reaction. As a result, many discrete nuclei are formed and particles grow from these nuclei until the reactants within the

immediate vicinity are consumed. With suitable modifications to the

plumbing, the flow reactor is equally adaptable to either batch or continuous powder syntheses as shown in figures 1 and 4b, respectively.

Synthesis and Characterization : The reaction of a solution of Y, Ba, and

Cu acetates with solution of potassium carbonate can be written

schematically as follows:

^3 + + 2Ba'2 + + 3Cu2 + CO32-

HoO

[Y(0H)C03 :2BaC03 :2Cu(0H)2 :CuC03 ] . (4)

The details of conducting this coprecipitation in the conventional manner

and evidence for the composition of the hydroxycarbonate product shown in

eq(4) have been described previously [ 13]

.

A schematic diagram of the chemical flow reactor arranged for batch

coprecipitation is shown in figure 1.

Pi Figure 1. Chemical Flow Reactor Schematic

A reactant reservoir pj^ mainstream pump E helical mixing element

B slurry reservoir P2 sidestream pump c stream confluence area

C reactaint reservoir T mixing tube A water solution of K2CO3 with [COg^"' = 0.30 M is contained in reservoir A

and delivered to the reactor, T, by gear pump pj^ at a velocity of 56 cm/s

.

Simultaneously, a water solution of Y, Ba, and Cu acetates with [Y'^'^] =

129
0.05 M, [Ba2+] = 0.10 M and [Cu^*] = 0.15 M, is contained in reservoir C

and delivered to the reactor by gear pump P2 at a velocity of 110 cm/s

.

The two process streams are combined in an area of confluence, c, at the

leading edge of the first helical mixing element, E. The solution volumes

and flow rates from reservoirs A and C are such that they become empty at

nearly the same instant. At this point, the product slurry which has been

accumulating in reservoir B, has reached a sufficient height to initiate a self -siphoning action into reservoir A. Pump p^ now serves to circulate

the slurry through the reactor, while P2 is shut down. The recirculation

process is continued for some minutes before the slurry is removed. The

coprecipitated hydroxycarbonates are separated from the liquid phase by

centrifugation, and washed several times with water followed by several

ethanol washings.

The washed coprecipiates from both conventional and flow reactor routes

are initially air dried and finally calcined at temperatures between 300-

500 °C for various periods of time(cf. fig. 5). Samples of the as- precipitated and calcined materials were subjected, as appropriate, to

examination by SEM, Transmission Electron Microscopy (TEM) , and by High

Resolution Analytical Electron Microscopy (HRAEM)

.

Results and Discussion : The SEM micrograph shown in figure 2a shows that

the coprecipitated material generated at low, ionic concentrations is

highly segregated. Moreover, earlier work indicates that this segregation

is independent of pH in the range from 8 -11 [14] Energy Dispersive X-Ray

(EDX) analyses indicate that the spherical particles are rich in both Y and

Cu while the dendritic structures are rich in Ba. This data, when combined

with previous evidence available on the precipitation forms of Y, Ba, and

Cu, indicates that the spherical particles are a yttrium-copper

hydroxycarbonate , while the dendrites are barium carbonate [ 13 ]

.

Figure 2b shows material coprecipitated at higher ionic concentrations. No

morphological distinctions between particles are apparent and, to the

resolution of the SEIM EDX (~1 micrometer) , no elemental segregation is

detected

.

Figure 3 shows the morphology of Y, Ba, and Cu hydroxycarbonates as coprecipitated in the chemical flow reactor. Note that the individual

particles, although catenated, are 20-30 nm in size in contrast to the 100-

300 nm sized particles derived from the conventional precipitation route

and depicted in Figure 2b.

Figure 4 shows the results of HRAEM Y, Ba, and Cu mapping on conventional

and flow reactor coprecipitated powders, respectively. Since the

resolution of the HRAEM is of the order of 10-20 nm, different portions of

individual small particles may be assayed for composition. While EDX

mapping of powder conventionally coprecipitated at high ionic

concentrations shows no segregation, the higher resolution data from HRAEM

indicates substantial inhomogeneities in this material. In contrast, the

flow reactor powder exhibits a rather close grouping of the data points,

indicating a higher degree of homogeneity.

130
i (a) (b)

Figure 2. SEM micrographs of hydroxycarbonates conventionally

coprecipitated at low- and high- ionic concentrations

a. pH = 10, [Y-^"*"] = 0.0033M, [Ba^"^] = 0.0066M, [Cu^"*"] = 0.0099M, [COg^"] = 0.02M

b. pH - 10, [Y^"^] = 0.05M, [Ba^+l = O.IOM, [Cu^+J = 0.15M, [COg^"] = 0.30M

I

Figure 3. TEM micrograph of Y, Ba, and Cu hydroxycarbonates as coprecipitated in the chemical flow reactor.

131
Reactant A

CP

Reactant B

Coprecipltatad Produce

Static HlK«r Elements

Rtactant Jt

• •'• jlpioauet tloriy

ii

Mactant

(a) (b)

Figure 4. Apparatus schematics and respective Y/Cu and Ba/Cu

concentration ratios (wt %) as determined by HRAEM.

a. Conventional coprecipitation b. Flow reactor coprecipitation

Flow reactor configured for continuous powder

synthesis

A consideration of the precipitating forms of the respective metallic

elements in a highly basic carbonate environment leads to some interesting

speculations. As shovm in equation (4), both Y and Cu precipitate as hydroxycarbonates , for which no solubility data have been found in the

literature. Barium precipitates predominantly as barium carbonate. In the

light of a recent publication by Livage[15], it is suggested that Y and Cu

ions undergo "olation" processes, i.e., the linking of these ions through

the formation of hydroxyl bridges. Thus, Y and Cu readily coprecipitate

.

However, Ba ions, with their strong affinity for carbonate, are not

chemically incorporated into the Y-OH-Cu network. This differential

precipitation phenomenon is manifested macroscopically at low- ionic

concentrations as shown in figure 2a. At higher ionic concentrations, or under the very efficient mixing regime offered within the flow reactor, the same phenomenon persists but on a very different scale. Thus, one can

132
envision a Y-OH-Cu network encapsulating smaller and smaller clusters of

barium carbonate according to the relative rates of precipitation

prevailing under the ionic concentrations and the mixing regime used.

Figure 5 a-d summarizes the results of HRAEM studies on flow reactor

powders calcined under various conditions. The initial high degree of

homogeneity deteriorates as calcining times and temperatures increase,

suggesting that "flash" calcining procedures may be advantageous for this

system. However, recent in situ phase development studies by hightemperature x-ray techniques [ 16 ] have shown that even flow reactor powders

with a calcining history similar to that shown in figure 5d, convert

readily to tetragonal YBCO at low

800 °C.

O2 partial pressures between 600 and

Ba/Cu

Figure 5. Y/Cu and Ba/Cu concentration ratios (wt %) for as- precipitated and calcined flow reactor powders.

a. as coprecipitated; b. 315°C, 5 min; c. 500°C, 5min; d. A50°C, 16. 5h

SUMMARY

This presentation has indicated how molecular chemical routes to ceramic

precursors are of practical utility to obtain desired ceramic phases often

at low to modest temperatures. Methods such as RDEH and alkoxide

hydrolysis, which have the potential for protostructure control, are quite

complex and difficult to study. In the area of alkoxide hydrolysis, for

example, much of the work has focussed on silicon-based materials [ 17 ], with

very little emphasis on the more complex transition metal alkoxides or mixed-element alkoxides. As a result, progress towards "designing"

133
protostructures in multicomponent precursors has been minimal. Both the

RDEH reactions and alkoxide hydrolyses are thought to be of the

nucleophilic displacement type. The recent work of Livage(15) outlining

procedures for estimating partial charges on metal atoms in molecules as a function of substituents should be quite useful in predicting at least the

initial stages of these displacement reactions.

Organometal complexing has been shown to be useful in the synthesis of

lithium- doped lanthanum nickelate. Although the procedure does not produce

a high- surface area material directly, acetic acid treatment improves the

surface area by a factor of five and produces a textured structure. The

textured structure may offer subsequent benefits when the nickelate is

fabricated into electrodes.

Homogeneity in the coprecipitation of Y, Ba, and Cu as hydroxycarbonates

has been shown to be a function of the precipitation method, with the most

homogeneous material obtained by using the chemical flow reactor.

Unfortunately this homogeneity deteriorates during calcining, but our experimental results suggest that "flash" calcining procedures may minimize

the segregation.

ACKNOWLEDGEMENTS

Research on conductive oxides was supported by the NASA-Lewis Research

Center, under contract number C - 30010 -M. The authors would like to thank

NIST scientists Dr. Eric Steel for the TEM micrograph and Dr. M. Epstein

for the FES analyses.

REFERENCES

1. P.P. Phule and S.H. Risbud, Mater. Sci.6e Eng., B3,(1989), 241.

2. P.P. Phule and S.H. Risbud, Mat. Res. Symp. Proc, 121, Materials

Research Society, (1988), p. 275.

3. David W. Johnson Jr., Am. Cer. Bull., 64,(1985), 1597.

4. J.J. Ritter, R.S. Roth, and J.E. Blendell, J. Amer. Cer. Soc . , 69,

(1986), 155.

5. R.S. Roth, J.J. Ritter, H.S. Parker, and D.B. Minor, J. Amer. Cer.

Soc. , 69, (1986) , 858.

6. J.J. Ritter, Advances in Ceramics: Ceramic Powder Science, G.L.

Messing, K.S. Mazdiyasni, J.W. McCauley, and R.A. Haber, eds ., Amer.

Cer. Soc, Westerville, OH, (1987), p. 21.

7. E. Earl Royals, Advanced Organic Chemistry, Prentice Hall Inc.,

NJ,(1959), p. 117.

8. Paul A. Lessing, Ceramic Bulletin, 68, (1989), 1002.

134
9. L.A. Chick, J.L. Bates, L.R. Pederson, and H.E. Kissinger, Proc . l^*' Int'l Symp.; Solid Oxide Fuel Cells, Subhash C. Singhal, ed. , (1989),

p. 170.

10. H.-M. Zhang, Y. Teraoka, and N. Yamazoe , Chem. Soc, Japan, Chem.

Lett., (1987), 665.

11. B. Delmon and J. Droguest, Fine Particles; 2"*^ Int'l Conf . , Wm. E. Kuhn, ed. , Electrochem Soc. Monograph, (1974), p. 24.

12. Chemical Reaction Engineering, 2"*^ ed. , Octave Levenspiel, J. Wiley &

Sons, Inc., NY, (1972), p. 107.

13. J. J. Ritter and J. F. Kelly, High Temperature Superconducting

Compounds: Processing and Properties, S. H. Whang and A. DasGupta,

eds . , Minerals, Metals and Materials Society, Warrendale PA, (1989),

p. 303.

14. J. F. Kelly and J. J. Ritter, to be published.

15. J. Livage, Sol-Gel Science and Technology, M.A. Aegerter, M.J.

Jafelicci Jr., D.F. Souza and E.D. Zanotto, eds., World Scientific,

N.J., (1989), p. 111.

16. J. P. Cline, R.S. Roth, and J.J. Ritter, to be published.

17. R.C. Mehrotra, Sol-Gel Science and Technology, M.A. Aegerter, M.J.

Jafelicci Jr., D.F. Souza and E.D. Zanotto, eds.. World Scientific,

N.J. , (1989) , p. 40.

135
DISCUSSION

R. Roth : I would like to make some comments about some of this superconductor

precursor Joe was talking about. We've done quite a bit of high-temperature x-ray

characterization of some of his products. As any of you who've been playing with the

copper oxide, 123 (or 213) compound know, there are problems with respect to carbon

dioxide in this compound. Starting with a carbonate precursor you might think that's not

the way to go if you're worried about carbonating the product. But we have found that

if you calcine this precursor in a high temperature x-ray, at 700 degrees in pure oxygen

you get a beautiful join of yttrium oxide and 2:3 barium cuprate. We know from earlier

work that the 2:3 barium cuprate doesn't form in the presence of carbon dioxide at all. Consequently, with the proper procedures you can get rid of all of the carbonate in this

mixture when you heat it. So it's really, in my opinion, a good way to go. I wouldn't

have thought so too long ago. Joe, you agree?

J. Ritter: I would say one thing in general. While we can synthesize good powders

chemically, we have a very poor understanding how to process them after we make them,

and that's Bob's point.

B. A. Scott : I'm curious about the reduction chemistry to form silicon carbide. It would

seem that you're sort of at the mercy of kinetics of those reduction processes, so you get a lot of wrong bonds as well as the ones you want.

J. Ritter: Yes.

B. A. Scott : So, is that something that you can circumvent?

J. Ritter : Let me put it this way. We demonstrated that the chemistry works. We have

not demonstrated how to control it. I think there is potential, but the approach needs a

lot of additional work.

P. Davies : With regard to Bob's point, on the stability of the Ba2Cu305 phase. If your

goal is to make homogeneous 123 with as small a particle size as possible you really need

to avoid firing anywhere below 700 or 750 °C. It seems desirable to heat straight to 800-

850 °C as quickly as possible; otherwise, you're going to undo all your initial homogeneity, start to form the 2:3 phase and yttria with reasonable particle sizes, which

will then have to resinter.

R. Roth : Oh, no, Peter, that depends on the Po2- If you do this at .5% oxygen in the

nitrogen mixture then you get your 123 to form at 600 °C.

P. Davies : Without seeing any 2:3?

R. Roth : Without seeing that phase at all.

136
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

ION EXCHANGE REACTIONS OF LAYERED PHOSPHATOANTIMONIC ACIDS : A ROUTE FOR NEW CATALYSTS AND LUMINESCENT MATERIALS

Yves PIFFARD

Laboratoire de Chimie des Bolides,

IPCM UMR CNRS n" 110, Universite de Nantes

2, rue de la Houssiniere, 44072 NANTES CEDEX 03, France.

I - INTRODUCTION

It is now more than thirty years since the ion exchange behavior

of zirconiuim phosphate gels was first reported [1,2]. Since that

time many investigations into the nature of this and similar com- pounds have continued, leading to the discovery of many other

layered groups (IV) and (V) phosphates and arsenates. The most important results dealing with these studies have been summarized in

various review articles and books [3-7]. As already pointed out

[8,9] the layered structure of a-zirconium phosphate (a-ZrP) and

related materials exhibit a remarkable analogy with that of smec- tite clays, leading to similar chemical behaviors. These analogies

have motivated considerable research on crystalline forms of

a-ZrP-like materials and in the recent past, exciting new developments have taken place in the areas of catalysis, electrochemistry

and fast ion conduction [7].

Prior to our work in this domain several attempts to prepare phosphatoantimonic acids have been reported [10-14] , but there was only one material : HSb(P0^)2 .XH2O for which the Sb/P ratio was well defined and whose structure, even not completely determined,

was known since it is closely related to that of a-ZrP [15-16].

However it must be mentioned that, for this compound prepared from

SbCl^ and concentrated phosphoric acid, some disagreements exist

about its characterization [10,12,17] (water content, decomposition temperature, amorphous character). As part of a search

for new phosphatoantimonic acids [18] HSb(P0^)2 .XH2O has been prepared through a new route along with an other layered phosphatoantimonic acid (LPA) : H2Sb2P20-]^^.xH20. This route involves the synthesis by solid state reaction of potassium phosphatoantimonates

and ion exchange reactions in acidic medium. After a description

137
of the crystal structure of the parent potassium compounds, the

preparation and characterization (thermal behavior, water content)

of the LPA will be examined. Then, it will be shown through two

examples how new catalysts and luminescent materials can be prepared by taking advantage of the ion exchange properties of these

acids

.

II - PARENT POTASSIUM PHASES

II. I KSb(P0^)2

Perspective drawing of two (Sb(PO, ). ) layers

A 2 n

Fig. 1 showing two potassium atoms in the interlayers space

This compound

crystallizes in the

rhombohedral system,

space group R3 , with

a = 4.7623(4)A, c =

25.409(4)A, Z = 3 [19]

.

It has a layered

structure (fig 1). The

(Sb(PO^) "2)^^ infinite

layers are built up

from SbOg octahedra and

PO^ tetrahedra sharing

corners. Three oxygen

atoms of each

tetrahedral phosphate

are linked to three an- timony atoms so that

each antimony is

octahedrally coordinated with six oxygen

atoms of six different

phosphate groups . The

fourth oxygen of each phosphate group is unshared and points into

the interlayer space. These layers are very similar to the

(Zr(P0^)2 ')^^ layers in the a-ZrP compound. Each potassium atom is

octahedrally coordinated to six of the unshared oxygen atoms (fig

1).

II.KjSbjP20j^^.xH20

This compound crystallizes in the rhombohedral system, space group

R3m, with a= 7.147(1)A, c= 30.936(6)A, Z = 3 [20]. It is a layered material (fig 2 and 3). The (^^2^2^!^ ^ ^^Y^^^ built

up from SbOg octahedra and PO^ tetrahedra sharing vertices. The

SbOg octahedra are linked together in the same way as the WOg oc- tahedra in the (001) plane of the hexagonal tungsten bronze. The

phosphate groups are linked to these layers of octahedra via three

138 1
of their vertices. The fourth, which is unshared, points into the

interlayer space wherein potassium atoms are situated. In the unit

cell, there are nine potassium atoms distributed over twelve possible positions which are therefore occupied at an average of 75%.

This compound, prepared by solid state reaction at 1000°C, is hydrated at room temperature (RT) [21]

Fig. 2 - ^3^'^3^2^1A ' P^°j °^ structure on (110).

3- Fig. 3 - [001] view of a (Sb P 0 ) layer.

Ill - THE LPA Hj^Sb^P203^_^3.xH20 (n=l,3) : PREPARATION, WATER

CONTENT AND THERMAL BEHAVIOR

111.1 - Preparation

The potassium compounds have been ion exchanged in a 8N nitric

acid solution at 50°C [22,23]. The solution to solid ratio was lOOml/g of alkali material. After 3 stages, with renewal of the

acidic solution between each 2hr stage, the extent of exchange is

at least 99% and the hydrolysis has proved to be very weak, with

less than 1% by weight. Electron and X-ray powder diffraction studies show unambiguously that the covalent layers which have been

evidenced in the potassium phases are still present in the LPA.

111. 2 - Water content, interlayer distance and protonic conductivity versus relative humidity (RH) at 20°C.

The LPA are hydrated compounds and their water content at RT

strongly depends on the partial pressure of water vapor. The RH

139
influence on the composition, interlayer distance and protonic

conductivity has then been studied at 20°C [24]. The main results

of these studies have been summarized on fig. 4. Since the lattice

parameters parallel to the layers are very close to those of the

parent potassium compound, one can calculate the increase of vo- lume per formula unit of the interlayer space simply by considering the evolution of the interlayer distance. It is then possible

to estimate if the corresponding increase of the water content can

be associated, entirely or partly, to crystallization water.

0 50 %RH 3 0 50 %RH

Fig. 4 - Adsorption isotherm of water, evolution of the water content x, of the interlayer

distance d and of the conductivity, as a function of the relative humidity at 20 °C for

a) HSbCPO, )„.xH„0 b ) H,Sb,P„0. , . xH„0. A2 2 332 14 2 For HSb(P0^)2 .XH2O the adsorption isotherm (fig. 4a) shows a sigmoid- like evolution and the volume increase of the interlayer

space with RH cannot account for the total increase of the water

content ; then one must assume that part of it refers to water mo- lecules adsorbed on the grain boundaries or at the surface . The

conductivity strongly depends on the water content but its variation is too important to be only due to the intercalation of water. Generally speaking the evolution of the conductivity and the

behavior towards hydration recall those observed for particle hydrates [25].

For H2Sb2P20]^4 •XH2O the protonic conductivity is closely related

to the water content (fig. 4b). Each plateau corresponds to a fixed X value and to a particular interlayer distance. Therefore,

this evolution shows a true "bulk- type" conductivity in agreement

with the fact that, from volume considerations, the water content

can be associated with crystallization water over almost the whole

range of RH.

140
III. 3 - Thermal behavior

In hydrates, properties related with the water content are strongly dependent on the temperature. The thermal behavior of the LPA

has then been studied with use of TGA, DTA, DSC and X-ray thermodiffraction techniques [22,23]. Typical TG curves for both LPA are represented on fig. 5. They show that the LPA are rather stable : the dehydration process is reversible up to 560° C for

HSb(PO^)2.xH20 and 300°C for H3Sb3P20-|^^ . XH2O

reversibla

range

TGA

2SbP05 •V2aSbjO|

Fig. 5 - Weight loss versus temperature for a) HSb(PO,A22 )„ . xH„0 b) H.,Sb.,P„0, , .xH^O. 332 14 2

IV - ACIDIC AND ION EXCHANGE PROPERTIES OF THE LPA

IV . 1 - Acidic character

Vibrational spectroscopies, infrared absorption and Raman diffusion, are well adapted to the characterization of inorganic acids

and especially to infer valuable information about protonic species. It is, however, essential to identify first the various vibrations corresponding to the atoms of the covalent layers. The

LPA have then been studied at 20° C in comparison with the parent

potassium compounds for which an assigment of the observed frequencies has been done [26,27] and more particularly for the

stretching P-0 and Sb-0 domains.

For both acids, only slight changes are observed [28] in the vibrational spectra of the covalent layers with respect to those of

the parent potassium phases, and this is true, in particular, for

frequencies assigned to the antisymmetric stretching vibration involving the terminal P-0 bonds ; i.e those bonds which point into

the interlayer space and correspond to unshared oxygen atoms. It

is then clear that the interaction between acidic protons, as well

as all protonic species, and covalent layers is relatively weak.

Strongly self -associated HPO^^' groups similar to those observed

in Q-ZrP do not exist in the LPA in which acidic protons belong to

H^O^ oxonium groups ; that means that they are stronger acids than

Q-ZrP and related compounds.

141
In addition the LPA exhibit

a swelling behavior similar

to that of smectite clays.

When suspended in water

they readily undergo infinite swelling leading to a sol- or a gel-like solution

from which a film can be

obtained after removal of

the solvent. This is a re- markable difference with

the a-ZrP like acids which

require a preintercalation

of amines to give colloidal

suspensions. Such a difference can be attribued

to the fact that the charge

of the anionic layer in

Q-ZrP is higher than that

of layers in the LPA or in

montmorillonite [8]

.

IV. 2 - Ion exchange

behavior

The ion exchange behavior

of the LPA is illustrated

by the titrations with

alkali hydroxide solutions

(fig. 6) [23,24] and also

by ion exchange isotherms such as those represented in fig. 7 for

H3Sb3P20-|^^.xH20 (isotherms with HSb(P0^)2 .XH2O are similar).

n«qOH7e of HjSbjPjO^ 2 1 MjO

Fig. 6 - Potentiometric titration curves for alkali

cations on a) HSb(PO, )^ .xH^O b) H^Sb.,P„0, , .xH„0.

2 332 14 2

Fig. 7 - a) Cs*/H^ b) Ag^/H^ ion exchange isotherm on H,Sb,F„0, , .xH„0.

3 3 2 l»i 2

142
The ordinate of these plots, X^, is the equivalent fraction of the

M"*" ion in the exchanger and the abscissa, X^, is the equivalent

fraction of the ion in the solution. Both curves show that the

selectivity coefficients of the Cs"'"/H''" and kg^/E^ systems with the

LPA are very high.

V - THE LPA AND SOME OF THEIR ION EXCHANGE DERIVATIVES AS ACIDIC

CATALYSTS

.

In order to evaluate their catalytic properties, the LPA and some of their ion exchange derivatives have been studied in a test

reaction which corresponds to the conversion of isopropanol (IPA)

into propene , diisopropylether and acetone according to the following reaction : CH^ CH = CH^ + H^O

CH^ - CH - CH^ > I (CH2)2CH - 0 - CEiCR^) ^ + U^O

I OH

CH^ - C - CH^ + H^

0

This reaction is not of industrial interest ; however since it can give rise either to dehydrations (intra- and/or inter -molecular) or to a dehydrogenation, it appears as a valuable catalytic

test

.

Cu''""'"-, Ni''"'^- and La''"''"''"- derivatives with various metal-ion loadings have been prepared by ion exchanging the LPA with solutions

containing either Cu(SO^) . 5H2O , Cu(N03) 2 . 3H2O , Ni(N03) 2 . 6H2O or La(N03)3.6H20.

The best results have been obtained with HSb(P0^)2 .XH2O and its

ion exchange derivatives. In all cases, compounds resulting from

the dehydration (propene and diisopropylether) are the majority

products, indeed even exclusive with a clear selectivity towards

propene. This is not surprising since acidic sites from the catalyst promote the carbocation formation occuring in the dehydration

process as follows : CH3- CH-CH > CH2-CH-CH3 ^CH^-CH-CH^ >CH2-CH=CH2

I +H"''(catalyst) | + OH ^0^ +H 0 -H

H H

The selectivity towards propene is better when ion exchange derivatives are used and the carbon yield is close to 1. This result can be explained by the fact that when metallic ions are introduced, the density of acidic sites at the surface of the catalyst is

decreased, therefore :

143
desorption of products is promoted and, as a consequence, their

polymerization (coke formation) is avoided.

a monomolecular dehydration process is favored with respect to a bimolecular one since the distance between acidic sites is sufficiently large.

Disactivation curves obtained with some

of our catalysts can be interpreted in a similar way. They represent the propene

yield (PY) as a function of IPA

conversion yield (CY) (fig. 8). The ideal curve (diagonal straight line with slope

1) corresponds to a situation in which

all the IPA is converted in propene. The

fact that the experimental curve is a straight line indicates that propene- and

coke -format ion occur on the same sites

and that the disactivation corresponds to

a decrease on the number of active sites.

When ion exchange derivatives are used

the disactivation process is very slow,

indeed even non observable, as in the case of Cuq •j5Sb(PO^)2 which has been

more extensively studied. For this catalyst the IPA CY has been

represented in fig. 9 as a function of time. Starting at about 95%

the CY first decreases down to 75% and then increases up to 100%

(after 120 hr) for a reaction temperature (Tj^) of 200°C.

Fig. 8 - Disactivation curve

(Propene Yield as a function of

the IPA Conversion Yield) for

conversion

locr

50

Tp = 200 C Tp =150 C

Tr=200(:

Tp =120X:

200 400 time (hr)

Fig. 9 - IPA Conversion Yield versus time with Cu^^ 35^''^^A^2'

This step allowing the catalyst to reach is working regime is

observed only if the catalyst is prepared from Cu(SO^).5H20 (not

with Cu(N02)2 • 3H2O) . This effect of residual SO^^' ions is not

understood but it is clear that it promotes the activity of acidic

144
sites since the specific activity, for the same copper content, is

three times as much as that of a catalyst prepared with

Cu(N03)3 . 3H2O (Table I). In the working regime at 200°C the CY

remains very stable (close to 100%). If Tj^ is decreased, the CY

decreases (fig. 10, see steps at Tj^=150°C and Tj^=120°C) . However,

coming back to 1^^=200° C a CY of about 100% is restored and

maintained throughout the test, i.e over more than 500 hr.

TABLE I - Influence of the Cu'^^ content (Cu H, „ Sb(PO, )„) y l-2y A 2 on the IPA conversion

y 0 0.35

•ie-Je 0.35 0.10** 0.01

Specific activity

(g IPA g-1 h"^)

0.287 1.24 0.432 0.619 2.18

Selectivity (C%)

Propene

Diisopropylether

Acetone

91

9

100 100 100 99

1

* prepared from CuSO^ . 5H2O

** prepared from Cu(N02)2-6H20

Conditions of the test - T„Total pressure

IPA mass flow

- W.H.S.V.

200°C

1.4 bar

0.134 g h

1.25 h''^

-1

Various tests involving different copper contents and other

ion exchange derivatives of HSb(P0^)2 .XH2O have been performed.

The main results have been summarized in tables I and II. It can

be seen from table I that even for small copper contents

(y = 0.01), the ion exchange derivatives exhibit a greater

specific activity and a better selectivity than those of the

acidic compound. It must be mentioned however that, for a given

selectivity, the evaluation of a catalyst is not limited to its

specific activity, but it includes the disactivation rate. If

traces of copper promote the activity, larger quantities are necessary to maintain it for a long time. Table II shows the

influence of the cation type on the catalytic reaction for a given

substitution ratio. The specific activities are similar but the

selectivities are different. In addition it must be mentioned that

the Ni- and La- catalysts disactivate rapidly (a few tens hr)

145
TABLE II - Influence of M (M^ H, Sb(PO, )„ with yz = 0.7) on the conversion

y 1-yz 2 (conditions for the test similar to those given in table I).

un ^^0.35 N^O.35 T o ^^0.23

Specific activity

(g IPA h"^)

0.287 0.432 0.435 0.426

Selectivity (C%)

Propene

Diisopropylether

Acetone

91

9

100 97.3

2.7

99.6

0.4

•OOa \M7

VI - THE LPA AS STARTING COMPOUNDS FOR

THE PREPARATION OF LUMINESCENT

MATERIALS

When ion exchange reactions with

HSb(P0^)2 ,xH20 are carried out with

Ln ^ cations (Ln=Nd,Eu) the extent of

exchange is at least 98% with solutions

containing an excess of these cations

[22]. When the amount of Ln^"*^ in the

solution is less than the theoretical

exchange capacity of the acid, the

cation uptake is maximum and leads to

(LnyH-|^ _ 3

) Sb ( PO^ ) 2 • x ' H2O compounds . Their thermal decomposition occurs

above 850° C leading, to the isomorphous

series ^^/2^^^^ {l-^) /l^^"^3/2^'^^i,h

whose structure is isotypic with that

of Bi-'--'--'-3^/2Sb^3/2(^°4)3 l^^'^ • '^^ ^^^^ latter compound the shortest Bi-Bi

distance is 8.5A. This feature is of

great interest since it corresponds to

the shortest Nd-Nd distance in

compounds for which the Nd concentration can reach 2.2x10^ cm"^ when x=l

.

Fig. 10 - Emission spectra of the

Nd .-Sb'^-'--^,, N,oSb\,o(PO, )- compounds (1) x=0.05, x/2 (l-x)/2 3/2 Ad

(2) x=0.10, (3) x=0.15, («) x=0.25, (5) x=0.30

4 4 corresponding to the transitions a) F, -> I. b) V

3/2 11/2

3/2 9/2

146
As a matter of fact, it has been shown that the degree of

concentration quenching in Nd- laser materials is determined by the

linkage between the NdO^ polyhedra [30], and especially by the

Nd-Nd distance.

Preliminary luminescence experiments have been undertaken on powder samples of (1-x) /2^^^3/2^^^^^ 3 compounds at RT.

Typical emission spectra, corresponding to the ^^2/2 ^^11/2

and ^2/2 ^^9/2 transitions of Nd are given in fig. 10. It can be seen that che Nd concentration strongly influences the intensity of the emission lines which reaches a maximum for x=0.25

thus giving a rather narrow band. According to this result, single

crystal corresponding to this particular composition (x = 0.25)

would be adequate for a mini laser application.

The LPA prepared by ion exchange in acidic medium are rather well

defined compounds in terms of structure, water content and thermal

behavior. Their swelling behavior similar to that of smectite

clays and their strong acidity make them efficient ion exchangers

from which new materials can be prepared. Those include catalysts

and luminophors but other ones such as pillared layered materials

could be thought up

.

REFERENCES

[1] K.A. KRAUS and H.O. PHILIPS, J. Am. Chem. Soc . , 1956, 78, p.

644

[2] C.B. AMPHLETT, L.A. DONALD and M.J. REDMAN, Chem. In.

(London), 1964, p. 1314

[3] V. VESELY and V. PEKAREK, Talanta, 1972, 19, p. 219

[4] A. CLEARFIELD, G.H. NANCOLLAS and R.H. BLESSING,

J. A. MARINSKY and Y. MARCUS Eds., M. Dekker, New York, 1973,

Chap . 1

[5] G. ALBERTI and U. COSTANTINO, J. Chromatogr., 1974, 102, p. 5 [6] G. ALBERTI, Acc . Chem. Res., 1978, 11, p. 163

[7] A. CLEARFIELD, C.R.C. Press Inc , Boca Raton, 1982

[8] A.DYER and D.LEIGH, J. Inorg. Nucl. Chem., 1972, 34, p. 369

[9] S.YAMANAKA and M. KOIZUMI, Clays Clay Min. , 1975, 23, p. 477

[10] A. WINKLER and E. THILO, Z. Anorg. Allg. Chem., 1966, 346, p.

92

[11] A. WINKLER and E. THILO, J. Polymer. Sci., 1969, 16, p. 4283

[12] J. P. JOLIVET and J. LEFEBVRE, Bull. Soc. Chim. , 1975, p. 2409

[13] J. P. JOLIVET and J. LEFEBVRE, Bull. Soc. Chim., 1975, p. 2420

[14] N.G. CHERNORUKOV, I. A. KORSHUNOV, N.P. EGOROV, A.I. ZABELIN

and T.A. GALANOVA, Izv. Akad. Nauk S.S.S.R., Neorg. Mater.,

1981, 17, p. 1058.

147
J.M. TROUP and A. CLEARFIELD, Inorg. Chem. , 1977, 16, p. 3311

J. ALBERTSON, A. OSKARSSON, R. TELLGREN and J.O. THOMAS,

J. Phys. Chem., 1977, 81, p. 1574

N. MIURA, Y.OZAWA and N.YAMAZOE, Denki Kagaku, 1985, 53, p.

338

M. TOURNOUX and Y. PIFFARD, French Patent 85-01839, 1985

Y. PIFFARD, S. OYETOLA, S. COURANT and A.LACHGAR, J. Solid

State Chem., 1985, 60, p. 209

Y. PIFFARD, A. LACHGAR and M. TOURNOUX, J. Solid State Chem.,

1985, 58, p. 253.

A. LACHGAR, S. DENIARD- COURANT and Y. PIFFARD, J. Solid State

Chem. , 1988, 73, p. 572

Y. PIFFARD, A. VERBAERE, A. LACHGAR, S. DENIARD -COURANT and

M. TOURNOUX, Eur. J. Solid State Inorg. Chem., 1989, 26, p.

113

Y. PIFFARD, A. VERBAERE, A. LACHGAR, S. DENIARD - COURANT and M.

TOURNOUX, Rev. Chim. Min. , 1986, 23, p. 766.

S. DENIARD -COURANT, Y. PIFFARD, P. BARBOUX and J. LIVAGE , Solid State Ionics, 1988, 27, p. 189

P. BARBOUX, R. MORINEAU and J. LIVAGE, Solid State Ionics,

1988, 27, p. 211

E. HUSSON, A. LACHGAR and Y. PIFFARD, J. Solid State Chem.,

1988, 74, p. 138

E. HUSSON, F. GENET, A. LACHGAR and Y. PIFFARD, J. Solid State

Chem. , 1988, 75, p. 305

E. HUSSON, M. DURAND-LE FLOCH, C. DOREMIEUX-MORIN , S. DENIARD-COURANT and Y. PIFFARD, Solid State Ionics, 1989, 35,

p. 133

S. OYETOLA, A. VERBAERE, D. GUYOMARD and Y. PIFFARD, J. Solid

State Chem., 1988, 77, p. 102

M.S. HWANG, H. Y.-P. HONG, M.C. CHENG, and Y. WANG, Acta

Crystallogr. Sect. C, 1987, 43, p. 1241.

148
DISCUSSION

P. Battle: Have you looked at the protonic conductivity at higher temperature, in the

reversible regime?

Y. Piffard : No, we didn't study the protonic conductivity at higher temperature. We

only looked at the conductivity at room temperature because I think that since it is very

likely that the protonic conductivity is related to a diffusion of protons on a hydrogen

bonded network within the interlayer space, if you remove water molecules you will

lower the conductivity very, very quickly. It's almost always true for this type of

mechanism.

k. Poeppelmeier: Following up on your earlier point where you convinced us that the

proton was not tightly bound to the terminal oxygen but associated with the water

between the layer. Through your dehydration studies it appeared you could reach a state where you have only HSb(P04)2.

Y. Piffard : Yes.

K. Poeppelmeier: If you do the infrared analyses on that compound is the proton then

tightly bound to the terminal oxygen?

Y. Piffard : In that case, the situation is probably different, yes.

K. Poeppelmeier: This is the sort of material that probably is important at elevated

ternperatures for reactions when you're studying solid acids.

Y. Piffard : At higher temperatures, yes.

R. Roth : When you're ion exchanging those layered structures with Rb and Cs, I would

assume the Rb and Cs ions go in place of the water molecules structurally. Do these

compounds then hydrate as much as the potassium phase (K3Sb3P20j4-5H20) does?

Y. Piffard : I don't remember exactly what is the number of water molecules, but I know

that it is smaller.

R. Roth : They still do hydrate?

Y. Piffard : Yes, but the water content is less.

A. Jacobson : You mentioned that, in the swelling behavior with water, they looked like,

and behaved like sfnectite clays?

Y. Piffard: Yes.

149
A. Jacobson : Have you looked at ion exchange reactions with large cations like

polyhydroxycations, for example, to see if you can pillar them?

Y. Piffard : We are looking at this problem at the moment. I can indicate that it's very

easy to ion exchange with amines, for example, even with long chain amines. We don't

need any preintercalation processes as in the case of a-ZrP-like acids.

K. Poeppelmeier: I have a general question; maybe for the audience, too. This

difference between sulfate and nitrate anions. One of the things we've been looking at

is something we've been calling salt inhibition and I was wondering if the geochemists

in the audience have observed this phenomenon. In these fairly open structures there are

sites for both anions and cations, and literally both are incorporated. It has been

observed in zeolites. I was wondering if this phenomenon wouldn't be enhanced on the near surface when you do these partial exchanges with copper and observe the effect of

the anion. That in fact, there is incorporation of both cations and anions and the sulfate

anion would be different from the nitrate anion. This isn't ion exchange but the

incorporation of a cation and corresponding anion. Do you have any thoughts about that?

Have others seen effects like this?

Y. Piffard : I don't think so because when a cation exchange reaction is concerned you

don't think, a priori, of any difference associated with the anion.

K. Poeppelmeier: But it would be very difficult, I think, to detect a small additional

incorporation of both cations and anions. It isn't really equivalent to the ion exchange

process. It might be one explanation for the difference in the effect on these anions.

A. Navrotsky : To address that from the geochemical point of view there are a number

of minerals that incorporate cations and anions in different amounts and certainly

cordierites and scapulites and so on will incorporate nitrate, sulfate and carbonate ions.

I think very little is known about the mechanisms of whether they're coupled, but I think

you're probably right. It's probably kinetically controlled.

K. Poeppelmeier: It may be an important reaction to think about when you are studying

these reactions.

150
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

MIXED METAL OXIDES WITH THE PYROCHLORE STRUCTURE AS

CATALYSTS FOR METHANE OXIDATIVE COUPLING

A. J. Jacobson, J. T. Lewandowski, C. A. Mims*, R. B. Hall and G. R. Myers

Exxon Research and Engineering Company, Annandale, NJ 08801

*Department of Chemical Engineering, University of Toronto, Toronto, M5S 1A4, Canada

ABSTRACT

A new series of bismuth tin pyrochlore oxides have been synthesized by

precipitation from aqueous potassium hydroxide and calcination at 400°C. The low

temperature synthesis method leads to the formation of compounds with compositions

Bi2(Sn2-xBix)07-x/2 (0.05<x<0.86) in which bismuth is present on both the A and B sites

of the pyrochlore structure. The bismuth rich pyrochlores decompose at temperatures

between 500°C and 700°C, depending on the specific composition, to Bi203 and

Bi2Sn207. The phase separated compositions are good catalysts for the direct conversion

of methane to higher hydrocarbons via oxidative coupling.

INTRODUCTION

A variety of metal oxides have been studied as catalysts for the direct conversion of

methane to higher hydrocarbons via oxidative coupling[ll. Following the initial work of

Keller and Bhasin[21, many different types of oxides have been found to be effective

catalysts at temperatures above 600°C. Examples include simple binary systems such as the alkaline and rare earth oxides, PbO, Sn02 and Bi203, often promoted with alkali metals

to improve selectivities[2 - 8]. In particular, lithium promoted MgO[5, 6], samarium

oxide[7, 8] and supported manganese oxides[9] have been examined in detail. More

complex mixed metal oxides, for example, LiNiO2[10, 1 1], the perovskites BaPb03[12]

and SrCe03[13], kentrolite Pb2Mn2Si209/Si02[14] and bismuth oxyhalides ABi304Cl2 (A = Li, Na, K and Ag)[15, 16] all show good selectivities (40 - 80%) at moderate

conversions.

Although considerable variations in specific activity and selectivity are observed

between different oxides, the reaction mechanism is thought to be similar for all catalysts.

The hydrocarbon products are formed by coupling of methyl radicals which are generated

by hydrogen abstraction from methane by an active oxygen site on the catalyst surface.

151
The active oxygen site is regenerated by oxidation and the elimination of water from the

oxide surface. In order to study systematically the influence of chemical substitution on the

chemistry of these reactions, we have studied mixed metal oxides of general composition

A2B2O7 with the pyrochlore structure which can be synthesized for a wide range of A,B

combinations provided the radius ratio, t/Jt^, is greater than approximately 1.4[17]. The

catalytic performance of the pyrochlore oxides, Ln2Sn207[18] has recently been found to show variations which were related to the ease of oxidation or reduction of the specific rare earth cation. In this paper, we focus on one specific set of pyrochlore oxides based on the

compound Bi2Sn207 and describe the synthesis and stability of a range of compositions

together with some preliminary data on their performance as catalysts for oxidative

coupling of methane.

The compound Bi2Sn207 was first reported by Roth[19] and subsequently further

studied by Gattow[20], Shannon et al. [21]and Vetter[22]. The compound has a tetragonal

structure related to pyrochlore at room temperature with cell constants 21.328 and 21.545

A. At 140°C the a form converts to the cubic p structure with a unit cell parameter of

2I.40A at 200°C. Both the a and p form have unknown non-centric crystal structures. At

680°C the p structure transforms to the y phase with the ideal pyrochlore structure and a cell parameter of 10.73A at 700°C. The phase transitions observed in Bi2Sn207 are sensitive to impurities or small amounts of dopants on either the A or B sites in the pyrochlore

structure. Impurities and dopants lower the transition temperatures and for many samples

the a to p transition is not observed. The high temperature phase diagram shows the

existence of the single stoichiometric composition Bi2Sn207.

High surface area compounds with the pyrochlore structure containing bismuth or

lead as the A site cation and ruthenium or iridium as the B site cation have been prepared

previously by crystallization at mild conditions (<100°C) from concentrated alkaU

hydroxide solutions [23, 24]. Compounds were prepared with Bi3+ occupying both the A

and B cation sites in the pyrochlore structure by control of the synthesis conditions

(solution composition, pH and temperature). The Bi2Ru2-xBix06.5 compositions have

been called 'expanded' pyrochlores because their cubic lattice constants increase Hnearly

with composition with increasing levels of substitution on the B cation site. A similar

series of phases is observed when bismuth is replaced by lead. We have extended this

solution synthesis approach to prepare a new series of bismuth tin expanded pyrochlore

compositions Bi2(Sn2-xBix)07-x/2 and examined their catalytic behavior. The end member

of the series, Bi2Sn207, has previously been found to be an average catalyst for methane

coupling[25]. Mixed Bi203 - Sn02 phases have also been studied at lower temperatures

(500°C) as catalysts for methane conversion using N2O as the oxidant and observed to be

selective for the formation of formaldehyde at low (<3%) conversion [26].

EXPERIMENTAL

In a typical synthesis, Bi(N03)3 5H2O (12.129g, Alfa) was dissolved in 100ml of

water and 30ml of concentrated niuic acid. Tin was added to the solution as SnCLj • 5H2O

152
(8.763g, Aldrich) followed by 250ml of 2M KOH solution. The pH and the total solution

volume were then adjusted by further addition of KOH solution and water. A white

amorphous precipitate was formed at this stage of the synthesis. The slurry was then

heated at 75 to 100°C for typically 2 to 5 days to complete the reaction. The fmal product

was separated by filtration, washed with water and dried at 100°C. The effects of changes

in the total reactant concentrations from 0.03 13M to 0.0625M in tin, the solution pH values

from 10 to 14 and Bi/Sn ratios from 1:1 to 3:1 were examined.

A Dupont 1090 thermal analyzer was used to determine the water contents of

samples by heating them in air to 1000°C at 10°C min"^. Elemental analysis for alkali

metals, bismuth and tin was carried out by Galbraith Laboratories. Surface area measurements were made by N2 adsorption using an Omicron Omnisorb 360.

Powder X-ray diffraction data were measured on an automated Siemens D500

diffractometer using Cu Ka radiation. Samples were loaded into 2mm deep aluminum

sample trays. Diffraction profiles were scanned in o) - 29 mode over the range 2.0 < 29 <

90.0° in steps of 0.02°, with count times of 2s at each point. In a few cases, data over a wider angular range (2.0<29<122.0) with higher statistical precision were obtained at count

times of 30s per 0.02° step (see Figure 3.). Lattice constants were obtained from the

powder X-ray diffraction data by least squares refinement of peak positions or for many of

the samples by full matrix Rietveld analysis of the diffraction profiles [27].

The full matrix least squares Rietveld refinements were achieved using a locaUy

modified version of the DBW3.2 code of Wiles and Young[28]. The background was

treated by linear interpolation between a set of estimated points that were updated

periodically. The peak shapes were described by the pseudo-Voigt function. Scattering

factors for the species Bi3+ , Sn4+ and O^^ were used. The space group Fd3m(No. 227

[29]) was assumed.

The thermal stabiHty of the expanded pyrochlore phases with different compositions was determined X-ray diffraction measurements both 1) at room temperature after

annealing to each of a series of temperatures for 16h, and also 2) at high temperatures. The

high temperature X-ray measurements were made on samples contained in quartz capillaries

using a wavelength of 0.97203A at the XIOA beam line at Brookhaven National

Laboratories

.

The performance of the materials as catalysts for methane oxidative coupling was

evaluated using a small downflow quartz reactor. Samples of the catalysts consisting of

80-120 mesh particles were prepared by pressing into pellets, followed by grinding and

sieving. A standard reactant gas mixture of CH4 : O2 : Ar = 2 : 1 : 5 (stoichiometric for

ethylene production) at one atmosphere total pressure was used. By variations in in the

reactant flow rate (50-150 ml/min) and sample size (0.3-3g), we were able to vary the gas

residence time , and therefore the methane conversion at constant temperature. Reaction

temperatures from 500°C to 900°C were investigated and the products were analyzed by gas

chromatography.

153
RESULTS

The final compositions of the bismuth tin pyrochlore phases when prepared by the

solution synthesis route depend on the pH of the syntheses, the reaction temperature, the

ratios of the reactants and the total concentration of bismuth and tin. The formation of a pure pyrochlore phase directly from solution at 100°C depends critically on the pH of the

synthesis solution. For example, at a ratio of reactants Bi/Sn =1 and a total solution

concentration 0.0625 molar in tin, a pure pyrochlore phase is obtained only in a narrow pH

range close to 1 3. At high pH values, a mixture of the pyrochlore phase and a second

crystalUne composition is obtained. At these conditions, the second phase (D*) is obtained

in pure form at pH >13.3. Close to pH = 14, Bi203 appears as a second phase together

with D*. As the concentration of the solution is decreased, D* persists to lower pH as a pure phase and coexists with the pyrochlore to a pH as low as 12.65 in the most dilute

reaction solutions (0.0313 molar). Below pH = 12.5, a new poorly crystalline phase is formed in addition to pyrochlore. On calcination, materials containing this poorly

crystalline phase convert to a mixture of a pyrochlore and an amorphous component.

The average composition of D* determined for four different samples by elemental

and thermogravimetric analysis corresponds to a Bi/Sn ratio of 2.5. Thermogravimetric

analysis data show the absence of coordinated water molecules but show a small (2.8 - 3.1%) reproducible weight loss at 300°C suggesting the presence of hydroxyls. On

complete dehydroxylation the product is a pure pyrochlore phase. The weight loss and the

chemical analysis data indicate a composition Bi5Sn209(OH)5. All the reflections in the Xray powder pattern of D* can be indexed with an orthorhombic unit cell and cell constants a = 7.613(2)A, b = 10.564(3)A and c = 6.734(2)A. The systematic absences in the powder

data suggest that the space group is P2i2i2i(No. 19[291). Further work to obtain crystals

of D* large enough for structural analysis is in progress.

On calcinarion at 400°C, pure pyrochlore phases are observed for all of the products

of the 100°C solution syntheses which are pyrochlore, D* or mixtures. The surface areas of the pure pyrochlore phases as prepared from solution are typically in the range of 50 to 100 m^/g and reduce to 30 to 50 m^/g on calcination at 400°C. In contrast, D* phases and

pyrochlore samples made from them have surface areas in the 1-6 m^/g range.

The extent of bismuth substitution in the pyrochlore phases Bi2(Sn2-xBix)07.x/2 can

be varied by adjusting the various synthesis parameters such as the pH and the

concentration of the reactants. The detailed influence of all of the various synthesis

parameters has not yet been fully evaluated though some of the general trends are clear. For example, Figure 1 shows the influence of the ratio and concentration of the bismuth

and tin components on the product composition. The solution volumes of 400, 600, and

800 ml indicated on the Figure, correspond to solution concentrations which are 0.0625,

0.0417 and 0.0313 molar in tin. The composition of the product becomes more bismuth

rich both as the reactant Bi/Sn ratio increases and as the synthesis solutions become more

dilute. The latter effect is due to the greater solubility of tin relative to bismuth in alkaline

solutions.

154
0.80 1.0 1.2 1.4 1.6 1.8 2.0 2.2

Bi/Sn Reactant Ratio

Figure 1. The dependence of the Bi/Sn ratio in Bi2(Sn2-xBix)07-x/2 on the ratio of the reactants and their concentrations used in the synthesis.

The analytical data for samples which after calcination in air at 400°C appear to be

pure pyrochlore phases by X - ray diffraction fall in the composition range 0.05<x<0.86.

The products also usually contain small amounts of residual potassium which corresponds

to 0.01 to 0.05 K/Bi. The excess potassium is presumably adsorbed on the surface of the

high surface area particles since it can be easily removed by further washing with dilute

nitric acid. The X ray diffraction data for all of the compositions show considerable line

broadening consistent with particle sizes of approximately 80 -lOOA and the measured

surface areas. Lattice parameters were determined from the positions of the strong fluorite

subcell peak positions or by fitting the entire profile. The lattice parameters as a function of

composition are shown in Figure 2. The scatter in the data is primarily due to uncertainties

in determining lattice parameters from broad diffraction peaks and inhomogeneities in

sample composition. Nevertheless, the data show a well defined trend of increasing lattice parameter with increasing bismuth content consistent with the substitution of the larger

Bi3+ cation for the smaller Sn"*"*" cation on the octahedral B site of the pyrochlore structure. A similar trend is observed for the ruthenium containing pyrochlores[23, 241.

The X-ray powder diffraction pattern for a sample calcined at 400°C with x = 0.64

is shown in Figure 3. The data were fit by Rietveld analysis to the simplest cubic

pyrochlore structure with space group Fd3m though little evidence is observed in the data

for superlattice reflections corresponding to the doubling of the fluorite sub-cell. Close

155
Figure 2. Lattice Parameter vs. Composition for Bi2(Sn2-xBix)07-x/2-

inspection of the data, however, does show the presence of weaJc reflections resulting from

the ordering of bismuth and tin cations on the A and B sites of the pyrochlore structure. Similar results are obtained across the entire composition range. The data for the sample

with X = 0.64 were refined assuming that all of the excess bismuth substitutes for tin on the

B sites and that the cation charge deficiency is compensated by oxygen atom vacancies.

The refinement converged to give values of 5.2%, 6.9% and 1.9% for the profile,

weighted profile and Bragg intensity R factors respectively. The calculated superlattice

reflections with this model are predicted to be very weak and appear to be broadened more

than the main fluorite subcell reflections suggesting significant disorder. The thermal

parameters are large as a consequence of disorder and possibly also reflect deviations from

the ideal pyrochlore structure, known to be present at the pure 1 : 1 composition. The

oxygen positional parameter refines to a value of 0.3230(9), close to the value of 0.3175

for an ideal octahedral B site. A model assuming a random distribution of cations between

both sites fits almost as well as the pyrochlore model, as expected from the weak intensities

of the superlattice reflections. In contrast, a random model fit with the ideal fluorite oxygen

position (x = 0.375) gives a significandy poorer fit. The presence of the superlattice

reflections is confirmed in samples which have been heated to higher temperature and have

consequently larger particle sizes and sharper lines. Figure 4 compares X - ray diffraction

data for a sample with x = 0.06 at 400°C and at 685°C. The superlattice reflections (with

h,k,l all odd) become sharper and more easily discemable at the higher temperature,

confirming that these phases are pyrochlores though substantially disordered and possibly

156
1 5

O

X

»

c

o>c

Bi / Sn = 1.93

Calcined at 400°C

u

I I II II II I II I II 1 1 1 1 1 11 I II I II I II n s

8 9 Two Theta

Figure 3. A room temperature powder X-ray diffraction pattern of a Bi2(Sn2-xBix)07.x/2

phase with Bi/Sn = 1.9. The observed and calculated data are shown together with a

difference plot. The vertical lines mark the positions of the contributing Bragg reflections.

of lower than Fd3m symmetry.

The compound Bi2Sn207 is the only stable phase at high temperature[22] and

consequently all of the expanded compositions decompose when heated to higher

temperatures. The decomposition temperature depends on the composition and varies from

approximately 700°C when x = 0.06 to 500°C at x = 0.86. A differential thermal analysis scan (10°C/min to 800°C) on a sample with x = 0.64 is featureless but on cooHng shows a sharp endotherm at 642°C. On reheating, a sharp exotherm is observed at 673°C, and on

recooling, the endotherm is observed to occur again at 642°C. The thermal behavior is characteristic of Bi203 which is formed by decomposition of the expanded composition on

the first heating cycle. X - ray diffraction patterns from the products confirm the presence

of bismuth oxide together with the stoichiometric pyrochlore phase Bi2Sn2C)7. The latter is observed to have the doubled cubic unit cell previously reported[21].

Catalytic results for one sample of Bi2(Sn2-xBix)07-x/2 with x = 0.71 are shown in

Figure 5. The top part of the figure shows the distribution of the major products CO, CO2,

157
c

o

c

400''C

(0c

o

685°C

( 22 2)

(440)

( 622)

1400)

J

( 844)

(444) (662)

(840)

(800) ( 10,22)

61 .25

(111) vj

(331 )(333)

(511)

1 1 \ 1 h

Two Theta 61 .25

Figure 4. Comparison of the X-ray diffraction patterns for Bi2(Sn2-xBix)07-x/2 with

Bi/Sn =1.06 at 430 and 685°C.

C2H6, C2H4 and higher hydrocarbons (mostly C3) at different conversions. The residence

times required to achieve these conversions are shown in the bottom part of the Figure.

The general form of the data is the same for other bismuth - tin pyrochlore based catalysts

examined, though there are some indications that the hydrocarbon selectivity increases by a small amount as x increases. For the specific catalyst used to obtain the data in Figure 5,

the selectivity is 60.2% at 1.4% methane conversion and rises to a maximum of 68.3% at 17.7% conversion. The maximum yield (selectivity x conversion) is 21.4 at 40.8%

methane and 87.3% oxygen conversions respectively for this set of data. In most respects

158
m C2-

oa C2=

C3:.

C02

CD

0.8

0.6

0.4

0.2

\1 kWWN

1.4 3.2 8.2 17.7 28.1

CH Conversion (%)

33.7 40.8

T

I I I I 0 10 20 30 40 50

CH^ Conversion (%)

Figure 5. Product distribution at different conversions for a Bi2(Sn2-xBix)07-x/2 catalyst

with X = 0.71 at 850°C and CH4 : O2 : Ar = 2 : 1 : 4 (top); residence time vs. conversion

for the same catalyst (bottom).

the data resemble results for other good oxide catalytic systems. Little CO is observed and

the major deep oxidation product is CO2. The ethylene to ethane ratio increases with

increasing conversion and the higher hydrocarbons reach about 10% of the total hydrocarbon product.

u

VMo

E

a>u

c

0)

'35 a

0.15

0.1

0.05

159
At the the reaction temperatures used, the expanded pyrochlore compositions are not stable and decompose to Bi203 and Bi2Sn207. However, physical mixtures of the two

components with comparable surface areas do not give catalysts with comparable

selectivities and the individual components have been shown to be only average catalysts.

At this time, the nature of the interaction between the phase separated bismuth oxide and the

Bi2Sn207 pyrochlore which leads to enhancement of the catalytic performance is not

understood.

The bottom panel of Figure 5 deserves further comment. The data show a nonlinear dependence of the methane conversion on the residence time. An initial period of

lower rate (below approximately 10% conversion) is followed by a period of decidedly

faster rate. This rapid increase is a consequence of the buildup of a gas phase flame radical

pool which augments the surface methane activation rate. Such autocatalytic behavior is common in gas phase hydrocarbon oxidations and pyrolysis[301. At longer residence

times and higher conversions the system becomes limited by the oxygen availability. At the

high temperature of the experiments (850°C), a model of the gas phase radical chemistry can provide an adequate description of most of the features of the hydrocarbon product

distribution[31]. The catalyst provides a source for methyl radical generation and catalyses

the oxidation of CO to CO2. A detailed model of the gas phase chemistry involved in

methane conversion chemistry and comparison with experimental data on product

distributions and isotopic labelling experiments will be presented elswhere[32].

CONCLUSIONS

A new series of bismuth tin pyrochlore oxides, Bi2(Sn2-xBix)07.x/2.

(0.05<x<0.86) have been synthesized by precipitation from aqueous potassium hydroxide

and calcination at 400°C. The bismuth rich pyrochlores decompose to Bi203 and Bi2Sn207

at temperatures in the range 700°C to 500°C. The compounds with higher x values

decompose at the lower temperatures. The phase separated compositions are good catalysts

for the direct conversion of methane to higher hydrocarbons via oxidative coupling.

REFERENCES

1. Hutchings, G.J., Scurrell, M.S. & Woodhouse, J.R.; Chem. Soc. Rev. 18, 251-283

(1989).

2. Keller, G.E. & Bhasin, M.M.; J. Catal. 73, 9-19 (1982).

3. Hinsen, W., Bytyn, W. & Baems, M.; Proc. 8th Int. Congr. Catalysis, Berlin, 3, 581-

592 (1984).

4. Hinsen, W. & Baems, M.; Chem.-Ztg 107, 223 (1983).

5. Ito, T. & Lunsford, J.H.; Nature 314, 721-722 (1985).

6. Ito, T., Wang, J.-X., Lin, C.-H. & Lunsford, J.H.; J. Amer. Chem. Soc. 107, 5062-

5068 (1985).

7. Otsuka, K., Jinno, K. & Morikawa, A.; Chemistry Letters 499-500 (1985).

160
8. Otsuka, K., Jinno, K. & Morikawa, A.; J. Catal. 100, 353-359 (1986).

9. Jones, C.A., Leonard, J.J. «fe Sofranko, J.A.; J. Catal. 103, 311-319 (1987).

10. Hatano, M. & Otsuka, K.; Inorg. Chim. Acta 146, 243-247 (1988).

11. Hatano, H. & Otsuka, K.; J. Chem. Soc. Faraday Trans. I 85, 199-206 (1989).

12. Kharas, K.C.C. & Lunsford, J.H.; J. Amer. Chem. Soc. 2336 (1989).

13. Machida, K. & Enyo, M.; J. Chem. Soc. Chem. Commun. 1639-1640 (1987).

14. Thomas, J.M., Kuan, X. & Stachurski, J. J.; Chem. Soc. Chem. Commun. 162-163

(1988).

15. Thomas, J.M. & Ueda, W.; J. Chem. Soc. Chem. Commun. 1148-1149 (1988).

16. Thomas, J.M. & Ueda, W.; Proc. 9th Int. Cong. Catalysis, Calgary, 4, 960-967

(1988).

17. Subramanian, M., Aravamudan, G. & Subba Rao, G.V.; Prog. Solid St. Chem. 15,

55-143 (1983).

18. Ashcroft, A.T., Cheetham, A.K., Green, M.L.H., Grey, CP. & Vernon, P.D.F.; J. Chem. Soc. Chem. Commun. 1667-1669 (1989).

19. Roth, R.S.; J. Res. Nat. Bur. Standards 56, 18-27 (1956).

20. Gattow, G. & Fricke, H.; Z. anorg. allg. Chem. 324, 287296 (1963).

21. Shannon, R.D., Bierlein, J.D., Gillson, J.L., Jones, G.A. & Sleight, A.W.; J. Phys.

Chem. Solids 41, 117-122 (1980).

22. Vetter, G., Queyroux, F. & Gilles, J.-C; Mater. Res. Bull. 13, 211-216 (1978).

23. Horowitz, H.S., Longo, J.M. & Lewandowski, J.T.; Mater. Res. Bull. 16, 489-496

(1981).

24. Horowitz, H.S., Longo, J.M., Horowitz, H.H. & Lewandowski, J.T.; ACS

Symposium Series, Eds. Grasselli, R.K. & Bradzil, J.F. 279, 143-163 (1985).

25. Roos, J.A., Bakker, A.G., Bosch, H., van Ommen, J.G. & Ross, J.R.H.; Catalysis

Today 1, 133-145 (1987).

26. Solymosi, F., Tombacz, I. & Kutsan, G.; J. Chem. Soc. Chem Commun 1455-1456

(1985).

27. Rietveld, H.M.; J. Appl. Cryst. 2, 65-72 (1969).

28. Wiles, D.B. & Young, R.A.; J. Appl. Cryst. 14, 149-151 (1981).

29. International Tables for Crystallography, D. Reidel, Dordrecht, Holland, (1983).

30. Pyrolysis: Theory and Industrial Practice; Albright, L.F., Crynes, B.L., Corcoran,

W.H., Eds.; Academic Press: New York, (1983)

31. Labinger, J.A. & Ott, K.C.; J. Phys. Chem. 91, 2682-2684 (1987).

32. Mims, C.A., Dean, A.M. & Hall, R.B.; to be published (1990).

161
DISCUSSION

S. Roth : It appears to me that based on the faster surface exchange reaction with CO2

than oxygen, a good way to carry out this reaction would be to use CO2 and then still get

methyl radical coupling but also make CO and is that possible?

A. J. Jacobson : The use of CO2 as an oxidant for methane conversion in this process has

been considered but I'm not aware of a good catalytic route.

A. Navrotsky : Still on the subject of rapid CO2 exchange I don't remember the details

precisely, but there has been a lot of geochemical work on oxygen exchange between

phases, and I remember considerable work that suggests that carbonates and silicates

exchange oxygen much more rapidly than silicate/ silicate pairs. That may be additional

evidence.

A. J. Jacobson : I'd be very interested to see that, Alex, because the exchange is an

interesting result which was unexpected.

A. Navrotsky : I'll try to dig that out.

162
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

SOFT CHEMICAL SYNTHESIS OF METASTABLE

TITANIUM, VANADIUM, AND MOLYBDENUM OXIDES

T.P. Feist and P.K. Davies, Dept. of Materials Science,

University of Pennsylvania, Philadelphia, PA 19104

ABSTRACT

Low temperature synthetic methods have been used to

prepare metastable transition metal oxides from a variety of layered precursors. Aqueous" ion exchange of

several alkali metal titanates followed by heating at

temperatures less than 500°C leads to the formation of

Ti02 (B) and other metastable titanates. Acid treatment

of layered brannerite- type phases AVMoOg (A=Li,Na) has

been used to synthesize new compounds

A'q 13VQ i2^0q 87^3 (A'=H,Na) with the "hexagonal

molybdenum oxide" structure. Methods of synthesis are

presented, and mechanisms of the transformations of the

various layered starting materials to new framework

structures are discussed.

INTRODUCTION

Many transition metal oxides are of great interest

for catalytic and electrochemical applications, and

metastable structures often demonstrate new or enhanced

properties compared to their thermodynamically stable

forms. These materials can often be prepared at low

temperatures by electrochemical intercalation, inorganic

ion exchange, or organic precursor methods. A specific

example is the preparation of Ti02 (B) , a metastable form

of Ti02, from the layered titanate K2Ti40g (fig. 1)

using soft chemistry. (1) The new material demonstrated

several interesting properties, including enhancement of

the catalytic activity of V2O5 when used instead of

anatase as the catalyst support and facile intercalation

i

163
Cs^gOii • "Step 5" Ti02(B)

KEY

9 alkali metal cation ^ MOg octahedron

Figure 1. The structures of several alkali mfal titanates and Ti02(B)

164
by alkali metal cations and hydrogen to form a mixed

valence titanium oxide bronze (2)

.

Subsequent work has shown that Ti02 (B) can be

prepared from a variety of other titanate structures

(3,4). but the mechanistics of the transformation from

layered titanate to Ti02 (B) have remained poorly

understood. It is the goal of the present work to

investigate the intermediate structures observed during

the dehydration of the different starting materials in

order to better understand the mechanistics of the

reaction and the stability of Ti02 (B)

.

In the second portion of the paper we report the

synthesis in aqueous acid of A'g i2'^0 13^°0 87^3

(A'=H,Na) from the compounds AVMoOg (A=Li,Na) which have

the layered brannerite structure, (fig. 2a) The

framework of the new structure is isotypic with the

hypothetical hexagonal tunnel structure of M0O3 (5)

shown in figure 2b. Conventional synthesis techniques

have been used to stabilize this structure by

incorporating large cations into the tunnels, but there

have been no reports of the preparation of the structure

containing empty tunnels. If such a material could be

prepared, it would likely be a candidate for a variety

of intercalation and catalytic reactions. We report

here the synthesis and characterization of new hexagonal

Mo03-type tunnel structures by a novel soft chemical

route of acid treatment and dehydration of alkali

brannerites

.

EXPERIMENTAL

Alkali metal titanates were prepared by solid state

reaction of Ti02 (anatase) and alkali metal carbonates

or nitrates as reported previously. (4) Acid exchange

of the layered titanates (fig. la) was achieved by

stirring the pulverized titanate in IM HCl (200 -fold

excess) for 24-72 hours at 60°C. Acid exchange of the

tunnel structure Na2Tig023 (fig. 1) was achieved by

immersing the powder in molten NH4NO3 for 5 days at

200°C. Thermogravimetric analysis was used to

investigate weight loss behavior of acid exchanged

materials, and TGA/Mass Spectroscopy confirmed that the

165
only species evolved during heating was H2O.

Differential scanning calorimetry (DSC) was used to

examine the energetics of the dehydration reactions, and

X-ray diffraction (XRD) was used to probe the structure

of intermediates isolated during the formation of

Ti02 (B)

.

Lithium and sodium brannerites were prepared by

solid state reaction of AVO3 (A=Li,Na) and M0O3 as

reported by Galy et al . (6) . The Na brannerite was

stirred in 0.25M HCl for 12 hours at 60°C, and the solid

was recovered by vacuum filtration. The process was

repeated to insure complete conversion of the brannerite

to the tunnel structure. LiVMoOg was immersed in 0.25M

HCl for 36 hours with stirring at 60°C. The solution

was then acidified with IM HCl and stirred for 2 hours

before isolating the product by vacuum filtration. The

products were analyzed by TGA, XRD, and atomic

absorption spectroscopy.

RESULTS AND DISCUSSION

(1) TITANATES

The TGA curves for each of the acid exchanged

layered titanates (fig. 3a) show multistep weight loss

behavior, indicating the possible formation of stable

intermediate structures. The final dehydration product

in each case is Ti02 (B) . X-ray diffraction confirms

that intermediate structures can be isolated during the

dehydration reactions. Figure 3b shows the powder

diffraction patterns of materials isolated during the

transformation of the trititanate H2Ti30'y to Ti02 (B) . The pattern of the first intermediate is very similar to

that of the condensed structure Na2TigO-|_3, and the

weight loss associated with the formation of the

intermediate is equal to 1 mole of water per 2 moles of

H2Ti307 . The reaction can be formulated:

2H2Ti307 ^ H2Ti5023 + H20T

which corresponds to a topotactic condensation of the

initial layered structure. Upon elimination of water,

layers join through the previously unshared oxygen atoms

at the corner of each "step". DSC shows that this

process is endothermic (fig. 4), which is consistent

166
(a)

H2TI3O7

H2TI4O9 H2O

H2TI5O, |- 2.5H2C

200 400 600 800 1000

Temperature (°C)

350°C; Ti02(B)

200°C; 1 month

140°C; 1 month

3 H2Ti307

Figure 3. (a) Thermogravimetric analysis of layered titanates.

(b) Powder X-ray diffraction of H2Ti307 dehydration

products.

167
with its topotactic nature. XRD and DSC indicate that

the other layered titanates behave similarly to H2Ti30'7

.

H2TigO-]_3, which has also been prepared by direct acid

exchange of the tunnel structure Na2TigO-j_3, does not

show any low temperature endotherm because the structure

is already condensed. (fig. 4)

X-ray diffraction shows that the next step results

in the formation of an intermediate which is

structurally similar to Ti02 (B) and is independent of

starting titanate structure. Although the powder

pattern could not be indexed, electron diffraction

indicated that the new structure has a C- centered

monoclinic cell with approximate lattice parameters of

a=:14A, ip=3.75A, and c=20A. P is not known. DSC shows

the formation of this structure, in contrast to the

previous topotactic step, is accompanied by a strong

exothermic heat effect. This leads us to postulate that

the intermediate which immediately precedes Ti02 (B)

forms by nucleation and growth from the condensed

structures. Ti02 (B) then forms in a single, final step

that shows no strong heat effect, presumably because the

structure of the last intermediate is very similar to

that of Ti02 (B)

.

Figure 4. Differential scanning calorimetry of acid exchanged

"step 3" titanates.

168
(2) VANADIUM -MOLYBDENUM OXIDES

The reactions which lead to the formation of the

hexagonal Mo03-type structure from alkali metal

brannerites are quite different than the chemistry of

layered titanates. Because vanadium and molybdenum have

appreciable solubility in dilute acid, the sodium and

lithium brannerites do not ion exchange, but instead

undergo "leaching" reactions. Vanadium is

preferentially dissolved from the brannerite, and a new

structure crystallizes. In the case of Na brannerite,

the process proceeds though a two -phase mixture of the

starting structure and the tunnel structure, with no

evidence of other phases. The final product,

^^0 . 13^0 . 13'^°0 . 87^3 * • 26H2O can be dehydrated at 400°C

but rehydrates immediately upon cooling, (fig 5a) The

tunnels, with a minimiam diameter of 3.2A, are too large

for Na"*" to occupy without coordination by water

molecules. Therefore, the rapid rehydration observed

during cooling suggests that sodium remains in the

tunnels. XRD confirms that the tunnel structure is

maintained throughout the dehydration/ rehydration

process

The Li brannerite behaves much differently. Upon

immersion in dilute HCl, the material dissolves, leaving

only a small amount of amorphous residue, (fig. 5) As

the solution continues to be stirred at 60°C, the

hexagonal tunnel structure begins to crystallize from

solution. Acidification of the solution is used to

remove any amorphous phase remaining after 3 6 hours, and

the final product is single phase with yellow needlelike particles 1-lOpm in length. Atomic absorption

shows that the Mo:V ratio is 6.7:1 and the material

contains little lithium. The final s toichiometry is

Hq 23^0 13^°0 87^3 ''^^2'^' Subsequent dehydration of the

hexagonal tunnel structure without residual alkali metal

ions allows preparation of the material with empty

tunnels. The TGA (fig. 5b) indicates that after the

loss of surface water, the water molecules which are

presumably located in the tunnels are removed, leading

to the formation of Hg 13^0 . 13M00 _ 87O3 at approximately

225°C. During the final weight loss, which corresponds

169
to 0.065(H2O), the residual protons are apparentlyabstracted from the structure with a small amount of

oxygen from the framework, leaving a sample with the

stoichiometry Vq ^ -]_3^*^0 . 87^2 . 93 5

* This yellow-green

material retains the hexagonal M0O3 structure, and

neutron diffraction refinements have confirmed the

absence of any species in the tunnels. Preliminary

experiments indicate that this new material can be ionexchanged with a variety of divalent cations, and it

also undergoes insertion reactions with H"*", Li"*", and

Na"*". Further experiments are currently being performed

to clarify the mechanisms of these unusual formation

reactions and to utilize the wide variety of reactions

exhibited by brannerites in order to form new compounds

with potential applications as catalysts and electronic

ceramics

.

101 1

1

0 50 100 150 200 250 300 350 Temperature (*C) Figure 5. Thermograviitietric analysis of acid treated sodium

brannerite (a) and lithium brannerite (b)

.

170
Ol

ol

CDI

COI

"i

cui

CD cn

l\ J

UJ iJ

brannerite + amorphous phase

-AjV— A 11 . -

LiVMoOg (brannerite)

5.0 10 15 20 25 30 35 40 45 50 55 60

Figure 6. Powder X-ray diffraction of acid treated LiVMoOg.

CONCLUSION

We have presented the formation of Ti02 (B) from

several different starting layered titanates and have

considered mechanistic aspects of the transformation

from layered structures to a three-dimensional

framework. We have also presented a novel synthesis of

a hexagonal vanadium/molybdenum tunnel structure with

empty tunnels. The electrochemical and catalytic

properties of this material are currently being studied.

ACKNOWLEDGMENTS

The authors gratefully acknowledge the support of

Exxon Research and Engineering Corp., International

Business Machines Corp., and the National Science

Foundation (DMR- 88 - 19885) .

REFERENCES

1. Marchand, R. , Brohan, L., and Tournoux, M. , Mat.

Res. Bull., 15, 1129-1133 (1980).

2. Tournoux, M. , Marchand, R. , and Brohan, L., Prog.

Solid St. Chem. , 17, 33-52 (1986).

3. Mocarski, S.J., Master's Thesis, University of

171
Pennsylvania (1986).

4. Feist, T.P., Mocarski, S.J. Davies, P.K., Jacobson,

A.J., and Lewandowski, J.T., Solid State Ionics, 28-30,

1338-1343 (1988)

.

5. Caiger, N., Crouch - Baker , S., Dickens, P., and James,

G., J. Sol. St. Chem. , 369-373 (1987).

6. Galy, J., Darriet, J., and Darriet, B., C.R. Acad.

Sc. Paris, Serie C, 264, 1477-1480 (1967).

172
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

PRELIMINARY SOLID-STATE MULTINUCLEAR NMR

OF TITANIUM AND ZIRCONIUM OXIDE

CERAMICS AND PRECURSORS

John J. Fitzgerald* and Sang-sub Han

Department of Chemistry

South Dakota State University

Brookings, SD 57007

and

Steven F. Dec, Mark F, Davis,

Charles E. Bronnimann and Gary E. Maciel

Department of Chemistry

Colorado State University

Fort Collins, CO 80523

ABSTRACT

Prehminary solid-state multinuclear NMR studies of selected titanium and

zirconium oxide electronic ceramic powders and precursors are described.

CRAMPS NMR has been used to examine the metal-hydroxide and metal-hydrate

protons sites in a range of hydrolyzed titanium(IV) and zirconium(IV) salts which

may serve as solution precursors of Zr(IV) and Ti(IV) oxide ceramic materials.

High-field (14,1 T) magic-angle spinning (MAS) ''^Ti, ''^Ti and ^^Zr spectra are reported for a range of Ti(IV) and Zr(IV) oxides including Ti02 (anatase), CaTi03,

SrTi03, BaTi03, SrZr03 and BaZr03. In addition, MAS ^^^Ba NMR results are also

reported for BaO, BaTi03 and BaZrOg. These initial investigations suggest the

potential of sohd-state NMR approaches to advance our fundamental knowledge of

the structure, metal ion environments, and reactions of a wider range of titanium and

zirconium oxide electronic ceramics and related precursors of current interest.

'Author to whom correspondence should be sent.

173
INTRODUCTION

A diverse range of electronic ceramics containing titanium(IV) and

zirconium(IV) ions in perovskite-type oxide lattices are of importance in materials

science. ^"^ Barium titanate (BaTiOj), and various alkaline earth titanates and

zirconates, have found important uses in micro-electronic packaging and thin film

capacitor applications/'^ The utility of titanium oxide-containing ceramics such as PZTs (lead zirconate titanate (Pbj^ZryTij_y03)) and PLZTs (lead lanthanum zirconate

titanate (Pbj.j^La^ZryTi^.^)j.j^/403)) powders lie in their widespread use as high quality

transparent electrooptic and piezoelectric materials.^"^'^'^'^^

The synthesis and processing of Ti(IV) and Zr(IV)-oxide containing ceramics

employ solid-state mixed oxide reactions, the thermal decomposition of metal

hydroxide and metal chelate precursors, and metal alkoxide sol-gel methods."*'^ For

example, BaTi03 powders and ceramics can be synthesized by numerous reactions

involving a variety of precursor materials:

BaO + Ti02 - —> BaTiOg (1)

BaC03 + Ti02 - —> BaTi02 + CO, (2)

Ba(OH)2 + Ti02 > BaTiOg + U^O (3)

BaTiO(Ox)2»2H20 —-> BaTi03 + 200, + 2H2O (4)

Ba(OR)2 + Ti(OR)4 > Sol-Gel Precursor (5)

Sol-Gel Precursor -"^-> -TM?--> BaTiOg

Similarly, the synthesis of PZT and PLZT has been reported^'^'^^ using solid- state oxide or carbonate mixtures of the various metal ion constituents and synthetic

approaches using various hydrolysis precursors such as ZrOClj* 8H2O or TiOS04

solutions/'^

The structural characterization of Ti(IV) and Zr(IV) oxide ceramics, their

synthetic precursors and the detailed examination of these synthesis reactions are of

major significance to our understanding of both their chemistry and functional

properties.

Solid-state multinuclear nuclear magnetic resonance (NMR) represents a unique and important spectroscopic technique with tremendous potential to advance

our fundamental knowledge of the structure and reactivity of titanium and zirconium,

oxide ceramics and their related precursors. This report describes the applications

of solid-state multinuclear NMR approaches to examine selected titanium and

zirconium oxide ceramic precursors and powdered materials. CRAMPS NMR^°

has been used to examine titanium and zirconium hydroxide-containing complexes.

MAS ^^Ti, '•^Ti, "^^Zr and ^^^Ba NMR has been used to examine the M(II) and M(IV)

metal ion environments in mixed metal oxides such as BaTi03 and BaZrOg^''^'.

174
EXPERIMENTAL

Solid-state MAS ''Ti, ^^Ti, ^^Zr and "^Ba NMR were obtained on a Bruker

AM-600 NMR spectrometer, using a home-built probe at spinning speeds of 5-6 kHz.

"^^Ti and '^^Ti NMR spectra were obtained at 33.8 MHz, ^^Zr NMR at 55.8 MHz and

^^^Ba NMR at 66.7 MHz using the external references TiCl4(f), BaZrOjCs) and

BaZr03(s), respectively, assigned as 0.00 ppm. Spectra were acquired using a Hahn

spin echo-pulse sequence in order to eliminate artifacts from probe "ringdown."

^H CRAMPS NMR were obtained on an extensively modified Nicolet NT-200

NMR spectrometer operating at 187 MHz using the BR-24 pulse sequence.

Samples were contained in glass tubes using a modified Gay spinner. The ^H NMR

chemical shifts were externally referenced to tetrakis(trimethylsilyl)methane

(TTMSM) and reported here relative to TMS.

RESULTS AND DISCUSSION

"^^Ti and "^^Ti NMR. '^^Ti and ''^Ti NMR spectra have been previously reported

for only BaTiOj.^^'^^ A wider range of alkaline earth titanates have been examined

here using MAS "^^Ti and '^^Ti NMR measurements. Titanium has two NMR-active

isotopes, '^^Ti (1 = 5/2, 7.28% abundance) and ''^Ti (1 = 7/2, 5.51% abundance), both

of which have very low gyromagnetic ratios and medium quadrupole moments (Q

C^^Ti) = 0.29xl0--^m-, Q ('^Ti) = 0.24xl0-^^m-). These parameters make these

nuclides unattractive for NMR study except at the highest available magnetic fields

(11.7 and 14.1 T).

Figure 1 shows the solid-state "^^Ti and '^^Ti MAS NMR spectra of CaTiOj,

SrTi03 and BaTiOj obtained at 33.840 MHz using a 14.0 T magnetic field and

spinning speeds of 5-6 kHz. The NMR reference is a neat solution of TiCl4(C) taken

as 0.0 ppm. No spectra were observed for MgTiOg under these conditions. The

47rj.-y49j- jsj]yjj^ spectra show only the central transition (1/2 <--> 1/2) resonance for

both nuclides. The "^^Ti resonance is the most intense and narrowest signal, being

about 270 ppm downfield from the '^'^Ti signal for SrTi03. The '*^Ti NMR peak is much broader for both CaTi03 and BaTi03 than that observed for SrTi03. The

chemical shifts of the "^^Ti resonance for CaTi03, SrTi03 and BaTi03 "^74, -843,

and -746, respectively, and are assigned to the 6-coordinate titanium(IV) ion.^"^"^^ The

variation in the "^Ti chemical shifts is probably due to the nature of the M(II) cation

in these lattices, whereas changes in the local symmetry about the Ti(IV) nucleus

probably significantly affect the linewidth of the '*''xi/'*^Ti central transition resonances due to second-order quadrupolar broadening.

The MAS ''^Ti and "^^Ti NMR spectrum of the anatase form of Ti02 is given

in Figure 2. The spectrum is dominated by the "^^i central transition, which shows

a second-order powder pattern with two peak maxima at -938 and -994 ppm relative

to TiCl4({). The '''^Ti resonance is upfield with observable intensity in the -1150 to - 1450 ppm range, and is significantly broadened relative to the ''^Ti signal. The

second-order powder pattern for the "^^Ti NMR resonance is due to the large electric

field gradient about the Ti(IV) ion in this axially symmetric 6-coordinate oxide lattice

site of anatase.

^"^'^^

175
^'Zr and ^'"Ba NMR. The ^'Zr nuclide is 11.23% abundant with a nuclear

spin of 5/2. To date, no solid-state NMR spectra have been reported for this

nucleus. Figure 3 shows the MAS ^^Zr NMR spectra of SrZr03 and BaZr03 taken

at 66.7 MHz using a 14.1 T magnetic field and spinning speeds of 5-6 kHz. No ^^Zr

NMR spectra were observed for MgZr03 and CaZr03 under these conditions. The

^^Zr NMR show only the central transition resonance, with BaZr03 exhibiting a sharp

signal taken as the chemical shift reference at 0.0 ppm. SrZr03 exhibits a broadened, unsymmetrical NMR resonance centered near -29 ppm. This resonance may be due to second-order quadrupole interactions or due to a range of chemically

distinct 6-coordinate zirconium(IV) ion sites in this sample.

The ^^^Ba nuclide (1 = 3/2) is 11.32% abundant and has Q = 0.28xl0"-W.

NMR spectra obtained for BaO, BaTi03 and BaZr03 at 66.7 MHz in a 14.1 T

magnetic field and 5-6 kHz spinning speeds are shown in Figure 4. BaZr03, which

has cubic 12-coordinate Ba(II) ion sites,^"* gives the sharpest signal and is used as the

reference standard. BaTi03, likewise, has a perovskite structure, although the 12-

coordinate Ba(II) ion sites are distorted in the tetragonal phase at room

temperature.^^ BaO has a NaCl lattice with six-coordinate barium(II) ion sites^'^ and

exhibits a ^^^Ba NMR signal at 481 ppm relative to BaZrOg.

^H CRAMPS NMR. ^H CRAMPS NMR have been reported for silica and

alumina/silica materials. Ceramic oxides of Ti(IV) and Zr(IV) ions are often

prepared from metal hydrates, hydroxides and various other complex hydrolysis

precursors containing M-OH, M-OH-M and M-OHj proton-containing structural

moieties.

Figure 5 shows the ^H CRAMPS NMR of various Ti(IV) and Zr(IV) salts

obtained at 187 MHz using the BR-24 pulse sequence and 2-3 kHz spinning speeds.

The spectrum of ZrO(OH)(C03), extensively hydrolyzed Zr(IV) complex,

probably containing primarily M-OH moieties, shows a single, broad resonance at 5.7

ppm. The ^H CRAMPS spectrum of ZrOCl2»8H,0, containing the tetrameric

Zr4(OH)8(H20)i/^ ion,^^ with a proton M-OH2/M-bH ratio of 4/1, exhibits two

resonances at 7.1 ppm and 9.5 ppm assigned to M-OH2 and M-OH protons,

respectively. The spectrum of crystalline Zr(S04)2» 4H2O, a structure with linear

chains of Zr(IV) ions with bridging OH groups,^^ gives a single, sharp resonance at 8.8 ppm. The ^H CRAMPS spectra of TiOS04» H2SO4 H.O, a complex material

containing a range of Ti-OH, Ti-OH-Ti, Ti-OH2 moieties, exhibits a H CRAMPS

NMR spectrum consisting of four sharp resonances in the 5-10 ppm range.

CONCLUSIONS

The initial ''^Ti and '^^Ti NMR spectra reported here demonstrate the

application of solid-state NMR to observe 6-coordinate Ti(IV) sites in cubic (SrTi03),

tetragonal (BaTi03), and axially symmetric titanium oxide (TiOj) materials. The ''^Ti chemical shifts of the '*^Ti NMR resonance for the alkaline earth titanates are dependent on the size and nature of the M(II) ion.

The ^^Zr NMR of SrZr03 and BaZr03 also demonstrate the capability of

high-field solid-state NMR to observe cubic (BaZr03, SrZr03) 6-coordinate Zr(IV)

176
ion sites in oxide lattices. The large linewidth observed for SrZr03 relative to BaZrOj possibly indicates a significant electric field gradient at the ^^Zr nucleus in

this perovskite lattice.

^^''Ba NMR reported herein allow one to distinguish between 6-coordinate

BaOg sites (in BaO) and 12-coordinate BaOi2 sites (in BaZr03 and BaTi03). The

distortions in the Ba(II) ion sites in tetragonal BaTi03 result in second-order

quadrupolar peak broadening.

The CRAMPS NMR results for a series of M-OH and M-OH2 containing

Ti(IV) and Zr(IV) salts suggested that this technique will provide an important

approach to examine both metal hydrolysis, metal chelate and sol-gel precursors of

Ti(IV) and Zr(IV) oxide ceramics.

ACKNOWLEDGEMENTS

The authors gratefully acknowledge partial support of this work by National

Science Foundation Grant RII-8902066 and the use of the Colorado State University

Regional NMR Center, funded by National Science Foundation Grant CHE-86 16437.

REFERENCES

1. B. Jaffe, W.R. Cook and H. Jaffe, Piezoelectric Ceramics. London, Academic

Press (1979).

2. R. Buchanan, ed.. Ceramic Materials For Electronics. Marcel Dekkar, Inc.,

New York (1986).

3. K. Wakino, "Piezoelectric and Pyroelectric Ceramics," in Fine Ceramics. S. Saito, ed., Elsevier, New York (1985), pp. 251-261.

4. A. Amin, M.A. Spears and B.M. Kulwicki, J. Am. Ceram. Soc. 66(10), 733-

738 (1983).

5. T.T. Fang and H.B. Lin, J. Am. Ceram. Soc. 72(10), 1899-1906 (1989).

6. W. Hertl, J. Am. Ceram. Soc. 71(10), 879-883 (1988).

7. J.J. Ritter, R.S. Roth and J.E. Blendell, J. Am. Ceram. Soc. 69(2), 155-162

(1986).

8. K. Kezuka, Y. Hayashi and T. Yamaguchi. J. Am. Ceram. Soc. 72(9), 1660-

1663 (1989).

9. G.H. Haertling and C.E. Land, J. Am. Ceram. Soc. 54(1), 1-11 (1971).

10. C.E. Bronnimann, B.L. Hawkins, M. Zhang and G.E. Maciel, Anal. Chem.. 60, 1743 (1988).

177
11. CE. Forbes, W.B. Hammond, N.E. Cipollini and J.F. Lynch, J. Chem. Soc.

Commun.. 433 (1987).

12. TJ. Bastow, J. Phys. Condens. Matter 1 . 4985-4991 (1989).

13. P. Duran and C. Moure, Am. Ceram. Soc. Bull. . 64(4), 575-579 (1985).

14. A.F, Wells, Structural Inorganic Chemistry. 5th Edition, Clarendon Press,

Oxford (1984), pp. 564, 588.

15. H.D. MeGaw, Proc. Roy. Soc. A, 189, 261 (1947).

16. H.F. Kay and P.C. Bailey, Acta. Cryst. . 10, 219 (1957).

17. P. Vousden, Acta. Cryst. . 9, 141 (1956).

18. I. Clearfield and P.A. Vaughan, Acta. Cryst. . 9, 555 (1956).

19. D.B. McWhan and G. Lundgren, Inorg. Chem.. 5, 284 (1966).

178
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

POLYMERIC SYNTHESIS OF PEROVSKITE POWDERS AND FILMS

H. U. Anderson, M. M. Nasrallah, F. D. Blum and M. S. Smith

University of Mi ssouri -Roll

a

INTRODUCTION

The electronic components industry is constantly striving to

make smallpr components with increased reliability and faster

information transfer. This requires the total integration of the

components into single devices. Improvements in integrated

electronic devices are directly linked to improvements in

properties which, in turn, are related to processing. This can be

achieved by utilizing the appropriate powder synthesis technique

and a thorough understanding of its chemistry.

Recent breakthroughs in the processing of materials have

increased the probability of success of electronic ceramics in

component integration. These include: improved sputtering and

chemical vapor deposition techniques, decreases in the sintering

temperatures of substrates e.g. AI2O3, development of relaxor

dielectrics with sintering temperatures around QOO'C, and the

development of organic derived sources of oxides which can be

processed at lower temperatures and can be used as precursors for

bulk and film devices. The latter, known as the "polymer

precursor" or "liquid mix" technique, is one of the most

significant and will be the focus of this work.

We have successfully used a modified liquid mix technique to

synthesize doped double oxide perovskite-type powders with

desirable sinterabil ity and electrical properties. Spun-on films

were also produced by this process. The synthesized perovskite

powders are of the general formula ABO3 where the A and B cations

occupy 12 and 6 coordinated sites, respectively. Their

electrical, magnetic and dielectric characteristics are strongly

dependant on their defect structure. This, in turn, is related to

the A/B ratio, the oxygen activity under which they are equilibrated, and the type and concentration of aliovalent dopants

that can be preferentially added on either the A or B site. The

liquid mix process appears to be very useful for production of

179
sinterable and homogeneous ABO3 type powders with the required

purity level and dopant content for the electronic industry.

Numerous compounds have been developed in our laboratories

including: donor doped Br and Sr titanates, counterdoped rutile,

acceptor doped Lanthanum and Yttrium chromites, Sr doped La

manganite, ferrites, cobaltites, niobates and superconductor

YBa2Cu307.3,. Our discussion here will be limited to the titanates,

manganites and the ¥88200307.^ compounds.

The technique has the advantage of producing single phase

ceramic powders with particle size as small as 50 nm and very

precise chemical compositions. Further processing of these fine

powders can be done at temperatures as low as BOO'C. The protocol

for the production of powders by this method was developed in the

1950' s^^\ however, no detailed study of the chemistry of these

systems has been reported.

POWDER SYNTHESIS AND CHARACTERIZATION

(i ) Liquid Mix Process

The liquid mix process utilizes complexing agents capable of

chelation, e.g. pyroxycarboxyl ic acids, to form polybasic acids

with various cations, e.g. Ti, Zr, Pb, Cr, Mn, Ba, La, etc. These

chelates can undergo polymerization in the presence of

polyhydroxyl alcohols or amines with a uniform distribution of the

"chelating" cations. The formed polymer precursor retains

homogeneity and stability (due to its high viscosity) and can be

calcined at relatively low temperatures to yield fine oxide

powders whose chemi stry has been carefully controlled. The

possible mechanism of polymer formation is as follows;

0 OH 0 0 OH 0

II I II .H^O II I

II HO-C-CH2-C-CH2-C-OH + HOCH2CH2OH -----> (0-C-CH2-C-CH2-C-0-CH2CH2)„

I I C - 0 C - 0

OH OH

Citric Acid Ethylene Glycol Polyester

The original Pechini method Involves mixing the required

amounts of soluble cation salts with ethylene glycol and citric

acid. The mixture is heated to about 90'C to insure complete

dissolution of the cation salts. The solution is then evaporated

to a polymeric residue. The metal ions are chelated to the

organic chain. The organic polymer turns into a black brittle

180
mass as it is heated to 400'C. This is ground, screened and

calcined at 600-800'C to yield a homogeneous fine particle size

oxide with the required composition.

(i i ) Ethylene Glycol Process

We have found that ethylene glycol oxidizes in the presence

of cold nitric acid to form oxalic acid. The latter consists of

two connecting functional carboxylic groups each of which loses a proton and may coordinate with the metal ion forming a ring

structure as shown:

CH2OH

CH2OH

Ethylene

glycol

CHpOH

ICHO

Glycol ic

aldehyde

CH2OH CHO

I » I

CO2H

CO2H

Glycol ic

acid

CO2H CO2H

O OH o o

Glyoxylic Oxalic

acid acid

\" \

C /

o o

This, "modified liquid mix" process^^', was found to possess

a wider range of application than was originally outlined by

Pechini. Titanates, zirconates, niobates, chromites, ferrites,

manganites, cobaltites and others have been successfully prepared

by this process. It offers the potential of synthesizing complex

double oxides with controlled stoichiometry, purity and particle

size. The powder produced is initially amorphous and as the

calcination temperature and time increase, crystallization

gradually occurs until the crystal structure of the desired oxide

is completely restored. This is detected by x-ray diffraction

techniques as a function of calcination temperature. The produced

powders are nearly spherical with about O.lum size as revealed by

scanning electron microscopy.

The weight loss and changes that occur as the polymeric

precursor is gradually heated are revealed by thermogravimetric

analysis (TGA) and differential thermal analysis (DTA). Both

techniques were carried out on the same sample. Results on LaMnOg

and BaTiOg are shown in figs. 1 a and b respectively. An

appreciable weight loss is observed during the early stages of

calcination, this also corresponds to the DTA exothermic peak

shown in fig. 1. Both samples were heated to IBO'C prior to

analysis to insure polymerization.

181
\00-

U) 80 3

60h

_Ja:

40

20

" n 1 I

- /

1 1 1

V-«-DTA

1 r 1

80

60

40

-20

Z

liJ

100 200 300 400 500 600 700 TEMPERATURE (°C)

(a)

800 100 200 300 400 500 600 TEMPERATURE (°C)

(b)

700 800

Fig 1. Relative Weight Loss (TGA) and Differential Thermal

Analysis (DTA) of Polymeric Precursors for (a) LaMnO,.

(b) BaTiOg.

(iii) Ethylene Diamine Process

A major limitation of the original Pechini process is the

possibility of precipitation during polymerization. This causes

inhomogeneities in the resulting oxide. It is therefore essential

that the mixture remains in solution throughout the polymerization

process. We have found that both citric acid and ethylene diamine

form chelates with the dissolved cations. The addition of

ethylene diamine not only retards precipitation (through pH

adjustments) but also enters the chelation process^^\ The

proposed reaction sequence is as follows:

o- M

-H2O

H2N NHH > O''

Ethylene diamine process,

182
Upon heating polyamidation occurs throughout the liquid

medium, resulting in a homogeneous solution of metal ions attached

to the organic matrix. This reaction indicates that not only

citric acid, but also ethylene diamine can form chelates with

dissolved metal ions increasing their concentration and the

stability of the system. This process has been successfully used

for synthesis of bulk and spun-on films of the superconductor

YBa2Cu307.,. The details of the process are published elsewhere^^\

Careful pH adjustment is critical to this process. The

synthesized powder was calcined at temperatures ranging from 300-

900'C whereby a change in structure from amorphous to the

orthorhombic YBa2Cu307., was observed. Fig 2 shows x-ray

diffraction patterns of YBajCUjOy., as a function of calcination

temperature. Below 300'C a diffused pattei-n was observed with no characteristic crystal structure, as the temperature increased the

crystalline superconducting phase appeared and became the

predominant phase above 700°C. The decomposition of the pclymsr

precursor to the desired oxide was also studied by simultaneous

DTA and TGA.

Both the ethylene glycol and the ethylene diamine processes

can be used for synthesis of a wide range of powders with improved

chemical homogeneity, increased reactivity, high purity, and

better compositional stoichiometry than other conventional

techniques. However, the mechanism of these processes is not

fully understood. This study is currently underway. One of the

techniques used for this purpose is nuclear magnetic resonance

(NMR) spectroscopy. It is based on the ability of interaction of

a nuclear magnetic moment with radio frequency waves (rf) in the

presence of an external magnetic field. This interaction is

affected by several factors including bonding, solvent,

temperature, concentration, structure, dipolar effects, etc. As a result of the interaction, one will observe a chemical shift of a particular NMR active species. The shift is defined as the

difference in resonance frequency between that of the active

species, due to the mentioned interactions, and that of a given

standard.

Pb-207 (22.6% natural abundance) has a very large chemical

shift range (-3500 ppm) and thus is very sensitive to changes in

chemical environment. Preliminary results have been obtained for

lead solutions in citric acid (CA) and ethylene glycol (EG). As

shown in fig 3, as the Pb"^^ concentration increases, a greater

change in chemical shift is observed with (CA) as compared to

(EG). Indicating that Pb"^^ ions interact more strongly with (CA)

than with (EG). When a mixture of both (CA+EG) is used, the

results indicate that the (CA) association is preferred over th.at

of (EG). This association can be interpreted in terms of Pb

183
chelation. We anticipate that this study will be extended to

other metal ion systems in order to better understand the

complexation process.

1 1

1 1 1

......Bo CO,

• CuO

- '

I

'-Co

sratO/jO, . . • . Ml s

I 450°C

TT ^ -1- . 1 , 20

2 e {degrees)

10

Fig. 2

X-ray diffraction as a function of calcination

temperature for

YBa^CUjO^.,.

250

fig. 3

Pb.207 NMR chemical

shift.

Moles of each ligand per mole Pb-MREFERENCES

1. M. Pechini, U.S. Patent No. 3,330,697, July 1967.

2. N. G. Eror and H. U. Anderson, Mat. Res. Soc. Symp. Proc,

Vol. 73, "Better Ceramic Through Chemistry 11", Eds. C.

Brinker, D. Clark and D. Ulrich, MRS, 1989.

3. C. C. Chen and H. U. Anderson, "Advanced Materials

Conference II, Proc", Ed. F. Smith, Adv. Matls. Inst.,

1989.

184
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

ELECTRODEPOSITION OF NANOMODULATED

ELECTRONIC CERAMIC THIN FILMS

Jay A. Switzer

University of Missouri-Rolla

Department of Chemistry and

Graduate Center for Materials Research

RoUa, MO 65401

ABSTRACT

Electrochemistry can be used for the atomic-level architecture of

ceramic materials. In this report, ceramic superlattices based on the

TlaPbbOc/TljjPbgOf system were electrodeposited with individual layer

thicknesses as thin as 3 nm. The superlattices were deposited from a single

aqueous solution at room temperature, and the layer thicknesses were

galvanostatically controlled. Substitution of TI2O3 into Pb02 appears to

stabilize a face-centered cubic structure with an average lattice parameter of

0.536 mn. The lattice parameters for the Tl^Pbi^Oc mixed oxides vary by

less than 0.3% when the Pb/Tl ratio is varied from 0.84 to 7.3. Because the

modulation wavelengths are of electron mean free path dimensions, this

new class of degenerate semiconductor metal-oxide superlattices may show

thickness-dependent quantum optical, electronic, or optoelectronic effects.

INTRODUCTION

We have recently demonstrated that it is possible to electrodeposit

nanomodulated ceramic superlattices based on the Tl^Pb^O^/Tl^jPheOf

system [1]. An idealized superlattice structure with square-wave modulation

of composition and/or structure is shown in Figure 1. The thicknesses of the

A and B layers are not necessarily equal, as long as the structure is periodic.

The idea of electrochemically depositing nanomodulated superlattices is not new, but it has not been applied previously to the deposition of

nonmetallic materials. Several research groups have shown that compositionally modulated metallic alloys can be electrochemically deposited from

185
Substrate

Figure 1. Idealized superlattice structure.

a single plating bath by cycling either the potential or current [2-5]. The

interest in nanomodulated metallic systems stems from their enhanced

mechanical and magnetic properties [3,6].

We have chosen the Tl3Pbi30^./TljjPbgOf system for our study for

several reasons: (1) during our previous work on the electrochemical and

photoelectrochemical deposition of thallium(III) oxide films, we found that

it was possible to deposit highly oriented films [7-9], (2) the deposition of

Pb02 [10,11] and PbgTl5024 [12,13] are well documented, (3) the anhydrous

oxides deposit directly at room temperature and require no heat treatment,

(4) there is a nearly isomorphous series of mixed thallium/ lead oxides that

should grow epitaxially, and (5) the device applications of these types of

materials have not been studied previously.

We expect that these materials will prove to have very interesting

optical and electrical properties. The end members of the series, Pb02 and

TI2O3, are both degenerate n-type semiconductors. They have the high

electrical conductivity of metals, with the optical properties of semiconductors. Because of their high majority carrier concentrations ( > 10^^/cm^),

they also have high reflectivity in the near-IR. Thallium(III) oxide, for

instance, is a degenerate n-type semiconductor with a bandgap of 1.4 eV and

a room temperature resistivity of only 70 ^ohm-cm [14]. Lead(IV) oxide

has a larger bandgap of approximately 1.8 eV [10]. These materials are, therefore, the metal oxide analogs of semiconductors such as GaAs and

AlGaAs, and they may function as multiple quantum wells when the

modulation wavelength is in the nanometer range [15,16].

186
EXPERIMENTAL

ElectrocJiemical depositions were performed using either a Princeton Applied Research (PAR) potentiostat/galvanostat, or a system consisting of a Stonehart BC 1200 potentiostat/galvanostat, PAR Model 175

universal programmer, and PAR Model 379 digital coulometer. Superlattices were deposited onto polycrystalline 430 stainless steel disks. The disks

were housed in a PAR Model K105 flat specimen holder, so that only the

front surface was exposed to the solution. The final polish on the stainless

was 0.05 /im alumina. Films were deposited from aqueous solutions of

0.005 M TINO3 and 0.1 M Pb(N03)2 in 5 M NaOH. The strong base is necessary to dissolve the Pb(N03)2, which precipitates in 1 M NaOH as PbO. The superlattices were deposited by galvanostatically pulsing between

either 5 mA or 1 mA and a lower current such as 0.05 mA. Each layer

thickness was proportional to the product of current and time. Elemental

analysis was done by EDS on a JEOL 35CF scanning electron microscope.

The L Hues of thallium and lead were used for the analysis because of the

overlap of the K lines.

RESULTS AND DISCUSSION

Since TI2O3 and Pb02 can both be electrodeposited from alkaline

solution, and there is considerable overlap of the deposition iV curves, it is possible to electrodeposit Tl^Pb^O^ films with compositions that are a function of applied potential or current [1]. We have carefully chosen the

reactant concentrations so that it is possible to vary the Pb/Tl ratio over a wide range. Since the TINO3 concentration is only 0.005 M, and the

Pb(N03)2 concentration is 0.1 M, the deposition becomes mass-transportlimited in [T1 + ] when the current is raised over about 1 mA. The Pb/Tl

ratio was varied from 0.84 to 7.3, when the current was increased from 0.05

mA to 5.0 mA (see Table I).

The x-ray diffraction patterns of the Tl^PbjjOj. films are very similar

to the TI2O3 pattern. An important difference between the patterns is that

mixed-index reflections such as (211) are not observed for the mixed oxide,

but they are present for TI2O3. The systematic absence of mixed index

reflections is consistent with an fee structure. This is in agreement with the

assignment of Sakai et al. [12], of electrodeposited Pb8Tl5024 as an fee

fluorite-type structure with a cubic lattice parameter of 0.53331 nm. For

comparison, TI2O3 has a bcc bixbyite structure with a = 1.05434 nm, and

Pb02 has either an orthorhombic or tetragonal structure [17]. The lattice

parameters that we measure for our mixed oxides vary by only 0.3% over the

187
Table I. Composition and values of the cubic lattice parameters for

electrodeposited T[^F\0^ films as a function of applied current

(fromref. [1]).

Applied Current

(mA)

Measured

Potential

(Vvs. SCE)

Pb/Tl

mole ratio

Cubic Lattice

Parameter

(nm)

0.05 0.06 0.84 0.536

0.5 0.12 1.9 0.535

1.0 0.14 2.5 0.537

5.0 0.27 7.3 0.535

entire composition range listed in Table I. Hence, these oxides are ideal for

the epitaxial growth of nanomodulated superlattices.

A modulated structure was prepared for SEM studies by alternately

pulsing the electrode at 0.05 mA for 8772 seconds and at 5.0 mA for 88

seconds. The layers were intentionally made quite thick, so that they could

be easily imaged in the SEM. Cross sections of a fractured film are shown

in Figure 2 for two different magnifications. The lower magnification

micrograph (top) shows the strong preferred orientation of the film. We

have previously reported similar columnar growth for electrodeposited

thallium(III) oxide films. The higher magnification micrograph (bottom)

clearly shows that the composition is modulated within each columnar grain.

Samples with much shorter modulation wavelengths and layer

thicknesses were prepared by using shorter dwell times during the doublepulse galvanostatic deposition. The superlattices with modulation wavelengths in the nanometer range were characterized by x-ray diffraction. The

periodicity of the superlattice manifests itself in two ways in the x-ray

diffraction pattern. Bragg reflections are observed at low angles which

correspond to the modulation wavelengths. These peaks at low angles

would be observed even with a modulated amorphous material. At high

angles the periodicity is seen as satellites around the main Bragg reflections

for the material. The analysis of our superlattices was done at high angles.

The modulation wavelength, a, can be calculated from Eq. (1), where > is the x-ray wavelength used, L is the order of the reflection, and d is the

diffraction angle. In Eq. (1), L = 0 for the Bragg reflection, while the first

188
Figure 2. Scanning electron micrographs of modulated TlaPb^Oj./

TljPbgOf films. Micrographs are low (top) and higher (bottom)

magnifications of a cross-sectional view of a fractured film. Film

was deposited by cycling electrode between 0.05 mA and 5 mA

in a solution of 0.005 M TINO3 and 0.1 M Pb(N03)2 in 5 M

NaOH. Layer thicknesses are intentionally quite thick (-0.5

/im) so that SEM imaging is possible.

189
satellite at low angle has the value L = -1, and the first satellite at higher

angle has the value L = +1.

(Li-L2)X

A = (1)

2 (sind

I

- sine 2)

The modulation wavelengths for several superlattices are listed in

Table II. The x-ray modulation wavelength, Ap, was calculated from

Faraday's law. For the Faraday's law calculation, we used p = 10.353 g/cm^,

n = 2, and M = 228 g/mol. This is only an approximation since the density

and formula weight varies somewhat with composition. We have found,

however, that for all of the compositions in Table I, approximately 1.1 to 1.2

)um of material are deposited per coulomb of anodic charge that is passed.

There is remarkably good agreement between Ap and A^ in Table II considering the approximation used for Ap. Superlattice satellites were not

observed for modulation wavelengths less than 5.9 mn.

Table 11. Comparison of modulation wavelengths calculated from

Faraday's law (Ap) and from x-ray satellite spacings (a^) for

ceramic superlattices deposited at various currents and dwell

times. The electrode area ranged from 0.78 to 0.83 cm^.

(mA)

ta

(s) (mA)

tb

(s)

Ap

(nm)

Ax

(nm)

0.05 40 5.0 0.4 5.9 5.9

0.05 50 5.0 0.5 7.4 7.4

0.05 60 5.0 0.6 8.3 8.0

0.05 70 5.0 0.7 9.7 9.9

0.05 80 5.0 0.8 11.5 11.6

0.05 90 5.0 0.9 12.3 12.9

FUTURE WORK

The effect of quantum confinement on the optical and electrical

properties of these materials is unknown. Quantum confinement should

produce thickness-dependent blue-shifting of the optical absorption spectra.

190
and the development of excitonic features. We are presently studying the

optical and electrical properties of these nanomodulated materials.

Another important area of future research will be to study epitaxial growth

in these systems by measuring nucleation and growth rate constants by

transient potential-step experiments for varying degrees of lattice mismatch.

These in situ studies should provide valuable information to architects of

future superlattice structures.

ACKNOWLEDGEMENTS

The author thanks Mitsubishi Corporation for financial support

through a Mitsubishi Kasei Faculty Development Award, and Unocal

Corporation for donation of all of the electrochemical instrumentation used

in this research.

REFERENCES

1. J. A. Switzer, M. J. Shane, and R. J. Phillips, Science, 247, 444 (1990).

2. U. Cohen, F. B. Koch, and R. Sard, J. Electrochem. Soc, 130, 1987

(1983).

3. D. Tench and J. White, Metall. Trans. A, 15, 2039 (1984).

4. J. Yahalom et al., J. Mater. Res., 4, 755 (1989).

5. D. S. Lashmore and M. P. Daniel, J. Electrochem. Soc, 135. 1218

(1988).

6. L. H. Bennett et al, J. Magn. Mater., 67, 239 (1987).

7. R. J. Phillips, M. J. Shane, and J. A. Switzer, J. Mater. Res., 4, 923

(1989).

8. J. A. Switzer, Am. Ceram. Soc. Bull., 66, 1521 (1987).

9. J. A. Switzer, J. Electrochem. Soc, 133, 722 ( 1986).

10. W. Mindt, J. Electrochem. Soc, 116, 1076 (1969).

11. J. C. G. Thomas and D. W. Wabner, J. Electroanal. Chem., 135, 243

(1982).

191
12. M. Sakai, T. Sekine, and Y. Yamazaki, J. Electrochem, Soc, 130. 1631 (1983).

13. W. Tillmetz and D. W. Wabner, Z. Naturforsch. Teil B, 39, 594

(1984).

14. V. N. Shukla and G. P. Wirtz, J. Am. Ceram. Soc, 60, 253 (1977).

15. C. R. Leavens and R. Taylor, Eds., "Interfaces, Quantum Wells, and

Superlattices," (NATO Series B: Physics, Vol. 179) (Plenum, New

York, 1988).

16. E. E. Mendez and K. von Klitzing, Eds., "Physics and Applications of

Quantum Wells and Superlattices," (NATO Series B: Physics, Vol.

170) (Plenum, New York, 1987).

17. P. T. Moseley, J. L. Hutchison, and M. A. M. Bourke, J. Electrochem. Soc, 129, 876 (1982).

192
III. SOLID-STATE SYNTHESIS

AND CHARACTERIZATION
i
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

PROPERTIES OF SOME MIXED URANIUM OXIDES

P.G. Dickens, A.V. Powell and G.P, Stxittard

University of Oxford, Inorganic Chemistry Laboratory,

South Parks Rd. Oxford U.K.

ABSTRACT

U^Og and stucturally related mixed uranium/transition metal

oxides UMO5 (M=Ti, V) are examined as hosts for intercalation

reactions. The location of lithium in LIq QU-^Og has been

determined by neutron diffraction. Changes in the electronic

properties of UVO^ on oxidation and of UTiO^ on intercalation

have been investigated by XPS and magnetic susceptibility

measurements.

INTRODUCTION

a-U-jOg is the structural archetype for a number of mixed

uraniiam/transition-metal oxides such as UVO^, UTiO^. CuU-^O^q

and UNb30]^Q, which have in common pillared-layer structures

consisting of planes of edge-sharing UO^^ and MOj^^ polyhedra

connected by metal-oxygen chains running perpendicularly to the

planes and providing an interplanar separation of ~4 A. a-UgOg

itself, (Figs 1(a) and 4) consists of edge-sharing UO^ pentagons

connected by U-O-U chains [1]; U-O bond orders in the chain

direction are higher than within the layers although no discrete

uranyl groups exist. The primitive unit-cell (dashed outline) has

pseudo-hexagonal symmetry (y= 119.3°). However, powder neutron

diffraction [2] indicates that above 210°C very slight atomic

displacements occxxr to give a phase, denoted a'-U^Og, which has

hexagonal syrmnetry. This phase reverts to the orthorhombic form

on cooling. The relatively open structure of cc-U-^Og provides a

host matrix for intercalation reactions with Li and Na which we

have described previously [3, 4].

195
Fig 1: (a) a-UgOg (b) UVO5

The structural relationship between a-U^Og and UVO5 is

shown schematically in Figs 1(a) and (b). In UVO^ edge-sharing

pentagons which form zig-zag chains are separated by single VO^

square-based pyramids forming layers of composition (UV04)j^ [5]

which in turn are linked by U-O-U chains in a similar way to those

in a-U30g. The V-O linkage to the oxygen at the vertex of the

VOiT square-based pyramid is very short (1.64 A) and gives rise to

BR absorption at 980 cm" characteristic of a vanadyl group. Again,

the U-O-U chains perpendicular to the UfV/O layers contain no

discrete uranyl groups. Oxygen coordination determined by powder

neutron diffraction is illustrated in Fig 2.

196
So far, the crystal structure of UTiO^ is not known in detail

but comparisons of powder X-ray data strongly suggest that it and

XJVO5 have similar planar metal atom arrangements which are probably also common to those of UMoO^, UNbO^ and UTaO^.

All appear to undergo similar intercalation reactions with alkali

metals leading to the formation of mixed-vaience compounds. In

this work we describe the synthesis and characterisation of some

new insertion compounds formed by U30g, UVO^ and UTiO^

together with an attempt to determine their electronic structures.

EXPERIMENTAL

a-U^Og was prepared, as previously described, by the thermal

decomposition of U04«2H20 at 720°C for 2 days. UVO^ was

prepared by conventional solid-state synthesis by reaction of ^20^,

V2O3 and U30g in a sealed silica tube at 700°C for 24 hours

whereas TJTiO^ was synthesised by heating a uranium-vanadium

co-precipitate of hydrated oxides in a modification of the method

described by Hoekstra and Marshall [7].

Insertion of both alkali metals and magnesium into the above

three oxides was performed chemically and electrochemically (in

aqueous and non-aqueous media). Chemical lithiation was

accomplished by use of lithium iodide in CCI4 in a manner that has

been described previously [8]. In a similar maimer, use of a

magnesium iodide monoetherate complex in CCI4 permitted the

insertion of magnesium at ambient temperature. Electrochemical

measurements in aqueous media were performed using a threecompartment glass cell in conjunction with a Princeton applied

research potentiostat operating in a galvanostatic mode. A platinum

coxmter-electrode was used whilst a calomel reference allowed all

voltages to be referred to the standard hydrogen scale. The cathode

was a mixture of the metal oxide, graphite and PTFE mixed in the

ratio (by weight) 80:10:10 and pressed onto a 13mm platinum

gauze disk. Electrolytes were IM LiCL or 2M MgCl2 solutions for

the insertion of lithium or magnesium respectively. Non-aqueous

electrochemical discharges were performed in an argon atmosphere

glove-box using a Teflon cell. The cathode was as described above.

The electrolyte was a solution of LiBF^ in propylene carbonate and

the anode a disc of lithium metal.

All products of insertion were handled in an argon

atmosphere dry box. Characterisation was by powder X-ray

diffraction, infra-red spectroscopy and TGA. Product compositions

were determined by reducing power titration [4]. XPS measuremnts

197
were performed using a VG ESCALAB spectrometer. An Oxford

Instruments Faraday balance was used to measure low temperature

magnetic susceptibilities and measurements at room-temperature

were made using a Gouy balance.

RESULTS AJSTD DISCUSSION

(i) Insertion Reactions

(a) U3O8

The insertion reactions undergone by a-U30g are summarised

in Fig 3. In all cases only small changes in lattice parameters

relative to those of the parent oxides are observed (Table 1),

strongly suggesting that insertion proceeds with retention of

structure.

U3O8 + Mgo.25Uo.75O2

(fluorite)

LixU308 (X ^ 2)

600'C

MgxU308

(X <, 0.5)

Li+, e- solv

Mgl2«Et20

Li0.9U3O8

(hex)

300*C

or Mg2+, 2e-, aq

U3O8

LiI/CCl4

Li0.9U3O8

(ortho)

Li+. e- aq

LixU308 (X ^ 0.9)

Fig 3: Insertion Reactions of a-U^Og

Table 1 : Lattice Parameters of Some Mixed Oxides and Insertion

Compoimds.

a/A b/A c/A

a-U30g 6.715 11.968 4.143

6.727 11.953 4.188

a -UgOg 6.812 6.812 4.142 (Y = 120°) ^ -^^0.88^303 6.849 6.849 4.197 (Y = 120°)

^§0.2^3^8 6.717 11.947 4.143

UVO5 7.207 12.364 4.123

Lii 1UVO5 7.28 6.22 4.14

UTi05 7.364 6.317 15.631 (P = 91°)

Lio9UTi05 7.412 6.241 16.034 (P = 91°)

Mgo.l5U^iC>5 7.387 6.270 15.667 (P = 91°)

198
In the case of LIq ggU^Og, a recent powder neutron

diffraction study [9] has located the lithium positions. For

material synthesised at ambient temperature the lithiums randomly

occupy interlayer sites (marked A, B, C in Fig 4) in the

orthorhombic a-UgOg structure. These sites are of approximately

trigonal bipyramidal coordination. However, on annealing in

vacuxmi at —573 K a transformation to the closely related hexagonal

form occurs. This structure is stable at all temperatures and does

not revert to the orthorhombic form on cooling. Refinement of

neutron data at 595 K demonstrated that the lithiums now

exclusively occupy sites of type A (2d in P62m). That this is a

stable high-temperature phase, with only a limited range of

stoichiometry,'Was demonstrated by its synthesis from the oxides,

Li20, U30g and UO2 in a sealed tube preparation at SOO'C. This

compound a'-Lij^U^Og (0.8< x< 0.9) appears to be identical with the

hexagonal phase reported by Kemmler-Sack and Riidorff [10]

obtained by high temperature synthesis and formulated as

(Li20)j^*U02*U03 which for x = 0.33 corresponds to LiU^Og a

composition which is achieved by half-occupation of the sites

marked 'A* in Fig 4. . o

UOU Chain

Oxygen

Fig 4: Sites for Insertion in a-U30g (z = 1/2)

Corresponding magnesium insertion compounds Mg^^U^Og

were prepared by reaction with Mgl2*Et20 (x< 0.5) or by

galvsmostatic discharge from aqueous media. The retention of

structtire of the orthorhombic a-U-^Og framework is apparent from

199
the data in Table 1. The thermal behavioiar of this phase is quite

different from that of LIq 88^3^8' heating in a sealed tube at

600''C, MgQ 2^3^8 disproportionates into U30g and the stable

cubic fluorite phase MgyU]^_y02 y- 0.33), which has been

described by Kemimler-Sack and Rvidorff [11], according to; Mgo.2U308 = 0.8 UgOg + 0.8 MgQ 25^0.75^2

(b) UTi05 and UVO5

The insertion reactions undergone by UTiO^ are summarised

in Fig 5. UTiO^ is a versatile host lattice landergoing intercalation

with the chemical reagents Lil, n-BuLi and Mgl2*Et20 and also

electrochemically in both aqueous and non-aqueous media . UVO5

also readily intercalates lithium with retention of structure.

LixUTiOs (x <. 0.5) LixUTiOs (x ^ 1)

Mgx(H20)nUTi05 (x ^ 0.3) < MgxUTiOs (x :S 0.25)

Fig 5 : Insertion Reactions of UTiO^

(ii) Electronic Properties

By analogy with other oxide insertion compounds it is

presumed that formation of A^^UMO^ involves transfer of valence

electrons from A to metal-based orbitals in UMO^.

A = Aj* + e'u or M

leading to consequential changes in electronic properties such as colour, conductivity, PE spectra and magnetism. The products

described above are not metallic and their properties are compatible

with localisation of the electrons at particular metal sites (small

polaron model). A qualitiative le~ energy level diagram can be

constructed for UMO^ (Fig 6) which illustrates some important

features of its electronic structure. In this diagram the 13 a

bonding orbitals derive from overlap between sp^ O hybrids with U

200
13 X a

.\ 2 X 7C - * 2 X JC

U

7 x5f

M {

7 X g (d^sp^)

6 X g (d?'sp3)

» A 1 \

W \ V

7 x5f(lJ)

.3 X 7CConducdon

'band'

_ 3 _ 4 X JCj. - 4 X g^

- 9 X g„

.2 X 7C' 2x71

p = O 'in plane* orbital

c = O 'in chain' orbitaJ

Fig 6: le" Energy-level Diagram for UMO5

201
(d^sp-^) and M(d sp^) sets, the A-tz bonding orbitals from M^j^-Op^

and U^^-Op^ chain interaction, and the Sn non-bonding (or weakly

bonding) orbitals from in-plane interactions. This places a nonbonding U(5f) set or a Md^y anti-bonding orbital above the filled

valence set as possible constituents of the conduction band. In

UTiO^ 40e" completely fill the valence band. Stoichiometric

XJTiO^ is an insulator. In UVO5, le" per UVO5 unit occupies either

the V3d^y level (V(IV)) or the uranium 5f set, U(V) and in

Li^UTiO^ xe" must occupy either U(V) or Ti(m).

In order to distinguish experimentally between these

alternatives some preliminary studies have been made using XPS

and magnetic susceptibility measurements on UTiO^, Li^UTiO^,

UVO^ and a fully oxidised compound TJVO^ ^ (^2^2^1 1^'

(a) XPS

Measurements were made by J. Kemp, S. Hawe and P. A. Cox

of this laboratory and a complete account will be published

elsewhere. Fig 7 shows the valence region of the XPS spectra for

the samples studied. The filling of the conduction band on passing

from UTi05 to Li^ 9X77105 and from the fully oxidised UVO5 5 to

UVO5 is clearly demonstrated. On oxidation, the conduction band

is removed or considerably reduced in intensity due to the loss of

mobile electrons .

T

, , 1

1 1 — T"

0 10 20 30 0 10 20 30

Binding Energy eV Binding Energy eV

Fig 7: Valence Region of XPS

202
In Table 2 are shown the binding energies of the core levels of

U4f-7^2' '^^P3/2 ^^P3/2 differences between relevant

pairs in the compounds studied.

Table 2: Binding Energies and Differences (eV)

Compound TiZp^p^ ^^^7/2

Lio9UTi05 459.7 382.0 77.7

UTi05 459.9 382.1 77.4

V2p3/2 U4f7/2

^

UVO5 517.1 380.5 136.6

^^^^5.5 526.9 390.8 136.1

Charging effects in insulating samples, in the absence of an

absolute binding energy reference, mean that only the differences

AE in the peak positions of the relevant transition metal/uranium

pairs are of significance. Oxidation raises the formal oxidation

state of one or other of the metals and so leads to a shift to

higher binding energy of the core line involved. The transition

metal 2p2/2 falls at a higher energy than does U4f'7y2 so that a

decrease in AE on oxidation implies an increase of the uranium

oxidation state. This is the observed direction of change for both

the pairs UVO5, UVO5 5 and LIq qUTIO^ and UTiO^ and so it is

concluded that in both it is the uranium which is principally

reduced leading to the formulation of Li^UTiO^ as Li^^U^^ Uj_ ^VI^-IVq^ and UVO5 as U^V^Og. Further support for this

conclusion is provided by the appearance in the reduced phases of

satellite lines at binding energies 7-8 eV above the U4f'7^2 core

line which is a characteristic feature of [12]. Such lines were

absent in UVO5 5 and UTi05.

(b) Magnetic Measurements

Magnetic susceptibilities of UTiO^, LIq 9^^^^5 and

MgQ 25UTi05 were measured at 293 K and effective Bohr

magneton numbers (per U^) evaluated by using nig££ = 2.82[(Xj^ - Xjn(UTi05))T]l/^. The values obtained; 0.93(Li UTi05) and

l.OSCMgj^UTiO^) are close to those found for other U compounds

[13] such as 1.06(LiUO3) and 1 . l(lJNb30iQ). They are much

lower values than expected for a magnetically dilute first-row

transition metal complex and support the contention that the

insertion compounds contain reduced uranium rather than Ti^^^

203
ACKNOWLEDGEMENTS

The assistance of J.M. Fleetwood in the preparation of samples is

gratefully acknowledged. AVP thanks the CEGB for a studentship.

REFERENCES

[I] B.O. Loopstra; Acta. Cryst. j/Z, 651, (1964).

[2] B.O. Loopstra; J. Appl. Cryst. 3, 94, (1970).

[3] P.O. Dickens, D.J. Penny and M.T. WeUer; Solid State Ionics,

18&19, 778, (1986).

[4] P.O. Dickens, A.V. Powell and A.M. Chippindale; Solid State

Ionics, 28-30, 1123, (1988).

[5] R. Chevalier and M. Gasperin; Bull. Soc. Franc. Minera. Crist.

93, 18, (1970).

[6] P.G. Dickens, R.G.J. Ball, S. Patat and S.Hull; In Preparation.

[7] H.R. Hoekstra and R.H. Marshall; J. Inorg. Nucl. Chem. 27,

1947, (1965).

[8] P.G. Dickens, S.J. French. A.T. Hight and M.F. Pye; Mat. Res.

BuU. 14, 1295, (1979).

[9] P.G. Dickens and A.V. Powell; Submitted to J. Solid State

Chem.

[10] S. Kemmler-Sack and W. Riidorff; Z. Anorg. Allg. Chem. 354,

255, (1967).

[I I] S. Kemmler-Sack 2ind W. Riidorff; Z. Anorg. Allg. Chem. 354, 264, (1967).

[12] C. Miyake, H. Sakurai and S. Imoto; Chem. Phys. Lett. 36,

158, (1975).

[13] B. Kanellakopulos, E. Henrich, C. Keller, F. Baumgaitner, E.

Konig and V.P. Desai; Chemical Physics, 53, 197, (1980).

204
DISCUSSION

M. Greenblatt: Have you tried any other group II elements besides magnesium?

P. Dickens: No. The only other metal we've investigated is sodium and, of course,

hydrogen.

M. Greenblatt: Sodium can be inserted as well?

P. Dickens : Yes but slowly. I should have pointed out that lithium in these compounds

is fairly mobile, it has a high diffusion coefficient. Sodium is much less mobile and other

metals, if they do insert, are unlikely to be very mobile.

R. Newnham: Will protons go in those positions?

P. Dickens: No, they don't go in the lithium positions, but actually attach themselves as OH groups. In all the oxide insertion compounds we've studied we find that hydrogen

is located as OH or OH2 groups. So if you prepare Hj^UTi05 the hydrogen appears as OH, but an electron, of course, is transferred to uranium to give a product isoelectronic

with Lij^UTi05.

P. Davies : Are you able to insert protons into these structures?

P. Dickens: Yes, at least for U3O8, UTiOj and UVO5. For example, for UVO5

something like HUVO5 is formed quite easily with retention of structure, but there comes

a point where it begins to become amorphous. In the particular case of UVO5, you have

3-connected oxygen and also short terminal vanadyl groups, and when the latter group

begins to be attacked, the metal oxygen distance is changed quite dramatically and the

parent structure is disrupted, leading to amorphization.

P. Davies : What about if you do subsequent low temperature heat treatment on a protonated compound, do you get any interesting changes?

P. Dickens : We haven't done that. We have just started doing ion exchange on some

of these phases. One attraction of these materials is that they're nearly totally insoluble,

particularly the niobium ones, so you can carry out reaction in aqueous media without

risk of dissolution.

M. Greenblatt: The lithium content in Lip 88U3O8 is a rather peculiar value, isn't it?

P. Dickens : Well, at Li^ oUaOg, the A sites would just be half filled. In fact, both

chemical and electrochemical preparations suggest that the phase limit for this phase lies at an X value in Lij^U30g below this and a completely ordered phase of Lij g^sOg is not

achieved. In the same way in Na^WOj the structural limit for the cubic phase Naj QWO3

is not quite reached under normal synthetic conditions.

205
M. Greenblatt: I see.

W. McCarroll : Can you use bond-strength, bond-length rules for checking oxidation

states in these compounds?

P. Dickens: Yes. In the case of the bond lengths obtained in our recent neutron

diffraction study of UVO5, bond strength sums are consistent with U^V^Oj.

206
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

INVESTIGATIONS ON THE STRUCTURAL, ELECTRICAL AND MAGNETIC

PROPERTIES OF Nd2-xMxNi04+5, M = Ca2+ AND Ba2+

Brian W. Arbuckle, Zhen Zhang, and Martha Greenblatt*

Department of Chemistry, Rutgers the State University of New Jersey

Piscataway, NJ 08855-0939

*Author to whom correspondence should be addressed.

INTRODUCTION

Ni

The Ln2Ni04+5 type compounds (Ln = La, Pr, Nd) are known to possess

K2NiF4-type structure, which can be described as alternating perovskite LnNi03

and rocksait LnO layers stacked along the crystallographic c-axis (Figure 1). The

Ni-O octahedral units form a planar array of Ni-O-Ni bonds in the ab plane. The

distance between the perovskite-layers along the c direction is double the distance

between the nickel ion nearest neighbors and therefore, the magnetic and electrical

properties of these compounds are highly anisotropic and usually confined within

the Ni-O basal plane (1-5). The influence of oxygen stoichiometry on structural,

electrical transport, and magnetic properties have been found to be considerable

(4,6,7). The excess oxygen has been attributed to the deviations in the metal-ion

stoichiometry (8), variable formal valence of the

transition metal (9), intergrowths of RuddelesdenPopper type phases (10), and the incorporation of

interstitial oxygen defects (11). Nonstoichiometric

La2Ni04+6 is tetragonal at ambient temperatures and

shows no long range ordering, but when reduced to nearly stoichiometric La2Ni04 it begins to exhibit

orthorhombic distortion and long range magnetic

ordering (2,4,6,12). Electrical conductivity studies

on La2Ni04+6 show a semiconductor-to-metal

transition between 500-620K (3,13). Orthorhombic

Nd2Ni04-(.5 and Pr2Ni04+6 undergo a phase

transition to tetragonal symmetry at 724K and 834K

respectively (14). Tetragonal La2Ni04+5 does not

undergo a phase transition up to 1073K. By

increasing the 6 value in Nd2Ni04+5 and Pr2Ni04+8

the long range magnetic ordering is suppressed,

which is attributed to the decrease in structural

distortion that arises with increasing 5 (4).

Ln

Figure I. Tetragonal K2NiF4 type

structure, Ln = La. Nd. or Pr

207
Substitution of Sr2+ into the Nd^"*" sites of Nd2Ni04+5 induces a phase

transition from orthorhombic to tetragonal symmetry (15). The orthorhombic

samples of the solid solution series Nd2-xSrxNi04+§ have large 5 values whereas

the tetragonal members (x > 0.2) are close to being stoichiometric in oxygen

(6 = 0). Samples with all values of x were semiconducting from room temperature

down to lOK except NdSrNi04, which was metalUc with a metal-to- semiconductor

transition at 190K. The room temperature resistivity of the samples decreases as a function of increasing x. The magnetic susceptibihty showed Curie-Weiss behavior

with long range magnetic ordering (Tn ~ 17K) for the intermediate members

(0.4 < x < 0.6). We have incorporated other group II divalent cations into

Nd2-xM xNi04 + 5 (M = Ca2+, Sr2+, Ba2+) with different sizes and

electronegativities to study these effects on the physical properties.

EXPERIMENTAL

Solid solutions of Nd2-xCaxNi04+5 (0 < x < 0.6) and Nd2-xBaxNi04+5

(0 < X < 0.5) were prepared by co-decomposition of the metal nitrates.

Stoichiometric quantities of Nd203 (Aldrich, 99.9%), NiO (Aldrich, 99.99%),

Ba(N03)2 (Aldrich, 99+%), and CaCOs (Aldrich, 99.9%) were mixed together and

dissolved in concentrated nitric acid. The excess nitric acid was removed by

gradual heating and the resulting green solid was calcined in air at 800° C. The

products were ground and heated in air at 1 150°C for 32h in alumina crucibles with

two intermittent grindings. The samples were cooled to 800°C in the furnace before

being quenched to room temperature. The products were ground and cold pressed

into pellets. The pellets were fired in air at 1150°C for 16h and slow cooled to room temperature.

Powder X-ray diffraction (XRD) data, collected on a SCINTAG PAD V

diffractometer with monochromatized CuKa radiation were used to monitor the

completion of the reaction. A least-squares refinement of the observed powder

diffraction data was used in the evaluation of cell parameters. The electrical

resistivity of samples was measured by the standard four-probe technique in a DE202 cryostat (APD Cryogenics) from room temperature to 15K. A SQUID

magnetometer (Quantum Design) was used to measure the magnetic susceptibility

of samples from 2K to 370K with an applied magnetic field of lOOOG. The total oxygen content was determined by hydrogen reduction using a DuPont 951

Thermogravimetric Analyzer and by iodometric titration. For TGA analysis, the

material was heated on a platinum pan in flowing 5% H2 in helium gas (O.lL/min)

to 900°C at 15°C/min. For iodometric titration a 125 mg sample, after being dried

at 100°C for 2 h, was dissolved in 25 mL of 0.4 mol/L aqueous HCl and 25mL of

0.15 mol/L aqueous KI solution. While under an inert atmosphere, the solution

was titrated with a standardized thiosulfate solution. A few drops of a 5% starch

solution were added just before the equivalence point in order to aid in the titration.

208
RESULTS AND DISCUSSION

5.45

•S5.40

S3S

(A)

ia(A)

ft (A)

? II

i

* {

§ 5 5 5 §

(B)

?5

Structural : Single phase specimens in the Nd2-xMxNi04+5 (M = Ca2 +

(0 < X < 0.6) and Ba2+(0 < x < 0.5)) solid solutions were prepared as evident by

X-ray diffraction patterns. Lehmann and Muller-Buschbaum (16) have reported

that Nd2Ni04 + § crystallizes with

orthorhombically distorted K2NiF4-type

structure (space group Abma). Arbuckle et

al. (15) reported that for x < 0.2 in the series

Nd2-xSrxNi04+5 the samples crystallized

with orthorhombically distorted

K2NiF4-type structure at ambient

temperatures and exhibited a phase transition to pseudo-tetragonal symmetry for x > 0.2.

The orthorhombic distortion manifests itself

in the splitting of some of the hkO type

reflections in the XRD pattern. The

tetragonal samples (x > 0.2), showed no

apparent splitting in the hkO reflections. The

The degree of divergence in the hkO

reflections of the orthorhombic specimens •§

served as an indicator of the degree of

distortion from tetragonal symmetry. All the

samples in the series Nd2-xCaxNi04+5

except for (0.2 < x < 0.3), appear to possess orthorhombic symmetry which is evident by the splitting of some of the hkO

reflections in the X-ray powder diffraction

pattern. The XRD pattern for 0.2 < x < 0.3,

showed no spHtting of the hkO reflections in

the range 10° < 29 < 60°. The incorporation

of small amounts of Sr2+ into the

orthorhombic samples transformed the

sample from orthorhombic to pseudotetragonal symmetry. The cell parameters as determined from least squares refinement of

the XRD data are presented in Table 1. The

changes in the unit cell parameters a, b, and

c as a function of x are depicted

in Figures 2A and 2B. As the Ca2+ content

increases in Nd2-xCaxNi04+5, the a and b

cell parameters converge to form a pseudotetragonal unit cell at x ~ 0.20 (Figure 2A).

12J6

1234

^ 12J2

1230

12.28

360

I

355 >

350

345

(C)

5?

^ ^ 5 S

0.1 0.3 0.5 0.7 0.9 1.1

Figure 2. A) Change in the lattice parameters a and 6 as a function of Ca^"*" content in Nd2.xCaxNi04+5, B) Change in the lattice parameter c as a function of Ca^*

content in Nd2-xCaxNi04+5, C) Change in the unit cell volume as a function of Ca2+

content in Nd2-jiCaxNi04+5.

209
Even though the 0.2 < x < 0.3 samples can be indexed as either orthorhombic or

tetragonal, only the orthorhombic unit cell parameters are listed in Table 1 . The a and b cell parameters diverge for x > 0.30 to form what appears to be an

orthorhombic unit cell. There is very litde change in the unit cell parameters for x >

0.6, which may suggest that x = 0.6 is the limit of solubility of Ca^"*" in Nd2-

TABLE 1 Strucmral and Physical properties of Nd2-xCaxNi04+5

Transition

Temperamre

Compound a(A) b(A) (^(A) volume(A)3 CO

Nd2Ni04.i8 ±0.03 5 378(1) 5.441(1) 12 355(1) 361.52(7) 575

Nd1.9Cao.1NiO4.10lO.O3 5 368(1) 5.443(1) 12 333(2) 359.19(5) 240

Nd1.8Ca0.2NiO4.1210.03 5 385(1) 5.378(1) 12 322(3) 356.84(6) —

Nd1.75Cao.25NiO4.08iO.O3 5 370(1) 5.373(1) 12 298(3) 354.77(6) —

Nd1.7Cao.3NiO4.01to.03 5 370(1) 5.368(1) 12 300(1) 354.56(4)

Nd1.6Cao.4NiO3.98io.03 5 328(1) 5.388(1) 12 266(1) 352.10(3) 380

Nd1.5Cao.5NiO3.96io.03 5 327(1) 5.387(1) 12 270(1) 352.05(5) 320

Ndi.4Cao.6Ni03.97±o.03 5 307(1) 5.356(1) 12 281(1) 349.09(3) 180

Ndi.3Cao.7Ni03.97+o.03 5 309(1) 5.369(1) 12 276(1) 349.92(4) 220

Nd1.2Cao.8NiO3.8OlO.03 5 308(1) 5.351(1) 12 284(1) 348.83(4) 180

Ndi.iCao.9Ni03.79+o.03 5 310(1) 5.366(1) 12 280(1) 349.88(4) 160

NdCaNiO3.68iO.O3 5 307(1) 5.351(1) 12 284(1) 348.83(4) 120

xCaxNi04+5, and therefore, the samples are not single phase material but mixtures

of compounds. Small impurity peaks of starting material are visible in the X-ray

pattern for the 0.7 < x < 1.0 samples. The c cell parameter decreases for 0 < x <

0.4 and then increases slightly for x > 0.4, while the volume decreases for 0 < x <

0.6 and thereafter remains fairly constant for x > 0.6. This is expected even though

the Ca2+ (I.32A, CN = (IX)), has a slightly larger effective ionic radius than Nd3+

(1.303 A for CN = (IX)). The decrease in the size of nickel ions when going from

210
Ni2+ to Ni3+, appears to have a greater influence on the cell parameters than does

the size effect of Ca2+ in the range 0 < x < 0.6.

5.435

5.415

5J95

The Nd2-xBaxNi04+5 samples with x > 0.5 show large amounts of an

impurity phase(s) in the XRD pattern which have not been properly characterized so

far. For this reason only samples with x < 0.5 are reported in the present study.

As discussed above the extent of

splitting in the hkO reflections of the

orthorhombic specimens served as an

indicator of the degree of distortion

from the tetragonal symmetry. From

the XRD patterns it is clear that for x =

0.2, the splitting of the hkO reflections

in the range 5° < 29 < 60° are absent,

indicating that a phase transition from

orthorhombic to pseudo-tetragonal

symmetry has occured. Even though

the low angle reflections did not show

any divergence in the hkO reflections,

the high angle data showed peak

broadening in the corresponding

reflections indicating that the actual

symmetry could be lower than

tetragonal. The cell parameters as determined from the least squares

refinement of the XRD data are presented in Table 2. For samples with

x > 0.2, the unit cell parameters could

be indexed as orthorhombic or

tetragonal, only the orthorhombic cell

parameters are listed in Table 2. The

changes in the unit cell parameters a, b,

and c as a function of x are illustrated in

Figures 3A and 3B. It is evident from

Table 2 and Figure 3A that the

orthorhombic to pseudo-tetragonal

phase transition in Nd2-xBaxNi04+s

solid solution is achieved via a rapid

compression in the b and an expansion in the a unit cell parameter (Figure 3A). The

c cell parameter and volume increase monotonically as a function of increasing x. In contrast, the ai cell parameter (the at for the orthorhombic samples were

calculated using the expression at = {{a+b)/2^), decreases slightly as a function of

increasing x for the orthorhombic samples (x < 0.2), but decreases more

dramatically for the tetragonal samples (0.2 < x < 0.5). Arbuckle et al. (15) and

12.60

12.50

12.40

12.30

{ (A)

1

0 (A)

• ft (A)

{

I

I

!

1 1 i

I

\

\

0.0 0.1 0.2 0.3 0.4 0.5 0.

Figure 3. A) Change in tlie hiuicc parameters a and ^> as a function of Ca^^ content in Nd2-

xCaxNi04+8, B) Change in the lattice parameter c as a function of Ca^^ content in Nd2-xCaxNi04+5.

211
Gopalakrishnan et al. (17) have reported similar trends in the at and c cell

parameters for Nd2-xSrxNi04+6 and La2.xSrxNi04+§ respectively. They attributed

the increase in at and c to the increased concentration of low-spin d^ Ni3+ (t2g^eg^)

and the Jahn-Teller distortion of the single electron being localized in the dz2 orbital.

Unlike the latter two systems, Nd2-xBaxNi04+6 compounds show an increase in

the unit cell volume with increasing x. This behavior is also observed for La2-

xBaxNi04+s (5). If a Jahn-Teller distortion mechanism is operative here it is clear

from the rate of increase in volume that the size effects of Ba2+ play the dominant

role in the expansion of the unit cell parameters.

TABLE 2

Structural and Physical properties of Nd2-xBaxNi04+§

Transition

Temperature

Compound a(A) b(A) c(A) volume(A)3 (°C)

Nd2Ni04.i8 ±0.03 5.378(1) 5.441(1) 12.355(1) 361.52(7) 575

Ndi.95Bao.05Ni04.i5+o.03 5.378(1) 5.439(1) 12.381(2) 362.16(7) 515

Ndi.9Bao.iNi04.i3+0.03 5.389(1) 5.426(1) 12.409(1) 362.91(5) 225

Ndi.85Bao.i5Ni04.xx±0.03 5.393(1) 5.417(1) 12.431(1) 363.16(6) 120

Ndi.8Bao.2Ni04.o8±0.03 5.400(1) 5.403(1) 12.465(1) 363.71(4)

Nd1.7Bao.3NiO4.03tO.O3 5.397(1) 5.392(1) 12.548(1) 365.17(5)

Nd1.6Ba0.4NiO3.9910.03 5.392(1) 5.393(1) 12.588(2) 366.04(4)

Nd1.5Bao.5NiO3.9110.03 5.391(3) 5.388(1) 12.608(1) 366.22(3)

Tamura has reported a high temperature structural transition in Nd2Ni04+5

at 847K (18), corresponding to a structural phase transition from orthorhombic to tetragonal symmetry. Subsequent studies have confirmed this transition from high

temperature X-ray diffraction and DSC studies (14). Incorporation of Sr2+ into the

Nd2Ni04+s structure, stabilizes the high temperature phase and thus lowers the

structural transition temperature (Tt) (15). In order to understand the influence of

Ba2+ substitution on the nature of this phase transition we have carried out high

temperature X-ray diffraction and DTA studies from room temperature to 850°C in

controlled oxygen atmospheres for samples with x < 0.2. Incorporation of Ba2+ in

Nd2Ni04 resulted in a dramatic decrease in the Tt corresponding to the structural

transformation (Table 2). We see no evidence for a structural transition above room

212
temperature for the pseudo-tetragonal samples corresponding to 0.2 < x < 0.5. It is apparent that the high temperature tetragonal phase of Nd2Ni04+5 is stabilized by

the incorporation of Ba^"*", although the exact mechanism responsible for this

stabilization remains unclear at present. Incorporation of Ca^^ also decreases the Tt

for X < 0.1. For x > 0.4 some of the hkO reflections appear to converge which is an indication of a phase transition from orthorhombic to a pseudo-tetragonal

symmetry. The orthorhombic-to-tetragonal Tt decreases anomalously as x

increases (Table 1). The trend in the T^ is probably due to subtle structural changes

with increasing x.

Oxygen content analysis:

In order to study the variations in the overall oxygen content as a function of

X in all the samples, we have carried out TGA studies in hydrogen atmosphere and

iodometric titrations. The oxygen content as determined by TGA measurement,

was estimated for all samples assuming that the reduction products were nickel

metal and the appropriate oxides (which was confirmed by X-ray analysis). The

corresponding values are given in Table 1 for Nd2-xCaxNi04+5 and in Table 2 for

Nd2.xBaxNi04+s.

The stability of K2NiF4 structures (A2BO4 type compound) are often

described in terms of the tolerance factor defined as (19):

t ^A-Q

where r(A-O) ^(b-O) refer to the sum of ionic radii of A-O and B-O bond in nine

and six coordination respectively. The K2NiF4 structure is stable over the range

0.85 < t < 1; this is consistent with the absence of rare earth nickelates with rare earth ions smaller than neodymium. The value of t is very close to the lower limit

for Ln2Ni04 (Ln = La, Pr, and Nd), thus explaining the deviations from the ideal

tetragonal geometry. Monoclinic or orthorhombic distortion has been attributed to the low tolerance factor for Ln = Nd and Pr (20). It is possible to increase the value

of t by either increasing the effective radii of A ions (through the substitution of

larger divalent alkaline earth ions for Ln), or by partially oxidizing the B cations to higher valence states by suitable substitutions at both A and/or B sites. As already

mentioned, the Nd2-xSrxNi04+§ (15) compounds have orthorhombic symmetry for

X < 0.2 and low tolerance factors. For 0 < x < 0.2 the compounds gained stability

by incorporating excess oxygen in the structure thus oxidizing some of Ni2+ to Ni3+ and therefore, increasing the tolerance factor by decreasing r(B-O)- At x ~ 0.2 an orthorhombic to tetragonal phase transition was observed and the 5 value

decreased to zero, keeping the oxygen stoichiometry constant at around 4.0 for 0.2 < X < 1.0. Even though the 5 value remained constant at zero, the tolerance factor

for these compounds increased as x increased because of the size effect of Sr2+

(1.45A for CN = (IX)) versus Nd3+ (1.303A for CN = (IX)). In

213
Nd2-xCaxNi04+6, the Ca2+(1.32A for CN = (IX)) and Nd3+ ionic radii are very

close in size, therefore, the size effect of Ca^"*" should not play a significant role in

influencing the tolerance factor, since the r(A-O) value will remain relatively

unchanged as a function of x. If 5 remains close to zero, then rs-o will decrease

and t will increase with increasing x. Based on the oxygen analysis of the samples

(Table 1), the oxygen content of Nd2-xCaxNi04+5 is high for 0 < x < 0.2 and then

it drops significantly for 0.2 < x < 0.3. The 5 value does remain close to zero for

0.3 < X < 0.7 and then becomes negative for x > 0.7. The oxygen deficiency

determined for the x > 0.7 samples is probably due to the impurity in these

samples.

For Nd2-xBaxNi04+§, it is clear from Table 2 that the 6 values deviate from

the ideal value of zero for the orthorhombic phase, and decrease, and approach zero

for the tetragonal phases. As was seen for Nd2-xSrxNi04+s for 0 < x < 0.2 (15),

the stabilization of the K2NiF4 structure with orthorhombic distortion is achieved

by incorporating extra oxygen in the structure. The tetragonal structure is stable

with the higher concentration of Ba^"*" and Ni^"*" ions and does not require the extra

oxygen. The oxygen discrepancy for the x = 0.5 sample may indicate that the

impurity seen in the x > 0.5 samples may be present even though no impurity can

be detected in the XRD pattern of the x = 0.5 sample.

Electrical Transport behavior:

The room temperature resistivity in the solid solution Nd2-xCaxNi04+s

(0 < X < 0.6) increases as x increases (Figure 4A). This is in contrast to Nd2-xSrxNi04+5 (15) and La2-xBaxNi04+5 (18). For the Nd2-xBaxNi04+§ sample

where x = 0.1 (Figure 4B), the room temperature resistivity increases with respect

to the parent compound Nd2Ni04+§ and then decreases for 0.05 < x < 0.5. For the

X = 0.6 sample which has large amounts of uncharacterized impurities present, the

room temperature resistivity increases dramatically relative to the x = 0.5 sample.

Both Nd2-xCaxNi04+5 and Nd2-xBaxNi04+5 solid solution series are semiconducting down to lOK. The activation energies estimated firom the plots of

0.70

2 0.50

= 0.30

0.10

Figure 4. Room temperature electrical resistivity as a function of: A) Ca^"*" content in Nd2-xCaxNi04+s, and B) Ba'^"'" content in Nd2.xBaxNi04+§.

214
logp vs 1/T (125-300K) show a gradual decrease with increasing x across both

series. The decreasing slope of logp vs 1/T seen at low temperature is characteristic of an increasing dominance of extrinsic effects on the transport

properties. For Nd2-xBaxNi04+6 the decrease in the activation energies in the high

temperature region appear to indicate the increased delocalization of Ni^"*" electrons.

Magnetic Properties:

The room temperature magnetic susceptibility for Nd2-xCaxNi04+§ and

Nd2-xBaxNi04+s decrease with increasing x for both compounds; this can be

understood by the fact that paramagnetic Nd3+ ions are being progressively

substituted by diamagnetic M^"*" ions. The behavior of X300K for both sets of

compounds correlate well with the result reported for Nd2-xSrxNi04+5 (15). The

susceptibility data in the temperature range 220-370K can be fitted to the CurieWeiss equation:

X = C/(T+0)+xo

Where C and 0 are the Curie and Weiss constants respectively, and the term xo

includes the temperature independent contribution of the Pauli paramagnetic

susceptibility, core diamagnetism, and the Van Vleck susceptibility.

2.0e-4

* °

NdrgCa 2Ni04 12 Ndi 7Bi3Ni04 (i3 NdMCafiNiOj 04 Nd|6Ba4Ni03.99

l.Oe-4-

Magnetic susceptibility as a function of temperature in the range 2-370K of

representive samples are shown in Figure 5 for Nd2-xMxNi04+5 (M = Ca2+,

Ba^"*"). The susceptibility values were not corrected for the core diamagnetism of

component ions because of the large

paramagnetic contribution from the

Nd^"^ ions. Long range magnetic

ordering is observed in Nd2 - xMxNi04+5 for x > 0.4 with M = Ca2+

(Tn = 12K), and x = 0.4, 0.5 with M = Ba2+ (Tn = 12K). The long range

magnetic ordering observed for these

samples are probably due to the Nd

sublattice. No long range magnetic

ordering is observed for samples with 0 < x < 0.3 in either series. The absence

of long range magnetic ordering in the

orthorhombic samples with 5 > 0, is consistent with observations of Buttery

et al. (4), who reported that magnetic

order in Nd2Ni04+5 is suppressed

with large excess of oxygen in the

structure.

100 200

Temperature

300 400

(K)

Figure 5. Magnetic susceptibility versus temperature of Nd, gCa 2Ni04 1 2,

Nd17Ba.3NiO4.03, Nd14Ca.6NiO3.98, and

Nd,.6Ba4Ni03.99.

215
Acknowledgements

The authors would like to thank Dr. Shu Li and Dr. K. V. Ramanujachary

for helpful discussions and critical evaluation of the manuscript. The work was

supported by National Science Foundation Solid State Chemistry Grant DMR8714072.

REFERENCES

1. J.M. Bassat, P. Odier, and F. Gervais, Phys. Rev. B, 35, 7126 (1987).

2. D.J. Buttrey, J.M. Honig, and C.N.R. Rao, /. Solid State Chem., 64, 287

(1986).

3. C.N.R. Rao, D.J. Buttrey, N. Otsuka, P. Ganguly, H.R. Harrison, C.J. Sandberg, and J.M. Honig, / Solid State Chem., 51, 266 (1984).

4. D.J. Buttrey, J.M. Honig, /. Solid State Chem., 72, 38 (1988).

5. A.B. Austin, L.G. Carreiro, J.V. Marik, Mat. Res. Bull. 24, 639, (1989).

6. R. Saez-Puche, J.L. Rodriguez, and F. Fernandez, Inorg. Chim. Acta.,

140,151 (1987).

7. R. Saez-Puche, and F. Fernandez, Solid State Commun., 72, 273 (1989).

8. J.T. L^wandowski, R.A. Beyerlein, J.M. Longo, and R.A. McCauley, J. Am. Ceram. Soc., 69, 699 (1986).

9. J. Drennan, CP. Tavares, and B.C.H. Steele, Mat. Res. Bull., 17, 621

(1982).

10. P. Odier, Y. Nigara, and J. Coutures, /. Solid State Chem., 56, 32 (1985).

11. J.D. Jorgensen, B.Dabrowski, Shiyou Pei, D.R. Richards, and D.G. Hinks,

Phys. Rev. B, 40, 2187 (1989).

12. J. Rodriguez-Carvajal, J.L, Martinez, J. Pannetier, and R. Saez-Puche, Phys.

Rev. B 38, 7148 (1988).

13. P. Ganguly and C.N.R. Rao, Mat. Res. Bull., 8, 405 (1973).

14. R. Saez-Puche, R. Fernandez, and W.S. Glaunsinger, Mat. Sci. Monogr.,

(React. Solid Pt.A), 28A, 607, (1985).

15. B.W. Arbuckle, K.V. Ramanujachary, Zhen Zhang, and Martha Greenblatt,

J. Solid State Chem., In Press.

16. U. Lehmann and Hk. Miiller-Buschbaum, Z. Natiirforsch., B35, 389

(1980).

17. J. Gopalakrishnan, G. Colsmann, and B. Renter, /. Solid State Chem.,

22, 145 (1977).

18. S. Tamura, /. Phys. Soc. Jpn., 33, 574 (1972).

19. J. F. Ackerman, Mat. Res. Bull., 14, 487 (1979).

20. B. Wilier and M. Dairre, C.R. Acad. Sci. (Paris), C, 267, 1482 (1968).

216
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

Janiiary 1991.

SYNTHESIS AND CHARACTERIZATION OF Lai-xSrxTiOa (0<x<0.5).

Susan M. Kauzlarich,* Joseph E. Sunstrom, FV* and Peter Klavins.''"

Departments of Chemistry* and Physics+

University of California

Davis, California 95616

ABSTRACT

A series of compounds Lai-xSrxTiOs (0 < x < .5) have been prepared and

characterized by powder X-ray diffraction, microprobe analysis, thermal

gravimetric analysis and temperature dependent magnetic susceptibility. These

pseudotemary compounds have been synthesized by arc-melting the stoichiometric

combinations of SrTiOs and LaTiOs under argon. The compounds Lai-xSrxTi03

(0 < X < .3) crystallize in the orthorhombic space group Pbnm. The compositions

with X = .4, .5 can be indexed with a tetragonal cell. The synthesis and magnetic

properties of these compounds are discussed.

INTRODUCTION

There were only a few examples of oxide superconductors ^ prior to 1986

and the discovery of superconductivity in the Cu-O system.2 Oxide

superconductors that have been discovered to date contain metals that are d^, s^ or d^ ions. There have been several recent investigations of the early transition metal

oxides in the search for new high temperature superconductors. ^ We are interested

in the lanthanide titanates, LnTi03, ^ series of distorted perovskite compounds,

because these materials show similar property dependence upon lanthanide cation

size reminiscent of the Ln2Cu04 series of compounds. The influence of the size of

the cation profoundly effects the magnetic and transport properties.'* The lanthanide

titanates, LnTiOs, crystallize in the orthorhombic Pbnm space group (GdFeOs

structure-type). 5 LaTiOs and CeTiOs are metallic (or semi-metallic) at room

temperature, PrTiOs and NdTiOs are small band gap semiconductors (.03 eV),

whereas the rest of the LnTiOs compounds are larger band gap semiconductors."*

Their low-temperature magnetic properties vary from antiferromagnetic ordering for

217
LnTi03 (Ln = La-Pr), no apparent ordering for NdTi03, and ferromagnetic

ordering for Ln = Gd-Yb, Y ^ There are no close Ti-Ti contacts in this structure so

that the properties are determined by the covalency of the Ti-O-Ti linkages. The

effect of the size of the lanthanide cations is seen dramatically in the magnetic and

electronic properties of the solid solutions La2/3+xTi036 and LaxYi.xTi037 The

substitution of an alkaline earth or alkali metal cation for the lanthanum cation will

effect the density of states at the fermi level and may give rise to a superconducting

state.

LaTi03 undergoes a metallic to semiconducting transition at 125 K.^ Below

125 K antiferromagnetic ordering occurs. The magnetic structure has been

determined by neutron diffraction to be type G with sublattice canting and a titanium

(III) moment of 0.45(5) |Ib.^ No structural changes are apparent at or below the

transition temperature. The properties of LaTi03 have been interpreted as a result

of itinerant d electrons in overlapping Hubbard bands.'*"^ A transition to an

itinerant antiferromagnetic state will give rise to a gap near the Fermi level and

hence the transition to semiconducting behavior.

Cation substitution of LaTi03 provides an optimal system for investigating

the effects of the cation size and covalency as well as the effects of the Ti valence

state upon the magnetic and transport properties in this system. The lanthanum

cation is not magnetic and substitution of an alkaline earth cation for the lanthanum

cation will result in holes in the valence band and metallic behavior. LaTi03 is apparently on the border of a metal-semiconductor transition and oxidation of the

compound may give rise to a superconducting state. This paper briefly outlines

some of our recent results on the structural and magnetic properties for the series of

compounds Lai-xSrxTi03 (0 < x < .5).

EXPERIMENTAL

Ti02 (99.999%) and SrC03 (99.999%) were obtained from Johnson

Matthey. La203 (99.999%) was from Research Chemicals. Ti203 (99+%) was

from Alfa Chemical. Ti203 and SrC03 were dried at 120°C for 24 hours prior to use. La203 and Ti02 were dried at 1200 'C for 24 hours prior to use. Ti203 was

characterized by thermal gravimetric analysis (TGA) and was determined to be

Ti202. 158(8}.

LaTi03 was prepared by arc melting stoichiometric amounts of La203,

Ti202. 158(8) and Ti02 in an argon atmosphere. SrTi03 was prepared by reacting a pressed pellet of well mixed SrC03 and TiC)2 in a platinum boat for 12 hours at nOO'C and then at MOO'C for an additional 12 hours. Several firings and

regrindings are usually necessary in order to obtain a pure material. A large

quantity of LaTi03 and SrTi03 (10 gms each) were prepared and used in the

preparation of all the Lai.xSrxTi03 samples. The Lai.xSrxTi03 samples were

218
prepared by arc melting stoichiometric amounts of LaTiOs and SrTi03 in an argon

atmosphere. TheLai-xSrxTi03 samples were dark blue-black in color.

X-ray powder diffraction data were obtained with copper radiation (Kai)

using an Enraf-Nonius Guinier camera equipped with a Johansson monochrometer.

Samples were ground and mixed with about 5-10% silicon powder and placed on a piece of cellophane tape. The positions of the five observed diffraction lines for

silicon were fit to known diffraction angles by a quadratic function. The lattice

constants were then calculated by least-squares fit to indexed 29 values. The

program POWDER was used to calculate the expected powder pattem for LaTi03

(space group = Pbnm, GdFeOs structure-type) and SrTiOs. The experimental

diffraction positions and intensities agreed very well with the calculated diffraction

pattem for LaTiOs and SrTi03. The lattice parameters for LaTiOs are similar those

reported from neutron diffraction data (a = 5.6247(1), J2 = 5.607(1) and £ =

7.9175(1))^ (see Table I). The lattice parameter for SrTi03 was determined to be a = 3.9048(6). Lattice parameters for Lai-xSrxTiOs (0 < x < .5) are given in Table I. Analysis of La, Sr and Ti was done on all samples using a Cameca SX50

microprobe. The data were consistent with the theoretical stoichiometry to within

2-3%. The oxygen stoichiometry was determined by TOA using a calibrated

DuPont 2100 Thermal Analyzer. All the TGA measurements were done in platinum

containers with an oxygen flow (50 cc/min). The heating cycle was from 20-

lOOO'C at a rate of 5 °C/min. Table 11 provides the oxygen stoichiometry for LaixSrxTiOs samples determined from the TGA data.

Magnetic measurements were carried out on a Quantum Design SQUID

magnetometer. The samples of composition 0 < x < .1 were cooled to 5 K in a

field of 10 kOe and the magnetization versus temperature data were taken at a field

of 10 Oe (5 - 170 K) and 10 kOe (170 - 300 K). Compositions with x = .1, .2 did

not show a magnetic transition and magnetization versus temperature data were

taken at fields of 1000 Oe for T = 5 - 170K and 10 kOe for T = 170 - 300 K.

Magnetic measurements on all the other compositions were taken at a field of 10

kOe.

RESULTS & DISCUSSION

A number of different reactions were performed in order to determine the

most reliable method for preparing single phase LaTiOs and Lai-xSrxTiOs

materials. Single phase LaTiOs was only obtained by combining stoichiometric

amounts of La203, Ti202. 158(8) and Ti02 and arc melting in an argon atmosphere.

Attempts to prepare LaTiOs by reacting a pressed pellet of the same composition at 1200 °C in an evacuated silica ampoule were unsuccessful. LaTi03 is somewhat air

sensitive and will react slowly with oxygen at room temperature to eventually

produce La2Ti207. The most satisfactory procedure for producing the Sr doped

LaTiOs compositions was to arc melt mixtures in an argon atmosphere according to the equation:

219
(l-x)LaTi03 +xSrTi03 Lai.xSrxTiOs

Table I shows the lattice parameters determined for Lai-xSrxTiOs (0 < x < .5). The

samples prepared with 0 < x < .3 could be indexed in the orthorhombic space

group, Pbnm. This is a distorted perovskite structure (GdFe03 structure type) with

the comer shared Ti06 octahedra tilted towards the rare earth ion. As the Sr cation

concentration is increased in these samples, the difference between the lattice

parameters a and b decreases ( 0 < x < .1) and then increases (.1 < x < .3). This is reflected in a change in slope for a plot of Volume vs. composition, x, at x = .1 indicating that there is a subtie change in the structure. The diffraction peak

positions and intensities agreed very well for all compositions with the calculated

powder diffraction patterns based on the Pbnm, GdFeOs structure type. For the

compositions x = .4, .5, the data could best be indexed as tetragonal. We are

currently using X-ray powder diffraction Rietveld refinement to better characterize

the structures.

Table I. Lattice parameters (A) for Lai-xSrxTi03 (0 < x < .5). x a b c

.00 5.627(1) 5.605(1) 7.914(2)

.05 5.610(3) 5.595(1) 7.912(1)

.10 5.592(5) 5.581(2) 7.900(2)

.20 5.586(4) 5.568(3) 7.885(3)

.30 5.577(5) 5.555(3) 7.861(3)

.40 3.918(1) 3.918(1) 3.898(5)

.50 3.9140(3) 3.9140(3) 3.989(2)

Microprobe were performed on all of the samples and it appears that the

materials can be prepared by this method fairly close to the desired composition.

From the back scattered image of the scanning electron microscope (SEM), it is apparent that there is a small amount (not detectable by Guinier powder diffraction

methods) of secondary oxidized phase at the grain boundaries. This indicates that

there is some inhomogeneity in oxygen content for these samples. It was noted that

the amount of this secondary phase appears to increase with Sr content. TGA data

indicates that LaTiOs begins to significantly uptake oxygen at about 350°C and the

strontium doped compositions at somewhat higher temperatures. As the amount of

Sr cation is increased, the temperature at which the sample uptakes oxygen

increases so that La.sSr sTiOs uptakes oxygen at about 650°C. Oxygen content

was determined by TGA from the oxidation of LaTiOs to La2Ti207 and

Lai-xSrxTi03 to La2Ti207 + SrTiOs. The samples appeared to be close to stoichiometry in oxygen (Table II). Several samples have oxygen stoichiometry

220
apparently greater than 3.0. This is consistent with a small amount of oxidized

secondary phase at the grain boundaries of these samples.

Table TT Oxvffen stoichiometrv forl^ai vSr-irTiO^ HptPTminftd from XfiA data

Theoretical Stoichiometry Experimentally Determined

Stoichiometry

LaTiOs LaiTi02.98

La.gSr.iTiOs La.9Sr iTiOs.oo

La.gSr^TiOs La.8Sr2TiO3.01

La.vSrjTiOs La.7Sr.3TiO3.02

La.6Sr4Ti03 La.6Sr.4TiO2.99

La.5Sr5Ti03 La.5Sr.5TiO3.01

Temperature dependent magnetic susceptibility was performed on all compositions. The polycrystalline samples of LaTi03 showed temperature

independent paramagnetism above 150 K as reported previously by several other

groups.'*"^ At 145 K, there is a sharp magnetic transition which has been attributed to a canted antiferromagnetic lattice.^ Except for the higher onset temperature (145

K versus 125 K) for the magnetic transition, our data is comparable with literature

values on single crystal samples.^ This transition temperature is very sensitive to the La stoichiometiy and small differences in either La or oxygen content could

account for the difference in temperature observed for this magnetic transition.^

For the composition, La.95Sr 05TiO3, this magnetic transition occurs at about 125

K. This magnetic transition is completely absent for the compositions with x > .1. Figure 1 shows the temperature dependent magnetic susceptibility data for the

samples Lai.xSrxTi03 (x = 0 - .1). The high temperature data are temperature

independent paramagnetic for the compositions 0 < x < .3 and x = .5. The

temperature dependent magnetic susceptibiUty for the composition La.6Sr.4Ti03 is more complex. The full details of the magnetic and electronic properties of these

materials will be described elsewhere.

221
T«fnp«r«t«r« (K)

Figure 1. Magnetic Susceptibility data for Lai-xSrxTi03 x = 0, .05, .1.

ACKNOWLEDGEMENTS

We thank Robert N. Shelton for use of the TGA and the SQUID magnetometer and

David J. Webb for useful discussion conceming the magnetic properties. We thank

Traci Y. Kuromoto and Peter Schiffman for technical assistance with the

microprobe data. This research is supported by the National Science Foundation

(DMR-8913831) and UCD Committee on Research.

REFERENCES

1. (a) (SrTiOs-x): Schooley, J. J.; Hosier, W. R.; Cohen, M. L. Phys. Rev. Lett. 1965, 75, 108. (b) (LiTi204): Johnston, D. C; Prakash, H.; Zachariasen,

W. H. Viswanathan, R. Mater. Res. Bull. 1973, 8, 111. (c) (AxWOs):

Sweedler, A. R.; Raub, C; Matthias, B. T. Phys. Lett. 1965, 75, 108.

2. (a) Bednorz, J. G.; Muller, K. A. Z. Phys. B: Condens. Matter. 1986, 84,

189. See for example the following recent reviews: (b) Sleight, A. W.

Science 1988,242, 1519. (c) Muller-Buschbaum, H. An^^w. Chem. Int. Ed. Engl. 1989, 28, 1089.

3. (a) Wani, B. N.; Rao, U. R. K.; Venkateswarlu, K. S. /. Mat. Sci. Let. 1990,

9, 277. (b) Yee, K. A.; Albright, T. A.; Jung, D.; Whangbo, M. -H. Angew.

Chem. Int. Ed. Engl. 1989, 28, 750.

222
4. (a) Grecdan, ].E. J. Less Comm. Metals 19S5, 111,335. (b) Greedan, J. E.

in The Rare Earths in Modern Science and Technology Vol. 3, McCarthy, G.

J.; Silber, H. E.; Rhyne, J. J. eds. Plenum Press: New York, 1982, p. 431.

5. Eitel, M.; Greedan, J. E. /. Less Comm. Metals 1986, 116, 95.

6. Bazuev, G. V.; Makarova, O. V.; Shveikin, G. P. Inorg. Materials 1983, 79,

92.

7. Goral, J. P.; Greedan, J. E.; MacLean, D. A. /. Sol. State Chem. 1982, 43,

244.

8. MacLean, D. A.; Greedan, J. E. Inorg. Chem. 1981, 20, 1025.

9. Goral, J. P.; Greedan, J. E. J. Magn. Magn. Mat. 1983, 27, 315.

10. Clark, C. M.; Smith, D. K.; Johnson, G. J. "A Fortran IV Program for

Calculating X-ray Powder Diffraction Patterns—Version V"; Department of

Geosciences, The Pennsylvania State University: University Park, PA, 1973.

223
DISCUSSION

M. Norton : Do you plan to put any of these materials on substrates and if so do you

think that you'll be able to expand the lattice parameter to get closer to 180 degrees

metal-oxygen-metal linkage?

S. Kauzlarich : Actually, we talked about it, with Dave Webb in Physics (UCD), and his

group is very interested in trying to put down films via sputtering. But we haven't

actually done anything yet. It is premature because the bulk materials are poorly

characterized at this point.

R. Roth : I have a little trouble believing your oxygen stoichiometry. SrTiOj, when arc- melted in argon, is going to be considerably less than O3. And yet you're going all the

way to 50% strontium titanate and saying that it is still very close to three and in fact

maybe over. What exactly are these oxygen numbers dependent upon experimentally?

S. Kauzlarich : That is a very good question. In fact, for lanthanum titanate itself we

tried really hard to reduce and it does not reduce at all. And, as you put in strontium,

because strontium titanate has such a wide oxygen content range, you would expect to be

able to vzu^y the oxygen content. And you would expect it to be deficient. What it could

be is that these materials are very air sensitive and they oxidize very rapidly. Much more

rapidly than we're being careful with or taking account of. Basically, we know looking

at the X = 0.5 sample that we have a more oxidized second phase in there, because we

saw that with SEM; whether or not the second phase came about on the grain boundaries

because of the way they polish it to look at it with SEM and microprobe, is a whole

debate on its own. But, we do see a second oxidized phase there and we haven't tried to reduce any of the strontium solid solution materials at all.

R. Roth : Tank argon has a Pqj about lO'-'. You can get it down considerably if you

put a little hydrogen into it. Have you tried any of that?

S. Kauzlarich : No. The only thing we tried to do is to reduce lanthanum titanate

because we were worried about nonstoichiometry in that system. If you go through the

literature there is a really wide disparity in the properties and so we were worried about

that a lot. We only tried 6% hydrogen in nitrogen so we really didn't try very hard to reduce it but we couldn't reduce it all.

K. Poppelmeier: It's just a comment. I agree with Bob Roth that you would think you

would have a lot of reduction; but one possibility that would keep it O3 is that when you

have the oxygen vacancy it would also be in the first coordination sphere of the

lanthanum/strontium. That may have some effect, with the higher charge of lanthanum

keeping the lattice from accommodating the oxygen defect just for local charge neutrality.

S. Kauzlarich : Yes, that is a possibility. Little is known about oxygen nonstoichiometry

in this system.

224
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

SYNTHESIS AND CHARACTERIZATION OF PHASES IN THE SYSTEM Ba-Au-Ag-02

R. S. Roth, C. J. Rawn and M. D. Hill

National Institute of Standards and Technology

Gaithersburg, MD 20899

ABSTRACT

Ba02 has been studied by heating in air, flowing O2 , and sealed Au or Pt

tubes with or without excess O2 and characterized by TGA and x-ray

diffraction of quenched products. Ba02 forms a nonstoichiometric solid

solution Ba02-x, with x <0.25 at ~800°C and transforms to BaO at -820°C

in oxygen. The c-axis of the barium peroxide, CaC2 -related structure,

decreases from 6.851A to ~6.666A with increasing x, although there is very

little change in the a- axis. Ba02-x has a c-axis of ~6 . 701A when quenched

in air from 725°C and transforms to BaO at 735 ± 2°C in air (Pq = 0.2).

The reaction between Ba02-x and Au or Ag which has been observed in air

up to -800° C is probably due to a hydrated melt formed by a reaction with

atmospheric moisture. However, when Ba(N03)2 is heated above its melting

point (~592°C) in contact with Au, several new phases are formed. The

most Ba-rich phase, presumably BagAu2 0g , hydrates readily on exposure to

air. A second phase apparently more Au-rich, appears to be more stable.

Liquid Ba(N03 )2 in contact with Ag does not decompose or react with Ag at

625°C (in flowing O2 ) . However, at 700°C in O2 , the BaCNOj )2 slowly

decomposes and forms single crystals of Ba02_x, and no clearly discernable

Ag containing phases can be identified. Ba02-x and Ba-aurates have been

characterized by powder and precession camera x-ray diffraction data, SEM,

EDX and DTGA data. The defect structure of Ba02-x Is currently being

investigated by neutron powder diffraction total profile (Rietveld)

analyses

.

INTRODUCTION

Both gold (Au) and silver (Ag) are known to react with and enter into the

lattice of the copper containing high T^ superconducting oxides. When BaO

is a major component of the phases, container and atmospheric reaction

problems with these metals become a very important factor in synthesis.

BaO reacts quickly with atmospheric moisture and various Ba-0H-H20 phases

are formed after a few moments exposure to air. BaCOg also forms by

reaction with CO2 in the atmosphere. The use of BaC03 as a starting

material generally results in CO2 contamination of the products, either

in the desired phase or in grain boundaries. For these reasons many

researchers have employed either Ba02 or Ba(N03)2 as starting materials.

However, Ba02 cannot be commercially obtained without about 2-3% BaC03

contamination, and Ba(N03)2 has the disadvantage of a very low melting

point ~590-595°C. This work is, therefore, the result of an effort to

understand the chemical reactions which take place when either Ba02 or Ba(N03)2 is heated in contact with Au or Ag.

225
NONSTOICHIOMETRY OF BaOo

Several commercial sources of Ba02 were examined by x-ray powder

diffraction. All had similar diffraction patterns, essentially identical

to that published [1] with a body centered tetragonal unit cell, a = 3.8154 (face centered equivalent reported as 5.3958 [2]) and c = 6.8513A

at 25°C. Previous reports in the literature [3] refer to the possibility

of an unknown phase between BaO and Ba2Cu03 [4] when synthesized in O2

.

Attempts to examine this region of the BaO-CuO-02 system always resulted

in Ba2Cu03 plus a phase similar to but not equivalent to Ba02 . When pure

Ba02 was heated in O2 at 800° C and "furnace cooled" by turning off the

furnace and allowing the specimen to cool to 250°-300°C in several hours

the same phase resulted, namely a Ba02-like phase with c = 6.786A. It

seemed most likely that the explanation for the change in the value of the

c-axis was due to a loss of oxygen. Therefore, a series of experiments

was conducted to characterize the results from various heat treatments of

Ba02 .

Several DTGA examinations were made of Ba02 in pure O2 (passed through a 'gettering' device to remove CO2 and H2O) and in 80/20 N2/O2 (figure 1)

.

At 800° C in O2 Ba02 was found to lose weight to a value corresponding to

about BaOj^ ; -BaO-j^ g , depending on whether the calculation is made on the

original loss or on the regain of weight, respectively. As the x-ray

pattern does not return to the original c-axis value, ~6.824A, the first

number must be low. However, obvious BaCOg and amorphous material in the

commercial Ba02 starting material indicate the larger number is probably

not low enough. The average value of BaO^^ 75 probably represents the best

available estimate of the minimum oxygen content in this phase before the

phase transition to BaO takes place. This is shown by the DTA peak and

greater weight loss at 820° C in O2 . The total weight loss in a mixture

of 80/20 N2/O2 is less than in O2 and the transition to BaO takes place

almost 100°C lower, at ~735°C. This data differs slightly from the values

which can be derived from other data [5] obtained by varying the oxygen

pressure at constant temperature.

A number of experiments were performed in a quench furnace in order to

measure the unit cell dimensions of Ba02-3j under various conditions of

Pq . The specimens were generally contained in 2 , 3 or 5 mm diameter Au

tufees, suspended in the hot zone by a thin Pt wire and after heating for

2 or more hours, dropped into a Ni crucible which was cooled by He flowing

through a Cu coiled tube immersed in liquid N2 . Several experiments were performed containing BaOg in a Au tube inside a larger diameter Pt tube

containing a weighed amount of Pt02 calculated to produce an O2 atmosphere

from -1.0 to -10 atmospheres depending on the temperature and volume of

the tube. All such experiments showed that the transition to BaO occurred

well above 850°C or even 900°C when the total O2 pressure exceeded 5-10

atmospheres. The smallest c-axis value or the greatest value of x in

Ba02 - X was found when Ba02 was heated in a sealed Pt tube with no Pt02

,

at 800° C and corresponded to c = 6.666A in equilibrium with some BaO.

226
Weight chonge %

102

200 300 400 500 600 700

Temperoture

Figure la. DTA/TGA for Ba02 . Heating and cooling rates =5"C/min.

Experiment performed in one atmosphere O2 , heated to 900°C, held

for one hour, cooled to 500°C, held for one hour and cooled to

250°C.

Weight chonge %

102

200 300 400 500 600 700

Temperoture

Figure lb. DTA/TGA for Ba02 . Heating and cooling rates =5°C/min.

Experiment performed in 80 N2/2O O2 heated to 800° G at 5°G/min,

held for one hour, cooled to 500°C, held for 3 hours and cooled to

250°C.

Figure 2 shows the values of the c-axis measured from specimens quenche

from various temperatures in air. A portion of the characteristic

x-ray powder diffraction patterns for some of the quenched specimens ar shown in figure 3

.

227
CO,

1

1

11 ,

730°C

725''C 700°C J

65o°c TV -^-^^yi 6oo°c ___7v__^nA'

550°C soo-c yv^.^ ^/

450''C 004 ;

50 ^1 ^ si* ^ sIb ^ ^ ^ e) TWO-THETA (DEGREES)

Figure 3. Portion of the x-ray powder diffraction

patterns of quenched Ba02 - ^ heated at various

temperatures, 450° - 735°C. Note change in the

position of the (004) peak.

228
Attempts to prepare a large batch of Ba02-x ^otc neutron diffraction

analysis have not been very successful as Ba02-x ^ii^ will either

hydrate and/or carbonate slightly at <700°C, and this causes a reaction

rim to form on the Au foil. Such experiments are currently being

performed in a pure oxygen atmosphere, but quenching a large batch in an enclosed atmosphere is a formidable experimental procedure. The x-ray

powder diffraction pattern of the reaction product which contains unknown

phases in addition to BaC03 and Ba02_x is shown in figure 4.

''s i5 iS 25 S 35 S 35 45 s5 S so TWO-THETA (DEGREES)

Figure 4. X-ray diffraction powder pattern of the

reaction product obtained by heating a 40g batch of

Ba02 in air on Au foil at 700°C - 20 min.

DECOMPOSITION-REACTIONS OF Ba(NO. )o

Due to the impurities (generally BaC03 ) always found in commercial Ba02

,

attempts were made to decompose Ba(N03)2 in O2 to form pure Ba02 . It was

found, however, that Ba(N03)2 is perfectly stable in one atmosphere of

flowing O2 at 575°C and below its melting point shows no sign of

decomposition or reaction with containers of Ag, Al or Au. However, at

600°C the molten Ba(N03)2 is very corrosive and completely destroys the

Al dish.

Reactions with Au

Molten Ba(N03)2 can be contained on Au foil or in Au tubes at 600°C, but

after -18 hours at 625°C in pure O2 , no Ba(N03)2 or (BaO, Ba02 , etc.) is

present in the x-ray powder diffraction pattern. If care is taken to

avoid reaction with atmospheric moisture, etc., after removal from the

furnace, an x-ray powder diffraction pattern can be obtained from the

material separated from the Au foil (figure 5) which is quite different

from any previously observed pattern. "Single crystal" diffraction

229
patterns were also obtained from some preparations, but these crystals

appear to be fiber bundles rather than real single crystals, as can be seen from various x-ray precession patterns (figure 6) obtained looking

dovm the fiber axis. At the present time it is not known if this phase

is really hexagonal or orthorhomic (figure 6c) . The phase is

characterized by a very strong peak at -9.75° 26 (~9.1A). Energy

dispersive X-ray (EDX) data (figure 7) indicate that the crystals contain

about one half as much Au as Ba. The composition, therefore, may be

Ba5Au20g , although it is not possible to rule out the presence of some nitrate ions in the structure.

A second phase occurs when liquid Ba(N03 )2 has flowed over the outside of

the Au tube. This phase presumably has a higher Au content than

"BagAu209", and may possibly be BaAu02 5 . "Single crystal" x-ray

precession patterns (figure 8) and a powder diffraction pattern (figure

9) show a characteristic phase with a strong low angle diffraction peak

at -8.1° 26 and possible tetragonal symmetry. Again this may well be only

a p sue do - symmetry as this compound has not been obtained as a single phase

in a large enough volume to obtain a detailed x-ray powder pattern. The

x-ray diffraction pattern shown in figure 9 also contains both Ba(N03 )2

and the Ba-Au-0 phase with the 9.75° 26 peak. This second phase with the

characteristic peak at 8.10° 26 is much more resistant to attack by

atmospheric moisture and often will show up as the last remaining

crystalline phase in a mixed phase batch after 2-3 days exposure to air.

li iS iS S S 35 45 ^ 60 TWO-THETA (DEGREES)

Figure 5. X-ray powder diffraction pattern of the

reaction between molten Ba(N03)2 and Au foil at

650° C in one atmosphere of O2 . This pattern shows

no Ba02-x. Ba(N03 )2 or Au.

230
Figure 6a. "multi

twinned" or "fiber

bundle" parallel

to needle axis

.

Figure 6c.

"single" crystal

parallel to needed

axis

.

Figure 6. X-ray precession photographs of "single crystals" formed by

the decomposition/reaction of Ba(N03)2 in a Au tube or on Au foil at

600°C in O2

.

Figure 6b. "single"

crystal perpendicular

to needle axis

.

231
Figure 7a. 40X. Figure 7b. 200X.

Ba

Figure 7. SEM/EDX data (courtesy J. Kelly) for "BagAujOg" crystals

prepared by heating Ba(N03 )2 in O2 at 600°C - 24 hours on Au foil.

232
Figure 8a. Perpendicular to

the needle axis.

Figure 8b. Parallel to the

needle axis

Figure 8. X-ray precession photographs of "single crystals" of the phase

"BaAu02 5" formed by a thin coating of molten Ba(N03)2 decomposition/

reaction with the outside wall of a Au tube at 600° C in 0,

-•JleX-J —

i

-4

"

1 ] 1

i j

1

1

] —\-

i

i

-

—!—

i

1

,

1 1 1 I

\ 1

1

1

1

!

I

!

i

i

!

!

tH—

1

i

r i; 9 10 11 12

20 Figure 9. Low angle portion of x-ray powder

diffraction pattern of reaction between a thin

film of molten Ba(N03 )2 and the outer wall of a Au tube, showing two phases "Hex"- "BagAu3 0g"

(9.75° 29) and "Tet" - "BaAuO2 . 5 (8.1° 20).

233
Reaction with Ag

When a Ba02 pellet is heated in air on Ag foil to -yOO'C an obvious

reaction rim is observed similar to that found with Au, suggesting that

Ba02 might have a melting point of about 450° C [6] . However, the DTA data

illustrated in figure 1 show that Ba02 does not melt. The reaction rim

therefore must be due to impurities present in the atmosphere, as such

"melting" phenomena do not occur in pure O2 or a mixture of 80/20 N2/O2

.

Mixtures of 50/50 Ba02/BaC03 in both Au and Ag tubes showed no sign of

reaction with the containers when heated in O2 to 700°C. It is therefore

concluded that this reaction rim or "melting" phenomena is due to the

presence of moisture and the formation of hydroxy or hydroxy/carbonate

containing phases

.

When Ba(N03 )2 is heated in Ag tubes to TOO^C in pure O2 no readily

identified extra phase is found. X-ray powder diffraction patterns show

mostly Ba02-x plus Ba(N03 )2 and a few small extra peaks which may be due

to hydration or some unidentified Ag-containing phase. However, this

procedure results in the formation of quite large single crystal fragments

of Ba02 _ X • The best crystals were obtained by slow heating, 3°C/hr, from

590°C to 700°C, holding 24 hours and cooling at 6°C/hr to 500° and again

holding for 24 hours. Apparently, crystal growth is initiated by the

formation of Ba02-x nucleation sites in the Ba(N03 )2 melt and the growth

of quite large platelet crystals with rectangular shapes 2-5 mm wide and

0.2-0.5 mm thick, figure 10. The crystals often show perpendicular twin

boundaries, and further twinning can be induced by pressing on the crystal

face with the point of a needle. This later phenomenon indicates that the

Ba02_x crystals are probably ferroelastic . Single crystal x-ray

precession photographs were used to characterize these crystals (figure

11) and show that the rectangular edges are (110) directions and the

c*-axis is perpendicular to the flat face. Crystals often show splitting

of spots due to the tendency to form low angle grain boundaries either

caused by growth boundaries or cleavage of the crystals perpendicular to

the flat face. Twinning can often be found with 2nd and 3rd twins

oriented along hkO planes with the c*-axes characteristically inclined at

-76° (figure lid)

.

234
\ \ 1 /

\ \ /

X i

' 1 / f \

/ / \ \ \

\

Figure 11a. hkO Figure lib, hOl

Figure 11c. hkl

• •

> «

Figure lid. hOl , unfiltered

radiation, showing typical

growth habit of multi (twinned?)

crystals with the c"*-axis always

inclined at the same angle (-76°)

Figure 11. Single crystal precession photographs (Mo radition) of

crystals of Ba02-x prepared by decomposition of m.olten Ba(N03 )2 from

590°C to 700°C in 0^ .

235
SUMMARY

Reversible nons toichiometry in Ba02.x was identified and characterized by

quenching and DTGA experiments. Ba02_x was found to transform to BaO at

~735°C in air and 820°C in one atmosphere O2 . Ba(N03)2 reacts with Au in

pure O2 to form several phases identified by "single crystal" and x-ray

powder diffraction and EDX analysis. No reaction with Ag could be

definitely confirmed, but single crystals of Ba02_x can be grown by

decomposition reaction of molten Ba(N03)2 in Ag tubes and a pure O2

atmosphere

.

REFERENCES

1. S. C. Abrahams and J. Kalnaj s , "The Crystal Structure of Barium

Peroxide", Acta Cryst. 7, 838-847 (1954).

2. H. E. Swanson, N. T. Gilfrich, and M. I. Cook, "Standard X-Ray

Diffraction Powder Patterns", NIST Circular 539 [6], 18-19 (1956).

3. K. G. Erase and D. R. Clarke, "Phase Equilibria and Crystal

Chemistry in the System Ba-Y-Cu-0", Adv. Ceram. Mat. 2 [3B],

295-302 (1987).

4. W. K. Wong-Ng, K. L. Davis, and R. S. Roth, "X-Ray Powder Study of

2BaO«CuO", J. Amer. Ceram. Soc . 71 [2] C-64-C-67 (1988).

5. 0. V. Kedrovskii, P. V. Kovtunewko , E. V. Kiseleva, and A. A.

Bundel Zh. Fiz . Khim. 41 [2], 414-419 (1967). (Russ.)

6. R. C. Weast, Ed. "Handbook of Chemistry and Physics". The Qeiical

Rubber Co., Cleveland, Ohio.

236
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

STRUCTURE AND PROPERTffiS OF REDUCED EARLY TRANSITION

METAL OXIDE SINGLE CRYSTALS GROWN FROM BORATE FLUXES

S.A. Sunshine, B. Hessen, T. Siegrist, A.T. Fiory & J.V. Waszczak

AT&T Bell Laboratories, 600 Mountain Avenue, Murray Hill, NJ 07974

Reduced ternary oxides of the early transition metals exhibit a variety of

interesting physical and structural properties. Superconductivity (LiTi204 [1],

tetragonal and hexagonal Mo and W bronzes [2]) and charge density wave

behavior (Mo 'blue' bronzes [3]) have been observed in these systems.

Structures range from simple perovskites like SrxNb03 [4] to complicated

structures involving metal-metal bound clusters [5]. Single crystals are highly

desirable for a better understanding of strucmre-property relations. In systems

having low-melting high oxidation state ternary precursors (like A2MO4 and

A2M2O7; A = alkali metal, M = Mo, W) molten salt electrolysis provides a convenient route to single crystals. For other systems (particularly niobates)

other methods have to be employed. Here we describe the use of borate fluxes

to produce single crystals of various reduced ternary niobates and titanates in

inert atmosphere. Borate fluxes have been regularly employed in the crystal

growth of ferroelectric s and garnets [6], but rarely used for reduced

compounds. An interesting aspect of the borates AnOm*xB203 is that they can

act both as a source and as a buffer for A^On, allowing the product formation

to be influenced by the flux composition when the A-cation is also part of the

reduced ternary system. An advantage of the Ba, Sr, and La-borates is their low

volatility, which allows their use under high-vacuum conditions.

THE BA-NB-0 SYSTEM

Heating a mixture of Nb02 and BaO-3B203 to 1100°C (10"^-10""^ Torr),

followed by slow cooling (8°C/h) produces golden crystals of the new niobate

Ba2Nbi5032 (average Nb oxidation state 4+), as well as crystals of Ba3Nb50i5

[7] (Nb"^-^"*") and BaNbgO^ [8] (Nb^-^^^) which form in a disproportionation

side-reaction. The structure of Ba2Nbi5032 (Fig. 1) contains a layer (A) of

isolated octahedral [Nb60i2]06 metal-metal bound clusters, interconnected by

Nb-atoms in octahedral holes in the lattice. These layers are separated by layers

of Nb30i3 groups (B), and layers with Ba-ions and isolated NbOg octahedra

(C). Single crystal resistivity measurements indicate semiconducting behavior.

237
o

Figure 1. X-ray structure of Ba2Nbi5032, projected along the c-axis (a), and

perpendicular to it (b). Oxygen atoms omitted for clarity.

Figure 2. Single crystal resistivity Figure 3. Magnetic susceptibility

data for Ba2Nbi5032. for Ba2Nbi5032 (uncorrected

for Curie 'tail').

with a semiconductor-to-semiconductor transition at 170 K (Fig. 2). This

coincides with a magnetic transition from a weakly paramagnetic state to diamagnetism below the transition temperature (Fig. 3).

The compound Ba3Nb50i5 has the centrosymmetric tetragonal tungsten

bronze (TTB)-type structure (Fig. 4), as observed by Magneli for K3W5O15 [9].

In contrast to the isostructural (metallic) Mo and W analogues, the resistivity of

Ba3Nb50i5 shows a semiconducting temperature dependence, and a stronger

anisotropy (p^c = 55 p//c at 300 K).

The perovskite BaxNb03 [10] has been less investigated than its Sranalogue and appears to have a much smaller single-phase region in ceramic

samples. Red metallic single crystalline Bao.95Nb03 (a = 4.0853(2) A) was

obtained by recrystallizing a ceramic with nominal composition BaNb03 from

the Ba-rich flux BaO B203. The refined Ba-site occupancy and cell parameter

compare well with those obtained by Svensson from profile analysis on a ceramic sample [11].

238
Figure 4. Polyhedral representation

of the X-ray structure of

BasNbsOis.

THE SR-NB-O SYSTEM

A broad solid-solution range, Sr^NbOs, is known in this system [4]. A

single phase cubic perovskite exists for 0.70 < x < 0.97, while for 0.60 < x <

0.70 a two-phase mixture containing the TTB-type Sr3Nb50i5 [12] is found.

By modification of the SrO nB203 flux composition and the ternary precursor,

single crystals of the perovskite« Sro.^2Nb03 (blue, a = 3.9865(2) A) and

Sro.86Nb03 (purple-red, a = 4.(X)77(6) A) were obtained, the first from ceramic

Sr5Nb40i4 [13] and SrO-2B203, the second from ceramic SrNb03 and

SrO-B203. Some disproportionation was observed, yielding multiply twinned

crystals of SrNb80i4 [14], especially in relatively Sr-poor mixtures. Crystals of

Sr5Nb40i5_x (dark hexagons) [13], and Sr2Nb207_x (black platelets) [15] were

observed as side products.

Formation of the TTB-type phase in the melts is apparendy less favorable

for Sr than for Ba. Small crystals of Sr3Nb50i5 could however be obtained by

grain growth in a ceramic pellet with addition of some borate mineralizer (a =

12.3609(7) A,c = 3.8974(3) A).

THE SR-TI-O SYSTEM

In ceramic studies, the reduced part of the Sr-Ti-O phase diagram was

found to be heavily dominated by the perovskite solid solution SrTi03_x and

binary oxides like Tin02n-i» although a strongly Sr-deficient species was

suspected to exist [16]. Crystallization of a SrTi02.5 ceramic from SrO-2B203

produces purple-blue crystals of a compound with Sr « Ti (from EDAX),

together with some Tin02n-i. Single crystal X-ray diffraction showed the

compound to be triclinic, and to have the stoichiometry SrTiiiO20. Its structure

(Fig. 5) can be described as consisting of two interconnected layers, A and B,

perpendicular to the a-axis. The A layer contains two crystallographically

independent 'double rutile' chains [17], that are sheared after every block of six

TiOg octahedra. These chains, running parallel to the fe-axis, are interconnected

239
BAB

Figure 5. X-ray structure of

SrTiii02o: (a) unit cell projected

along the ft-axis; (b) A-layer, and

(c) B-layer, projected along the aaxis.

®

by single TiOg octahedra creating cavities for the Sr-ions. The B-layer contains

a single rutile chain that is sheared after every block of four TiOg -octahedra.

These chains cover the cavities with the Sr-ions in the A-layer. It is unclear if this phase is identical to the proposed "SrTii20i9" from ceramic studies [16], as there are significant differences in the reported X-ray powder pattern and

that calculated for SrTinO20- The average oxidation state of Ti in SrTinO20

is 3.45+, which implies that some oxidation of the product has occurred. Indeed

all reduced titanates in molten fluxes are extremely good getters for oxygen and

water. Recrystallization of SrTi02.5 from the Sr-rich flux SrO-B203 even in

vacuum produces small cubes of SrTi03_x which, as a function of x, range in

color from black to blue to colorless.

240
OTHER TITANATES

The oxophilicity of reduced titanates mentioned above is especially

prominent when more volatile fluxes are used which necessitate the use of

atmospheric pressure. Recrystallizing the superconducting spinel LiTi204 from

LiB02 flux under a purified Ar atmosphere produces well-shaped blue-black

octahedral crystals. The cell parameter (a = 8.3618(8) A) indicated, however,

that the product is not very reduced, and non-superconducting [1].

Lanthanide borates generally have a much higher melting point than the

Sr and Ba borates, but can still be used in vacuum over a limited temperature

range. Heating ceramic LaTiOs in La203 -36203 to 1250°C followed by

cooling to 1150°C yields a rather complicated mixture, which includes needles

of a new lanthanum titanate, Lao.3Ti204, with a monocUnically distorted

hollandite structure (Fig. 6). The tunnel sites are partially occupied by La.

Attempts to prepare this phase as a ceramic have not been successful, but

interestingly a recrystallization of a ceramic with nominal composition

"Lao.3Ti204" (not containing hollandite phases) from K2B4O7 produced

needles of the tetragonal hollandite KxTi204 [18].

Polyhedral representation

X-ray structure of

c CONCLUSIONS

Borate fluxes are interesting media for the crystal growth of reduced

ternary transition metal oxides. Modification of the flux composition provides a way of influencing the product formation. The fluxes are fairly easily removed

using dilute acids at ambient temperature, whereas most of the products are

inert under these conditions. In a number of cases the formation of new phases

was observed. The method is however very sensitive to the quality of the

atmosphere. A high-quality atmosphere control or the use of sealed

molybdenum containers should enable a more efficient use of this technique.

241
ACKNOWLEDGMENT

BH thanks the Netherland Organization for Scientific Research (NWO)

for the award of a fellowship.

REFERENCES

[1] (a) D.C. Johnson, H. Prakash, W.H. Zachariasen, and R. Viswanathan,

Mat. Res. Bull. 8 (1973) 777; (b) M.R. Harrison, P.P. Edwards, and J.B. Goodenough, Phil. Mag. 52B (1985) 679; (c) Y. Ueda, T. Tanaka, K.

Kosuge, M. Ishikawa, and H. Yasuoka, /. Solid State Chem. 11 (1988)

401.

[2] (a) Ch.J. Raub, A.R. Sweedler, M.A. Jensen, and B.T. Matthias, Phys.

Rev. Lett. 15 (1964) 746; (b) A.R. Sweedler, Ch.J. Raub, and B.T.

Matthias, Phys. Lett. 15 (1965)108; (c) A.W. Sleight, T.A. Bither, and

P.E. Biersted, Solid State Comm. 1 (1969) 299.

[3] (a) M. Greenblatt, Chem. Rev. 88 (1988) 31; (b) C. Schlenker (ed.),

"Low-Dimensional Electronic Properties of Molybdenum Bronzes and

Oxides", Kluwer, Dordrecht (1989).

[4] D. Ridgeley, and R. Ward, /. Am. Chem. Soc. 11 (1955) 6132.

[5] A. Simon, Angew. Chem. Int. Ed. Engl. 21 (1988) 159.

[6] D. Elwell, and H.J. Scheel, "Crystal Growth from High-Temperature

Solutions", Academic Press, London (1975).

[7] A. Feltz, and H. Langbein, Z. Anorg. Allg. Chem. 425 (1976) 47.

[8] S.J. Hibble, A.K. Cheetham, J. Koehler, and A. Simon, /. LessCommon Met. 154 (1989) 271.

[9] A. Magneli, Ark. Khemi 1 (1949) 213.

[10] (a) E.L Krylov, and A.A. Shamin, J. Gen. Chem. U.S.S.R. 25 (1955)

1637; (b) R. Kreiser, and R. Ward, J. Solid State Chem. 1 (1970) 368.

[11] G. Svensson, and P.E. Werner, Mat. Res. Bull. 25 (1990) 9. [12] M.M. Ninh Nguyen, J. Choisnet, and B. Raveau, Ci?. Acad. Sci. Paris

C 282 (1976) 303.

[13] J.M. Carruthers, and M. Grasso, J. Electrochem. Soc. 117 (1970) 1426.

[14] J. Koehler, A. Simon, S.J. Hibble, and A.K. Cheetham, J. Less-Common

Met. 142 (1988) 123.

[15] N. Ishizawa, and F. Marumo, Acta Cryst. B31 (1975) 1912.

[16] G.J. McCarthy, W.B. White, and R. Roy, /. Am. Ceram. Soc. 52 (1969)

463.

[17] A.F. Wells, "Structural Inorganic Chemistry", 5th ed.. Clarendon Press,

Oxford (1984), p. 214.

[18] T. Vogt, E. Schweda, C. Wuestfeld, J. Straehle, and A.K. Cheetham, /.

Solid State Chem. 83 (1989) 61.

242
DISCUSSION

A. Navrotsky: Do you ever get any ternary or higher borates formed other than the

starting material?

S. Sunshine: That is highly possible. We tend to disregard anything that remains white.

If there are borates, they are oxidized borates. In these reactions what you end up with

is a black reduced material and some sort of white flux or transparent flux that's left over.

The strontium borate we started with is obviously chemically altered in some way because

in these reactions the flux itself is actually being used in the reaction. We've really never

analyzed the left over borate after reaction. We have no evidence so far of any borate

incorporation in the materials that I've described or any of the reduced materials we've

looked at so far.

B. McCarroll: What did you use for a container?

S. Sunshine: In general most of these reactions were done in molybdenum foil; foil boats

that we folded ourselves. I started by doing reactions in alumina, but alumina is not a good choice. I made a lot of reduced niobium aluminates that way. In general, with

molybdenum we find no evidence for molybdenum incorporation in our products. If you were worried about that you could easily switch from molybdenum foil, to niobium foil,

and so forth, but we've been doing the reactions in molybdenum foil.

B. McCarroll : Were you pumping on these systems?

S. Sunshine: These reactions have been done, either in inert flow or in a vacuum system.

When they've done in a vacuum system they'are being pumped on continuously, even at melt temperatures. That is the advantage of the strontium, barium and lathanum borate

fluxes; even at those temperatures their volatility in a vacuum is relatively low.

H. Zandbergen : If you start with potassium titanium oxide is it possible to take out all potassium leaving titanium oxide hollandite?

S. Sunshine: Well, I think that there is in general a small amount of potassium left over.

I think the numbers are something like Kq 04TiO2 or something to that effect. That's not

work that I have done, so I don't know how closely they tried to analyze those numbers.

They claim it to be essentially a potassium free titanium oxide, although, I think they do

claim that there is a possibility for some residual potassium.

R. Shannon: How stable are these materials to oxidation?

243
S. Sunshine: They are very stable to oxidation at room temperature. Even when adding

indium contacts at < 200 degree centigrade, they seem to be stable. The compounds that

you have the most problems with are the ones that have a tunnel or layered structure in

which case they are much more susceptible to oxidation or reaction with water. But in

general if the ion, strontium, barium or lanthanum, is well encased in the structure then

they are actually quite stable.

244
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

THE KINETICS AND MECHANISM OF THE CRYSTALLIZATION OF

Mg2Al4Si50i8 FROM MgAl204 AND Si02 IN THE PRESENCE OF A

BISMUTH OXIDE FLUX

Ryan W. Dupon, Adam C. Tanous and Mark S. Thompson*

Corporate R&D, Raychem Corporation, 300 Constitution Dr., Menlo Park, CA 94025

ABSTRACT

Quantitative kinetics and mechanistic interpretation are presented for the reaction of

colloidal Si02 and MgAl204 to give Mg2Al4Si50j gin the presence of 2 atom% bismuth

oxide. The active flux agent appears to be comprised predominantly of bismuth oxide and

silicon oxide. The reaction proceeds through dissolution of the MgAl204 in the siliceous

flux to produce an aluminum substituted quartz stuffed with magnesium ions as an

intermediate. The stuffed quartz then converts to cordierite by a first order process. The

rate constant for this conversion is 4.4xl0'^s"^ at 1000^ with an activation energy of 70

kcal/mole

INTRODUCTION

The preparation of cordierite(Mg2Al4Si50| g) ceramics has attracted considerable attention

due in large part to the small thermal expansion coefficient, good strength, and low

dielectric constant of Mg2Al4Si50^ g. Reported preparation methodologies have included

the crystallization of glasses with the cordierite composition^"^ and the hydrolysis of

mixed, reactive intermediates such as metal alkoxides^'^ with the cordierite composition

followed by crystallization to the cordierite phase. The conversion mechanism from

starting materials to cordierite for each of these processes has been investigated in some

detail. Schreyer and Schairer have performed a comprehensive study of Uie observed

245
phases in the crystalhzation of glasses in the MgO-Al203-Si02 system around the

composition of cordierite. ^ Chowdry^ and Bemier et al^ have followed the conversion of

hydrolyzed precursors to Mg2Al4Si50|8 by X-ray diffraction. As preparative processes

both the glass crystallization and sol-gel route have some drawbacks. The glass

crystallization process uses common oxide starting materials but requires high temperatures

and is inherently a two step process. The sol-gel process uses expensive, high energy

intermediates which makes it an expensive and impractical route for bulk oxide materials.

We recently reported the preparation of Mg2Al4Si50i8 below lOOO'C by the reaction of

Si02 with MgAl204 in the presence of a small amount of a bismuth oxide flux^, while in

the absence of bismuth oxide there is no observed reaction. In that paper we described

several distinctive features of that reaction. It was observed that there was an induction

period before the appearance of cordierite that was both temperature and bismuth oxide

concentration dependent. Further, the induction period went to infinity at around 960*0,

below which temperature no cordierite was formed. As the presence of bismuth oxide

provides imique reaction paths to convert simple oxides to useful materials such as

cordierite it was of interest to us to follow the phase evolution of this reaction in an attempt

to understand why the bismuth oxide exerts such a strong influence on the reaction of these

oxides. Additionally, we sought to obtain quantitative kinetic information on the reactions

and phase transformations as they occurred. Quantitative kinetic analyses are not typically

done of reactions in the solid state; however, such information could be valuable to advancing the state of preparative art in oxide systems. In order to do this, an x-ray

diffractometer was fitted with a platinum hot stage to enable the collection of timedependent diffraction data near lOOO'C. The quantitative study of the kinetics for the

reaction of colloidal Si02 with MgAl204 to make Mg2Al4Si50j g in the presence of a bismuth oxide flux is presented in this paper. For purposes of comparison, data are also

presented for the crystallization of a pure, non-bismuth containing glass of the cordierite

composition.

EXPERIMENTAL

Powder Synthesis: A complete preparative description for this family of powders and

resultant ceramics appears elsewhere.^

The non-bismuth containing glass with the cordierite composition was obtained from

Gordon Brown of Stanford University and was prepared by conventional melt and quench

processing followed by grinding to a powder.

High Temperature XRD Description: X-ray diffraction data were collected with an automated powder diffraction system equipped with a high temperature attachment . ^

The cordierite precursor powder was spread in a thin layer onto a resistively heated

platinum band. APt/PtlO%Rh thermocouple, spotwelded to the heating band, was used

for temperature feedback to a Haider controller. The sample was heated with a ramp of

lO'C/min to the reaction temperature, and then held at that temperature until all measurements were completed. Diffraction data were collected using CuKaj radiation

(50kV, 36ma), a single crystal germanium incident beam monochromator,^^ and a position

sensitive detector.^^ The diffractometer was scanned from 9 to 39 degrees 2-theta at a speed of 2.5 degrees 2-theta/min after reaching the reaction temperature. Isothermal scans were repeated every 20 to 60 minutes imtil the reaction was complete.

246
The raw data were subjected to a trend oriented peak search to obtain peak positions

and integrated intensities.

Kinetics Determination: All data for kinetic analyses were obtained by scanning

repetitively at the desired temperature between 9° and 39°. Several peaks were selected for quantitative work. The peaks at 10.8^100), 26.6^112), 28.6^202), and 29.6't211) were

used to foUow the appearance of cordiente. The disappearance of the magnesium

aluminatespinel was tracked at 36.7°. The presence of the Si02 intermediates Bcristobalite and the substitutionaUy modified quartz were monitored at 35.7°and 26.5°

respectively.

For rate constant calculations the integration of count intensity over angle two theta(I) was

used as an expression of concentration. Simple first order rate law provided the best fit of

the data and was used to determine a rate constant for the appearance of cordiente. By

assuming a simple A to B reaction mechanism from a modified quartz

(1) [Intermediate]—^—cordiente, d(cordierite)/dt = k[Intermediate]

intermediate(vida infra) to cordiente as shown in equation (1), the rate expression

ln(l-Ij/I^)=-kt was derived and used to determine first order rate constants for the

appearance of cordierite. RESULTS and DISCUSSION

The results of tracking the powder diffraction pattern as a function of time at lOOO'C from

starting materials through Mg2Al4Si50j g formation appears in figure (1).

(2) 2MgAl204 + 5Si02-^^^^Mg2Al4Si50ig

Several features of these diffraction data should be noted. In the initial state (ie, lOOO'C,

t=0), the only crystalline species present is MgAl204, as the Si02 used

6000 I

rFigure 1 Figure 2

shows only an amorphous halo. As reported in the original paper/ there is an induction

247
period of several hours at lOOO'C before crystalline Mg2Al4Si50jg begins to appear. By

12 hrs reaction time, the conversion to Mg2Al4Si50ig is complete and there are no other

crystalline phases present. Two transient species are observed in the intermediate scans. The intermediate with the highest abundance indexes as 6-cristobalite and the less abundant

intermediate best indexes as a substituted quartz phase. Discussion of the significance of

the presence of these intermediates is left until later in this paper. To monitor the progress of the reaction and to allow the determination of reaction rates, plots of integrated peak intensity as a function of time for reflections of the spinel starting

material, cordierite product, and both quartz intermediates were made. These plots are shown in figure (2), with the indicated degree two theta. Qualitatively, several features are apparent from these plots. The spinel begins to disappear immediately, and then plateau

before resuming its disappearance. After some induction period B-cristobalite is observed

to grow in, reach a maximum, and then disappear. The stuffed quartz also grows in and

then disappears, but with an earlier maximum. Finally, the Mg2Al4Si50ig grows in

smoothly after an induction period of several hours.

The crystallization of cordierite from a glass at 960"^ as a function of time is illustrated in figure (3). This presents the growth of the 100 reflection of cordierite at 10.8°. The time

dependance of this crystallization is very different from that of the cordierite formation by

the bismuth oxide reaction. The glass crystallization shows two distinct stages; one with a half life of about 1 8 minutes and the other with a half-life of about 2 hours. One

interpretation of this is that the fast process is a nucleation step, while the slow step is the

consolidation of small crystalline domains into larger crystallites which scatter more

efficiently.

Figure 3 Figure 4

Using a first order rate law, rate constants for the appearance of Mg2Al4Si50^ g were

obtained from the integrated intensity vs time data for the reaction containing bismuth

oxide. Typical first order plots for Mg2Al4Si50jg appearance are shown in figure (4). Rate constants for the reaction at 975'^, lOOOX:, lOlSV, and 1075^ appear in table (1 ). It can be observed that the rate constants for Mg2Al4Si50ig appearance at different 20 are generally in good

248
TABLE 1

°20 (hkl) k(s-b T(^) °20(hkl) k(s-l)

975 10.8 (100) 4.0x10-5 1025 10.8 (100) 1.6x10-4

26.6 (112) 4.7x10-5 26.6(112) 1.9x10-4

28.6 (202) 4.7x10-5 28.6 (202) 1.5x10-4

29.6 (211) 4.2x10-5 29.6 (211) 1.8x10-4

1000 10.8 (100) 9.5x10-5 1075 10.8 (100) 2.4x10-4

26.6 (112) 9.2x10-5 26.6 (112) 3.4x10-4

28.6 (202) 1.0x10-4 28.6 (202) 4.9x10-4

29.6 (211) 1.1x10-4 29.6 (211) 4.5x10-4

agreement. The measured activation energy is 70 kcal/mole assuming an Arrhenius

behavior.

980 1 000 1 020 1 040 1 060 TEMPERATURE (C) Figure 5

The temperature dependence of the induction

period for the appearance of Mg2Al4Si50| g

is also shown graphically in figure (5). This curve

has a roughly exponential shape, with induction

period decreasing with increasing temperature. It was previously determined that below 960"^ this reaction does not proceed; that is, the induction

period goes to a very large value at 960°. These

two observations have some mechanistic

implications for this reaction. A likely reason that the reaction to make cordierite does not proceed

below 960° is that the bismuth silicate eulytite

(Bi4(Si04)3) melts between 955° and 960°. With

an excess of Si02 present it is expected that all the bismuth is tied up as this phase in the

solid state. The exponential decrease in induction period with temperature is also consistent with the picture of a bismuth silicate flux acting as a reactive solvent in the

reaction. Viscosity and diffusion are both exponential functions of temperature,^ ^ thus a reaction that relied on the transport through a flux to proceed would be expected to display an induction period that decreased exponentially with increasing temperature, as this one

does. There is also evidence that the flux is bismuth and silica rich. Where Si02 is left out of the reaction there is no effect on MgAl204 by bismuth oxide alone. On the other hand,

bismuth oxide is effective at crystallizing the amorphous Si02 to B-cristobalite. Thus it is the authors' conclusion that the induction period is due to the dissolution of MgAl204 in the bismuth silicate flux and the diffusion required to make a suitably substituted silicate precursor to crystallize to Mg2Al4Si50jg.

(3) Si02/Bi203(melt) + MgAl204-^[(Sii.xAlx)Mgx/202]-^Mg2Al4Si50ig + Si02

The distribution of intermediates in this system is somewhat different from either the case of sol-gel or glass ceramic prepared cordierite. Both the crystallization of a glass and the

249
crystallization of hydrolyzed precursors with the cordierite composition show the formation

of a modified, substituted quartz^'^ as an intermediate. This quartz solid solution is described as an aluminum substituted Si02 with magnesium in the channels to compensate

for the excess negative charge. This modified hexagonal quartz has a wide compositional

range around the cordierite composition with cell parameters varying smoothly with

composition. The reaction described here appears to have this substituted quartz as an

intermediate. This intermediate quartz does not, however, appear to be of the cordierite

composition. The observed intennediate has reflections essentially identical to a hexagonal, substituted quartz^^ with cell parameters of a=5.006A and c=5.459A. These

cell parameters match very well with those of a quartz solid-solution from reference(l)

measured at 750^ with aluminum and magnesium content of roughly half of

Mg2Al4Si50jg, which is to say it is very Si02 rich. This is consistent with the dynamic

x-ray measurements where less than half the MgAl204 disappears in the early stages of the

reaction where the intermediate forms . There is also a second, more abimdant intermediate

that appears in this reaction that does not appear in the other glass-ceramic prepared

cordierite. This intermediate indexes as B-cristobalite. It seems unlikely that this is a

reactive intermediate on the reaction path to cordierite.

The mechanism by which MgAl204 and Si02 convert to Mg2Al4Si50jg in the presence

of the bismuth oxide flux thus appears to be a complex one. The initial disappearance of

the MgAl204 is consistent with its dissolution in the flux. The bismuth oxide containing

flux likely is quite rich in Si02, as the lower temperature limit of this process of 960^

suggests the melting of eulytite, Bi4(Si04)3, which is in the range of 955"^ to 960X1.

Thus the process stops below about 960'C, where there is no melt present. The stuffed

quartz is clearly implicated in the conversion of starting materials to Mg2Al4Si50j g. In

addition to being an intermediate in the aforementioned glass and sol-gel routes to Mg2Al4Si50jg, the stuffed quartz concentration reaches a maximum just before the

appearance of crystalline cordierite in the x-ray powder pattern. As this modified quartz

does not appear to be of the cordierite composition, then the crystallization products must

be Mg2Al4Si50ig and free Si02. The other transient species, indexed as B-christobalite,

is likely not a reactive intermediate on the path to cordierite. On the contrary, the time

dependent x-ray data suggest that the formation of this crystalline species may inhibit the

reaction.The initial disappearance of MgAl204 and its subsequent reaction with Si02 are believed to yield the stuffed quartz phase. The sharp decrease in the rate of spinel

disappearance coincides with the appearance of crystalline 6-cristobalite. This is also

coincident with the plateau of the appearance of the stuffed quartz. The slowdown in the

rate of reaction is likely due to the conversion of high surface area, reactive Si02 to a low

surface area crystalline form.

Structural analysis is in progress to determine the form of the cordierite (disordered and

hexagonal or ordered and orthorhombic) and to locate bismuth in the final structure. ACKNOWLEDGMENT Thanks to Stefan Justi and Ken Schwartz for assistance with

the hot stage x-ray diffraction data collection.

250
REFERENCES

1) Schreyer, W. and Schairer, J. F., "Metastable Solid Solutions with Quartz-Type

Structures on the Join Si02-MgAl204," Zeitschriftfur Kristallographic, 1961, 116, tQ82

2) Gregory, A. G. and Vearey, T. J., "The Crystallizaton of Cordierite Glass,"/. Mater.

Sci., 1971,6, 1312-21

3) Zdaniewski, W., "DTA and X-ray Analysis Study of Nucleation and Crystallization of

MgO-Al203-Si02 Glasses Containing Zr02, Ti02, Ce02," J- Am. Ceram Soc,

1975,5S, 163-69

4) Bridge, D. R.; Holland,D.; and McMillan,?. W., "Development of the Alpha-Cordierite

Phase in Glass Ceramics for Use in Electronic Devices," Glass Technol. ,1985,26, 286-

92,

5) Watanabe,K. and Giess,E., "Coalescence and Crystallization in Powdered Highcordierite(2MgO-2Al203-5Si02) Glass,"/. Am. Ceram. Soc, 1985,68, C-102-3

6) Moyer.J. R.; Prunier, A. R.; Hughes,N. N.; and Winterton,R. C, Mater. Res. Soc.

Symp. Proc, 1986,73, 117-21

7) Gensse,C. and Chowdry,U., "Non-Conventional Routes to Glass-Ceramics for

Electronic Packaging," Mflfer. Res. Soc. Symp. Proc, 1986,75, 693-703

8) Bemier,J. C; Vihninot,S.; Rehspringer,J. L.; El Hadigui,S.; and Poix,P., "Sol-Gel

Processes and Synthesis of Dielectric Powders for Multilayer Ceramics," High Tech

Ceramics, P. Vincenzini, ed, Elsevier, Amsterdam, 1987, pp 1443-50

9) Dupon,R. W.; McConville.R. L.; Musolf,D. J.; Tanous,A. C; and Thompson.M. S., "Preparation of Cordierite Below lOOO'C via Bismuth Oxide Flux," /. Am. Ceram. Soc,

1990,75,335-339,

10) JCPDS file#12-708

11) R. H. Perry and C. H. Chilton ed.,Chemical Engineers Handbook, 5tt edition,

McGraw Hill, 3-247, 1973

12) Siemens D-500

13) Anton Paar, K.G., P.O. Box 17, Graz, Austria

14) Haider Elektronik, Munich, FRG

15) Huber Diffraktionstechnik, D-8219 Rimsting, FRG

16) PSD-OED50M, Braun, Gutenbergstrasse 3, D-8046 Garching, FRG

251

CRYSTAL CHEMISTRY

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

SOLID STATE CHEMISTRY OF PEROVSKITE AND RELATED OXIDES

CONTAINING Fe'*'^, Ni^"^, and Cu^"^

Mikio TAKANO

Institute for Chemical Research, Kyoto University,

Uji, Kyoto-fu 611, Japan

1. INTRODUCTION

Perovskites (AMO3) and related oxides allow 3d transition

metals (M) to be stabilized in unusually high formal valencies:

Iron is oxidized to Fe'*^ in SrFe03. And mixed valence states are rather easily produced by substituting La^"*" for Sr^"*", for example,

or by changing oxygen content. The 3d levels of an M ion lowers in

energy as the atomic number of M increases and as the formal

valency of M increases. This makes the M-0 bond covalent, and thus

the role of the oxygen lattice and, in turn, the role of the A

cations bound to the oxygen atoms become important in determining

the electronic properties of a material.

SrFeOg is a metallic conductor, while CaFeOg is semiconductive, though the cell volume is smaller for CaFeOg. Discrete

vacancy-ordered phases SrFeOg.j^^j^ (n=8,4,2) are all semiconductive, even in their high temperature disordered phase.

Low dimensional M-0 lattices formed in intergrowth structures

generally show polaronic electric conduction. There is, however, a semiconductor-metal transition in Lag.^^Sr^^NiO^ . The transition

temperature T^ decreases monotonically with increasing x but with

a change in dT^/dx at x~0.6. Correlated with this is an anomalous

dependence upon x of structural details, especially of the

La(Sr)-0(2) bondlength. Composition x=1.4 remains metallic at

least down to lOK.

Oxidation and reduction of [Cu02]^~ layers in intergrowth

structures can give rise to high T^ superconductivity and the T^

varies depending upon details concerning composition and structure. Among the experimental parameters controlling the phase

formation and crystal structure, influences of high pressures have

relatively been unknown. Examples of layered phases stabilized

only under high pressures will be shown.

2. PEROVSKITES CONTAINING Fe'*'^

255
1) SrFeOg and CaFeOg

SrFe03 retains cubic perovskite structure (a=3.850A at 300K)

and am etallic conductivity of ~10~^Qcni down to 4K [1]. The

magnetic moment estimated from magnetization measurements on SrFe .jjCOjjOg , which is ferromagnetic for 0.2<^x<_l, is

// pg4+=3.72/^ g [2]. A screw spin structure with a propagation

vector along the <111> direction has been found below Tj^=134K [3].

This structure results from a competition of the ferromagnetic

nearest neighbor interaction and the antiferromagnetic second- and

fourth-nearest interactions [4] : The angle between the nearest Fe

moments is only about 40° [3].

CaFeOg crystallizes in a tetramolecular tetragonal cell with

a=5.325A and c=7.579A, and Tj^=115K [5]. In contrast with SrFeOg,

semiconductive behavior has been found below room temperature,

though there is some uncertainty caused by the small size of

samples prepared under 2GPa. Further details are under

investigation.

2) Microscopic electronic state of Ca-j^.j^Sr^jFeOg

Mossbauer spectroscopy has proved useful to elucidating how

the electronic state depends upon composition. The spectra of

SrFe03 [6] are simple and beautiful in which the unique electronic

state and high crystal symmetry are reflected. A narrow single

peak with an isomer shift (IS) of 0.05mm/s relative to a -Fe

appears at room temperature, and a single magnetic pattern is seen below Tf^. The hyperfine field (Hi) is 33. IT and IS=0.146mm/s at

4K. Any quadrupole interaction has not been detected.

CaFeOg shows a single peak at 300K with IS=0.073mm/s just as for SrFe03, while, at 4K, two sets of magnetic patterns, I and II,

with (Hi, IS)=(41.6, 0.34) and (27.9, 0.00) appear. And the

intensity ratio of the two sets is 1:1 within experimental error

as seen in Fig. 1(a). Quadrupole interactions are absent. The

great difference of the parameters between the two components

clearly indicates that an electronic phase transition

discriminating iron into two different states occurs below 300K.

Considering that both these parameters are sensitive to the d

electron density of an Fe ion, a disproportionation model

2Fe4+ (tgg^ cy*^)5>Fe3+(t2g\2)+Fe^+(t2g^)

has been proposed, components I and II corresponding to Fe and

Fe^"*", respectively [7]. There are three important consequences of

this model. The first is that high spin state was assumed for

Fe'*'*'. The (nearly) delocalized character of the 6* electron of

eg-orbital parentage has been understood to yield no quadrupole

interaction. The spherical configurations of Fe^"*' and Fe^'*" have

also been done so. This assumption was strongly supported by the more recent magnetic measurements on BrFe-j^.j^COj^Og mentioned above.

The second is the high content of Fe^"*". Iron in this high

oxidation state was known at that time only as an impurity of ppm

256
(a) 300K

H 1 1 1 1 1 1 3 2 10 1 2 3 velocity/mms'

* ; *

• '

' .

• • >'•' • .

I I ' —

1

(I)

-I—t—I—I—I—I—I—I—f- -8 -6 -A -2 0 2 A 6 8

velocity/mms -1

4 0.0

30.0 -

I 1 r- 0.5

-1 1 . '0

[ -A A

aoe?-— D II

Fig. 1 Mossbauer spectra of CaFeOg (a) and IS and Hi of

Caj^.jjSrjjFeOg at 4K (b).

order in SrTiOg [8]. The Hi value for component II of CaFeOg at

4K just coincides with the corresponding EPR data for the

^'^Fe^'^'/SrTiOg sample. Later Buffat et al . succeeded in

stabilizing LaLig 5Fe^''"Q 5O3 under a high oxygen pressure [9]. The

third is the transition mechanism.

Ca-j^.jjSrjjFeOg can be prepared for the whole range of x. The

cell volume monotonically increases with increasing x, while the

structure symmetry becomes cubic for X20.25. Shown in Fig. 1(b) are the composition dependence of the parameters at 4K [10]. To be

noticed is that the two components gradually become similar to

each other as x increases, while the mean parameter values,

(I+II)/2, remain almost constant. The iron valencies may be

expressed as Fe^^"'^ and Fe^^"*"*^ 6 varying continuously from

S =1 at x=0 to <S =0 at x=l. Charge density wave (CDW) may be more appropriate in terminology than disproportionation.

Just the same kind of disproportionation has been found also

in Sr-j^.yLayFeOg [11,12] as can be typically seen in the spectra

for y=0.3 reproduced in Fig. 2. As a result of a rapid electron

exchange the Fe ions are in an average valence state at room temperature for ylO.6 so that IS linearly varies with y as IS(y)=(l-y)IS(0)+yIS(l),

where IS(0)=0.05 mm/s (SrFeOg) and IS(1)=0.39 mm/s (LaFeOg). On

the other hand, at 4 K, the spectrum for x=0.3 is composed of two

components with (Hi, IS)=(46.0, 0.36) and (26.9, -0.05) with a relative intensity of 1:11=66:34. The low temperature state can therefore be formulated as SrQ yLaQ sFe^'^'g 55^® 0 35^3*

257
-4-2 0 2 4

V (mm/s)

.—I 1 1 1 1 1 1 1 »— -8-6-4-2 0 2 4 68

Velocity (mm/sec)

Fig. 2 Mossbauer spectra of STq yLaQ 3Fe03,

3) CDW, why ? It is quite natural to suppose that the disproportion should

be accompanied by a lattice distortion: The oxygen octahedron

containing Fe^"*" should be enlarged and the one containing Fe^"*" should be compressed. However, diffraction measurements using Xray (XRD) and neutron (ND) [13] have failed to detect any sign of

such a distortion. The present author has suspected that the tran

sition is set in microdomains which are so small that these diffraction measurements detect only seemingly high, averaged symmetry. Electron diffraction (ED) and transmission electron microscopic (TEM) studies at low temperatures are of great interest and

are now in progress.

A theoretical calculation has been done for monometric

([FeOg]^") and dimeric ( [FegO^j^]"-*-^) clusters by a DV Xa method

[14]. These clusters are embedded in a Madelung potential

reflecting the perovskite structure. For either octahedron in Fig.

4(a) the Fe-0 distance is 1.887A, half the cubic root of the cell

volume for CaFeOg. The effective charges and the magnetic moments

are given in the figure. The electron configuration is essentially

such that both the d^2_y2 and orbitals are half occupied,

which is high spin state. These orbitals each contain an oxygen 2p

component weighing about 28%, while the weight is about 14% for

the t2g orbitals. The Fe ions carry a moment of Z.lju^, in good

agreement with the extrapolated magnetization data for SrFe03. On

the other hand, the Fe(I)-0(l,3) distance is elongated but the

Fe(II)-0(2,3) distance is shortened, each by 0.053A, in Fig.

3(b). As a result of this displacement the configurations are

258
'^['ll^

• ^-887 A d(Fe(I)-0) : 1.940 A

' ' 3.7ug d(Fe(n)-0) : 1.834 A

Fe(II)29^ 3.2UB

0(1)1-77- )1-88'

0(3)^-3''-

Fig. 3 [FegO-j^-j^] '^ cluster before (a) and after (b) an anion

displacement

modified as a^2_^2^a^2^ for Fe(I) and d 2_y2^a^2^-^ for Fe(II).

The electron transfer from Fe(II) to Fe(I) reflects the opposite

shifts of the level schemes for these ions. The moments differ by

0.9// g, and the hyperfine fields do by IIT according to a preliminary estimation. The experimental value is 13. 7T for CaFeOg and

19. IT for Srp yLag 3FeO^. The composition dependence of the d

electron densities seen in Ca-j^.^Sr^^FeOg and Sr^.yLa^jFeOg can thus

be assigned to a smooth, composition-dependent anion displacement.

How the disproportionation can be stabilized has been ex- plained by Goodenough [15]. The d-orbital contribution to the

Fe-0 interaction energy, 0 , is assumed to have the Morse form.

And the change in binding energy due to the anion displacement is

divided into two parts as A £ =-(A-B)u2,

where u is the anion displacement normalized to half the Fe-Fe

separation, coefficient A is proportional to d <j> I , and Bu is

the elastic energy contribution. For the stabilization of

disproportionation A>B is necessary. For CaFeOg, in comparison

with SrFeOo, the Fe-0 distance is shorter and the cr

' band is

narrower because of the more acidic nature of Ca . These factors

have been argued to be more advantageous for CaFe03 than for

SrFe03 to satisfy the above criterion of the disproportionation

stabilization.

4) CaFeOg and SrQ yLag 3Fe03 in the transition temperature range

As seen in Fig. 4(a), the paramagnetic single peak of CaFeOg

begins to be divided into two below 290K [16]. The splitting has

259
been interpreted to indicate the disproportionation taking place

in the paramagnetic state. The separation reaches 0.32mm/s at

140K (Fig. 4(b)), which just coincides with the difference in IS

between components I and II at 4K. Another reason why this

splitting should not be considered as a temperature-dependent

quadrupole interaction is that components I and II are not

perturbed by quadrupole interactions at 4K. Below Tjq a pair of

magnetic patterns appear and the Hi's vary as shown in Fig. 4(c).

Thus, the disproportionation proceeds gradually with decreasing

temperature, and the magnetic transition follows.

Quite a different type of transition occurs for

SrQ yLaQ 3Fe03 and also for CaQ c^STq 5Fe03. For either phase a paramagnetic single peak and a

' pair* of magnetic patterns are superimposed in a very wide temperature range indicating a firstorder transition between the paramagnetic average valence state

and the antiferromagnetic disproportioned state.

In conclusion Ca-j^_jjSrjjFe03 and Sr2_yLayFe03 show very interesting phase transition behavior depending upon composition. The

cr electrons in SrFe03 tend to become localized by forming CDW,

the amplitude being larger as Sr is replaced more by Ca and La.

It should be noted that the replacement by these elements reduces

the covalency of the Fe-0 bond.

Experiments in progress include measurements of XPS,

Mossbauer spectroscopy under high pressures, and preparation of

single crystalline films to be used for precise electric and

magnetic measurements.

Fig. 4 Mossbauer spectra of CaFeOo (a) and the parameters (b).

260
5) SrFeOg.i/n

The system SrFeO^, 2 . 5 <.z <.3 , forms a continuous solid

solution at temperatures T>.T^(z), but a series of discrete

ordered-vacancy phases BrFeOg,/-, /^^s (n=8,4,2) appear below

T^(n8)=523K, T^(n4)=598K, and T^(n2)=1103K. The order-disorder

transition, ordered structures, n-dependence of the electronic

state were reported in refs. 18 and 19.

3. Lag-xSr^NiO^COixlLS)

Incorporation of excess oxygen and substitution of Sr for La

both introduce a mixed valency in the Ni02 planes by lowering the

Fermi energy. La2NiO^^^ shows an orthorhombic-tetragonal

transition and a change in thermal expansion coefficient from

weakly negative (T<T^) to strongly positive at T^~650K [20]. And a semiconductor-metal transition in the basal plane has been

observed at (or near) T^ [21]. The metallic temperature region is

widened by the Sr-for-La substitution [22].

1) La-2Ni04^-|_/2n " O^^Sen ordering

According to a recent ND study [23], there exists a two phase

region in the compositional range 0.02<d<0.13 and the excess oxygen in the (5>0,13 phase has been located as an interstitial

oxide ion in the tetrahedral site of the La202 layers. However, a more microscopic study using ED and TEM [24] has revealed

that there are two types of homologous series of oxygen-ordered

phases. These are 2a>4caxlc (k,l=n or 2n depending upon n; the most

frequently observed sets are k/2=l/2=n=2 and k=l/2=n=3) and

y2ax pv^2ax qc(most frequently p = 2 and q=l), where ax ax c stands

for the KgNiF^ type cell.

The most probable ordering models are illustrated in Fig. 5. In Fig. 5(a), the interstitial oxygen atoms are aligned along the

[111] direction within a (Oil) plane. The rows of interstitial

atoms are separated from each other by 2a along the a axis. And

these specific (Oil) planes appear with a period of na along the b

axis. According to this model, <5=l/2n, i.e. (5=0.25 for n=2 and

5=0.17 for n=3. In Fig. 5(b), the interstitial atoms are separated from each other by V~2a along the [llo] direction and

by p/2a along the [110] direction within a basal plane. For p=2,

the [llo] excess oxygen rows are shifted by y~2a along the [110]

direction in adjacent basal planes to from a body-centeredorthorhombic lattice of >r2ax 2>r2axc. For p>2 a monoclinic

superlattice seems to result. Excess oxygen content is again

5=l/2p. This type of excess oxygen ordering along the [110]

direction is rather similar to that causing the structure

modulation in the superconductors of the Bi-Sr-Ca-Cu-0 system.

There are experimental suggestions that excess oxygen atoms

are rather mobile and become ordered at low temperatures, not much

261
Fig. 5 Model structures of the 2axkaxlc type (a) and the

y2ax 2y2ax c type (b). Large and small open circles in

(a) are the interstitial and normal oxygen atoms,

respectively, and solid circles and squares indicate La

and Cu atoms, respectively. In (b) large and small

circles reprisent interstitial oxygen atoms at z=0 and

c/2, respectively.

higher than room temperature. So, special attention should be paid

in studying properties as a function of oxygen content and

temperature

.

2) La2_xSrjjNi04

Very recently samples with x<_1.6 were prepared and examined

by various physical measurements [25]. The oxygen content was controlled by using suitable oxygen pressures (10-10"^ Pa) depending

upon X. The Ni-0 and La(Sr)-0 bondlengths, determined by Rietveld

analysis of powder XRD data, are shown Fig. 6. Ni-0(2) (apical

oxygen) bondlength decreases essentially linearly with increasing

X throughout the interval 0<.x<.1.4, whereas the La(Sr)-0(2)

bondlength increases linearly with x only in the interval

0lx_<0.6. Thus a maximum in the c-axis and the lattice axial

ratio c/a appears at x=0.6 reflectng the very anisotropic change

in the La(Sr)-0 bondlength in the interval 0lx<.0.6. The axial

ratio Ni-0(2)/Ni-0(l) of the NiO^ octahedra, on the other hand,

decreases linearly with increasing x throughout the interval

0<_x<1.4 with no obvious anomaly at x=0.6. An important shift in

the relative strengths of the La(Sr)-0(2) and Ni-0(2) bonding is

thus suggested.

262
o

L 1.0

»; 2.0

z 1.9

"1 I I I

1 1 1 1 r

(a) .

N1-0(2)

Q o o OOOOO

N1-0(1)

-J I I ] I I I I L. 0.0 0.4 0.8 1.2 1.6 X

A

2.8

2.7

2,5

2.5

2A

2.3

2.2 -

T—I

1 1 1 1—I—I—

(b)

^ i ^ * °°Ooo

La(Sr)-0(2)//t110]

^ ^ ^ a A A La(Sr)-0(1

)

O I)

La(Sr)-0(2)//c

0.0 0.^4 O.J X1.2 1.5

Fig. 6 Variation of the Ni-0 distances and the tetragonal ity of

the NiOg octahedron (a) and the variation of the La(Sr)-0

distances (b).

0.5

0.0

-0,5

-1.0

-1.5

-2,0200

O A

-I 1 1 r- (a)

O XrO.O A X:0,2

X=0,'t

• X:0.8

oo oooo

400 600 T(K)

800 1000

Fig. 7 Resistivity vs. T (a) and T.^. vs. x (b) data for

La2_xSrxNi04

263
5) SrFe03_i/ji

The system SrFeO^ , 2.5<.z<.3, forms a continuous solid

solution at temperatures T>^T^(z), but a series of discrete

ordered-vacancy phases SrFeOg./ -i /^^v (n=8,4,2) appear below

T^(n8)=523K, T^(n4)=598K, and T^(n2)=1103K. The order-disorder

transition, ordered structures, n-dependence of the electronic

state were reported in refs. 18 and 19.

3. La2_xSrxNiO4(0<x<1.6)

Incorporation of excess oxygen and substitution of Sr for La

both introduce a mixed valency in the Ni02 planes by lowering the

Fermi energy. La2Ni04^^ shows an orthorhombic-tetragonal

transition and a change in thermal expansion coefficient from

weakly negative (T<T^) to strongly positive at T^~650K [20]. And a semiconductor-metal transition in the basal plane has been

observed at (or near) T^ [21]. The metallic temperature region is

widened by the Sr-for-La substitution [22].

1) La2Ni04^-|^/2n * Oxygen ordering

According to a recent ND study [23], there exists a two phase

region in the compositional range 0.02<<5<0.13 and the excess oxygen in the (5>0.13 phase has been located as an interstitial

oxide ion in the tetrahedral site of the La202 layers. However, a more microscopic study using ED and TEM [24] has revealed

that there are two types of homologous series of oxygen-ordered

phases. These are 2ax kax ic (k,l=n or 2n depending upon n; the

most frequently observed sets are k/2=l/2=n=2 and k=l/2=n=3) and

V2axpv'~2axqc (most frequently p=2 and q=l), where ax ax c stands

for the KgNiF^ type cell.

The most probable ordering models are illustrated in Fig. 5.

In Fig. 5(a), the interstitial oxygen atoms are aligned along the

[111] direction within a (Oil) plane. The rows of interstitial

atoms are separated from each other by 2a along the a axis. And

these specific (Oil) planes appear with a period of na along the b

axis. According to this model, d-l/2n, i.e. 5=0.25 for n=2 and

(5=0.17 for n=3. In Fig. 5(b), the interstitial atoms are separated from each other by /2a along the [110] direction and

by p/2a along the [110] direction within a basal plane. For p=2,

the [110] excess oxygen rows are shifted by /"2a along the [110]

direction in adjacent basal planes to from a body-centeredorthorhombic lattice of /'2ax 2/'2axc. For p>2 a monoclinic

superlattice seems to result. Excess oxygen content is again

d=l/2p. This type of excess oxygen ordering along the [110]

direction is rather similar to that causing the structure

modulation in the superconductors of the Bi-Sr-Ca-Cu-O system.

There are experimental suggestions that excess oxygen atoms

are rather mobile and become ordered at low temperatures, not much

264
Typical resistivity data are given for T>300K in Fig. 7(a).

The data for x=0,6 is clearly anomalous above 400K and is

transitional between compositions exhibiting a rather abrupt

semiconductor-metal transition at T^. in the interval 0lx<0.6, a smooth transition in the interval 0.6<x<_1.2. A plot of T^ vs. x, Fig. 7(b), shows two straight-line segments intersecting near

x=0.6. This feature and the structural anomaly mentioned above

suggest that there may be a significant change in the electronic

state of the system at and below room temperature near the

composition x~0.6 [25].

These results have been interpreted as follows. The semiconductor-metal transition at T^ in LagNiO^^^ is due to a localized-electron to itinerant electron transition among the

antibonding cr electrons. Below T^, this material is a smallpolaron conductor and has a negative thermal-expansion coefficient

consistent with increasing electron delocalization with increasing

temperature. The electronic transition and the structural

transition may be separated in oxygen-rich samples, the structural

transition occurring at T^'<T^. In the Lag.xSrj^NiO^ system, the

position of the 0(2) atom shifts at room temperature away from its

La(Sr) nearest neighbor toward its Ni nearest neighbor in the

interval 0_<x<.6 where the resistivity remains polaranic at room temperature; this shift stops for x>0.8 where the roomtemperature conduction is metallic. This unusual behavior is con- sistent with an 0(2) -displacement and polarization controlling a transition from localized-electron to itinerant-electron character

at T^ in the interval 0<.xi0.6.

Hall effect measurements for samples in the range 1.0<.x<.1.4

indicated that higher-mobility hole carriers are introduced; the

Fermi energy drops into the n band at room temperature for

X21.0 in accordance with equation Ni^''"+0^~

^

Ni^"'"+0~ being

shifted to the right.

4. HIGH PRESSURE SYNTHESIS OF NEW Cu OXIDES CRYSTALLIZING IN

LAYERED STRUCTURES

All the high Tc superconducting cupric oxides contain twodimensional (2D) CuO^ sheets irrespective of the sign of the

carrier [26, 27]. It is an important task for solid state chemists

to continue efforts to find new cupric oxides crystallizing in

layered structures. For this purpose the author's group has studied how high pressure treatments influence the phase formation and

crystal structure. Fine oxide powders containing appropriate metal

elements were treated under 6GPa at 1300K.

It has been found that ACUO2 (A: Ba-|^y3Sr2/3~Sr-|^y3Ca2/3

)

crystallizes in a 2D structure made of alternately stacked A and

[Cu02]oo layers [28]. This structure was reported for the first

time for Car, o^Srr, icCuOo [29]. Under ambient pressure, however,

265
only a narrow composition range around Ca:Sr~9:l can be stabilized

in this structure. For example, Cu-O-Cu zigzag chains are formed

in SrCuOg in which CuO^ squares share their sides, not corners. The high-pressure phase of SrCu02 has a higher specific density.

As seen in Fig. 8(a) the lattice changes its size very

anisotropically as a function of r^. At relatively small r^'s both

the Cu-0 and A-0 bonds are stretched linearly as rA increases,

while the rates are d(Cu-O)/dr^=0. 271 and d(A-0)/dr^=0.896,

reflecting the relative stiffness of the Cu-0 bond. The

anisotropic lattice expansion results from this difference. The

expansion tends to be suppressed for r^>1.2lA, more strongly for

the Cu-0 bond. The maximum length reaches 1.97A, which coincides

with the "ideal" length calculated from the table in ref. 30. If

CaCuOg were stabilized, the lattice constants would be a=3.849A

and c=3.160A, i.e. the Cu-0 bondlength would be 1.92A.

RgCuO^'s containig relatively small R^"*^ ions Y, Dy, Ho, Er,

and Tm have also been found to be stabilized under the pressurized

conditions [31], though these are not formed at all under ambient

pressure. The structure is the Nd2Cu04 (T') type with lattice

constants plotted against the R ion radius in Fig. 8(b). It is

interesting to notice that two straight-line segments intersect at

R=Sm. For the R ions smaller than Sm, the [CuOgloo layers probably

become slightly distorted for the sake of lattice matching with

the R2O2 sheets: The tolerance factor is too small to maintain

genuinely tetragonal symmetry. In support of this ED and TEM

studies of Y2CUO4 have revealed a complicated superlattice

structure. And magnetization measurements on the R=Y and Ho phases

have indicated weak ferromagnetism due to antisymmetric exchange

0.400-

'a^ Ionic radius of R (nm)

Fig. 8 Lattice constants and the A-O bondlength of ACUO2 (a) and

the lattice constants of Rr^CuO^ (T' ) (b).

266
interactions caused by the distortion [32]. The saturated weak

ferromagnetic moment is 9xio~^//g/Cu for Y2CUO4. If an atomic

moment of 0.5/Zg/Cu is assumed, the canting angle is calculated to

be 0.05°.

The tolerance factor would approach unity, if the [Cu02]oo

layers are contracted by oxidation. Injection of carrier holes,

not electrons, may thus be possible. If the data for the two

different structures in Fig. 8 are compared with each other, one would notice that the Cu-0 bondlength in these structures overlap

each other for a considerably wide range of R and A ions.

^2\i-1^^^2m+2 ""^y ^® formed from appropriate combinations.

REFERENCES

J. B. MacChesney et al., J. Chem. Phys. 43 (1965) 1907.

T. Takeda et al., J. Mag. Mag. Mater, 31-34 (1983) 797.

T. Takeda et al .

, J. Phys. Soc. Jpn. 33 (1972) 967.

T. Takeda et al., "FERRITES: Proc. Int. Conf. Sept. -Oct.

1980 Japan", (Cen. Academ. Pub., Japan, 1981) p. 385.

Y. Takeda et al.. Mat. Res. Bull. 13 (1978) 61.

P. K. Gallagher et al., J. Chem. Phys. 41 (1964) 2429.

M. Takano et al., Mat. Res. Bull. 12 (1977) 923.

K. A. Muller et al., Solid State Commun. 9 (1971) 1097.

B. Buffet et al., C. R. Acad. Sci. 292, Ser. II (1981) 509.

Y. Takeda et al., J. Phys. Colloq. 40 (1979) C2-331.

M. Takano et al., J. Solid State Chem. 39 (1981) 75.

P. D. Battle et al., J. Solid State Chem. 77 (1988) 124.

P. D. Battle et al., J. Solid State Chem. 84 (1990) 271.

H. Adachi et al ., in preparation.

C. Gleitzer and J. B. Goodenough, Structure and Bonding 61

(1985) 1.

M. Takano et al., J. Phys. Colloq. 40 (1979) C2-313.

T. Shinjo et al. "FERRITES: Proc. Int. Conf. Sept. -Oct.

1980, Japan" (Cen. Acad. Pub., Japan, 1981) 383.

Y. Takeda et al., J. Solid State Chem. 63 (1986) 237.

M. Takano et al., J. Solid State Chem. 73 (1988) 140.

C. P. Tavares, Mat. Res. Bull. 20(1985) 979.

C. N. R. Rao et al., J. Solid State Chem. 67 (1984) 266.

M. Khairy et al., J. de Physique 47 (1986) Cl-831.

J. D. Jorgensen et al., Phys. Rev. B38 (1988) 11337.

Z. Hiroi et al., Phys. Rev. B41 (1990) 11665.

Y. Takeda et al.. Mat. Res. Bull. 25 (1990) 293.

G. Bednorz and K. A. Muller, Z. Phys. B64 (1986) 188.

Y. Tokura et al., Nature 337 (1989) 345.

M. Takano et al . , Physica C 159 (1989) 375.

T. Siegrist et al.. Nature 334 (1988) 231.

R. D. Shannon , Acta. Cryst. A32 (1976) 751.

H. Okada et al., Physica C 166 (1990) 111.

H. Okada et al., Phys. Rev. B in press.

267
DISCUSSION

M. Norton : Have you studied the pressure dependence of CaFe03 to find if it becomes

metallic?

M. Takano: Mossbauer measurements under pressure are being done now. I suppose

that the charge disproportionation will fade away and that a metallic state results.

A. Navrotsky : Are the oxygen contents of SrFe03 and CaFe03 controlled by the

synthesis conditions? Did you ever get more than three oxygens?

M. Takano: CaFe03 was prepared from CaFe02 5 (Ca2Fe205) sealed in a gold tube

together with an oxygen releaser, KCIO4: the tube was heated at about 1000°C under 20

kb. SrFe03 needs an oxygen pressure above 500 b. For the both the oxygen content

does not go beyond three.

R. J. Cava: What are the magnetic properties of CaFe03?

M. Takano: It is an antiferromagnet with a Tj^ = 115 K. The neutron diffraction

pattern is quite similar to that of SrFeOj. It suggests that CaFe03 also has a screw spin

structure resulting from a competition between the ferromagnetic nearest neighbor

interaction and the antiferromagnetic second- and fourth-nearest neighbor interactions as

in SrFe03.

A. Nath : Did you observe Fe'*"'" in low spin state in any other compound?

M. Takano: Yes, in an oxygen deficient phase, SrFe02 75. An oxygen vacancy creates a pair of Fe^"*" ions on either side of it, while the Fe ions coordinated octahedrally

strongly attract oxygen ions and the spin state becomes low spin state. Even in the

magnetically ordered state of SrFe02 75 the low spin Fe"^"*" ions show only a broad single

line in the spectrum.

P. D. Battle : How do you recognize Fe'*"'' in low spin state?

M. Takano: From the (nearly) nonmagnetic character seen in the Mossbauer spectrum.

A low spin Fe^"^ has S = 1 and L = 1, and a nonmagnetic ground state with J = 0

results.

P. D. Battle : But the moment would increase quickly as a function of temperature,

wouldn't it?

M. Takano: Mossbauer measurements on SrFe02 75 were made only at 300 K and 4 K.

At 4 K the low spin state ions shows only a broad single line.

268
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

IMPURITY INCORPORATION MECHANISMS IN LiNb03

Dunbar P. Birnie, III

Department of Materials Science and Engineering

University of Arizona, Tucson, AZ 85721

ABSTRACT

Impurity incorporation mechanisms in lithium niobate

are discussed. It is concluded that the simplistic

electrical charge misfit considerations may in fact

control site occupancy for impurities. Divalent Fe, Mg,

and Mn are found to sit on lithium sites. Trivalent Fe

sits on lithium sites and hexavalent W sits on niobium

sites. The general observation of 3-fold axial defects

in EPR is attributed to the internal electric field

aligning point defect clusters on cooling.

INTRODUCTION

The impurity incorporation mechanism often controls

the behavior of impurities in crystals. This is

especially important for impurities such as Fe and Mg

which have technicological importance in holographic

storage and optical waveguide device studies.

Iron is interesting because of its two valence

states that can be changed using oxidation and reduction

treatments. This behavior activates the photo-refractive

effect (PRE) that is important for optical damage in

waveguide devices.

Magnesium is important in reducing optical damage

in optical waveguide devices and for growing high quality

single crystals without striations.

Information about other impurities may help shed

light on general trends related to impurity

incorporation. Therefore, this paper will review what is

known about impurity substitution in lithium niobate.

Important in this discussion is knowledge about the

native point defects that dominate this system. This will

give us a basis for understanding what possible charge

compensation mechanisms can occur when defects dissolve

into lithium niobate.

269
UNDOPED LITHIUM NIOBATE

The low temperature, ferroelectric, phase of lithium

niobate is a derivative of the hexagonal corundum

structure of Al^Oj. The lithium and niobium atoms

substitute for the two octahedral Al atoms in the

structure. The substitution is performed such that each

near-neighbor Li-Nb pair are aligned in parallel. Thus,

along the c-axis there are columns of octahedral sites

which are 2/3 filled [1]. The arrangement causes the

crystal to have some net spontaneous polarization

characteristic of ferroelectric materials.

All three of the possible octahedral sites (lithium

site, niobium site, and vacant site) along the c-axis are

quasi-octahedral and thus would form quite similar

spatial environment for impurity atoms.

Lithium niobate exhibits a large range of

nonstoichiometry . This behavior is indicative of large

levels of point defects being present in the crystal.

One interesting aspect of the nonstoichiometry is that

it occurs only to the Nb-rich side of the phase diagram

[2-5]. The extent of nonstoichiometry can be up to

several percent, depending on the temperature of

equilibration

.

The mechanism for the nonstoichiometry in this

system has been determined using single crystal X-ray

diffraction [6]. It was found that at room temperature

the niobium excess is accomodated with niobium atoms

sitting on lithium sites (NbLi) combined with vacant

niobium sites (V^j,) . The concentrations of the two defects

are present in levels such that global charge neutrality

is maintained in the crystal. Both defects may be present

in large concentrations because of the wide single phase

field. The niobium antisite and niobium vacancy have net

charges of +4 and -5 respectively.

One interesting consequence of these large native

point defect levels is that under most circumstances

lithium niobate will be intrinsic; impurity defects

present at low concentrations will often be insignificant

compared to the native defect concentrations. Only when

the doping level becomes larger than the nonstoichiometry

will the impurity start to define the defect chemistry.

This is important because for low level doping we will

expect that the impurity will substitute for only one of

the two sites rather than a mixture of the two sites.

It is conspicuous that there is little or no

nonstoichiometry to the Li-rich side of the phase

diagram. This is compatible with the above stated defects

being present and other defects being difficult to form

270
in the lattice.

With these large concentrations of charged point

defects we may expect some large degree of defect

association during cooling from high temperature. This

rearrangement and clustering of point defects in crystals

is quite common in ionic crystal structures. Clustering

can occur between the native point defects and also with

the impurity atoms at whatever level they are present.

The ferroelectric nature of lithium niobate may

affect the orientation of defect clusters that may form

on cooling. The parallel alignment of the Li-Nb pairs in

the structure defines some spontaneous polarization along

the c-axis direction in the lattice. This generates an

internal electric field in the opposite direction which

could orient defect clusters. This can occur because a

clustered defect pair will consist of one positive and

one negative defect, bound electrostatically which will

respond as a small electrical dipole. The positive

defects will tend toward the negative c-axis direction

and the negative defects will tend oppositely.

In summary, we expect significant concentrations of

niobium atoms sitting on lithium sites that are charge

balanced by vacant niobium sites. These will be charged

positive and negative, respectively. For low doping

levels these two defects will define the bulk charge

neutrality at room temperature. But, impurity defects

that may be present in low levels may become

electrostatically bound to these defects, creating

cluster defects after cooling from the crystal growth

condition. Also, it is possible that some preferential

orientation of defect clusters will occur such that

dipolar defects are aligned along the c-axis direction

in the crystal.

The following sections will describe information

relating to the preferred incorporation site for several

impurities in lithium niobate. Clustering and charge

balance considerations will be discussed for all cases.

The most conclusive data comes from X-ray diffraction

crystal structure determinations.

IRON

The most abundantly studied impurity in lithium

niobate is iron. This is because of its importance as an

identified cause of optical damage in waveguide devices.

Part of the optical damage problem that arises with iron

is that it has two stable oxidation states (either 2+ or

3+) . In principle these two ions could substitute in

271
different locations, however, it appears that both occupy

lithium sites in the lattice. Field annealing has shown

that both ferrous and ferric ions migrate toward the

negative electrode, showing that they are indicates that

they both substitute for Li [7].

For either valence of iron, substitution onto the

lithium site creates a net-positive defect. At low levels

of iron present the material is intrinsic so no change

in the bulk charge neutrality condition will occur.

However, there could be some electrostatic clustering of

the positive iron defects with the negative defects

that will be present. If the internal electric field

caused these to align primarily oriented along the caxis, then the many of axially symmetric EPR

identifications can be explained.

MAGNESIUM

Little information about the magnesium atom site is

available because it has only one stable valence state

which is not paramagnetic. However, one single-crystal

X-ray diffraction structure determination has shown that

the magnesium also substitutes for lithium [8]. The

sample used had a magnesium concentration that was above

"threshold" and showed magnesium atoms on lithium sites

charge balanced by vacant niobium defects. The above

threshold concentration indicates that the Mg^i defects

were controlling the defect chemistry rather than being

intrinsic

.

MANGANESE

Although there are many spectroscopic investigations

of Mn in lithium niobate, most techniques can only say

that the manganese ions sit at axially symmetric sites

(which could be Li or Nb sites). However, an EXAFS

determination has shown that Mn substitutes for lithium

[9]. EXAFS is sensitive to some radial distribution

function of neighboring atoms in the structure, including

second neighbors. Therefore the Li and Nb sites can be

differentiated, even though their local symmetry is

identical. Charge balance should be similar to that found

for low levels of iron as discussed above.

TUNGSTUN

One X-ray diffraction determination of a tungstuncontaining solid solution exists [10]. The W atoms were

found to occupy the niobium sites in lithium niobate.

The peak intensities in powder diffraction were very

272
similar to those found for pure lithium niobate, whereas

if the W had substituted for Li the peak intensities

would have been significantly different.

This 6-valent impurity will also make a netpositive type defect when substituted for niobium.

Therefore, at the large concentrations found in their

crystals, the material was no longer intrinsic; likely

charge balancing defects would be niobium vacancies.

DISCUSSION

The most likely simple model for site selection is

based on simple considerations of charge misfit; low

valence ions are often attributed to substitute for Li

and high valence ions are often attributed to sit on Nb

sites. Based on the few identifications above, this seems

to be qualitatively true. Atom size arguments can

probably be ignored because the Li and Nb sites are both

octahedral and of similar size.

The charge misfit argument leaves unclear the

preferred location for normally trivalent impurities

which will be net two charges negative or positive,

depending on substitution site.

One convenient substitution mechanism for trivalent

atoms could be simultaneous substitution onto the lithium

and niobium sublattices, forming near-neighbor chargebalanced pairs on adjacent Li and Nb sites. This may infact be happening for doping levels where the material

is no longer intrinsic.

In both low and high doping levels there will be

point defects present that have the opposite sign to the

impurity defect. As discussed above, these can form

cluster defects on cooling to room temperature. And, the

internal electrical field in the ferroelectric phase can

cause alignment of these dipoles along the c-axis of the

crystal [11]. This can explain the axially symmetric

defects often observed in various spectroscopic

techniques

.

Of the above impurity defects all divalent atoms

(Fe, Mg, and Mn) substitute onto the lithium site. This

is compatible with the charge misfit argument for site

substitution. This trend might easily apply to many of

all other divalent cations.

The only trivalent atom is Fe^* which is found to

substitute for lithium also, although the behavior at

high concentrations has not been fully characterized.

Since there are no other trivalent impurity atoms well

enough characterized, we cannot conclude that many or

most will also substitute for lithium.

273
There are no 4-valent atoms that have well

characterized locations (although Ti is often

qualitatively assigned to the niobium site). Finally,

the 6-valent substitution is found on the niobium site,

again in line with simple electrostatic considerations.

CONCLUSIONS

The divalent impurities, Mn, Fe, and Mg are all

found to substitute for lithium. The trivalent Fe also

replaces lithium. And, W replaces niobium. All of these

observations are compatible with simple electrostatic

charge misfit arguments. However, this still leaves

uncertainty about all other ion substitutions, especially

in the intermediate valence range.

REFERENCES

1) S. C. Abrahams, J, M. Reddy, and J. L. Bernstein,

J. Phys. Chem. Solids 27., 997-1012 (1966).

2) J. G. Bergmann, A. Ashkin, A. A. Ballman, J. M.

Dziedzic, H. J. Levinstein, and R. G. Smith, Appl

.

Phys. Lett. 12, 92-94, (1968).

3) P. Lerner, C. Legras, and J. P. Dugas, J. Crystal

Growth 3, 231-35, (1968).

4) L.O. Svaasand, M. Ericksrud, G. Nakken, and A. P.

Grande; J. Crystal Growth 22, 230-2, (1974)

5) S. G. Boyer and D. P. Birnie, III, Ceramics and

Inorganic Crystals for Optics ^ Electro-Optics and

Non-Linear Conversion . pp73-80, SPIE Proc. 968,

(1988)

.

6) S.C. Abrahams and P. Marsh; Acta Cryst. B42 . 61-8

(1986)

.

7) G. E. Peterson, A. M. Glass, and T. J. Negran; Appl.

Phys. Lett. 19, 130 (1971).

8) H.-R. Tan, C.-F. He; Ceramics and Inorganic Crystals for Optics^ Electro-Optics and Non-Linear

Conversion, pp48-53, SPIE Proc. 968, (1988).

9) C. Zaldo, F. Agullo-Lopez , J. Garcia, A. Marcelli,

and S. Mobilio, Sol. St. Comm. 71, 243 (1989).

10) G. Blasse and A. D. M. dePauw, J. Inorg. Nucl. Chem.

32 , 3960 (1970).

11) G. H. Jonker and P. V. Lambeck, Ferroelectrics 21,

641 (1978).

274
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

ALKALINE EARTH NITRIDES AND HYDRIDES

Nathaniel E, Brese

Department of Chemistry

Arizona State University

Tempe, AZ 85287-1604

INTRODUCTION

Nitride and hydride materials have been overlooked in favor of

oxide materials, since many are difficult to synthesize and are quite sensitive to humidity. However, their polymorphism,

crystal chemistry, and physical properties (including

superconductivity in transition metal nitrides such as MoN

approaching 30K) are fascinating. The structures of alkaline

earth nitrides and hydrides are reviewed here. The strontium

system has proved the most controversial; our recent results are therefore emphasized as they are illuminating. The PbCl2

structure adopted by most of the hydrides is discussed in

terms of cation packing and bond valences.

NITRIDES

The composition M3N2 is expected for all alkaline earth nitrides as this would provide for normal oxidation states for all the

atoms. The local coordination geometries required by the

stoichiometry include 4- and 6-coordinated or 6- and 9-

coordinated M and N atoms, respectively. The antibixbyite

structure affords tetrahedral sites for the metal atoms and

octahedral sites for the nitrogen atoms.

Beryllium nitride is polymorphic, although the normal form

crystallizes with the antibixbyite structure (1, 2, 3, 4). The

275
transition temperature is roughly 1400°C (5) and the hightemperature form is based on close-packed layers of Be atoms

in the sequence CCHCCH; the nitrogen atoms fill the octahedral

holes (6). The local coordination geometries are preserved, but

the material is no longer cubic. The high-temperature

compound reported by Eckerlin contained trace amounts of Si;

whether the presence of Si is an integral part of the beta form

is unknown. Microscopy on the Be3N2-Si3N4 system revealed

intergrowth structures in the BeSiN2-Be3N2 system (7).

Magnesium nitride forms with the antibixbyite structure. It is

easily prepared from the elements or from reactions of Mg

metal with ammonia at elevated temperatures (1, 8, 9, 10).

The normal, brown calcium nitride (a-Ca3N2) has the

antibixbyite structure (11). However, black hexagonal P-Ca3N2

and yellow orthorhombic Y-Ca3N2 are also known by their

powder patterns (12). A pernitride, Ca3N4, has been reported

(13, 14). The structure of a green subnitride, Ca2N, was solved

from single-crystal X-ray data, although the crystal was lost

and impossible temperature factors resulted, since no

absorption correction could be made (15, 16). A reddishbrown compound with composition CauNg has been prepared

from Ca3N2 at 1050°C (17); the refinement of the X-ray data

was poor. Since CauNg is extremely air-sensitive, we believe it may be an oxynitride. A nitride hydride phase with the

rocksalt structure (random distribution of hydrogen atoms

determined by neutron diffraction) is also reported (18, 19).

The black imide, CaNH, has the rocksalt structure (20).

The strontium nitrogen system does not exhibit as many

phases as the calcium one. Three nitrides have been reported,

Sr2N, Sr3N2, and SrN. The primary difficulty in characterizing

these materials is determining the levels of oxygen and

hydrogen impurity (16, 21).

Sr2N has the anti-CdCl2 structure (22) and has been the subject

of considerable discussion (16). Since the conducting behavior

276
of Sr2N was in question (23, 24), many assumed that hydrogen

was located between the layers analogous to Ca2NCl and CaiNBr

(stuffed CdCl2 structure) (25).

Using a very clean system, we recently prepared Sr2N and

showed by neutron diffraction that there is no hydrogen

between the layers, further that Sr2N is a weak metallic

conductor, and that it exhibits temperature-independent

paramagnetic behavior (26, 27). The slabs of edge-sharing

NSr6 octahedra are squashed while the vacant octahedra

between slabs are elongated. (This is exactly converse to the

situation in Ag2F in which elongated FAg6 octahedra are rationalized by Ag - Ag bonding.) We find that Sr2N is the only

stable phase in the binary system.

SrN is reported with the rocksalt structure (28), but its lattice

parameter is similar to both SrO (29) and SrNH (20). Brice

proposed that SrN is really a solid solution SrNi-xHx-nSrO (0.02

< n < 0.08) (16); it may also be a simple oxynitride (30). This

seems likely, since the only time we observed a rocksalt phase

was when the nitriding apparatus was contaminated with

moisture.

The expected composition Sr3N2 has been observed by several

workers (2, 31, 32), and its diffraction pattern has been

indexed on a centered monoclinic cell (21, 33, 34). The pattern,

however, closely resembles a mixture of Sr2N and SrN

(elemental analysis of an equimolar mixture would have the

composition Sr3N2). Motte believes Sr3N2 to contain

significant amounts of hydrogen (21). Gaude and Lang have

been unable to prepare this composition as well (35). As we

have been unable to prepare this compound, I feel that the

existence of Sr3N2 is still an open question.

Ba3N2 can be prepared as a stoichiometric compound (2, 31, 36,

37). Numerous workers have collected diffraction patterns (38,

39), but only Linke was able to index a hexagonal cell (13).

Linke's data, however, differ significantly from those collected

277
by others and by us. The main problem lies in the broad peaks

found on the diffraction pattern. Our attempted neutron

diffraction analysis suffered from the same problem. Attempts

to recrystallize BasNi at temperatures from 700-1000°C and 4-

20 kbar have so far been unsuccessful.

HYDRIDES

By comparison, the hydrogen-alkaline earth system is simple.

MgH2 crystallizes with the rutile structure (40) and CaH2, SrH2,

and BaH2 crystallize with the PbCh structure (41, 42, 43, 44).

The rutile structure allows for no variation in bond length, but

the PbCh (C23) structure type allows for unusually short and

long bonds (45). Although SrH2 was originally assigned its own

structure type (C 29) on the basis of an x-ray structure

determination in 1935 (41), our neutron experiments show it to be better described as C23.

The PbCl2 structure is generally described as cations centering

tricapped trigonal prisms. However, we find that describing

structures in terms of anion-centered polyhedra can be

revealing (46, 47). We use SrD2 as an example.

Figure 1 shows the traditional view of SrD2 (cation centered

polyhedra) while Figure 2 displays the hexagonal eutaxy of the

cations with tetrahedral sites filled by anions.

The ratio of apparent valence (bond valence sums) of the

tetrahedral and octahedral anion is much greater than unity for

the alkaline earth hydrides. This ratio gets higher as one

progresses down the periodic table (1.65-1.81 for Ca-Ba). It is

suggested that nonbonded metal - metal interactions contribute

to the stretching of the octahedral bonds (giving a low apparent

valence at the octahedral site). One expects more metal - metal

interactions with larger metal atoms, and this idea is supported

by the valence ratio trend.

278
Fig. 1 SrD2 showing Sr-centered Fig. 2 SrD2 showing D(l)Sr4

tricapped trigonal prisms, tetrahedra. Large balls

D atoms are at 0 and 1/2. are D(2) atoms.

NITRIDE HYDRIDES?

The only definite compounds observed in the Sr-N-H system

are those with N-H bonds, the imide (SrNH) and the amide

(SrNH2). Nitride hydrides, M2NH (M=Ca,Sr,Ba), have been

suggested with structures analogous to Ca2NCl (stuffed CdCl2-

type structures) and rocksalt (18). The only quantitative

structure determination was that for Ca2NH which reportedly

crystallizes with a superstructure of rocksalt, although H atoms

partially occupy general positions (19).

ACKNOWLEDGEMENTS

This material is based upon work supported under a National

Science Foundation Graduate Fellowship to NEB and is part of a continuing program in crystal chemistry supported by the

National Science Foundation (grant number DMR 8813524 to M.

O'Keeffe).

279
REFERENCES

1. B. Neumann, C. Kroger and H. Haebler, Z. anorg. allg. Chem.

204, 81 (1932).

2. A. V. Antropoff and K. H. Kruger, Z. Phys. Chem. A 167, 49

(1933).

3. M. V. Stackelberg and R. Paulus, Z. Physik. Chem., Abt. B 22,

305 (1933).

4. P. Chiotti, /. Amer. Ceram. Soc. 35, 123 (1952).

5. P. Eckerlin and A. Rabenau, Z. anorg. allg. Chem. 304, 218

(1960).

6. D. Hall, G. E. Gurr and G. A. Jeffrey, Z. anorg. allg. Chem. 369,

108 (1969).

7. T. M. Shaw and G. Thomas, /. Solid State Chem. 33, 63

(1980).

8. J. David, Y. Laurent and J. Lang, Bull Soc. fr. Mineral.

Cristallogr. 94, 340 (1971).

9. J. David, Rev. Chim. Miner. 9, 717 (1972).

10. D. R. Glasson and S. A. A. Jayaweera, /. Appl. Chem. 18, 77

(1968) . 11. P. Y. Laurent, J. Lang and M. T. LeBihan, Acta Crystallogr. B

24, 494 (1968).

12. Y. Laurent, J. David and J. Lang, C. R. Acad. Sci. Paris C259,

1132 (1964).

13. K. Linke and H. Lingmann, Z. anorg. allg. Chem. 366, 89

(1969) . 14. K. Linke and H. Lingmann, Z. anorg. allg. Chem. 366, 82

(1969).

15. E. T. Keve and A. C. Skapski, Inorg. Chem. 7, 1757 (1968).

16. J.-F. Brice, J.-P. Motte and J. Aubry, Rev. Chim. miner. 12,

105 (1975).

17. Y. Laurent, J. Lang and M. T. L. Bihan, Acta Crystallogr.

B25, 199 (1969).

18. J.-F. Brice, J.-P. Motte and J. Aubry, C. R. Acad. Sci. Paris

C276, 1093 (1973).

19. J. Brice, J. Motte, A. Courtois, J. Protas and J. Aubry, /. Solid

State Chem. 17, 135 (1976). •

280
20. H. Hartmann, H. J. Frohlich and F. Ebert, Z. anorg. allg. Chem.

218, 181 (1934).

21. J. Motte, J. Brice and J. Aubry, C. R. Acad. Sci. Paris C274,

1814 (1972).

22. J. Gaude, P. L'Haridon, Y. Laurent and J. Lang, Bull. Soc. fr.

Mineral. Cristallogr. 95, 56 (1972).

23. S. M, Ariya, E. A. Prokofyeva and L L Matveeva, /. Gen.

Chem. USSR 25, 609 (1955).

24. F. Hulliger, "Structural Chemistry of Layer-type Phases," D.

Reidel, Dordrecht, Holland (1976).

25. C. H. H. Herdejurg, Z. anorg. allg. Chem. 545, 177 (1987).

26. N. E. Brese and M. O'Keeffe, J. Solid State Chem. 87, 134

(1990).

27. S. M. Ariya, M. S. Erofeeva and G. P. Mochaloy, J. Gen Chem

USSR 27, 1806 (1957).

28. J. Gaude, P. L'Haridon, Y. Laurent and J. Lang, Rev. Chim.

Miner. 8, 287 (1971).

29. W. Primak, H. Kaufman and R. Ward, /. Amer. Chem. Soc.

70, 2043 (1948).

30. J. Gaude and J. Lang, Mat. Res. Bull. 7, 193 (1972).

31. A. Guntz and F. Benoit, Ann. Chim. 20, 5 (1923).

32. J. Gaude and J. Lang, C. R. Acad. Sci. Paris C271, 510

(1970) . 33. H. Emons, D. Anders, G. Roewer and F. Vogt, Z. anorg. allg.

Chem. 333, 99 (1964).

34. J. Gaude and J. Lang, Rev. Chim. Miner. 7, 1059 (1970).

35. J. Gaude and J. Lang, Rev. Chim. Minir. 9, 799 (1972).

36. Y. Okamoto and J. C. Goswami, Inorg. Chem. 5, 1281 (1966).

37. R. C. Blair and Z. A. Munir, /. Chem. Eng. Data 16, 232

(1971) . 38. K. Torkar and H. T. Spath, Monat. Chem. 98, 2020 (1967).

39. J. Gaude and J. Lang, C. R. Acad. Sci. Paris C274, 521

(1972) . 40. W. H. Zachariasen, C. E. J. Holley and J. F. J. Stamper, Acta

Crystallogr. 16, 352 (1963).

41. E. Zintl and A. Harder, Z. Elektrochem. 41, 33 (1935).

42. J. Bergsma and B. O. Loopstra, Acta Crystallogr. 15, 92

(1962).

281
43. A. F. Andresen, A. J. Maeland and D. Slotfeldt-Ellingsen, /.

Solid State Chem. 20, 93 (1977).

44. W. Bronger, S. Chi-Chien and P. Muller, Z. anorg. allg. Chem.

545, 69 (1987).

45. M. O'Keeffe and N. E. Brese, in prep (1990).

46. M. O'Keeffe and B. G. Hyde, /. Solid State Chem. 44, 24

(1982).

47. M. O'Keeffe and B. G. Hyde, Structure Bonding 61, 77

(1985).

282
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

OXYGEN DIFFUSION IN YoOo-CONTAINING TETRAGONAL ZIRCONIA

FOLYCRYSTALS (Y-TZP)

Yasuro Ikuma

Kanagawa Institute of Technology, Atsugi, Kanagawa 243-02, Japan

ABSTRACT

Oxygen diffusion coefficient in YoOg-containi ng tetragonal

ZrOg polvcrystals (Y-TZP) (2.75% Y2O3) was determined at 600-800'C

using 0 as a tracer. The measurement was performed on the

specimens with different grain sizes, 0.35-2.6/im, which was resulted in almost identical oxygen diffusion coefficients

suggesting that there was no fast diffusion along the grain

boundaries in Y-TZP. It was also found that the oxygen diffusion

in Y-TZP is in good agreement with the electrical conductivity

data of the same material. However, it was faster than the oxygen

diffusion in tetragonal Zr02 doped with 14 mol% Ce02

.

INTRODUCTION

A discovery of superplastici ty in Y-TZP [1,2] has introduced

a new type of processing method into ceramics; In some cases ceramics could be formed by plastic deformation. This

superplastici ty is thought to be caused by grain boundary sliding

[3]. Since diffusion is related to the grain boundary sliding and

zirconia is an oxygen ion conductor, we would anticipate that the

diffusion plays an important role in the superplastici ty of this

material. However, oxygen lattice diffusion coefficient of Y-TZP

has never been determined. The objective of this study is to determine the oxygen lattice diffusion coefficient of Y-TZP and to compare with the oxygen lattice diffusion in tetragonal and cubic

zirconia.

EXPERIMENTAL

Zr02 doped with 2.75 mol% Y2O0 was sintered at 1400*'C for 2h

and then it was pressed by HIP at 1400° C for 1.5h in an atmosphere

that consists of 96% Ar and 4% O2 . Some of the specimens were

283
annealed at 1500 or 1600*C in air in order to grow the grain size.

Grain sizes of these specimens were determined by examination of

fractured surface under SEM or the observation of thin section

under TEM. The x-ray diffraction analysis revealed that the

original specimen consisted of 100% tetragonal ZrO^ (t-ZrO^).

After annealing, however, the specimen consisted of 89% and 86%

tetragonal Zr02 for 0.42|jm and 2.6)^m, respectively. The specimens

were cut into spheres of 2mm in diameter. Their surface was finished with #1200 diamond.

Density of specimens used in the diffusion experiment was 6.08g/cm . The fracture stress of the specimen was 1450MPa.

Mi cros tructure of the specimen without annealing is shown in Fig.

1. The grains are uniform in size and are equiaxed. No pores are seen in the figure indicating that the specimens are closed to the

theoretical density.

Specimen was placed in a Pt crucible and heated to 600-800°C

in ^^©2 atmosphere until no weight change could be observed. The

specimen was cooled down to room temperature and 0 was introduced into the system. It was reheated to the same temperature. An exchange reaction shown by:

t-Zr^^02 + ^^02 = t-Zr^^02 + ^^02 (1)

1 o took place, which was followed by a subsequent diffusion of 0

into the bulk of specimen. The whole process was monitored by a microbalance. Readers are referred to elsewhere [4,5] for the

experimental detail.

Figure 1. Transmission electron micrograph of Y-TZP (2.75 mol%

Y2O3, grain size = O.GSnm).

284
RESULTS AND DISCUSSION

Typical results of diffusional annealing of specimens were shown in Fig. 2 for the specimens with 0.35j>(rn in grain size. Mj. in the figure is the total amount of 0 in a specimen at time t and My^ is the corresponding quantity after an infinite time.

Consequently if the diffusion was completed, M|./M^would be unity.

The data points are concave up within a short annealing time.

This implies that the exchange reaction at the surface of specimen

was not fast enough as compared to the diffusion process in the

bulk. The consequence of this is that we have to use an equation

[6] for diffusion from a well-stirred solution of limited volume

into a sphere when the equilibrium at the surface proceeds at a finite rate. We tried to fit the data to the equation. The curves in the figure were the results of such calculation. Since

the calculated results were in good agreement with experimental

data, the value of D (diffusion coefficient) used for the

calculation of the curve was assigned [7] to the oxygen diffusion

coefficient at the experimental condition.

The diffusion coefficients obtained in this manner are plotted in Fig. 3. The diffusion coefficients from various grain

sizes are shown in the figure. It is evident that the grain size

has little effect on the diffusion coefficient of this material

[81. This is one of the experimental results which indicate that

the grain boundary diffusion in t-Zr02 is not enhanced over the

lattice diffusion coefficient. A similar observation was reported

[9,10] for cubic ZrOo (c-ZrOo). The result in this study.

285
consequently, is the oxygen lattice diffusion coefficient (Dj q)

in Y-TZP. There is some scatter in the data at low temperatures

(600-650°C). According to the phase diagram [11] of this system,

600°C at 2.75mol% Y2O3 is really in a contradictory region. The

system may have crossed into the two-phase region (tetragonal and

monoclinic). The concentration of Y9O3 in monoclinic Zr02 (mZr02) would be smaller than the concentration of it in tetragonal

Zr02. This might be one of the factors that contribute to the low

oxygen diffusivity in m-ZrO-) . That could be the cause of the

scatter in the data. A 1 eas t^square fit to the data yields:

Dj^Q = 1.24x10"^ exp[-137kJmorVRT3 cm"s~^ (2)

It is the oxygen lattice diffusion coefficients in Y-TZP.

T/°C

1A00 1000 800 600 AOO 200

-5

-6

-12

-13

T—1

1 1 r

c-ZrOz + IZVoYzOz

o 0.35iJLm

A 0.42;im

2.6;im

c-Zr02 + 14.2%CaO

. c-ZrOz + 12V.Y203((T.)

- t-Zr02 14%Ce02

t-ZrOi 3.3y.Y203(<r.)

J i_ 1

6 8 10 12 U 16 18 20

10* TVK '

Figure 3. The diffusion coefficients in Y-TZP. Also shown are tracer diffusion in t-Zr02 doped with 14% Ce02 [12], c-Zr02 doped

with 12% Y2O3 [10] and c-Zr02 doped with 14.2% CaO [9] and

diffusivity calculated from electrical conductivity of Y-TZP (3.3

mol% Y2O3) [14] and c-Zr02 doped with 12% Y2O3 [13].

286
In Figure 3 also shown are the oxygen lattice diffusion

coefficients in cubic ZrOo doped with Y2O3 [10] and CaO [9] and

tetragonal Zr02 doped with Ce02 [12]. The diffusion coefficients

calculated from the electrical conductivity data of cubic zirconia

doped with Y20g [13] and tetragonal zirconia doped with Y2O3 (YTZP) [14] are also included in the figure. The results of current

study are in good agreement with the electrical conductivity data

for Y-TZP (3.3 mol% Y2O3) [14]. In calculating the diffusion

coefficients with the use of Nernst-Einstein equation, it was assumed that correlation factor was unity. That may have shifted

the diffusivity from the electrical conductivity to higher value

by a factor of about 4. Considering that the measurements were

not performed over the same temperature range, the agreement of

these data is excellent.

Extending the diffusion data of t-Zr02 doped with 14% Ce09

[12] to the lower temperature we found that it is smaller than the

results of this study. Doping Zr02 with Y2O3

Y2O3 = 2Y2/ + 30o^ + Vq-- (3)

the concentration of Vq* ' would be increased [15]. That would in

tern enhance the oxygen lattice diffusion. Doping Zr02 with Ce02

will not have this kind of effect. Some cerium atoms in Zr02

could be in 3+. However, the results in the figure 3 suggest that

most of cerium atoms are in 4+.

Oxygen lattice diffusion in c-Zr02 doped with Y2O3 [10] was determined at temperatures higher than lOOO^C. Extending their

results to lower temperatures, it agrees well with the data of

Stratton et al . [13] who have determined the electrical

conductivity of c-Zr02 doped with 12 mol% Y2O3. The diffusivity

calculated from the conductivity for c-Zr02 agrees fairly well

with the current results. This is rather surprising because the

concentration of impurities are quite different among these two

oxides; 2.75% vs. 12%. Since the oxygen lattice diffusion in

calcia stabilized Zr02 at 14.2 mol% CaO [9] is also in the same orders of magnitude, there might be a saturated value for the

diffusion of oxygen in Zr02. It will be increased to a higher

value if zirconia is doped with Y2O3. However, the concentration

of Y2O3 seems to have little effect on diffusivity once certain

amount of Y2O3 is exceeded. This is true for any kinds of

additives. At higher concentration of additives, a defect complex

would be formed. It is one of the factors which limit the

diffusivity at high concentration of additives.

287
CONCLUSIONS

The oxygen lattice diffusion coefficient in Y-TZP is

determined over the temperature range, 600-800°C. It did not

depend upon the grain size of the specimen suggesting that there

is no fast diffusion along the grain boundaries. The diffusivity

is in good agreement with the diffusion coefficient calculated

from the electrical conductivity measurement in Y-TZP. It was also in the same orders of magnitude as the diffusion of c-ZrOo

doped with YoOo or CaO.

ACKNOWLEDGMENTS

The author acknowledges Dr. T. Masaki of Toray Industries,

Inc. for providing the samples and Y. Tsubaki for carrying out the

experiments. This work was supported in part by the Japanese

Ministry of Education, Science and Culture under contract of

Priority Areas "Advanced Inorganic Materials."

REFERENCES

1. F. Wakai , S. Sakaguchi, Y. Matsuno, Adv. Ceram. Mater., i[3]

259 (1986).

2. F. Wakai, S. Sakaguchi, H. Kato, Yogyo-Kyokai-Shi , 94C8] 721

(1986).

3. F. Wakai, T. Nagano, J. Mater. Sci., in press.

4. W. Komatsu, Y. Ikuma, Z. Phys. Chem. N. F. , Hi, 79 (1982).

5. Y. Ikuma, W. Komatsu, Adv. Ceram., iO, 464 (1984).

6. H. S. Edwards, A. F. Rosenberg, J. T. Bittel, Report No. ASDTDR-63-635 (1963) Wright-Patterson Air Force, Ohio.

7. Y. Ikuma, W. Komatsu, Mater. Sci. Forum, 29, 199 (1988).

8. Y. Ikuma, Y. Tsubaki, J. Ceram. Soc. Japan, submitted.

9. L. A. Simpson, R. E. Carter, J. Am. Ceram. Soc, 49[3] 139

(1966)

.

10. Y. Oishi, K. Ando, in "Transport in Nonstoichiometric

Compounds", eds. G. Simkovich and V. S. Stubican (Plenum

Publ,, 1985), p. 189.

11. H. G. Scott, J. Mater. Sci., i0[9] 1527 (1975).

12. K. Ando, S. Morita, R. Watanabe, Yogyo-Kyokai-Shi, 94C8] 732

(1986).

13. T. G. Stratton, D. Reed, H. L. Tuller, Adv. Ceram., i, 114

(1981)

.

14. N. Bonanos, E. P. Butler, J. Mater. Sci. Lett., 4, 561 (1985).

15. Y. Ikuma, K. Komatsu, W. Komatsu, Adv. Ceram., 24, 749 (1988).

288
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

A STRUCTURAL BASIS FOR THE DIFFERENT DIRECTIONS OF

OXYGEN NONSTOICHIOMETRY IN La^CuO, AND Nd^CuO,

2 4 2 4

D. J. L. Hong and D. M. Smyth, Materials Research Center

Whltaker Lab #5, Lehigh University, Bethlehem, PA 18015

ABSTRACT

La2CuO^ is an oxygen-excess, p-type semiconductor while

Nd^CuO^ is oxygen-deficient and n-type. Since they contain the same amount of the same oxidizable and reducible cation, their

different behavior must be attributed to their different crystal

structures. It is shown that the major distinction is that there

is a favorable site for interstitial oxygen in the La CuO^

structure, but not in the closely-related structure of NcT^CuO^.

INTRODUCTION

La^CuO^ and Nd^CuO^ have played important roles in the

development of high temperature superconductors based on complex

Cu oxides. Ba-doped La CuO. was the first such material

discovered (1), and was T:he forerunner of an extensive family of

cuprate superconductors whose chemistry is dominated by

acceptor-doped, oxygen-excess, p-type compositions. Superconductivity has subsequently been discovered in Nd^CuO^ that was donor-doped and reduced (2) . This was the first n-type copper

oxide to exhibit superconductivity.

The different treatments given to La„CuO, and Nd^CuO, to

2 4 2 4 achieve a superconducting state are also reflected in their

different tendencies to become nonstoichiometric . Their high

temperature, equilibrium conductivities as a function of oxygen

activity are displayed in Fig. 1. It is seen that the

equilibrium conductivity of La^CuO^ increases with increasing

oxygen activity indicating oxygen-excess, p-type behavior, while

that of Nd^CuO^ increases with decreasing oxygen activity

indicating oxygen-deficient, n-type behavior. Clearly, the

stoichiometric compositions for these two oxides are displaced

laterally along the oxygen activity axis such that Nd^CuO^ is

stoichiometric at a much higher oxygen activity than is La^CuO^.

The superconducting states are achieved by doping ancT

equilibrating these materials in a way that augments their

natural tendency toward nonstoichiometry

.

289
Figure 1. The equilibrium electrical conductivities

of undoped La-CuO, and Nd„CuO, measured

as a function of oxygen activity.

The stoichiometric composition of an oxide corresponds to

the condition that the electron and hole concentrations are equal

(n = p) , and is the boundary between oxygen-excess

nonstoichiometry with excess holes, and oxygen-deficient

nonstoichiometry with excess electrons (3). The equilibrium

conductivity will have its minimum value at the oxygen activity

corresponding to the stoichiometric composition, insofar that the

electron and hole mobilities are equal. It is possible to derive

an expression for the equilibrium oxygen activity at the

stoichiometric composition from the mass-action expressions for

the equilibrium oxidation and reduction reactions. If it is

assumed that anion Frenkel defects are the most favorable lattice

defects, and that the ionic defects are fully ionized, the

oxidation reaction can be written as

+ 1/2 0^ ^ 0'^ + 2h' (1)

and the corresponding mass-action expression is

-AH /kT , [0'^] p2/ [V^] = K^^e PCO^)'^^ (2)

290
The appropriate reduction reaction is

0 1/2 0^ + V" + 2e' (3)

o o

with the mass-action expression

-AH

red

/kT -1/2 V[0^] = K

red

e (A)

The two mass-action expressions can be equated for the

stoichiometric condition n = p, and solved for the corresponding

equilibrium oxygen partial pressure, 'PiO^)°, i.e. the oxygen

partial pressure at the conductivity minima (4)

where the superscript zero denotes the situation at the

stoichiometric composition. The term in braces is a small factor near unity. It is seen that the dominant factor is the relative

magnitude of the enthalpies of oxidation and reduction. If the

enthalpic cost of oxidation is large relative to that of

reduction, then the stoichiometric composition will occur at

relatively high oxygen activities, and the material will favor

oxygen-deficient, n-type behavior at oxygen partial pressures of

1 atm and below.

The magnitudes of the enthalpies will depend on two main

factors

1 . The ease with which the ionic components of the material

can be chemically oxidized or reduced, i.e. a redox factor.

2. The ease with which the ionic defects that result from

nonstoichiometry can be accommodated within the lattice, i.e. a structural factor.

If the individual enthalpies can be taken as the sum of

these two components, then AH = h + h. , where the latter two

terms represent the, enthalpic , , . oxcontributions p, . mt from ^ the, formation r, .J.of

holes and of oxygen interstitials , and aH , = h + h , where

, , ^ red . n^ -iVac

the latter two terms correspond to the formation of electrons and

oxygen vacancies. Comparison of the exponential terms in Eq. (5)

for La CuO^ (unprimed symbols) and Nd^CuO^ (primed symbols) can be written as

exp [AH -AH ,) /kT] ox red (5)

291
h + h. - h - h vs. h' + h! - h' - h'

p mt n vac — p int n vac

(6)

Since both compounds contain exactly ;^e same amount of the same oxidizable and reducible species, Cu , the h and h terms

should be similar for both and the comparison i^s reduced to only

the structural contributions

h. - h mt vac —vs. h! - h' mt vac

(7)

Their different behavior would then seem to result primarily from

their different crystal structures. Even though the two

compounds have the same generic formula, and their rare earth

ions differ in size by less than 5%, they do have different

structures, as shown in Fig. 2. It is instructive to examine

these structures in some detail to see if there is any apparent

reason for a difference in the enthalpic cost of the lattice

defects

.

^^^^ ^^^^^

NaCi

NaCi

La 2 CUO4 Ndz Cu O4

Figure 2. The crystal structures of La^CuO and Nd^CuO^. Small

solid circles are Cu; large solid circles are rare earth cations; open circles are oxygens. The

alternate NaCl and perovskite (P) layers are indicated

for La^CuO,

.

292
La„CuO

Undoped La^CuO^ has a stoichiometric excess of oxygen when

equilibrated in an oxidizing atmosphere such as air or pure

oxygen, and can be represented by the more general formula

La^CuO^ . Recent thermogravimetric studies by Li and Alcock,

using nydrogen reduction to LajD^ and Cu to obtain absolute

compositions, showed that y > 0 ror oxygen partial pressures between 10 and 1 atm over the temperature range 750-950°C (5)

.

Appelman et al. found by iodometric titration that y = 0.007 for

samples equilibrated in air at 1000°C (6) . The equilibrium con- ductivity is in good agreement with these findings, as shown Ail

Fig. 1. The observed dependence of the conductivity on PCO^)

corresponds to the mass-action expression for the oxidation

reaction, Eq (2) , when this reaction is the major source of

charged defects such that p = 2[0l|!]. This will be discussed in

greater detail elsewhere.

A stoichiometric excess of oxygen can be incorporated either

as interstitial ions, as assumed in Eq. (1), or by the formation

of cation vacancies. The latter would require vacancies of both

cation species, or the exsolution of the oxide of the cation

whose vacancy is less-favored. Experimental evidence to be

discussed shortly shows that interstitial oxygen is the favored

mode of incorporation. The structure of La^CuO,, sometimes

referred to as the K„NiF^ structure, can be viewed as a coherent

intergrowth of the NaCl and perovskite structures (7). The

corresponding parts of the structure are indicated in Fig. 2.

However, interstitial anions are never observed as majority

defects in compounds having either the NaCl or perovskite

structures, so their presence in La^CuO^ appears to be anomalous.

This has been resolved recently by powder neutron diffraction

studies ^2 isostructural La^NiO^ by >J^2rgensen et al. (8).

Since Ni is more easily oxidized th^n Cu , the nickel analog

has much higher contents of excess oxygen, with 0.13 < y < 0.18.

This proved to be sufficient to locate the excess oxygen, and it was found to be in the NaCl part of the structure, as shown in

Fig. 3. This site corresponds to the tetrahedrally coordinated

site in the NaCl part of the structure. In the ideal NaCl

structure with close-packed anions of radius R , the

tetrahedrally coordinated site will accommodate an ion of radius

0.225R . This should be much too small to accommodate an interstitial anion. However, in La^CuO^, the NaCl part of the

structure is far from the close-pacKed prototype in which the

octahedral sites contain an ideal cation of radius 0.414R . In

~f"3 3- fact. La is approximately the same size as 0 , and the anion

sublattice is forced to expand far beyond the close-packed case. Thus instead of having the ideal NaCl packing as shown in

293
Figure 3. La2NiO^ showing the interstitial oxygen

(large solid circle) and the distorted cube

of nearest neighbors (double circles)

.

Adapted from Ref. (8).

Fig. 4a, the situation is more accurately represented by Fig. 4b.

As the cation in the NaCl structure approaches the anion in size,

the tetrahedrally coordinated sites expand by more than a factor

of 3 to a radius of 0.732R . In other words, the tetrahedrally

coordinated site takes on t^e character of the body-centered site

in a simple cubic structure with the cube corners alternately

occupied by cations and anions. This relieves the size

constraints on the incorporation of interstitial anions.

Jorgensen et al. also found that the surroundings of the

interstitial anion are distorted so that the four cations are closer (average distance 0.2395 nm) than the four anions (average

distance 0.2625 nm) so that the electrostatic environment is more

hospitable for the interstitial anion. The observed

oxygen-excess, p-type behavior of La^CuO^ is related to the

presence of an accessible interstitial site for excess anions.

294
a. No CI

(220)

b. Log Cu O4 c. Ndj Cu

Figure 4. (a) Side view (220 plane) of the ideal NaCl structure with R = 0.414R . (b) Side view of NaCl-type

part of the^La^CuO^ structure with = R . (c) The

lateral shift of a pair of ions that gives the more

effective space-filling of the Nd.CuO structure.

When La2CuO^ is acceptor-doped, as material is even more easily oxidized to

state, because of the ease of filling

vacancies

V" + 1/2 0^ - ^- 0 + 2h' (8) o Z o

This oxidation reaction actually reduces the concentration of

lattice defects and is thus energetically more favorable than the

oxidation of the undoped material where the number of lattice

defects is increased, as shown in Eq. (1). Thus while Fig. 1 shows that there is essentially no temperature dependence to the

equilibrium conductivity for the undoped material, indicating

that aH

^

0, the oxidation of the acceptor-doped material is ox exothermic, so that the oxygen content increases with decreasing

temperature. This is another example of the importance of the

enthalpic cost of forming the lattice defects.

Nd^CuO,

The equilibrium conductivity of undoped Nd_CuO^ shows

oxygen-deficient, n-type behavior as shown in Fig. 1. The

conductivity increases with decreasing oxygen partial pressure as expected for the loss of oxygen according to Eq (3)

When the cations and anions are about the same size, the

NaCl structure is not an effective way to fill space, as shown in

Fig. 4b. More efficient packing can be accomplished by

in La_ Sr CuO, the

^—2^ x/Z an oxygen-excess, p-type

the extrinsic oxygen

295
translating adjacent layers so that the ions in one layer fit

into the interstices of the next, as shown in Fig. 4c. This is

approximately the response of the Nd CuO^ structure where the

anions on the cell edges in one of the NaCl layers in La^CuO^

move onto the faces where they can move up between the Targe

cations in the adjacent layer. Since this increases the net

anion occupancy in this layer to 2 (4 anions on faces instead of on edges) , the apex anions from the CuO^ octahedra in La^CuO^

have dropped into this anion layer in Nd^CuO^, leaving the Cu in

four-fold, planar coordination. Schematic cross-sections through

these anion planes are sho\^ for the two structures in Fig. 5.

In La2CuO^, there is a La in the center of the four edge

oxygens. In Nd2CuO^, the oxyg^^s on the faces intrude further

toward the center, and the Nd is extruded from the remaining

restricted space in this plane toward the adjacent Cu-0 plane.

As a result of the rearrangement of the NaCl pattern in La CuO^

to the more effective packing of Nd^CuO, , the c-axis of the

latter is shorter^^y 7.6%, more than can be ac^unted for by the

replacement of La by the slightly smaller Nd

Net Oxygen

Content

Edges Faces

Fig. 5. (a) Cross-section through a NaCl-type layer in

La^CuO^ with oxide ions on the cell edges and

lanthanum in the center, (b) The corresponding

shifted layer of oxide ions in Nd^CuO^ that have

moved to the faces without a central cation.

In Nd CuO^, the space that was available for interstitial

anions in La2CuO^ is now occupied by the shifted oxygens and can no longer accommodate excess oxygen. By this reasoning, one might look for a possible interstitial site back at the vacated

edge in Nd CuO^. However, as shown in Fig. 6, this site is now electrostatically very unfavorable because it is surrounded by 8

other anions and a single Cu in the adjacent plane. The large

296
cations "see" the interstitial site through the space between

four oxygens and are thus quite well screened from the site.

Thus this interstitial site is much less favorable than the one

in La„CuO,

Nd, Cu O4

Figure 6. A possible interstitial site (X) in Nd^CuO^

near a site normally occupied by oxygen in

La^CuO,

I 4

One might also consider the apical oxygen sites that are filled in the perovskite layers of La^CuO^, but that have been

vacated in Nd^CuO^, as possible interstitial oxygen sites in the

latter compound. However, the adjacent anion layer in Nd^CuO^

has twice the anionic charge it has in La2CuO^ , and it has moved

closer to the apical site, making this interstitial site much

less favorable. Thus the potential sites for interstitial oxygen

in Nd-CuO, are much less favorable than the site in La„CuO, . As

/ q } L\ a result, h! for Nd„CuO, in Eq . (7) is much larger than h, int 2 4

-I ' o int

for La^CuO, . Assuming that the h terms are comparable, fnxs

2 4 vac will displace the stoichiometric composition for Nd^CuO^ to much

higher oxygen activities so that in the normal experimental range

below 1 atm, this material shows only oxygen-deficient, n-type

behavior

297
VThen Nd2CuO^ is donor-doped, the stoichiometric composition,

corresponding to a combination of stoichiometric binaries, will

be of the type Nd„ Ce CuO, . ,_. Since there is no favorable

site in the structure for the excess oxygen, it is expelled under

usual conditions, leaving behind electrons

(Nd 0 ) ZCeO^ > ^^^Nd ^ ^°o ^ ^2 ^^-^

This gives the characteristic n-type behavior of the donor-doped

material.

RELATED COMPOUNDS

It has recently been shown that Nd_NiO, , normally exists

with a stoichiometric excess of oxygen, i.e. with positive values

of y (9). Since this compound is isostructural with La.CuO,

,

this is to be expected according to the arguments of this paper,

and confirms that the oxygen-deficient, n-type behavior of

Nd„CuO, is not related to its Nd content, but to its structure.

The major difference between La2CuO^ and Nd2NiO^ is that the

am^^nt of excess oxygen is much greater^^n the nickelate, because

Ni is more easily oxidized than is Cu

Lightfoot et al. recently determined that

La^ 25^^0 J5^^^3 75^0 5 so-called T* structure, in which

one* end o*f the unit cell has the La„CuO^ structure while the

other end has the Nd„CuO, structure (iO) . Thus the material has

favorable sites for interstitial oxygen, but only half as many as in La^CuO , . The material was found to exist with a stoichiometric excess of oxygen which was determined by powder

neutron diffraction to exist as interstitial oxygen only in the

NaCl part of the La^CuO^-end of the unit cell. Thus when the

unit cells contain elements of both types of structure, the

availability of a favorable interstitial site for excess oxygen

determines the basic redox balance of the material.

CONCLUSIONS

Since La^CuO^ and Nd^CuO have the same type of formula, and

contain the same amount of tne same oxidizable and reducible

cation, the chemical tendency to oxidize or reduce should be the same. Thus the different directions of oxygen nonstoichiometry

observed for the two undoped materials must be attributed to

their different crystal structures. It is shown that the primary

distinction is that there is a favorable site for interstitial

oxygen in the La CuO lattice, but not in the Nd CuO. lattice.

298
ACKNOWLEDGEMENT

This research was supported by the Division of Materials

Research of the National Science Foundation.

REFERENCES

1. J. G. Bednorz and K. A. Muller, Z. fur Physik B - Condensed

Matter, 64, 189 (1986).

2. E. K. Takayama-Muromachi, Y. Matsui, Y. Uchida, F. Izumi, M.

Onoda, and K. Kato, Jap. J. Appl. Phys., 2_7(12), L2283

(1988) .

3. D. M. Smyth, in Selected Topics in High Temperature

Chemistry: Defect Chemistry of Solids, 0. Johannesen, and A.

G. Andersen, Eds., Elsevier, Amsterdam (1989), pp. 1-27.

4. N.-H. Chan, R. K. Sharma, and D. M. Smyth, J. Am. Ceram.

Soc, 65(3), 167 (1982).

5. B. Li, Ph.D. Dissertation, University of Notre Dame, March

1990.

6. E. H. Appelman, L. R. Morss, A. M. Kini, U. Geiser, A.

Umezama, G. W. Crabtree, and K. D. Carlson, Inorg. Chem.

,

26, 3237 (1987).

7. D. M. Smyth, in Ceramic Superconductors II, Research Update,

1988, M. F. Yan, Ed., The American Ceramic Society, Inc.,

Westerville, OH (1988), pp. 1-12.

8. J. D. Jorgensen, B. Dabrowski, S. Pei, D. R. Richards, and

D. G. Hinks, Phys. Rev. B, Condensed Matter, 40(4), 2187

(1989) .

9. B. W. Arbuckle, K. V. Ramanuj achary , Z. Zhang, and M.

Greenblatt, to appear in J. Sol. State Chem., 88 (1990).

10. P. Lightfoot, S. Pei, and J. D. Jorgensen, Physica, C169, 15

(1990) .

299
DISCUSSION

R. Cava: So, Don, you would predict that the neodymium cuprate or any neodymium

cuprate structure type cannot be made into a P-type superconductor,

D. Smyth : No, I wouldn't say that. It's just that everything is displaced in oxygen

pressure space, and you would have to equilibrate the material in high oxygen pressures.

B. Wuensch: Don, this is not a sling or an arrow. This is support. I don't think you

have to be quite so apologetic for an interstitial oxygen in a rock salt-type oxide. There

is fragmentary evidence, but a lot of evidence, from creep data and diffusion data that

there may be oxygen interstitial as the predominant effect, and also there is the peroxicenter, for example, in barium oxide.

D. Smyth : The peroxide ion is quite a different thing. What type of materials for the

creep data?

B. Wuensch : MgO, MgO types, and solid solutions.

D. Smyth : There are processes, like creep or sintering, that require transport of a minority defect. However, I know of no case in which interstitial anion are majority

defects in the rock-salt structure.

A. Macker: These materials are also potential candidates for high temperature reagents

in solid oxide fuel cells. While we have studied the lanthanum strontium cuprate, but the

neodymium cerium cuprate, is also that material which has no phase transition below

room temmperature.

D. Smyth : Well, there is some phase separation behavior at low temperatures. We've

done our conductivity measurements above 500°C where I don't think that's the case.

I don't know whether those types of phase separation at lower temperatures would affect

your properties or not.

B. Scott : I want to mention that the Johnson group at Ohio State reported a lanthanum

copper oxide, which they get to in quite an unusual way, but reported lanthanum copper

oxide with excess oxygen ions in the structure, and it has as much as, I think, 4.33

oxygen.

D. Smyth : I believe that this material is made by a low temperature process, and that the excess oxygen is not thermodynamically stable.

B. Scott: Yes.

300
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

RUDDLESDEN-POPPER PHASES An+iMnOsn+i:

STRUCTURES AND PROPERTIES

J.S. Kim, J.Y. Lee, J.S. Swinnea, and H. Steinfink

Materials Science and Engineering, Department of Chemical Engineering,

The University of Texas, Austin, TX 78712,

W.M. Reiff, Department of Chemistry,

Northeastem University, Boston, MA 021 15, and

P. Lightfoot, S. Pei, and J.D. Jorgensen,

Argonne National Laboratory, Argonne, IL 60439

INTRODUCTION

Intergrowth crystal structures in inorganic systems have provided a fertile field of solid state chemical research for many years and the discovery of high-Tc

superconducting cuprates has focused attention on intergrowth structures based on

perovskite blocks. Ruddlesden and Popper (R-P hereafter) (1) have described

intergrowth structures with the general formula Srn+iTin03n+i where n is the

number of perovskite layers separated by SrO layers. Indeed, La2Cu04

corresponds to the n = 1 phase in which the decreased ionic charge of the octahedral

ion is compensated by the increased ionic charge of the dodecahedral ion.

Numerous structures have been derived from these parent compounds by

the replacement of Ti"^"*" with lower valent transition metals and charge

compensation with higher valent ions on the alkaline earth site. In this manner the

series of compounds Ln2AM207, Ln^"*" = La, Nd, Sm, Eu, Gd, Tb, A2+ = Sr, Ba,

M^"*" = Mn, Fe, has been synthesized and they are isomorphous with the n = 2 R-P

phase (2,3). The replacement of Ti"^"*" by Cu^"*" to form La2SrCu206 requires the

removal of an oxygen ion for overall charge neutrality. The apical oxygen bridging

the two octahedra is missing resulting in layers of copper pyramids whose bases

face each other (4). The lowering of the valence of the octahedral ion from 4+ to 3+ is charge compensated by trivalent rare earth ions and maintains the octahedral

double layer. However, the replacement of the 4+ ion with Cu2+ introduces a modification in the structure by the removal of the apical O^- to create separate

copper pyramids. The replacement of 2Ti^''" in the parent compound by Fe3"'"Cu2+, to create an intermediate valence state +2.5, again causes the removal of an oxygen

ion, but now the unshared apex of the octahedron is lost. The structure of

YBaPeCuOs contains double layers of apex-sharing Fe/Cu pyramids separated by a

single layer of Y in 8-fold coordination - a fluorite unit (5). Its structural

relationship to the R-P parent compound with n=2 is not apparent from the

stoichiometry. To emphasize it requires the formula to be written as

301
(YBao)(FeCu)07^, but no cation vacancies exist in this structure. We synthesized

a compound that contains the apex-sharing double pyramidal layer as a structural

motif and has all the alkaline earth sites fiUed.

Relatively few compounds with n=3 have been investigated. We have

synthesized and investigated the properties of LaSr3Fe309+s and report here some

of the results.

Y2SrFeCu06.5

The substitution of 2Y and FeCu for 2Sr and 2Ti in Sr3Ti207 led to a new

phase Y2SrFeCu06.5. Its crystal structure and magnetic properties were

investigated using single crystal X-ray diffraction, powder neutron diffraction,

Mossbauer spectroscopy, and magnetic susceptibility measurements.

Y2SrFeCu06 5 is orthorhombic, Ibam, a = 5.4036(8) [5.4149(1)] A,

b = 10.702(1) [10.7244(1)] A, c = 20.250(2) [20.2799(2)] A; values in square

brackets are neutron data. The structure is derived from the R-P phase

Srn+iTinOsn-i-i with n = 2. The Fe/Cu atoms occupy randomly the approximate

center of oxygen pyramids. The pyramids share the apical bridging oxygen and

articulate laterally by comer sharing of oxygen to form a double pyramidal layer

perpendicular to c. The pyramidal slabs are separated by double layers of Y that are

in 7-fold coordination to oxygen, forming a defect fluorite unit. Fig. 1.

Sr^Ti.O^, Y^SrCuFeO.

J z 7 z o

Fig. 1. The evolution of the Y2SrCuFe06.5 from the parent compound. Large

circles are O, smaller circles are Y, and the smallest are Fe/Cu. The defect fluorite

unit is outiined around one Y.

302
The oxygen layer between the Y layers contains a vacant site. The surrounding

oxygen ions within the layer and oxygen ions firom the pyramidal bases situated

above and below the vacant site are displaced toward it and distort the pyramids.

The average bond length Sr-10 O is 2.793(10)A and for Y-7 O is 2.331(6)A. The

Fe/Cu-0 distances vary from 1.940(6) to 2.002(2)A.

Mossbauer spectra indicate a unique iron environment and magnetic

ordering at about 265 K. The paramagnetic phase coexists with the magnetic phase over an approximate temperature range 300 - 263 K, characteristic of low

dimensionality (1 or 2D) magnetic interactions. The temperature dependence of the

molar magnetic susceptibility is consistent with the Mossbauer data and also

suggests low-dimensionality with a broad maximum in Xj^ vs. T at 260 K. The

isomer shift and quadrupole spUtting are consistent with Fe^"*" in 5-fold coordination

and Hint values also indicate classic high spin Fe3+, Table 1, Fig. 2.

Table 1. Isomer shifts 5*. quadrupole splittings AE, and internal fields Hint for Y2SrFeCu06 5.

Temperature (K) AE (mm/sec) 5 HintOcG)

306 0.56 0.26 0

295.7 0.57 0.25 0

295 0.54 0.25 0

285 0.61 0.28 >0

253.4 252

241.2 270

216.8 333

199.3 362

181.7 385

181.1 382

145.7 418

120.7 440

77 467

4.8 499

Relative to natural iron metal foil.

TlK)

Fig. 2. Internal magnetic

field vs. temperature

for Y2SrFeCu06.5

303
Powder neutron diffraction data at lOK show several new peaks that were absent in

the room temperature data. They could be indexed on the basis of a primitive unit

cell of the same dimensions as the original body-centered cell. Numerous models

for the magnetic structure were tried but none accounted satisfactorily for the

observed magnetic line intensities. It is interesting to note that no magnetic

structural studies of R-P related phases have been reported. These structiu'es

(AO-2ABC)3) may be regarded as intermediate between the 3D perovskite (ABO3)

and 2D Sr2Fe04 (A0-AB03) structure types and would be of interest in studying

pseudo-2D magnetic properties.

This phase does not display detectable deviation from the 1 : 1 Fe-Cu ratio.

Substitution of 0.20 Ca for Y is possible. This phase can also be synthesized with

Gd and Er. It is a semiconductor.

LaSr3Fe309+5

Compounds based on the parent structure with n - 3, Sr4Ti30io, have not

been investigated extensively. Brisi and Rolando (6) suidied the system Sr-Fe-O

and report an oxygen deficient Sr4Fe3Oi0-x phase. We prepared LaSrsFesOx by

standard ceramic techniques as single phase material by reacting the oxides and

carbonates at about 1450°C. A single crystal from the reaction product was

selected for the determination of the crystal structure by X-ray diffraction. The

compound is isostructural with Sr4Ti30io, Fig. 3, but is oxygen deficient, with

composition LaSr3Fe309+5. y.

La Srg

shown as small non-bonded

circles.

Fig. 3. The structure of

LaSr3Fe309±§. The oxygen

vacancies are primarily in the

equatorial plane of Fe 1 . The

crystallographicaily distinct

atoms are labled. Sr are

304
The value of 5 is approximately ±0.25. The oxygen vacancies are concentrated

around the central Fel octahedron and oxygen can be reversibly intercalated. The

equatorial Fe-O lengths for the octahedra are equal, i.939(5)A, but the apical Fel - 02 lengths for the central octahedron are significantly shorter, 1.8785(2)A, than for

the flanking octahedra. The Fe2 - 02 length of 2.0982(1 5)A to the shared 02 is also significandy longer than Fe2 - 03, 2.058(8)A, to the unshared apex. One

La/Sr is in a dodecahedral site and the other crystallographically independent La/Sr

is in 9-fold coordination, a monocapped square antiprism. Bond-valence

calculations (7) using the expression V = I exp (^^"^y^)

ro (Fe3+) = 1.759, ro La3+ + 1 Sr2+) = 2.132, yield +3.92 for Fel, +3.33 for

Fe2, +2.28 for La/Srl, and +2.20 for La/Sr2. If the oxygen stoichiometry for this

phase can be increased to 9.5, the Fel most likely becomes the tetravalent ion. The

valence sums for La/Sr are equal to the weighted value for the site occupant in the

structure.

Mossbauer spectra and magnetic susceptibility data indicate 3D

antiferromagnetic ordering temperatures of ~75K for the low oxygen content

material and 1 50 K for the high oxygen content phase. The compound decomposes

when oxygen content is less than about 8.6. We have determined that the Sr/La

ratio is 3:1 within the limits of second-phase detection by powder X-ray diffraction.

This phase forms with La, Nd, Pr, and Gd but not with Er, Ca and Ba do not

substitute for Sr. ACKNOWLEDGMENTS

The first four authors acknowledge the support of the Robert A. Welch

Foundation of Houston, Texas, and the Microelectronics and Computer

Technology Corporation (MCC) of Austin, Texas. W.M.R. thanks the

Northeastern University Research and Scholarship Development and Biomedical

Support Funds for support to purchase a cobalt-57 Mossbauer spectroscopy

source. The research by P. L., S.P. and J.D.J, is supported by the US Department

of Energy, Basic Energy Sciences/Materials Sciences under Contract No.

W-31-109-ENG-38 and by NSF Science and Technology Center for

Superconductivity Grant No. DMR88-09854. We thank Dr. Joel S. MiUer and Mr.

Scott McClean of the Central Research and Development Laboratory, DuPont Co.,

Wilmington, DE for the magnetic susceptibility measurements.

305
References

1 . S. N. Ruddlesden and P. Popper, Acta Cryst., 1 1, 54 (1958).

2. D. Samaras, A. Collomb, and J. Joubert, J. Solid State Chemistry 7, 337

(1973).

3. M. N. Deschizeaux Cheruy and J. C. Joubert, J. Solid State Chemistry, 40,

14 (1981).

4. N. Nguyen, L. Er-Rakho, C. Michel, J. Choisnet, and B. Raveau, Mat. Res.

Bull. 15, 891 (1980).

5. L. Er-Rakho, C. Michel, Ph. Lacorre and B. Raveau, J. Solid State

Chemistry, 73, 531 (1988).

6. C. Brisi and P. Rolando, Ann. Chim. (Rome), 59, 385 (1969).

7. I. D. Brown and D. Altermatt, Acta Cryst. B41, 244 (1985).

306
SUPERCONDUCTORS

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

MICROSTRUCTURES IN HIGH TEMPERATURE SUPERCONDUCTORS

H.W. Zandbergen

National Centre for HREM, Laboratory of Materials Science, Delft University of Technology,

Rotterdamseweg 137, 2628 AL Delft, The Netherlands

ABSTRACT

Electron microscopy results of a number of research projects are reviewed, showing the important role electron microscopy can play: studies of

several defects in YBa2Cu307, (Bi,TI)2(Sr,Ba)2CanCui +n06 + 2n. and

Pb2Sr2(Cai -xYx)nCu2 + n06 + 2n+3 have led to successful synthesis of a number of compounds; a high density of [100] 90° rotation twin boundaries in YBa2Cu307 leads to a decrease in Tc and the superconducting fraction, structural

features at the grain boundary can explain the cause of the low critical currents

in polycrystalline materials; using detailed analysis of superstructures in (Bi,Pb)2Sr2CaCu208+3 models were postulated, which were confirmed with

other techniques.

INTRODUCTION

Microstructures, such as grain boundaries and defects, determine for a major part the properties of materials. For instance grain boundaries in YBa2Cu307 appear to have a strongly negative effect on the critical current. Unique data can be obtained with electron microscopy, in particular when high

resolution electron microscopy (HREM) is combined with element analysis on a nanometer scale.

From different types of research on superconductors quite a number of

questions are raised, which can at least partly be answered using electron

microscopy. A number of these questions are listed in Table 1

.

Table 1. Several questions concerning (micro) structures of superconductors.

Question Type of EM research required

Which compositions have a change to be stable as new single phase materials? Can specific physical properties be correlated

to a given defect? What are the reasons for the low Jc's in polycrystalline materials? Which defects have good pinning properties?

What is the nature of thin film/substrate

interfaces What is the cause of the modulated structures, and how can we influence it? What are the solubility ranges?

What is the distribution of elements over

the sample?

What are the exact atom positions?

What is the valency of the ions?

HREM and Eelement analysis

(determine defects)

TEM, HREM (quantization of defects in well characterized samples)

HREM (study of grain boundaries)

TEM, HREM (determine density of defects)

HREM and element analysis

HREM, electron diffraction

element analysis, TEM, HREM

element analysis

detailed HREM, image processing

EELS (but difficult)

309
For many applications the performance of large scale manifactured high

temperature oxide superconductors is still too poor. Thin films, obtained by

laser ablation or off-axis sputtering, meet the requirements, but these

techniques are often too time-consuming and too expensive. Single crystals are

interesting for fundamental research studies but will also be too expensive and

too complicated for application. Consequently it is very worthwhile to try to understand the reasons of the poor performance of many of the superconducting

materials, in particular polycrystalline materials. Concerning the latter it is important to understand why it shows to have such low critical currents and to find a method of grain boundary engineering to obtain high Jc polycrystalline

material.

HREM was carried out with a Jeol 200CX (200 kV, ±10° tilt), a Jeol

4000EX (400 kV, ±20° tilt) or a Philips CM30ST (300 kV, ±25° double tilt). The last electron microscope i^s equipped with a Link EDX element analysis

system. Thin foils for the study of grain boundaries were obtained from pellets

by mechanical thinning with dry grinding paper to about 80 nm and further

thinning by ion milling. Ion milling was performed on liquid nitrogen cooled

specimens with Argon and was started with 20°, 5-6 kV, 2mA and finished with

12°, 3-4 kV, 0.8 mA. For other electron microscopy studies thin specimens

were obtained by mounting grounded crystal fragments on carbon coated

perforated polymer films supported by a Cu or Au grid.

DEFECTS AS INDICATIONS FOR EXISTANCE OF NEW COMPOUNDS

In Table 2 compounds are listed which were observed as defects in superconductors before the compound itself was synthesized. Only defects in the

first two compounds will be discussed here.

Table 2. Defects occurring in parent compounds indicating a new stable compound

Parent compound Defect observed New compound References

YBa2Cu408 (CuO)2 double layers YBa2Cu408 1,2,3

Bi2Sr2CaCu02 extra Ca and Cu02 layers Bi2Sr2Ca2Cu30io 4,5

Tl2Ba2CanCui+n06+2n single TIO layer TlBa2CanCui+n05+2n 6,7

Pb2SrLaCaCu308+a Ca, Cu02 layers absent Pb2SrLaCu206+a 8,9

PbO, Cu layers absent Pbi-xSr2+x(CaY)Cu206 10,11

The most frequently observed planar stacking fault in YBa2Cu307 is a (CuO)2 double layer formed by insertion of an extra CuO layer at the CuO layer

between the two BaO layers [1,12,13]. Detailed analysis of HREM images [1] showed that at the defect Cu-O-Cu-0... chains occur in the direction of the

displacement required to accommodate the extra CuO layer. These planar stacking

faults were observed first as defects formed in YBa2Cu307 when it is exposed to

air at room temperature for some time [1,12] and it was suggested that the

formation of these defects is due to the thermodynamical instability of this

material in the temperature range 20-500°C. Recently data were published

showing YBaaCusOy to be unstable below 800°C. Regular intercalation results

in a superstructure, which has been observed first in thin films [14-16] and

310
films [14-15] and near grain boundaries in dense sintered material [14]. Later,

it has been found possible to synthesize materials having a regular incorporation

of extra CuO layers. Bordet et al [16] reported the existence of YBa2Cu3.507.5,

in which every second CuO layer is replaced by a (CuO)2 double layer. The

existence of such a structure was predicted in one of our early papers [12].

YBa2Cu408, in which every CuO layer is replaced by a (CuO)2 double layer, has

also been synthesized [2,3,17].

When we started our research [4,18] on the superconducting phases in the system Bi203-SrO-CaO-CuO only the compounds with n=0 (Tc=20K) [19]

and n=1 (Tc=80K) [20,21] were known. HREM on a sample partly containing

material having a Tc of 110 K, showed the existence of regions having a longer c- axis (3.7 nm), suggesting the existence of a n=2 type material with a Tc of 110

K, which was confirmed by many researchers. Analytical electron microscopy on

this sample showed the existence of crystals with a range of compositions, in agreement with the intergrowth of the n=1 and n=2 phase. Further research has

revealed that the compounds in these series show a strong intergrowth of

lamellae of each of the phases in this system.

STRUCTURE/PROPERTY RELATIONS

Direct evidence of structure property relations in oxide superconductors

is still quite scarce. In this paper an example will be given of the influence of

[100] 90° rotation twins (see figure 1) on the superconducting properties.

Another example reported in literature [22,23] is the presence of many edge

dislocations due to interruptions of planar stacking faults, which occur in many

thin films and which seem to lead to enhanced flux pinning.

Figure 1. Micrograph of YBa^CusOy.a showing [100] 90° rotation twin boundaries. For

A, B, and C regions the c axis is out of the plane, vertical and horizontal respectively.

311
0.0 -

c-0.4

-0.8 -a Q a

I I

I I

I i_

• sso-c O 790°C

20 40 60

T(K)

80 40 42 44 46

26(degrees)

48

Figure 2. Magnetic susceptibility and -X-ray diffraction data of samples of YBa2Cu307.a

obtained by heating at various temperatures resulting in different densities of [100]

90° rotation twin boundaries. For high densities the superconducting fraction is very

small and the material remains tetragonal (bottom line in X-ray plot).

CuO

® o o •

® o o • ® o o • ® o

® o •

• o ®

o • o

• o ®

o • o

• o ® o • o

• o ® o • o

• o ®

o • o

• o ®

o • o

o ®

c axis

9 O • o • o

® o • o • o

® o • o • o

9 O • o • o

® o • o • o

® o • o © o

® o •

ro • • o o ®

•o ®

e o o • • o

o ®

•o ®

• o o •

(Ba,Y)0 BaO

o • • o

o ®

•o ® • o

o • • o

o ®

•o ®

• o

o •

o • o • • o • o

O ® O 9

• • o ® o ®

• o • o

o • o • • o • o

O 9 O ®

• • o ® o ®

• o • o

o • o •

90«0® »0»0

twin boundary

o « o • o

® o • o ®

o • o • o

® o • o ® o • o • o 9 O • O 9 o • o • o

9 O • O ® o • o • o

9 O • O ®

O • O • O • 9 O • O ®

c axis

o • o •

® o • o

• o

9

• o

9

• O

9

• O

• O • o • o

• 09

o • o

• O 9

O • O

9 O 9

9 O • O

O 9 O 9

® O • O

r

CuO,

• o • o • O 9 O 9 O

® O ® O 9

• • 9 O ® O 9

O 9 O 9 O

• O • O 9 O

9 O 9

• 9 O 9

O 9 O

• O •

O •

twin boundary

o

Figure 3. Atomic arrangement at [100] 90° rotation twin boundaries. From HREM

images it cannot be determined which one of both structures occurs.

[100] 90° rotation twins occur when the c/a ratio in YBaaCuaOy -like materials is close to 3. An electron microscopy study has shown [24] that the

number of [100] 90° rotation twins increases with decreasing calcination

temperature. These twins are formed during the growth of the (tetragonal)

crystal are not the result of impinged grain growth.

The presence of pre-existing [100] 90° rotation twins in the tetragonal

phase has significant implications. Such twin boundaries act as impediments to the complete transformation the high temperature tetragonal phase to the low

temperatur orthorhombic one. This is mainly due to the mismatch between the a

axis on the one hand and the b and c/3 axes on the other hand, which will lead to a

high number of dislocations, as is observed experimentally [24]. The required

312
strain energy will at least reduce the free energy decrease accompanying the

tetragonal-to-orthorhombic transformation and can completely inhibit the

transformation for sufficiently high [100] 90° rotation twin densities. Indeed

the larger the twin density the poorer the superconductivity properties are found

to be, as can be in figure 2. When the synthesis [25] is done in nitrogen at about

650°C good superconducting orthorhombic YBa2Cu307-9 can be obtained after annealing at about 450°C. This shows that the formation of these twins is not due

to low diffusion rates at 800°C but must be due to a difference in the structures

with various oxygen contents.

The atomic arrangement of these twin boundaries is given in figure 3. Across the twin boundary the sublattice of the Cu ions is continued [26]. Many

grain boundaries on a axis oriented thin film have a similar atomic arrangement

because many grains have a [100] 90° twin relation. Because the sublattice of

the Cu ions are continued across the grain boundary with only a change in direction of the (Ba,Y) sublattice and the oxygen sublattice, this type of grain

boundary is expected to be not a weak link. Indeed still high critical currents are observed for a axis oriented thin films [27], even when these films consist of

multilayers of superconducting YBa2Cu307 and non-superconducting

PrBa2Cu307 [28]. Babcock et al [29] report similar results for a bicrystal

with one [100] 90° rotation twin boundary.

THE CAUSE OF LOW Jc's IN POLYCRYSTALLINE MATERIALS

Grain boundaries can act as strong pinning centers if the Lorentz force is almost exactly perpendicular to the grain boundary, but they can also act as weak

links for currents crossing the grain boundary. Small deviations, inherent to grain boundaries, can strongly influence the performance of superconducting

materials. Since the coherence lengths of superconductivity in copper oxides are very small (e.g. 0.4 nm parallel to the c axis and 3 nm perpendicular to the c

axis) the wave function of the condensate will decay very rapidly at the grain

boundary. Consequently the grain boundary has to be studied on an atomic scale,

in order to determine its possible influences on superconducting properties.

O G © © © Y ® o ® o ® O ® o ® o ® CuO^ o • o • o • o • o • o BaO

-• o • o • o • o • o • CuO o • o • o • o • o • o BaO ® o ® o ® o ® o ® o ® CuOj

© © © © © Y ® o ® o ® o ® o ® o ® CuO,

o • o • o • o • o • o BaO

• o • o • o • o • o • CuO / • o • o BaO

grain boundary -

Figure 4. Representation of the atomic arrangement at a (001 ) grain boundary. The unit

cell in perspective view and in [100] projection are given on the left. The grain boundary is either between the BaO and CuO planes or between the Cu02 and BaO planes.

313
Electron microscopy studies on grain boundaries in YBa2Cu307-3 have

been reported by many groups [30-34]. During sintering no amorphous phase is formed [30-32], but such a phase can be formed during annealing [35]. Also

agreement exists on the occurrence of dislocations at a number of grain

boundaries [32-33], the presence of planar defects near the grain boundaries

[35], and cracks, being perpendicular to [001], with a width of 2-8 nm

[30,34].

Dimes et a! [36] have suggested that dislocations present at low angle

grain boundaries are the cause the large drop in Jc of maximum 50 times, as observed in studies of thin films of YBa2Cu307-a on SrTiOa bicrystals. Babcock

et al [29] have reported similar results. It should be noted that the grain

boundaries Dimes et al investigated occur only scarcely in polycrystalline

YBa2Cu307-9. Most of the grain boundaries in YBa2Cu307-3 are parallel to (001) of one of the adjacent grains. Also the drop, observed by Dimos et al, is mud less than the one found for polycrystalline materials (about 1000 times).

Consequently other structural features in randomly oriented polycrystalline

material must be the cause the very considerable extra drop in JcIn randomly oriented YBa2Cu307-a the connecting grain boundaries are

for more than 90% parallel to (001) of one of the adjacent grains (see figure

4). It is very plausible that it is this type of grain boundary indeed strongly

deteriorates the critical current [35], because the coherence length along the c

axis, and thus over (001) grain boundaries, is very small, so that such grain

Figure 5. Micrograph of a grain boundary in Lai PsBai 75CU307.3, heated in nitrogen

at 950°C for 24 hours, showing the presence 0f'La2-xDaxCuO4 at the grain boundary.

When this material is heated in oxygen atmosphere no second phase is formed.

314
boundaries may very well act as a weak links. Also grain boundaries are effected

by a decomposition due to the thermodynamical instability of YBa2Cu307-3

below 800°C. This is seen most pronounced for Lai +xCai -xBaCuaOy-a, which

after annealing at 450°C in air or insufficiently dried oxygen, shows formation

of amorphous and crystalline material at many of the grain boundaries (see

Figure 5). Similar effects, although less pronounced, are observed for

YBa2Cu307-3. The thermodynamical instability can also lead to the formation of

planar defects, caused by the intercalation of an extra CuO layer.

In this respect the question arises whether grain boundary engineering

will finally lead to materials with critical currents high enough for applications.

As is shown above, the low critical currents in polycrystalline materials are

probably caused by a combination of factors: a dead layer at (001) grain

boundaries, dislocations at grain boundaries, and degradation at grain boundaries.

It will be difficult to overcome all problems completely. Dominancy of (001)

grain boundaries can possibly be prevented using Lai +xCai-xBaCu307-a, or

sintering at lower temperatures. The structure degradation at grain boundaries

is due to the thermodynamic instability below 800X and is difficult to prevent

completely. But a careful choice of annealing temperature, time and atmosphere

is expected to lead to an acceptable minimization of the degradation. Dislocations

at grain boundaries will be rather difficult to prevent in polycrystalline

materials and will be also generated by difference in thermal expansion along the

crystal axes. STRUCTURE ANALYSIS USING HREM IMAGES

In general analysis of defectless structures can be done best with X-ray

single crystal diffraction or neutron powder diffraction. However HREM can be

very useful in situations where no good diffraction data can be obtained e.g. because no single crystals are present and the material is not suited for powder

diffraction because the crystallites are too small, multiply twinned, or contain a

superstructure (often not observable with powder diffraction techniques). This

will be discussed using the superstructures in the compounds Bi2Sr2CaCu208+3

and Bii.5Pbo.5Sr2CaCu208+a.

Bi2Sr2CaCu208+a. Only

Figure 6. Schematic

representation of the

modulated structure of

the cations are shown

(black=Bi, white = Cu,

small shaded=Ca, large

shaded=Sr).

315
Since superconductivity at high temperatures was discovered in the

systems Bi203-SrO-CaO-CuO [37] much structural research has been done on

these compounds. It was quickly established that Tc's of 80 and 110 K were

related to the phases Bi2Sr2CaCu208+a [38,39] and Bi2Sr2Ca2Cu30i o+a

[40], which are the n = 1 and n = 2 members of the series

Bi2Sr2CanCui +n06 + 2n + 8- The structures of the Bi containing high

temperature superconductors are modulated (see references in [41]), which can

be described with a modulation vector, ki , which can vary between 3.5b and

5.0b, b being the b axis. It is schematically represented in Figure 6. Concensus

is growing on the position and the occupation of the oxygen atoms in or near the

Bi double layer. Most of the latest results [42-46] seem to confirm the presence

of extra oxygen atoms in the BiO layers, which was first suggested by us [47]

from electron microscopy observations. We explained the superstructure with an

insertion of extra oxygen in the BiO layer to adjust partially the mismatch of the

BiO double layers and the perovskite block between the BiO double layers.

Figure 7. Schematic

abed

representation of [100] diffraction patterns of Bi2-xPbxSr2Ca1

yYvCu20 8 + a with a) x=0,y=0.6, b) x=0.3,y=0.3, c) x=0.3,y=0.0, and d)

x=0^6,y=0.0, showing a) the Bi-type of modulation, b) and c) the mixed Bi/Pb type of modulation, d) the Pb-type of modulation.

Pb substitution for Bi is reported [48] to have a positive effect of the

formation of the n=2 phase, which has a Tc of about 11 OK. A number of groups

[41] have shown that partial substitution of Bi by Pb leads to the formation of a

superstructure which is different from that of the pure Bi compounds. Schematic

representations of the diffraction patterns are shown in Figure 7. The modulation can be represented with a vector k2, being approximately 8b, and will be called

hereafter the Pb-type of modulation. No sudden change occurs from the Bi-type

to the Pb-type of modulation with increasing Pb content. Over a considerable

region both modulations coexist. Both modulations interact as is evident from the

presence of combined modulation vectors ki+k2 and k-|-k2 [see ref 41]. The

presence of the combinations of qi and q2 indicates that both modulations occur

in the same region.

The model we proposed for the modulated structure consists of two

aspects. First, incorporation of clusters of extra oxygen in the BiO layer, as

316
depicted in figure 8, is assumed in order to get a better lattice matcli of the

perovskite block and the BiO layers. The clusters of neighbouring BiO layers

within one {BiO)2 double layer are located at the same positions along the b axis,

this forming double clusters. This leads to the modulation in neighbouring BiO

layers to be in anti-phase (see figure 9). The oxygen rich clusters form strings

along the a axis. The insertion of the oxygen rich clusters will lead to a local

expansion of the BiO layers, which allows contraction elsewhere, thus resulting

in a modulation of the structure. Second, a shift of the atoms in the BiO layers in the [100] direction is suggested to be caused by the lone pairs. In that model both

aspects were not directly correlated to each other.

BiO

ribbon

BiOj

ribbon

BiO

ribbon

b

Q one O Q ^ * H ^' ^'^'^ P^'*^

Figure 8. Schematic representation of the effect of the introduction of extra oxygen in the BiO layer. In the oxygen rich region the lone pairs point no longer to an oxygen of the adjacent BIO layer allowing a shorter BIO-BIO layer distance as is Indicated in b).

The basic idea in the model for the Bi modulation is the compensation for the lattice mismatch by patches with extra oxygen. The model of the Pb type of

modulation should be based on the same principle. Because the substitution of

Bi3+--> Pb2+ will probably result in less oxygen in the (Bi,Pb)0 layers the

Pb-type of modulation must be correlated to a smaller amount of oxygen rich

clusters. These oxygen rich clusters cannot be ordered as in the Bi-type of

modulation, in which double clusters occur, because the oxygen rich regions

would be too far apart. This leads to a model for the Pb-type of modulation in which oxygen rich regions in adjacent BiO layers are separated as much as

possible resulting in the model as shown in Figure 9. Electron diffraction patterns (see Figure 7) show superreflections qi-i-q2

and similar combinations and also that q2/qi is either close to 2/3 or 1/2.

Interaction between both modulations without any changes in these modulations Is

317
difficult to imagine, because the displacements along the c axis of the two types of

modulations are incommensurate with each other. However interaction can still occur in a matrix in which the atoms are displaced in the c direction according to a Pb-type of modulation. This situation is schematically represented in figure

9c. With such a model of the mixed modulation one would expect that for a dominant Pb-type of modulation the amplitude along the c axis of the Bi-type of

modulation is much smaller than that for the pure Bi-type of modulation. This is because the c components of both modulations are not compatible. This

expectation is in agreement with a recent publication of Gao et al [49] on a X-ray

single crystal diffraction study on Bi2-xPbxSr2Cu06+3 with x=0.3 showing

that the c component of the Bi type modulation in the BiO layer is almost zero.

Bi-type

—o— —omixed-type

Pb-type

Figure 9. Schematic representation

of models for the modulation in Bi2-xPbxSr2CaCu208+3 with a)

x=0, and b) x=0.5. The spheres

represent oxygen rich regions and

the curved lines the BIO layers.

Unit cells are indicated by

straight lines, c) represents a

situation of intimate mixing of both

modulations.

DISCUSSION

The analysis of HREM images has to be done very carefully, because the

observed image is not a direct image of the structure, but an image, modified by

the electron microscope. This is done partly by the imperfections of the

microscope and partly by the defocus of the microscope, which is chosen by the

microscopist, because specific structural features are observed best at well

defined defocus values. For experimental images the exact defocus value is not

known exactly. At best the defocus can be determined with an accuracy of about

10 nm. This shows that careful calculations are necessary, and that conclusions

318
regarding the positions of atoms should never be done using one experimental

image, but always using several images of a through focus series, for which the

defocus step is known.

Mostly the comparison between calculated and experimental images is done by eye. However, the printing of the experimental images as well as the

calculated images is mostly not a linear process. Thus, comparison by digitized

gray values is much more accurate, an example of which is given in ref. 1. Nevertheless, using comparison by eye several conclusions can be made in combination with other data from the high resolution images. In the first place

the positions of the mirror planes is not effected by the printing. Also symmetry

properties of the crystal can be lost because of misalignments of the microscope

but not gained. Consequently the symmetry in the experimental image is the

minimum symmetry of the structure. By comparison of experimental and

calculated images it is very often possible to correlate the existing mirror planes

unambiguously. Thus, when defects or intergrowth occur, the symmetry in the

experimental image can be used to rule out a large number of locations of these

defects. Secondly in the case of planar defects, the increase or decrease in the

period perpendicular to the planar defect also restricts the number of models and

image calculations can help with a further restriction. Often the high resolution

image often allows one to determine the number of layers which are added or omitted in the defect.

Since the structure determination using image calculations is a trial and

error method, whereby accurate atomic positions are not obtained, it is evident

that it is tried to obtain a direct image of the structure. This is in principle

possible if one knows how the microscope transforms the image. Knowing the

transfer function of the microscope, the experimental image can be deconvoluted,

to obtain the wave function at the exit of the crystal. Next the projected potential

has te be retrieved using a sophisticated computer program, tracing the electron

back to the entrance of the crystal. Having thprojected potential, atomic positions can determined much more accurately. It is expected that such a procedure will be available in near future making structure determination with high resolution

electron microscopy much more reliable. When one is able to retrieve the

projected potential, the only important characteristic of the microscope will be

the information limit. The advantages of the possibility to do electron microscopy with EDX

element analysis are large in the search for new superconductors and in the

determination whether intended substitutions are realized. For instance, many

papers have been published on the substitutions in YBa2Cu307, without a

verification whether the substitution was completely successful. Performing EDX

element analysis on crystals, which have the 1-2-3 structure according to electron diffraction, can show very easily whether the substitution was complete.

In this paper some clear examples of structure/property relations are

given. However, this relation between many defects of various nature and their

influence on the properties have not been established. The determination of such

relations will be difficult because a given material will contain various defects.

Small changes in composition, not detectable with electron microscopy, will further complicate such a determination. Evidently very careful studies are necessary on well characterized samples with different physical properties.

319
Acknowledgement

The author wish to thank Prof. G. van Tendeloo (University of Antwerp) and Prof.

R. Gronsky (Lawrence Berkeley Laboratory) for the use of some electron

microscopes and valuable discussions.

REFERENCES

1. H.W. Zandbergen et al, Nature 331, 596 (1988)

2. P. March et al, Nature 334 141 (1988)

3. E. Kaldis et al, Physica C 159, 668 (1989)

4. H.W. Zandbergen et al, Nature 332, 620 (1988)

5. H.W. Zandbergen et al, Solid State Comm. 66, 397 (1988)

6. H.W. Zandbergen et al, AppI Phys. A46, 233 (1988)

7. R. Beyers, et al, AppI. Phys. Lett. A 53, 432 (1988)

8. H.W. Zandbergen et al, Physica C 158, 155 (1988)

9. H.W. Zandbergen et al, Physica C 159, 81 (1989)

10. H.W. Zandbergen, and G. van Tendeloo, Phys. Rev B40, 11300 (1989)

11. T. Roulllon et al, Physica C 159, 201 (1989)

12. H.W. Zandbergen. R. Gronsky and G. Thomas, Phys. Stat. Sol. (a) 105, 207 (1988)

13. D.J. Li, et al, Physica C, 156, 204 (1988)

14. H.W. Zandbergen and G Thomas, Phys. Stat Sol (a) 107, 825 (1988)

15. A.F. Marshall, Phys. Rev B, 37, 9353 (1988)

16. P. Bordet et al. Nature 334, 596 (1988)

17. R.J. Cava et al. Nature 338, 338 (1989)

18. H.W.Zandbergen et al, AppI. Phys. A46, 153 (1988)

19.. C. Michel et al, Z. Phys. B68, 421 (1987)

20. M.A. Subramanian et al. Science 239, 1015 (1988)

21. J.M. Tarascon et al, Phys. Rev. B57, 9882 (1988)

22. R. Ramesh et al. Science 247, 57 (1990)

23. LA. Tietz, J. Mater. Res. 4. 1072 (1989)

24. H.W. Zandbergen et al, Physica C, 166, 255 (1990)

25. M.S. Horowitz et al. Science 243, 66 (1989)

26. H.W. Zandbergen, J. van der Berg, and P.H. Kes, Cryogenics 1990, 633 (1990)

27. R.Feenstra et al, AppI. Phys. Lett. 54, 1063 (1989)

28. C.B. Eom et al, presentation at MRS Spring Meeting San Francisco 1990

29. S.E. Babcock et al, Proc. Int. Conf. Cryogenic Materials, Garmisch Partenkirchen,

May 9-11 (199)

30. H.W. Zandbergen and G. van Tendeloo, Mat. Res. Soc. Proc. Vol 156, 209 (1989)

31. T. Kogura et al, Physica C, 156, 707 (1988)

32. S.E. Babcock, D.C. Larbalestier, J. Mater. Res. 5, 919 (1990)

33. H.W. Zandbergen, and G. Thomas, Acta Cryst A44, 772 (1988)

34. T. Yamamoto et al, Supercond. Sci. Technol. 1, 153 (1988)

35. H.W. Zandbergen et al, Cryogenics 1990, 628 (1990)

36. D. Dimos et al, Phys. Rev. Lett. 61, 219 (1988)

37. H. Maeda, Y. Tanal<a, M. Fukutomi, T. Asano, Jap. J. AppI. Phys. 27 (1988) L209

38. J.M. Tarascon et al, Phys. Rev. (1988)

39. H.G. von Schnering et al, Angewante Chemie 27, 574 (1988)

40. H.W. Zandbergen et al. Nature 332, 620 (1988)

41. H.W. Zandbergen et al, Physica C, 168, 426 (1990)

42. V. Petricek, Y. Gao, P. Lee, P. Coppens, preprint 43. G. Calestani, C. Rizzoli,M.G. Francesconi.G.D. Andreetti, Physica C 161 (1989) 598

44. A. Yamamoto, M. Onoda, E. Takayama-Muromachi, F. Izumi, T. Ishigaki, H. Asano,

Phys. Rev. B preprint 45. Y. le Page et al, Phys. Rev. B40, 6810 (1989)

46. J.M. Tarascon et al. Phys. Rev. B39, 11587 (1989)

47. H.W. Zandbergen, Physica C, 156, 325 (1988)

48. S.A. Sunshine, et al. Phys. Rev. B 38 (1988) 893

49. S.M. Green, C. Jiang, Y. Mei, H.L Luo, C. Politis, Phys. Rev. B 38 (1988) 5016

50. H.W. Zandbergen and W.A. Groen. Physica C 166 (1990) 282

320
DISCUSSION

A. Sleight : Many people have, of course, made 123 at low temperatures and have seen

the Meissner effect disappear, even though it looks like it's still exactly 123-oxygen 7. But when we did specific heat measurements on those samples we found they were

completely superconducting. The Meissner effect disappears because the particle size gets

smaller than the penetration depth of the magnetic field. So I would assume that your

materials are really superconducting although the Meissner effect is not indicating

superconductivity. It is disappearing just because the particles are becoming smaller than

the penetration depth. -

H. Zandbergen : In our experiments with changing preparation temperatures we tried to keep the particles at the same size by variation of the time of preparation. Still I expect

that the material prepared at low temperature will contain many weak links because of the

many [100] 90 degree rotation twin boundaries. The magnetic field can penetrate the

particles through the weak links which will result in a situation as if you have very small

particles.

A. Sleight : The field can penetrate?

H. Zandbergen : Can penetrate, yes, along the [100] 90 degree rotation twin boundaries.

A. Sleight: I think the only way to check the presence of superconductivity in such

samples is by doing specific heat measurements.

H. Zandbergen : I agree.

M. Marezio: It's very dangerous to determine by measuring the Meissner effect whether

a compound is superconducting, as Sleight has already said. I suggest you check this

with other types of measurements.

I would like to ask you what the valency of the Bi atoms is when you put extra

oxygen in the BiO layer and you make Bi02 patches. The Bi-0 distances are very short

and my guess is that the Bi is probably five-valent. In that case you don't need a lone

pair chemistry since Bi^"*" does not have a lone pair. The Bi-O distance has been

determined with X-ray diffraction, but the errors are so big that one cannot decide on the

valency of that Bi atom. My feeling is that in the model with the extra oxygens the BiO

distances are rather short.

R. Cava: Yes, they are. I have no idea about the lone pair effect. What I do know is that the structure is slightly or even more complicated.

H. Zandbergen : We cannot isolate the bismuth oxide layer from the rest of the structure. We know that Bi can partially substitute Sr. If you put extra oxygen in the BiO layer one would expect that the Bi in the SrO layer is situated around the position of the extra

321
oxygen. Thus you do not need very much charge compensation for the extra oxygen if there is extra Bi on the Sr sites near that regon. Concerning the exact distances I have no good values because they will depend on the amplitude of the modulation.

M. Marezio: But the Bi-O distance will be close to half of the diagonal of the average

structure and that is short.

H. Zandbergen : Yes, it's short but the distance from a Bi atom to its nearest neighbor

oxygen atom in the SrO layer is also very short, so it is a funny structure anyway.

R. Roth : Has anybody done calculations of the bond valence sum around Bi based on good X-ray data?

D. Brown: Because of the incommensurate structure of the Cu compounds no one has

been able to look at it properly. I can't remember exactly but I must have calculated the

bond valence sums. I don't think there is any reason to believe the Bi atoms are anything

other than trivalent. The extra oxygen results in extra charge on the CUO2 layer and not on the BiO layer.

M. Marezio: X-ray absorption edge measurements have been done and no Bi^"*" is observed. The value of this result depends of course on the accuracy of the technique.

D. Smyth : If you look at the actual detailed compositions of these materials which Art

Sleight has described several times and add up the cations you find that it is positive

charge deficient relative to the ideal composition that you show here, which probably

means that it is oxygen deficient. So it should be relatively easy to slip some more

oxygen into this material.

M. Marezio: It is a combination of the two. This has been discussed at Bordeaux during

the International Meeting of Crystallography, Coppens, who observed the cation

vacancies with X-ray diffraction, and others, who described the superstructure based on

X-ray diffraction, agreed that the complicated structure is due to a mixture of excess oxygen and cation vacancies. On the other hand there are quite a number of reports

which state that the various sites are fully occupied.

H. Zandbergen : The problem with the determination of the occupancies with diffraction

techniques is the variety of possible substitutions: Ca on the Sr site, Bi on the Sr and Ca

sites, Sr on the Ca site. So it is very difficult to determine whether cation vacancies do

occur.

R. Cava: But there is virtually no crystallographic study which is consistent with the

cation ratio that you give here.

H. Zandbergen : No, that is right.

R. Cava: Virtually none.

322
H. Zandbergen : Right. I and many others are only using this notation for reasons of

simplicity.

P. Davies: I thought that AEM work has shown very convincingly that the Bi phase is quite Sr deficient.

H. Zandbergen : That's why I said that Sr is partly substituted by Bi thus probably

stabilizing the structure and leading to Sr deficient material.

P. Davies: My other question relates to the Pb doped phase, when you go from the "inphase" to the "out-phase" type of modulation due to the reduction in the number of

oxygen ions in the Bi layer. When you take a little oxygen out of the Bi material do you

see any change in the nature of the modulation?

H. Zandbergen : You cannot take enough oxygen out of pure Bi compounds to change

the type of modulation. But the oxygen content can be increased easily by partial

substitution of calcium by yttrium or another trivalent atom. With such a substitution you

have to put in extra oxygen in the structure. Indeed with this substitution the modulation

vector is decreasing which is consistent with an increase of oxygen in the bismuth oxide

layer.

A. Nath : There are many groups which have been heating the 123 compounds between

500 and 700°C for long times. Could you eleborate on your statement that the compound

is not stable in oxygen below 800°C? What kinds of effects do you expect?

H. Zandbergen : I am expecting mainly amorphization. Apart from many clean grain

boundaries we have also seen many grain boundaries (in particular for LaBa2Cu307) with

contained amorphous phases at the grain boundary and some small CuO crystallites.

323

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.CRYSTAL CHEMISTRY OF OXIDE SUPERCONDUCTORS

Arthur W. Sleight

Oregon State Univ., Corvallis, OR

INTRODUCTION

This paper discusses some recent results for copper oxides

which contain the Cu02 sheets that are present in aU of the copper

oxide based superconductors. Another aspect of this paper is to

discuss oxidation state stabilization and oxidation state fluctuations as they might relate to a mechanism for superconductivity.

NEW MATERIALS

Bednorz and Miiller discovered (1) superconductivity in doped

La2Cu04. However, our attempts to produce superconductivity in

several closely related materials have been met with many

disappointments. The R2CUO4 (R = Pr - Gd) phases could not be

doped p-type to become superconducting (2), but they were eventually

doped n-type to become superconducting (3).

The two layer analog of La2Cu04 has been known for several

years (4), but the doping of La2CaCu206 to produce superconductivity

was reported only very recently (5). There had been a suggestion (6)

of structural reasons for this apparent inability to produce

superconductivity, but these reasons were not substantiated by

detailed structural studies (7).

Another structure very closely related to La2Cu04 is that of the

A2CUO2X2 series where A can be Ca or Sr and X may be CI or Br (8).

Attempts to produce superconductivity by either p- or n-type doping

have failed to date. We have recently (9) expanded this oxyhalide

family to include the two-layer versions, Ca3Cu204Cl2 and

Ca3Cu204Br2 (Figure 1). It is interesting that attempts to prepare

three-layer structures of the type A4CU3O6X2 or A4CU3O8 have so far

failed. Our attempts to dope Ca3Cu204X2 phases p-type have not

succeeded. However, Ca3.xYxCu204X2 phases have been prepared

which are presumably n-type (9). Superconductivity has not yet been

observed in this two-layer structure.

MIXED VALENCY AND THE PARTIALLY FILLED BAND

A feature common to both bismuth and copper oxides is the

availability of three consecutive oxidation states: Cu^, Cu^^, Cu^ and

Billl, Bi^V, BiV. All of these oxidation states are well known to

325
chemists, with the exception of Bi^^. We know that cations such as Bi^V are unstable when concentrated. Thus, we do not normally find

discrete compounds where bismuth is entirely in the tetravalent

state. However, we do find Bi^^ in dilute systems, both insulating

and metallic systems (10).

326
The lowest oxidation states for copper, CuKSd^^), and for bismuth, Bi^C^s^), are diamagnetic. The middle valency, Cu^KSd^) or Bi^(6s^), is always a one electron state. The highest oxidation state

is always diamagnetic in the case of bismuth, Bi^(6s^). However, the

3d^ state for Cu^^^ is paramagnetic for cubic (octahedral or tetrahedral) copper but is diamagnetic for square planar copper. It would

seem significant that all copper oxide superconductors have a square

planar arrangement of oxygens around copper. There may be a fifth, or a fifth and sixth oxygen, ligand; however, these are always at

much longer distances. Thus, the mixed valency analogy between

bismuth and copper is complete for the structures where superconductivity is observed. Both the lower and upper oxidation states are diamagnetic and the middle oxidation state is spin one half.

The mixed valency in any given superconductor may be viewed

as involving just two oxidation states: Cu^^'^^^ in La2-xSrxCu04,

CuI'II in Nd2-xCexCu04 and Bi^>^ in Bai-xKxBiOa. It is essential for

metallic properties and for superconductivity that the two oxidation

states be present on one crystallographic site. Thus, YBa2Cu306 is

neither metallic nor superconducting because Cu^ and Cu^^ occupy

distinct crystallographic sites and there is no mixed valency on either

site. On the other hand, YBa2Cu307 is metallic and superconducting, and this is related to Cul^'^ mixed valency in the Cu02 sheet

sites.

The mixed valency situation may be equally well described in

terms of bands. The band to be considered in the case of bismuth is

the Bi6s - 02p band. The band in the case of copper is the Cu3dx2.y2 - 02p band. Both of these are a* bands. For the lowest oxidation state

(Cu^ or Billl), these bands are filled; for the highest oxidation state

(Cu^^^ or Bi^), these bands are empty. According to simple band

structure considerations, we would expect metallic properties for any

partial filling of these a* bands in bismuth and copper oxides. While

metallic properties are indeed observed for most of the intermediate

band fillings, insulating properties are always foimd for the halffilled band in both cases, i.e., the situation for pure Bi^^ or pure Cu".

There are two competing mechanisms to produce an insulating state for the*half-filled a* band. One is through valency disproportionation, and this is the case for BaBiOa, i.e., 2Bi^^ -> Bi^^^ -i- Bi^.

If Bi^V had not disproportionated in BaBiOa, we would expect this

compound to be metallic because it would have a half-filled band.

Many have been confused by this disproportionation description

because of confiision about the meaning of valent states as opposed to

real charges (10).

327
Disproportionation of Bi^V in BaBiOa leads to two distinct sites

for Bi^ and Bi^, Thus, we formulate this compound as Ba2BimBiV06. The Bi-0 distances are those expected of these oxidation states (10). The real charges on Bi^ and Bi^ are very much

lower than the oxidation states due to covalency effects. Thus, the

real charge difference between Bi^ and Bi^ is much less than two.

An equivalent description of the disproportionation situation in

BaBiOs is the term charge density wave (CDW),

There is another way to achieve an insulating state for the

half-filled band, and that is through antiferromagnetic ordering.

This happens for all of the copper oxides which contain copper only

as Cu^I. Thus, spin pairing of electrons on Cu or Bi sites is always

related to the insulating state of the Cu'^ and Bi^ oxides. For the

copper oxides, it is spin pairing between adjacent Cu^^ sites. In the

case of Bi^V, spin pairing occurs on alternate Bi sites creating BilH

and BiV. The magnetic situation in the Cu^I oxides may also be

described as a spin density wave (SDW).

We will refer to the two competing mecheinisms for localization

at the half-filled band as disproportionation and magnetic. We could

say that if BaBiOs did not become insulating through disproportionation, it would have become a BaBi^^Oa antiferromagnetic insulator.

However, muon spin rotation experiments have completely ruled out

that imlikely possibility (11). We could also say that if Cu^^ oxides did

not become antiferromagnetic insulators, disproportionation into Cu^

and Cu^I^ would have occurred. This is, of course, exactly what

usually happens for Ag^I and Au^I. For AgO, the oxidation state of

silver is clearly not two. This compound is instead Ag^Agm02.

Schematic band structures for the copper and bismuth oxides

just discussed are given in Figure 2. The hypothetical half-filled

band metals do not exist due to electron-electron correlations which

cause band spliting. The n* band is placed so that it is always below

the Fermi level. This is consistent with all band structure calculations. However, band structure calculations are not reliable on this

issue because they ignore the electron-electron correlations which

split the o* band and which could cause the filled portion of the a*

band to drop below the top of the k* band. The primary assurance

that the ti* band always lies below the Fermi level is that an obvious

symmetry exists at the Fermi level, i.e., doping either p- or n-type

produces essentially the same result (Figure 3). This would not be

expected if the t: band overlaped the Fermi level. The consensus from

nuclear resonance studies is also that there are carriers in only the

G* band (12).

328
Metal

La2Cu04

BaBiOa

(b)

Insulator

La2Cu04

BaBiOg

(c) (d)

Superconductor Superconductor

(La. Sr)2Cu04 (Nd. Ce)2Cu04

(Ba. K)(Bi. Pb)03 Nd2Cu{O.F)4

Figure 2

Schematic band structures for pure and doped BaBiOa and La2^04. The metallic

compoimds are hypothetical materials without disproportionation or magnetic interactions

which cause band splitting when the band is halffilled. Both the p-type (c) and n-type (d)

conductors are shown.

329
STABILIZATION OF OXIDATION STATES

We know that the 0"2 ion is not stable in the gas phase. Such a

species would spontaneously loose an electron £ind become 0"l.

Nonetheless, we correctly conclude that oxygen is a divalent in nearly

all of its compounds, including O2 and 02"^. This is not surprising in

view of the differences in real charges, oxidation states and valent

states just discussed. However, there are two very different ways to

describe the stabilization of divalent oxygen. One way has a covalent

approach, the other an ionic approach. In highly covalent compounds such as O2, CO2, SO3, and OSO4, we would refer to oxygen as

divalent even though the real charge on oxygen is zero or just slightly

negative. Thus in this case, the problem of stabilizing the O"^ ion

never occurs because the real charge on oxygen never approaches

two.

In oxides of high ionicity, it is useful to employ another

approach to rationalize the existence of divalent oxygen. This alternative approach has particular relevance to the copper and bismuth

oxides despite the covalency of the Cu-O and Bi-0 bonds. Although

the 0"l ion is not stable in the gas phase, it can be stabiHzed as a fully

doubly negatively charged O"^ ion in a solid through Madelung

forces. The cations surrounding oxygen in a solid produce an electrostatic field intense enough to capture the second electron. This

means that we have so stabilized the oxygen 2p levels that the second

electron may be added despite the strong on-site Columbic forces

which repel this second electron.

One of the common features of the copper and bismuth oxides

superconductors is the presence of highly electropositive cations such

as Ba'I. These cations may have several important roles such as

enabling certain framework structures and increasing the covalency

of the Cu-O or Bi-0 bonds (13). One very crucial role of the

electropositive cations is the stabilization of oxygen 2p levels to an

extent which allows for imusually high oxidation states for copper

and bismuth. We know that in the binary Cu/0 and Bi/0 systems

neither Cu2^''03 nor Bi2^05 can be produced even at high oxygen

pressure. However, introduction of highly electropositive cations

stabilize oxygen 2p states to such an extent that Cu^^^ and Bi^ can

exist. Lacking the electropositive cations, Cu^^^ or Bi^ would readily

oxidize O^' to 02"^^ or O2. Good examples of Cu^^^ and Bi^ are in

KCu02 and KBiOa where potassium serves the role of stabilization of

oxygen 2p states sufficiently to allow Cu'^^ or Bi^ to exist. Highly

330
electropositive cations might also play a role in stablizing Cl"^ in the

presence of Cu^ or of 1'^ in the presence of Cu^^.

The disproportionation of Bi^V and Cu^^ is shown schematically in Figure 4. This disproportionation can only be understood if

one takes into account the high covalency and the stabilization of Bi^

and Cu^^^ by highly electropositive cations such as Ba^^. The main

point is that their is a redistribution of electrons in the Bi-0 and Cu-0

bonds. Although a charge density wave is produced, the magnitude

of the charges is much less than would be implied in an ionic model.

The situation is in fact very much analogous to what occurs in

BaTiOa on going through its ferroelectric transition (Figure 4). The

main difference is that disproportionation is an antiferroelectric,

rather than ferroelectric, process.

MECHANISM FOR HIGH Tq

There is still no consensus for a mechanism for the high

Tc in the cuprate superconductors. Nonetheless, we have

learned much about the electronic structure of such materials.

Some proposed theories may now be discarded. The discovery

of the n-type cuprate superconductors was the clinching evidence needed to discard theories based on some unique feature

of an oxygen 2p band, a tc* band, or overlapping bands. The

central question now for the cuprate superconductors is the

fate of spin correlations as the Cu^^ based insulators are doped.

Certainly, these spin correlations are diminished with either

n-type or p-type doping. Long range magnetic order disappears, and the materials become metallic. One line of thinking is that it is a vestige of this magnetic state which is

indirectly responsible for superconductivity. A more traditional view is that magnetism and superconductivity tend to be

mutually exclusive phenomena. The first role of doping is

then to weaken the spin correlations sufficiently so that charge

fluctuations can then dominate aind produce superconductivity. Only this latter view is compatible with both the Bi-based

and the Cu-based superconductors, and only this view will be

further discussed here.

Some of the features common to both the copper-based

and bismuth-based superconductors have been discussed

earlier in this chapter. One might say that these two types of

superconductor differ in their isotope effects in a way that suggests a different mechanism, but this is not really the case. It

is now generally agreed that the oxygen isotope effect in

YBa2Cu307 is finite but quite small. However, it is also known

331
Disproportionation of Bi and Cu"

IV IV IV rv IV IV —Bi —O—Bi —O —Bi —O —Bi —O —Bi —O—BiV i "I I V A I" I V i III -O-Bi-O-T-Bi-t-O-Bi-O-l-Bi-i-O-Bi-O-l-Biin A I

I III 4 It III i I -O-Cu-O-I-Cu-t-O-Cu-O-l-Cu-t-O-Cu-O-t-CuBond Disproportionation in BaTiO^id^)

- + - + - + - + - + - +

0-:--Ti-:.-0-:.-Ti-:.-0-:.-Ti-:.-0-:.-Ti-:.-0-:.-Ti-:.-0---Ti-:.

-+ -+ -+ -+ -+ - +

-0::Ti-:-0::Ti— :-0::Ti-:-0::Ti-:-0::Ti-:-0::TiDistortions with filled d shell

-O-Sn-O-Sn-O-Sn-O-Sn-O-Sn-O-Sn-O-

• « «•

••

• •

Figure 4

Disproportionation and distortions in the perovsldte structure. Arrows represent the spin of

antibonding electrons in the BiDLo^D Cu^-O-D bonds. For BaTiOa, the dots represent

bonding electrons, a and n. For Sn and Cu, the pair of dots represent a pair ofnonbonding

electrons resulting from bond bending.

332
that this fact alone cannot be used to discount a possible critical

importance of the electron-phonon interaction in such superconductors. The oxygen isotope effect is pronounced in both

La2-xAxCu04 and Bai-xKxBiOa superconductors (14,15).

Thus, we can be rather certain that the electron-phonon interaction is important for all these superconductors, including

the cuprate superconductors with Tc's above 100 K.

I continue to believe that the mechanism for high Tc lies

in the tendency for disproportionation of the middle valence

state, i.e., Cu^^ or Bi'^ (Figure 3). Such a tendency is well

known in chemistry, as previously discussed in this chapter.

A problem has been, and remains, that this tendency to disproportionate is not well understood in terms of physics.

Nonetheless, this is a real phenomenum which inherently

gives the electron-electron attraction required for superconductivity. Such an approach to forming Cooper pairs necessarily

means that there will be an electron-phonon interaction. Still,

there may be important charge fluctuations that are not so strongly coupled to phonons, e.g., on site polarizibility. This

may also make a significant contribution to increasing Tc . A

serious problem in developing this view is that charge fluctuations in ferroelectrics such as BaTiOa (Figure 4) have never

been well understood from a physics perspective.

References

1. Bednorz, J. G. and Miiller, K. A. Z. Phys. B64:198-193

(1986).

2. Gopalakrishnan, J., Subramanian, M.A., Torardi,

C.C., Attfield, J.P., and Sleight, A.W., Mat. Res. Bull.,

24:321-330, (1989).

3. Tokura, Y., Takagi, H. and Uchida, S. Nature 337:345

(1989).

4. N. Nguyen, N., Er-Rakho, L., Michel, C, and Choisnet, J.,

Mat. Res. Bull. 15:891,(1980).

5. Cava, R.J., Batlogg, B., van Dover, R.B., Krajewski, J.J.,

Waszczak, J.V., Fleming, R.M., Peck, Jr., W.F., Rupp, Jr.,

L.W., Marsh, P., James, A.C.W.P., and Schneemeyer, L.F.,

Nature, 345:602-603 (1990).

333
6. Hervieu, M., Maignan, A., Martin, C, Michel, C, Provost, J.,

and Baveau, B., J. Solid State Chem., 75:925 (1988).

7. J. Huang and A.W. Sleight, Physica 0,169:169-173 (1990).

8. Grande, Von B. and Miiller-Buschbaum. A. Anorg.

Allg. Chem., 429:88 (1977).

9. Huang, J., Hoffmann, R.-D., and Sleight, A.W., Mat.

Res. Bull., in press.

10. Sleight, A.W., Valency Considerations in Oxide

Superconductors , Proc. Welch Conf. on Chem. Res.,

XXXIL Valency, p. 123, 1989.

11. Uemura,Y.J., Sternlieb,B.J., Cox, D.E., Brewer, J.H.,

Kadono, R. Kempton, J.R., Kiefl, R.F., Kreitzman,

S.R., Luke, G.M., Riseman, T., Williams, D.L.,

Kossler, W.J., Yu, X.H., Stronach, C.E., Subramanian,

M.A., Gopalakrishnanan, J., and Sleight, A.W.,

Nature, 335:151(1988).

12. Walsted, R.E. and Warren, W.W., Science, 248:1082

(1990).

13. Sleight, A.W., High-Temperature Superconductivity in

Oxides in Chemistry of High-Temperature Superconductors, (Nelson, Whittingham, George, eds.) p. 2,

Amer. Chem. Soc, Wash. D.C.,(1987).

14. Crawford, M.K., Kunchur, M.N., Farneth, W.E.,

McCarron, E.M., and Poon, S.J., Phys. Rev. B., 41:282

(1990).

15. Hinks, D.G., Richards, D.R., Dabrowski, B., Marx, D.T.,

and Mitchell, A.W., Nature 335: 419 (1988); Batlogg, B.

Cava, R.J., Rupp, L.W., Mujsce, A.W., Krajewski, J.J.,

Remeika, J. P., Peck, W.F., (hooper, A.S., and Espinosa,

G.P. Phys. Rev. B 61:1670 (1988).

334
DISCUSSION

H. Zandbergen : You showed a correlation between and modulation periodicity for

the Bi2Sr2.xYjjCaCu20g^y type phases, but what was the actual compositional variation?

A. Sleight : Both yttrium content and oxygen content are varied to obtain the entire range

of modulation periodicity from 5.0 to 4. 1. The phase with the periodicity of 5.0 contains no yttrium and is obtained by extracting oxygen from the lattice. Increasing the yttrium

content results in a decrease in the modulation periodicity, but 4.0 is never quite

achieved.

H. 2^dbergen: The most important thing you are changing is probably the Cu-0

distance along the c axis.

A. Sleight : Perhaps, but we do not have the adequate crystallographic data to comment

on that. We might infer the Cu-0 distance within the CUO2 sheets from the a and b cell dimensions. This approach indicates the expected decrease in this Cu-O distance as the

copper is oxidized. However, this is complicated by the fact that the CUO2 sheets are not

flat. Thus, the a and b cell dimensions do not accurately reflect the Cu-0 distances in

the sheets.

B. Dobrowski : In the case of EXAFS data from the Ba(Bi,Pb)03 system, you still see two Bi-0 distances for the metallic BaPbg j^^^o 25^3 composition. How then can you

explain the metallic properties?

A. Sleight : I agree that I would not have expected two Bi-0 distances in the metallic

region. Perhaps the BaPbo js^^o 25^3 sample has regions where bismuth has

disproportionated and regions where it has not disproportionated. Homogeneity is an

issue here. Some optical studies also suggest that the charge density wave extends into

the metallic region.

B. Dabrowski: We did the same EXAFS experiment for the (Ba,K)Bi03 system. There

you see disproportionation in the nonsuperconducting region but only one Bi-0 distance

for the superconductors.

A. Sleight : That is what I would expect to see. We are planning to reexamine the

EXAFS of superconducting Ba(Pb,Bi)03 on another sample.

P. Davies : Do you actually know the carrier type in your oxychlorides?

A. Sleight : Our only evidence is based on the changes in unit cell dimensions on

substitution of yttrium. Because Y^"^ is smaller than Ca^"*", we see the expected decrease

in the c axis on replacing some Ca with Y. However, this substitution causes an expansion of the a axis and thus the Cu-0 distance. This expansion is only expected if

335
the average copper oxidation state has decreased below two. This is probably a more

significant result than a Hall or Seebeck measurement.

M. Takano: This is not a question but a comment. I want to say something on

disproportionation. We have looked at the disproportionation of iron(IV) into iron (III)

and iron(V), and found the effective charges of iron(III) and iron(V) to be +2.7 and

+2.9, respectively.

A. Sleight : Right. I liked your comment on what happens when covalency is introduced.

The reason that oxidation state arguments work, even though real charges are totally

dfferent than oxidation state values, is because the covalent mixing gives new orbitals

with the same symmetry as before mixing. Thus, the fact that real charges are very much

smaller than oxidation state values does not perturb our argument about

disproportionation

.

You find a real charge difference of only 2.0 for an oxidation state difference of two.

That seems about right to me given such high covalency.

336
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

OUT OF PLANE DISPLACEMENTS OF OXYGEN FROM THE

CUO2 SHEETS IN CA 85SR 15CUO2 BY

ATOM-PAIR DISTRIBUTION FUNCTION ANALYSIS

S. J. L. Billinge, P. K. Davies and T. Egami

Department of Materials Science and Engineering

University of Pennsylvania, Philadelphia, PA 19104-6272.

C. R. A. Catlow

Royal Institution, 21, Albemarle Street, London WIX 4BS,U.K.

ABSTRACT

Atom-pair distribution function (PDF) analysis of neutron powder

diffraction data has been carried out on the layered material Ca.85Sr.i5Cu02.

The result indicates that the Cu02 sheets in this material are not flat but

distort locally with oxygens displacing up and down from their average

positions in the plane by 0.1 ± 0.03A. A static lattice energy minimisation

calculation, using central force pair potentials, indicates that, in this ionic

approximation, such a distortion is energetically favourable with respect to

the flat planes. Thermal factors have been refined from the PDF for the case

of displaced oxygens and their temperature variation is in good agreement

with Debye theory. The distortions can be explained using ionic arguments

and are predicted by bond valence theory. The importance of these

distortions and the in-plane bonding environment, to the superconducting

cuprates is discussed.

INTRODUCTION

Ca.85Sr.i5Cu02 has a structure of infinitely repeating Cu02 planes,

separated by the alkaline earth ions Ca-^'*' and Sr-^'''.^ These planes

are essentially the same as those that occur in all of the high Tc

cuprate superconductors. Ca.85Sr.i5Cu02 is the endmember of the

A2B2Can-i Cun04+2n (A=Tl,Bi, B=Ba,Sr, n=l,2,3...) family of materials

2LS n —> 00. It has no complicated intergrowths, which are characteristic

of the superconductors, and the unit cell contains just four atoms. It is thus of interest to be studied as a model compoimd for the Cu02 planes.

The average structure is tetragonal with space group P4/mmm with a =

b = 3.86IIA c = 3.2052A and has been refined from an x-ray single crystal

337
study. ^ In common with the superconductors, large thermal factors (up to

<u-^>= 13xlO~'^A-) were required on the in-plane copper and oxygen sites, even though the average structure is apparently so simple.

The purpose of this work is to study the intrinsic local structure of the

Cu02 plane. A local probe of structure, the PDF analysis of neutron powder

diffraction data has been used. This analysis presumes no periodicity and

yields information about short range correlations between atoms. A brief

description of the analysis procedure will be given, before the results are presented. The structiire has been modelled using a static lattice energy

minimisation routine and a summary of these results is also presented.

PAIR DISTRIBUTION FUNCTION ANALYSIS

The method is widely used for structiural studies of glasses and liquids

but more recently has been successfully applied to crystalline systems,

where its utility is in elucidating non-periodic order in the structure. ^The

relationship of the PDF to scattered intensity is through a fourier transform.

The raw diffraction data is reduced to the total scattering function, S(Q),

after correcting for background, absorbtion and detector efficiency. This is then fourier transformed to get the PDF.

These PDFs can then be compared to simulated PDFs, calculated from

model structures, which are determined in real space, directly from the atom

positions in the model, with no recourse to scattering theory. An agreement

factor (A-factor) was defined to describe the fit of simulated PDF to data;

[/>dat(r) - /?sim(r; ak)]' dr] ^ A — min

Pdat(r) is the value of the measured PDF and /?sim(r; ak) is the value of the

simulated PDF, at r. This function was minimised with respect to variation

of thermal factors, ak, to produce a best-fit of model to data. An examf)le

of PDFs from the lOK data and model are shown in fig 1. EXPERIMENTS AND RESULTS

Ca.85Sr.i5Cu02 was prepared in powder form by solid state reaction of

CaCOs, SrCOa and CuO. This mixture was decarbonated at 850°C for eight

hours, then baked at temperatures up to 980° C for up to 60 hotirs, with

frequent grindings.

Neutron powder diffraction data were collected at the Intense Pulsed

Neutron Source (IPNS) at Argonne Natl. Lab. using the time-of-flight,

Special Environment Powder Diffractometer (SEPD).^ Fifteen grams of

338
sample were sealed in a vanadium sample holder with He exchange gas and

cooled using a Displex closed loop He refrigerator. Data were collected at

lOK, 150K and 298K.

r(A) Figure 1. PDFs for Ca.85Sr.i5Cu02 from data taken at lOK (crosses)

and best-fit model (line).

A Rietveld full profile refinement was carried out on the data.

Refinements were carried out in two space groups; P4/mmm and P4m2.

In the latter group, displacements of atoms out of the plane are allowed. No

improvement in Rw was obtained in the P4m2 group, though oxygen atoms

did refine off their in-plane sites by a small amount. There is little evidence

from Rietveld results for any deviations from planarity of these planes though

the possibiHty that this is occurring is not discounted by these results.

The neutron powder data were then converted to PDFs. A model

structure Wcis constructed and atoms displaced from their a\'erage positions

systematically. Various possible distortions are suggested by the large

thermal factors refined during conventional crystallographic analyses. Each

of these was tried in the current modelling. The distortions were: A shift

of oxygen up and down parallel to z, a shift of copper up and down parallel

to z, a shift of oxygen in-plane along directions perpendicular to the Cu-0

bond and a shift of oxygen parallel to < 101 > type directions. The last

distortion is suggested by the possibility of the disorder being some kind of

lattice dilation to accommodate the oversize strontium ions. In each case atoms were displaced up and down on alternating sites by a value ±A and

the thermal factors refined by least squares minimisation of the A-factor.

A was then increased in a stepwise fashion, repeating the process at each

step. In this way graphs of best-fit A-factor vs. A were plotted for each

distortion.

339
The A-factor curves diverge when peaks in the model are shifted to

such an extent with respect to the experimental PDF that only a poor fit

is possible. This provides a stringent test for the maximum extent of a

distortion. It was found that the best-fit A-factor diverged for A < 0.04A

for all of the distortions tried except when oxygen was displaced up and

down along the z-direction. This will be referred to as the Oz distortion.

Thus all of the distortions introduce peaks into the PDF which are not

compatible with the experimentally derived PDF, except for the Oz case. The results for Oz are shown in figs. 2a and 2b. Fig. 2a plots the best-fit

A-factors determined when isotropic thermal factors were used on all sites.

The advantage of this constraint is that the number of fitting parameters is only three in each refinement. The curves do not diverge before A = 0.11 A

and fall to a minimum before diverging. From this we can infer that the

PDFs indicate that a distortion of 0 along z up and down by 0.10 ± O.OSA

from the plane is probably occurring.

o

<

ino

d A

CM

O

q

6

3

CNI

uj '—

o>-X

o

o

d

figure 2. (a)Best fit A-factors for fit of model PDF to data at three

temperatures for the case of modelling using isotropic thermal factors,

(b) Lower curve is the A-factor curve for lOK data using anisotropic

thermal factors. The upper curve shows the refined value of <U33> from

the oxygen site as a function of static displacement of oxygen along z. The modelling was then repeated with anisotropic thermal factors and

the results for the lOK data are shown in fig. 2b. In this case there are seven independent thermal parameters. There is no obvious minimum in the

best-fit A-factor curve however it does show a plateau extending beyond

A = O.lA which diverges sharply thereafter. Also plotted are values of the

oxygen <U33> mean square displacement parameter refined from the PDF.

340
This falls by a factor of two from the case of no oxygen displacement to a

that of a 0.1 A Oz displacement.

The truly dynamic nature of a refined thermal factor may be gauged by

its temperature dependence, an estimate of which is possible using Debye

theory.'* The isotropically averaged mean square displacement values refined

on oxygens sites from the PDF, were plotted as a function of temperature

for the cases of no distortion and for an Oz distortion of 0.1 A. Using the

Debye expression for < u-^ >, the lOK data point was fitted to determine

a Debye temperature 0d and the expected temperature dependence was

plotted. The behaviour of the <u^> values refined with the Oz distortion

accord much better with the theory than do those refined for no distortion.

A static lattice energy minimisation calculation has been carried out

to investigate the nature and extent of atomic reorganisations around the

strontium ions which may be the cause of the observed disorder. An

ionic model was used with two-body central-force pair potentials of the

Buckingham form with the shell model of Dick and Overhauser^ used to

describe polarizability. The potential parameters were transferred from

CuO, CaO and SrO. They predicted well the phonon density of states of

the real material and thus appear to provide a good approximation to the

real potentials. Strontium ions were then introduced into a nominally pure

CaCu02 compound which took the layered structure of Ca.85Sr.i5Cu02.

The results indicated that the strontium ion could be accommodated with

minimal atomic displacements around it. Nearest neighbour copper and

oxygen ions were seen to move by no more than 0.06A and atoms beyond

this first shell moved by an order of magnitude less. These deviations are much less in extent than those observed in the PDF.

A perfect lattice energy minimisation calculation was carried out on an

idealised CaCu02 structiire. It was found that this had a lattice energy

minimum, E falling from -74.253 to -74.430eV/formula unit, when oxygen

displaced up and down from the plane, along z by O.SA, in a correlated

fashion from site to site. Thus the calculation predicts the Oz distortion,

though it overestimates the effect. DISCUSSION AND CONCLUSIONS

A driving force for the Oz distortion must also explain the lattice

calculations which predicted it within an ionic approximation. Such a driving

force may derive from the fact that the Ca-0 bond length in the material

is longer than the ideal CaO bond length as determined from ionic radii

considerations. The ideal Ca-O distance is calculated to be 2.47A^and the

distance in the crystal is 2.5lA. If the shape of the crystal potential is

341
convex in this region then a spUtting of the bonds to become equally longer

and shorter would produce an overall saving in energy since the shorter

bonds would gain a greater saving in energy than the longer bonds would

lose. For a value of A = 0.1 A the short Ca-0 bond becomes 2.44A and

the long bond becomes 2.57A, placing the shorter bond close to the ideal

Ca-0 ionic separation. This can also be thought of in terms of bondvalence theory.'For Ca^+ in Ca.85Sr.i5Cu02 in ideal 8-fold coordination

the average bond valence will be 0.25 valence units. The non-linearity of

the relationship between bond valence and bond length has been used to

predict the occurrence of distortions in situations where ions cannot take up

their ideal bond length, since an equal lengthening and shortening of bonds

leads to an increase in the average bond valence. In the present case the

undistorted bond length R = 2.5lA gives s = 0.230 valence units, somewhat

lower than the expected value. If the oxygens displace up and down by O.SA

then there are four bonds of length R = 2. 33A, s = 0.341, and four bonds

of R = 2.7lA, s = 0.151 valence units, thus the average bond valence is raised, simply by introducing the distortion, to 0.246 valence units. The real

distortion is closer to A = 0.1 A, however this still raises the average bond

valence towards the ideal value.

The properties of the plane, if not the detailed structure, will be similar

in all of the Cu02 planes in the superconducting cuprate materials. These

studies indicate that there is an ionic driving force tending to shift oxygen

off its average position in the plane. In the more complex superconducting

structures it is likely that other structural elements in the intergrowth regions

between the planes will dictate the exact configuration taken up by the

planes. It appears, however, that there is no strong restoring force holding

the oxygen in the plane and thus both the static and dynamic nature of

the planes are able to be significantly affected even by small second order

interactions.

REFERENCES

[1] T. Siegrist et. al.. Nature 334, 231 (1988).

[2] T. Egami et. al.. Materials Transactions, JIM, 31, 163 (1990).

[3] J. D. Jorgensen, et. al, J. Appl. Cryst. 22, 321 (1988).

[4] P. P. Debye, Ann. Phys. Lpz. 43, 49 (1914).

[5] see "Computer Simulation in SoUds" (C. R. A. Catlow, W. C. Mackrodt,

eds.). Springer-Verlag, (1982).

[6] R. D. Shannon and C. T. Prewitt, Acta. Cryst. ,925, 946 (1969).

[7] I. D. Brown in "Structure and Bonding in Crystals, Vol IF (M. O'Keefe

and A. Navrotsky, eds.) Academic Press (1981).

342
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

CRYSTAL CHEMISTRY AND OXYGEN ACTIVITY EFFECTS ON

T, T AND T* PHASE STABILITIES IN (La,RE)2Cu04 SYSTEMS

JOSEPH F. BRINGLEY, BRUCE A. SCOTT AND STEVEN S. TRAIL

IBM Research Division, Thomas J. Watson Research Center Yorktown

Heights, NY 10598

ABSTRACT

La2_xRExCu04 (RE= Nd-Y) solid solution systems have been

investigated to determine the factors stabilizing the T (La2Cu04), T'

(Nd2Cu04) and hybrid T*-type structures. A simple ionic model with a

perovskite-like tolerance factor is found to accurately define the existence

field of each. A remarkably sharp boundary is found to separate the T

and T' stability fields, while the T* structure occurs in a narrow region

adjacent to this boundary. The formation of the T* phase is viewed as re- sulting from a thermodynamic competition between the T and T' phases,

due to a tendency toward T/T' site ordering. Oxygen activity is found to

have a strong influence on phase selection in (RE)^Cu04 systems to the

extent that two different structure-types can be stabilized at a fixed cation

composition, but slightly different oxygen stoichiometrics. It is found that near the phase boundaries in La2.xREj^Cu04, the T and T* structure-types

are stable for S > 0, while the T structure is stable for ^ < 0.

INTRODUCTION

There are three closely related phases with stoichiometry RE2CUO4

(whose structures are shown in Fig. 1), all of which can be made superconducting under the appropriate doping and/or processing conditions

(1-3). The cation positions in these structures are to a first order approximation identical, so that the only difference between the three structuretypes is in the arrangement of oxygen about the cation framework. The

different oxygen arrangements determine the Cu-O coordination framework which varies from either 4-fold (T'-structure), 5-fold (T*-structure) or 6-fold (T-structure) Cu-O coordination. The RE-O framework changes

from nine-fold coordinate, rock-salt-like layers in T, to 8-fold, fiuorite-like

layers in T' or a mixture of the two (T*). Most striking, however, is the

varying substitutional preferences of these materials and the manifestation

of these preferences in their physical properties. The T-structure (Fig. la)

343
becomes superconducting when hole-doped, and furthermore, can only be

doped with holes (and not with electrons). In contrast, The T'-structure

(Fig. lb) becomes superconducting when electron-doped, and can only be

doped with electrons. The T* phase, which can be considered a hybrid of

the T and T' phases, might then be expected to exhibit both hole and

electron superconductivity, but has thus far been found to exhibit only

hole superconductivity (4,5).

Lo2Cu04(T) Nd2Cu04(T') T*

Fig. 1. The T, T' and T* structures, Cu = solid circles,

O = open circles and RE = shaded circles.

In an effort to understand their crystal chemistry and phase stability, and to gain insight into the nature of the specific substitutional pref- erences of the T, T' and T* phases, solid solutions and phase relationships

in the systems La2.xR-EjjCu04 (RE = Nd-Y) have been investigated.

These mixed rare earth systems allow us to evaluate how ionic size and

ordering affect structure and stability. The results are explained within the

framework of a simple ionic model in which a perovskite-like tolerance

factor (t) is found to be a remarkable predictor of T, T' and T* stability

limits (6). In addition, we have examined the effect of oxygen activity

(PO2) on solid solutions of the type La2.xNd^Cu04 and La2.^Dy CUO4 (Ce

and Th-doped). We have found that subtle changes occur in the oxygen

sublattice depending upon temperature and oxygen partial pressure (Po^)

during processing. These observations allow the stabilization of two different structure types (either T/T' or T'/T*) at a fixed cation composition,

and allow the stability fields of the T, T' and T* phases to be manipulated

considerably (7).

344
EXPERIMENTAL

Powdered mixtures of rare-earth and copper oxides, or coprecipitated

hydroxides, were heated in Pt crucibles for 24h, ground, pelletized, and

reheated, at temperatures from 900-1 150°C (48-72h). They were cooled

slowly or quenched from reaction temperature. Experiments were performed in air, oxygen, or in a high-pressure oxygen environment using an

externally heated Rene alloy vessel. Powder X-ray diffraction data were

obtained on a Siemens D500 diffractometer using CuKa radiation. Full

experimental details are provided elsewhere (6).

RESULTS & DISCUSSION

Extensive study of solid solution formation was performed on several

La, RE CUO4 systems, with particular attention paid to La2xNd CUO4.

ougn there is some disagreement in the literature (6-9), we find that

the equilibrium phase diagram of La^.^Ndj^Cu04 consists of a binary system containing the T-structure for 0.0 < x < 0.4, and the T' structure for

0.5 < X < 2.0, separated by a narrow two-phase region. In addition, a non-equilibrium T* phase near x = 1 and a fourth phase denoted T" near X = 0.5 have also been reported (6,9). For all the rare earths examined

(6), however, it is found that the T-structure of La2Cu04 becomes unstable

and a transition occurs to either a single-phase T'-structure or a two phase

T/T' mixture as the RE radius decreases from La-^Yb. In fact, if the

limiting solid solution concentrations of RE in the T-structure of

La2.xRE^^Cu04 from this work and from other literature sources are compared (6), it is found that the transition occurs sharply at an average

cation radius of 1.204(4) A. Below this value, the T'-structure becomes

stable for the stoichiometry RE2CUO4. These trends can be understood

within the framework of the following simple ionic packing model.

A complete summary of the results of the present work are collected in

Fig. 2, where solid solution compositions are plotted vs. a tolerance factor

(t), calculated in all cases for a "reference state" of 9-fold RE-O coordination. The tolerance factor, defined as t = (r* + Vq)I-J2{t^ + Tq), has

long been used to describe the stability of the related perovskites, and is a measure of the periodic match of the RE-O and metal oxide layers which

comprise these structures (6). It is found that T, T' and T* phases fall into

separate regions having quite discrete boundaries or transition points. The

T -^T' transition takes place sharply at a tolerance factor t = 0.865. This

defines the lower stability limit of the K2NiF4 (T-type) structure. Further,

if all of the known metallates having the T-structure are tabulated, it is found that the structure is stable over the range 0.865 < t < 0.993 (6).

However, for clarity, only the regions where the cuprates occur are emphasized here. The T-structure orthorhombically distorts at the lower

boundary of its range, 0.873 ^ t ^ 0.865 (dashed line of Fig. 2), before

undergoing transformation to the T'-form, which is stable over the interval

345
0.865 > t > 0.832. Combinations of copper, oxygen and rare earth cations

with t < 0.832 do not form RE2CUO4 compounds. RE2CU2O5 phases, in

which the rare earth is 6-fold coordinated, are the only known cuprates to

form in these mixed oxide systems (11), although the T' phase has been

stabilized at high pressures for the smallest rare earth cations (12).

L02CUO4

0.82 0.84 0.86 0.88

TOLERANCE FACTOR.

t

Fig. 2. Compositions of discrete phases and solid solutions

having the T, T' and T* structures. Boundaries are based on

the present results; for clarity, not all points are represented.

The data of Fig. 2 show the hybrid T*-phase falling in a very narrow

region near the T/T" phase boundary. This is not surprising in view of its close relationship to both the T and T' structure-types. We have investigated the systems La2.;^REjjCu04 (RE = Sm - Yb, Y) near x = 1, to ex- amine how rare-earth size and cation ordering govern the formation of the

T* phase. It was found that metastable T* phases form for RE = Sm - Gd but they could not be isolated as single phases. However, T* does becomes quite stable for the smaller rare earths Tb (13) and Dy and can be

isolated as single phases for 0.6 :^ x < 1 .0 and 0.8 < x < 1 .0, respectively.

We have recently shown that the T* phase occurs only when there is an

appreciable disparity in size between La and the smaller RE cation with

346
which it is matched (6), and that this provides a driving force for A-sitc

cation ordering.

Oxygen Activity Effects

All three phases in the RE2CUO4 system are known to exhibit nonstoichiometry in the anion sublattice. We have examined the effect of ox- ygen activity on solid-solutions of the type La2.xNdj^Cu04 and

La2.xDy^Cu04 (Ce and Th doped) to determine its effect upon the crystal

chemistry and phase stability in RE2CUO4. We find that subtle

stoichiometry changes occur in the oxygen sublattice depending upon oxygen activity and temperature during processing. These small variations in

oxygen content influence the precise location of the T/T' or T'/T* phase

boundaries in these systems, thereby stabilizing two different structuretypes at a A-site cation content.

The solid solution limit of Nd in the T-phase of La2.xNdjjCu04 in air

at 1050°C occurs at x = 0.4. However, if the samples are prepared at

910°C in 400 bar O2, the solid solution limit extends to x = 0.7. Figure 2

shows the X-ray powder diffraction patterns of (a) La, 4Nd 6Cu04 + ^ prepared in 400 bar O2 at 910°C, and (b) the same sample heated in air and

quenched from 1050°C. The pattern of the sample prepared in 400 bar

62 is that of a single phase material having the T structure, while the

phase prepared in air is T'. The oxygen stoichiometrics of the two samples

differ by only 6 = 0,06. Thermogravimetric analysis of solid-solution

compositions close to the phase boundary indicate that, in general, the

T-phase is stable for ^ > 0 and T' for ^ < 0, with the T T' transition

occuring for S x 0. The transitions are also found to be reversible (7).

A similar result was observed in solutions of the type

Laj 2^y.s-x^x^^^4 + 6 ^ ^h)' where unique cation compositions

could be prepared in either the T' or T* structure-type depending upon the

oxygen activity. For example, a sample of composition

Laj 2Dy 7Ce ,Cu04^.^ was found to exhibit the T* structure when prepared in 1 atm O2, but was converted to the T' structure when heated at

950°C in pure N2.

The oxygen activity effects observed in our experiments occur for compositions adjacent to the transition boundaries determined in air. Since

two phases will have similar free energies near their phase boundary, it is perhaps not surprising that the resulting small shifts in oxygen

stoichiometry are so effective in steering phase selection. A

substoichiometric oxygen content is destabilizing to the T structure with its higher cation-oxygen coordination numbers, while this favors the T'

structure. A similar argument can be made for T* vs T'. These effects are also consistent with the observed doping preferences of the T, T' and

T* phases and may be be related to electronic factors.

347
T

1 1 1 1 1 1 1 -r

in 'c

s

xi

wO

(a) La|4Ndo.6Cu04 04

Mi, Afc k1 1

—1 1 1 1 1 1— 1

' "

(b) La|.4Ndo.6Cu03 98

>». Jl. 1. |1 . ll :

z

UJ

Iz

20 60

2(9)

Fig. 3. X-ray powder diffraction profiles of (a) the T and

(b) the T' phase at x = 0.6. The oxygen stoichiometrics are given.

SUMMARY AND CONCLUSIONS

The following picture emerges from our study of La2.^RE^Cu04

systems. As the value of x increases, the average rare earth ionic radius

declines and the 9-fold coordinated site of the T structure becomes unstable. Structural changes are first manifested by the transition in T-phase

symmetry from tetragonal to orthorhombic. With further increase in x, the

average rare earth cation radius becomes too small to support the

348
orthorhombically-distorted structure, and a major transformation takes

place to T'. Near this boundary both T and T' are of comparable

thermodynamic stabiHty, as is T*, so this hybrid structure also makes an

appearance. At first, T* is metastable, but it begins to compete favorably

with T' when the mismatch in RE size becomes sufficient to provide a driving force for cation ordering. We estimate that approximately a 10%

difference in average rare earth radius between the 9- and 8-fold coordinated sites is required to fully stabilize T*. At larger values of x (or for the

smaller rare earths) this difference is exceeded and T' becomes the stable

phase. Eventually, the average RE radius becomes too small to support

8-fold coordination, lowering the tolerance factor further, and RE^CuOd

phases are no longer found at ambient pressure. The stability fields, ana

therefore the phase boundaries, can be manipulated considerably by careful control of temperature and oxygen activity, thereby allowing two different structure types to occur at a fixed cation composition, albeit a

slightly different oxygen stoichiometry. Near the phase boundaries in

La2.,RE^Cu04, the T and T* structure-types are stable for ^ > 0, while

the T' structure is stable for ^ < 0. These results, in combination with the

tolerance factor model, can be very useful in the design of selective synthetic routes to the T, T' and T* phases.

REFERENCES

1. H. Miiller-Buschbaum, Angew. Chem. Int. Ed. ]6, 674 (1977).

2. Y. Tokura, H. Takagi, and S. Uchida, Nature 337, 345 (1989).

3. A. C. W. P. James, S. M. Zahurak and D. W. Murphy, Nature 338,

240 (1989).

4. H. Sawa et al. Nature 337, 347 (1989).

5. M. F. Hundley, J. D. Thompson, S.-W. Cheong, Z. Fisk, R. Scharwz

and J. E. Schirber, Phys. Rev. B. B40, 5251 (1989).

6. J. F. Bringley, S. S. Trail and B. A. Scott, J. Solid State Chem. 86, 310

(1990).

7. J. F. Bringley, S.S. Trail and B.A. Scott, J. Solid State Chem. in-press

(1990).

8. K. K. Singh, P. Ganguly and C. N. R. Rao, Mat. Res. Bull. ]7, 493

(1982).

9. A. Manthiram and J.B. Goodenough. J. Solid State Chem. in-press

(1990).

10. R. Shannon, Acta Cryst. A32, 751 (1976).

11. N. Kimizuka, E. Takayama and S. Horiuchi, J. Solid State Chem. 42,

322 (1982).

12. H. Okada, M. Takano and Y. Takeda, Physica C 166, 111 (1990).

13. T. Kenjo and S. Yajima, Bull. Chem Soc. Jpn. 50, 2847 (1977).

349
DISCUSSION

B. Dabrowski: Have you seen other members of this family, namely the RuddlesonPopper phases, in high pressure samples?

J. Bringley : No, we didn't. We tried to prepare them directly, but were never able to.

B. Dabrowski: But they have been reported in the literature?

J. Bringley : Yes, there are two reports. Frankly, I don't believe that they exist although

they may be present as intergrowths in some samples.

R. Cava: But they're from I.B.M.

J. Bringley : I think they observed them by electron microscopy; correct? So they appear

to exist as intergrowths.

H. Zandbergen : Have you any idea about the reason why the Demazean phase of

LaCu03 is rhombohedral, because it's prepared at the same temperature?

J. Bringley : Well, it appears to be the high pressure polymorph of LaCu03

prepared at 65 kilobar, versus a hundred fold less pressure in this case. So I think that

would perhaps explain why it has different symmetry. Actually if you look at the two

structures they are almost identical, the tetragonal and rhombohedral distortions are very

small. Both are quite close to cubic symmetry.

H. Zandbergen: Is there a difference in volume?

J. Bringley: Yes, the volume of the tetragonal form is about five percent larger as you

would expect for a low pressure polymorph.

A. Navrotsky: You mentioned two phase regions. Do you definitely see two phases

coexisting in the x-ray patterns?

J. Bringley : Yes, we do. And let me just comment that I was a little bit hesitant to show

that slide because it is not really a phase diagram at all. The way we came up with that

diagram is that we determined lattice parameters from a collection of samples, which were prepared under different conditions, so it really is only to show the evolution of the

lattice parameters and the approximate homogeneity ranges of these phases. But, in fact, one can, by playing tricks at lower temperatures prepare, for example, the tetragonal

phase with a much lower oxygen stoichiometry. So those two-phase regions, yes, we did

see both phases growing in under these conditions.

350
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

Janxiary 1991.

ROGM TEMPESA1URE MTTRTT.TTy and DIFFUSION OOEFFICIE24T of

CSflfGEN in PaLYtVYbTALLINE YBa2Cu30y_^

Yosef Soolnik , Michael Schwart:^, Eyal Sabatani,

and David Cahen

the Vfeizmann Institute of Science, Rehovot , Israel 76100

KESTRMJF. The practical diffusion coefficient of oxygen in

room tenperature, as measured by a current decay

method in a wet electrochemical set-up was found to be 4 ±

2x10" cm /sec, essentially independent of grain size (in

polycrysLalline thin films or bulk pellet) and of surface

prc^jerties. Therefore it is concluded that the rate-determining

step in the diffusion process, the movement of the oxygen inside

the grain, is much faster than vihat would be expected on the

basis of extrapolation of hig^i tenperature data. This explains

our ability to use RT reduction to form metastable 123 variants.

jnh^oduction

Oxygen-deficient YBa2Cu30y_jj^( 123 ) has been prepared by several

methods, mostly by reoucing at, or quenching frcm high

tenperatures. Such hig^ tenperature methods of ( de )Qxygenation

change not only content but also the arrangement of oxygens (1).

This makes it difficult to distinguish between structural effects

and electronic ones (change of number of charge carriers due to

change in oxygen content) on the superconducting prcperties of

123. Earlier we have shown (2,3) that 123 can be reduced at RT

in a wet electrochemical cell and that this provides a way to

s^)arate the electronic and structural effects. This can lead to

unique behaviour including material showing a 20"'K T^ (3)

[recently confirmed by the elegant experiments of Jorgensen et

al . ( 4 ) ] . In view of this possibility of RT reduction the

question of oxygen mobility in this, relatively low, tenperature

regime becomes inportant.

EXPERIMEmTlL

Material Preparation: 123 was prepared using in situ

prepared nitrate precursors, as described in refs. 2 and 3. Xray powder diffraction showed the orthorhcmbic pattern of 123.

The material showed reproducibly an oxygen content of 6.97 ± 0.02

as determined by iodcmetric titration. Lag ^Stq ^CcO^^^ was

351
prepared accxDrding to ref . 5 and was single phase as determined

by X-ray pcwder diffraction. The parameter 'x' was determined to

be 0.01 by iodcmetric titration. The thin films used were

prepared at Tel Aviv University by e~- gun sputtering of 123

targets, using alumina (a-corundum) substrates. The films were

non-epitaxial, polycrystalline and 1pm thick. In seme of the

experiments special treatments were done on the pellets before

electrochemical reduction. The density and average grain size of

pellets could be changed by varying the heat treatment during the

final sintering stage. The effect of porosity was examined using

pellets v*iich were reground and repressed without sintering. Porosity was further increased by preparing pellets v^ch

contained tetra-butyl ammcnium perchlorate (tBAP) vrtiich was

removed by soaking. Porosity was decreased artificicilly by

covering the working side of the pellet with a micrcporous

membrane. Porosity was estimated ir/ SEM. The surfaces of seme

pellets were modified by etching with a Br2/ethanol solution (1%

by wei^t) for 10 to 40 minutes, to remove insulating hydroxide

and carbonate formed by air exposure (6). Also, Ag was

evaporated onto the pellets because of the reported decrease in

activation energy for oxygen out-diffusion from sanples coated

with Ag (7).

The wet electrochemical reductions were performed in a standard

3-electrode cell. Generally, propylene carbonate/ (O.IM tBAP) was

used as the electrolyte. In seme cases we used dry acetonitrile

as the solvent. Pressed pellets were 8 mm in diameter and 0.5 mm

thick, with a density of 4.6-5.0 g/cxtr. Contacts to all sanples

were prepared by attaching a Cu wire to the sanple with Ag paint

and insulating epoxy cement. Two Pt wires were used as counter

and quasi-reference electrodes. Ar was bubbled throu^ the

solution before the reduction. During the esqjeriment a stream of

Ar, blanketing the solution surface, was maintained. Standard

electrochemical equipment including a coulcmeter, was used.

Magnetic susceptibility measurements were done using the a.c. mutual inductance technique at a frequency of 75 Hz and a field

of approximately 1/3 Gauss.

We performed cyclic voltammogramms of pellets treated with

ethanol/Br2 and inpedance measurements to determine the porous

diffusion coefficient of the pellets. The experiments were done

in a 3-electrode system with a platinum counter electrode, a Ag/O.IM Ag"*" reference electrode and a pellet as the vrorking

electrode. For this e:q)eriment we used an acetonitrile solution

that contained O.SrrM tetramethyl ammonium methyl

ferrocene( TmamFe*'^ )

.

352
Method of analysis The diffusion cxDefficient of oxygen in 123

was determined using the method of Van Buren et al. (5) for

polycrystalline sanples, in wtiich the time decay of the current

is measured. This method can be used for the determination of

oxygen diffusion coefficients in porous pellets. It is based on

a bounded 3-dimensional diffusion model under potentiostatic

conditions. A plot of I vs. l//t was made and from the intercept

on the l//t axis (=> t^) D could be determined from the following

equation:

X = a//(Dto) (1)

v*iere a is the average grain size (determined by SEM) and X is a

shape parameter taken to be 2 as intermediate between a sphere

and a cube. In order to test the applicability of the method to

our set-up and the organic solutions, the esqjeriraent was also

performed with LaiQ ^Srg 5C0O3 „ in aqueous KM as reported in

ref.(5). The diffusion coefficients that we obtained were

similar to those reported in ref. 5. Inpedance measurements were

in the form of ccnplex inpedance plots and were analyzed by the

method of ref. 8.

RESULTS

EVIDENCE FOR ELECTROCHENICAL REDUCTION OT 123: There are several

indications that the currents measured in these reduction

esq^eriments are indeed due to oxygen reduction, and are not the

result of seme corrosion or surface catalysis process.

Stnxrtural evidence: The X-ray diffraction patterns show

shifts, especially in the region of the (200) and [(020) + (006)]

reflections, that are similar to those reported for high

tenperature reduction of 123, in the case of thin films. For the

pellets there are some differences and the reflections are

broadened ijpon reduction, as corpared to those for non-reduced

ones (cf. ref. 3). In the thin films this broadening did not

occur. Since the X-ray patterns for pellets are measured on

reground pellets, we concluded that we measure a phenomenon

represQitative of the bulk of the pellets.

- Stochianetric evidence: A second result vdrLch SL^ports oxygen

removal (in the form of 0" ) from the bulk of the pellet is the

exact agreement between the quantity of oxygen removed as determined by coulometry and the final oxygen content of the

sanple as determined by iodcmetric titration. This analysis

could be done only on pellets. When the esqjeriment was done in

the same way with a pellet of YBa^Cu30g (1236) essentially no

current could be passed, again indicating that the reaction in

123(7) is not a surface phenomenon.

353
- Atomic Absorpticn ^lectrosoopy was used to check for cc^per in

the electrolyte solutions after reduction. Because the amount of

oc^jper that we found could account for less than 0.2% of the

charge that was passed, corrosion of the material cannot be

responsible for the current during the experiment.

- Magnetic susceptibility of reduced sanples: The transition to

the superconducting state is broadened in reduced pellets

relative to starting material. The degree of broadening

increases with increasing degree of reduction (for a full

description of these and electrical resistivity results, see

refs. 2 and 3).

- Impedance measurements: The frequency plots that we obtained

are very close to those expected fran theory (8). The calculated

porous diffusion coefficient is of the order of 10" -10"

cm /sec, vAiich shews that porous diffusion is NOT limiting.

DIFFUSION OOEFFICIENT : The practical diffusion coefficient,

determined by potentiostatic current decay, was found to be:

4 ± 2 X lO"-'-^ cm^/sec

for pellets and thin films. In order to test vdiether the

measured effective diffusion coefficient was actually due to

oxygen diffusion within a grain or due to seme surface or

solution process, the preparation conditions were varied (as

described in the Experimental Section) to obtain pellets with

modified jiiysical prcperties. The grain size was varied from

5-10 ium (as shown by SEM) and pellets with densities of 4.2-5.0

g/cm were used. The changes in grain size and density resulted

in pellets with different porosities as observed by SEM.

Porosity was also varied with the use of a micrcporous membrane

on the working surface. In addition, the working voltage was

varied from -1.0 to -0.5 V in various experiments. Pellets were

also used in v*iich the surfaces had been modified by an etching

process and by Ag evaporation as described above. Cyclic

voltammograms clearly showed an influence of the etching process

on the surface prcperties of the pellets (as exgressed by

improved electrode behaviour). In all these experiments the

measured diffusion coefficient of 123 at rocm torperature was

found to be independent of the different treatments. Sanples

that were partially reduced by hi^ tenperature methods, to

oxygen contents similar to those obtained by RT reduction, gave

similar results, within the scatter of the euqjerimental data.

The thin film results were similar to those obtained for the

pellets. Their analysis was done assuming a grain size of Ijm.

The polycrystalline nature of these thin filins is v*iat allowed us

to use this current decay method for them.

354
DISCUSSION

It was alreac3y found that the diffusion coefficient for oxygen in

this material is very hi^. For exanple a value of 5x10"

cair/sec was reported at 550 "C (9). ExtrapDlation of hic^

tenperature results to roan tenperature gives a value for D of

approximately 10" an /s. Such a small diffusion coefficient

would limit oxygen loss to a very narrow surface region and

cannot esqDlain our results. As stated before our results give a

much hi^ier practical diffusion coefficient. Th^e results

should be ccnpared to the values of: 1.4 x 10" cm /s and 5 x

10" an /s at 25 °C for D in other perovskites, namely

Lag ^Stq 50o03_„ and Ndg ^SrQ ^OcO^-^ (5). These values are also

orders oi magnitude larger tl^an tliose obtained by extrapolation

from high tenperatures. We note that Yoshida et al. (10), using

a plasma-oxidation technique, found a diffusion coefficient of

the order of lO"-"-^ anVsec for 123 films below 160 °C.

OGNCLUSIONS

The effective diffusion coefficient of oxygen in polycrystalline

123 at roan tenperature, as found by the current decay method,

was found to be independent of the sanple form (polycrystalline

thin films or bulk pellets) and of the various treatments given

to the pellets. In particular the changes in porosity, density

and surface treatments do not affect the measured effective

diffusion coefficient. These results lead to the conclusion that

the rate-determining step in the diffusion process is the

diffusion of oxygen inside the grain, rather than a surface

reaction, and that at roan torperature this diffusion is

unexpectedly facile. This explains the ability to modify the Ocontent at room tenperature, a process that allows the formation

of metastable variants of 123.

Acknowledgements : We thank the US-Israel Binational Science

Foundation, Jerusalem, Israel, for financial support. We are

grateful to G. Deutscher, U. Dai and M. Has fron Tel Aviv

Iftiiversity for providing us with thin films.

# Present adress: Xsirius, Hamada 4, Jerusalem , Israel

* Authors for correspondence

355
REETRENCES

1. Manthiram, A.; Swinnea, J.S.; Sui, Z. T.; Steinfink, H.;

Goodenough, J. B.; J. Am. Cheta , Soc . 1987, 109 , 6667.

2. Schwartz, M. ; Ra^paport, M. ; Modes, G.; Cahen, D.; Physica C

1988, 153-155, 1457.

Schwartz, M. ; Cahen, D.; Rappaport, M. ; Hodes, G. ; Solid

State Ionics 1989, 32-33, 1137.

3. Schwartz, M. ; Rappaport, M. ; Hodes, G.; Reich, S.; Cahen,

D.; Mat. Lett. 1989, 7, 411.

4. Jcongensen, J. D.,* Pei, S.; Li^tfoot, P; Shi, H.; Paulikas,

A. P.; Veal, B. W. ; Physica C

1990 167,571.

5. van Buren, F. R.; Broers, G. H. J.; Bouman, A. J.; Boesveld,

C; J. Electroanal . Chem. 1978, 88, 353.

6. Vasquez, R. P.; Hunt, B. D,; Foote, M. C; Pippl . Phys . Lett.

1988, 53, 2692.

7. Schrott, A. G.; Tu, K. N.; Yeh, N. C; Singco, G. ; Levi, A.;

Tsuei, C. C; Phys . Rev. B 1989, 39 , 2910.

8„ De Levie, R.; Electrochimica Acta 1963, 8 ,751.

9. Ginley, D. S.; Nigrey, P. J.; Venturini, E. L.; Morosin, B.;

Rwak, J. F.; J. Mater. Res . 1987, 2, 732.

10. Yoshida, A.; Tamura, H.; Morciiashi, S.; Hasuo, S.; P^ppl . Phys . Lett. 1988, 53, 811.

356
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

PHASE SEPARATION IN Nd2 -xCexCu04

B. Dabrowski,. P. Lightfoot, D.R. Richards, D.G. Hinks,

J.D. Jorgensen, Shiyou Pei, Y. Zheng, D.T. Marx and A.W.

Mitchell, Materials Science Division, Argonne National

Laboratory, Argonne, IL 60439, USA

ABSTRACT

A series of Nd2 - x*-^® x^^O 4 samples with starting

compositions 0<x<0.25 have been prepared in several

different oxygen partial pressures at 990-1100°C. The

high resolution neutron and synchrotron X-ray

diffraction data show that single-phase material is

formed only for the x=0 . 0 and x~0.165 compositions. For

the x=0.165 sample there is a strong dependence of

cerium solubility on synthesis temperature.

INTRODUCTION

The Nd2 - x^ ® x*^^ O 4 system is unique among the

superconducting copper oxides because the charge carries

are considered to be electrons instead of holes [1]. The

assignment of electrons as charge carriers was based on

structural cons iderat ions [ 2 ] , chemical analysis [3] , X-ray absorption spectroscopy [ 4 ] and Hall and Seebeck

coefficient measurement s [ 1 , 5 ] . However, several

experimental facts remain ambiguous. For example, some

authors reported conflicting results for the sign of the

Hall coef f icient [ 6 ] and for the total oxygen content of

the samples [7]. The presence of two or even more phases

in the same sample were reported in electron diffraction

studies [8]. Moreover, reducing synthesis conditions (as

used for the synthesis of superconducting Nd2-xCexCu04

and try to dope the system with electrons via oxygen

vacancy generation) are now frequently used for

preparation of hole doped superconducting systems like

Sr-Ca-Y-Pb-Cu-0 [ 9 ] and to improve the superconducting

properties in the Bi-Sr-Ca-Cu-0 system[10].

We have recently shown [11] that consideration of

phase separation is required before the superconducting

properties of Nd2-xCexCu04 can be uniquely determined.

Our high resolution neutron and synchrotron X-ray

diffraction data show that single phase material is

357
formed only for the x=0 . 0 and x~0.165 compositions.

Other samples contain a mixture of the x~0.165

composition, which is superconducting, and other

compositions which are nonsuperconduct ing . In this

article we review our previous results and provide some

additional data for the temperature dependence of the

observed phase separation for samples with fixed

starting cerium content (x=0.165).

EXPERIMENTAL

Powder samples of Nd2-xCexCu04 with 0<x<0.25 were

prepared from the metal oxides in several different

oxygen partial pressures in the temperatures range

990-1100°C, followed by quenching. Annealing in high

purity argon at 900°C followed by quenching produces

superconducting material. Neutron powder diffraction

data were obtained using the Special Environment Powder

Dif f ractometer (SEPD) at Argonne National Laboratory's

Intense Pulsed Neutron Source. High resolution X-ray

data were collected on the powder dif f ractometer on line

X7A at Brookhaven's National Light Source. The

superconducting properties were studied by AC

susceptibility using a Lake Shore Cryotronics, Inc.

Model 7000 AC svisceptometer

.

RESULTS AND DISCUSSION

Our synthesis conditions for samples with different

cerium content were typical (fired in air at 1080°C and

annealed in argon at 900°C) to that used in previous

studies [ 1-3 ] . We found that firing in air at higher

temperatures frequently led to partial melting,

especially for low cerium content samples. Our results

for T^;, and superconducting phase fraction, shown in Fig.

la, are also similar to the previously reported data [12]

and show bulk superconductivity occurring only in a narrow range of composition (Ax-O.Ol) centered at

x=0.16. The superconducting fraction was determined from

the zero field cooled AC susceptibility at 4 . 5K and is

normalized to the largest fraction occurring for x=0.16

sample (which is about 0.6 of the full diamagnetic

signal) . The superconducting T^, determined from the

onset of the diamagnetic signal, appears to remain

almost constant for x<0 . 1 6 at 24K, then decreasing by

only a few degrees at larger cerium concentrations.

358
25

20

1 5 o H 1 0

5

0

3.10

' ' 1

' (a)

O

o

cnco

r3c

oo

0)Q.

T

""03

0.00 1

I I I

'

I 'I 'I I I I 1

t I

t 0.00 0.05 0.10 0.15 0.20 0.25

X

Fig. 1. Superconducting T^-. and volume fraction (a),

lattice parameters ratio c/a (b) and volume phase

fraction (c) as a function of starting cerium content

for Nd2-xCexCu04

.

359
The neutron diffraction data, analyzed initially by

the Rietveld method based on a single-phase model, have

been used to refine the structure in the tetragonal

l4/mmm space group as previously reported [13] . However,

substantially better fits, which accounted for the

distinct shoulders on many of the Bragg diffraction

peaks, were obtained when a two-phase model was used for

the x-0.05, 0.1, 0.13 and 0.15 samples. The model

consisted of two tetragonal phases for which only the

lattice parameters and the phase fractions were allowed

to refine. For other samples the multiphase behavior was

not clearly evident in the diffraction data. The refined

lattiq^e parameters ratio c/a and the volume phase

fraction for each of the two phases are shown in Fig. lb

and Ic, respectively. The c/a ratio displays an almost

linear decrease with x for one phase while remaining

practically constant for the other phase. The volume

fraction of the phase with the constant c/a ratio peaks

sharply for x=0.16 and 0.17 samples coinciding with the

peak in the superconducting phase fraction (Fig. la) . This behavior suggests that only one composition,

x~0.165, is superconducting and traces of

superconductivity in other samples result from the

presence of a small amount of that phase.

The high-resolution X-ray data confirmed, in

principle, these initial findings. However, while X-ray

data show relatively sharp peaks, broad shoulders are

observed extending to higher values of x. This may be

interpreted as the presence of several phases extending

to the solubility limit

.

To understand the phase separation process better,

we have performed a structural study as a function of

firing temperature and oxygen partial pressure for

samples with fixed starting cerium content (x=0.165)

.

Figure 2 shows portion of the Rietveld refinement

profiles for a sample fired in air at 1000°C (a) and at

1080°C (b) followed by quenching. Firing at 1000°C leaves

some unreacted Ce(Nd)02 oxide leading to the lower final

cerium content in the compound. The sample also shows

phase separation in agreement with the previous

observations (shoulders are clearly visible on several

of the diffraction peaks) . All the cerium is

incorporated into the compound at 1080°C and the sample

is single phase (the two-phase refinement shown on Fig.

2b is only marginally better than single-phase

refinement)

.

Although the shape of the observed miscibility gap

is not yet fully determined, several interesting

360
Fig. 2. Portion of the Rietveld refinement profiles for

sample fired in air at 1000°C (a) and at 1080°C (b)

followed by quenching.

361
conclusions can be made. The miscibility gap appears to

have one limit around ~0. 05-0. 10 and the other limit at

0.15 for temperatures near 1000°C. The phase separation

is an intrinsic characteristic for this system, thus,

similar behavior must be present for samples studied in

other laboratories. Any conclusions regarding continuous

variation of the normal and superconducting state

properties as a function of cerium content must be

reconsidered. The presence of two phases in most samples

may be the cause of disagreement reported in the

literature because properties of the compounds on both

sides of the miscibility gap may be very different, and,

in particular, one compound may be electron doped while

the other hole doped.

CONCLUSION

We have examined a series of Nd2-xCej^Cu04 samples

with nominal compositions 0<x<0.25 using high resolution

neutron and synchrotron X-ray powder diffraction and AC

susceptibility. Bulk superconductivity, with onset

T(-;~24K, as determined by AC susceptibility, occurs only

in a narrow range of composition (Ax~0.01) centered at

x=0.16. The diffraction data show that samples of

Nd2-xCexCu04 are not single phase except near x=0 and

x~0.165. Two-phase Rietveld refinements indicate that

other starting compositions contain a fraction of the

superconducting composition in an amount proportional to

the observed diamagnetic fraction. The cerium solubility

is strongly dependent on the synthesis temperature. Our

results imply that interpretation of past results has to

be taken with caution and great care has to be taken

before the superconducting properties of Nd2-xCexCu04

can be uniquely determined.

*Supported by NSF-Office of Science and Technology

Centers under contract #STC-8809854 (BD, PL, SP, DTM)

and U.S. Department of Energy, BES, under contract

#W-31-109-ENG-38 (DGH, JDJ, YZ, AWM)

.

REFERENCES

1. H. Takagi, S. Uchida, and Y. Tokura, Phys . Rev. Lett.

62, 1197 (1989)

.

2. Y. Tokura, H. Takagi, and S. Uchida, Nature 337, 345-

(1989)

.

362
3. J.T. Market, E.A., Early, T. Bjonholm, S. Ghamaty,

B.W. Lee, J.J. Neumeier, R.D. Price, C.L. Seaman and

M.B. Maple, Physica C158, 178 (1989).

4. J.M. Tranquada, S.M. Heald, A.R. Moodenbaugh, G.

Liang and M. Croft, Nature 337, 720 (1989); E.E. Alp,

S.M. Mini, M. Ramanathan, B. Dabrowski, D.R. Richards

and D.G. Hinks, Phys Rev. B40, 2617 (1989).

5. Y. Hidaka and M. Suzuki, Nature 338, 635 (1989).

6. Z.Z. Wang, T.R. Chien, N.P. Ong, J.M. Taraskon and E.

Wang, (preprint)

.

7. E. Wang, J.-M. Tarascon, L.H. Greene, G.W. Hull and

W.R. McKinnon, Phys. Rev. B41, 6582 (1990).

8. F. Izumi, Y. Matsui, H. Takagi, S. Uchida, Y. Tokura,

and H. Asano, Physica C 158, 433 (1989); C. H. Chen,

D. J. Werder, A. C. W. P. James, D. W. Murphy, S.

Zahurak, R. M. Fleming, B, Batlogg, and L. F.

Schneemeyer, Physica C 160, 375 (1989); T. Williams,

Y. Maeno, I Mangelshots, A. Reller, and G. Bednorz,

Physica C 161, 331 (1989)

.

9. R.J. Cava at al . , Nature 336, 211 (1988).

10. S. Koyama, U. Endo and T. Kawai, J. Appl Phys. 27,

L1861 (1988) . 11. P. Lightfoot, D.R. Richards, B. Dabrowski, D.G.

Hinks, S. Pei, D.T. Marx, A.W. Michell, Y. Zheng and

J.D. Jorgensen, Physica C168, 627 (1990)

.

12. G. Liang, J. Chen, M. Croft, K V. Ramanu jachary , M.

Greenblatt, and M. Hedge, Phys. Rev. B 40, 2646

(1989)

.

13. H. K. Muller-Buschbaum and W. Wallschlager , Z. Anorg.

Allg. Chem. 414, 76 (1975)

.

363

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

NEW RARE-EARTH CUPRATES WITH THE NaCu02 STRUCTURE

P.K. Davies, E. Caignol, and T. King

Department of Materials Science and Engineering,

University of Pennsylvania, Philadelphia, PA 19104-6272.

INTRODUCTION

With the reports of n-type superconductivity in Ce and Th-doped Nd2Cu04,

the isolation of new Nd-containing cuprates is of considerable scientific interest.^

Our original aim in this work was to investigate the formation of double- Cu sheet

compounds with the T-type Nd2Cu04 structure. In the corresponding lanthanum

systems it is well known that a simple "Ruddlesden-Popper" T-type phase,

La2MCu206, in which Cu ions in square-pyramidal coordination form 2-

dimensional double sheets, is stabilized by the intra-sheet cation M, where M=Ca or

Sr.2 In this structure the pairs of Cu-sheets are separated by a single LaO rocksaltrelated layer and the 9-fold coordination of La is essentially identical to that in the

La2Cu04 T-type structure. Extensive p-doping of La2MCu206 is quite difficult,

however recent work by Cava et al has shown that by using high pressure

techniques superconducting transitions on the order of 60K can be produced in

samples with (Lai.6Sro.4)CaCu206.-^ The T' analog of this compound, in which

the double Cu-sheets - with Cu in square-planar coordination - would be separated

by fluorite-related intergrowths with Nd retaining the cubic coordination adopted in

Nd2Cu04, should also have the nominal stoichiometry Nd2MCu206. For M=Sr a

single phase compound does form, however it crystallizes with a T-type structure,

with Nd in 9-fold coordination, and is isostructural with the corresponding La

compound.

365
The apparent instability of a double-sheet T cuprate with the Nd2SrCu206

stoichiometry, is most likely due to an effective increase in the size of the ions on

the "Nd" site resulting from a distribution of Sr on both the intra-Cu sheet site and

the Nd site. To reduce the effective size of the ions in the Nd-site we synthesized a

series of compositions of Nd2MCu206 with M = (Sri-xCax), hoping that the

smaller Ca ions would favor the formation of a T'-type Nd2MCu206 phase.

Despite the apparent logic of this approach, no evidence was found for a doublesheet T cuprate phase. As the concentration of M approached pure calcium the x- ray patterns were multi-phase, and contained reflections from a new insulating

phase, with a non-perovskite-related structure, with the stoichiometry

Ca2±xNd2±xCu50io. Isostructural phases were also found in the corresponding

Y203-CaO-CuO and Gd203-CaO-CuO systems. The structures of these new

phases, which are closely related to that of NaCu02, are the subject of this paper.

STRUCTURAL STUDIES OF Ca2±xM2±xCu50io (M = Nd, Gd,Y)

All the specimens were prepared by standard solid state techniques at

1(X)0°C in air. The samples were investigated by powder x-ray diffraction using an automated Rigaku DMaxB diffractometer with CuKa radiation generated at 50kV

and 40mA, and by TEM using a Philips 400T electron microscope operated at 120kV. For M = Y, a single phase solid solution is formed for 0<x<0.8. X-ray

patterns collected from a series of samples within this range are shown in figure 1

.

With the exception of those indicated in the figure, all of the remaining peaks could

be indexed using an orthorhombic cell with, for x=0, a = 2.817(1)A, b =

6.185(1)A, c = 10.594(1)A and V = 184.63(7)A3. This cell, space group Fmmm,

is the same as that reported by Seigrist et al"^ for the sub-cell of the low temperature

phase Cao.8Cu02, which is isostructural with NaCu02. This structure, first

reported by Hoppe^' consists of one-dimensional chains of edge-sharing Cu square

planes, linked by sodium ions in octahedral coordination, see figure 2. In NaCu02

the octahedral sites are fully occupied; however, both Cao.8Cu02 and the ternary

Ca2±xM2±xCu50io cuprates are cation deficient, and the strong additional

366
25.0 30.0 35.0 4O0 450

29

50.0

Figure 1. X-ray patterns for Ca2±xY2±xCu50io with x = 0, 0.4, 0.75. Shaded

peaks due to incommensurate superstructures.

• Cu

O 0 # Na

Figure 2. Schematic of NaCu02 structure. Edge-shared Cu square planes form 1- D chains along a axis. These are linked by larger cations located in channels.

367
reflections in the x-ray patterns result from commensurate and incommensurate

ordering of the Ca/M cations in the channels of the NaCu02-type cell.

The positions of the reflections corresponding to the orthorhombic sub-cell

show httie variation as the composition of the solid solution is changed. However,

the positions of the extra "un-indexable", superlattice diffraction peaks exhibit a

large and systematic variation across the solid solution region - see figure 1. By

using fractional Miller indices based on an orthorhombic "NaCu02-type" sub-cell,

we have found that it is possible to index all the diffraction peaks in the powder

patterns of the Ca2±xM2±xCu50io phases. For example, the two strong reflections

close to 20 = 30°, can be indexed as (1 -5a, 1,1 -5c) and (l-5a,l,l+5c) respectively.

Using this type of indexing scheme we find that for Ca2Y2Cu50io, 5a = 0.201(1)

and 5c = 0.243(3); relative to the orthorhombic sub-cell, this corresponds to a

supercell with a = 5asub' c = 4.12csub- ^ more complete indexing scheme for

Ca2Y2Cu50io is shown below.

Ca2Y2Cu5Oi0: sub-cell; a=2.817, b = 6.185, c = 10.594 A

5a = 0.201, 5c = 0.243; a - 5asub» c = 4.12csub

2eobs 2ecalc I h k 1

16.65 16.72 6 0 0 2

28.85 28.84 13 0 2 0

29.83 29.83 24 l-5a 1 l-5c

31.03 31.00 20 l-5a 1 l+bc

33.53 33.52 100 0 2 2

33.81 33.81 30 0 0 4

35.99 36.01 23 1 1 1 37.53 37.52 18 l-5a 1 3-5c

40.37 40.33 17 l-5a 1 3+bc

43.61 43.59 42 1 1 3

This indexing scheme was successfully extended to the powder patterns of

all the compositions in the solid solution series. For example, for a Ca-rich sample

with X = 0.75, 5a = 0.171(1) and 5c = 0.227(1), which corresponds to an

368
incommensurate orthorhombic supercell with a = 5.848asub' c = 4.41csub- A

summary of the range of superstructures characterized for the Nd, Y and Gd

systems is shown below.

System Ca/M l/5a l/5c

Ca2+xY2-xCu5Oi0 1.0 5.0 4.03

1.5 5.405 4.17

2.25 5.848 4.405

Ca2+xGd2-xCu50io 0.77 4.524 5.0

1.0 4.878 4.566

1.5 5.263 4.484

Ca2+xNd2-xCu50io 0.625 4.0 0.0

1.0 4.76 5.714

1.33 5.0 5.0

The validity of this indexing scheme was further confmned by electron

diffraction. Although a detailed description of the electron diffraction patterns will

be described elsewhere, in figure 3 we show electron diffraction pattems collected

from the same samples shown in figure 1 . These patterns were collected along the

b axis of the orthorhombic sub-cell, and clearly show the superlattice reflections

originating from the ordering along the a and c directions of the sub-cell. The

periodicities of the supercells, which are indicated on the pattems for the a*

direction, are in excellent agreement with those calculated from the x-ray pattems.

STOICHIOMETRY RANGE, VALENCE CHANGES.

The limits of single phase formation for the Ca2±xM2±xCu50io systems are

a function of the chemistry of the trivalent M ions. In all three systems Ca-rich

samples can be prepared at 1000°C, with the sohd solution moving toward a Ca4Cu50io end-member. It was found that the range of solution increases with

Y>Gd>Nd. For all the Ca-rich samples the oxygen stoichiometry remains

369
Figure 3. [010] diffraction patterns for a series of samples in the

Ca2±xY2±xCu50io system.

constant, implying a stabilization of the 3+ state for Cu. However, in all cases the

samples remain non-conducting. Trivalent-rich samples could only be synthesized

for Nd and Gd, and the range of solution increased in accordance with the ionic

radius of the cation, with Nd>Gd. In contrast to the stoichiometric 225 and Ca-rich

compositions, for the M3+-rich samples the (Ca-i-M):Cu ratio was < 4:5, and the

composition of the solution proceeded toward an M3.33CU5O10 (i.e. Nd2Cu306)

end-member, with the copper ions remaining in the divalent state. These

compositional variations were reflected by an abrupt change in the lattice parameters

of the sub-cells in the Nd and Gd systems at Ca/M<l, with a in particular becoming

shorter. At the limit of the Nd solution, with an approximate composition

Nd2.5Cai.25Cu50io, i.e. Nd2CaCu408, the supercell is commensurate with a=

4asub» c = c, see figure 4.

370
JU1

1

1 1 1 25.0 3Q0 35.0 40.0 45.0 50.0

29

Figure 4. X-ray pattern of Nd2.5Cai.25Cu50io. Commensurate superlattice, peaks

shaded, with a = 4asub, b=b, c=c.

SUMMARY.

In summary, we have been unable to stabilize double-sheet T' analogs of

the La2SrCu206 structure. . Phase equilibria studies conducted at 1000°C in air

indicate that the M203-CaO-CuO systems (M=Y,Nd,Gd) are dominated by a new

ternary phase Ca2±xM2±xCu5Oi0. The structures of these layered cuprates are related to that of NaCu02 and contain linear chains of edge- shared Cu square

planes. These structures exhibit complex x-ray and electron diffraction patterns that can be indexed using incommensurate supercells. It is possible to stabilize valences

of Cu >2+, however all of the phases are non-metallic.

We acknowledge support from the National Science Foundation through

grants DMR-88- 19027, and through the MRL program, grant DMR88- 19885.

REFERENCES.

(1) Y. Tokura, H. Takagi, and S Uchida, Nature, 337, 345 (1989)

(2) N. Nyugen et al.. Mat. Res. Bull., 75, 891 (1980)

(3) R.J. Cava et. al.. Nature, 345, 602 (1990)

(4) T. Siegrist, R.S. Roth, C.J. Rawn, and J.J. Ritter, Chemistry of Materials, 2,

192 (1990)

(5) K. Hestermann and R. Hoppe, Z. Anorg. Allg. Chem., 367, 261 (1969)

371

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

THERMODYNAMIC MEASUREMENTS IN THE Y-BA-CU-0

SYSTEM

Fernando H. Garzon and Ian D. Raistrick

Los Alamos National Laboratory

The discovery of three high temperature superconducting phases: YBa2Cu307.x,

YBa2Cu3.50i5-x and YBa2Cu408 within the ytrrium-barium-copper-oxygen

system has aroused interest in the thermochemistry of these compounds and other

coexisting, non-superconducting phases. Measurement of thermodynamic values

of the formation of these compounds such as enthalpies (AH°), entropies (AS°) and

Gibbs free energies (AG°) can lead to a better understanding of the phase equilibria,

reactivity and chemical bonding within these novel materials. Enthalpies of

formation are of particular value as they may be evaluated together with heat

capacity data to provide AG° and AS°. The enthalpy of formation of the

YBa2Cu307.x phase has been measured by Morss et al. using isoperibol solution

calorimetry [1]. The focus of this work is on enthalpic measurements of

compounds that exist within the Y-Ba-Cu-0 quaternary system using precision

isothermal solution calorimetry.

EXPERIMENTAL

Samples of YBa2Cu307-x , Y2BaCu05, and BaCu02 were synthesized from Y2O3,

CuO and BaC03 by grinding the components in a mortar and pestie, and firing at 900°C for 10 hours in flowing dry oxygen. The samples were then quenched to room temperature, reground, pelletized and re-annealed for 10 more hours. The

YBa2Cu307-x sample was reheated at 400°C in oxygen for three hours then slowcooled to room temperature. The YBa2Cu408 sample was produced by mixing

YBa2Cu307.x and CuO and heating for 10 days at 850°C at 1 atm of oxygen

pressure. X-ray diffraction was used to check the phase purity of the samples. The

only impurity found was trace amounts of CuO in the YBa2Cu408 sample. Oxygen

stoichiometry analysis of the YBa2Cu307.x sample was performed by hydrogen

reduction of the compound in a Perkin Elmer TGA7 thermogravimetric analyzer.

The mass change of the sample corresponded to an initial oxygen content of 6.96.

A Tronac 558 isothermal solution calorimeter, modified for perchloric acid

calorimetry, was used for the determination of the enthalpies of solution of

YBa2Cu307-x , Y2BaCu05, YBa2Cu408 and BaCu02in 4.00 M perchloric

acid(Baker Ultrex grade) at 25°C . The Tronac calorimeter consists of a stirred

teflon and platinum isothermal calorimetric vessel that maintains bath temperature

373
by a combination of thermoelectric cooling and resistive heating. Heat events

occurring within the vessel are compensated by an automatic adjustment of the

resistive heater power, thus providing a direct method of measuring the magnitude

of the heat event. The calorimeter operates submerged in a high stability isothermal

bath (±0.001°C/week). The advantages of this technique over the more commonly

used quasi-adiabatic calorimetry (isoperibol) are: a) During the course of the

experiment, the calorimetric vessel temperature does not deviate appreciably from

the bath temperature thus vessel heat leak corrections are almost negligible, b) The

heat flux is sJso measured directly rather than back-calculated from the temperature

rise of a quasi-adiabatic vessel.

The calorimetric measurement procedure consists of the following steps. The

samples (3-5 determinations per compound) are fu-st encapsulated in borosilicate

glass ampoules, which are attached to the stirring rod of the isothermal vessel that

contained 25 ml of 4.00 M HCIO4. The calorimetric vessel is submerged in the

bath and allowed to reach thermal equilibrium. An electrical calibration is made by

the addition of a known amount of joule heat from a calibration resistor immersed in

the solution within the vessel. The sample-containing ampoule is broken by a

translation of the stirring rod downward against the bottom of the isothermal vessel.

The heat of dissolution is then determined by automatic electrical compensation.

After the system reaches thermal equilibrium a second joule heat calibration is made. As AH° is a state function, standard- state enthalpies of formation may be

calculated by the combination of enthalpies of appropriate chemical reactions whose

net summation yields the elemental formation reaction for the compound at standard

state conditions.

RESULTS

The enthalpies of dissolution of the measured compounds are listed in Table 1. The

enthalpy of solution of the YBa2Cu307.x compound was in excellent agreement

with Morss et al.[l]. Thermodynamic data for heats of solution of other

compounds used in the thermodynamic cycles were taken from ref. [1].

Compound A//^29S solution (kJ/mole)

YBa2Cu307.x fx=0.04) -796.1±5

YBa2Cu408 -847.5 ± 5

YiBaCuOs -655.5 ± 19

BaCu02 -249. 1± 0.2

Table 1. Heats of dissolution of listed compounds in 4.00 M HCLO4 at 298 K.

374
YBa2Cu307.x :

{ 1 } Y(ClO4)3ra^;+2Ba(ClO4)2(tz<7;+3Cu(ClO4)2ra<7;+6.5H2O^a^;+0.23O2

YBa2Cu306.96 +13HC104ra^)

AH{i}=796±5kJ/mole

{2} ¥^5; +3/402(gj =0.5Y2O3(,; ^ AH{2} =0.5(-1905.6±2.0 kJ/mole)

{3} 0.5Y2O3(5j+ 3HC104(a^; =Y(C\04)3(aq)+ l-5U20(aq)

AH{3} = 0.5(-365.7±2.0 kJ/mole)

{4} 2Bsi^s)+2C\2(g) + 22/202(o)+6R2(g)= 2Ba(C104)2«3H20^^;

AH{4} =2(-1691.6±3.0 kJ/mole)

{5} 2Ba(C104)2 •3H20^,; = 2Ba(C104)2ra^; +6H20^a^;

AH{5) =2(+32.1±3.6 kJ/mole)

{6} 3Cu^^;+ 3Cl2(gj + 2102(g; + im2(g) = 3Cu(C104)2»6H20^^;

AH{6] =3(-1928.4±1.0kJ/mole)

{7 } 3Cu(C104)2'6H20(^; = 3Cu(C104)2(fl^; + im20(aq)

AHn) =3(+5.7± 0.2 kJ/mole)

{ 8 } 10HClO4fa<?; = 5C\2(g) + 20O2fg; +5ll2(g)

AH{8) =10(+127.1±1.0 kJ/mole)

{9} l9U20(aq) =l9U2(g) + 19/202(^;

19(+286.10±0.5 kJ/mole)

AH°{f}i23 =Y(,; + 2Ba(,; + 3Cu^,;+ 6.96/202(g; = YBa2Cu306.96f5;

AH^ifjns = ZAH{i-9}= -2718 ±20 kJ/mole

YBa2Cu408:

{ 1 } Y(ClO4)3(a<7j+2Ba(ClO4)2ra^j+4Cu(ClO4)2fa^;+7.5H2O(a<y;+0.5O2 = YBa2Cu408(,; + l5UC\04(aq)

AH{i} = 847.5 ±5kJ/mole

{2} Y^,; +3/402(g) =0.5Y2O3(,;

AH{2} =0.5(-1905.6±2.0 kJ/mole)

{ 3 } 0.5Y203f,j + 3HCIO4 (aq) = Y(C104)3fa^;+ 1 .5H20^a^;

AH{3} = 0.5(-365.7±2.0 kJ/mole)

{ 4 } 2Ba,',;+ 2C\2(g) + 221202(g) +6H2fg; = 2Ba(C104)2 •3H20(^;

AH{4) =2(-1691.6±3.0 kJ/mole)

375
{ 5 } 2Ba(C104)2 •3H20(,; = 2Ba(C104)2(a^j +6H20(aq)

AH{5} =2(+32.1±3.6 kJ/mole)

{6} 4Cu(s)+ 4C\2(g) + 2m2(g) +24H2(gj = 4Cu(C104)2*6H20(5;

AH{6} =4(-1928.4±1.0 kJ/mole)

{ 7 } 4Cu(C104)2-6H20^,j = 4Cu(C104)2fa^j + 24H20(a^j

AH{7} =4(+5.7± 0.2 kJ/mole)

{ 8 } 12HC104ra^; = 6C\2(g) + 2402(g; +6U2(g)

AH{8} =12(+127.1±1.0 kJ/mole)

{ 9 } 24H20^^) = 24U2(g) + 24/202f^;

24(+286.10±0.5 kJ/mole)

AH°{f}i24 =Y(s) +2BsL(s) +4Cu(s)+ 402(g) = YBa2Cu408f5)

AH°{f} 124 = ZAH{ 1.9}= -2906.4 kJ/mole ±25 kJ/mole

Y2BaCu05:

{ 1 } 2YiC\04)3(aq)-^ B?i(C\04)2(aq)+Cu(Cl04)2(aq)+ 5H20(aq) = Y2BaCu05r.j + l0UClO4(aq)

AH{i} = 655.5±19kJ/mole

{2} 2Y^,; +3/202fg) = Y203f,)

AH{2} =-1905.6±2.0 kJ/mole

{ 3 } Y203f,; + 6HCIO4 (aq) = 2Y(C104)3fa<?j+ 3H20(a^)

AH{3} =-365.7±2.0 kJ/mole

{ 4 } Ba(^;+ C]2(g) + 1 V202(g) +3H2(g) = Ba(C104)2 •3H20(^;

AH{4} =-1691.6±3.0 kJ/mole

{ 5 } Ba(C104)2 •3H20(,) = Ba(C104)2(a^j +3U20(aq)

AH{5) =+32.1±3.6kJ/mole

{6} Cu^,;+ C\2(g) + 102(g) +6H2(g) = Cu(C104)2-6H20(,;

AH{6) =-1928.4±1.0 kJ/mole

{ 7 } Cu(C104)2-6H20^,) = Cu(C104)2f^<y) + 6H20^^^j

AH{7) =+5.7± 0.2 kJ/mole

{ 8 } 4llC\04(aq) = 2C\2(g) + ^2(8) +2U2(g)

AH{8) =4(+127.1±1.0 kJ/mole)

376
{9} m20(aq) =m2(g) + 7/202(g)

7(+286.10±0.5kJ/mole)

AH°{f}2ll5 =2Y(s) +B^s) +Cu(s)+ 5/2 02(g) = Y2BaCu05f,j

AH°(f}2ii5 = IAH( 1.9}= -2686.9 ±22 kJ/mole

BaCu02:

{ 1 } Ba(C104)2ffl^j+Cu(C104)2fa<?j+ 2H20(aq) = BaCu02f,) + 4HC104(a^;

AH{i} =249.1±0.2kJ/mole

{ 2 } B2i(s) + C\2(g) + 1 l/202(g) +3H2(g) = Ba(C104)2 •3H20(,;

AH{2) =-169L6±3.0kJ/mole

{ 3 } Ba(C104)2 •3H20(,j = Ba(C104)2(fl^j +3H20(a^)

AH{3} =+32.1±3.6kJ/mole

{4} Cu(s)+ Cl2(g) + 702(g) +6n2(g) = Cu(C104)2-6H20(,j

AH{4) =-1928.4±1.0 kJ/mole

{ 5 } Cu(C104)2-6H20(,; = Cu(C104)2fa^j + 6H20(aq)

AH(5) =+5.7± 0.2kJ/mole

{ 6 } 4HC\04(aq) = 2C\2(g) + ^02(g) +21l2(g)

AH{8) =4(+127.1±1.0 kJ/mole)

{7} 7H20(^<^; =7H2fg) + 7/202fg)

7(+286.10±0.5kJ/mole)

AH°{f}oii2 =Baf^j +Cu(s)+ 02(g) = BaCu02('5j

AH°{f}oii2 = lAH{i.7}= -822.3±7.9 kJ/mole

We have analyzed the following two possible decomposition reactions at the 1 :2:4 and the 1:2:3 cation ratio compositions using a value of -155.2 kJ for the heat of

formation of CuO[2] and a partial molar enthalpy of oxidation of -190 kJ/mole O2

for the YBa2Cu307-x compound[l] :

YBa2Cu307.x + CuO +x/2 O2 = YBa2Cu408 (I)

4YBa2Cu307-x + (2x-l/2)02 = 2YBa2Cu408+Y2BaCu05+ 3BaCu02 (II) The decomposition reaction represented by reaction (I) at 298 K and 1 atm does

not involve oxygen gas at x=0; under these circumstances AG(i) ~ AH(r). The value

obtained (-38 kJ/mole) clearly indicates that YBa2Cu3C)7.x and CuO will transform

into YBa2Cu408 at low temperatures under equilibrium conditions.This has been

experimentally observed by a number of researchers [ 3-7]

377
Reaction (II) is a completely solid-state reaction at the x= 0.25 composition. This is the equilibrium YBaiCusOv-x stoichiometry at approximately 873 K and 1 atm of

oxygen pressure [8]. The AGrx = A//rx value of -44 kJ/mole indicates that, at this

pressure and temperature, YBa2Cu307.x is thermodynamically unstable with

respect to YBa2Cu408, Y2BaCu05, and BaCu02. At higher temperatures reactions

(I) and (II) will no longer be favorable due to the increasingly positive entropic

contribution of the change of state of oxygen gas to lattice oxygen. This effect will

eventually change the sign of the AGrx and thus YBa2Cu307.x will become stable, as is observed experimentally.

The implications of the thermodynamic instability of the YBa2Cu307.x structure

limit the ability to synthesize the YBa2Cu307.x compound at low temperatures and

may effect the electronic properties of the materials by the appearance of

microscopic defects within the YBa20i307.x lattice as the material is cooled below

its boundary of thermodynamic stability.

The authors would like to thank Dr. David S. Ginley of Sandia National

Laboratories and Dr John W. Halloran of CPS Corporation for supplying the

YBa2Cu408 material and helpful discussions.

REFERENCES

[1] L. R. Morss, D. C. Sonnenberger, R. J. Thorn, Inorg.Chem. 27, 2106-2110

(1988)

[2] O. Kubaschewski, C. B, Alcock, Metallurgical Thermochemistry 5th Ed.,

Pergamon Press, Oxford, 282, (1979)

[3] R.J. Cava, J.J. Krajewski, W.F. Peck, B.Batlogg, L.W. Rupp, R.M.

Fleming, A.C.W.P. James, P. Marsh , Nature 334 328-330 (1989)

[4] D.E. Morris, N. G. Asmar, J. H. Nickel, R. L. Sid, J. Y. T.Wei, J. E. Post,

Physica C, 159, 287-294 (1989)

[5] J. Karpinski, S. Rusiecki, E. Kaldis, B. Bucher, E. Jilek, Physica C, 160,

449-457 (1989)

[6] J. Karpinski, S. Rusiecki, E. Kaldis, B. Bucher, E. Jilek, Physica C, 161,

618-625(1989)

[7] S. Jin, H. M. O'Bryan, P.K. Galagher, T.H.Tiefel, R.J. Cava, R.A.

Fastnacht, G.W. Kammlott, Physica C, 165, 415-418 (1990)

[8] K. Kishio,, T. Hasegawa, K. Suzuki, K. Kitazawa, IC Fueki, Mat. Res.

Soc. Symp. Proc. 156, 91-97(1989)

378
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

CALORIMETRIC STUDIES OF CERAMICS

Alexandra Navrotsky

Department of Geological and Geophysical Sciences

and Princeton Materials Institute

Princeton University

Princeton NJ 08544

ABSTRACT

High temperature reaction calorimetry is used to

study a variety of ceramic problems , Lead borate

solution calorimetry near 973 K and transposed

temperature drop and scanning calorimetry to 1773 K are

complementary techniques. This paper illustrates two

diverse applications: the energetics of oxide

superconductor phases in both the YBCO and doped

lanthanum cuprate systems and in situ studies of melting

and crystallization in a glass-forming silicate,

CaMgSi206.

CALORIMETRY

Recent advances, in both isothermal and scanning

experiments (see Table 1), make possible direct study

of ceramic processes and processing at 873-1773 K.

Three types of calorimeters are in use in our

laboratory. Custom built Calvet twin microcalorimeters

operate isothermally , typically near 973 K. Two sample

chambers, roughly 2.5 cm in diameter by 15 cm high, each

surrounded by sensitive thermopiles, are embedded in an

alloy block maintained at constant temperature (1) . Chemical reactions (e.g dissolution of a sample into

molten lead borate solvent) take place in the sample

chamber and heat flow is monitored. The relatively large

size of the sample chamber permits the introduction of

sample into solvent through a mechanical stirring

device. The high sensitivity allows the measurement of

small heat effects (0.3 to 1 J per experiment) and the

use of small samples (10-50 mg of sample dissolving into

20-30 g of solvent) . Samples can also be dropped into

this calorimeter from room temperature. When no solvent

is present, a transposed temperature drop experiment

379
Table 1 High Temperature Calorimetric Experiments

Calorimeter essentially isothermal

Solvent present

Solution calorimetry: sample equilibrated in

hot calorimeter, then dissolved.

Differences in heats of solution give

heat of reaction at calorimeter T

Drop solution calorimetry: sample dropped from

room T. Sample may be encapsulated in Pt,

pyrex glass (which dissolves) or lead

borate glass (same as solvent), or as

unencapsulated chunks. Differences in

heats of drop-solution give heat of

reaction at room T.

Differential drop solution calorimetry: sample

in lead borate capsule dropped on one

side simultaneously with matched empty

capsule on other side, improved

sensitivity for 5 mg samples

Solvent absent

Transposed temperature drop calorimetry

No permanent change in sample: heat

content measurement, includes heat of

any reversible phase change

Sample changes oxidation state:

difference between first and second

drop gives heat of redox reaction

Sample loses H2O, CO2 or other volatiles:

difference gives heat of

devolat ilizat ion

Sample undergoes irreversible phase

change : difference gives heat of

transformation at room T.

Sample undergoes change in degree of

order or other structural parameters

difference gives ordering energy.

Calorimeter changes temperature

At constant rate

Scanning calorimetry: generally

semiquantitative measure of heat capacity

and phase change or decomposition

In steps separated by isothermal periods

Step scanning calorimetry: best measurement

of heat capacity and reversible phase

transition enthalpies

380
measures the heat content of the sample and any

transformation which occurs rapidly at calorimeter

temperature (e.g. oxidation, dehydration, rapid phase

transition, melting, or structural change) (1,2,3) . When

solvent is present, the drop solution experiment

measures the sum of heat content and heat of solution

(2) . Drop solution calorimetry is useful when the sample

in question can not be equilibrated at calorimetric

temperature prior to dissolution because of

decomposition, as is often the case with metastable or

readily oxidized or reduced materials (2,3).

A commercial Setaram HT 1500 calorimeter also uses

a thermopile to measure heat flow. Rather than being

embedded in a massive alloy block, this much smaller

calorimetric detector hangs free in a furnace, and can

perform transposed temperature drop experiments at any

temperature from about 873 to 1773 K (4) . It can also be

operated in scanning mode (continuous heating or

cooling) or step-scanning mode (rapid temperature change

by a small increment followed by an isothermal

equilibration period) (5) . For the latter experiments,

using about 1 g of sample, we have made significant

modifications to the original sample assembly and

general operating procedure, as well as written our own

software

This calorimeter is too small for easy use for

solution calorimetry. Furthermore, its calibration

factor is sensitive to sample geometry and to detector

positioning in the furnace, which changes slightly from

run to run because the entire detector must be removed

to change samples. We are presently constructing a calorimeter which should overcome some of these

difficulties, combining features of our solution

calorimeters with scanning capability using a thermopile permanently positioned in a furnace.

A commercial Setaram DSC 111 uses a very

miniaturized thermopile detector in a silver block for

differential scanning calorimetry from about 200 to

1073 K. It can measure heat capacities on 50-100 mg

samples and detect phase transitions, oxidation, and

other changes. Optimum heat capacity measurements using

this instrument take considerable attention to detail.

We use our own software rather than the commercially

available package.

381
The DSC 111 is most useful as a semi=quant itative

exploratory tool, as is the HT 1500 when used at a rapid scan rate (10 K/min) . Once transitions are

located, they can be studied in quantitative detail.

Enthalpies of formation of complex oxides from their

binary oxide components and of mixing in solid solutions

are best determined using solution or drop solution

calorimetry (1) . Enthalpies of phase transitions can be

determined by these two methods if both phases in

question can be quenched to ambient (1,2) . If a phase

transition is rapid and reversible, scanning or stepscanning can be used (6) . The ability to scan to 1773 K

opens many new ceramic applications (e.g. melting, glass

transitions, rapid displacive phase transitions, orderdisorder . ENERGETICS OF PHASES IN DOPED LANTHANUM CUPRATE

AND YBCO SYSTEMS

Systems La2-xAxCu04 -y (A = Ba, Sr, Ca, Pb)

In these phases with structures derived from the

K2NiF4 type, the substitution of a divalent ion for

trivalent La is compensated by oxidation of copper from

formal 2+ to formal 3+ oxidation state and/or by the

creation of oxygen vacancies. We performed solution

calorimetry on a series of single phase materials for

Ba, Sr, Ca, and Pb, after having characterized their

cryst allographic parameters, metal ratios and oxygen

contents (8)

.

X-ray diffraction data show that at room

temperature, an orthorhombic to tetragonal phase change

is observed at x=0 . 1 for Ba, Sr, and Ca . Johnston et al.

reported that La2Cu04 undergoes a second order

orthorhombic to tetragonal phase transformation between

506-536 K (7) . Since solution calorimetry was performed

at 977 K , the data obtained pertain to the tetragonal

K2NiF4 structure for all samples. Oxygen contents,

obtained by iodometric titration and/or

thermogravimet ry , for samples equilibrated in air at

973 K are shown in Fig. (la) . For Ba, Sr, and Ca

.

Initial substitution retains the stoichiometric oxygen

content (y<0.02 for x<0.10), but increasing substitution

causes oxygen vacancy formation to approximately the

same extent for all three cations. Lead substitution

results in oxygen loss for all values of x to the extent

382
Fig.l Doped Lanthanum Cuprates

(a)

O Ba

A Sr

D Ca

O Pb

0.0 0.2 0.4 0.6 0.8 1.0 1.2

+2 Ion Content (x)

0

-10 -

-20'

-30

-40

Heat of formation for observed

oxygen stoichiometry, (y)

i J I L

(b)

O Ba

A Sr

H Ca

O Pb

0.0 0.2 0.4 0.6 0.8 1.0 1.2

•i-2 Ion Content, (x)

(C)

A Sr

O Pb

0.0 0.2 0.4 0.6 0.8 1.0 1.2

+2 Ion Content (x)

383
the enthalpy of formation of the Sr-subst ituted

compound (AH4) from the binary oxides remains very

similar to that of pure La2Cu04 . AH4 in turn reflects

the balance of two factors: (a) differences in the

intrinsic properties of La and Sr (ionic size, charge,

"basicity") which affect the energetics of the

substitution of Sr for La, and (b) the enthalpy of

oxygen vacancy formation. The nearly constant value of

AH4 suggests either that these two effects cancel each

other in the Sr-La substitution, or that both terms are

small

.

If the same calculation is done for Ba-La using

AH3 derived from the Sr system, then values of AH4 are

slightly negative. This may be because AH3, is

actually more exothermic or that AH4 is more exothermic

for the Ba system.

If only enthalpies are considered, all these phases

should have y values of zero because oxidation is

clearly exothermic. Oxidation involves a large negative

entropy due to fixing molecular oxygen into a crystal

structure. A typical value for the entropy of oxidation

for transition metals is -92 J/mol-K (9). At 977 K,

this entropy can compensate the large negative enthalpy

to yield non-zero y values. This entropy argument also

explains why, as T increases (at given imposed external

oxygen fugacity) the degree of oxidation of a given

phase generally decreases.

Thermodynamics of Phases in the Y-Ba-Cu-0 System

High temperature reaction calorimetry has been done

on YBa2Cu30x and related phases in the Y-Ba-Cu-0 system.

(3,14) . Preliminary analysis of the oxidation state

of Cu in lead borate indicates that the reaction during

drop solution calorimetry is : YBa2Cu30x (298 K) + (6.5-x)/2 O2 (g, 975 K) = {1/2 Y2O3 + 2 BaO + 3 CuO} (soln, 975 K) (5)

384
that no formal oxidation states greater than Cu^"*" (and

Pb2+, o2-) are detected.

Heats of solution represent the reaction (at 977

K) : La2-xAxCu04-y (xl) = x AO (soln) + (l-x/2) La203 (soln)

+ CuO (soln) + ((x-2y)/4) O2 (g) (1)

From these values and the heats of solution of the

component oxides the heat of formation at 977 K can be

calculated:

(l-x/2) La203 (xl) + (x)AO (xl) + CuO (xl)

+( (x-2y) /4)02 (g) = La2-xAxCu04-y (xl) (2)

Fig. (lb) shows the heat of formation, AHf, at 977

K as a function of divalent substitution, x. As x increases to about 0.15, AHf becomes more exothermic,

with the magnitude of the effect decreasing in the order

Ba > Sr > Ca > Pb. At higher concentrations, the heats

of formation for Ba and Sr level off, while those for

Ca and Pb become more endothermic.

These observations can be interpreted as follows.

The enthalpy of formation (Eq. (2)) can be divided into

two steps : AxLa2-xCu04- (x-2y) /2 + (x-2y) /4 O2 = AxLa2-xCu04 (3)

X AO + (l-x)/2 La203 + CuO = AxLa2-xCu04- (x-2y) /2 (4)

Eq. (3) represents oxidation of an oxygen deficient

phase containing only Cu2+ to a fully oxidized phase

free of oxygen vacancies and containing the maximum

amount of (formal) Cu-^"*". The enthalpy associated with

this process has been estimated as -157 ± 10 kJ per mol

of O2 ) by comparing the transposed temperature drop

enthalpies of two samples, each having different oxygen

contents, of both Lai . 25Sro . 75Cu04-y and of LaSrCu04-y . The enthalpy of step (4) can then be calculated by

subtracting AH3 from AH2 . The results (see Fig.(lc))

show that, when the effect due to oxidation is removed.

385
Data for the same system were obtained previously by

transposed temperature drop calorimetry to 1057 K (3) . Both show a linear dependence of enthalpy on oxygen

content, x. The two lines are almost parallel, and give

a partial molar enthalpy of oxygen (oxidation enthalpy)

,

AHqj^ = -209±7 kj mol~l per O2 (present data) and AHqj^ = -190±5 kJ/mol (3) . The intercalation of oxygen is

exothermic and energetically stabilizes structures with

higher oxygen contents. The constancy of AHqx with

composition implies that the oxygen energetics are

insensitive to the orthorhombic tetragonal transition

and that the exact distribution of valence states

changes continuously, with no titration end-point at

x=6.5 (3)

.

Measured enthalpies of formation of YBa2Cu30x and

related phases are listed in Table 2. Large exothermic

enthalpies of formation from simple oxides are due to

the basicity of BaO . Y2CU2O5 has a small endothermic

Table 2 The Enthalpies of Formation of YBa2Cu30x and

Related Phases from Oxides and from Elements

Compound AHf (kJ/mol)

from oxides from elements

YBa2Cu30x (295 K)

x=6.33 -79 . 4±7 3

a -2575±9

x=6 .45 -76 . 8±4 9 a -2590±7

x=6 .48 -74 . 1±6 1 a -2592+8

x=6 . 64 -86 . 8±7 .2 b -2608±9

x=6.87 -101 .8±6 . 4 b -2623±9

x=6. 92 -110 .3±9 .5 b -2631+11

Y2BaCu05 (295 K) -52 9±3 . 0

c -2 6 64+5

Y2BaCu05 (975 K) -59 3±3 . 0 c Y2CU2O5 (295 K) + 9 . 2+3 . 2

c -2211±5

Y2CU2O5 (975 K) + 6. 6±2 . 9

c BaCu02.oi (295 K) -83 . 8±3 . 1 b -7 8 9±4

YBa2Cu408 (295 K) -145 . 8±10 . 4 b -2824+12

a. from BaO Y2O3, CuO and CU2O

b. from BaO, Y2O3, CuO and O2

c. from BaO, Y2O3, CuO

386
enthalpy of formation from the oxides, implying that it

is entropy-stabilized.

To estimate the stability of various phases to

possible decomposition, we choose reactions involving

neither oxidation nor reduction, so that enthalpy can be

taken as an estimate of free energy, since entropy

changes of solid-solid reactions are usually small. For

the reaction:

YBa2CU307 + CuO - YBa2CU408 (6)

AH298 = -78.1±5.3 kJ/mol. This implies that the 1:2:4

phase is indeed more stable at its own composition than

a mixture of fully oxygenated 1:2:3 and CuO.

For the reaction:

1/2 Y2BaCu05 + 3/2 BaCuO2.01 + CuO = YBa2Cu306 . 515 (7)

AH298 = 28.9+4.5 kJ/mol . This suggests that the 1:2:3

compound with intermediate oxygen content near x=6.5 may

be unstable with respect to a mixture of "green phase"

and "BaCu02", and may explain why these phases are often

found as impurities in the 1:2:3 material. We are

investigating whether the fully oxygenated 1:2:3 phase

has any stability field in T- f(02) space.

DIRECT MELTING AND CRYSTALLIZATION OF CaMgSi20 6

CaMgSi206, diopside, is a common mineral and

constituent of natural and synthetic silicate melts and

glasses. Its nominal melting point is 1665 K but some

incongruency below that temperature is now established

(4,5) . Rapid cooling (>100 K/min) produces a glass;

slower cooling (1-100 K/hr) produces crystallization of

a mixture of phases at 50-100 K below the liquidus

Using step-scanning mode in the HT 1500 calorimeter, we

traced the melting of diopside and determined that, for

a stoichiometric crystalline sample, incongruent

melting begins at 1606 K. (5) . The first 20% of melting

occurs over a 44 K interval, and the remaining 80%

within 15 K. Scanning calorimetry is a very sensitive

tool for detecting small amounts of melting (see Fig.

(2)) . The enthalpy of fusion, 137.7 kJ/mol, (see Fig.

(2)) is in excellent agreement with previous

measurements by transposed temperature drop calorimetry

387
Fig. 2

Diopside melt ing : step scanning calorimetry

O

E

Q.

o

15000

10000-

5000

Heat capacity

2000

1400 1 500 1 600 1700 1800

Temp (K)

o

E

Q. O

start of

melting „^ yliquidus

J I L

Heat capacity,

expanded scale

1 550 1 600 1 650 1 700 1 750

Temp (K)

500

o

E

==- 400 -

CO

a>

CM

300 -

200

Heat content

1 400 1 500 1600 1700 1800

Temp (K)

388
(5) and by thermochemical cycles involving the glassy

phase (10) . The heat capacities derived for the

crystalline and liquid phases are in good agreement with

previous values obtained by conventional drop

calorimetry (10) . Step scanning provides a sensitive and

accurate way of determining the heat capacities of

silicate liquids . We are currently studying a series of

Fe203-bearing molten silicates.

We performed crystallization experiments in the

calorimeter using continuous cooling at different rates

(5) . An initial experiment at 30 K/hr starting at 1715 K

produced a single large exotherm at 1545 K. The product

was fully crystalline, consisting of two phases, a magnesium-rich diopside solid solution Cai-xMgi+xSi206

and a slightly magnesian wollastonite Cai -xMgxSi03 . Three experiments were done at a cooling rate of 10

K/hr, starting at 1691 K. Each produced single

exotherms, but, though the conditions were nominally

identical, the exotherms occurred at somewhat different

temperatures, 1576, 1534, and 1566 K, Our experiments

are in general agreement with the kinetic

crystallization study of Kirkpatrick et al. (11),

although we could not confirm the small amounts of

forsterite, Mg2Si04, which they reported.

In situ studies of crystallization in glass

ceramics should now be possible. Since the enthalpy

released is a sensitive indicator of crystallization,

scanning calorimetry can be used directly to construct

time-temperature-transformation (T-T-T) curves. When a metastable phase is initially formed (e.g. beta-quartz

in a cordierite ceramic) , its subsequent annealing and

transformation can be followed. The incubation or prenucleation stages of crystallization (during which no

crystallization is detected by bulk methods such as

diffraction) may nevertheless involve significant

changes in glass structure. Since much of the energy

difference between glass and crystal may come from

changes in local environments rather than from long

range order per se (12), we propose that a significant

part of the enthalpy of vitrification may be released

before crystallization becomes detectable by

conventional means. This hypothesis should be tested in

glasses for which homogeneous rather than heterogeneous

nucleation is suggested.

389
ACKNOWLEDGEMENTS

This work was supported by the National Science

Foundation, the Department of Energy, and IBM

Corporation. The following people participated in these

studies at Princeton: C. Bennett, J. Bularzik, J.

DeYoreo, J. DiCarlo, R. Lange, K. Mocala, M. Parks, L.

Topor, Z. Zhou, D. Ziegler. The following collaborators

were involved: J. Bringley, P. Davies, A. Jacobson, B.

Scott, E. Takayama-Muromachi, S. Trail.

REFERENCES

1. Navrotsky, A. Phys . Chem . Min . 2, 89 (1977).

2. Ashida, T., Kume, S., Ito, E., Navrotsky, A.

Phys. Chem. Miner. 16 239 (1988).

3. Parks, M. E., Navrotsky, A., Mocala, M.,

Takayama-Muromachi, E., Jacobson, A., and Davies, P. J

.

Solid State Chem. 79 . 53 (1989) . Erratum (in press) . 4. Ziegler, D., Navrotsky, A. Geochim. Cosmochim.

Acta 2461 (1986) . 5. Lange, R.A., DeYoreo, J.J., Navrotsky, A. Amer

.

Mineral

.

(in press)

.

6. Redfern, S.A.T., Salje, E., Navrotsky, A.

Contrib. Mineral. Petrol. 101, 479 (1989)

.

7. Johnston, D.C., Stokes, J. P., Goshorn, D.P.,

Lewandowski, J.T., Phys. Rev. B 36 . 4007 (1987).

8. Bularzik, J., Navrotsky, A. DiCarlo, J.

Bringley, J., Scott, B., and Trail, S. submitted. Fall

1990 Materials Research Society Meeting.

9. Navrotsky, A., in "M. T. P. Reviews of Science:

Inorganic Chemistry Series 2" (D.W.A. Sharp, Ed.), P.

29, Butterworth-Univ. Park Press, Baltimore, MD (1974)

.

10. Stebbins, J.F., Carmichael, I.S.E., Weill,

D.F., Amer. Mineral. 68 . 717 (1983).

11. Kirkpatrick, R.J., Kuo, L.C., Melchoir, J.

Amer. Mineral. M, 223 (1981) . 12. Navrotsky, A., Geisinger, K.L., McMillan, P.F.,

Gibbs, G.v. Phys. Chem. Mineral. 11, 284 (1985)

13. Zhou, Z., Navrotsky, A., Scott, B., Trail, S.,

submitted Fall 1990 Materials Research Society Meeting.

390
DISCUSSION

P. D. Davies: One thing that we both well know is the advantage of calorimetric

sensitivity to probing short-range order, which is quite difficult to quantify using

structural techniques. The ferroelectric relaxors discussed by Bob Newnham and other

speakers exhibit quite complex microstructural ordering phenomena. When these are made at very low temperatures, say 6(X)°C, by sol-gel techniques you might expect the

short range order to be different to that in samples prepared by conventional techniques

at high temperatures. Do you have any plans to examine these compounds to see whether

different synthesis methods lead to differing degrees of short-range order?

A. Navrotsky: Well, in fact, that was sort of behind some questions that I raised in my

talk. Certainly something like magnesium and niobium ordering, although the

electrostatic argument says the size of the domain is going to remain small, could well

vary with T. The ratio of ordered to disordered regions may be strongly dependent on

the preparation techniques.

Niobium is not an easy element to dissolve in high temperature oxide fluxes. We've

actually made some progress in this technique. Interestingly enough, this progress shows

the interplay of different fields. We first got started dissolving lanthanum, niobium and

titanium, for looking at glass structures, looking at the behavior of highly-charged cations

in glasses, partly for chemical systematics and partly for rad waste reasons. Yes, we can do it now and, yes, if somebody had some well-characterized samples, we'd be delighted

to try.

391

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

SYNTHESIS AND PROPERTIES OF

Ba-FREE SUPERCONDUCTIVE (Eu,Ce)4(Eu,Sr)4Cu6_xMx0z (M:Fe,Co,Al)

Ataru Ichinose, Takahlro Wada, Yuji Yaegashi, Akiko Nara

H. Yamauchi and Shoji Tanaka

Superconductivity Research Laboratory, ISTEC

10-13, Shinonome 1-Chome, Koto-Ku, Tokyo 135, Japan

Ba-free "446" [(Eu, Ce)4(Eu, Sr)4Cug02-type] samples were synthesized by partially

substituting Fe. Co or Al for Cu. We prepared Ba-free

ti(^^l/2^T^2/Z^ i^H-x^x'^z samples (M=Fe. Co and Al and x=0. 5, 1.0. 1.5 and 2.0). All the samples were of near single phase and the crystal structure

was tetragonal. As x increased, the length of the a axis increased, while that

of the c axis decreased. The samples with x=0. 5 and M=Fe and Co exhibited su- perconductivity with zero-resistance temperatures at 23 and 28 K. respectively.

None of the samples with M=A1 exhibited zero resistances, but those with x=0. 5 and

1.0 exhibited superconductivity onsets around 20 K. The appearance of superconductivity for the series of the samples with Al was explained in terms of the carrier density. However, the appearance of superconductivity for both series of

samples with Fe and Co was indifferent of the carrier density.

Cu(l) • :Nd, Ce(A1) O : r4d, Ba (A2)

• :Cu O :0

I . INTRODUCTION

Sawa et al.[l] discovered cuprate superconductors of a new type with Tc's around

40 K. These cuprates are represented by a general formula, (Ln,Ce)4(Ln,Ba)4Cug02

(Ln=Nd, Sm and Eu), and are termed as 446

oxides in the present paper. The structure of this family of cuprates is

perovskite- and f luor ite-related and

similar to both the triperovskite structure

of YBa2Cu302 and the Nd2Cu04 structure.

The crystal structure is illustrated in

Fig. 1. In the unit cell of this structure, two different types of sites exist

for lanthanide ions: the "Al" sites are oc- cupied by small ions, i.e., Ce^*, and

medium-sized ions, i.e., Nd^*, Sm*^* and Eu*^"^, while the "A2" sites

are occupied by large ions, i.e., Ba^*, as well as by the mediumsized ions. The oxygen site on the Cu(l)-0(1) plane was found to

be partially occupied[l]. That is, oxygen vacancies are randomly

Fig. 1. Crystal structure of the 446 conpound.

393
distributed on the Cu(l)-0(1) plane such that the crystal system

of this cuprate be tetragonal. Previously, we reported on 446

materials with Sr^* which partially substituted for Ba^* in the

A2 site[2]. We also reported that 446 oxides with two different

ions for medium-sized ion sites, i, e. , (Ln, Ce)4(La, Ba,

Sr)4Cu602 (Ln=Eu, Gd, Dy, Ho and Y),[3],[4] were able to be

prepared. In the present work, we synthesize Ba-free 446 oxides

by substituting metallic ions such as Fe, Co or Al ions for copper ions. The samples studied are given by the formula

(EU3 /4Ce2^/4)4(EUj^/3Sr2/3)4Cug_xMj^02, where M=Fe, Co and Al, and

x=0.5, 1.0, 1.5 and 2.0.

n. EXPERIMENTAL PROCEDURE

The starting materials were EU2O3, Ce02, SrC03, CuO, Fe203,

C0CO3 and A1(N03)39H20 powders. The powders were weighed to

nominal compositions of (Eu3/4Cew4)4(EU]^/3Sr2/3)4Cug_jjMx02 in

which M=Fe, Co and Al and x=0.5, 1.0, 1.5 and 2.0. The mixed

powder was calcined at 1000-1020 °C in air, and then pulverized

and pressed into parallelepiped bars. The bars were sintered at

1040 °C in O2 gas flow and slowly cooled. These samples were re- heated at 400-600 V at a high oxygen partial pressure, P(0 2)=200

atm.

The lattice constants were determined by x-ray powder diffraction using CuKa radiation. A curved graphite monochrometor was placed in the scattering beam path. The oxygen contents of the

samples were determined by a coulometric titration method[5] , [6]

.

Electrical resistivity was measured from room temperature down to

the He boiling point by a conventional dc four-probe method.

in. RESULTS & DISCUSSION

Figure 2 shows of the x-ray

diffraction pattern for

( Eu3/4Cei/4 ) 4 ( Eui/3Sr2/3 ) 4CU4

C02O2 annealed at P02=200 atm.

The peaks in the pattern are indexed for a tetragonal unit

cell with the lattice constants, a=3.860 X and c=27.95

X. The peaks which are not

identified as those of 446

phase are marked with asterisks. All the samples

0 CuKa

0

<

0 rS 10 IS 20 25 30 35 40 45 50 55 60 65

29 (deg)

Fig. 2. X-ray diffraction pattern for (Eu3/^Cej/^)^(Euj^Sr2/3)4Cu^Co20^ annealed at high oxygen partial pressure of 200 at*. were of near single phase and the crystal system of each sample

was tetragonal. The peaks marked with asterisks in Fig. 2 were larger for the samples with the smaller values of the metallic ion

content, x. The samples with x=0 were of multi phases and peaks

for a 446 phase were hardly observed. The relationship between

the oxygen content, z, and the metallic ion content, x, is given

in Figs. 3(a), (b) and (c) for three different series of samples

394
0.5 1.0 1.5 2.0 M Contents, x

0.5 1.0 1.5 2.0 M Contents, x

Fig. 3.

0.5 1.0 1.5 2.0 M Contents, x etallic ion content. x. for Relationship between oxygen content, z, and

(Eu3/4Cej/4)4(Euj/3Sr2/3)4Cu5_jjMjj02 with (a)M=Fe, (b)M=Co and (c)M=Al. Close circles are for as-sintered

samples and open circles for annealed samples at high oxygen partial pressure of 200 ata.

28.3

28.1 -<

27.9 £ o

c

m 27.7

incoO

27.5

a> 3.88

<a _l .—

3.86

(c)

3.82 0.5 1.0 1.5 2.03.82

3.84 -

3.82 0.5 1.0 1.5 2.0 0.5 1.0 1.5 2.0 M contents, x M contents, x M Contents, x

Fig. 4. Relationship between lattice constants and letallic ion content, x. for {^n-^/^^y^ \(^'^\/':^'^i/i> t^%-

j^O^ with (a)M=Fe. (b)M=Co and (c)M=Al. Close circles are for as-sintered saaples and open circles for annealed saiples at high oxygen partial pressure of 200 at*. containing Fe, Co and Al ions, respectively. The oxygen content

increased as x increased for each of the three series. For the

series of samples containing Al, as x increased, the oxygen con- tent, z, was initially low and leveled off at a constant value

around 18.2. The oxygen contents of the high-02-pressure an- nealed samples were larger than those of as-sintered samples for

the series containing Fe and Co. On the contrary, for the series

of samples containing Al, the oxygen contents were nearly identical before and after high-02-pressure annealing. These observations suggested that the oxygen contents for the series of Samples

containing Al were hard to be increased by high-02-pressure an- nealing. This was probably because the ionic radius of

Al3*(0.535 X)[7] was much smaller than that of Cu2*(0.73 X)[7].

Figures 4(a)-(c) show the relationships between the lattice con- stants, a and c, and the metallic ion content, x. For the series

395
of 446 samples containing Fe and Co, as metallic ion content, x, in- —

creased, the c-axis length decreased o monotonlcally , while the a-axis 2.

length Increased. It should be >. noted that both a and c axes were %

shorter after high-02-pressure an- « nealing for these two series of S

samples (Figs. 4(a) and (b)). Both

lattice constants, a and c, for the

series of samples containing Al

abruptly changed between x=1.0 and

1.5. However, the changes in the

a- and c-axis length before and "g" after high-02-pressure annealing^

were minimal (Fig. 4(c)), Thus, it"^

is likely that the changes in the^Slattice constants were caused by thc:^

oxygen added by the high-02-pressure

annealing procedure. ir The temperature dependence of

electrical resistivity is given in

Figs. 5(a), (b) and (c) for the

three series of the high-02-pressure

annealed samples containing Fe, Co ^

and Al for M, respectively. Among 0

the samples with M=Fe, superconduc- 9.

tivlty clearly appeared for the ^

sample with x=0.5 and the zero-

resistance temperature, Tp'^"^, for «

1010'

10'

10-* 10-

• X r 0.5

• X = 1.0 ° x = 1.5

• x = 2.0

50 100 150 200 250

Temperature (K)

100 150 200 250

Temperature (K)

10'

10*

10'

10'

10'

lO"

10"

10"

» • \ (c)

1 1 1

1 • X = 0.5 \

x = 1.0 " Ir1.5

» Xr2.0

50 100 150 200 250

Temperature (K)

Fig. 5. Teiperature dependence of resistivity for

(Euj/^Cej/^) 4 (Eu^/3Sr2/3) 4Cu5.^M^0^ annea led at high oxygen partial pressure of 200 ata with (a)M=Fe. (b)M=Co and (c)M-Al.

C • ^""^ i2 this sample was 23 K. The mag- S

nitude of resistivity of each sample was lower than that of the same

sample without high-02-pressure an- nealing. The zero-resistance temperature, T^,'^"^, and the superconductivity onset temperature, T,,®^,

of the sample with M=Co and x=0.5

increased respectively from 10 to 28

K and from 25 to 39 K by high-02-pressure annealing. The magnitude of resistivity of each sample in the series containing Co was lowered by high-02-pressure annealing. For the series of

samples containing Al, even if the samples were annealed at high02-partial pressure, the superconducting transition temperature

and the magnitude of resistivity in normal state were not much altered.

Sunshine et al.[8] reported that in YSr2Cu3_xMx02 (123

structure), Fe, Co or Al for M occupied Cu(l) sites located between two Sr sites along the c axis. Because of the similarity

396
of the 446 structure to the 123 structure, we anticipated that Fe,

Co or Al when doped into the 446 structure would occupy the Cu(l)

sites (Fig. 1). Then the oxygen occupancy for the 0(1) sites

which surrounded a Cu(l) site should be altered to maintain the

local electro-neutrality condition when Fe, Co or Al ions are substituted for Cu(l) ions. Therefore, if the valence of a Cu(l)

ions is assumed to be 2+, and the metallic ion content, x, in- creases by 1.0, the oxygen content, z, increases by 0.5 for M with

the valence of 3+, while z remain unchanged if the valence of M is

2+. From this point of view, "3+" is resulted for the valence of

a Fe ion (which is one of the most possible valence numbers for a Fe ion), because the gradient of the z-vs-x curve given in Fig.

3(a) is approximately equal to 0.5. Similarly, a valence between

2+ and 3+ is resulted for the valence of a Co ion (which is possible for a Co ion in a resonating valence state between 2+ and

3+), since the gradient of the z-vs-x curve given in Fig. 3(b) is

—0.2. Therefore, it is likely that the metallic ion (M) substitution for copper in the Ba-free 446 compounds,

(Eu3/4Ce2^/4)4(Euj^/3Sr2/3)4Cug_jjMx02 , does not alter the carrier

density for x up to 2.0 when M=Fe or Co. For each of the series

of samples with M=Fe and x up to 1.5 and those with M=Co and x up

to 2.0, the magnitudes of resistivity at room temperature are nearly identical. It seems that the carrier density for every

sample of each of the series is nearly constant regardless the

value of x. If so, we cannot employ the carrier density as the

crucial factor the appearance of superconductivity. On the con- trary, for the series of samples with M=A1 as shown in Fig. 3(c),

the oxygen content did not increase enough to offset the positive

charge due to the substitution of Al*^"*^ at the Cu(l) site. The excess positive charge would reduce the hole (carrier) concentration. Therefore, for the samples with 1VI=A1, it is likely that

the disappearance of superconductivity is due to the decrease in

the carrier density.

For the series of samples with M=Fe or Co, the disappearance of

the superconductivity might be caused by the magnetic interference

due to the substituted metallic ions, i.e., Fe or Co ions.

However, presently we have no experimental evidence to support

this speculation.

So far, two cases were distinguished for the disappearance of

superconductivity in 446 compounds with metallic ions which were partially substituted for Cu ions : (1) Disappearance of superconductivity in the 446 compounds with Fe or Co ions doped was hardly

related to the change in the carrier density. (2) In the 446 com- pounds with Al ions doped, superconductivity disappeared when the

carrier (hole) density was reduced. Case (2) corresponds to

the mechanism proposed by Tokura et al.[9] Case (1) seems to

remain puzzled. There have been reported experimental data that

the a-axis dimension which represents the Cu-0 bond length on the

Cu-0 plane is crucially important for the appearance of supercon397
ductivity in such systems as (La]^_j^Mx)2Cu04 (M=Ba, Sr or Ca)[10]

and (Pb,Cu)Sr2(Y,Ca)Cu202[ll] . Actually, among the samples with

M=Fe or Co, superconductivity showed only for those which had the

a-axis length shorter than ~3.84 %, and among the samples with

M=A1, superconductivity was observed for those whose a-axis length

was shorter than 3.85 %. From these observations, it may be said

that for the present 446 compounds containing Fe, Co and Al, su- perconductivity showed when the a-axis length was shorter than a threshold value around 3.845 X.

IV. CONCLUSION

We successfully synthesized Ba-free 446 [ (Eu,Ce)4(Eu,Sr)4Cug02-

type] samples by partially substituting Fe, Co or Al for Cu. It was demonstrated that without substitution of metallic ions such

as Fe, Co and Al ions for Cu ions , Ba-free 446 structure was not

formed. The a axis increased and the c axis decreased with increasing X for (Eu3/4Ce]^/4)4(Euw3Sr2/3)4Cug_jfM^02. The samples

with x=0.5 and M=Fe and Co exhibited superconductivity. The su- perconductivity transition temperature of these samples were

raised when the samples were annealed at a high oxygen pressure of

200 atm. The superconductivity of the sample with x=0.5 and M=A1 was significantly improved by annealing at high oxygen partial

pressure. For the samples with M=A1, the disappearance of superconductivity was explained by the reduction in the carrier density. For the samples with M=Fe and Co, the disappearance of

superconductivity was not related to a change in the carrier density which rather remained constant. For all the samples studied

in this work, it seemed that there might be a threshold value for

the dimension of the a axis : only samples with the a axis shorter

than the threshold value (3.845 K) exhibited superconductivity.

ACKNOWLEDGMENT

We are indebted Dr. S. Tajima of SRL for her helpful discussions .

REFERENCE

[I] H. Sawa. K. Obara, J. Akiiitsu. Y. Matsui and S. Horiuchi, J. Phys. Soc. Jpn.. 58. 2252(1989).

[2] A. Ichinose, T. Wada. Y. Yaegashi. H. Yawuchi and S. Tanaka, Jpn. J. Appl. Phys.. 28. L1765(1989).

[3] T. Wada. A. Ichinose. Y. Yaegashi. H. Yawuchi and S. Tanaka, Phys. Rev. B37. 5912(1989).

[4] T. Wada. A. Ichinose. Y. Yaegashi. H. Yawuchi and S. Tanaka. Jpn. J. Appl. Phys.. 29. L266(1990).

[5] K. Kurusu. R Takaai and K. Shintoni, Analyst. 114. 1341(1989).

[6] I*. Kosuge and K. Kurusu, Jpn. J. Appl. Phys.. 28. 1810(1989).

[7] R. D. Shannon. Acta Crystal logr. A32. 751(1976).

[8] S. A. Sunshine. L. F. Sct»iee»eyer. T. Siegrist D. C. Douglass. J. V. Waszczak. R. J. Cava. E. M. Gyorgy

and D. W. Murphy. Chei. Mater.. 1. 331(1989).

[9] Y. Tokura. J. B. Torrance, T. C. Huang and A. I. Nazzal, Phys. Rev. B38, 7156(1988).

[10] K. Kishio. K. Kitazawa, N. Sugii. S. Kanbe. K. Fueki, K Takagi and S. Tanaka. Chem. Lett, 635(1987).

[II] T. Maeda, K. Sakuyaaa. S. Koriyana, H. Yaaauchi and S. Tanaka. to be published.

398
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

EXCESS OXYGEN DEFECTS IN LAYERED CUPRATES

P. Lightfoot""-, Shiyou Pei-*-, J. D. Jorgensen^, X.-X.Tang"^,

3 3 . A. Manthiram and J. B. Goodenough , (1) Science and

Technology Center for Superconductivity and (2)

Materials Science Division, Argonne National Laboratory,

Argonne, IL 60439 and (3) Center for Materials Science

and Engineering, University of Texas at Austin, Austin,

TX 78712.

ABSTRACT

Neutron powder diffraction has been used to study

the oxygen defect chemistry of two non-superconducting

layered cuprates, La^ 25Dyo . 75^^*-^3 . 75^0 . 5, having a T*-

related structure, and La^ 85^^! . 15^^206 . 25 ^ having a structure related to the that of the newly discovered

double-layer superconductor La2_xSrxCaCu20g . The role

played by oxygen defects in determining the

superconducting properties of layered cuprates is

discussed

INTRODUCTION

Defects are known to play a critical role in the

synthesis and superconducting properties of the high T^

oxide superconductors (1) . The unique role played by

excess oxygen doping has been elegantly demonstrated in

the La2Cu04+5 system by Jorgensen et al . (2), who showed

that superconductivity in this system in fact results

from a subtle phase separation into an oxygen-rich

superconducting phase and a stoichiometric nonsuperconducting phase, both having orthorhombic

structures based on the K2NiF4, or T structure. Due to

the relatively small defect concentration, it was not

possible to locate the excess oxygen defect in the

neutron powder diffraction study of La2Cu04+5. However,

a later study of the isostructural La2Ni04+y (y ~ 0.18)

unambiguously determined the excess oxygen to be located

in an interstitial site in the La202 rock-salt like

layers of the structure (3) . Two structures closely related to the T structure

are the T' and T* structures, differing only in the

positions of the oxygen atoms within the lattice.

399
leading to square planar and square pyramidally

coordinated copper, respectively. In view of the

similarities between these structures, it may be

expected that similar interstitial oxygen defects could

also occur in these systems, and may indeed play a critical role in determining the electronic behaviour

of these systems, both of which have been made

superconductive (4,5) . In this paper we describe the

observation, by neutron powder diffraction, of such a defect in the T* phase L^i . 25^70 . 75CUO3 75F0 . 5 (6) . We

also describe the identification of the excess oxygen

location in the double-layer phase La^ ^ ss^^i . 15*^^206 . 25

(7), and discuss this in the light of the related,

recently discovered (8) superconducting phase

Lai.6Sro.4CaCu206.

EXPERIMENTAL

A powder sample of La^ _ 25Dy 0 . 75*-'^*-'4 obtained

from stoichiometric quantities of La203, Dy203 and CuO,

fired at 950°C for 12h and then at 1050°C for a further

20h. Fluorination was achieved by a low-temperature

anion exchange reaction, by reacting La^ ^S^Yo .75*-^^04 with

ZnF2 in an evacuated sealed tube at 280°C for 48h.

1 . 8 1^ 1 . 1 5CU2O g _ 2 5 obtained by firing

intimately mixed La203, SrC03 and CuO in 1 atm O2

according to the following scheme: 1000°C (20h) , 1050°C

(20h), 500°C (12h), 450°C (12h), 400°C (12h), 350°C

(12h) and 300°C (12h) . This procedure was found to

maximise the oxygen content for a 1 atm O2 anneal.

Neutron powder diffraction data were collected on the

Special Environment Powder Dif f ractometer at Argonne '

s

Intense Pulsed Neutron Source. Data were analysed by the

Rietveld method.

RESULTS AND DISCUSSION

Laj _ 25^yo . 75^"^3. 75^0 . 5 • ^ur refinement (6) shows

that this phase adopts an orthorhombically distorted

structure based on that of the T* phases (9) . Incorporation of the excess anion content occurs in an

interstitial site in the La202 rocksalt layers of the

400
(a) (b)

Figure 1. (a) Idealized tetragonal T* structure, showing

alternating T' (top half) and T (bottom half) layers,

(b) The distorted T* structure of Lai,25Dyo.75Cu03,75Fo.5,

showing incorporation of the interstitial defect, 0(5).

Favorable coordination of 0(5) to four M(l) and four

displaced anion, 0(4), sites occurs. Note that only half

of the true unit cell (which is rotated by 45° around

the c-axis, relative to (a) ) is shown, for ease of

comparison of the two structures

.

401
structure, and not in the vacant "apical" position in

the T' -type half of the structure. This type of

interstitial defect is directly analogous to that

previously observed in La2Ni04 (-^^ that thought to

be the doping mechanism responsible for

superconductivity in La2Cu04+5 (2) .Accommodation of the

defect requires displacement of the four neighboring

apical oxygens on to new sites, in order to provide

favorable anion-anion contacts. A comparison of the

undistorted (tetragonal) L^i . 25'^yo . 75^^*-'4 the

distorted ^a^ _ 25^yo .75^^^3 .75^0 . 5, showing incorporation

of the anion defect, is presented in Figure 1.

Interestingly, high oxygen partial pressures are often

found necessary to make the T* phases superconducting

(10,11) . Hence, although the present phase does not

superconduct , it is not unreasonable that a similar type

of defect may again turn out to be the doping mechanism

responsible for superconducting behavior in the T*

phases

.

1 . 85^^l . 15^^2^6 . 25 '^^^ Structure adopted by this phase

(7) is an oxygen-deficient variant of the well-known

Sr3Ti207 structure (n=2 member of the Ruddlesden-Popper

series A0.nAB03) . A view of the structure is shown in

Figure 2. Previous refinements of the structures of

cuprates of this type, La2SrCu206 (12) and La^ _ gCa^ _ 1CU2O6

(13) , have suggested a zero, or very small occupancy of

the 0(3) site between neighboring CUO5 pyramidal layers.

For the present case, where a significantly higher

oxygen content has been obtained, we have shown

unambiguously that the inter-layer site can accommodate

a high concentration of excess oxygen. Indeed, our work

suggests transfer of a small, but significant amount of

the in-plane oxygen (01) to the inter-layer site. It is

interesting to compare this phase with the recently

discovered (8) superconducting system La2_xSrxCaCu20g

(maximum = 60K at x = 0.4) . Izumi et al . (13) showed

that in La^ _ gCa^ iCu20g, Ca has a strong preference for

the smaller, M(l) site between the CUO2 layers. Thus it

would be expected that in La2-xSrxCaCu20g the Ca

would also adopt the M(l) site, thus leading to a

relatively short (~3.3A) Cu-Cu distance, and inhibiting

the occupancy of the 0(3) site. This system, therefore.

402
Figure 2. Structure of Lai .gsSri .15CU2O6 .25 . Note that the

0(3) site, which is completely vacant in the case of

Lai . gCai . 1CU2O6, is partially occupied in the present

case .

403
easily maintains the pseudo-two-dimensional nature

believed to be required for superconductivity. In

^ ^ 1 . 8 5 ^ ^ 1 . 1 5 ^ 2 6 . 2 5 ' Cu-Cu distance is

correspondingly greater {3.756A), facilitating the

occupancy of the inter-layer 0(3) site, and thus leading

to the partial breakdown of the 2D nature of the

structure. Further careful structural work will show

whether such ideas may help to explain the absence of

superconductivity in L^i , eS'^^l . l5C^2'-*6 . 25 related

phases (14)

.

CONCLUSIONS

We have presented two examples of the way in which

oxygen defects play a critical role in the structural

chemistry of the layered cuprates, and suggest that such

defects may have a profound effect on the electronic

properties of these systems. Clearly, before any

definitive statements relating defect chemistry to

superconductivity may be made, more structural data will

be required. We have so far only examined the simplest

systems. Neutron powder diffraction will play a central

role in the elucidation of the defect structures of more

complex systems. However, it is also clear, for example

in the Nd2-xCexCu04 system (15, 16), that the differences

between superconducting and non-superconducting samples

is very subtle, and defect chemistry on a local scale

may provide the key to the understanding of these

systems in the future.

ACKNOWLEDGEMENTS

This work was supported by the NSF-Office of

Science and Technology Centers under contract #STC8809854 (PL, SP) , U.S. Department of Energy, BES, under

contract #W-31-109-ENG-38 (JDJ) and the R.A. Welch

Foundation, Houston, TX (XXT, AM, JBG)

.

REFERENCES

1. J. D. Jorgensen and D. G. Hinks, Neutron News, 1, 24

(1990) .

404
2. J. D. Jorgensen, B. Dabrowski, S. Pei, D. G. Hinks,

L. Soderholm, B. Morosin, J. E. Schirber, E. L.

Venturini and D. S. Ginley, Phys . Rev. B., 38 . 11337, (1988) . 3. J. D. Jorgensen, B. Dabrowski, S. Pei, D. R.

Richards and D. G. Hinks, Phys. Rev. B., 40, 2187

(1989) . 4. Y. Tokura, H. Takagi and S. Uchida, Nature, 337 . 345

(1989)

5. J. Akimitsu, S. Suzuki, M. Watanabe and H. Sawa,

Jpn. J. Appl. Phys., 22, L1859 (1988).

6. P. Lightfoot, S. Pei, J. D. Jorgensen, X.-X. Tang,

A. Manthiram and J. B. Goodenough, Physica C, 169 , 15 (1990) . 7. P. Lightfoot, S. Pei, J. D. Jorgensen, X.-X. Tang,

A. Manthiram and J. B. Goodenough, Physica C, in

press . 8. R. J. Cava, B. Batlogg, R. B. van Dover, J. J.

Krajewski, J. V. Waszczak, R. M. Fleming, W. F. Peck

Jr. L. W. Rupp Jr., P. Marsh, A. C. W. P. James and

L. F. Schneemeyer, Nature, 34 5 , 602 (1990)

9. H. Sawa, S. Suzuki, M. Watanabe, J. Akimitsu, H.

Matsubara, H. Watabe, S. Uchida, K. Kokusho, H.

Asano, F. Izumi and E. Takayama-Muromachi, Nature,

337 , 347 (1989) . 10. M. F. Hundley, J. D. Thompson, S. W. Cheong, Z.

Fisk, R. B. Schwartz and J. E. Schirber, Phys. Rev.

B. , 4^1, 5251 (1989) . 11. Z. Fisk, S. W. Cheong, J. D. Thompson, M. F.Hundley,

R. B. Schwartz, G. H. Kwei and J. E. Schirber,

Physica C, 162-164 . 1681 (1989)

12. V. Caignaert, N. Nguyen and B. Raveau, Mater. Res.

Bull. , 199 (1990) . 13. F. Izumi, E. Takayama-Muromachi, Y. Nakai and H.

Asano, Physica C, 157, 89 (1989)

14. J. B. Torrance, Y. Tokura, A. Nazzal and S. S. P.

Parkin, Phys. Rev. Lett., M, 542 (1988).

15. P. Lightfoot, D. R. Richards, B. Dabrowski, D. G.

Hinks, S. Pei, D. T. Marx, A. W. Mitchell, Y. Zheng

and J. D. Jorgensen, Physica C, 168 , 627 (1990) . 16. B. Dabrowski et al . (this meeting).

405

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

CAN Co(Fe) SUBSTITUENT IN YBa2Cu307.5 MIGRATE BACK

AND FORTH BETWEEN Cu(1) AND Cu(2) SITES?

AmarNath, Zoltan Homonnay, Guang-Way Jang, Sandor I. Nagy, Yen Wei

and Chi C. Ctian

Department of Chemistry, Drexel University, Philadelphia, PA 19104

Cobalt and iron substitute preponderantly at the Cu(1) site

especially if present in low concentrations C). In the oxygen-deficient 1-2-3

compound containing less than O6.4, the Cu2+ ions in the sheets order

antiferromagnetically (2). Cu on Cu(1) site is monovalent and non-magnetic

and therefore do not participate in the magnetic ordering. The neighboring

Cu2+ ions along the sheets are strongly coupled antiferromagnetically,

while there is weaker coupling between the sheets along the c-axis.

57Fe{57co) on Cu(1) sites do not sense the magnetic field as the chain

sites are sandwiched between the two neighboring sheets and they are

symmetrically disposed with respect to the spins aligned antiparallel.

Consequently, observation of a Zeeman split sextet in oxygen-deficient

1-2-3 compound is generally regarded as a signature for Mossbauer probe

being situated on the sheet (i.e. Cu(2)) site(3-i4) However, we had

mentioned earlier of the possibility that the Mossbauer probe be situated at

neither the Cu(1) nor Cu(2) site but occupy an interstitial position (12,13)

We created a large amount (-50 %) of the magnetic species by

thermal treatment of YBa2Cu3(57Co)07.5 (containing -30 ppm 57co

substituent) at 800 °C for 5 h under Ar flow. This species is created at the

expense of the square planar species A situated at Cu(1 ) site (Fig. 1 ). It manifests itself as a Zeeman split sextet (designated M) in the

deoxygenated compound, and as a doublet (M') in the oxygenated sample ~ the oxygen loading being carried out at a relatively low temperature of

450 °C. The doublet M' has Mossbauer parameters characteristic of ionic

Fe3+. The Mossbauer spectra change drastically after the thermal

treatment at 800 °C and followed by oxygen loading (Fig. 1). However, it may be mentioned that superconducting properties of this thermally treated

sample are identical to the original one, not having undergone any thermal

treatment. This was verified by non-resonant microwave absorption

studies. M' (or M) consitutes about 50 % of the total area of the spectra.

The other major component is the square planar species A (Fig.1 ). If M'

407
really represents 50 % of the 57Co being situated at the Cu(2) site, then it is a great blessing. We can probe both copper sites simultaneously for

varying oxygen contents, and increasing amounts of substituents like Fe,

Co, Ni and Zn in the fully oxygenated state. These investigations could

provide very valuable information regarding the role of Cu-0 chains

vis-a-vis the Cu-0 sheets, especially with respect to the charge transfer

between the two sites. This apparent migration of Co or Fe substituent from

Cu(1) to Cu(2) site during thermal treatment at high temperatures in an

inert ambient was originally proposed by Takayama-Muromachi et al. C^)

and later observed by Mossbauer researchers(3.i2-i4) vve report here that

a thermal treatment at 700 °C under oxygen flow converts the M' doublet to

species A mainly (Fig. 1).

Figure 1. Emission Mossbauer spectra of YBa2Cu3(57co)O6.87±0.02 (obtained

by heating fully oxygenated material at 410 °C under Ar flow to convert 5-fold coordinated

species to the 4-fold one) with differing thermal history, (a) The sample was subjected to thermal treatment with O2 flow at 950 °C and then oxygenated at 450 °C; (b) thermal treatment of sample (a) with Ar flow at 800 °C (to generate the magnetic species) and then

oxygenation at 450 °C; (c) thermal treatment of sample (b) with O2 flow at 700 °C for -40 h

followed by oxygenation at 450 °C.

-2 -1 O 1 VELOCITY (mm/s)

2

408
So the question is whether Fe and Co substituents can really

migrate back and forth between Cu(1) and Cu{2) sites with the help of

different thermal treatments. Takayama-Muromachi and co-workersC^)

reason that the chains would lose all their oxygen above 750 °C in an inert

ambient, and that Fe originally present in the chains would seek the

oxygen rich environment (5-fold coordination) of the Cu(2) site.

Our detailed investigations do not bear out this scenario. On the

other hand, our observations discussed below, lead us to the conclusion

that species M(M') is situated at an interstitial site and not at the Cu(2) site.

Firstly, the species M' (doublet) or M (sextet) hardly shows any

change in isomer shift on deoxygenation of the 1-2-3 compound. It is generally believed that during deoxygenation, electrons are transferred

from the chains to the sheets, thereby, weakening the hole

superconductivity in the sheets. The probe site is apparently oblivious to the charge transfer between the two Cu sites. If one assumes that the holes

are primarily localized on the oxygen atoms and that the oxygen band is the one responsible for the metallic and superconducting behaviorC^),

then the microprobe on Cu(2) site would not sense any charge transfers,

and neither would it sense the transition from metallic to insulating state.

However, recent photoemission studies indicate that there is significant Cu

3d character in the density of states at the Fermi level suggesting strong O

2p and Cu 3d hybridization^"^). Therefore, the absence of changes in the

isomer shift of species M'(M), in contrast with that of species A at Cu(1 ) site

where significant changes are observed (Fig. 2), strongly suggest an

interstitial site for M'(M).

Secondly, thermal treatment at 700 °C in presence of O2 converts

species M'(M) into species A which is situated at Cu(1 ) site. Under these

conditions, Fe and Co can form only 4-fold coordinated species at Cu(1)

site with the limited amount of oxygen available. The question arises as to why the substituent if situated at Cu(2) site would want to forgo the 5-fold

coordination site and migrate to a 4-fold coordination Cu(1 ) site. On the

other hand, if species M(M') constitute interstitials, they can diffuse to nearby Cu(1) site and form 4-fold coordinated species A, which partly

converts to 5-coordinated species during cooling in O2 ambient.

Thirdly, the outermost peaks of the sextet are quite broad and their

widths decrease as one proceeds inwards. This is a signature for

distribution of hyperfine magnetic fields, which would be expected of

interstitial occupancy.

The following scenario emerges from the above observations. The

dopant 5^Co or ^^Fe with 4 or 5 oxygen neighbors in oxygen rich

microdomains at the chain site, when subjected to further deoxygenation at temperatures > 700 °C under Ar flow would tend to break loose, as it can

not form stable 2-coordinate linkages with oxygen, unlike Cu. When the

409
substituent is present in macroscopic amounts e.g. 1 % or more (as is generally the case for 57Fe substitution, as absorption Mossbauer

spectroscopy is much less sensitive than emission Mossbauer

spectroscopy, where 20 to 30 ppm substitution of 57co may be adequate),

clustering occurs and 5-coordinate linkages are formed which are more

reluctant to part with the bound oxygen and resist dissociation. The

dissociated Co or Fe ion migrates to a nearby potential well and is trapped

there during cooling. The interstitial 57co2+ (57Fe3+) is situated on a site

which can not be equidistant from the two neighboring sheets and would

sense a net magnetic field in an oxygen-deficient sample. It would also

probe the Neel temperature characteristic of the lattice, viz. 415±5 K for

Oq q stoichiometry, which is in harmony with the one determined by

neutron diffraction(2). The interstitials acquire some mobility above 600 °C

and in the presence of oxygen would tend to settle down at nearby Cu(1

)

site as 4-coordinated species.

A detailed account of this work will be published elsewhere.

I I

I 1 I I I I 150 250 350

Temperature (K)

Figure 2. MGssbauer isomer shifts (relative to a-Fe) for cobalt species A and M'(M)

in YBa2Cu3(57Co)06.27 (open) and YBa2Cu3(57co)06.87 (solid) as a function of tennperature. Shaded areas are a guide to the magnitude of changes.

REFERENCES

(1) L H. Greene and B.G. Bagley in "Physical Properties of High

Temperature Superconductivity H",D. M. Ginsberg, edit., p. 509, World

Scientific, Singapore (1990)

410
2) J. M. Tranquada in "Earlier and Recent Aspects of Superconductivity",

J. G. Bednorz and K. A. Muller, eds. p. 422, Springer-Verlag (1990)

;3) Z. Q. Qiu, Y. W. Du, H. Tang and J. C. Walker, Journal of Magnetism &

Magnetic Materials 78, 359 (1 989)

4) R. A. Brand, Ch. Sauer, H. Lutgemeier, P. M. Meuffels and W. Zinn,

Proc. ICAME'89, Hyperfine Interactions ,in press.

;5) L Bottyan, R. A. Brand, J. Dengler, B. Molnar, D. L. Nagy, G. Ritter, and

J. Schober, in "Proc. XIV Zakopane School on Physics ", World

Scientific Publ. (1990).

[6) T. Shinjo and S. Nasu, in "Mechanism of High Temperature

Superconductivity", H. Kamimura & A. Oshiyama, eds., p. 166 (1989)

7) S. Nasu, M. Yoshida, Y. Oda, T. Kohara, T. Shinjo, K. Asayama, F.E.

Fujita, S. Katsuyama, Y. Ueda and K. Kosuge, J. Magn. Mag. Mat., in press

;8) L. J. Swartzendruber, L. H. Bennett, M. Z. Harford and M. Rubinstein, J. Supercond.^,2^9 (1988)

[9) Yu. T. Pavlyuhin, N. G. Hainovsky, Y. Y. Medikov and A. I. Rykov,

Pramana 31, L 445 (1988)

[10) K. Donnelly, J. M. D. Goey, S. Tomlinson and J. M. Greneche, Physica

C156, 579(1988)

[11) 1. Nowik, M. Kowitt, I. Reiner, and E. R. Bauminger, Phys, Rev. 0 38,

6677 (1988).

[12) A. Natli, S. Nagy, M Barsoum, S. D. Tyagi and Y. Wei, Solid State

Commun. 68, 181 (1988)

[13) Z. Homonnay, S. Nagy , G. W. Jang, Y. Wei, S. Tyagi and A. Nath,

Proc. IGAME'89, Hyperfine Interactions, in press

[14) S. Katsuyama, Y. Ueda and K. Kosuge, Physica C 165, 404 (1990)

[15) E. Takayama-Muromaclii, Y. Uchida and K. Kato, Jap. J. Appl. Phys.

26, L2087 (1987)

[16) "Earlier and Recent Aspects of Superconductivity", J. G. Bednorz and

K. A. Muller, eds., Springer-Verlag (1990)

[17) R.S. List, A.J. Arko, R.J. Bartlett, C.G. Olson, A.B. Yang, R. Liu, C. Gu,

B.W. Veal, J.Z. Liu, K. Vandervoort, A.P. Paulikas and J.G. Campuzano,

J. Magn. Mag. Mat. 81, 151 (1989)

411

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

ELECTROCRYSTALLIZATION IN THE Ba-K-Bi-O SYSTEM

Michael L Norton and Horng-Yi Tang

Department of Chemistry, University of Georgia, Athens, GA.

and

Judith Eglin

Chemistry Department, Michigan State University, East Lansing,

Michigan

ABSTRACT

A technique for the low temperature, isothermal, isopotential

electrochemical deposition of superconducting, highly crystalline

materials in the Ba-K-Bi-O system has been developed. The

current status of this method of producing large (2mm diameter),

well facetted crystals in this system is presented in detail.

INTRODUCTION

The Bismuth based system (Ba-M-Bi-O) (1,2), given sufficient

development, may provide higher temperature superconductors.

This conclusion is supported from two quite different approaches.

First, the electron-phonon interaction is so strong in BaBi03 that

the frozen phonon mode or commensurate charge density wave is stable to 800 K (3). The Tc of a system may be considered to have as its upper limit the temperature of this phase transition.

A second approach utilizes the relationship between average

electronegativity and Tc (4) to predict that the Tc of the 3-D

members of this system should approach 180 K (5). Few other

systems currently display this high probability of future rises in

Tc.

The Ba-K-Bi-O system displays several modes of instability.

At high temperatures, the potassium oxide component is volatile

under standard oxide ceramic synthesis conditions (6). At

intermediate temperatures, the potassium becomes insoluble in the

perovskite phase (7), and at lower temperatures, the oxygen

content becomes unstable (6). This system is thus not amenable

to high temperature crystal growth. The utilization of low

temperature synthetic techniques allows for the exploration of new

regions of the phase diagram of chemical systems, and allows the

discovery 9f materials not accessible via high temperature

techniques due to the presence of large activation energies

intrinsic to massively reconstructive phase transformations.

The molten salt electrocrystallization method has a number of

potential advantages to offer in the synthesis of novel conducting

materials. Conducting phases are preferentially synthesized in

electrocrystallization because electrochemical potential losses ( IR

drop) are minimal in such phases. Growth of non-conductors is self limited by a process commonly referred to as electrode

passivation. There is nearly instantaneous nucleation of the desired

phase under proper conditions. Because nucleation and growth

413
occurs only on the selected electrode area, there are only a

limited number of nucleation sites, leading to a reduced number of

nucleated crystals, in comparison to many crystallization

techniques. The relative supersaturation can be set electronically

(potentiostatically), and the reaction can be stopped or quenched

quickly. All crystalline product can be readily and quickly

extracted from the growth solution.

The anodic electrocrystallization technique is exceptionally

well suited for the preparation of large single crystals in the BaK-Bi-O system for at least two reasons. First, a successful flux

or solution based growth technique requires significant solubility

of the reagents to enable diffusion to provide significant mass

transport. The product phases (BKBO's) are not very soluble in

KOH, and therefore would cause recrystallization of BKBO to be a slow crystal growth process in this system. However, Ba(OH)2 and

Bi203 are relatively soluble in KOH. The electrochemical

oxidation of Bi from an average formal +3 to a >+4 oxidation state yields an insoluble product which grows on the substrate. The

large change in solubility accompanying the rather large change in

average formal Bi oxidation state may readily be understood in

terms of the associated change in Madelung energy. The second

most important feature impacting crystal growth in this system is the intrinsic necessity of oxygen or oxidizing agent transport to the growing crystal/solution interface. Such oxidative power is readily delivered electrically, but chemical oxidation requires

diffusion of oxygen across the ambient gas/condensed phase

interface. This results, in our experience, in formation of a solid

polycrystalline surface barrier film of the oxidized product phase,

which impedes further oxygen transport. In the electrochemical

system, oxidative power is delivered only to the desired

crystal/solution interface.

PREPARATION

Low temperature electrodeposition experiments are performed

in readily obtainable Teflon sample containers (Cole-Parmer).

Teflon is the only container material we have identified which

does not contaminate the melt with long term contact. Potentiostatic deposition is performed utilizing a three electrode, one compartment cell geometry. A high current synthetic

potentiostat (Bioanalytical Systems, SP-2) system was employed for

these studies. The working electrode (positive, or biased

anodically) can be platinum, gold, or silver. The deposited film

protects this electrode from anodic dissolution. A 1 mm diameter

silver wire electrode (Aesar, 99.9%) is most commonly used. The

counter electrode (negative or biased cathodically) used is a 1 mm

diameter platinum wire (Aesar, 99.99%), which is rapidly coated

with a thick polycrystalline layer of bismuth metal during use. The reference electrode, which is placed in close proximity to the

working electrode, appears to be relatively stable and is particularly appropriate to this system, consists of a polycrystalline Bismuth metal rod, cast to a diameter of 1 mm.

414
The stability of this reference electrode's potential is related to the p02 and to the Bismuth salt concentration in the solution.

This electrode system does not introduce impurities detectable by

EDX (energy dispersive X-ray analysis) into the product phase.

The temperature of the cell, its envelope, and a large heat

capacity metal thermal ballast cup holding the assembly is maintained using a controller (Omega) to drive a hotplate. In an

average deposition, 20 grams of KOH is placed in a 15 cc teflon

crucible and melted at the reduced melting point of 180 C. The

lowered melting point is presumed to result from the known high

water content of commercial KOH. After raising the temperature

to 200 C, the other components of the melt, 1.60 grams of

Ba(OH)2*8H20 and 2.37 grams of Bi203 are added with stirring to aid dissolution. A quiescent solution is maintained during

deposition. Maintenance of an inert atmosphere is necessary in

order to avoid "spontaneous" or air oxidation of the chemical

system. This "spontaneous " oxidation can be quite rapid at elevated temperatures. A quartz cell holder contains the cell,

supports the electrodes, and allows maintenance of a water vapor

saturated nitrogen cover gas flow over the cell. The constant

PH2O may influence the long term stability of the reference

electrode potential.

Upon sedimentation of excess, undissolved 61203, electrolysis

of the crystal clear solution is begun at an applied potential of

0.7 V vs the Bi reference electrode. Deposition begins

immediately, and is terminated by withdrawing the crystal laden

electrode from the melt after any desired period of time. The

useable time is limited by the depletion of Bi from the solution

through deposition of Bi metal (cathodic) and Bi oxides (anodic),

to periods on the order of 48 hours for the system described here.

Samples are air cooled, rinsed and ultrasonically cleaned in

distilled water, then allowed to air dry.

CHARACTERIZATION

Morphological observations are made using an inverted optical

metallograph (Leco).

Scanning electron microscopy and energy dispersive x-ray

analyses are performed using a Phillips model 505 SEM equipped

with a Tracor Northern 5500 data aquisition system and

semiquantitative analysis software. Multiple sites are probed with

25kV excitation and 100 nm spot size. Data collected over the

range 0 - 25 kV are collected. The M line of Ba, the K line of

K, and the L line of Ba are fit to yield atomic percentages of

these elements.

X-ray diffraction patterns are acquired using a Scintag XDS

2000 Theta/Theta X-ray Powder Diffraction system. Data were

taken for two-theta values in the range 15 - 60 degrees, and

analyzed on the systems Microvax. Profile analysis utilized a Rietveld Analysis program modified by Hill and Howard (8). Magnetommetry was performed using a SQUID Magnetometer

(S.H.E.).

415
RESULTS AND DISCUSSION

Optical Microscopy: Materials prepared with deposition times

of 6 hours have been observed to grow as dendrites along 111.

The nuclei and earliest grown crystals are metallic bronze in

color. The later grown crystals are bright blue in reflection. It could be expected that these colors reflect the electronic

properties of the material, which are governed by their K content. SEM/EDX: One of the larger crystals grown is shown in

figure 1.

lmm2 4.9kU 5,3 0E1 8030/04 BKBO7 0

Figure 1 Scanning Electron Micrograph of BKBO Crystal

EDX analysis shows that the bronze colored early crystalline

material displays a composition roughly corresponding to Ba gK 2Bi03. The later material, which composes the bulk of the

material displays the stoichiometry Ba 5K 4Bi03, the well known

superconducting composition. Although mapping of the composition

along the growth direction has not been performed, it is anticipated that there is a K concentration gradient along this

line.X-ray Diffraction: Analysis of the X-ray diffraction pattern

indicates significant broadening of the peaks, which are readily

indexed as cubic. The pattern is not readily refined, and indicates some spread in lattice constants, consistent with the EDX and

optical observations that the material is not homogeneous.

Magnetic Susceptibility: AC Squid measurements indicate that

the product does contain substantial ammounts of superconducting

material (ca 20%). The magnetic transition, displayed in Fig. 2,

416
(BaK)Bi03 zero field cooled

50G run 0.1209 grams

6 1

' 8 1^ ' Te""" ?o ' 25" Temperature (K)

Figure 2 Susceptibility Plot for BKBO Polycrystalline Mass

is not sharp, and the onset (ca 18 K) is far below transition

temperatures reported for powders of this material (30 K) (7). These findings support the conclusions that the crystalline product

is inhomogeneous, contains nonsuperconducting stoichiometries, and

likely is strained by the gradient in lattice parameter. The lack

of evidence for the highest Tc material is unexpected, and

difficult to explain.

DISCUSSION AND CONCLUSION

The electrodeposition reaction can be viewed as a Bismuth

oxide polymerization process resulting in a 3-D polymer stuffed

with cations. That is, a polymer crosslinked in two dimensions

perpendicular to the growth direction is formed, and incorporates

ions from the melt as growth progresses. The reason for utilizing

potentiostatic control of the electrode process rather than the

constant current deposition previously employed (9) was to maintain a constant rate of polymerization, and therefore to maintain a steady state doping rate. It would appear that, over

relatively long deposition times, this control breaks down. Solution

or local depletion of Ba content may be responsible for the dopant

profile. Higher temperature deposition experiments in larger

volumes of solution will lead to higher diffusion constants for Ba,

and less perturbation of solution Ba content during cell operation.

Sample homogenization via thermal annealing does not appear

practical in view of the thermal sensitivity of these materials.

Inhomogeneity appears to be a general problem in the Ba-K-Bi-O

system, and has been observed in crystals ranging from .3 mm to 5 microns in diameter (10).

As well as the graded concentrations inferred in these

materials, it may be possible to form block copolymers and dopant

superlattices by modifying the deposition rate " by controlling the

deposition potential.

417
Large regions (lOO's of microns) of the grown crystals do

appear to be homogeneous via EDX analysis, and small single

crystals grown quickly may be quite homogeneous. Therefore,

although this method will require further development to deliver

extraordinarily large, perfect crystals, the products readily

obtainable may be of great value for optical studies, which are highly surface sensitive, and temperature dependent X-ray

structural studies, which require reasonably small samples.

This work was partially supported by SURA. We acknowledge

the use and staff support of the Center for Ultrastructural Studies

at UGA, and X-ray diffraction studies performed by B. Chakomaukos at Oak Ridge National Laboratory.

REFERENCES:

1. A.W. Sleight, J.L. Gilson, and P.E. Bierstedt, Solid State Commun. 17, 27 (1975).

2. L.F. Mattheiss, E.M. Gyorgy, and D.W. Johnson,Jr., Phys. Rev.

B 37, 3745 (1988).

3. D.E. Cox and A.W. Sleight, Acta Cryst. B 35, 1 (1979).

4. D.A. Nepela and J.M. McKay, Physica C 158, 65 (1989).

5. D. Tseng and E. Ruckenstein, J. Mater. Res., 5, 742 (1990).

6. D.G. Hinks, B. Dabrowski, J.D. Jorgensen, A.W. Mitchell, D.R.

Richards, S. Pei, and D. Shi, Nature, 333, 836 (1988).

7. D.G. Hinks, J.D. Jorgensen, D.R. Richards, B. Dabrowski, Y.

Zheng and S. Pei, preprint, Mol. Cryst. Liq. Cryst. (1990).

8. R.J. Hill and C.J. Howard, program LHPMl, CSIRO Div. of

Mineral Chemistry, P.O. Box 124, Port Melbourne 3207.

9. M.L. Norton, Mat. Res. Bull., 24, 1391 (1989).

10. L.F. Schneemeyer, J.K. Thomas, T. Siegrist, B. Batlogg,

L.W. Rupp, R.L. Opila, R.J. Cava, and D.W. Murphy, Nature, 335,

421 (1988).

418
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

STRUCTURAL DIVERSITY IN OXYGEN-DEFICIENT PEROVSKITES

J.T. Vaughey and K.R. Poeppelmeier

Department of Chemistry and The Science and Technology Center for

Superconductivity, Northwestern University, Evanston, IL 60208

The discovery of high temperature superconductivity in several perovskiterelated copper oxides which exhibit variable oxygen stoichicmetry has initiated

renewed interest in the structures of oxygen deficient perovskites and related

compounds^'^. For example, it has been suggested that in the YBajCujOy.j system,

the reported plateau in the plot of oxygen content versus T^ at 60K is due to an

ordering of the oxygen vacancies^ Therefore the determination of possible

structures originating from the removal of oxygen and ordering of the anion vacancies

is essential. There are several known perovskite-related oxides where the ordering

of vacancies have been well determined by either X-ray single crystal or neutron

powder diffraction structure determinations. In this paper, we will compare the

structures of LaSrCuA105^ Ca2Mn205^ LaSrCuGaOj^ and YBaCuFeOj^. These

examples demonstrate that compounds which exhibit a large fraction (» 17%) of

anion vacancies adopt nonrandom vacancy distributions that reflect the coordination

chemistry of the B-cation(s). Each exhibits a unique way to remove one oxygen from

a double perovskite (AjBjO^). The synthesis and X-ray single crystal structure for

YBaCuFeOj are reported and compared with an earlier neutron powder diffraction

study^.

CRYSTAL STRUCTURE

Single crystals of the compound YBaCuFeOs were isolated as small black plates

from a flux of the general composition YBa4CuioFeOj that was slowly cooled (6°C/hr)

from 1000°C in an alumina crucible. Energy Dispersive X-ray Analysis (EDX)

confirmed the presence of all four cations in an approximate 1:1:1:1 ratio. No

contamination from the alumina crucible could be detected. The unit cell was

determined from 25 independent reflections collected on a Enraf-Nonius CAD4

diffractometer collected at -120°C. The cell was determined to be tetragonal with a = 3.865(1) A, and c = 7.6425(8) A. There was no evidence from Weissenberg

photos of a supercell in agreement with the unit cell determined by the

diffractometer. From analysis of the data collected (h k 1 and ¥ Ic T), the compound

419
was found to be acentric with no observable systematic absences. These reflection

conditions lead to five possible space groups. Structural solution was undertaken

using the X-ray crystallography package TEXSAN*. The final space group chosen was P4mm (#99) based on the refinement from a double perovskite model where the

Y^^/Ba+^ and Fe^VCu^^ could be ordered in alternate AO and BOj layers,

respectively. To determine the position of the oxygen vacancy, each B-cation was

assumed to be six coordinate and the occupancies of each oxygen were allowed to

vary. The vacant oxygen sites occur at (1/2, 1/2, 1/2), which results in square

pyramidal coordination of both the copper and iron sites. Refined positional

parameters and temperature factors are in Table 1. The final R-factor was 0.092, R^ = 0.103. An ORTEP of the structure is shown in figure 1.

TABLE 1, Positional and temperature factors for YBaCuFeOs

Atom X y z B(A^)

Ba 0 0 0 [0.0226]* 0.61(3) [0.82]

Y 0 0 0.4910(8) [0.5112] 0.43(3) [0.58]

Cu 0.5 0.5 0.7157(7) [0.2738] 0.68(7) [0.47]

Fe 0.5 0.5 0.254(1) [0.7387] 0.26(7) [0.86]

0(1) 0.5 0.5 0.990(5) [0.0149] 0.9(2) [1.68]

0(2) 0.5 0 0.301(2) [0.3331] 0.3(2) [1.09]

0(3) 0.5 0 0.669(3) [0.7028] 0.5(2) [0.89]

( * = brackets indicate neutron diffraction study^)

The refined structure is similar to that reported by Er-Rakho, Michel, Lacorre,

and Raveau when the placement of barium at the origin is taken into account.'

Although copper and iron could not be distinguished based on X-ray diffraction

methods, the copper was placed on the z — 0.75 site because of a more distorted

coordination on that position consistent with the Jahn-Teller d' copper cation. In

contrast, in the neutron diffraction study, iron was placed predominantly on the z =

0.75 site (62%). The neutron scattering lengths for copper and iron are 0.772 x 10 *^ cm and 0.954 x 10 cm, respectively. Mossbauer spectroscopy showed iron was

present on both sites. The distorted coordination environment found for iron (-f-3) was consistent with other compounds, such as LaSrFeO/.

DISCUSSION

A wide variety of physical methods, including high resolution electron

microscopy, single crystal X-ray diffraction, and powder neutron diffraction have been

used to investigate oxygen deficient perovskites. This paper will concentrate on

420
421
compounds with an oxygen deficiency that has manifested itself in some well ordered

arrangement of oxygen vacancies and is still easily compared to the perovskite

structure. In general, the ordering of anion vacancies is due to the coordinational and

redox chemistry of the B-cation'°. One of the earliest discovered and most studied

members of this class of compounds is the mineral brownmillerite, Ca2FeA105

(ABOj.^; X = 0.5)'^ It has been found that the aluminum and iron unequally mix

over the tetrahedral and octahedral sites, with aluminum having a preference for the

tetrahedral site^^. In oxygen-deficient perovskite compounds, the structures of

possible compounds are limited by the coordination environment of the B-cation. The

four compounds Ca2Mn205, YBaCuFeOj, LaSrCuGaOj, and LaSrCuAlOj will be used

to illustrate the structural variation that has been observed in oxygen deficient

perovskites with the composition ABO25 (A2B2O5). Two involve only square

pyramidal coordination of the B-cation while two others involve tetrahedral and

octahedral coordination. Figure 2 shows the relationship of the four structures to be

compared to the parent perovskite structure. The crystallographic relationships of the

four compounds and perovskite are listed in table 2. TABLE 2. Relation of unit cells to perovskite

a (A) b(A) c(A) V(A^) V(A')

LaSrCuGaOj 16.383 (4ap) 5.533 (/2ap) 5.330 (/2ap) 483.1 60.4

LaSrCuAlOj 7.922 (2ap) 11.020 (2/2ap) 5.424 (/2ap) 473.5 59.2

YBaCuFeOj 3.865 (ap) 3.865 (ap) 7.643 (2ap) 114.2 57.1

Ca2Mn205 5.43 (/2ap) 10.24 (2/2ap) 3.74 (ap) 208.0 52.0

* = volume per one perovskite formula unit

Ca2Mn205 was first reported in 1982 as the end product of the reversible

oxygenation of CaMn03^. Under a reducing atmosphere the compound CaMn03.,

(0^6^ 0.5) was found to evolve oxygen and form a series of stable, oxygen deficient

compounds. Using powder neutron diffraction, the structure of CaMnO^s (or

Ca2Mn205) was determined to consist of square pyramidal manganese (d* - Mn , (t2g)' (e^)*). The structure of CaMnO^s is characterized by oxygen vacancies

generated by the reduction of Mn"*^ to Mn^^ that order as shown in figure 2. The

intermediate members of the series contain various ordering patterns of square

pyramids and octahedra'^

YBaCuFeOj was first reported in 1988^. One description of the structure is that the oxygen vacancy occurs in the AO layer to form square pyramids of copper

422

(d^ - Cu^^) and iron (d^ - Fe'^). YBaCuFeOj is of considerable interest because it can also be viewed as the first member of the structural family that includes

YBa2Cu307_j, the 93K superconductor^*. However, the structure of YBaCuFeOs

is closely related to Ca2Mn205. To illustrate the similarity of Ca2Mn205 and

YBaCuFeOj the oxygen vacancy should be viewed to be in the mixed (Fe,Cu)02

layer. In CaMnOis one-half of the oxygen atoms in alternate [110] rows are missing,

whereas in YBaCuFeOs, one-quarter of the oxygen atoms sites are vacant in every

[110] row. A reason for the different ordering patterns is the ionic radii of the Acations. Whereas yttrium (1.16A) and calcium (1.26A) are approximately the same

size and often substitute for one another, barium (1.56A)^^ is larger, and requires a higher coordination number. As predicted^, the coordination of the A and B cations

together determines the overall structure. In YBaCuFeOs, the yttrium site is 8-

coordinate and the barium site is 12-coordinate, in contrast to CaMnO^s, where the

calcium is 10-coordinate.

LaSrCuGaOj was recently discovered^ and shown to have the brownmillerite

structure with octahedral and tetrahedral coordination of the mixed B-cations,

respectively. The oxygen vacancies are ordered within the gallium - oxygen layer as shown in figure 2. This arrangement of vacancies is probably the most commonly

seen in oxygen deficient perovskites. A large number of compounds have been

isolated and studied with this general structure encompassing several late transition

metals (Co, Fe) and the group IIIB cations (Al, Ga, In)*^

.

LaSrCuAiPs is an another example"* where oxygen vacancies order to form

tetrahedrally coordinated aluminum. However compared with LaSrCuGaOj, the

vacancies are in the same location layer to layer with respect to the copper plane. It

is closely related to the brownmillerite structure, but as is evident from table 2 and

figure 2, there is a doubling of the a-axis, instead of the quadrupling seen in the

LaSrCuGaOs structure. CONCLUSION

The structures of four distinct variations that result from the removal of one

oxygen from the double perovskite formulation A2B2O5 (i.e. A2B2O5) have been

discussed. The importance of the electronic structure, the coordinational preferences

of the B-cation, and the coordinational number of the A-cation have been shown to be the factors in understanding the arrangement of oxygen vacancies that each

structure has adopted. These results present a clear picture of the possible structural

variatioas, and provide the basis for understanding structure/property relationships in areas as diverse as catalysis and superconductivity.

ACKNOWLEDGMENTS

Support for this research was from the Science and Technology Center for

Superconductivity (NSF DMR-8809854).

424
REFERENCES

1. Bednorz, J. G.; Miiller, K. A. Z. Phys. 1986, B64, 189.

2. (a) Nguyen, N.; Choisnet, J.; Hervieu, M.; Raveau, B.J. Solid State Chem.l9»l, 39,

120. (b) Goodenough, J.B.; Demazeau, G.; Pouchard, M.; Hagenmuller,P. /. Solid

State Chem. 1973, 8, 325.

3. (a) Renault, A.; Burdett, J.K.; Pouget, J.P. /. SoM State Chem. 1987, 77, 587.

(b) Lin, Y.P.; Greedan, J.E.; O'Reilly, A.H.; Reimers, J.N.; Stager, GV.; Post, M.L.

J.Solid State Chem. 1990, 84(2), 226.

4. Wiley, J.B.; Sabat, M.;Hwu, S.-J.; Poeppelmeier, K.R.; Reller, A.; Williams, T.

/. Solid State Chem. 1990, 88(1), 250.

5. Poeppelmeier, K.R.; Leonowicz, M.E.; Longo, J.M. J.Solid State Chem 1982, 44(1),

89.

6. Vaughey, J.T.; Wiley, J.B.; Poeppelmeier, K-R. Z.anorg.allg.Chem, 1990, submitted.

7. Er-Rakho, L.; Michel, C; Lacorre, P; Raveau, B. /. Solid State Chem. 1988, 73,

531.

8. TEXSAN"; Molecular Structure Corp.: College Station , TX 77843.

9. Soubeyroux, J.L.; Courbin, P.; Fournes, L.; Fruchart, D.; Le Flem, G.J. Solid State Chem. 1980, 31, 313.

10. (a) Poeppelmeier, K.R.; Leonowicz, M.E.; Scanlon, J.C.; Longo, J.M.; Yelon, W.B.

/. Solid State Chem. 1982, 45(1), 71. (b) Leonowicz, M.E.; Poeppelmeier, K.R.;

Longo, J.M. /. Solid State Chem. 1985, 59, 71.

11. Hansen, W.C.; Brownmiller, L.T. Amer. J Sci. 1928, 15, 224.

12. Smith, D.,Acta Cryst. 1962, 15, 1146.

13. (a)Chiang, C.K.; Ph.D. Thesis, Northwestern University, Evanston, IL, 1989.

(b)Reller, A; Thomas, J. M.; Jefferson, D.A Proc. R. Soc. London 1984, A394,

223.

14. Cava, R.J.; Sci. Amer., 1990, 263(2), 42.

15. Shannon, R.D.Acta Cryst. 1976, A32, 751.

425
16.(a) Schenck, V.R.; MuUer-Buschbaum, H. Z. anorg allg. Chem. 1973, 395, 280.

(b) Battle, P.D.; Gibb, T.C.; Nixon, S. /. Solid State Chem. 1988, 73, 330.

(c) Harder, R.; Muller-Buschbaum, H. Z. anorg allg. Chem. 1980, 464, 169.

(d) Colville, A.A; Geller, S.Acta Cryst., 1971, B27, 2311.

426
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

NEW LaCuOj^ PEROVSKITES PREPARED

AT HIGH OXYGEN PRESSURE

J. F. BRINGLEY, B. A. SCOTT, S. J. LA PLACA, R. F. BOEHME

M. W. McElFRESH, T. M. SHAW, S. S. TRAIL and D. E. COXf.

IBM Research Division, Thomas J. Watson Research Center, Yorktown

Heights, NY 10598. tNational Synchrotron Light Source,

Brookhaven National Laboratory, Upton, NY 11973.

INTRODUCTION

The transport and physical properties of the perovskite-related superconducting cuprates are known to be strongly correlated with hole concentration (p), i.e., the oxidation state of the Cu-O matrix (1, 2). Although

a large number of cuprates with intermediate hole concentrations, p = 0.0

-0.5 (Cu2+ - Cu^-^"*"), have been extensively characterized, relatively little work has been reported on materials with higher hole concentrations, p = 0.5 - 1.0. Although a number of such cuprates are known, in which

copper occurs formally as Cu^"*" (3, 4), most of these are structurally unrelated to the superconductors, and a direct comparison of their chemical

and physical properties may not be appropriate. The synthesis of a cuprate exhibiting a wide oxygen stoichiometry range, while maintaining

the basic perovskite structure, would therefore be of great value for elucidating the nature of Cu/O hole states, determining how magnetic and

transport properties change with carrier concentration, and for establishing a complete set of "standards" for the spectroscopy of Cu^"^, Cu^"^ and

intermediate valence states. In this paper, we report the synthesis, structure and preliminary characterization of a new oxygen-defect perovskite,

LaCuOj.^, related directly to the superconducting cuprates. Although

superconductivity has not been observed in this system, its oxygen

stoichiometry range, 0.0 < ^ < 0.5, is the widest yet found in a cuprate,

allowing the formal oxidation state of copper to be varied from +2 to +3

within the same basic perovskite framework. As a result of these features

and of those described below, this may represent a model system with

which to address the above questions and to probe the subtle interplay of

427
the crystal-chemical and physical properties of the superconducting and

related cuprates.

The perovskite LaCuOi can be considered the fully-oxidized, threedimensional prototype of the layered high-T^, materials. This compound

was first prepared by Demazeau et al. at 900 °C and 65 kbar using

KCIO3 as an oxygen buffer (5). Its x-ray powder diffraction pattern was

indexed on the basis of a rhombohedrally-distorted perovskite cell and it was reported to be fully stoichiometric and to display metallic properties

down to 77 K. Because of the very high pressure and other complexities

of its synthesis, however, further characterization of this important phase

has not been forthcoming. In examining whether LaCu03 can be prepared

under milder synthetic conditions, we have discovered a new series of ox- ygen defect LaCuOj.^ perovskites, with 0 < 8 < 0.5.

EXPERIMENTAL:

High-pressure oxygen reactions were performed in an externally heated

Rene alloy vessel pressurized at or below room-temperature with highpurity gas. Reaction temperatures were measured at a shallow recess in

the reactor wall at the sample position, and pressure measured with an

uncalibrated gauge. Phases in the LaCuOj.^ system were prepared at 900

°C at oxygen pressures from 0.2 - 1 kbar using coprecipitated La and Cu

hydroxide precursors. The precursors were decomposed by slow heating

to 650''C under O2 prior to reacting them in high pressure oxygen. Exclusion of CO2 and other organic impurities from the starting materials

and the oxidizing atmosphere was necessary in order to obtain singlephase, homogeneously oxidized materials. Compositions were established

by thermogravimetric analyses with an accuracy in 5 of + 0.02. Highresolution X-ray powder diffraction data for the various phases in the

LaCuOj.^ system were collected at the Brookhaven National Synchrotron

Light Source. Structural refinements were performed with a modified

Rietveld program.

RESULTS AND DISCUSSION:

Syntheses performed under oxygen pressures ranging from 0.4 - 1 kbar

yield a tetragonal phase with 0.0 <> 0 <, 0.2. Its structure is a simple

tetragonal distortion of a cubic perovskite with a = 3.81875(5) A and c = 3.97270(7) A at ^ = 0.05, space group P4/m. The tetragonal distortion

is brought about by a slight Jahn-Teller-like elongation of the CuO^

octahedra, in which there are four in-plane Cu-O bonds at 1.909 A, and

two apical Cu-O bonds at 1.986 A. This is in contrast to the

rhombohedral form in which all Cu-O distances are equal at 1.94 A (3).

The lanthanum cations are twelve-fold coordinated by oxide ions with

428
bond distances ranging from 2.70 -2.75 A. In addition, electron diffraction

data indicate a very weak and diffuse supercell with multiplicity 2c, however, this was not observed in the synchrotron data and is not included in

the structural refinement. The unit cell volume of the tetragonal phase is ^ 5% larger than that of the rhombohedral form (5), which suggests that

the latter is the high-pressure and the former the low-pressure polymorph

of LaCuOj.

The tetragonal form of LaCuOj.^ can be made oxygen deficient either

by lowering the synthetic pressure or by reducing it in 1 atm O2 or N2 at 400 - 600 °C. Below its homogeneity range, 0.0 <, d <, 0.2, the ordering

of oxygen vacancies results in the formation of a monoclinic perovskite

stable over the range 0.2 <, 5 <, 0.4. The structure of this phase is shown

in Fig. 1. The vacancies order along the pseudo-cubic {021} axes, giving

rise to a new a x ^Sa x s/5a cell with lattice parameters a = 3.83076(2)

A, b = 8.62884(5) and c = 8.65149(5) A for (5 = 0.33, space group

P2/m. The synchrotron diffraction peaks of this phase are extremely sharp

(about 0.025° FWHM) so that the monoclinic distortion from

orthorhombic symmetry can be readily distinguished, although it is very

small (/?= 90.214(4)°). The unit cell composition is LasCusOjjjj.

Fig. 1. Projection onto the (100) plane of the crystal structure

of La5Cu50j3+ , with the unit cell indicated.

La = large open circles, Cu = filled circles.

429
The structure of this phase closely resembles that of BaLa4Cu50j^ +

(6), although the latter is tetragonal, space group P4/m. It is comprised

of one-dimensional chains of CUO5 octahedra oriented along c which share

corners with four CUO5 square-pyramids in the (100) planes. The CUO5

pyramids share corners with four other square pyramids but are oriented

with their basal faces at 90° to one another. In contrast to BaLa4Cu50,3,

the structure is significantly distorted from tetragonal symmetry, and tne

Cu-O-Cu bonds adjoining octahedra and pyramids are bent considerably

below their ideal 180° bond angle.

Yet a third phase can be isolated in the LaCuO^.^ system by heating the

tetragonal form to 725 °C in N2. This perovskite is orthorhombic and exhibits a homogeneity range of 0.43 < b < 0.5. Rietveld refinement shows

that oxygen vacancy ordering in the (010) plane, which is complete at h = 0.5, creates a cell with lattice parameters related to those of the pseudo

cubic perovskite by a ci; 2-j2a^, b a and c -J^a^. Cell parameters

obtained from the refinement are a = 10.4782(2) A, ^ = 3.87956(7) A and

c = 5.5491(1) A for a sample of composition LaCu02 53. The orthorombic

perovskite, whose structure shown in Fig. 2 is isostructural with

CaMnOjs (7) and is entirely composed of corner-shared CUO5 pyramids

which form pairs by sharing oxygen vacancies in the (010) planes. The

structure can be derived from a pseudo-cubic perovskite by placing an oxygen vacancy on the face of the cubic cell and rotating the cube 90°

around its four-fold axis, thereby generating all possible pyramid orientations with a repeat of 4a.

Fig. 2. Projection onto the (010) plane of the crystal structure

of LaCuOjs- The unit cell is drawn, see note to Fig. 1.

430
A summary of the phase and crystallographic data for LaCuOj.^ are

collected in Fig. 3. The phase fields shown were obtained from samples

prepared directly, or modified by reduction of a ^ = 0.1 composition in

N2 or O2 (1 atm) at 300 - 750°C. Therefore, Fig. 3 is not an equilibrium

phase diagram. Indeed, we have found it is possible to'cHange the ranges

somewhat by varying the preparation or processing conditions. Depending

upon composition, three distinct ordered arrangements of oxygen (or vacancies) are observed. A tetragonal perovskite with a homogeneity range

0 ^ 5 <, 0.2 occurs at 900°C and Po =0.4-1 kbar. At lower oxygen

< ^ < 0.5 and appears to be an

pressures, Pq = 0.2 - 0.4 kbar, the degree of oxidation decreases leading

to the formation of a monoclinic phase stable over the interval 0.2 <> S <, 0.4. Both forms are metallic, with room temperature resistivities of 100

fiQcm and 300 fiQcm for S = 0.14 (tetragonal) and S = 0.33 (monoclinic),

respectively. Superconductivity has not been observed in either case down

to 4.2 K. Lastly, an orthorhombic perovskite can be prepared in this system by heating the tetragonal form to 700-800 °C in N2. This phase has

a homogeneity range between 0.43

insulator.

3.96 -

X

<

I o

UJX

<

I CQ

3.92 -

3.88

3.86 -

3.84

3.82 i3.80 -

0.0

8 .aCu03_8

Fig. 3. The variation of the perovskite subcell parameters with oxygen

content in LaCuOj.^. The subcell <2-axis of the orthorhombic phase (3.70

A) is not shown for clarity, 2<^ = approx. two phase region.

431
As oxygen is removed from the tetragonal S = 0 phase, two CUO5

pyramids and two formally "Cu^"'' sites" are created in the form of

square-pyramidal pairs per each vacancy inserted. The result is that Cu

is found only in octahedral and/or square-pyramidal coordination

throughout the homogeneity range of LaCuOj.^, and that fortuitously, the

number of holes (i.e., "Cu^"^ -sites") will be formally equal to the number

of octahedral coordination sites in the system. This aspect of the structure

means that the +3 and +2 valence states of Cu in LaCuOj c are "tagged," a feature which can be used to great advantage in unravelling the electronic structure of the cuprates. For example, this fact has been exploited

to gain insight into factors governing the "missing" electron paramagnetic

resonance in the superconducting cuprates (8, 9). A detailed study of the

effects of systematic valence changes on the electrical and magnetic properties in the LaCuOj.^ system will be reported elsewhere.

REFERENCES

1. M. W. Shafer, T. Penney and B. L. Olson, Phys. Rev. B. 36, 4047

(1987).

2. J. B. Torrance, Y. Tokura, A. I. Nazzal, A. Bezinge, T. C. Huang and

S. P. Parkin, Phys. Rev. Lett. 61 1127 (1988).

3. K. Hestermann and R. Hoppe, Z. Anor^. A//g. Chem. 367, 249

(1969).

4. N. ER. Brese, M. O'Keefe, R. B. Von Dreele and V. G. Young, ./. Sol.

St. Chem. 83, 1 (1989).

5. G. Demazeau, C. Parent, M. Pouchard, and P. Hagenmuller, Mat.

Res. Bull. 7, 913 (1972).

6. C. Michel, L. Er Rahko, M. Hervieu, .1. Pannetier and B. Raveau, ./. Sol. St. Chem. 68, 143 (1987).

7. K.R. Poeppelmeier, M.E. Leonowicz, J.C. Scanlon and J.M. Longo,

/. Solid State Chem. 45, 71 (1982).

8. F. Mehran and P. W. Anderson, Solid State Comm. 71, 29 (1989).

9. F. Mehran, J. F. Bringley, B. A. Scott and T. R. McGuire, submitted

to Nature.

432
SURFACE CHEMISTRY

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

SURFACE ENERGY BARRIER FORMED

BY ADSORBED OXYGEN IN POROUS ZnO

Satoru FUJITSU, Kunihito KOUMOTO" and Hiroaki YANAGIDA'"""' Department of Materials Science and Ceramics Technology,

Shonan Institute of Technology, 1-1-25 Tsujido-Mishikaigan,

Eujisawa, Kanagawa 251, Japan

Department of Industrial Chemistry, the University of Tokyo,

Z:z3-1 Hongo, Bunkyo-ku, Tokyo 113, Japan

Research Center for Advanced Science and Technology,

the University of Tokyo, 4-6-1 Komaba, Meguro-ku, Tokyo 154, Japan

The surface energy barrier formed by the chemisorbed oxygen

on porous ZnO has been measured. The height of the barrier de- creases by holding the sample at room temperature. This is due to

the change in the oxidation state of the chemisorbed oxygen. The

doping of CoO into porous ZnO enhances the saturated amount of

cheinisorbed oxygen and stabilizes the high temperature type of

oxidation state into the low temperature.

Introduction

It is known that ZnO based varistor is the typical functional

ceramics using the energy barrier along the grain boundary[l].

This characteristic depends on the atmosphere, especially oxygen,

during its fabrication process. When the inoterial is heated in

the reducing ambient, the large non-ohmic characteristic cannot be

observed, wliile this property is restored by in^c';.i,rg in air or oxygen

The authors have pointed out that the acceptor state to form

sui,!i an energy barrier was brought about by the excess ionized

oxygen along the grain boundaries[2-4] . Though such a state along

the grain boundary cannot be detected, the analogous state can be

recognized on the surface as the chemisorbed oxygen.

'Jhen oxygen atom or i.iolecule adsorbs on n-type semiconducting

cxica as ZnO, it traps a conduction electron from the bulk to

acquire a negative charge. The carrier density then decreases,

and, to compensate for the negative charges on the surface, the

electron depletion layer is generated simultaneously. This depletion layer acts as the conduction barrier. In the case of a

435
Depltion

layer

Neck

Eg

Ef

porous body (see Fig.l), the average

neck radius connecting the grains is

supposedly very small, and the effect of

a depletion layer on total conduction

behavior cannot be ignored[2, 5, 6] . It

is also considered that the energy

barrier along the grain neck formed by

penetrated oxygen contributes to control

the total conduction behavior.

The authors reported that the

electrical conductivity of porous ZnO

having 50% porosity was controlled by

the energy barrier formed by the

chemisorbed and penetrated oxygen at the

neck area. And the non-ohmic properties

were observed at rooivi teriiperature when

the sample was quenched after annealing

>400*'C in air[7]. These phenomena can

be explained by the proposed theoretical

model to estimate the surface energy

barrier formed by chemisorbed oxygen.

The effect of the penetrated oxygen has

not been clarified yet.

In this study, the formation of the surface energy

and its stability are exarnined from the viewpoint of the

sorption of oxygen and the change in its oxidation state. The

effect of CoO doping is also examined on the energy barrier which

is an important dopant to modify the varistor.

Fig.l Cross Section of

a connectiihj

porous ZnO

energy band.

neck

and

1 n a

i t s

barrier

chemiExperimental Procedure

(l)Saniple Preparation

Nominally pure ZnO powder was dispersed in an aqueous solution of AKNO^)^ (and Co( f'iO^)^ ) » and c'ri;-:c'. The resulting powder

was calcined at 600" C to decompose aluminum nitrate into 0.5mol%

Al^O^ oxide (and cobalt nitrate into 0.5ii!ol% CoO oxide). The

powder was then pressed into a rectangular bar in a iteel die and

sintered at 300^0 for 1 h in air. This type of Al ^O-^-doptc' sample

will be denoted as S-n. And Al^O^-CoO-dopcd sample will be called

S-Cc, Aluminum oxide was added to make the ZnO body porous. When

pure ZnO was sintered by the same procedure, the relative density

attcined >80/o. Of course, trace amount of Al^O^ dissolved into

ZnO cs a donor. The prepared sarnplc-s i'lad surface areas, of

3.2m /g(S-n) and 2.8m /g(S-Co) and 50% relative density (S-n and

S-Co). Ho CoG, Al^Oo and the other con;pounds than ZnO were

indicated by X-ray diffraction or EPMA. It is expected from the

repOi'ted phase diagrcin that CoO dissolved into ZnO completely

while considerable amount of Al^O^ precipitated in the grain

' cjr^iary.

436
(2) Non-ohinic Properties

After annealing at 500^C in air for 1 h, the samples were quenched into liq.N^. An In-Ga alloy was painted on each sample

to achieve ohmic contact. The current-voltage (I-V) and capacitance-voltage (C-V) characteristics were then estimated at 25 C

using a curve tracer (Kikusui 5802) at 50Hz and an LCZ meter (HP4274A) at lOOkHz, respectively. The samples were enclosed in the

quartz tube during heating and quenching. And they were put into

dry air during the electrical measurements.

(3) ESR measurement

The paramagnetic centers on the surface and in the bulk of

the sample after heating, quenching and holding were measured

using an ESR spectrometer at X-band f requency( JEOL JES-PE-3X),

with the pressure in the sample tube below ISOPa and a field

modulation of lOOkHz. The spectra were recorded at 20 C. The

intensity of the signal was corrected by that of Mn-niarker which

was inserted into the cavity with the sample tube.

Results and Discussion

(1 )Non-ohmic properties

The S-n and S-Co both quenched from 500 C in air showed

non-ohmic exponent, a, expressed as I = (V/C)°', of 2.8 and 6.2,

respectively. The energy barrier heights estimated by C-V characteristic were 4900 and 29500V/cm for S-n and S-Co, respectively.

At 500 ° C, the saturated amount of chemisorbed oxygen was

4.8ylSTP/cm'^ for S-n and that was 7.8ulSTP/cm2 for S-Co.

The height of the surface energy barrier i<t)^) can be ex- pressed as[8]

V(2eN,/)/£e,N^) (1)

where N^^ is the surface state density, which should be proportional to the amount of chemisorbed oxygen in the present case; Nj

is the carrier density; and the other symbols have their usual

meanings. Since the chemisorbed oxygen forms the electron trap

state on the surface, the high energy barrier is generated when

the large amount of oxygen chemisorbs on the sample.

The height of energy barrier and the saturated amount of

chemisorbed oxygen for S-Co were larger than those for S-n in

otlier temperature region. [7]

Figure 2 shows changes in a and current density for samples

held at 25''C in dry air: The non-ohmic exponent decreased and the

current density under the bias of 1 V/cm increased with increased

holding time. Since the I-V relation was non-ohmic, conductivity

could not be determined using Ohm's law. The non-ohmic exponent

of the S-n decreased to 1 within 100 h, whereas the value for S-Co

437
decreased more slowly. This

decrease in the non-ohm ic

exponent and increase in

current density are due to the

decrease in barrier height

predicted by the C-V relation.

Equation (1) indicates

that either increased carrier

density or decreased surface

state density vyith holding

time at room temperature

should explain decreases in

barrier height. The decrease

in the carrier density and the

desorption of oxygen were not

observed.

The decrease in the

barrier heic,ht was due to the

change in the oxidation state

of c[iemisorbed oxygen. At

high temperature (>400^C), the

oxidation state of the

chemisorbed oxygen can be

represented as 0 . According

to the 1 i terature[9] , 0 , 0~,

and Oo~ are stable above

300''C, at 150'' to 300''C, and

belov/ 150° C, respectively.

Even if the pmount of oxygen

c h e ill i s o r b e G at h i c h temperature is frozen by

quenching, its oxidation state

may become more stable, with

fewer electrons leading to a decrease in barrier height,

w !i G n the s a n! p 1 e is held at

room temperature, as expressed

by

0^"-^ 0"+e (2)

02-+(l/2)02(g) 202+e (3)

where e represents the

released conduction electron.

(2)ESR Measurement

Change in the oxidation

scate of the chemisorbed

-

O O O—

^

^^^-~A_

— • • • •

1

1

1 100 1000 10000

t (min)

Fig. 2 Change in the non-ohmic

sxponent( A , A ) and the currant

density under the bias(lV/cm)

(0.#) with 'loldiny time at 25 °C

in dry air. Open and closed

i!i a r i< s s h 0 \: the S - n and S - C o , respectively. The sample were

prepared by qiienciiing from 500°C

into liq.N2

10 100 1000 10000

t (min)

Fig. 3 Change in the spin density

resulting from Go" and 0~ for Sn(0) and S-Co(#).

438
oxygen were observed directly by ESR spectroscopy. The ESR

signals assigned to the chemi sorbed oxy^an (02~[10] and C~[ll]) were observed in both S-n and S-Co, and its intensity increased

with increased holding time at room temperature. To evaluate the

increased intensity of the signals, the spin concentration was calculated by double integration of the spectra. Since the

signals could not be clearly divided into 0~ and O^", only the

total spin concentration was estimated in the present calculation.

As shown in Fig. 3, the spin concentration increased with holding

time at room temperature, similar to the changes in the non-ohmic

exponent and current density in Fig. 2. This result indicates that

the decreased barrier height resulted from the reactions expressed

in Eqs. (2) and (3), in which trapped electrons were released,

hence decreasing the surface state density. The slow decrease in

S-Co energy barrier height with holding time at room temperature

apparently corresponds to the slow increase in spin concentration.

The considerable difference in ESR spectra for S-n and S-Co

is the signal at a g-value of 1.96. Though this signal can be observed in S-n after any treatment, that was never observed in SCo. The 1.96 line is attributed to mobile electrons in the con- duction band and/or a trapped hole in the shallow donor band,

which probably arises from defects such as interstitial zinc

and/or oxygen vacancies[ 1 2] . Since the difference of the electrical conductivity in

both samples was not as great as that in the

intensity of the ESR

signals of the two

samples with g-1.96, this

signal should be assifjned

not to the conduction

electron but to the

paramagnetic defects as oxygen vacancies, zinc

atoms at the interstitial

site trapping one hole.

The substituted aluminum

ion is a candidate of the

parcrnagnetic center, but

aluminum ion was dopsd in

both samples. Probably,

CoO Hoping forms the

donor state that is

undetectable by ESR

spectroscopy as doubly

ionized oxygen vacancies

and/or the zinc atom at

the interstitial site.

As shown in Fig.

Ou (b)

(c)

.1, 1 1 1 1 1

1 1 800 1300 1800 2300 2800 3300 3800 4300

H (gauss)

Fig. 4 ESR spectra for S-Co

(a) spectruin for ths sample evacuated

at 600*C for 10 h. No si.jnal was

o;-;served in this spectrum.

(b) spectrum for the sample annealed

at ^-"O^C for 1 h in 130Pa of oxygen

pressure and (c)spectrum for the

Su.M^le exposed in 1 30Pa of oxyyen

nressure at 25''C after evacuation.

439
(a), evacuated S-Co exhibited no signal by ESR measurement at room temperature. Oxidized S-Co showed ths signals assigned to

chemisorbed oxygen and also a broad band at 1500 gauss (g 4.25).

Since this broad band was observed in the sample oxidized at room temperature where the oxidation of CoO in grain bulk cannot be

expected, the formation of such a signal is resulted from the

change in the oxidation state of Co ion near the surface. The

enhancement of the saturated amount of chemisorbed oxygen and the

stability of 0 should be supported by the formation of complex

of Co ion and cheiiii sorbed oxygen. In the present experiment, it

is impossible to assign this unknown signal at 1500 gauss.

References

[I] K.Mukae et al., Jpn. J. Appl . Phys. , 16,1361(1977).

[2] S.Fujitsu et al., Bui 1 . Chem. Soc. Jpn. , 61,1979(1988).

[3] S.Fujitsu et al., Solid State Ionics., 32/33,482(1989).

[4] S.Fujitsu et al., J . Am. Ceram. Soc. , 70,0-71(1987).

[5] N.M.Beekmans, J. Chem. Soc. , Faraday Trans. 1, 74,31(1978).

[6] M.Takata et al., Yogyo Kyokaishi, 87,19(1979).

[7] S.Fujitsu et al., J. Am. Ceram. Soc. , to be published.

[8] S.R.Morrison, Advanced Catalysis, 7,259(1955).

[9] M.Sedaka et al., Bui 1 . Chem. Soc. Jpn. , 38,1414(1965).

[10]J.0.Cope et al., J. Chem. Soc. , Faraday Trans. 1, 69,1(1973).

[II] N. B.Wong et al., J. Chem. Phys. , 60,2148(1974).

[ 1 2]G. Neumann, "Paramagnetic Properties", p. 269 in Current Topics

in Material Science ed. by E.Kaldis 7,(1981), North-Holland

Amsterdam.

440
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

SURFACE PROPERTIES OF BaTiOj AT ELEVATED TEMPERATURES

J. Nowotny^ and M. Sloma^

(1) Australian Nuclear Science & Technology Organisation

Advanced Materials, Menai , NSW 2234, Australia

(2) Max-Planck-Institute for Solid State Research

7000 Suttgart 80, FR Germany

ABSTRACT

Surface electrical properties of undoped BaTiOj have

been studied by using the High Temperature Kelvin Probe.

Changes in work function have been measured during oxidation and reduction processes under controlled oxygen activity. Experimental data are considered in terms of

changes in nonstoichiometry of the BaTiOa surface layer

and structural reconstructions within this layer. Impact

of the observed phenomena on properties of ceramic materials based on BaTiOa is discussed.

INTRODUCTION

Segregation causes the chemical composition, crystalline ordering and electrical properties at interfaces,

such as external surfaces and grain boundaries, to differ

from those of the bulk. The effects of segregation have

been the subject of extensive studies for metals and alloys and have been confirmed for ceramic compounds such

as binary metal oxides and their solid solutions [1].

Thus it is expected that segregation at interfaces will

also affect the properties of ternary oxides such as

BaTiOs [2].

Several properties of BaTiOs, which are important for

application, such as dielectric properties and temperature characteristics of resistivity, are strongly influenced by interfaces. So far, however, little is known about

the effect of segregation on surface properties of BaTiOj

and its impact on interface properties.

The predominant lattice defects in BaTiOg are oxygen

vacancies [3,4]. Their bulk concentration is determined

by oxygen activity in the gas phase. The relationship

between oxygen partial pressure and oxygen nonstoichiometry of BaTiOg.^, in equilibrium state, is described in

several publications. Similar relationships may also be

expected for the surface layer. However, as a result of

segregation the concentration of defects in this layer
may be higher than that in the bulk. If this is the case,

the resulting enrichment may lead to the formation of

low dimensional surface structures.

Electrical properties are very sensitive to nonstoichiometry and related defect structure. The most frequently applied methods in studies of defect structure

are electrical conductivity and thermopower. Unfortunately, both these methods are bulk sensitive and show very

little, or no sensitivity to the surface. On the other

hand, the work function is a very surface sensitive parameter. It can be determined using a High Temperature Kelvin Probe [5] which enables studies of solids at elevated

temperatures and under controlled gas atmosphere.

The purpose of this work is to study surface properties of BaTiOa at elevated temperatures under controlled

oxygen activity by using work function measurements.

WORK FUNCTION AND DEFECT STRUCTURE

The work function (WF) of BaTiOa was measured by the

Kelvin method [5].

The Kelvin method involves taking measurements of

the contact potential difference (CPD) between two

surfaces (1) and (2):

CPD = 1/e ((p,-(p,) (1)

where e is the elementary charge and (p is WF. Accordingly, WF changes may be expressed as:

A(Pi = e(CPD) + A(P2 (2)

Changes in the WF of semiconducting compound caused by

the variations in nonstoichiometry can be directly

related to the concentration of defects. For a p-type

semiconductor we have:

A(p = -kT In (N/[h-] (3)

where N is the density of states, [h ] is the concentration of electron holes, k is the Boltzmann constant and

T is absolute temperature. Work function changes can also

be related to the Fermi energy level of the surface

layer:

A(p = -Ep. ' (4) .

The relationship (3) is only valid for WF data taken at

elevated temperatures when an equilibrium between the gas

phase and the solid has been established. If this condi442
tion is satisfied, then differences in the WF can be

directly related to the defect structure of the boundary

layer

The defect structure of nonstoichiometric oxides is

determined by the oxygen activity of the gas phase. Then

changes in oxygen partial pressure results in variations

of defects concentration which can be considered either

as acceptor or donor centers. Usually, increases in piOz)

results in either increases in the concentration of

acceptors or decreases in the concentration of donors.

Both result in the decrease of and, consequently, in

the increase of WF assuming that the crystal ordering

remains unchanged.

Changes in the oxide nonstoichiometry and resulting

concentration of electronic defects within the p-type

composition range with pCOj) can be expressed by the

following relationship

[h ] = const pCOJ^/" (5)

where n is the parameter related to the valency of

defects. For doubly ionized oxygen vacancies, n is 6. If

predominant lattice defects are singly ionized oxygen

vacancies or acceptor-type impurities then, according to

the mass action law, in both cases the oxygen exponent is

equal 1/4

Combination of Eq. (3) and (5) results in the

following dependence between the exponent 1/n and WF:

1/n = (l/kT){a(p/ainp(0,)} (6)

Hence the parameter n is related to the surface layer.

It has been reported that the parameter n for undoped BaTiOs within the p-type region is equal to 4 [4].

This value is consistent with several defect models including: the intrinsic model, which is based on singly ionized oxygen vacancies, extrinsic model based on acceptor

type impurities and the cation nonstoichiometry model [ 4 ]

EXPERIMENTAL

The High Temperature Kelvin Probe [5] enabled the

performance of WF measurements on both single crystals

and polycrystalline specimens up to 1273 K under gas

atmosphere of controlled oxygen activity. The measurements involve the determination of the CPD between the

studied specimen and Pt reference electrode. The oxygen

partial pressure in the gas phase was imposed by using a

Ar-Oj mixture. Its activity was determined by using a

zirconia gauge.

443
BaTiOa specimens were prepared by coprecipitation of

Ba and TiO oxalates. Details of the preparation were

described in ref. [6].

WF measurements were performed on specimens initially standardized at 1123 K in argon of oxygen activity

p(02)=10^ Pa and then cooled down to the temperature of

the measurements. The applied experimental procedure

involved the isothermal determination of WF changes as a

function of time after pCOa) was suddenly changed to a new

value (6.9x10^ Pa). The measurements were performed within

the range 473-1019 K.

RESULTS

Figs. 1 and 2 illustrate the WF changes while

cooling (first part of figures) and WF changes during

successive oxidation and reduction runs (second part of

figures) at 3 25 and 590°C, respectively. The WF changes

on cooling are caused by several processes including the

chemisorption of oxygen and surface nonstoichiometry . The

sign of measured A(p values for the two runs at both

temperatures are different. At 325°C; during the oxidation

and reduction runs, A(p is positive and negative, respectively. At 590°C a reverse effect is observed.

DISCUSSION

Measured WF changes at 590°C (Fig. 2) are in accordance with theoretical predictions described above. One

may therefore presume that the oxidation run causes

oxygen incorporation and results in the formation of

donor centers such as oxygen vacancies

z

0 V + zh (7)

0 0

Corcondantly , the reduction run leads to evolvment of

oxygen and disappearance of the donors. It can therefore

be assumed that the equilibrium (7) is shifted into the

right and left by oxidation and reduction runs, respectively. The model is fully consistent with experimental

data illustrated in Fig. 2. In contrast, the WF data in

Fig. 1 suggest, that surface electrical properties at

3 25°C are controlled by some other surface process. These

data can be explained in terms of crystalline reordering

induced by changes in oxygen activity. One possible

explanation involves the tetragonal/cubic (t/c) transition. In air bulk BaTiOa transforms at 120°C. However, it

has been reported that the grain boundary layer of BaTiOj

exhibits the cubic structure at room temperature [7]. It

is not known at what temperature the t/c transition may

take within the boundary layer but one may expect that

444
UJo

z

HI

q:

111 u.

u.o

o

Q.<o

o

BaTi O3 1 T=325'C

0 - P02= 1 87 X 10^ Pq 1

1 COOLING (eo'/h) 1

1

1 OXID(69«10^Pq)

-100

1 1

1 \ r

-200 1 V 1

1

1

1 RED(15 "lO^Pa)

' 1 1 1 1 1 1 800 600 AOO TEMPERATURE I'Cl

6 10 K 18 TIME IhJ

22

Fig. 1. CPD vs. T (cooling) and vs. time at 325°C

LUO

2

UJ

a:

UJ

u.

if

oQ.<

o

u

-ICQ

-200 -

BaTiOa T=590*C

P02=

I 1.87xlo2pa I "COOLING I (80*/h)

TIME [h]

TEMPERATURE ['C]

Fig. 2. CPD vs. T (cooling) and vs. time at 590°C

the Curie temperature for the interface layer depends on

the oxygen activity. Thus, the effect observed in Fig. 1

may be ascribed to the interfacial structural reordering

between the tetragonal structure to the cubic structure

stable at higher oxygen activity. The tetragonal/monoclinic reordering has also been taken into account [8].

There is a lack of experimental data which could

confirm the oxygen induced interface transition model

described above. Electron microscopy of BaTiOj standardized under controlled oxygen activity may be helpful

in this matter.

The sign of WF changes taken at 590°C is in agreement

445
with the model of nonstoichiometry for BaTiOj which involves oxygen vacancies as predominant defects. The value

of the parameter n determined here (n=4) can be considered along with several defect models mentioned above.

CONCLUSIONS

Work function studies indicate that the oxygen

pressure exponent for the surface layer at 590°C is essentially the same as that reported for the bulk of BaTiOj.

Its absolute value (1/n = 1/4) is in agreement with

several defect models considered for this compound.

WF measurements at 325°C are inconsistent with the

above defect models and indicate that changes in oxygen

activity result from structural reordering of the surface

layer. It is expected that the reordering involves the

t/c transition which is limited to the boundary layer. On

the basis of the results in this study, it is likely that

the structural reordering takes place when BaTiOa and

BaTiOj-base dielectric materials are cooled below 325°C.

Up until now, the interface properties of the dielectrics

have been considered in terms of oxygen chemisorption and

nonstoichiometry. The results in Fig. 1 indicate that the

structural evolution of interfaces during cooling also

must be taken into account. It is well known that the

crystalline structure at interfaces changes when the

density of defects exceeds some critical value. Thus it

is probable that the reordering which occurs at the

surface of BaTiOa is related to segregation.

AKCNOWLEDGEMENTS

This work has been supported by the European

Community within the BRITE/EURAM Project, Grant No BE3595/89. Thanks are due to Kath Smith for proofreading.

REFERENCES

[1] J. Nowotny, "Surface & Interfaces of Ceramic Materials", L.-C. Dufour et al., Eds., Kluwer Acad.,

Dordrecht, 1989, p. 205-239

[2] S. Desu, "Interfacial Segregation in Perovskites and

Applications to Boundary Layer Devices", PhD thesis,

University of Illinois, Urbana, 1982

[3] A.M.J.H. Seuter, Philips Res . Rep . Supl . , 3, 1-84

(1974)

[4] J. Nowotny and M. Rekas, J . Europ . Cer . Soc . , 5, 173-182

(1989)

[5] J. Nbwotny, M. Sloma and W. Weppner, J.Am.Ceram.Soc.

,

72 [4] 564-70 (1989)

[6] M. Rekas and M. Wierzbicka, Solid State Ionics 20,

51-54 (1986)

[7] J.-C. Niepce, Ref. [1], p. 521-533

[8] K. Smith, unpublished results.

446
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

A SCANNING TUNNELING MICROSCOPY STUDY OF SINGLE

CRYSTAL ZnO AND Ti02 SURFACES

Gregory S . Rohrer

Carnegie Mellon University

Dept. of Metallurgical Eng. and Materials Science

Pittsburgh PA 15213

and

Dawn A. Bonnell

University of Pennsylvania

Dept . of Materials Science and Engineering

Philadelphia PA 19104

ABSTRACT

The surfaces of ZnO and Ti02 single crystals have

been imaged in ult rahigh-vacuum with a scanning

tunneling microscope. Constant current images of the

polar ZnO faces show that single crystal surfaces

cleaned by vacuum heating have numerous steps consisting

mainly of (1010) and (1120) faces. Images of the

reduced TiO2(110) surface indicate that crystallographic

shear planes in the bulk strongly influence the surface

structure. Local variations in stoichiometry ranging

from TisOs to TigOn have been determined by directly

measuring the spacing between the periodic intersections

of adjacent shear planes with the TiO2(110) surface.

The results in this paper demonstrates that it is

possible to resolve atomic-scale features on the

surfaces of wide bandgap semiconductors using the

scanning tunneling microscope.

INTRODUCTION

The reactions at transition metal oxide surfaces

which underlie complex phenomena such as heterogeneous

catalysis [1] and gas sensing [2] are strongly

influenced by the geometric and electronic surface

structure. The application low energy electron

diffraction (LEED) [3] and high resolution transmission

electron microscopy [4] have provided the first insights

447
into the structures of these surfaces, but the current

information is still rather limited and many systems

have not yet been investigated. We have used a scanning

tunneling microscope (STM) , an instrument which has had

enormous impact on the study of metal and covalently

bonded semiconductor surfaces, to study the surfaces of

two transition metal oxides. From topographic maps of

these surfaces, we have been able to identify the types

of facets which form on the ZnO polar faces and

determine the influence of bulk, planar defects on the

atomic structure of the TiO2(110) surface.

EXPERIMENTAL DETAILS

The anisotropic etching of the ZnO (0001) and

(0001) faces was used to determine the polarity of the

single crystal. The polar faces were mechanically

polished until there were no optically visible defects

by sequentially reducing the size of abrasive alumina

grinding media down to 0.05 micron. After heating the

sample in ultrahigh-vacuum (UHV) at 600-700 °C for 2

hours, sharp LEED patterns were observed and Auger

electron spectroscopy indicated that the surfaces were

free of carbon contamination. Previous work has

indicated that such a treatment produces steps and

facets on the surface. [5] The titania surface,

prepared in a similar way, was first polished and then

reduced by heating in UHV at a temperature above 900 K

for 36 hrs . Following this treatment, the dark grey,

opaque crystal was argon ion milled and then annealed at

823 K for 30 min in 1 X 10""^ torr of O2 in order to reoxidize the surface and anneal out the damage produced

by ion bombardment . LEED patterns indicated that this

preparation method produced a well ordered, relatively

defect free surface. [6] The crystal was then exposed

to air for several days before being introduced into the

UHV STM chamber, where it was heated at approximately

673 K for 2 min and then 'flashed' at approximately 973

K for 5 sec. Following this treatment, a (1X1) LEED

pattern was found, identical to that observed following

the initial preparation.

The microscope consisted of a commercial STM head

[7] controlled by 'home made' feedback electronics and

software of conventional design. The images were

collected in the constant current mode using sample

biases of -1.5 to -3.0 V and tunnel currents between 0.1

and 1.0 nA and are presented after the subtraction of a

448
background plane which eliminates the tilt of the sample

that would otherwise overwhelm small variations and make

atomic features undetectable. Other experimental

details have been described previously. [8]

RESULTS AND DISCUSSION

The constant current image of the ZnO(OOOl) oxygen

terminated face, shown in Fig. 1, demonstrates that

surfaces prepared by in situ annealing have a high step

density. However, unannealed surfaces appeared even

more disordered, presumably due to polishing damage

Atomically smooth flat areas are separated from one

another by either single or multiple steps which are

aligned predominantly in the either the [1010] or [1120]

direction. This is reasonable, since these are the

nonpolar faces and are known to be thermodynamically

stable. [2] The height of the step which forms the

large hexagonal plateau in the lower left part of the

image is 5.21 A, which agrees well with the ideal

spacing between the oxygen planes along the [0001]

direction which is 5.25 A. Similar results were

obtained for the (0001) surface. Spatially resolved

tunneling spectroscopy has been used to detect mid-gap

defect states associated with these steps, a result

described in a more detailed paper. [9] In the majority

of the regions imaged, the atomically smooth areas were

only about 100 A in diameter. This high density of edge

sites is likely to affect the electronic properties and

reactivity of this surface.

No atomic-resolution images of the ZnO surfaces

have yet been recorded. However, we have examined only

intrinsic (not intentionally doped) crystals with low

electronic conductivities and the resolution of the

images was undoubtedly space charge limited. [10] This

was not the case for the reduced titania crystal which

we have also examined.

Fourier analysis and direct measurement of the

images of the TiO2(110) surface (see Fig. 2a) indicate

that the rows oriented in the [111] direction have an

8.5 A periodicity and a 1 A corrugation height. The

corrugation directed along the rows has a 3.4 A

periodicity and a 0.2 A amplitude. The width of these

rows seems to vary somewhat and there are also numerous

"missing cells". The bulk terminated rutile lattice

449
Fig. 1. Constant current image

of the ZnO(OOOl) face. Steps

are visible which form (1010)

and (1120) faces. The image is

771 A X 771 A ahd was acquired

at -1.5 V sample bias and 0.3

nA. The vertical resolution is

36 A from black to white.

{0010]

<

[0100]

[1000]

Fig. 2. a) 109 A X 109 A

image of titania (110)

.

acquired at -2.0 V and

0.1 nA. Vertical

resolution is 2 A from

black to white and the

rows are oriented in

the [ill] direction,

b) 273 A X 108 A image

of a different region

acquired using the

same conditions. The

vertical resolution is

9 A and the rows are

oriented in the [113]

direction

.

[113]

vectors have been added for comparison. Other images of

this area show ledges which occur approximately every

100 A and some regions of the images are either

disordered or covered with adsorbates . The image in

Fig. 2b, acquired on another area of the surface, shows

rows oriented along the [113] direction with an

approximately 20 A spacing. It is clear from these

observations that the surface structure can not be

450
described by a simple termination of the bulk unit cell

or some combination of unit cells.

The orientation and spacing of the rows on the

crystal surface suggest that they are caused by the

intersection of crystallographic shear (CS) planes with

the (110) surface. The (121) family of planes are among

those along which the CS displacements can occur. [11]

The (121) plane intersects the (110) surface along the

[111] direction, the same direction that the rows _^in

Fig. 2a are oriented. The intersection of the ^121)

plane and the (110) surface is in the direction [113],

the direction in which rows in Fig. 2b are aligned.

Because the CS displacement is always of the type

^/2[0Tl], adjacent sections of the crystal are

vertically displaced 1.5 A with respect to one another

along the line normal to the (110) surface. Because the

plane subtraction applied to each image removes any

global slope, a stepped structure appears saw-toothed,

so that the intersection of each CS plane with the

surface appears as a depression or dark spot on the

image. The spacing of the rows in Fig. 2a suggests that

every sixth anion (121) plane is a CS plane and that the

local composition in this area is TisOs, while the

spacing of the rows in Fig. 2b suggests that CS occurs

at every twelfth anion (1^1) plane and that the local

composition is TigOn. The 0.2 A corrugations within the

rows of Fig. 2a have a 3.4 A spacing which is very close

to the spacing of titanium atoms (3.5 A) along this same

direction in the ideal rutile unit cell.

It should also be noted that the presence of excess

titanium along the CS planes should result in an

electronic density of states corrugation which is

opposite to the geometric effect. However, corrugations

in the electronic density of states are usually some

tenths of angstroms, so this effect should not be able

to cancel the larger geometric effect. In fact, the

observation that the corrugations in Fig. 2a are about 1 A, less than the expected value of 1.6 A, suggests that

the presence of excess titanium in these positions

compensates somewhat for the geometric effect. Although

previous models of the reduced TiO2(110) surface have

not considered the possible existence of CS planes or

their effect on the titania surface structure [12], the

existence of these planar defects has been proposed by

Firment et al. [13] to explain angle-resolved UPS data

451
from the reduced MoO3(010) surface, another transition

metal oxide which forms a variety of CS structures.

CONCLUSION

This work demonstrates that the STM is capable of

determining the surface structure of ionic materials,

such as Ti02 and ZnO, and that these oxides have unusual

surface structures which undoubtedly play a role in

determining their activity as catalysts and sensors.

ACKNOWLEDGEMENTS

The authors acknowledge J. Vohs for providing the ZnO

crystal. Also, V. Henrich and Z. Zhang are acknowledged

for preparing and providing the rutile sample. This

work was funded by IBM and the Nation Science

Foundation

.

REFERENCES

[I] W. E. Farneth, E. M. McCarron III, A. W. Sleight,

and R. H. Staley, Langmuir 3, 217 (1987)

.

[2] W. Gopel, Prog. Surf. Sci . 20,9 (1985).

[3] V. Henrich, Rep. Prog. Phys . 48, 1481 (1985)

.

[4] Z. C. Rang, D. J. Smith, and L. Eyring, Surf. Sci.

175, 684 (1986)

.

[5] M. Grunze, W. Hirschwald, and D. Hofmann, J. Cryst

.

Growth 52, 241 (1981)

.

[6] H. R. Sadeghi and V. E. Henrich, Appl . Surf. Sci.

19, 330 (1984) . [7] The STM head was manufactured by WA technologies,

England

.

[8] G. Rohrer and D. Bonnell, J. Am. Ceram. Soc . , in

press.; G. Rohrer, V. Henrich and D. Bonnell,

submitted to Science.

[9] G. S. Rohrer and D. A. Bonnell, in preparation

[10] F. Flores. and N. Garcia, Phys. Rev. B 30, 2289

(1984)

.

[II] L. A. Bursill and B. G. Hyde, Prog, in Solid State

Chem . Reiss and McCaldin, Ed (Pergamon Press, New

Jersey, 1972), Vol. 7, p. 177.

[12]W. G(5pel, G. Rocker, and R. Feierabend, Phys. Rev. B

28, 3427 (1983)

.

[13] L. E. Firment, A. Ferretti, M. R. Cohen, and R. P.

Merrill, Langmuir 1, 166 (1985) .

452
DISCUSSION

J. Nowotny: In the Max Planck Institute for Solid State Research in Stuttgart, there has

been research on the electrical properties of titanium oxide involving measurements of

electrical conductivity as a bulk property and work function as a surface property under

the same conditions. What has been observed is that the same material, under the same

conditions is n-type in the bulk and is p-type at the surface. So we have p-type conditions

coming from the bulk to the surface. Is it known how the surface might become the type

as you have mentioned?

G. S. Rohrer: Yes, under some circumstances I think it's quite possible for there to be

an accumulation layer at the surface that could actually transform the material from one

type to the other. However, that did not appear to be the case during these

measurements.

D. Cahen : Could it be possible that when you try to compare the results from UPS, XPS

and STM, you see problems that are associated with desorption and adsorption of

oxygen? Because unless you work consistently at very low pressures you will get oxygen

very quickly, on most of these surfaces.

G. S. Rohrer: I have to say that's always a possibility, because I don't know exactly the

chemistry of the edge sites. But I'd also have to point out that these surfaces stay very

clean in the UHV pressure range. In fact, they only interact weakly with most gases, so

that once you clean them off they are very stable. These experiments were carried out

in roughly the 5x10"^^ torr pressure range.

R. Newnham: What are the prospects of using STM for real ceramics?

G. S. Rohrer: For nonconductive ceramics, the prospects are very poor right now. We've made preliminary efforts to do this at Penn, by photo-exciting surfaces. Another

approach is to dope an insulating surface and make it conductive, but then, of course, you

might change the chemical properties of the surface.

We've already put out some papers on polycrystalline zinc oxide ceramics and there is a lot of promise for probing grain-boundary interactions. But the key thing right now is that the crystal has to have some kind of conductivity.

D. Cahen: I'm a little bit concerned about your remark that you have clean surfaces,

because in all papers which I know there's always a question of oxygen and water

absorption. How do you know that you have clean surfaces?

G. S. Rohrer: Well, with respect to oxygen and water absorption, it's difficult to say

absolutely. However, I followed instructions from the literature for cleaning procedures

which produced one-by-one surfaces as determined by LEED.

453
As far as the chemistry and possible absorption at the edge sites, I have to try to deduce

that from the experimental results. However with respect to carbon-containing species,

my experiments indicate that once I cleaned the surfaces, they stayed clean in UHV.

454
VII. THEORY AND MODELING

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

FACTORS DETERMINING THE DIELECTRIC CONSTANTS

OF OXIDES AND FLUORIDES

R. D. SHANNON

E. I. Du Pont de Nemours and Co., Inc.

Central Research and Development Department

Experimental Station

Wilmington, DE 19880-0356

ABSTRACT

A preliminary set of 16 ion dielectric polarizabilities was derived from the dielectric constants of 63 oxides and 4 fluorides

using a least squares fitting procedure and the ion additivity rule

in conjunction with the Clausius-Mosotti equation and accurately

determined dielectric constants and molar volumes. Assuming an arbitrary value for anCB"^"*") = 0.05 anion polarizabilities of

2.01, 2.18, and 1.63 A^ are found for 0=, OH", and F",

respectively. Typical cation polarizabilities are 0.79, 0.85, 1.31

and 3.15 A^ for A13+, Si4+, Mg2+ and Ca2+, respectively. The

agreement between measured dielectric polarizabilities as determined from the Clausius-Mosotti equation and those calculated

from the sum of oxide polarizabilities and the sum of ion

polarizabilities is -~ 0.5% for 18 oxides. The effects of

composition, structure, polar impurities, mixed valence and defects

on total polarizability are demonstrated.

INTRODUCTION

Recent efforts to replace AI2O3 as an electronic packaging

material have generated an interest in a better understanding of

the factors which determine the dielectric constant of oxide

materials (Chowdhry and Sleight, 1987; Kawakami et al., 1987; Kondo

et al., 1987; Tummala, 1988). Olhoeft (1981), in a study of a wide

variety of silicates found the dielectric constant, k' = 1.91^

where d = density in g/mL. However, plots of k' vs. d for a variety of borates, aluminates, silicates and phosphates show

considerable scatter.

Another approach used to systematize dielectric constant data

for many decades involves the use of polarizabilities. The

457
dielectric polarizability , a^, is related to the measured

dielectric constant by the Clausius-Mosotti equation:

«D = 1/b [(\) (k'-1)/(k'-.2)] (1)

where is the molar volume in A^, b is assumed to be 47r/3, and

k', the real part of the complex dielectric constant, is measured

in the range 1 KHz to 10 MHz (Roberts 1950,1951). The dielectric

polarizability includes both ionic , ot^, and electronic , a^,

components. The electronic polarizability, a^, is related to the

refractive index, n, by the Lorenz-Lorentz equation (Lorentz 1880;

Lorenz 1880):

«e

= [{\) (n2 -l)/(n2 +2)] (2)

Although the Clausius-Mosotti equation is strictly valid only

for compounds where the molecule or ion has cubic symmetry

(Szigeti, 1949; Bosman and Havinga, 1963; Dunmur, 1972; Megaw,

1975; and Roberts, 1949,1950,1951), it has been shown to be

approximately valid for a number of noncubic crystals by Roberts

(1949, 1951) and Lasaga and Cygan (1982).

The concept of additivity of molecular polarizabilities

whereby the molecular polarizability of a complex compound A2BX4

can be broken up into the molecular polarizabilities of the simpler

compounds AX and BXg according to:

aD(A2BX4) = 2 acCAX) + ^^(BXs) (3)

is useful in systematizing dielectric constant data and has been

discussed by Heydweiller (1920) , Cheng (1940) , Jonker and Van

Santen (1947), Roberts (1949, 1950, 1951), and Lasaga and Cygan

(1982) . It has recently been demonstrated that the oxide

additivity rule, in conjunction with simple oxide polarizabilities,

accurate dielectric constants and molar volumes, results in

agreement between observed and calculated polarizabilities of ~ 1%

for a wide variety of oxides (Shannon and Subramanian, 1989;

Shannon et al., 1989; Shannon et al., 1990a, b).

Molecular polarizabilities of complex substances can also be

broken up into ion polarizabilities according to:

a(A2BX4) = 2 a(A2+) + a(B4+) + 4 a(X=) (4)

Ion polarizabilities are potentially more useful and easier to use

for predictive purposes than simple oxide polarizabilities. Many

sets of empirical electronic ion polarizabilities, a^, have been

derived from the C-M equation and the additivity rule using the

458
alkali halides and alkaline earth chalcogenides (Tessman, Kahn and

Shockley (TKS) , 1953; Pirenne and Kartheuser, 1964; Wilson and

Curtis, 1970; Boswarva, 1970; Coker, 1976; Claro, 1978) and using a variety of minerals (Lasaga and Cygan, 1982) . These

polarizabilities were found to be useful in predicting refractive

indices. A set of 18 dielectric ion polarizabilities j_ ac> > derived

from cubic compounds by Roberts (1949), assuming o!p(0~) = 2.387 A"^ was found to give agreement between observed and calculated

polarizabilities of -~ 5%. A second set of 11 dielectric ion

polarizabilities was derived f or_ application to minerals by Lasaga

and Cygan (1982), assuming ap(0~) = 2.37 A^. Using these ion

polarizabilities agreement between observed and calculated total

polarizabilities of 24 common minerals was ~~ 10% with some individual deviations as large as 20%.

In this paper we derive a set of preliminary dielectric ion

polarizabilities that can be used to predict total polarizabilities

of a variety of oxides and fluorides to an accuracy of about 0.5%.

We then summarize recent data showing the validity of the oxide

additivity rule and compare the results from the oxide additivity

rule with those from the ion additivity rule. Finally, we show

that large deviations from additivity can sometimes be traced to

certain physical characteristics such as ionic or electronic

conductivity, the presence of dipolar impurities or piezoelectric

resonance

.

PROCEDURE

The data set includes only single crystal dielectric constants

of estimated precision < 1%; data from noncubic crystals is

included only if the dielectric anisotropy was determined. Sources

include Westphal and Sils (1972) , Young and Frederickse (1973)

,

Olhoeft (1981), Shannon et al. (1989, 1990a, 1990b, 1990c), Shannon

and Subramanian (1989) and Subramanian et al. (1989). Many of the

data found in the earlier compilations (Keller, 1966; Westphal and

Sils, 1972; Young and Frederickse, 1973; and Olhoeft, 1981) were

not useful because of poor accuracy, poor knowledge of chemical

composition or cell dimensions, polycrystalline nature of the

samples or lack of complete anisotropic data for noncubic

materials. Where mineral dielectric constants were used, efforts

were made to include only data on gem quality crystals with

accurate cell dimensions and whose composition was not in question.

A few dielectric constants obtained from infrared reflectance data

were used; accuracies of ~~ 1% were assumed (Gaskell and Johnson,

1974) . Dielectric data from ionic conductors and f erroelectrics as well as data from polycrystalline samples were avoided.

I

459
If it is assumed that the total molar dielectric

polarizability of a compound can be calculated (acalc) ^ simple

linear combination of individual ionic polarizabilities,

a (ion) , then:

N

«calc = ^ n-aoCion) (5)

1=1

where, I varies over the total number (N) of types of ions (I) in

the formula unit and n is the number of ions of type I in the

formula unit, and Q;i5(ion), can be used as a set of refinable

parameters in a least squares procedure which minimizes the

function:

M

E ''i ^"obs - «calc)^

i=l

Here i varies over the number of measurements of cti^ for a variety

of compounds and w^ = c^'^ , where O"^ is the estimated % error in

the experimental dielectric constant.

The standard deviation of a refined value of ai5(ion) is

calculated as:

<r[aD(ion)] = (7)

where Mj is the diagonal element of the inverted normal matrix

corresponding to the ionic polarizability for ion "I" varied in the

least squares procedure.

In practice it is necessary to fix one polarizability as done

for electronic polarizabilities by TKS {a^(lii*)) , Wilson and Curtis

(ag(Li*)), Coker (a^CLi*)), and Boswarva (ag(Mg2+)), and for

dielectric polarizabilities by Roberts (o;d(^~))* ^® have

arbitrarily assumed the value for ^^(B^"'') = .05 A^. This

assumption results in a value for oc^{0~) = 2.01 A^, close to values

obtained by assuming an a(0~)/r3 law (Shanker et al., 1973) with

either crystal radii (aoCO") = 2.15-2.87 A^; Shannon, 1976) or bonded radii (aoCO^) = 1-97-2.55 A^; Gibbs and Boisen, 1986).

Empirical and theoretical studies of electronic

polarizabilities show that cation electronic polarizabilities are

460
constant and anion polarizabilities vary approximately with lattice

constant and anion molar volume (TKS, 1953; Kirsch et al., 1974;

Coker, 1976; Mahan, 1980; Fowler and Madden, i984a,b). In the

refinements carried out here, a slight linear dependence of q;q on oxygen molar volume was found. However, because of the limited set

of data used, this dependence was ignored in the set of ion

polarizabilities reported in Table 1.

RESULTS

The results from refinements of ai5(ion) made using

experimental polarizability data from 63 oxides and 4 fluorides are given in Table 1. The ratio of observables to variables is only

2.4; desirable values are of the order of 3-6. Most of the

calculated uncertainties in refined ion polarizabilities are small,

i.e., 0.001-.005 A^. These uncertainties are unrealistically low; a more reasonable estimate of the uncertainty is probably ± 0.02 A"^

for the smaller ions such as Al^* and Si^"*" and ± 0.05 A"^ for larger

ions such as Ca2+, Y*^"*" and Zr^"*".

Table 1

Ion Dielectric Polarizabi lities

Ion a

F-1 1.63 0.001

0-2 2.01 0.002

OH-1 2.18 0.005

Be+2 0,30 0.003

Mg+2 1.31 0.002

Fe+2 2.22 0.005

Mn+2 2.64 0.005

Ca+2 3.15 0.002

B+3 0.05

Al+3 0.79 0.003

Ga+3 1.49 0.003

Sc+3 2.81 0.004

Y+3 3.79 0.003

Gd+3 4.36 0.003

Si+4 0.85 0.004

Zr+4 3.27 0.013

Table 2 compares total molecular polarizabilities of a variety

of borates, aluminates, gallates and silicates determined from the

measured dielectric constants using the Clausius-Mosotti relation

and from the oxide and ion additivity rules. In general both the

461
oxide and ion additivity rules are obeyed to ± 0.5%. The rules

hold for simple oxides such as BeAl204 and Be2Si04 as well as for more complex compositions such as CeiB2Si208 and Y3Sc2Ga30]^2 •

Table 2.

Comparison of Observed and Predicted Single Crystal

Dielectric Polarizabilities

Pred.a Pred.^ exp*=

Compound <k'> a-Y a-p a-p

Al ggFe oiOOH 8.44 4.99 5.02 + 0.6

Mg' QftFe'ooAlBO^ 8.075 10.24 10.21 10.18 - 0.6 - 0.3

CaBoSioOo 6.73 21.13 21.05 21.29 + 0.8 + 1.1

BeAl204 8.93 9.84 9.92 9.87 + 0.3 0.5

BegBOaOH 4.82 8.84 8.80 0.5

BeAlSi040H 6.64 12.11 12.17 12.16 + 0.4 0.1

MgAl204 8.18 10.96 10.94 11.10 + 1.3 + 1.4

Y3AI5O12 10.60 39.79 39.48 39.38 1.0 0.2

Y3Sc2Al30i2:Nd 12.20 43.64 43.60 43.48 0.4 0.3

Gd3Sc2Al30i2 12.71 45.73 45.23 45.35 0.8 + 0.3

^^3 . 03^^4 . 97^12 12.18 44.83 44.78 44.66 0.4 0.3

Y3Sc2Ga30i2 12.94 45.30 45.63 45.42 + 0.3 0.5

Be2Si04 6.21 9.31 9.49 9.35 + 0.5 1.5

Mg2Si04 7.00 11.54 11.53 11.54 0.0 0.0

7.27 11.78 11.76 11.80 + 0.2 + 0.3

Fe2Si04 8.77 13.28 13.24 13.24 0.3 0.0

ZrSi04 11.5 12.16 12.17 12.12 0.3 0.4

<A> = 0.5% 0.5%

a. Oxide additivity rule

b. Ion additivity rule

c. Using the Clausius-Mosotti relationship: polarizabilities in

DISCUSSION

The set of dielectric polarizabilities found in Table 1 is

significantly different from the sets of Roberts (1949) and Lasaga

and Cygan (1982) . The differences arise primarily because of the

large difference in (0=): 2.01 a3 (our set) vs. 2.387

462
(Roberts; Lasaga and Cygan) . Their higher value occurred because

they had no data from borates or beryllates; use of aD(0~) = 2.387

results in negative polarizabilities for B^"*" and Be2+. Our

polarizabilities are more reliable because of more accurate

dielectric constants from a larger number of compounds involving

more ions. The problem of relative values of cation and anion

polarizability still remains, but based on the a^/r^ relationship

and the concept of bonded radii, the values obtained from the least

squares procedure seem reasonable.

The Clausius-Mosotti equation shows explicitly how dielectric

constant depends on composition and crystal structure through

polarizability and molar volume. Low dielectric constants are clearly favored in compounds containing ions of low polarizability.

Figure 1

DIELECTRIC CONSTANT VS. MOLAR VOLUME OF Si02

K = (3V + 87iaD)/(3V - 4mp)

K

24

22

20

18

16

14

12

10

8

6

4

2

0

1

- ostishovite

1 1 1

amorphous

quartz •

tridymite MC

1

1 1 1 1 25 30 35 40

Molar Volume, A

^

45 50

463
Crystal structure is generally important in determining k only

insofar as it determines molar volume. The effect of the molar

volume term is shown for the Si02 polymorphs in Figure 1 where Vj^

is plotted vs. the calculated dielectric constant, k (calc) . There

is a rapid decrease in K(calc) as V^^ increases. Thus, the more

open the structure, the lower the value of the dielectric constant.

The dielectric constant of stishovite, although not yet determined,

would be predicted to be approximately 22. The more open tridymite

and amorphous structures have lower k than the more dense quartz

polymorph.

The excellent agreement between observed and calculated total

polarizabilities in Table 2 implies a good degree of predictability

using the C-M equation for well-behaved compounds. Some of the

poor agreement between calculated and predicted k's in more complex

structures indicated in Table 3 can be traced to factors which are known to increase k: ionic and electronic conductivity, HgO or CO2

Table 3.

Deviations from the Additivity Rule

Type and Composition <K •> Pred.» Bxp.b

<Kt Ref

.

Li6a02 LiGa02 7 27 6.47 6 96 +7 0 1

fayalite Fej . ggUn . 07^8 . 02^1 . oi^ii .O1O4 14 1 13.25 15 0 +11 8 2

cordierite Na os'fel 07^^ 03^l3 87^^

[(H20).28(C02y.O8]

. O3Si50l8- 5 81 46. 4<:

57.

55

55

1 1 1

00

00

+16

+2

8

1

3

fluorite CaF2:Er (3%) 8 28 6.48 6 90 +6 1 4

akermanite Ca2CoSi207 11 29 23.88 28 58 +16 4 5

pyrope Ugi . QgFe . ggCa, osM^i . oiAli . 98Si30l2 12 20 33.54 35 72 +6 1 6

almandine Fej _ ^sMgj _ le^^. 20^. 04^^ 2 . OoSi30i2 11 96 33.65 35 99 +6 5 6

grossular Ca2.98Mn_oiAli g7Fe_02^i2 . 94O12 8.53 37.59 35 67 -5 4 6

a. Oxide additivity rule

b. Using the Clausius-Uosotti relationship; polarizabilities in A-^ c. Omitting contributions of H2O and CO2

d. Including contributions of H2O and CO2: «d(^2^) = **1 * P^/3kT = 3.5 + 28.4 = 31.9 A^/mole H2O

1. Shannon et al., 1990a 2. Suwa, 1964 3. Shannon and Subramanian,

unpublished 4. Fontanella and Andeen, 1976 5. Shannon et al. (1990d)

6. Shannon et al. (1990c).

464
in channels or large cavities. Thus, poor agreement in the case of

LiGa02 is probably caused by Li mobility (Shannon et al., 1990a).

Dielectric constants of Fe2Si04 vary widely i deviations from a mean

value of about 8.8 are associated with high losses and can be

associated with the presence of Fe^"*" and space charge polarization.

Poor agreement between ^^(calc) and ^^(exp) in cordierite is caused

by lack of consideration of the presence of HgO and CO2 in the

channels (Shannon et al., 1989). The increased polarizability of

CciFg'Er is caused by the formation of Er3+-Fi dipoles (Fontanella

and Andeen, 1976). The unusually large deviation of the observed

dielectric constant of Co akermanite from the calculated value is

probably associated with piezoelectric resonance and possibly also

the incommensurate structure and possible local displacements of

the Ca ions from their mirror planes (Hemingway et al.,1986).

Positive deviations of pyrope and almandine and the negative

deviation of grossular appear to be related to structural

peculiarities of the garnet family. Similar deviations from

additivity were found by Holland (1989) in a study of the

thermodynamic quantities of silicate minerals, who found the

entropy contributions of Mg2+ and Fe2+ in pyrope and Ca2+ in

grossular to be larger and smaller, respectively, than expected for

these cations in eight-coordination.

Garnet structural anomalies include: (1) an unusual amount of

edge sharing and (2) octahedral edges shared with the dodecahedral

edges which are longer than the unshared edges (Novak and

Gibbs,1971). These constraints cause changes in what are normally

invariant bond distances. For example, the mean Al-0 distance,

d<Al-0> is 1.886 A in pyrope, but is 1.924 A in grossular. The

mean <Al-0> distance varies continuously in the series

pyrope-almandine-spessartite-grossular (Novak and Gibbs, 1971) and

is inversely proportional to the garnet polarizability deviations.

The differences between shared and unshared 0-0 distances lead to

bond strain as defined by Novak and Gibbs. The deviation from

additivity of polarizability is approximately proportional to the

bond strain. Thus, instead of anomalies only at the pyrope and

grossular end members, the polarizability deviations are seen to be

a gradual function of both d<Al-0> and bond strain. The

polarizability additivity deviations seem, therefore, to be related

to the rigidity of the garnet framework.

ACKNOWLEDGEMENTS

I wish to thank A. J. Owens and J. B. Parise for the least

squares program, M. A. Subramanian, R. A. Oswald and R. W. Shiffer

for sample preparation and dielectric constant measurements, J. T.

Armstrong and G. R. Rossman for electron microprobe analyses, and

C. Foris and A. Brown for cell dimensions.

465
REFERENCES

Bosman AJ, Havinga EE (1963) Phys. Rev. 129:1593-1600.

Boswarva, IM (1970) Phys. Rev. Bl : 1698-1701

.

Cheng CK (1940) Phil. Mag. 30:505-515.

Chowdhry U, Sleight AW (1987) Ann. Rev. Mater. Sci. 17:323-340.

Claro FH (1978) Phys. Rev. 618:7058-7061.

Coker H (1976) J. Phys. Chem. 80:2078-2084.

Dunmur DA (1972) Molec. Phys. 23:109-115.

Fontanella J, Andeen C (1976) J. Phys. C. 9:1055-1064.

Fowler PW, Madden PA (1984) Phys. Rev. 829:1035-1042.

Fowler PW, Madden PA (1985) J. Phys. Chem. 89:2581-2585.

Gaskell PH, Johnson DW (1974) J. Noncryst. Solids 20:153-169.

Gibbs GV, Boisen MB (1986) Better Ceramics Through Chemistry II,

Mat. Res. Soc. , Pittsburgh 73:515-527.

Hemingway BS, Evans HT, Nord GL, Haselton HT, Robie RA, McGee JJ

(1986) Can. Miner. 24:425-433.

Heydweiller A (1920) Z. Phys. 3:308-317.

Holland TJB (1989) Amer. Mineral. 74:5-13.

Jonker GH, Van Santen JH (1947) Chemisches Weekblad 43:672-679.

Kawakami K, Takabatake M, Minowa T, Chiba J, Sasaki M (1987) Adv.

Ceram. 19:95-102.

Keller JV (1966) Handbook of Physical Constants, Eds. SP Clark Jr.

Geological Soc. Amer. Memoir 97, Geological Soc. Amer. pp.

553-577.

Kirsch R, Gerard A, Wautelet M (1974) J. Phys. C: Solid State Phys.

7:3633-3644.

Kondo K, Okuyama M, Shibata Y (1987) Adv. Ceram. 19:77-87.

Lasaga AC, Cygan RT (1982) Am. Mineral. 67:328-334.

Lorentz L (1880) Ann. Phys. Chem. 11:70-103.

Lorenz HA (1880) Ann. Phys. Chem. 9:641-645.

Mahan GD (1980) Solid State Ionics 1:29-45.

Megaw HD (1975) Ferroelectricity in Crystals. Methuen, London, pp.

165-178.

Novak GA, Gibbs GV (1971) Amer. Mineral. 56:791-825.

Olhoeft GR (1981) Physical Properties of Rocks and Minerals . Y.S.

Touloukian et al. Eds. McGraw-Hill, New York. pp. 257-329.

Pirenne J, Kartheuser E (1964) Physica 30:2005-2018.

Roberts R (1949) Phys. Rev. 76:1215-1220.

Roberts R (1950) Phys. Rev. 77:258-263.

Roberts R (1951) Phys. Rev. 81:865-868.

Shanker J, Kumar N, Verma MP (1973) Ind. J. Pure Appl. Phys.

11:644-647.

Shannon RD (1976) Acta. Cryst. A32: 751-767.

Shannon RD, Subramanian MA (1989) Phys. Chem. Minerals 16:747-751.

Shannon RD, Subramanian MA, Mariano AM, Rossman GR (1989) Proc.

Mater. Res. Soc. Symposium on Materials for Magneto-Optic Data

Storage, San Diego, April 24-27.

466
Shannon RD, Subramanian MA, Allik TH, Kimura H, Kokta MR, Randies

MH, Rossman GR (1990a) J. Appl. Phys. 67:3798-3802.

Shannon RD, Subramanian MA, Hosoya S, Rossman GR (1990b), to be

published.

Shannon RD, Subramanian MA, Rossman GR (1990c), to be published.

Shannon RD, Subramanian MA, lishi K, Rossman GR (1990d) , to be

published.

Subramanian MA, Shannon RD (1989) Mater. Res. Bull. 24:1477-1483.

Subramanian MA, Shannon RD, Ghai BHT, Abraham MM, Wintersgill MC

(1989) Phys. Chem. Minerals 16:741-746.

Suva K (1964) J. Earth Sci. Nagoya Univ. 12:129-134.

Szigeti B (1949) Trans. Farad. Soc. 45:155-166.

Tessman JR, Kahn AH, Shockley W (1953) Phys. Rev. 92:890-895.

Tummala RR (1988) Ceram. Bull. 67:752-758.

Westphal WB, Sils A (1972) Dielectric constant and loss data. US

National Technical Information Service AFML-TR-72-39

.

Wilson JN, Curtis RM (1970) J. Phys. Chem. 74:187-196.

Young KF, Frederickse HP (1973) J. Phys. Chem. Ref .Data 2:313-409.

467
DISCUSSION

B. Wuensch : Bob, based on a hazy recollection of what lies behind the Clausius-Mosotti

equation, I believe that it hinged on an assumption by Lorentz of polarization of point

dipoles. Your oxide additivity relations nicely finesses the variation of individual ion

polarizabilities from structure to structure. Yet, in other calculations such as the MottLittleton-type calculations of defect formation energies (Yamashita and Kurosawa, J. Phys. Soc. Japan 89,944,1954) problems were encountered with ionic polarizabilities

when attempting to extend the calculations from simple ionic compounds such as the

alkali halides to oxides. Is there any correlation between the extent to which the original

Lorentz assumption is followed and your deviations from additivity?

R. Shannon: The Yashita-Kurosawa calculations showed that it was necessary to introduce variable oxygen polarizability in order to obtain meaningful results in

calculations of defect formation energies in MgO, CaO and SrO. Variable oxygen

polarizability was found by Tessman, Kahn and Shockley (Phys. Rev. 92,890,1953) and

also by others (Mahan, Solid State Ionics, 1,29,1980; Fowler and Madden, J. Phys.

Chem. 89,2581,1985) in studies of electronic ion polarizabilities. However, my studies

of dielectric polarizabilities show little or no dependence of oxygen polarizability on oxygen molar volume. At the present time I do not see a clear correlation between the

additivity relationship deviations and oxygen molar volume.

M. O'Keeffe: In pyrope the bond valence sum at Mg is much less than two and in

grossular the bond valence sum is much greater than two. Such deviations from ideal

values may serve as warnings that the additivity relations will not hold.

R. Shannon : I agree. The unusual bond valence sums around the X ions in silicate

garnets are a result of structural constraints imposed by the unusual extent of edgesharing. This results in bond distances and bond valences that give rise to unusual

polarizabilities that may be characteristic of "rattling" Mg ions and "compressed" Ca ions.

M. O'Keeffe: (1) Dielectric constants depend strongly on frequency and temperature.

Did you carry out yout experiments at a standard temperature and frequency? (2) The

electronic part of the dielectric constant depends on collective properties, i.e., charge.

Perhaps an electronegativity correction might be useful.

R. Shannon : The experiments were carried out at room temperature and over a frequency range of 30 KHz to 3 MHz but the results are reported at 1 MHz. Over that

frequency range there is generally less than 0.2% deviation for the samples that I have

looked at so far.

D. Cahen: Did you also note only a 0.2% variation in the dielectric constant of LiGa02

over the frequency range 30 KHz to 3 MHz? If your explanation for the additivity

deviation in LiGa02 is correct, ionic migration will be expressed in a variation of K with

frequency.

468
R. Shannon : The dielectric constants reported for LiGa02 were obtained at BTL and were not reported at different frequencies. However, Grant et al. (J. Am. Ceram. Soc.

60,226,1977) showed a variation in K with frequency at 482°C.

R. Newnham: In going from the total polarizabilities for the simple oxides such as Si02

and CaO to the polarizabilities for individual ions, do we have the same problem here that

we have with ionic radii, where we must arbitrarily fix the radius of one ion? If so, do

these individual polarizabilities mean that much?

R. Shannon : We have all the same problems with ion polarizabilities that we have with

ionic radii. Although it is necessary to mjike some arbitrary choices, there are limits.

For example, it is not possible to choose the polarizabilitiy of oxygen to be 2,4 A'' because that results in negative polarizabilities for B^"^ or Be^"*". I have picked an

arbitrary value for the dielectric polarizability of B^"*" (.005 A^) that seems reasonable

when you look at calculated values for electronic polarizabilities. I don't know if that's a good assumption but at the present time it is the best I can do.

R. Newnham: So there might be something like a 10% limit for the polarizabilities or

is it higher than that?

R. Shannon: It could be larger than 10%.

469

National Institute of Standards and Technology Special I*ubIication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

INTERNAL STRAIN IN PEROVSKITE RELATED MATERIALS

I.D.Brown

Institute for Materials Research, McMaster University,

Hamilton, Ontario, Canada L8S 4M1

1. INTRODUCTION

Many simple inorganic structures can be built from AO and BO2 layers (Figure

1), where A and B are cations (metals) and O is an anion which in this paper is assumed to be oxygen. Five additional types of layers may be created by

removing either oxygen atoms (to give the layers A, B and BO) or cations (to give

the layers O and O2).

X Bq:^—V >: -B-O-rS-O-B-O-Bw ! : V 0 0 0 0

W^V 6 6 ^6 6

A A A -B-O-B-O-B-O-BFigure 1 By stacking the layers together in different sequences, a large variety of different

structures can be generated. Rocksalt has the sequence (AO.AO.AO), fluorite the

sequence (A.O2.A.O2) and perovskite the sequence (BO2.AO.BO2.AO). The

number of different sequences that are found in nature is limited by crystal

chemical constraints. The sequence must be electrically neutral for the given

composition and the cations must be in sites with appropriate coordination

environments. The layers must also have the same in-plane spacing which, given

that no two cations have exactly the same size, means that the bonds in some

layers will be stretched while those in others will be compressed. A variety of

mechanisms is available for relaxing these strains depending on the size and nature

of the cations present, and it is this variety that leads to the rich crystal chemistry

that these systems display. In some cases the strain is not completely relaxed and

the resulting residual strain itself may have important effects on the properties.

The three most important crystal chemical constraints are discussed in the next

section which is followed by a detailed description of the relaxation mechanisms.

Four examples are then considered and the paper concludes with speculations on

the relation between the strain and the superconducting transition temperature.

2 THE CRYSTAL CHEMICAL CONSTRAINTS

1. The sequence must provide a site whose coordination number is close to the

Average Observed Coordination Number (AOCN) of the cation.

471
The average observed coordination numbers (AOCN) for many cations

surrounded by oxygen have been determined by examining all the

compounds stored in the Inorganic Crystal Structure Database . In the

majority its of compounds, a cation is found with a coordination number

that lies within 1, or possibly 2, of its average, i.e. its AOCN.

Cation environments with almost any desired coordination number can

be created in structures composed of AO and BOj layers by selecting a

suitable sequence. Examples of coordination nun^bers and the sequences

that give rise to them are:

2 AO.B.AO

4 AO.BO.AO A.BO2.A

5 AO.BO2.A AO.AD.B

6 AO.BO2.AO AO.AO.AO

7 A0.A07B0

8 BO2.A.BO2 BO2.A.O2

9 AO.AO.BCl

10 BO2.AO.BO

12 BO2.AO.BO2

In each case the coordination number refers to the cation in the center of

the triplet.

2. The compound must be charge neutral.

This is an absolute constraint for the whole compound but local charge

neutrality is also expected. For example, BaBiO^ is composed of

alternating BaO and Bi02 layers, each of which is charge neutral, but

Cu'*"'"02 carries a double negative charge, and is expected to be surrounded

by positively charged layers. A perovskite sequence (CuO2.AO.CuO2.AO)

cannot occur since no quadrivalent A cations have a coordination number

of 12. The sequence (CUO2.A.CUO2.A) requires only a divalent ion with a coordination number of 8. Ca (AOCN = 7.3) is a good candidate and is the

cation most often found in this sequence.

3. Each layer in the sequence must have the same in-plane lattice spacing.

If the layers are to stack to form a commensurate crystal, they must all have exactly the same in-plane lattice spacings (a and b). Since each

cation has its own distinctive size, commensurability can only be achieved

by compressing some and stretching other bonds. The fascinating and

extensive crystal chemistry shown by these compounds largely results from

the different ways in which these strains are (or are not) relaxed. The

different relaxation mechanisms are discussed in more detail in the next

section.

3. FITTING THE LAYERS TO A PROCRUSTEAN BED

Bond valences provide a quantitative method of analyzing the bonding

requirements of different cations. They also provide a convenient tool for

modelling the relaxations. The crystal is treated as a network of bonds that link

nodes (atoms). The important properties of this network are the valences

(oxidation states) of the atoms and the valences associated with the bonds. These

are related by two Kirchhoff-like network equations:

472
1 . The sum of bond valences at each node equals the atomic valence of the

node

2. The sum of bond valences around any loop is zero. Bond valences are found empirically to be inversely related to bond length^. By

solving the Kirchhoff equations for a given bond network, bond valences can be

predicted and from these the expected bond lengths can be calculated. Even if the

bonding network is not known, it is possible to predict the average length of the

bonds formed by a cation by recognising, from rule 1, that the average bond

valence is equal to (atomic valence)/(coordination number). Average bond lengths

calculated in this manner have been used to determine the ideal in-plane lattice spacings for various cations in the coordinations close to their average observed

coordination number (Fig. 2).

^0

<i> E

o

- Cu3*

3-5

31

Pb

fig

Sr Sr

SfPb la Eb la

La

La Co if

Ca Y ai X

Tl YIL

I I I 456 6789 10 —B— A

Coordination Number

Figure 2. Lattice parameters expected for various AO and BO^

layers with A and B cations in coordinations close to their ACXTN

(underlined).

BaO and SrO layers match the Cu02 layers in size, but most cations have ideal inplane spacings that are considerably smaller than those of the CUO2 layers. A

variety of mechanisms is available for compressing or, more usually, stretching a layer to provide a match.

1. Buckling. Buckling a layer results in non-planar bonding and brings the

B (or A) atoms closer together. The effect is not large, a 0.3A

displacement of the O atoms perpendicular to the layer reduces a by only

O.OSX. This is generally not an effective way of reducing the in-plane

spacing of a layer because buckling is usually controlled by stronger

constraints on the inter-layer bonding.

473
2. Changing the inter-laver spacing. Any coordination environment can be

expanded within a layer by compressing it perpendicular to the layer (and

vice versa) . There are two ways in which the bonding environment of a

cation can be compressed to increase the in-plane lattice spacing:

i) For octahedrally coordinated A cations shortening the bonds

between layers results in lengthening the bonds within the layers.

The effectiveness of this mechanism is shown in Fig. 3 (top) for A

_ in octahedral coordination. The broken line indicates

regular coordination. To match the spacing of the CUO2 layer

(3.85A), the inter-layer spacing would have to be reduced by 03K

leaving Tl with a coordination that most chemists would describe as

linear.

I I I I I 1 Lo 30 AO A

Introlayer Spacing (a)

Figure 3

ii) For 8 -coordinate A cations, changing the bond angles at the

cation can change the ratio of in-plane (a) to inter-layer spacing

(d) without changing the bond lengths. The effect, which depends

only on geometry, is shown in Fig. 3 (bottom) for A = Y, Ca and

Sr. Changing the inter-layer spacing is effective in changing the in-plane

spacing when only one kind of bond bridges the layers, as found, for

example, between two TIO layers or around a Ca layer that contains no

anions. When two different kinds of bond bridge the layers, they are usually in competition, one bond attempting to decrease, the other

increase, the inter-layer spacing (see, for example. Fig. 6).

3. Moving^ the cation off the center of its coordination sphere. As shown

elsewhere^, if the bonds formed by a cation are changed from their

average value by making some longer and some shorter, the average bond

length increases. If a cation is moved within its layer away from the

center of its coordination sphere, the ligand atoms will move further apart.

Rule 1 can be used to calculate how the bond valences, and hence the bond

lengths, vary as the environment is distorted. Its effectiveness is shown by

474
the broken line in Fig. 4, which shows the changes in the bond lengths and

lattice spacing produced by a trigonal distortion of octahedrally

coordinated Bi . The increase in lattice spacing is negligible for small

distortions but is significant for the kinds of distortions typically shown by

cations with stereoactive lone pairs of electrons in their valence shells (e.g.

T1+, Pb++ and Bi+++).

4. Charge transfer between cations . Increasing the charge on a cation

shortens the bonds that it forms. Since charge is conserved, increasing the

charge on one cation means decreasing the charge, hence lengthening the

bonds, of another cation in the structure. This is modelled in Fig. 5 which

shows the effects of moving positive charge from Tl''"''"'' to Cu*'^. The

effect is large but this mechanism is restricted to compounds that have two

different cations both capable of changing oxidation states.

(Valence of short bond)

Figure 4 Figure 5

5. Addition of interstitial oxygen atoms. Adding one oxygen atom to every

fifth or sixth unit cell in the layer will lead to a superstructure with a

larger average spacing between the cations. The negative charge of the

additional oxygen must be balanced by additional positive charge at the

cation sites.

4. CONSTRUCTING THE LAYER SEQUENCE

Local charge neutrality will favour certain sequences. Neutral layers (e.g.

A''"'*'0) can be stacked in any order but, in the copper oxides with which this

paper is chiefly concerned, the layers usually contain copper and therefore

carry a charge of -2, more or less, "depending on the Cu oxidation state. This

requires the presence nearby of positively charged layers such as A"*"*"^©, A"'"''"''" or

A"*"*". Among the sequences that provide charge neutrality are (Cu02.A"^+.Cu02.A'^'^) and (A'^'^"''0.Cu02.A"*"*+0), both of which are commonly

founcT. The coordination number of Cu is 4 in the first sequence and 6 in the

second, both acceptable coordination numbers for copper. The coordination

number of A is 8 in first sequence (Ca and Sr are candidate cations, see Fig. 2)

and 9 or 10 in the second, the obvious candidate for this site being La (AOCN =

8.5).

Sequences of AO layers give rise to a rocksalt structure and can be referred to

475
as 'An' sequences where n is the number of layers in the sequence. The

alternating CUO2 A.CUO2.A.CUO2 sequences are perovskite-like and can be

referred to as 'Bm' sequences where the total number of CUO2 layers in the

sequence is m and the number of A layers sandwiched between the CUO2 layers is (m-1). A third type of sequence (A.O2.A) can be derived from an A2 sequence

by rearranging the O atoms. It has the fluorite structure and can be labelled 'Fn',

where n is the number of A layers in the sequence. To maintain the fluorite

structure, an F sequence will always be bounded by BO2 layers. Most copper

oxide layer compounds are composed of combinations of A, B and F sequences

and can be labelled 'An-Bm' etc. Any such sequence must be electrically neutral

and must contain coordination environments appropriate to the cations present.

Some examples of observed sequences are: A2-BI AO.AO.BO2 La2Cu04

A2-B2-F2-B2 AO.AO.BO2.A.O2.A.BO2 NdSrCu04

A3^-B2 AO.BO.AO.BO2.A.BO2 Ba2YCu307

A3-B3 AO.AO.AO.BO2.A.BO2.A.BO2 TlBa2CaCu309

A4-B2 AO.AO.AO.AO.BO2.A.BO2 Tl2Sr2Ca2Cu20g

The superscript in the third example indicates that one layer in the AO sequence

has been replaced by a BO layer. The A cations in the B sequences are 8

coordinate (from Fig. 2, Ca"^"*^, Sr"*""^, Pb"*"^, La"^"^"*" and Y^"^"^ are candidates). The

A cations at the ends of the A sequence are 9-coordinate (La"'"'"'*', Sr''"'' and Ba''"*' are candidates). Those in the middle of the A sequence are 6-coordinate (Tl''^''"'", gj+++^ Pb''"'' and Ca''"'^ are candidates). The 6-coordinate layers, which

contain the smallest cations, will be the most strained by the requirements of

commensurability, and any cations (e.g. Ca and Y) that can occupy sites with

higher coordination numbers will tend to do so.

5. A SIMPLE MODEL OF THE RELAXATION: THE A4-B2 STRUCTURES

All the layers are subject to some strain as a result of the different cation sizes,

but the strain is seen most dramatically around Bi and Tl. Although both cations

have oxidation states of +3, Bi contains a lone pair of electrons in its valence shell and so responds to the strain differently from Tl. Fig. 2 shows that the TIO layer

has to expand by 0.64A to match the lattice parameter of the CUO2 layer. In Tl2Ba2CaCu20g the largest relaxation observed is a reduction of the TIO-TIO

inter-layer spacmg from 2.25A to 2.00A resulting in an expansion of the in-plane

spacing by 0.40A (Fig. 3). Transfer of about 0.20 electrons from Cu to Tl results

in a further accommodation of 0.1 OA (Fig. 5). The remaining relaxation is achieved by the Tl atom going off center in its coordination sphere in a way not

yet fully understood, an effect that is both assisted and complicated by the

presence of Tl''' produced by the charge transfer.

The macroscopic effects of the strain relaxation are thick, chunky crystals

resulting from the strong interplanar bonding, and superconductivity resulting

from the holes on the CUO2 layers. In the observed structure the bond valence

sum around Ba is larger than expected (2.2 valence units (v.u.) instead of 2.0) and

that around Cu is smaller, suggesting that there is a residual strain in which the Ba

is in compression and the Cu in tension.

476
The change required in the lattice spacing of the nominally isostructural

Bi2Sr2CaCu20o 2 is smaller, totalling only 0.50A (Fig. 2). The lone pair

distortion oT the Bi environment relaxes about 0.2A of this (Fig. 4) and the rest is relaxed by the incorporation of about one extra O atom in every tenth unit cell. The additional oxygen raises the oxidation state of Cu to +2.2 and hence leads to superconductivity, but, in contrast to the Tl compound, the lone pair environment

of Bi leads to very weak bonding between the BiO layers making the crystals thin

and flaky.

These two examples show that even in nominally isostructural compounds the

strain can be relaxed in different ways, leading, in general, to different physical

properties. The similarity in the superconducting properties is deceptive; quite

different mechanisms are responsible for its occurrence.

6. A MORE SOPfflSTICATED APPROACH TO RELAXATION: THE

KIRCHHOFF EQUATIONS.

The above analysis is rather crude, being based on a model in which all the

bonds around a cation are initially assumed to be equal. This will not, in general,

be the case even in a completely unconstrained crystal because the bonding around

the anions will be different. A better set of starting bond lengths can be

determined using the Kirchhoff rules given in Section 3 above. This approach

will be illustrated with two examples.

6.1 THE A2-B1 STRUCTURE: La, o-rSrn ,«jCuO^

a b

Figure 6

La^CuO^ (Fig. 6a) is a stoichiometric compound in which the copper is divalent.

The bond valence sums (c.f. Table I Col. C) indicate that it is highly strained with

the lanthanum in tension and the copper in compression. The trivalent lanthanum

can be substituted by the larger divalent strontium atom, simultaneously increasing

the oxidation state of Cu, both of which effects help to relieve the strain. The

Kirchhoff equations are used to predict the bond lengths and bond valence sums shown in Col. A of Table I for the doped sample.

The strain is relieved in part by bringing the two LaO layers closer together. The

consequences of shortening the inter-layer La-O bond to the observed length and

477
adjusting the in-plane La-02 bonds to give the correct valence sum around La are shown in Col. B. The in-plane La-02 bond is lengthened^ but only by about a third of the required amount. The observed bond lengths are shown in Col. C

and the residual strain in Col. D. It is seen that considerable strain remains; the

in-plane La-02 bonds are stretched by 0.14 A and the Cu-Ol bonds compressed

by 0.08A.

Table I A B C D

Bond Lengths

La/Sr-02 2.53A 2.37A 2.37A O.OOA

La/Sr-Ol (x4) 2.65 2.65 2.64 0.01

La/Sr-02 (x4) 2.53 2.59 2.73 0.14

Cu-Ol (x4) 1.97 1.97 1.89 0.08

Cu-02 (x2) 2.40 2.40 2.40

Valence Sums

La 2.92 2.92 2.58 -0.34

Cu 2.15 2.15 2.62 +0.47

Ol 2.00 2.00 2.28 +0.28

02 2.00 1.85 1.61 -0.39

compression

tension

compression

(assumed)

(Col C - Col A)

Table L Bond lengths and valence sums in La2.05SrQ i^CuO^.

Col A: Predictions of Kirchhoff equations (the long Cu-02 bond is fixed

at the observed value to simulate the Jahn-Teller distortion. Col B: After

relaxation of the short La-02 bond. Col C: Observed . Col D: Residual

strain (Col C - Col B).

6.2 THE A3^-B2 STRUCTURE, Ba2YCu30j^ (x = 6-7)

Modelling this structure (Fig. 6b) is more complex than modelling La2Cu04

because, in order to use the Kirchhoff rules, one needs to make some assumptions

about the distribution of charge between the two Cu atoms .

Cul

Figure 7. Modelled in-plane lattice spacings for the different layers of

Ba2YCu20j^ assuming different distributions of the charge, a) for x = 7

478
and b) for x = 6. The observed spacing is indicated by the broken line.

Fig. 7a shows the in-plane lattice spacings predicted for the BaO, CuO (Cul) and

CUO2 (Cu2) layers for various distributions of charge between Cul and Cu2. A

distribution with a charge of +2.33 on all copper atoms leads to complete

commensurability between all three layers and this is the structure observed.

However, a similar modelling of Ba2YCu20g (Fig. 7b) gives a different result. The charge transfer required to achieve commensurate layers would leave the

2-coordinate Cul with a charge of +2.3 and the 5-coordinate Cu2 with a charge of

+1.35, both of which are incompatible with the coordination environment. In this case the observed oxidation states remain close to those expected from the

coordination numbers (+1 for Cul and +2 for Cu2). As a result the bonds are

strained. This is shown by the average bond valence sum around copper being 0.2

v.u. too large (Cu-O in compression), and that around Ba being 0.2 v.u. too small

(Ba-O in tension). In diis case there is no effective relaxation mechanism and a

residual strain remains . Between these two extreme compositions, one can predict the charge

distribution by noting that, as each oxygen atom is added to the Cul layer, it forms Cu-O bonds that contribute only about +1 v.u. to the oxidation state of Cul

while adding -2 to the charge of the crystal. The additional positive charge is supplied by Cu2. The predicted oxidation states of the two copper atoms (the

lines) are compared with those observed (crosses and open figures ) in Fig. 8.

2-5r—I—I—1—I—I—I—I—I—I

\

Figure 8 Figure 9

7. CONCLUSION

La2Cu04 becomes superconducting when holes are introduced into the CUO2

layer oy doping the crystal with Sr, but in some structures the size of the cations,

and hence the natural incommensurability of the different layers, results in

relaxations that also introduce electron holes into the copper planes. Since the

presence of holes is a sine qua non for superconductivity there is a natural

relationship between the expectation of strain in the bonds and the presence of

superconductivity.

479
Studies of several systems^' ' ' have shown that Tc reaches a maximum at a hole concentration of about 0.2 to 0.3 holes/Cu, falling off as the hole

concentration is either decreased or increased from this value (see Fig. 9).

Recently de Leeuw et al. have shown that the maximum value of Tc (Tc(max))

in these systems is higher for those compounds in which the bond valence sums

around the Cu atom are smaller than expected and lower for those compounds in

which the valence sums are larger. The structures of many of these crystals are not accurately known and the calculation of the valence sums by de Leeuw et al. was fairly crude, but, since the valence sums should all be close to 2.2 - 2.3, the

differences which they calculate (ranging from 2.06 to 2.53) suggest that Tc(max)

is determined by the presence of residual internal strain . If the CuO^ layers are

in tension (valence sums around Cu low) Tc(max) will be above 90K; ir they are

in compression (valence sums at Cu high) Tc(max) will be lower. As discussed

above, Ba2YCujOy, which has Tc(max) = 90K, has its strain fully relaxed. On the

other hand Tl2Ba2CaCu20g, with Tc(max) = 125K, has residual strain in which

the Cu02 layers are in tension. The valence sum around Cu in Laj c5Sr0 i^CuO^

is 2.62 (Table I), a surprisingly high value that is indicative of considerable

compression. Tc(max) for this structure is only 35K.

7. ACKNOWLEDGEMENTS

I would like to thank the Natural Science and Engineering Research Council of

Canada for a research operating grant.

8. REFERENCES

1. LD.Brown Acta Crvst. B44 (1988) 545-553.

2. LD.Brown, Phvs. Chem. Miner. 15 (1987) 30-34.

3. LD.Brown and D.Altermatt Acta Crvst. B41 (1985) 244-247.

4. LD.Brown and R.D.Shannon, Acta Crvst. A29 (1972) 266-282.

5. A.W.Hewat, E.A.Hewat, J.Brynestad, H.S.Mook and E.D.Specht, Phvsica C 152

(1988) 438-444.

6. J.M.Tarascon, Y.LePage, P.Barboux, B.G.Bagley. L.H.Greene, W.R.McKinnon,

G.W.Hull, M.Giroud and D.M.Hwang, Phvs. Rev. B37 (1988) 9382-9389.

7. Y.LePage, W.R.McKinnon, J-M.Tarascon and P.Barboux, Phvs.Rev. B40 (1990)

6810-6816.

8. S.Pei, J.D.Jorgensen, D.G.Hinks, B.Dabrowski, P.Lightfoot and D.R.Richards,

Phvsica C (1990) submitted.

9. LD.Brown J. Solid State Chem. (1990) submitted

10. J.D.Jorgensen, B.W.Veal, A.P.Paulikas, L.J.Nowicki, G.W.Crabtree, H.Claus

and W.K.Kwok. Phvs. Rev. B41 (1990) 1863-1877.

11. R.J.Cava, A.W.Hewat, E.A.Hewat, B.Batlogg, M.Marezio, K.M.Rabe,

J.J.Krajewski, W.F.Peck Jr. and L.W.Rupp Jr. Phvsica C 165 (1990) 419-433.

480
12. J.B.Torrance, Y.Tojura, A.I.Nazzal. A.Bezinga, T.C.Huang and S.S.P.Parkin,

Phvs. Rev. Letts 61 (1988) 1127-1130.

13. Y.Tokura, J.B.Torrance,T.C.Huang and A.I.Nazzal, Phvs. Rev. B38 (1988)

7156-7159.

14. W.A.Groen, D.M.de Leeuw and L.F.Feiner, Phvsica C 165 (1990) 55-61.

15. D.M.de Leeuw, W.A.Groen, L.F.Feiner and E.E.Havinga Phvsica C 166 (1990)

133-139.

16. LD.Brown, Phvsica C (1990) in press.

481
DISCUSSION

R. Cava: There are two experimental observations that we made in oxygen deficient 123.

These involved discontinuities in two things that are independent of the model. One of

these things is the bond length between the copper in the plane and the apical oxygen, and

the other thing is the variation of T^ as a function of oxygen content, both of which show

discontinuities at places where T^ changes, either an abrupt change where it goes away

or a gradual change where it goes from 90 K to 60 K. And I know you're aware of that. The picture you showed for your corrected values of copper valence implies a continuous

variation of bond valence sum at the Cu(II) site as a function of x. You must have some

explanation as to why the properties are discontinuous when you have an apparently

continuous change in valence.

I. D. Brown : There are three circles in Fig. 8 that are critical and which don't lie on that

line, the points for 6.0 < x < 6.5. And the reason is that you were studying a rather

ordered sample. What happens is that there is a phase at x = 6.5 in which the oxygen

atoms are ordered. And because your sample is particularly ordered it means that you

create this ordered phase. And that ordered phase will have a higher T^. (60K). This

phase comes in, in your case, a little earlier and stays around a little longer as x is varied. In the Jorgenson study the samples were prepared a different way, and are more

disordered. And, in fact, I have modeled the T^^s based on the idea that the presence of

the ordered phase is responsible for the discontinuity.

P. Lightfoot: In support of this picture, if you look at our samples which have been

prepared without allowing the lattice to relax at room temperature, we do not see evidence for a discontinuity until we heat the samples back to two to three hundred

degrees, so probably your model is closer to such an unrelaxed state.

I. D. Brown: That's right. The model is completely unrelaxed, it represents a

completely disordered state.

M. O'Keeffe: Your prediction for the change of T^, with stress or strain can be mirrored

by the effect of the pressure on the T^; is that, in fact, true?

I. D. Brown: Yes, we looked at that, but the problem is that you're working in one

system. That other graph that I showed (not reproduced in the printed text) represents

only the maxima in Fig. 9 plotted against the bond valence sum at Cu. If you take any one sample and you subject it to pressure (for instance, if you subject Ba2YCu30^ to pressure) you change the charge distribution. The distortion theorem predicts that what

will happen when you compress the 123 compound is that the oxygen that bridges the

planes and the chains, which is asymmetric to begin with, will actually become more

symmetric. The effect of that is to transfer holes into the planes and therefore to move

you to the right in Fig. 9. And if you make a sample in the left hand part of the curve

then you get quite a large change in T^,. If you make a sample near the top of the curve

the change in T^, with pressure is very small because you're on the top of the curve.

482
A. Navrotsky : Going back to your theorem that if the bond lengths deviate from the

average, the average will then deviate positively to a longer bond length, that seems to me to fundamentally mirror the nature of the interatomic potential and its anharmonicity;

that, in fact, it's easier to stretch than to compress. Is there any way of getting

quantitative data on the nature of that potential from looking systematically at bond length

variations?

I. D. Brown: If you look at the expression relating bond valence and bond length you

can see that the bond valence is very like a Bom repulsion potential. So basically, in a

sense, what we're looking at here is the repulsive term as we bring two atoms up to each

other. If we put more electrons into the bond we compress it. We press the two atoms

closer together and the distance is determined really by the repulsion between the atoms.

So there is some relationship here. Bob Shannon and I, in the first paper in this series

(Acta CrySt. A29, 266 (1973)), showed that there was a relationship between that

exponent and the Bom exponent from pressure measurements.

B. Dabrowski: Relating to dependence of and the amount of ordering, there are

several papers which showed the same behavior. In a recent paper, Jim Jorgenson and

coworkers (Physica C 167. 571 (1990)) showed room temjjerature relaxation after

disorder, changing T^. For this composition it takes a period of 20 to 30 days, and it involved local ordering of oxygen around copper in the chain. There is agreement

between copper band-valence differences when the data was taken after the samples

relaxed.

The other problem is oxygen content, x. One of the Xs in Fig. 8 suggests there is some

impurity. Low oxygen content samples may therefore shift in the graph and make it much steeper.

B. Cava: It is well established that there is a distinct difference between samples that are prepared by quenching from a high temperature and samples that are annealed over long

periods of time at low temperature. We're not the only people that now see this very

sharp disappearance of superconductivity. There's data from Switzerland, as well, that

shows extremely sharp changes of T^ vs x when the samples are annealed for a very long

time. It all has to do with the perfection of the ordering.

I. D. Brown: I think that's right. So the graphs I've shown of T^ vs x (not given in

text, see Fig. 8 in I. D. Brown, J. Solid State Chem. in press) are what one would expect

just for a completely disordered stmcture. What you're saying is, when you allow these

things to relax they order and this discontinuity begins to appear. Is that right?

B. Dabrowski: Right. I do observe this taking place, also. At room temperature it takes

about 30 days. It may take two months.

I. D. Brown : Yes, it takes longer.

483

National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

EMPIRICAL METHODS IN OXIDE CRYSTAL CHEMISTRY

Michael O'Keeffe

Chemistry Department, Arizona State University,

Tempe, AZ 85287, U.S.A.

INTRODUCTION

It is clear that Solid State Chemistry in general,

and oxide chemistry in particular, will remain an experimental science for the foreseeable future. We cannot

predict what compositions will occur as stable compounds

and we cannot predict what crystal structures they will

adopt when they do occur. Our deficiency in the latter

regard has been called a "continuing scandal" (1) perhaps

without a full appreciation of the daunting complexity of

the crystal structures sometimes found for compounds with

simple chemical compositions. While awaiting the millennium, we must make do with empirical and semi-empirical

rules that allow some generalizations about crystal

chemistry; it is with a review of some of these from a personal, possibly minority, perspective that this paper

is concerned.

In particular, although the goal of structure

prediction is still over the horizon, we can make some

statements about expected bond lengths, and to a lesser

degree about bond angles and coordination numbers in

crystal structures . -

485
THE BOND VALENCE METHOD

Until recently bond lengths in crystals were

described as sums of a bewildering variety of radii of

limited and specialized applicability {covalent , ionic,

metallic, etc. radii) . More recently the bond valence

method, which has a long history, has come to the fore as

it presents a number of distinct advantages over the use

of radii both for the prediction and interpretation of

bond lengths in crystals (2) . These advantages become

particularly apparent in crystals with irregular coordination polyhedra and/or mixed anions. The method utilizes

two empirical observations.

The first of these, associated particularly with the

names of Pauling and of Zachariasen, is that the bond

valence, defined such that the sum of the valences of the

bonds from a given atom is equal to the atomic valence,

is a unique function of the bond length. Recent work

(3,4) has used a "standard" form for the bond length dij

of a bond of valence v±j:

dij = Rij -blnvij (1)

Here i? is a "universal" parameter equal to 0.37 A, and

Rij is a parameter unique to the pair of atoms i and J

forming the bond. Rij is, in a formal sense, the length

of a single (unit bond order) bond; however the range of

applicability of eq. (1) is strictly limited to the range

of bond valences normally encountered in crystal structures. Extensive tables of Rij are now available (3,4) .

Except in the simplest cases, the constraints that

bond valence sums are equal to atomic valences are insufficient to uniquely determine individual bond valences.

An important step was taken by Brown (5) in introducing

486
the second principle {Brown's Ansatz) that, subject to

those constraints, bonds between pairs of atoms of a given kind will be as nearly equal as possible. For

example in a complex magnesium oxide, bonds from a given

Mg atom to crystallographically-dist inct O atoms will be

as nearly equal in valence as possible. It transpires

that this principle always results in an unique set of

expected valences (2) and a computer algorithm for deriving them for complex crystals has been described (6) .

IMPLICATIONS OF BROWN'S ANSATZ

The principle that bond valences for bonds from an

atom to others of the same chemical kind are as nearly

equal as possible suggest the following corollary: structures in which bonds from an atom to chemically-equal

atoms are forced by bond valence constraints to be very

unequal are unlikely to be stable. I (2) consider this to

be an updated and more-generally applicable restatement

of Pauling's celebrated electrostatic valence rule. There

is space for just one simple example of its use.

In the sesquioxides M2O3 the average coordination

number of O is 2/3 that of the metal atom M. Thus unless

the coordination number of M is 3, 6, 9,... there must be

oxygen atoms with more than one coordination number. In

particular for 4 -coordinated M (coordination number

indicated by a preceding superscript roman numeral) one

must have ^^M2-^-'--'-0 ( 1 ) 2'^"'"0 (2 ) . The valence sums require

that M-0{1) bonds have valence 2/3 and that the M-0(2)

bonds have a valence 1 . It is noteworthy that the only

known sesquioxide with tetrahedrally-coordinated metal is

in the high-pressure form of B2O3 which indeed has has

long and short bonds in the BO4 tetrahedron. (Note in

passing that the ionic radius approach would predict four

bonds of equal length and that the zero pressure

487
structure of B2O3 has 3-coordinated E) . On the other hand

there are at least 33 sesquioxides with 6-coordinated

metal atoms. There are also five compounds with 7-

coordinated metal atoms {A rare earth structure) but

these have very anomalous bond lengths (2)

.

EXCEPTIONAL ELEMENTS

Sufficient examples have now been adduced in the

literature to show that the principles outlined above

successfully predict bond lengths in a wide variety of

crystals . It is therefore rewarding to identify the cases

where the Brown ansatz fails to predict observed bond

lengths. Most of these cases can be identified with three

kinds of cation.

(a) " Jahn-Teller" cations. These are notably d^ and

d^ electronic configurations of the transition elements

[e.g. Mn(III) and Cu(II)]. Here configurations such as

octahedra with four short and two long bonds are found

for generally well-understood reasons.

(b) "Lone pair" cations. These are characterized by

a ns^ valence configuration found in lower oxidation

states of the heavier elements of groups 13-16. T1(I),

Pb(II) and Hi (III) are common examples.

(c) "d^" cations. This category refers to the early

transition elements (groups 3, 4, and 5) in their maximum

oxidation states. Examples are Ti(IV), Nb(V), W(VI).

Oxides of these elements often (but not always) contain a wide range of bond lengths even when not required by

valence sum constraints. A striking example is afforded

by the low temperature structure of WO3 with each O

bonded to two metal atoms, but which contains a wide

488
range of W-0 bond lengths in contrast to the topologically-identical structures of ReOs (d^) and TeOs (d^^) in

which all bond lengths are equal.

The predilection of cations in this third category

to irregular coordination is manifest in their ubiquity

in ferro- and ant iferro-electric crystals with distorted

MOq octahedra . It is also in accordance with their occurrence in structures with oxygen atoms with different

coordination numbers. Notable among these last are the

"shear" structures such as Tij50277-i (with ^^^O and ^"^O)

and W^Osn-i (with ^^O and ^^^0) . It possibly relevant also

to oxidation catalysts which require an easy transition

between M=0 and M-O-H at the surface

.

BONDING RADII

It has recently been found (7) that bond valence

parameters Rij for a wide variety of bonds can be

expressed to a good approximation in terms of two atomic

parameters and Ci

:

Rij = ri + rj - [rirj(Vci-Vcj)2] / (ciri+cjrj) (2)

An expression of this form was proposed in a somewhat different context by Ray, Samuels and Parr (8) in

which the parameters c were the atomic electronegativities. It was indeed found that the empirically-determined

c parameters did in fact correlate very well with atomic

electronegativities. The parameters r in Eq. (2) are

interpreted as bonding radii, with the last term being a correction for bond shortening in heteropolar bonds . This

last term is generally rather small, with an average

value of 1% of -Rij and a maximum value of 11% for Cs-F

bonds

.

489
In determining the parameters in Eq (2), no data

were included for metallic bonds, nevertheless it was

found that the derived parameters predicted the bond

lengths in elemental metals rather well (largest error

10%). This suggests that the bonding radii, together with

the "electonegat ivity" parameters, can be used for all

types of bond (covalent, ionic, metallic) . It should be

noted that as data for homopolar bonds are included,

there is not an arbitrary division into "cations" and

"anions", and one has an unambiguous scale of "size".

Typical bonding radii, r, are (in A)

Li Be B C N O F

1.00 0.81 0.79 0.78 0.72 0.63 0.58

Na Mg Al Si P S CI

1.36 1.21 1.13 1.12 1.09 1.03 0.99

NON-BONDED INTERACTIONS

It might generally be conceded that coordination

numbers in crystals are controlled to some extent by

next-nearest (non-bonded) interactions. However, the

relative importance of cat ion ... cat ion and anion ... anion

interactions in crystals such as oxides remains a somewhat contentious issue. It is my opinion that the role of

cation ... cat ion interactions has generally been greatly

discounted: some reasons for this opinion are summarized

here .

Early attempts (9) to characterize non-bonded

interactions focussed on A. . .B distances in A ^ B moities

("geminal" or "one-angle" configurations) in molecules.

It was found that, to a good approximation, the distances

could be expressed as a sum of non-bonded radii.

490
Extensions to crystals (JO) followed; a set of such radii

for comparison with the bonding radii given above follows

(in A) .

.

Li Be B C N O F

1.5 1.35 1.26 1.25 1.14 1.12 1.10

Na Mg Al Si P S CI

1.68 1.66 1.62 1.53 1.46 1.45 1.44

The scale of these radii has now been provided with

a sound theoretical basis (11) . As discussed below, I feel that they are now mainly of historical interest,

although they do illustrate the same periodic trends as

the bonding radii. In particular O and F are again the

"smallest" atoms.

Another observation that leads to the belief that

metal. . .metal interactions are of importance in oxide

crystal chemistry is that many oxide structures, including some apparently complex ones, may be elegantly

described as a "stuffing" of simple cation arrays {12) . Indeed, in far too many cases to be merely a coincidence,

the metal arrays in complex oxides are those of common

alloy types. [The noted crystal psychologist, P. B. Moore

(13) has remarked that, if these observations are to be

ascribed to coincidence, we must "declare that Nature's

dice are loaded"]

.

Particularly when the crystal- and thermo-chemist ry

of compounds rich in alkali and alkaline earth elements

is examined, one is lead to the observation that these

elements behave as if they strongly repel each other.

This is manifest in unusually long (14) and weak (15)

bonds in oxides and related crystals involving these

elements. This idea, coupled with a realization of the

491
important role of stoichiomet ry , leads to an explanation

(16) of the occurrence of unusual oxidation states and

coordination numbers in such metal-rich compounds, as

well as the apparent nonexistence of compounds such as

NasN. The importance of large "cations" in stabilizing

high oxidation states in copper oxide superconductors has

also been emphasized (17) .

SIMULATING CRYSTAL STRUCTURES

This not the place for a review of the extensive

literature on simulation of crystal structure. However,

it needs to be stated that the information content of

structure is generally less than commonly supposed. Thus

the apparently complex crystal structure of the garnets

A2B2C2O12 is completely determined by four parameters

(the cubic cell edge and oxygen coordinates) and one can

already get a good approximation (better than many

"sophisticated" simulations) to the structure simply by

using metal-oxygen distances from tables of ionic radii

{10) . To test simulation methods, particularly if they

are to lead to any useful generalizations, one must do

very well in predicting both interatomic distances and

lattice parameters. A possibly useful approach to the

problem, still actively under development is outlined

here

.

As a starting point, for a given topology, one can

predict expected valences and hence bond lengths. In

exceptional cases, these are sufficient to determine a crystal structure (see the contribution by I. D. Brown to

these proceedings) . In this case one could minimize the

difference between calculated valences Vq and predicted

valences Vp . This corresponds to minimizing for all

bonded distances dk the sum

492
Ijt I exp [ (Kij-djt) /b] - Vp (3)

More generally there will be more structural parameters than bond lengths. A simple example is afforded by

the familiar corundum structure of AI2O3 which has two

lattice parameters and an 0 parameter, but from the point

of view of the bond valence method, only one predicted

Al-0 distance. There are in fact two independent (and in

this case significantly different) Al-0 distances, and to

reproduce them one needs to include the effect of 0...0

and A1...A1 repulsions.

Non-bonded repulsion energies are commonly, and with

some theoretical justification, expressed as ill):

Here djc is the kth non-bonded distance and A^j and a±j are

constants for the atom pair under consideration and have

been determined for a number of more electronegative

element pairs empirically from molecular properties [see

(11) for references] . Minimizing the repulsion energy can

be expressed as minimizing the sum

Where Sij ^ aijlnAij. The expression (5) is written in

this way to illustrate the suggestive similarity to (3)

.

To determine an optimal structure (3) and (5) have to be

minimized simultaneously.

One needs to find parameters S and a for atom pairs

that reproduce known crystal structures, and a start has

been made. A large number of structures has to be examined and a consistent set of parameters extracted - a large task. To test whether the proposed method has any

Ek = Aijsxp i-dk/aij) (4)

Ijcexp [ iSij-dk) /aij] (5)

493
merit, to date a somewhat arbitrarily chosen constant

value of a = 0.25 A has been chosen (in the same spirit

as a constant b = 0.37 A is used for bonded interactions)

and S found for non-bonded interactions. Typical values

are SaiaI = 3.0 A and Sqq = 2.7 A.

As those who have attempted to model structures by

similar methods (e.g. so-called DLS methods which are

close in spirit) will know, the derived structures can be

very sensitive to the values of the input parameters. It

is observed that for many oxide structures, especially

those with several different types of cation, that the

valence sums at O are very close to 2 . 0 even though the

sums at the individual cations deviate significantly from

the atomic valences. I have found that the structure

simulation procedure for oxides is much more robust (and

significantly more difficult to implement) if one

constrains the valence sums at O atoms to be exactly 2.0.

As a hint of the' possible utility of the method

very preliminary results (Table 1 . ) are given for the

simulation of three A3B2C3O12 compounds with the garnet

structure by this method. In this work a constant value

of 5 = 3.2 A was used for all metal ... metal interactions

and S = 2.7 A for 0...0 interactions. Despite the

simplicity of the model several distinctive features of

the structures are reproduced. For example the short Al-0

bond lengths in Mg3Al2Si30i2 and the long Al-0 bond

lengths in Ca3Al2Si30i2 (compare the contribution of R.

D. Shannon to these proceedings) are reproduced. To be

noted also is that the two independent A-0 bonds in the

structure are calculated to have significantly different

lengths in accord with observation, even though the

simple bond valence method predicts that they have the

same valence. Finally the low valence sum at Mg and the

high valence sum at Ca is reproduced.

494
Table 1. Comparison of observed and simulated (calc)

structures for three garnets, a and x, y, z are the

lattice parameter and oxygen coordinates. Va is the

valence sum at Mg, Ca or Sr

Mg3Al2Si30i2 Ca3Al2Si30i2 Sr3Y2Ge30i2

calc obs calc obs calc obs

a 11 .412 11 .546 11 . 912 11 .846 13 .081 13 .087

X -0 .0347 -0 .0329 -0 .0433 -0 .0381 -0 .0434 -0 .0395

y 0 .0502 0 .0502 0 .0459 0 .0449 0 .0542 0 .0517

z 0 . 1547 0 . 1534 0 . 1510 0 . 1514 0 . 1591 0 . 1591

d(A-O) 1 2 .198 2 .197 2 .390 2 .320 2 . 603 2 .551

d(A-O) 2 2 .340 2 .342 2 505 2 .503 2 .661 2 . 684

d(B-O) 1 .898 1 .887 1 949 1 . 933 2 .270 2 .250

d(C-O) 1 . 603 1 . 634 1 624 1 . 653 1 .748 1 .768

1 .72 1 .72 2 21 2 .48 2 .00 2 . 12

It is interesting that in this example, as in many

others, it is found that the contribution of

cat ion ... cation and anion ... anion repulsions to the total

repulsion energy is comiparable, and that omitting one or

the other invariably leads to unsatisfactory structures.

Thus the answer to the question of whether cation repulsions or anion repulsions are important in oxide chemistry is (as is so often the case with such questions)

that both are important and that both must be considered.

495
Acknowledgment . This work was supported by a grant

(DMR 8813524) from the National Science Foundation.

REFERENCES

1. J. Maddox, Nature, 335, 201 (1988).

2. M. O'Keeffe, Structure and Bonding, 71, 162 (1989).

3. I. D. Brown and D. Altermatt, Acta Crystallogr. B41,

240 (1985)

.

4. N. E, Brese and M. O'Keeffe, Acta Crystallogr. in

press

.

5. I. D. Brown, Acta Crystallogr. B33, 1305 (1977).

6. M. O'Keeffe, Acta Crystallogr. A46, 138 (1990).

7. M. O'Keeffe and N.E. Brese, to be published.

8. N. K. Ray, L. Samuels and R. G. Parr, J. Chem. Phys

.

70, 3680 (1979)

.

9. L. S. Bartell, J. Chem. Phys. 32, 827 (1960).

10. M. O'Keeffe and B. G. Hyde, in Structure and Bonding

in Crystals (M. O'Keeffe and A.Navrotsky, eds . ) Wiley,

New York (1981) . 11. M. A. Spackman, J. Chem. Phys. 85, 6579 (1986).

12. M. O'Keeffe and B. G. Hyde, Structure and Bonding,

61, 77 (1985)

.

13. P. B. Moore, Mineralogical Record, 20, 339 (1989).

14. N. K. McGuire and M. O'Keeffe, J. Solid State Chem.

54, 49 (1984)

.

15. M. O'Keeffe and J. A. Stuart, Inorg . Chem. 22, 111

(1983) . 16. M. O'Keeffe and B. G.Hyde, Nature, 309, 411 (1984) . 17. M. O'Keeffe and S. Hansen, J. Am. Chem. Soc . 110,

1506 (1988)

.

496
DISCUSSION

R. Cava: Why is it that the D zero ions have this peculiar property of not working?

M. O'Keeffe : I don't know really, however, I can give you a glib answer. In a TiOg

regular octahedron there are empty states, the D manifold, which are non-bonding in

octahedral symmetry. If the symmetry is reduced these states mix in with the bonding

states, the unbonded energy states go up and the bonded energy states go down,

producing a 2nd order Jahn-Teller effect. The exact distortion type cannot be predicted,

but any distortion from octahedral symmetry will make the non-bonding states mix with

the bonding states. Then if there is one electron in the d-band the whole statistics of the

occupied states goes up. Presumably as you go further across the d-level the anti-d-states

get higher and higher energy.

A. Nath : The deviations which you mentioned in garnet, for example, would they be

explained by the features which you mentioned in your last slide?

M. O'Keeffe: No, I think they're fully explained just by the repulsions within the

crystal, the metal-metal and oxygen-oxygen repulsions.

A. Nath : That would explain all the deviations?

M. O'Keeffe: That's what I tried to show in my simulations. In this calculated structure,

the valence sum at oxygen is constrained to be two. That is the only arbitrary nature of

the calculation. We have predicted bond lengths for all the bonds in the structure, or

predicted valences, and we also have non-bonded interactions. What we do is minimize

the deviation from the predicted valences and the non-bonded repulsions together. What

you come up with is the observed structure.

I. Brown : In that last calculation do you use the idea of equal valences between the

bonds?

M. O'Keeffe: Yes. For example, calcium is eight coordinated and all the bonds should

be the same. Each bond should have a valence of a quarter, so the deviations from a

quarter are what is implicated.

I. Brown : Then you constrain the oxygen to be two and after that put in the repulsions.

M. O'Keeffe: Yes. Where the parameters come from, of course, is important. We have

looked at a large number of structures like corundum to find out what sort of aluminumaluminum and oxygen-oxygen repulsions must be put in to get the observed structure. This is well defined, and these parameters are simply then transferred to more complex

crystals.

497
D. Cahen : Your earlier slides of magnesium oxide showed first a cut through the

electron density along the magnesium oxygen bond. I assume it was the valence electron

density?

M. O'Keeffe: No, total electron density. That was on a logarithmic scale.

D. Cahen: So afterwards you took the total electron density in your equation to get the

potential?

M. O^Keeffe: Yes.

D. Cahen : I wonder what would happen if you can, in a pseudo-potential like way, just

take the valence electron density.

M. O'Keeffe: You get very similar plots actually. The core electrons are so close to the

nucleus that the effect of their electrostatic potential would be very small. So you get a very similar curve.

B. Dabrowski : In your equation the single parameter is expressed as a function of Rj

,

R2; are you including <dso the dependence on thermodynamic functions like temperature

and pressure?

M. O^Keeffe: No.

B. Dabrowski: Can you explain phase transition temperatures?

M. O'Keeffe: These parameters are average room temperature zero pressure parameters.

Bond lengths don't change that much with temperature except for alkali metals where

there is very large bonding. All parameters are for zero pressure. I think that's all I can

say.

R. Roth : I would like to make a comment regarding some recent work from NIST and

DuPont on the coordination number of the Bi"'"-' ion, that relates to the subjects discussed

by O'Keeffe and Brown as well as the earlier talk by Shannon. The crystal structures of

three Bi"''-' containing phases have recently been solved by x-ray single crystal structure analyses and neutron powder diffraction Rietveld analyses. In Ca4BigOi3 (Chem. Mater.

2 454, 1990), although 2/3 of the Bi"^^ ions are in 5-fold coordination (with the lone pair

electrons presumably in the 6th direction to complete an octahedron) 1/3 of the Bi"^^ ions are in 3-fold coordination with oxygen. In Ca2Bi205, 1/3 of the Bi''"'' ions are in 5-fold

coordination with 2/3 of the Bi"^^ ions in an unusual 'saw-horse' 4-fold configuration,

with the 5-fold polyhedron bridging two 4-fold ions in a U-shaped BijO^ unit. In Sr2Bi205 all of the Bi ^ ions are in 3-fold coordination. Thus the coordination polyhedra

formed by the Bi"*"^ ion are not 6-, 8-, or 12-fold as might be expected but rather form

strange and entirely unexpected patterns. Many more such structures need to be

examined before we can feel that we understand and can predict the oxygen coordination

of the Bi + ^ ion.

498
National Institute of Standards and Technology Special Publication 804. Chemistry of Electronic Ceramic

Matenals, Proceedmgs of the International Conference held in Jackson, WY, August 17-22 1990 issued

January 1991.

COMPUTER SIMULATICN STUDIES OF ELEdKONIC CERAMICS

A.N. Cormack

NY State College of Ceramics

Alfred University, Alfred, NY 14802

INTRODUCTION

The detailed chemistry of electronic ceramics is often rather

complicated, either by "accident" or design. By "accident," we

mean that a compound with a complex stoichiometry, such as the

YBa^Cu Og or Bi Sr^CaCu^O high T ceramic superconductors, is

found to nave useful properties. By design, on the other hand, we mean that a basically simple composition, such as ZnO or BaTiO^

,

has a number of other components added to it in order to control

more precisely the desirable electronic properties.

As so often with ceramic materials, processing is a key issue

and control of the chemistry and stoichiometry is an essential part

of the processing conditions. From a crystal chemical point of

view, incorporation of minor components, or maintenance of a complex stoichiometry, each require control and hence an understanding of the defect properties of the material. The

objective of computer simulation studies is to provide a theoretical basis from which to understand defect properties,

especially in situations v^ere experimental access to these

properties is limited in one way or another. In this short review,

we will describe the methodology behind the calculations and aim to

demonstrate the efficacy of the approach using a nimber of recent

exaitples and illustrations.

METHODOLOGY

The basic physics of our type of simulations is embodied in

the Born Model of the solids, supplemented by a representation of

ion polarisability for v^ich the Shell Model, due to Dick and

Overhauser, is used. The Shell Model treats each ion in terms of

two point components: a core in which the mass of the ion is

contained on a mass-less shell which may be loosely taken to

represent the valence electrons. The total formal charge on the

ion is divided between the core and the shell which are coupled

499
together via a harmonic spring so that the polar isability (which

depends on the shell charge and the harmonic spring constant) is

modeled correctly. (In fact, what one strives to reproduce is the

dielectric behaviour of the solid which is governed by the ion

polarisabilities)

.

In addition to the coulombic term, there is a contribution to the

overall interatomic potential from short-range effects, notably the

repulsions that come into play when ions approach each other too

closely. Since most of the chemistry—by which I mean the

differing behaviour of different ions or indeed compounds, such as MgO and CaO—is predicated on the validity of the interatomic

potential the way in which the short range terms are modeled is of some importance. There is no single, correct formula; indeed, one could use computed values stored numerically in the computer. The

most common approach, however, is to use a parametrised analytical

form usually of the Born-Mayer or Buckingham type:

V(r) = A.exp (-r/p) - Cr"^

where the parameters A, p and c are appropriate for any given

interaction.

The lattice energy, from which the properties such as structural equilibrium, elastic and dielectric tensors may be

derived is then just the sum of pair potential interactions:

Defect energies are obtained, at least in principle, by summing the

interactions of the defect species with all the other ions in the

crystal. Clearly, this is impossible in practice, so a two-region

strategy is used. In this approach an inner region I, which is

treated atomistically using appropriate interatomic potential

models, is embedded in a dielectric continuum in which the

polarisation due to the net charge on the defect may be evaluated

according to Mott-Littleton theory.

A key point, both with respect to the defect energy

calculations and in determining structural equilibrium is the use

of energy minimisation (with respect to ion position) to ensure

that, in the case of the defect calculation, the structure around

the defect is in equilibrium with the defect and that, for perfect

lattice calculations, the whole unit cell is the equilibrium.

A more complete description of the computational and

theoretical details may be found in ref (1).

ALKALINE EARTH EXCESS ALKALINE EARTH TITANATES

The alkaline earth titanates are widely used in the electronic

ceramics industry in a number of different applications. A key

property to be controlled during their processing is the

500
TABLE 1 BASIC POINT DEFECT ENERGIES (IN EV)

IN ALKALINE EARTH TITANATES

Defect BaTiO^ SrTiO^ CaTi03

Vo 17.79 18.95 18.96

\ 16.73 19.82 21.15

81.77 86.52 86.26

\i 66.30 66.29 63.79

59.75 62.52 63.79

A/Ti cation ratio. Of interest in this regard is the ease with

v^ich these compounds may become nonstoichiometric and by what

mechanism. A review of the literature suggests that BaTiO^ , SrTiOj

and CaTiO each behave somewhat differently with respect to

incorporation of excess alkaline earths. In summary, BaTi03

tolerates virtually no excess Ba, SrTiO^ accommodates some via

point defect formation and more through the formation of planar

defect structures, but for CaTi03 , the situation is less clear,

with conflicting data.

Our interest here was to see whether these differing

observations could be explained in terms of a single defect model

common to the three compounds or whether different mechanisms might

be operating. Accordingly, using a consistent set of interatomic

potentials^ we calculated the basic defect energies for the three

compounds. Note that these are the energies required to produce

the defect. For example, oxygen vacancy formation involves the

removal of an oxide ion from the structure to infinity and this is

the energy recorded in table (1). Using these energies, one may

obtain the energies of the various defect reactions that may occur,

allowing one to determine which processes may be the more

favourable.

For instance, excess alkaline earths may be incorporated via

two principle mechanisms according to the reactions given in

Table 2. The first involves Ti vacancies whilst the second

produces A^,^ defects. Our results indicate that reaction (1) is

TABLE 2 REACTION ENERGIES (IN EV) FOR EXCESS

ALKALINE EARTH INCORPORATION

Reaction BaTiO^ SrTi03 CaTiO,

(1) AO -> A, + 0^ + . + 2V„ 4.14 10.25 10.11 ^ ' A 0 T 1 O (2) 2A0 ^ A^ + + 20p + 2.25 0.91 4.80

501
energetically unfavourable for all three titantates, whereas for

SrTiOj reaction (2) has a low energy of 0.91 eV associated with it.

In BaTiOj and CaTiO^ , P^^ defects are not expected to form, again

becasue of high reaction energies (see Table (2)).

Another way to consider solubility limits is to examine the

energetics of second phase formation. In the case of the

orthotitanates, A^TiO it is interesting to note that whilst

Ba^TiO^ adopts the K^SO structure, Sr^TiO^ has the K^NiF^

structure whereas Ca^TiO^ is not known as a single phase. We

calculated the energies of orthotitanate formation assuming, for

each compound, both structure types. For Ba^TiO^ we predict that

the K^SO structure will be preferred whereas for Sr^TiO^^, the

K^NiF^ structure will be adopted. For Ca^TiO^ we find a positive

reaction energy for both structure types, indicating that either

will be energetically unstable with respect to disproportionation

to CaO and CaTiO^ . All of these results concord with the

experimental observations and taken together with the point defect

calculations indicate that there are sufficient differences in the

energetics of each alkaline earth to account for their varying

behaviours.

DOPED MANGANESE FLUORIDE

In this example, the simulations were aimed at addressing two

principle questions arising from some experimental observations.

Ling and Nowick, measuring the ionic conductivity, noted a degree

of ani sot ropy that was less than expected given the tetragonal

symmetry of the rutile structure which is adopted by MnF . This

was ascribed to the mechanism of ion transport in the material but

a detailed and quantitative explanation could not be extracted from

their data. Secondly, in doped samples, but not in pure sanples,

low temperature dielectric loss peaks were observed that have

associated energies too small to be assigned to ions jumping

between lattice sites. Since these peaks were associated with

impurity (or dopant) defect structures, it seemed natural to

undertake a set of defect calculations, with the objective of

uncovering a suitable migration mechanism and elucidating the

nature of the defect structure causing the dielectric loss„

In summary, our findings^ were as follows: The intrinsic

disorder is predicted to be of anion Frenkel type, and,

interestingly, that the anion interstitial would adopt a "split

interstitial" structure, that is an anion in an interstitial site

would displace an adjacent lattice anion into another interstitial

position. Thus the interstitial would be split between two sites

bridged by a vacant lattice anion site. This has iirportant

consequences for both of the experimental observations described

above. Firstly, it provides for a low energy pathway in the basal

plan in addition to the empty channels in the c direction, thus

502
suggesting an explanation for the lack of ani sot ropy seen in the

conductivity. Secondly, the calculations reveal that when bound to an dopant ion such as y or Er^""^ sitting substi tut ionally on an Mn site, the split interstitial configuration is able to relax

further into a low synrmetry arrangement in which one of the

interstitial fluorine moves to a site above the basal plane, whilst

the other moves to a position below the basal plane with the dopant

remaining in the basal plane. A "flip-flop" orientation of this

defect conplex, switching the F ions which lie above and below the

basal plane, accounts for the dielectric loss measurements.

Note that the dopant is required for the split interstitial

structure to go off-symmetry: a somewhat subtle chemical influence

which has been elucidated from the computer simulations.

Our last example concerns a rather complex system with the

potential for an extensive range of nonstoichioimetries; hence, an

ideal subject for attack by atomistic computer simulations.

NONSTOICHIOMETRY IN LaMgAl^^ ^ 0^ ^ LaMgAl O^g has the magnetoplumbite structure, which is

related to the 3-alumina family of structures in that both consist

of slabs with the spinel structure stacked together. The basic

differences occur in the structure of the inter-slab region. These

differences can be crucial: magnetoplumbites do not exhibit cation

conductivity v^ereas ^-aluminas are very good fast ion conductors.

In this review, we will focus on just one aspect of the

crystal chemistry of these materials. The following formulae have

been variously suggested to describe the nonstoichiometry in

LaMgAl, , 0, ^ : La, MgAl, , ^ 0, „ ; LaMg Al, , 0, „ ^ ; La, Al, ~ , ^ , , . 0, „ ; 11 19 1-x ~ 11+x 19' 11 18+x' 1-x 11+2/3+x 19'

La, MgAl,, , ,,0,0 ; LaMg Al, , ^ . , , , , .0,. and 1-x ~ 11 + 5X/3 19 + x' -"x ll + 2/3(l-x) 19

La, MgAl,, ,0, „. E&ch formula represents the creation l-x-'y ll + x + 2/ 3(l-y)19

of a different set of point defects and thus, using computer

simulations to calculate the energies of these different point

defects, it is possible to give some quantitative indication of the

relative ease of producing these nonstoichiometries. The results

are not altogether surprising. For example, changing the La to Al

ratio occurs relatively easily by substituting Al on La sites, and

this nonstoichiometry can be facilitated by losing magnesium. The

formation of defect complexes can, however, change the energetics

somewhat as we discuss next.

An interesting result of our calculations concerns the

incorporation of excess oxygen. The basic principles of crystal

chemistry dictate that substituting an anion in place of a cation

is not a sensible thing to do. However, when Stevels first

503
proposed that oxygen would substitute for lanthanum, he noted that

there were likely to be additional defect interactions because of

the diffuse scattering he observed. We have found that, not

surprisingly, attempts to replace a single lanthanum by oxygen is

not energetically favourable; indeed, it is difficult to get the

oxygen to remain in the same vicinity as the lanthanum site.

However, if additional defect species are introduced, such as those

depicted in figure (1), then the 0^^ defect can, in fact, be

stabilised, and moreover, the binding energy of such a defect

cluster renders this mode of nonstoichioraetry energetically

compatible with other likely mechanisms. Thus, the corrpater

simulation studies, in addition to discriminating between possible

modes of nonstoichiometry, can also lend support to the structure

of defect complexes inferred from x-ray scattering studies.

References

1. C.R.A. Catlow and W.C. Mackrodt, "Computer Simulation of

Solids" Lecture Notes in Physics, vol. 166, Springer, Berlin

(1982).

2. K.R. Udayakumar and A.N. Cormack, J. Phys. Chem. Solids, 50, 55

(1989).

3. S. Ling and A.S. Nowick, Phys. Rev. B40, 3266 (1989).

4. A.N. Cormack, C.R.A. Catlow and S. Ling, Phys. Rev. B40, 3278

(1989) . 5. Like Xie and A.N. Cormack, J. Solid State Chem. 88, (October)

(1990) .

(I) (b) (c)

© O^Xff. iniMsiiMil oi)fCD Al vacjoc) • AJ Li vacancy • inifisnml Al m minor p'jne S spmfl b'.Ki Figure 1. Defect complexes associated with the stabilisation of

0^^ defects. Note that displacement, but not complete removal, of

a mirror plane Al, increases the stability of the cluster.

504
proposed that oxygen would substitute for lanthanum, he noted that

there were likely to be additional defect interactions because of

the diffuse scattering he observed. We have found that, not

surprisingly, atteitpts to replace a single lanthanijm by oxygen is

not energetically favourable; indeed, it is difficult to get the

oxygen to remain in the same vicinity as the lanthanum site.

However, if additional defect species are introduced, such as those

depicted in figure (1), then the 0^^^ defect can, in fact, be

stabilised, and moreover, the binding energy of such a defect

cluster renders this mode of nonstoichiometry energetically

coitpatible with other likely mechanisms. Thus, the computer

simulation studies, in addition to discriminating between possible

modes of nonstoichiometry, can also lend support to the structure

of defect complexes inferred from x-ray scattering studies.

References

1. C.R.A. Catlow and W.C. Mackrodt, "Computer Simulation of

Solids" Lecture Notes in Physics, vol. 166, Springer, Berlin

(1982).

2. K.R. Udayakumar and A.N. Cormack, J. Phys. Chem. Solids, 50, 55

(1989)

.

3. S. Ling and A.S. Nowick, Phys. Rev. B40, 3266 (1989).

4. A.N. Cormack, C.R.A. Catlow and S. Ling, Phys. Rev. B40, 3278

(1989) . 5. Like Xie and A.N. Cormack, J. Solid State Chem. 88, (October)

(1990) .

505
DISCUSSION

(Unknown') : My first question concerns the model that you used: Is it a pure ionic

model? I may have missed that. But my real question is that you showed a list with

defect formation energies. How do they compare to experimental? They seemed to be

awfully large.

A. Cormack: The individual defect formation energies bear no resemblance to what you

measure experimentally. What you measure, if you can, is the Scottley or Frenkel

disorder energy or some antisite reaction energy. These energies involve a combination

of the raw energies given in that table.

(Unknown) : There are experimental data that you can compare to.

A. Cormack: Not directly for the alkaline earth titanates. I think one thing we can compare here is what we would predict the defect concentration such as strontium or

titanium sites in strontium titanate, and our energies predict the concentration which is in the same ballpark as seen experimentally.

Just to comment on the first question about the model; I prefer to use the term Bom

model instead of a purely ionic model because the model consists of both point charges

and short range repulsive interactions and is not just purely an electrostatic model.

506
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued

January 1991.

COMPUTER SIMULATION OF DOPANT SUBSTITUTION

IN YBa2Cu307

M.S. Islam and C. Ananthamohan

Department of Chemistry, University of Surrey

Guildford GU2 5XH, UK

ABSTRACT

Atomistic simulation techniques are used to investigate dopant

substitution on the Cu sublattice in YBa2Cu307. The techniques are based on the Mott-Littleton methodology which effectively treats

lattice relaxation around defects or impurity species. Solution

energetics are calculated for both divalent and trivalent dopant ions.

Correlations between particular bond distances and Tc are also

examined.

INTRODUCTION

Since the discovery of the YBa2Cu307 superconductor, the effect of

substitution by dopant ions on its properties has been extensively

studied to investigate the origin of the superconductivity. Such studies

have demonstrated that rare-earth substitution for Y has almost no

effect on the critical temperature (Tc) [1], whereas substitution for the

Cu sites by 3d transition elements results in a decrease in Tc, the

magnitude being dependent on the particular ion [e.g. 2-7]. As is now

well known, YBa2Cu307 adopts a perovskite-like structure with two

distinct copper sites : Cu(l) chain and Cu(2) plane. In evaluating the

influence of substituting ions it is apparent that the precise lattice

location is important. In some cases, however, the preferred site has

not been clearly established, particularly for the Cu(l) versus Cu(2)

sites. This partly arises from differences in experimental conditions

and difficulties in characterizing the exact oxygen stoichiometry.

In this paper the energetics of dopant incorporation are examined

using atomistic simulation techniques. These methods have been

successful in describing defect properties of ceramic oxides [8-10] and

have been applied to the cuprate superconductors [11-14]. This study

extends our recent studies of defects [12] and oxygen migration [13] in

YBa2Cu307, by focusing on cation substitution within the Cu

sublattice.

THEORETICAL METHODS

The present account of these widely used techniques will be brief

since detailed reviews are given elsewhere [9,10]. The calculations are formulated within the framework of the Born model and are based on two-body potentials to represent short-range forces, which include the

shell description of ionic polarization. The essential feature of the

507
lattice simulations is the relaxation of ions in the vicinity of the defect or dopant. The Mott-Littleton approach is to divide the crystal into

two regions so that ions in the inner region surrounding the defect are relaxed explicitly, whereas the remainder of the crystal is treated by

continuum approximations. The potentials for YBa2Cu307 are the

same as in the previous studies [12,13], which reproduced the structure

to within 0.03 A. Full details of all the potential parameters and

computed crystal properties are given by Baetzold [12].

RESULTS AND DISCUSSION

The principal dopants or impurities in YBa2Cu307 are divalent and

trivalent cations. The simulation approach is based on the calculation

of the energetics of dissolution in the oxide, which are obtained by

combining appropriate cohesive energies with energy terms

accompanying the formation of the substitutional species.

We first consider the dissolution of divalent impurities into the

Cu2+ sublattice, which can be represented by the defect equation:

MO + Cu^^ —• M^^ + CuO (1)

The calculated energies of solution, AE, for the two nonequivalent

copper sites are listed in Table I, where our sign convention is such

that a positive value indicates an endothermic process. The

calculations suggest that the transition-metal cations Fe^*, Co2+, Ni2+,

Zn2+ and Cd2+ preferentially substitute for Cu(2). Most notably, the

highly favourable solution energies for NiO and ZnO are in line with

the observation that these oxides dissolve readily in YBa2Cu307. The

results generally agree with diffraction studies that find Ni2+ and Zn2+ on the Cu(2) plane. As expected the larger alkaline-earth ions Ca2+

and Sr2+ have high and unfavourable energies. Note that

experimentally, Fe and Co are considered to be in the 3+ charge state,

which will be examined below.

TABLE I. Calculated energies of solution (in eV per dopant ion).

MO ^^Cu(l) ^^Cu(2)

MgO 1.20 0.18

CaO 1.28 1.15

SrO 2.28 2.94

FeO 0.91 0.02

CoO 0.85 -0.10

NiO -1.86 -3.19

ZnO -1.96 -3.18

CdO -0.82 -1.23

508
In Fig. 1 the calculated energies of solution are plotted versus the

ionic radius of the dopant ion; a strong correlation between the two is found with an energy minimum observed at Ni2+. Such variation

with ion size is explained in terms of the relaxation of the surrounding

ions, and suggests that changes in local structural ordering may be

important in the observed depressions in TcPig. 1. Calculated solution energies versus divalent dopant radius

The simulation program used in this study also provides detailed

information on the relaxed positions of host-lattice ions. From our

structural analysis we find significant ion displacements around the

dopant ion. In Table II, the calculated M2+-Cu(l) and Cu(l)-0(4)

bond distances for the 3d dopant ions (M2+) are summarized. It has

been proposed that the Cu-0 bonds that link the chains to the planes

provide a sensitive probe of charge transfer effects. Indeed, from

neutron diffraction experiments correlations of certain average bond

lengths with Tc have been discussed. Therefore, to illustrate the

behaviour of our computed bond distances we plotted both the

M2+-Cu(l) and the Cu(l)-0(4) distances versus Tc, shown in Fig. 2. A trend towards decreasing Tc with increasing M2+-Cu(l) and

decreasing Cu(l)-0(4) bond distances is found. We recognize,

however, that the observed value of Tc is strongly dependent on the

precise dopant concentration. Nonetheless, diffraction measurements

[5] have similarly found Tc decreasing systematically with decreasing

Cu(l)-0(4) bond distance, for both Co-doped and oxygen-deficient

materials.

509
TABLE II. Calculated bond distances (in A) for divalent dopant

substitution on the Cu(2) site.

M2* M-Cu(l) Cu(l)-0(4)

Fe 4.2748 1.8430

Co 4.2806 1.8434

Ni 4.1936 1.8614

Cua 4.2560 1.8741

Zn 4.1835 1.8514

a Undoped lattice

(K)

100

80

60

40

20

1 .Cu

- -1 1

• NI

Fe . t Co

•

418 4-21 4-24 427

M^* - CuCI J bond disMnce CAj

OL1-84 1-85 1-86 1-87 1 88

CutU-OC4] bond dietancecAl

Fig. 2. Variation of Tc with calculated bond distances for dopants

on Cu(2). Observed values of Tc from [2]

For trivalent impurities substituting at Cu^* sites there are two

alternative charge-compensating defects: oxygen interstitials or copper

vacancies. The corresponding defect reactions are M2O3 + 2Cu^^ — 2M^^ + 0'1+ 2CuO (3)

or M2O3 + 3Cu^^ —t 2M^^ + V'! + 3CuO (4)

The former process involving oxygen incorporation is calculated to

have the lowest energy [12] and is, therefore, predicted to be the

majority compensating mechanism. This result is in accord with

thermogravimetric measurements that show an increase in oxygen

content with increasing A13+, Fe^* and Co^* dopant concentration [4]. Note that oxygen uptake accounts for the observed increase in

unit-cell volume even though AP*, Fe^* and Co3+ are all smaller than

Cu2*.

510
The calculated energies of solution for the trivalent ions Al, Sc, Fe,

Ga and Bi are presented in Table III. Examination of these energies

reveals that the trivalent dopants are predicted to favor substitution at

the Cu(2) site. Diffraction experiments have, in general, suggested

that Al and Fe substitute predominantly for Cu(l), although recent

neutron studies find significant Fe occupancy of Cu(2) sites even for

low dopant content [6,7]. Note that in this study we have considered

systems at high dilution i.e. the solution energies are for isolated

dopant ions. At higher concentrations the clustering of dopants may

play an increasingly crucial role, which could clearly effect the

calculated solution energies. These defect interactions obviously are

subjects for further study.

TABLE in. Calculated energies of solution (in eV per dopant ion).

M203 ^^Cu(l) ^^Cu(2)

A1203 0.10 -1.66

SC203 0.75 -1.44

Fe203 -0.31 -1.84

Ga203 -0.25

Bi203 4.23 4.22

a Calculation not converged

CONCLUSIONS

We have shown that the application of atomistic simulation

techniques can provide a useful way of investigating dopant phenomena

in the YBa2Cu307 ceramic oxide. While there may remain some

uncertainties in the calculated solution energies for this complex

system, our study demonstrates the ability of the method to

(fistinguish between alternative modes of dopant substitution and

compensation mechanisms. From our analysis of local ion

displacements immediately surrounding the dopant ion, we find a trend

towards decreasing Tc with increasing M2+-Cu(l) and decreasing

Cu(l)-0(4) bond distances, which is consistent with results of

neutron-diffraction measurements.

ACKNOWLEDGEMENTS

MSI would like to thank Roger Baetzold at the Kodak Research

Laboratories NY, for many useful discussions.

511
REFERENCES

[I] P.H. Hor, R.L. Meng, Y.Q. Wang, L. Gao, Z.J. Huang, J. Bechtold, K. Forster and C.W. Chu, Phys. Rev. Lett. 58, 1891

(1987) .

[2] G. Xiao, F.H. Streitz, A. Garvin, Y.W. Du and C.L. Chien,

Phys. Rev. B 35, 8782 (1987).

[3] Y. Maeno, T. Tomita, M. Kyogoku, S. Awaji, Y. Aoki, K.

Hosino, A. Minami and T. Fujita, Nature 328, 512 (1987).

[4] J.M. Tarascon, P. Barboux, P.F. Miceli, L.H. Greene, G.W.

Hull, M. Eibschutz and S.A. Sunshine, Phys. Rev. B 37, 7458

(1988) .

[5] P.F. Miceli, J.M. Tarascon, L.H. Greene, P. Barboux, F.J.

Rotella and J.D. Jorgensen, Phys. Rev. B 37, 5932 (1988).

[6] S. Katano, T. Matsumoto, A. Matsushita, T. Hatano and S. Funahashi, Phys. Rev. B 41, 2009 (1990).

[7] S. Katsuyama, Y. Ueda and K. Kosuge, Physica C. 165, 404

(1990).

[8] C.R.A. Catlow, R. James, W.C. Mackrodt and R.F. Stewart,

Phys. Rev. B 25, 1106 (1982).

[9] C.R.A. Catlow, Annu. Rev. Mater. Sci. 16, 517 (1986).

[10] C.R.A. Catlow, in Solid-State Chemistry: Techniques, edited

by A.K. Cheetham and P. Day (Clarendon, Oxford, 1987),

Chap. 7.

[II] M.S. Islam, M. Leslie, S.M. Tomlinson and C.R.A. Catlow, J. Phys. C 21, L109 (1988); C.R.A. Catlow, S.M. Tomlinson,

M.S. Islam and M. Leslie, ibid. 21, L1085 (1988).

[12] R.C. Baetzold, Phys. Rev. B 38, 11304 (1988); M.S. Islam and

R.C. Baetzold, ibid. 40, 10926 (1989)

[13] M.S. Islam, Supercond. Sci. Technol. (to be published).

[14] W.C. Mackrodt, Supercond. Sci. Technol. 1, 343 (1989).

512
National Institute of Standards and Technology Special Publication 804, Chemistry of Electronic Ceramic

Materials, Proceedings of the International Conference held in Jackson, WY, August 17-22, 1990, issued January 1991.

MOLECULAR DYNAMICS SIMULATIONS OF ION MOTION

IN DIVALENT AND MIXED MONOVALENT-DIVALENT

BETA' ' - ALUMINA

C. A. Lane and G. C. Farrington

University of Pennsylvania, Philadelphia, PA 19104

INTRODUCTION

The ability to design materials with specified

properties is an intriguing goal, one which may become

obtainable with the growing accessibility of fast

computers . To reach this goal, however, computer models

must be developed which can accurately reproduce

experimentally observed properties of known materials

before they can be trusted as design tools. Beta''-

alumina is a particularly good material for testing such

models. Well-known for its ability to undergo ion

exchange with a variety of mono, di, and trivalent

cations (1), beta''- alumina exhibits different

structural, optical, and transport properties depending

on its mobile ion content. Therefore, a reasonable

computer model for beta''- alumina should be able to

reproduce the properties of any of its isomorphs

.

Prior studies have been carried out on pure Na(I)-

beta''- alumina using molecular dynamics (MD) (2,3).

This work explores how well MD represents other beta' '

alumina isomorphs, in particular mixed Na ( I

) -Ba ( I I

) -, mixed Na ( I) -Sr ( II) -, and pure Sr(II)- beta''- alumina.

Simulation results will be discussed in light of a.c.

impedance measurements and structural data.

MD SIMULATION

The MD code used in this work is loosely based on a constant volume, number, and energy code written by

Walker (4) . Since the MD technique has been described

513
in great detail elsewhere (5) , only those aspects

particular to this study are presented here.

Interionic Potentials . All ions were treated as

rigid spheres to minimize computation time. The

coulombic contribution to the interionic potentials was

calculated by an Ewald summation. Short-range

potentials were in the Born-Mayer-Huggins form. The

values of the parameters, A, rho, and C, were taken from

several sources . The parameters for the pairs of ions

present in beta- alumina were taken from a fit to bulk

properties of Na(I)- beta- alumina (6), the parameters

for the Mg-0, Sr-0, and Ba-0 pairs wete taken from a fit

to bulk properties of their native oxides (7), and the

Sr- and Ba- related cation-cation potentials were

calculated using the electron gas approximation (8)

.

Equilibration . For the simulations of the mixed

monovalent-divalent systems, the MD box included 9 full

unit cells or approximately 800 atoms. The Sr(II)

simulations included 36 unit cells. Initially, the

mobile cations were placed randomly in the twodimensional conduction layers while all other ions were

placed in their crystallographic sites. The velocities

and lattice parameters were scaled periodically until a stable target temperature was reached and the internal

stresses were negligible. The volume of the box was

then held fixed for the rest of the simulation. The

equilibration period typically took several thousand

timesteps, where each timestep was 5x10"^^ sec long for

the mixed systems and 3.5x10"-'-^ sec long for the Sr(II)

system

.

Diffusion . The tracer diffusion coefficient was

obtained from the two-dimensional random walk equation.

The mean square distance travelled was calculated by

averaging over the number of mobile cations and over

time origins (9) . Values of the Haven ratio as a function of composition were taken from (10)

.

SIMULATION RESULTS AND DISCUSSION

Simulations of the mixed Na (I) -Ba (II) - system were

quite successful at reproducing expected microscopic

ordering processes and conductivity trends. Thomas et

al . have already reported preliminary MD simulations of

one composition of Na (I) -Ba (II) - beta'' - alumina, where

they examined the influence of a single Ba(II) on its

514
local environment (11) . Simulations presented here were

done at 927°C for Na-j^ 67-2x^^x'^'?0 61^-^10 33*^17' where x = 0.0-0.78.

Our simulations showed that the cations in the

conduction layers were quite mobile at low Ba(II)/Na(I)

ratios. However, near 80% replacement of Na(I) by

Ba(II), the cations and vacancies assumed an ordered

arrangement in which either a Na{I) or a Ba(II) occupied

one of the two possible tetrahedral, low potential

energy sites, known as the BR sites. Simulations of

this ordered structure indicated that its ionic

conductivity was much lower than that of compositions

containing more Na(I) . In fact, experimental

E

SIo

o3D

C

Ou

10

1

-

.1 -

.01 1

.001 =

.0001

o* •Ho*

O X

X >K X

T

20 40 60 80 100

% divalent exchange

Figure 1 Simulated and experimental conductivity as a function of composition in mixed Na(I)-Ba(II)- and mixed

Na(I)-Sr(II)- beta'' - alumina. Closed circles and X'

s

represent experimental data for the Na(I)-Ba(II) and the

Na(I)-Sr(II) systems respectively at 400°C. Open

circles represent Na(I)-Ba(lI) conductivity values

calculated from MD simulations at 927°C.

515
measurements of the ionic conductivity of Na (I) -Ba (II)

beta''- alumina as a function of composition showed that

the conductivity passes through a pronounced minimum at

this composition (12) . The simulated and experimental

conductivity data are shown in Figure 1.

Difficulties with the simulations arose near the

end compositions, pure Na(I)- and pure Ba ( 1 1

) -beta' '

alumina. The simulations at the high target temperature

we used became unstable at low Ba(II) content, below

about 20% exchange of Ba(II) for Na(I) . In addition,

above 80% exchange, the Ba(II) ions were essentially

immobile during the relatively short time span, 30 psec,

of the simulation. The resulting meager statistics made

it impossible to calculate a meaningful value of the

diffusion coefficient. Therefore, the simulations did

not predict the experimentally observed increase in

conductivity above 80% exchange.

The results of a.c. impedance measurements of mixed

Na ( I

) -Sr ( II ) -beta' ' - alumina as a function of

composition are also shown in Figure 1. In this case,

the pronounced minimum exhibited in the Na(I)-Ba(II)

system is not observed. It is conjectured that this is

because Sr(II) has been experimentally observed to

prefer an alternative octahedral site, the mid-oxygen

site, (14) and, therefore, would not be in competition

with the Na(I) cations for their preferred BR sites.

Simulations of 80% exchanged Na ( I) -Sr ( II) -beta' '

alumina, however, have incorrectly placed the Sr(II)

cations in the same locations as the Ba(II) cations and,

therefore, have incorrectly reproduced the conductivity

that has been observed experimentally.

Simulations of pure Sr (II) -beta' ' - alumina showed a small percentage of the Sr(II) cations either at the

mid-oxygen sites or off-centered from the BR sites . The

cations in these sites served as boundaries dividing

domains of ordered regions over which the Sr(II)

occupied one or the other of the two possible BR sites.

These boundaries were very short-lived as the Sr(II)

cations rapidly moved to the tetrahedral sites, and did

not convincingly reproduce the expected high occupation

of the octahedral sites.

CONCLUSION

The goal of this work has been to predict with some

accuracy the properties of several beta''- alumina

516
isomorphs, with a clear eye on the possibility of

ultimately being able to predict the properties of

unknown isomorphs. Our work illustrates the potential

opportunities as well as the current pitfalls which

accompany MD simulations.

First, it is gratifying that MD is quite successful

as a structural probe in the cases of Na ( I) - and Na(I)-

Ba ( II ) -beta' ' - alumina, but rather disconcerting that it

fails with the Sr(II) systems. It is likely that this

is due to both a problem with transporting cation-oxygen

potentials from their simple native oxides to the more

complex beta''- alumina structure, as well as a problem

with the simple assumption that all of the atoms are

fully charged ionic spheres . To remedy the latter

problem, we are in the process of simulating the mixed

Na(I)-Sr(II) system allowing for polar izability of

Sr(II) and neighboring oxygens by incorporating the

shell model

.

Second, the time-scale of MD simulations is

extremely short when compared to the time-scale of reallife ionic transport. As shown in this work, MD

performed quite well predicting the general decrease in

ionic conductivity for low Ba(II) concentrations in the

mixed Na(I)-Ba(II) system, but failed to predict the

subsequent increase in conductivity at the high Ba(II)

content end where fewer mobile cations are present and

jumps are infrequent

.

ACKNOWLED GEMENT S

This work was supported by the Office of Naval

Research and by the NSF-MRL program under Grant No. DMR8819885. Simulations were carried out primarily on the

Pittsburgh Supercomput ing Center Y-MP . C. Lane would

like to thank H. Hummel for providing the Sr- and Barelated cation-cation potentials.

REFERENCES

1. Y.-F.Y. Yao & J.T. Kummer, J. Inorg- Nucl . Chem.

.

29 (1967) 2453/ G.C. Farrington & J.L. Briant, Science,

204 (1979) 1371; G.C. Farrington & B. Dunn, Solid State

517
lonicS f 2 (1982) 267; G.C. Farrington, B. Dunn, & J.O.

Thomas, Appl . Phys . A , 2Z (1983) 159.

2. M.L. Wolf, J.R. Walker, & C.R.A. Catlow, Solid

State Ionics . 13 (1984) 33.

3. M.A. Zendejas & J.O. Thomas, Solid State Ionics ,. 28-30 (1988) 46; J.O. Thomas & M.A. Zendejas, submitted

for publication.

4. J.R. Walker, in: Computer Simulation of Solids , 166 , Lecture Notes in Physics , C.R.A. Catlow & W.C.

Mackrodt, (eds.). Springer, Berlin (1982).

5. M.J.L. Sangster & M. Dixon, Advances in Physics , 25 , 3 (1976) 247.

6. J.R. Walker & C.R.A. Catlow, J. Phys. C15 (1982)

6151 . 7. G.V. Lewis, Physica 13 IB (1985) 114.

8. R.G. Gordon & Y.S. Kim, J.Chem. Phys. , 56 , 6 (1972)

3122 . 9. M.J. Gillan, Physica 131B (1985) 157.

10. G.E. Murch & R.J. Thorn, Phil. Mag. , 36 , 3 (1977)

529.

11. J.O. Thomas, M. Alden, & G.C. Farrington, Solid

State Ionics . 9&10 (1983) 301.

12. C. Lane, G.C. Farrington, J.O. Thomas, & M.A.

Zendejas, Solid State Ionics , in press.

518
AUTHOR INDEX

Amren, R., 21

Ananthamohan, C, 507

Anderson, H. U., 179

Arbuckle, B. W.,207

Battle, P. D., 53

Batllo, F., 77

Beauger, A., 77

Bell, S., 21

Bhalla, A., 93

Billinge, S. J. L., 337

Bimie, D. P., 269

Blum, F. D., 179

Boehme, R. F., 427

Bonnell, D. A., 447

Brese, N., 275

Bringley, J. F., 343, 427

Bronnimann, C. E., 173

Brown, I. D., 471

Bursill, L. A., 67

Cahen, D., 351

Caignol, E., 367

Chan, C. C, 407

Cormack, A. N., 499

Cox, D. E., 427

Dabrowski, B., 357

Davies, P. K., 163, 337, 365

Davis, M. K., 173

Dec, S. F., 173

Dickens, P. G., 195

Dupon, R. W., 245

Egami, T., 93, 337

Eglin, J., 413

Farrington, G. C, 513

Feist, T. P., 163

Fiory, A. T., 237

Fitzgerald, J. J., 173

Fujitsu, S., 435

Garzon, F. H., 373

Goodenough, J. B., 399

Greenblatt, M., 207

Hall, R. B., 151

Han, S. S., 173

Hessen, B., 237

Hill, M. D., 225

Hinks, D. G., 357

Homonnay, Z., 407

Hong, J. L., 289

Huebner, W., 85

Ichinose, A., 393

Ikuma, Y., 283

Islam, M. S., 507

Jacobson, A. J., 151

Jang, G. W., 407

Jannot, B., 77

Jorgensen, J. D., 301, 357, 399

Jules, J. C, 77

JuLin, P., 67

Kauzlarich, S. M., 217

Kelly, J. F., 123

Kim, J. S., 301

King, T., 365

Klavins, P., 217

Kolar, D., 3

Koumoto, K., 435

Lane, C. A,, 513

LaPlaca, S. J., 427

Laurent, M., 77

Lee, J. Y., 301

Lewandowski, J. T., 151

Lightfoot, P., 301, 357, 399

Maciel, G. E., 173

Maglione, M., 77

Manthiram, A. , 399

Marx, D. T., 357

McElfresh, M. W., 427

Millar, D. M., 113

Mims, C. A., 151

Minor, D. B., 123

Mitchell, A. W., 357

Myers, G. R., 151

Nagai, A., 99

Nagy, S. I., 407

Nara, A., 393

Nath, A., 407

519
Nasrallah, M. M., 179

Navrotsky, A., 379

Negas, T., 21

Newbury, D. E., 123

Newnham, R. E., 39

Niepce, J. C, 77

Norton, M. L., 413

Nowotny, J,, 441

O'Keeffe, M., 485

Payne, D. A., 113

Pei, S., 301, 357, 399

Piffard, Y.,137

Poeppelmeier, K. R., 419

Powell, A. v., 195

Raistrick, I. D., 373

Randall, C, 85

Rawn, C. J., 225

Reiff, W. M., 301

Richards, D. R., 357

Ritter, J. J., 123

Rohrer, G. S., 447

Rosenfeld, H. D., 93

Roth, R. S., 225

Sabatani, E., 351

Sasaki, Y., 99

Schwartz, M., 351

Scolnik, Y., 351

Scott, B. A., 343, 427

Shannon, R. D., 457

Shaw, T. M., 427

Shrout, T. R., 105

Siegrist, T., 237

Sleight, A. W., 32*5 Sloma, M., 441

Smith, M. S., 179

Smyth, D. M., 289

Steinfmk, H., 301

Stuttard, G. P., 195

Sunshine, S. A., 237

Sunstrom, J. E., 217

Swinnea, J. S., 301

Switzer, J. A., 185

Takano, M., 255

Tanaka, S., 393

Tang, H. Y., 413

Tang, X. X., 399

Tanous, A. C, 245

Thompson, M. S., 245

Trail, S. S., 343, 427

Vaughey, J. T., 419

Wada, T., 393

Wang, S. P., 85

Waszczak, J. T., 237

Wei, Y., 407

Yaegashi, Y., 393

Yamauchi, H., 393

Yanagida, H,, 435

Yeager, G., 21

Yoshimoto, T., 99

Zandbergen, H. W., 309

Zhang, Z., 207

Zheng, Y., 357

520
ATTENDEES

Batllo, F. Negas, T.

Battle, P. D. Newnham, R. E.

Billinge, S. J. L. Norton, M. L,

Bimie, D. P. Nowotny, J. Brese, N. O'Keeffe, M.

Bringley, J. F. Owen, A. E.

Brown, I. D. Payne, D. A.

Bursill, L. A. Piffard, Y.

Cahen, D. Poeppelmeier, K.

Cava, R. J. Raistrick, L D.

Christoffersen, R. Rawn, C. J. Clark, T. Ritter, J. J. Cormack, A. N. Rohrer, G. S.

Dabrowski, B. Roserifeld, H. D.

Davies, P. K. Roth, R. S. Dickens, P. G. Roth, S. Dimos, D. Sasaki, Y.

Farley, T. Scott, B. A.

Feist, T. P. Shannon, R. D.

Fitzgerald, J. J. Shrout, T. R.

Fujitsu, S. Sleight, A. W.

Garzon, F. H. Smyth, D. M.

Greenblatt, M. Steinfmk, H.

Huebner, W. Subramanian, M, A.

Ichinose, A. Sunshine, S. A.

Ikuma, Y. Sunstrom, J. E.

Islam, M. S. Switzer, J. A.

Jacobson, A. J. Takano, M.

Kao, H. C. I. Tanzawa, T,

Kauzlarich, S. M. Thomas, T.

Klavins, P. Thompson, M. S.

Kolar, D. Trontelj, M.

Lane, C. A. Wilson, G.

Lessing, P. Wuensch, B. J. Lightfoot, P. Zandbergen, H. W.

Lukasiewicz, S. Macker, A.

Marezio, M.

McCarroll, W.

McHale, A.

Nasrallah, M. M.

Nath, A.

Navrotsky, A.

521

i^ U.S. GOVERNMENT PRINTING OFFICE : 1991 0 - 288-456 : QL3

NiST-1 14A U.S. DEPARTMENT OF COMMERCE

(REV. a^*?) NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY

BIBLIOGRAPHIC DATASHEET

1. PUBUCATION OR REPORT NUMBER ! NIST/SP-804

2. PERFORMINQ ORQANIZATION REPORT NUMBER

3. PUBUCATION DATE January 1991

4. TITLE AND SUBTITLE CHEMISTRY OF ELECTRONIC CERAMIC MATERIALS

5. AUTHOR(S)

Peter K. Davies and Robert S. Roth

6. PERFORMINQ ORQANIZATION (IF JOINT OR OTHER THAN NIST, SEE INSTRUCTIONS)

U.S. DEPARTMENT OF COMMERCE

NATIONAL INSTITUTE OF STANDARDS AND TECHNOLOGY

QATTHERSBURQ, MD 20899

7. CONTRACT/GRANT NUMBER

8. TYPE OF REPORT AND PERIOD COVERED Final

9. SPONSORING ORQANIZATION NAME AND COMPLETE ADDRESS (STREET, CITY, STATE, ZIP)

NASA, ONR, DuPont, Idaho National Engineering Laboratory, and University of Pennsylvania

10. SUPPLEMENTARY NOTES

DOCUMENT DESCRIBES A COMPUTER PROGRAM; SF-18S, FIPS SOFTWARE SUMMARY, IS ATTACHED.

IV ABSTRACT (A 200-WORD OR LESS FACTUAL SUMMARY OF MOST SIGNIFICANT INFORMATION. IF DOCUMENT INCLUDES A SIGNIFICANT BIBUOQRAPHY OR UTERATURE SURVEY, MENTION IT HERE.) The International Conference on The Chemistry of Electronic Ceramic Materials was held at the Sojourner Inn, in the Teton Village, Jackson Hole, Wyoming from August 17-22, 1990.

In an attempt to maximize the development of this rapidly moving, multidisciplinary field,

this conference brought together major national and international researchers to bridge the

gap between those primarily interested in the pure Chemistry of Inorganic Solids and those

interested in the Physical and Electronic Properties of Ceramics. With the many major

discoveries that have occurred over the last decade, one of the goals of this meeting was

to evaluate our current understanding of the Chemistry of Electronic Ceramic Materials, and

to assess the state of a field that has become one of the most important areas of Advanced

Materials Research.

12. KEY WORDS (6 TO 12 ENTRIES; ALPHABETICAL ORDER; CAPITAUZE ONLY PROPER NAMES; AND SEPARATE KEY WORDS BY SEMICOLONS)

crystal chemistry; dielectric ceramics; low temperature synthesis and characterization;

solid state synthesis and characterization; surface chemistry; superconductors; theory

and modelling

13. AVAILABIUTY UNUMITED

FOR OFFICIAL DISTRIBUTION. DO NOT RELEASE TO NATIONAL TECHNICAL INFORMATION SERVICE (NTIS). ORDER FROM SUPERINTENDENT OF DOCUMENTS, U.S. GOVERNMENT PRINTING OFFICE, WASHINGTON, DC 20402. X ORDER FROM NATIONAL TECHNICAL INFORMATION SERVICE (NTIS), SPRINGFIELD. VA 22161.

14. NUMBER OF PRINTED PAGES 513

15. PRICE

ELECTRONIC FORM

i X Technical Publications

Periodical

Journal of Research of the National Institute of Standards and Technology—Reports NIST research

and development in those disciplines of the physical and engineering sciences in which the Institute

is active. These include physics, chemistry, engineering, mathematics, and computer sciences. Papers cover a broad range of subjects, with major emphasis on measurement methodology and

the basic technology underlying standardization. Also included from time to time are survey articles on topics closely related to the Institute's technical and scientific programs. Issued six times a year.

Nonperiodicals

Monographs—Major contributions to the technical literature on various subjects related to the

Institute's scientific and technical activities. Handbooks—Recommended codes of engineering and industrial practice (including safety codes) developed in cooperation with interested industries, professional organizations, and regulatory bodies.

Special Publications—Include proceedings of conferences sponsored by NIST, NIST annual reports,

and other special publications appropriate to this grouping such as wall charts, pocket cards, and

bibliographies.

Applied Mathematics Series—Mathematical tables, manuals, and studies of special interest to physicists, engineers, chemists, biologists, mathematicians, computer programmers, and others engaged in scientific and technical work.

National Standard Reference Data Series—Provides quantitative data on the physical and chemical

properties of materials, compiled from the world's literature and critically evaluated. Developed un- der a worldwide program coordinated by NIST under the authority of the National Standard Data

Act (Public Law 90-396). NOTE: The Journal of Physical and Chemical Reference Data (JPCRD)

is published quarterly for NIST by the American Chemical Society (ACS) and the American Insti- tute of Physics (AIP). Subscriptions, reprints, and supplements are available from ACS, 1155 Six- teenth St., NW., Washington, DC 20056.

Building Science Series—Disseminates technical information developed at the Institute on building

materials, components, systems, and whole structures. The series presents research results, test methods, and performance criteria related to the structural and environmental functions and the

durability and safety characteristics of building elements and systems.

Technical Notes—Studies or reports which are complete in themselves but restrictive in their treat- ment of a subject. Analogous to monographs but not so comprehensive in scope or definitive in treatment of the subject area. Often serve as a vehicle for final reports of work performed at NIST

under the sponsorship of other government agencies.

Voluntary Product Standards—Developed under procedures published by the Department of Commerce in Part 10, Title 15, of the Code of Federal Regulations. The standards establish nationally

recognized requirements for products, and provide all concerned interests with a basis for common

understanding of the characteristics of the products. NIST administers this program as a supplement

to the activities of the private sector standardizing organizations.

Consumer Information Series—Practical information, based' on NIST research and experience, covering areas of interest to the consumer. Easily understandable language and illustrations provide useful background knowledge for shopping in today's technological marketplace.

Order the aboye NISTpublications from: Superintendent of Documents, Government Printing Office,

Washington, DC 20402.

Order the following NISTpublications—FIPS and NISTIRs—from the National Technical Information

Service, Springfield, VA 22161.

Federal Information Processing Standards Publications (FIPS PUB)—Publications in this series col- lectively constitute the Federal Information Processing Standards Register. The Register serves as the official source of information in the Federal Government regarding standards issued by NIST

pursuant to the Federal Property and Administrative Services Act of 1949 as amended. Public Law

89-306 (79 Stat. 1127), and as implemented by Executive Order 11717 (38 FR 12315, dated May 11, 1973) and Part 6 of Title 15 CFR (Code of Federal Regulations).

NIST Interagency Reports (NISTIR)—A special series of interim or final reports on work performed

by NIST for outside sponsors (both government and non-government). In general, initial distribu- tion is handled by the sponsor; public distribution is by the National Technical Information Service,

Springfield, VA 22161, in paper copy or microfiche form.
U.S. Department of Commerce

National Institute of Standards and Technology

(formerly National Bureau of Standards)

Gaithersburg, MD 20899

Official Business

Penalty for Private Use $300