BL1A

Mo L3-XANES Study of Silylation Effect for Methane
Dehydroaromatization over Mo/H-MFI Catalysts
H. Aritani, K. Kuramochi and S. Yoshinaga
Faculty of Engineering, Saitama Institute of Technology, Fukaya 369-0293, Japan

A Mo-modified H-MFI zeolite (Mo/H-MFI)
catalyst is one of a typical catalyst for direct GTL(Gas to liquid), because it shows high activity fordehydroaromatization of methane in absence ofoxygen. In this reaction, reduction of Mo species isbrought about in contact with methane in initial step,and reduced Mo ions react methane to form carbide
and/or oxycarbide species in next step. It has been
accepted that the carbide and/or oxycarbide species
are the active center for dehydroaromatization of
methane. However, deactivation cannot be avoided bycarbon deposition. Hydrogen co-feed with methane is
one of an effective method for suppression of coking.
However, excess hydrogen may affect a reduction of
Mo species during the reaction, and decrease of active
Mo-oxycarbide species may give low reactivity. On
the other hand, strong acid sites in external H-MFI
surface possibly relates to the coking directly. Thus,
selective covering of the coking site may be effective
for durable activity [1]. Silylation of one of a uniquemethod for selective coverage of external acid sites
onto H-MFI. This study focused on the effect ofsilylation on methane dehydroaromatization over
Mo/H-MFI catalysts. Mo LIII-edge XANES studies
were introduced to characterize the active Mo species
on bare and silylated H-MFI.
Catalysts were prepared by impregnation of H-MFI(Si/Al2=67) synthesized hydrothermally with
MoO2(acac)2íCHCl3 solution, and followed bydrying overnight and calcination at 773 K.
MoO3-loading amount is 5.0 wt% in whole catalysts.
For silylation of Mo/H-MFI catalysts, heptane
solution of teiethoxyvinylsilane (Mo/Si=10) was
employed for impregnation, and followed by drying
and calcination at 873 K. Mo LIII-edge XANES
spectra were operated in BL1A of UVSOR-IMS in a
top-up mode. InSb double crystal monochromator
was used for the XANES measurements in a
total-electron yield mode. The value of photon energy
was calibrated by using Mo metal-foil at Mo LIII-edge,
and normalized XANES spectra and their second
derivatives were obtained for characterization.
Catalytic activity was evaluated in a fixed bed flow
reactor. Each catalyst (0.250 g) was placed in aquartz-tube reactor, and pretreated in HeíCO(0-2%)
flow (30 mL min-1) at 973 K for 1 h. Then
CH4(20%)í H2(1%)íHe reactant gas was fed at 973
K (30 mL min-1: SV = 7.2 L g-1 h-1). Products wereanalyzed by online GC.
Figure 1 shows the Mo LIII-XANES spectra ofsilylated/unsilylated Mo/H-MFI catalysts after
dehydroaromatization at 973 – 1073 K for 3 h. As

109

Fig. 1. Mo LIII-edge XANES spectra of (A) bareand (B) Silylated- MoO3(5.0wt%)/H-MFI catalysts
after methane dehydroaromatization with (1)CH4(20%)-He or (2) CH4(20%)-H2(1.0%)-He atvarious reaction temperature.
shown in Fig. 1 (A-1 and B-1), According to a higher
reaction temperature, the edge energy of XANES
threshould becomes lower. It indicates the reduction
and/or carbonization of MoOx species to MoOxCyones. Excess reduction of Mo species brings about thedeactivation of the catalysts because of excess carbon
deposition [1, 2]. In case of silylated Mo/H-MFI [Fig.
1 (B-1) and (B-2)], the reduction degree of Mo
species is almost as same as bare Mo/H-MFI [Fig. 1(A-1) and (A-2)]. On these catalysts, high and durable
activity for methane dehydroaromatization issignificantly shown by silylation. The effect ofsilylation is possibly due to the depression of coking
from methane, and the reduction of Mo species isalmost independent. For Mo/H-MFI catalysts,
formation of active MoOxCy species (partiallycarbonized MoOx) is one of a key role for highactivity. And thus, deactivation of coking sites in
H-MFI extrapores is very effective to obtain durable
activity. In this study, it is suggested that silylation byusing teiethoxyvinylsilane act as the depression ofactivity on the coking sites because coking amountbecomes low. Detailed study about the silylation
effect on Mo species is now in progress.
[1] H. Aritani, S. Shinohara, S. Koyama, K. Otsuki, T.
Kubo, A. Nakahira, Chem. Lett. 35 (2006) 416.
[2] H. Aritani, H. Shibasaki, H. Orihara, A. Nakahira,J. Environm. Sci. 21-6 (2009) 736.

BL1A

XANES Analysis of Na-Incorporated E-Tricalcium Phosphate

K. Kawabata1 and T. Yamamoto2
1
Department of Holistic Human Science, Kwansei Gakuin University, 662-8501, Japan
2
Faculty of Science and Engineering, WasedaUniversity, 169-8555, Japan

110

Intensity (arb. units)

NaCl

E-TCP:Na

1070

1080

1090

Energy (eV)Fig. 1. Observed Na-K XANES spectra of
Na-incorporated E-TCP (Ca2.7Na0.3(PO4)2) and NaCl.

Intensity (arb. units)

Bioceramics based on the calcium phosphate are
now widely used for repairing the bone defects. As it
is well known that the main component of the
inorganic part of human bone consists ofhydroxyapatite (HAp). HAp itself is, of course, usedas bioceramics, but other calcium phosphates are also
employed, in which E-tricalcium phosphate (E-TCP)
is one of the most active one. It was reported thatincorporation of trace elements in such bioceramics
accelerates the creation of HAp in the human body.
To understand the influence of trace elements, it isessential to understand the local environment of traceelements in an atomic scale. We have systematicallyinvestigated the local environment of divalent ions,
i.e., Mg2+ [1], Mn2+ [2] and Zn2+ [3], by the XANES
measurements with the aid of the first-principles
calculations. In the present study, local environmentof monovalent Na ions in E-TCP is investigated by
the Na-K XANES measurements.
Sample specimens were fabricated by theconventional solid-state reaction method. High purity
powders of CaHPO4, CaCO3 and NaHCO3 weremixed and ground in an agate mortar, which were
calcined at 1273 K for 6 hours in air. Na-K XANES
spectra are measured at BL1A in UVSOR by the total
electron yield method. Sample powders were
mounted on the first dinode of the electron multiplierusing the carbon adhesive tape.
Observed
Na-K
XANES
spectrumofNa-incorporated E-TCP is shown in Fig. 1 together
with that of NaCl for comparison. It is difficult todetermine the local environment of Na ions in E-TCP
only from the comparison with the spectra of standard
materials. Then the theoretical calculations aremandatory for the quantitative analysis of the local
environment of incorporated Na ions. Here thefirst-principles calculations were carried out to obtainthe theoretical Na-K XANES spectra of Na in E-TCP
by using the WIEN2k package [4], in which thecore-hole effect was directly included by removing
the Na 1s electron. Several models by changing thesubstituted sites of Na incorporation were constructed.
Calculated spectrum using the model, in which Na issubstituted at Ca(4) site in E-TCP, is compared withthe observed one in Fig. 2. Observed Na-K XANES
profile of Na-incorporated E-TCP is quantitatively
well reproduced by this model. This result confirmedthat the incorporated Na ions are substituted at Ca(4)site in E-TCP.

Expt.

Calc.
1070

1080

1090

Energy (eV)Fig. 2. Comparison between observed and calculated
Na-K XANES spectra of Na-incorporated E-TCP.
[1] K. Kawabata, H. Sato and T. Yamamoto, UVSOR
Activity Report 35 (2007) 108.
[2] K. Kawabata, H. Sato and T. Yamamoto, J. Ceram.
Soc. Jpn. 116 (2008) 108.
[3] K. Kawabata, T. Yamamoto and A.Kitada, Physica
B 406 (2011) 890.
[4] http://www.wien2k.at

BL1A

Evaluation of Hydrotalcite-Based Nanosheet by Soft Chemical Process

1

A. Nakahira1, S. Yamamoto1, M. Sato2, Y. Takamatsu1, S. Misu1 and H. Aritani3
Faculty of Engineering, Osaka Prefecture University, Gakuencho, Sakai 599-8531, Japan
2
Osaka Center, IMR, Tohoku University, Gakuencho, Sakai 599-8531, Japan
3
Saitama Institute of Technology, Fukaya 369-0293, Japan

Introduction
Layered double hydroxide (LDH) is one of uniqueclay minerals with a layered structure, its composition
formula of [M2+1-X M3+X (OH)2࡮A n-X/n࡮nH2O].
These LDH and its related materials are expected to
be one of novel materials because of their potential
applications in wide areas, such as catalysis,
adsorption, and nanocomposite and drug delivery.
Recently, the nanosheets from the exfoliation oflayered host compounds have received increasing
attention because the resulting unilamellar materials
or nanosheets have various attractive aspects.
Furthermore, the nanosheets from the LDH are
expected to be especially useful among novel
nanosheet related materials.
In the present study, we focused on the structural
evaluation of the nanosheets from LDH basicallycomposed of Mg2+ and Al3+, which is calledhydrotalcite, “MgAl-Tal”. Furthermore, it is known
that these nanosheets are important for the novel
nanostructured materials synthesized by thelayer-by-layer methods. However, it is not clarified
about the detailed structure evaluation of thenanosheets from MgAl-Ta. In the present study, thestructure evaluation of the nanosheets from LDH
composed of Mg2+ and Al3+ (MgAl-Tal: hydrotalcite)were performed.

Results and Discussion
MgAl-Tal samples prepared by the co-precipitation
method was identified to be a layered doublehydroxide consisted of Mg2+ and Al3+ by powder
X-ray diffraction analysis. Then, the delamination ofobtained products were successfully prepared by thesoft chemical process. The products obtained fromthe delamination of MgAl-Tal indicated the tyndall
phenomenon, suggesting that nanosheets wereobtained. It was suggested that products obtained
from the delamination of MgAl-Tal had broad peaks
derived from XRD results.
Figure 1 shows the results of XANES of Al-K edgeof the nanosheet products obtained from thedelamination of MgAl-Tal. For the comparison,
MgAl-Tal samples prepared by the co-precipitation
method were similarly evaluated. The spectrum of
he nanosheet products obtained from thedelaminationof
MgAl-Talcorresponded
approximately to that of MgAl-Tal samples preparedby the co-precipitation. It was elucidated that Al-K
edge XANES spectra of the nanosheet products
obtained from the delamination of MgAl-Tal weresimilar to that of MgAl-Tal as a starting material.

Experiments
From 0.2 mol/dm3 MgCl2 aqueous solution and 0.1
mol/dm3 AlCl3 aqueous solution, mixing solution was
prepared by the co-precipitation method. MgAl-Tal
with M2+/M3+ =2~4 were synthesized by adding
mixing solution into 0.05 mol/dm3 NaHCO3 at room
temperature. 1 mol/dm3 NaOH was simultaneously
added into the aqueous solution in order to keep pH
10. Products for MgAl-Tal were aged at room
temperature for 2 hours. They were separated andsufficiently washed by deionized water and finally
air-dried at 50 ͠ for 24 hours. The delamination ofobtained MgAl-Tal products were prepared by soft
chemical process. The products were obtained fromthe delamination of MgAl-Tal.
The local structures around Al for the nanosheets
were characterized by measuring X-ray adsorption
near edge structure (XANES) at BL1A in UVSOR
with KTP.
Fig. 1. Results of XANES of Al-K edge of the
MgAl-Tal nanosheet.

111

BL1A

Evaluation of Local Structure of SBF-Type Mesoporous Silica Bulks

A. Nakahira1,2, S. Yamamoto1, M. Sato2, Y. Takamatsu1, S. Misu1, Y. Kawabe1,
Y. Nishio1 and H. Aritani3
1
Faculty of Engineering, Osaka Prefecture University, Gakuencho, Sakai 599-8531, Japan
2
Osaka Center, IMR, Tohoku University, Gakuencho, Sakai 599-8531, Japan
3
Saitama Institute of Technology, Fukaya 369-0293, Japanproducts of SBF-type bulk were characterized bymeasuring X-ray adsorption near edge structure(XANES) at BL1A in UVSOR.

Introduction
Much attention has been paid to the mesoporous
silica families, MCM, FSM, and SBF and so on.
Since SBA-15 powders synthesized by the sol-gel
method with triblock copolymer surfactants possess ahuge specific surface area and range of relatively
uniform pore size and furthermore contain a uniqueinternal structure with uniform arrays of hexagonalmesopore, it is expected to be potentially useful as
catalysts supports and adsorbent and so on. In
particular, SBA-15 powders have the thicker silica
walls around mesopores than that of MCM-type
mesoporous silica (MCM and FSM) made withconventional cationic surfactants, leading to thestability for SBA-15 at high temperature and/or under
severe conditions. However, the development for the
preparation of bulky mesoporous silica SBA-15 isstrongly desired, because powdered SBA-15 materials
are restricted in their practical application in many
fields. We successfully prepared the dense SBA-15
bulk by the modified hydrothermal hot-pressing
process. In this study, the local structures for
SBF-type mesoporous silica bulks were investigated
by XAFS measurements.
In the present study, we focused on the structureevaluation of the bulks of SBF-type mesoporous silica
by the modified hydrothermal hot-pressing process. In
special, the structure evaluations of the SBF-type bulk
were performed.

Results and DiscussionDense SBA-15 bulks were successfully
synthesized by the modified hydrothermal
hot-pressing possess. These dense SBA-15 bulks withmesoporous structure were similar to that of
powdered SBA-15, suggesting that the dense SBA-15
bulks with mesoporous structure as well as powdered
SBA-15 were successfully synthesized by themodified hydrothermal hot-pressing possess.
Figure 1 shows the results of X-ray adsorption nearedge structure (XANES), Si-K edge, for SBA-15 bulk
prepared by the modified hydrothermal hot-pressing
at 150 ºC. For the comparison, SiO2, FSM and
SBA-15 powder were similarly evaluated. X-ray
adsorption near edge structure (XANES) of the dense
SBA-15 bulks indicated the similar pattern of
powdered SBA-15. It was obvious that dense SBA-15
bulks obtained by the hydrothermal hot-pressing
possess hexagonal regularity similar to powdered
SBA-15 from these Si-K edge XANES results.

Experiments
SBA-15 was synthesized by the following method.
Silica and block copolymer mesophases weresynthesized at room temperature as follows. Pluronic
P123 (Mav=5800, EO20PO70EO20) as a block
copolymer was dissolved in water and HCl solution
with stirring at room temperature. Then, TEOS as asilica was added into that solution with stirring atroom temperature. The mixture was aged at 80 ºC for
12 h without stirring. The solid product was filtered,washed, and air-dried at room temperature.
Calcination was carried out, slowly increasing
temperature from room temperature to 500 ºC andheating at 500 ºC for 6 hours. Subsequently, theconsolidation of this SBA-15 powder was carried out
by the modified thermal hot pressing method with theuniaxial pressure of 30 MPa at 110 ºC to 150 ºC for 2
h. After the modified thermal hot pressing treatments,
obtained bulks were heat-treated at 600 ºC in the air
atmosphere. The local structures around Si for the

Fig. 1. The results of X-ray adsorption near edge
structure (XANES), Si-K edge, for SBA-15 bulks.

112

BL1A

1

Investigation of Local Structure of Mg K-Edge in
Chitosan and Mg Co-Addition Calcium Phosphate Materials

T. Onoki1, Y. Nishio1, T. Moriguchi1, Y. Kawabe1, M. Sato2 and A. Nakahira1,2
Graduate School of Engineering, Osaka Prefecture University, Sakai 599-8531, Japan
2
Osaka Center for Industrial Materials Research, Institute for Materials Research,
Tohoku University, Sakai 599-8531, Japan

Human bone has an organic/inorganic composite
structure. Bone replacing biomedical materials, forexample hydroxyapatite (HA) ceramics, has beenrequired more biocompatibility and optimalmechanical properties. Recently, bioceramics were
fabricated by some hydrothermal techniques at low
temperature [1]. Magnesium (Mg) is known to be an
important trace element in bone and teeth. Indeed,
despite its low concentration (generally between 0.5
and 1.5 wt%), it plays a key role in bone metabolism,in particular during the early stages of osteogenesis
where it stimulates osteoblast proliferation, and its
depletion causes bone fragility and bone loss.
Relationships have been suggested between themagnesium content in enamel and the development ofdental caries. Given the biological relevance ofmagnesium, many research teams have worked on the
preparation of apatite and calcium phosphate implantmaterials containing relatively low levels of Mg,
which has been shown to improve their bioactivity [2].
According to the literature, the replacement ofcalcium by magnesium in HA is difficult.
Chitosan was one of origanic materials derived
from crab shell. And has function of accumulatingmineral species. So we conducted for making a Mg
adding HA/chitosan composite ceramic for morebiocompatibility. Calcium hydrogen phosphate
di-hydrate (DCPD) is one of starting materials for HA
ceramics through a hydrothermal hot-pressing
technique at around 100 ºC [1]. As a first step, Mg
adding DCPD/chitosan composite powder wereprepared. Co-precipitation method was employed inorder to prepare Mg adding DCPD/chitosan
composite materials. The composite materials werebased on the DCPD preparation [3], which 1.0M
calcium nitrate solution(99.0%; Ca(NO3)2࡮4H2O,
KANTO
CHEMICAL,
Japan)
and1.0M
di-ammonium hydrogen phosphate solution (98.5%;(NH4)2HPO4 ; KANTO CHEMICAL, Japan) weremixed. Magnesium chloride (99.0%; MgCl; Wako,
Japan) Chitosan powder (LLWP Grade, KIMICA,
Japan) was dissolved in water with a little acetic acid.
The chitosan and MgCl solution was added toCa(NO3)2࡮4H2O solution with various concentration
before mixing to the (NH4)2HPO4 solution. And thenthe fabricated powder were converted to HA by ahydrothermal treatment (150 ºC, 2h).

113

Mg K-edge XANES spectra for the fabricatedpowders were obtained in a total electron yield modeat room temperature using a KTP double-crystal
monochrometer at BL1A station of UVSOR. The
spectra were collected in the photon energy range from1295-1330 eV at intervals of 0.05eV with a dwell timeof 1s. Mg-K edge XANES of the DCPD/chitosan
composite and the HA converted powder were shownin Figure 1. Firstly, we could not detect Mg in the HA
sample without chitosan. On the other hand, a little
amount of Mg can be observed in both of DCPD and
HA with chitosan. These result indicate that chitosan
play an important role for Mg adding to calcium
phosphate materials. It is shown that chitosan is agood carrier for doping non-HA mineral elements to
HA.

Fig. 1. XANES spectra of Mg K-edge of the
DCPD/chitosan composite and the HA converted
powder with Mg.
[1] S. Ishihara, T. Matsumoto, T. Onoki, T. Sohmura
and A. Nakahira, Mater. Sci. Eng. C 29 (2009) 1885.
[2] S. Gomes, G. Renaudin, E. Jallot and J. M.
Nedelec, ACS Appl. Mater. Interfaces 1 (2009) 505.
[3] T. Onoki, K. Hosoi and T. Hashida, Scr. Mater. 52(2005) 767.

BL1A

1

Investigation of Effect of Ions in Seawater for Ettringite Formation

M. Sato1, S. Yamamoto2, Y. Nishio2, Y. Takamatsu2 K. Kumadani2 and A. Nakahira1,2
Osaka Center for Industrial Materials Research, Tohoku University, Sakai 599-8531, Japan
2
Graduate School of Engineering, Osaka Prefecture University, Sakai 599-8531, Japan

Ettringite (3CaO࡮Al2O3࡮3CaSO4࡮32H2O) is oneof calcium sulpho aluminate (CSA) system, and it hasa needle-like microstructure belongs to hexagonal
system. It is a naturally occurring mineral found inGermany for the first time. This mineral ischaracterized by the very high content of watermolecules. Since ettringite is important material toprovide the rapid set and/or high early strength forconstructional materials, more effective material
design can be expected by investigating the generation
behavior in detail. However, there are few reports forthe generative mechanism of ettringite. In this study,
we prepared ettringite using different aqueous
solutions including Mg2+, Na+, CO32-, SO42- ions etc.
and investigated the generation behavior in different
ion enviroments.
In order to obtain the ettringite, tricalcium
aluminate (C3A, 3CaO㨯Al2O3) were firstly preparedby CaCO3 and Al2O3 powders as starting materials.
After mixed up each powder, it was calcinated at 1573
K for 36 hours in air. And then, obtained C3A andCaSO4㨯2H2O were mixed and stirred at 200 rpm from
5 minutes to 24 hours in the different aqueous
solutions. After stirring, samples were filtered anddried.
Na-K, Mg-K, Al-K and S-K edge XANES spectraof obtained ettringite powders were measured in a total
electron yield mode at room temperature using Beryl,
KTP and InSb double-crystal monochrometer at BL1A
station of UVSOR, respectively.
Figure 1 shows Mg-K edge XANES spectra of
ettringite prepared in 1 mol/l MgCl2 aqueous solution
at room temperature for several aging times. In thecase of MgCl2 solution with 0.01 mol/l concentration,
no peaks derived from Mg compounds were observed.
However, spectra derived from Mg compounds wereobserved in the case of 0.1 and 1 mol/l MgCl solution
as shown in Fig. 1. From XANES measurement, it issuggested that the Mg based layered double hydroxide(LDH) is generated within 5 minutes, and thegeneration of LDH affects the formation of ettringite.
Similar trend that the ettringite formation is affected at
the concentration of solution more than 0.1 mol/l was
obtained in other solution such as NaCl, KCl andNa2SO4.

Normalized Intensity ( a.u.)

24 h

1h
5 min
M g(OH)2

M gO

1300

1310

1320

Energy / eV
Fig. 1. Mg-K edge XANES spectra of ettringite
prepared in 1 mol/l MgCl2 aqueous solution at room
temperature for several aging times.

114

BL1A

1

Investigation of Local Structure of P-K Edge in Fe Doped HAp

M. Sato1, S. Yamamoto2, Y. Nishio2, Y. Takamatsu2 and A. Nakahira1,2
Osaka Center for Industrial Materials Research, Tohoku University, Sakai 599-8531, Japan
2
Graduate School of Engineering, Osaka Prefecture University, Sakai 599-8531, Japan

So far, many kinds of magnetic beads such asparamagnetic
iron
oxide
particles
andsurface-modified ones by polymer have been widelyused in a medical field. However, those iron oxidebased magnetic beads have been pointed out the risk ofthe elution of iron by long term use. The resolvabilityof iron strongly depends on its particle size. Since theparticle size of magnetic beads is decided by a service
space, control of resolvability is one of the problems
that we should overcome to use it in a living body.
Because the elution of iron affects proliferation andmetabolism of cell, there is a possibility to exhibit
some toxicity. Therefore, development of non toxic
magnetic bead has been strongly desired.
Hydroxyapatite (Ca10(PO4)6(OH)2, HAp) is a goodcandidate for drug delivery system materials due to its
excellent ion exchange ability and protein adsorption
ability. By using these unique abilities of HAp, nontoxic magnetic bead can be developed.
In this study, Fe doped HAp was prepared byconventional wet process and local structure around P
atom was investigated by XAFS measurement.
0.1 mol/l Ca(NO3)2, (NH4)2HPO4 and FeCl3
solutions were used as starting materials. These
solutions were mixed up each other becoming the ratioof Ca to P of 1.67 and the ratio of Fe from 0 to 20
mol% for Ca solution. Mixed solution was aged for 1 h
at room temperature and then it was filtered and driedat 323 K for overnight. P-K edge XANES spectra of
Fe doped HAp powders were measured in a total
electron yield mode at room temperature using InSb
double-crystal monochrometer at BL1A station of
UVSOR.
Figure 1 shows P-K edge XANES spectra of Fe
doped HAp prepared by conventional wet process atroom temperature. Obtained P-K edge XANES spectraof samples were similar to that of commercial HAp,
and a change of the spectral pattern was not clearlyobserved within any ratio of Fe. This result indicatesthat the local structure of around P atom of Fe doped
HAp samples is almost same with commercial HAp.

Normalized Intensity / a.u.

20 mol% Fe

10 mol% Fe

5 mol% Fe

1 mol% Fe
0 mol% Fe

Comm.. HAp

2140

2145

2150

2155

2160

Energy / eV
Fig. 1. P-K edge XANES spectra of Fe doped HAp
prepared by conventional wet process at room
temperature.

115

BL1B

Optical Spectroscopy of ZnO Thin Films

N. Kashiwagura, Y. Kubo, T. Hashida, Y. Hayashi and M. Yamaga
Department of Mathematical and Design Engineering, Gifu University, Gifu 501-1193, Japan
Today the most widely adopted TCO (Transparent
Conductive Oxides) films for use as electrodes ofsolar cell and liquid crystal display is ITO (Indium
Tin Oxides) films. However, ITO films include sosparse Indium that the substitutive materials are
expected. ZnO thin films are one of the hopefulmaterials as TCO films.
ZnO thin films were deposited on a fused silica glass(10 mm x 20 mm x 0.5 mm) by the DC sputtering
method. The sputtering time, the gas (Ar) pressure
and the substrate temperature were 240 minutes, 0.40
Torr. and 200 ͠, respectively. The induced electric
power was 60 W for sample 1, and 45 W for sample
2.
Optical absorption, luminescence and excitation
spectra were measured using the BL1B beam in thetemperature range of 10-300 K.
Figure 1 shows the optical absorption spectra for
ZnO thin films. The sharp edge at 150 nm for sample
2 is due to the absorption of glass substrate. The
several broad bands below 360 nm (3.3 eV) are due tothe band-to-band transition of ZnO. The sample 2 isefficiently transparent as TCO thin film above 380
nm (visible light region), while the sample 1 is not so.
Figure 2 shows the luminescence spectra withvarious excitation wavelengths in sample 2 at 300 K.
The band-to-band excitation below 360 nm produces
two broad luminescence bands. In these spectra, theadditional peaks around 540 and 590 nm are theinfluence from outer lights, which could not be
perfectly eliminated. As the excitation wavelength
increases (the energy decreases), the intensity ofluminescence band around 570 nm decreases, whilethat around 420 nm still remains.
Figure 3 shows the luminescence spectra in sample
2 at 10 K. The spectrum with the 100 nm excitationshows no dent around 480 nm, in other words, bothbands around 420 nm and 570 nm get broader andcloser. With increasing the excitation wavelength, theluminescence intensities decrease. As a consequence,
the band around 420 nm remains, whereas that around
570 nm almost disappears.
The band around 420 nm is estimated to be
intrinsic luminescence (for example, self-trapped
exciton). The broader band around 570 nm is
estimated to be the luminescence associated withdefects (for example, donor/acceptor recombination).
The relation between the luminescence andelectrical conductivity will be examined. In this study,
optical properties of pure ZnO thin films were
investigated. It seems to be more important toinvestigate Ga or Al doped ZnO thin films in the
point of electrical conductivity.

Fig. 1. Absorption spectra in ZnO thin films at 300 K.

Fig. 2. Luminescence spectra in sample 2 at 300 K.

Fig. 3. Luminescence spectra in sample 2 at 10 K.

116

BL1B

Optical Spectroscopy of Wide Gap Semiconductors

M. Yamaga, H. Uno, Y. Ogoshi, A. Kanetake, S. Takano and T. Nohisa
Department of Material and Design Engineering, Gifu University, Gifu 501-1193, Japan
InxGa1-xN and Cd1-xMnxTe crystals are wideband-gap semiconductors. Band-gap energies ofindium gallium nitride (InxGa1-xN) mixed crystals
decrease with increasing x. Absorption edges of
InxGa1-xN, corresponding to the band-gap energies,
vary from the ultraviolet to visible ranges. Then,
InxGa1-xN mixed crystals are a promising material for
multi-junction thin film solar cells. On the other hand,
Cd1-xMnxTe crystals have magneto-optical properties
induced by Mn2+ ions. They are useful as an opticalisolator material.
InxGa1-xN thin films were deposited on SiO2
substrates by rf-sputtering at room temperature.
Sputtering targets were In and Ga metals with various
weight ratios. The X-ray diffraction (XRD) patternsof the InxGa1-xN thin films consisting of several peaksin the angle range between 30 and 70 degrees indicate
microcrystalline thin films.
Cd1-xMnxTe single crystals were grown by thevertical Bridgeman method. The composition of thecrystal is the same as the melting composition.
Figure 1 shows absorption spectra observed in
In0.2Ga0.8N and Cd0.5Mn0.5Te at 300 K. The absorption
edge of the GaN thin film was observed at 360 nm atroom temperature, whereas that of In0.2Ga0.8N thinfilm is shifted toward a longer wavelength of about
450 nm as shown in Fig.1. The absorption spectrumof Cd0.5Mn0.5Te consists of several broad bands in therange from 100 to 600 nm.
It is reported that GaN crystals show sharp exciton
luminescence around 360 nm and substitution of In in
GaN crystals leads to red-shift and line-broadening of
the luminescence. UV excitation for the In0.2Ga0.8N
thin film produces broadband luminescence withdouble peaks at 460 and 560 nm as shown in Fig. 2.
The luminescence for the GaN thin film is similar tothat for In0.2Ga0.8N thin film. This result suggests thatthe 460 and 560 nm luminescence broadbands may be
associated with N deficits and impurity O ions.
Figure 3 shows temperature dependence of theluminescence spectra excited at 100 nm for
Cd0.5Mn0.5Te. The peak energy of the luminescence isjust below the band edge energy in Fig. 1. The
luminescence intensities markedly decrease above
120 K. The luminescence may be due to weakly
localized excitons.
Identification of these luminescence centers
InxGa1-xN and Cd1-xMnxTe crystals gives information
on structure and energy levels of the localized
luminescence centers.

6

Absorbance

Cd0.5Mn0.5Te

4

In0.2Ga0.8N

2

0
100

200 300 400 500
Wavelength (nm)600

Fig. 1. Optical absorption spectra at 300 K.

Intensity (arb. units)

In0.2Ga 0.8N

1.0

Oex=260 nm
Oex=300 nm

0.5

0

400

500
600
Wavelength (nm)700

Fig. 2. Luminescence spectra in InGa0.8N.
1.0

Intensity (arb. units)

Cd0.5Mn0.5Te
O =100 nm
ex

0.5

0
300

300K
120K
90K
60K
8K

400 500 600 700
Wavelength (nm)800Fig. 3. Luminescence spectra in Cd1-xMnxTe.

117

BL3U

Soft X-Ray Emission Spectroscopy in the C KĮ ofGraphene-Nanoflakes (GNFs) Functionalized with Nitrogen

J. W. Chiou1, M. T. Liou2, W. F. Pong2, H. Yamane3 and N. Kosugi3
1
Department of Applied Physics, National University of Kaohsiung, Kaohsiung, Taiwan
2
Department of Physics, Tamkang University, Tamsui, Taiwan
3
Institute for Molecular Science, Okazaki 444-8585, Japan
Figure 1 shows normal x-ray emission spectra(NXES) and Figure 2 shows resonant inelastic x-ray
scattering (RIXS) spectra excited at different photon
energies in the spectra of the C KĮ of HOPG and
GNFs:N deposited on Si(100) substrate, preparedusing microwave plasma enhanced chemical vapor
deposition process and subsequently functionalizedwith nitrogen plasma by the electron cyclotron
resonance process. Spectral features of these GNFs:N
films are essentially similar, consisting of a broad
main feature around 277.4 eV (A) and a high-energy
shoulder feature near 281.7 eV (B), to those observedby Muramatsu et al. [1]. However, the peak-to-heightratio of the shoulder-feature to the main-featuredepended on the nitrogen at.% (0 to 20 at.%). These
spectral features therefore suggest that local electronic
structures and chemical bonding states are essentiallysimilar among the GNFs:N films, but that their ratios,for example of sp2 and sp3-configured carbon atoms,
differ depending on the nitrogen at. %. Comparing thespectra of GNFs films and reference HOPG, weconclude that the main feature (A) is in same position
with HOPG and the high-energy shoulder feature (B)of the GNFs films are not in the same position asslightly shifts towards higher energy than in HOPG.
These features A and B are the ı and ʌ orbitals,
respectively. The feature height and their ratios aretabulated in Table 1. It is very difficult to draw someconclusion on the (ʌ/ı) ratio presently because on
nitrogen functionalization initially it decreases andthen increases indicating the change of sp2/sp3 ratio.
Figure 2 shows that the spectral features are variedwith the excitation energy, indicating sp2 and/or sp3
carbons are resonantly excited at certain excitation ofenergy. The redistribution of intensities in the spectra
indicates changes in the occupancies of different
molecular orbitals, possibly due to changes in electron
density of states upon nitrogen incorporation in
GNFs:N. The NXES spectral profile has been used asa means to identify certain chemical states ofincorporation of nitrogen in the GNFs structure.
However, it is also observed in Fig. 2 that the C KĮ
RIXS spectra strongly depend on excitation energy
and the dependence is believed to stem from theresonant inelastic scattering and related to the band
structure of the GNFs:N. This interpretation isimportant for all the emission experiments with highresolution synchrotron radiation and implies thepossible use of photon-in-photon-out technique forelectronic band structures. This soft NXES/RIXS

revealed a new picture of the core excitation and theexistence of strong vibronic coupling effects in GNFs
and GNFs:N.

Intensity (arb. units)

C KD NXES
Eex=310 eV

A
B

HOPG
GNFs:N (at.%)
0
7
10
20

265

270 275 280
Energy (eV)285

Fig. 1. NXES spectra of HOPG and GNFs:N.
HOPG
GNFs:N(at.%)
020

Intensity (arb. units)

C KD RIXS
310.0 eV291.9 eV288.4 eV285.4 eV284.8 eV265

270 275 280
Energy (eV)285

Fig. 2. RIXS spectra of HOPG and GNFs:N.
Table 1. Intensity of ı and ʌ and their ratios.

HOPG
GNFs
GNFs:N(7at.%)
GNFs:N(10at.%)
GNFs:N(20at.%)Peak B
(ʌ)
26.3
36.2
34.3
26.2
58.2

Peak A
(ı)
37.5
41.2
47.7
31.2
66.4

(ʌ/ı)
ratio
0.70
0.88
0.72
0.84
0.88

[1] Y. Muramatsu et al., Carbon 39 (2001) 1403.

118

BL3U

Probing Electronic Structure of Graphdiyne
by Resonant Soft-X-Ray Emission Spectroscopy

J. Zhong1, T. Xie1, X. H. Sun1, S. D. Wang1, Y. L. Li2, H. Yamane3, N. Kosugi3 and J. H. Guo4
1
Jiangsu Key Laboratory for Carbon-Based Functional Materials and Devices, Institute ofFunctional Nano and Soft Materials Laboratory (FUNSOM), Soochow University,
Suzhou 215123, China
2
Institute of Chemistry, Chinese Academy of Sciences, Beijing 100190, China
3
Institute for Molecular Science, Okazaki 444-8585, Japan
4
Advanced Light Source, Lawrence Berkeley National Laboratory, Berkeley, CA 94720, USA
The structural study of carbon is a main
fundamental interest and is expected to be widely
applied to nanosciences and technology. Due tocarbon’s three hybridization states (sp3, sp2 , and sp),having numerous combinations by which atoms ofthis element can be bonded to each other to produce
many carbon allotropes; such as graphite (sp2),
diamond (sp3) in the nature, and many novel carbon
allotropes such as fullerene (sp2), carbon nanotube(sp2), and graphene (sp2), are successfully synthesized.
The design and synthesis of new carbon allotropes
with definite structure and property is a significant
and ongoing challenge within new materials scienceand there are still a large number of new forms ofcarbon to be discovered. Graphdiyne proposed is a
novel structure in the carbon family and predicted tobe the most stable of the various diacetylenic
nonnatural carbon allotropes and is one of the most“synthetically approachable” [1]. Recently, thesynthesizing of a novel, stable graphdiyne form ofcarbon has been reported and in this work, we probe
the electronic structure of graphdiyne in comparison
with that of highly oriented pyrolitic graphite(HOPG), single-walled carbon nanotube (SWNT) andgraphene by using resonant soft-x-ray emission
spectroscopy (XES).
X-ray absorption spectroscopy (XAS) probes theunoccupied density of states (DOS) and XES probes
the occupied DOS. In addition to the inherent
elemental selectivity of X-ray spectra, energy
selective excitation of the emission spectra allows
separation of features that pertain to different atoms
or same element in different chemical environments.
In previous report [2], the quantum confinement of
SWNT with decreased tube diameters has beenreported which can be observed in resonant emission
spectra by significant changes in the ʌ and ı bands.
Figure 1 shows the C K-edge resonant emission
spectra of HOPG, SWNT, graphene and graphdiyne at
the excitation energy of 285.5 eV. The spectrum forgraphdiyne in the ı band shows great changes
compared to that for HOPG and SWNT, while similar
to that for graphene, indicating a similar hybridization
as graphene. Figure 2 shows the C K-edge normal
emission spectra at the excitation energy of 305 eV. A
difference in the ı band for graphdiyne can be found119

compared to all other samples, indicating the specificelectronic structure of graphdiyne.

Fig. 1. C K-edge resonant emission spectra of HOPG,
SWNT, graphene and graphdiyne at the excitationenergy of 285.5 eV.

Fig. 2. C K-edge normal emission spectra of HOPG,
SWNT, graphene and graphdiyne at the excitationenergy of 305 eV.
[1] G. X. Li, Y. L. Li, H. B. Liu, Y. B. Guo, Y. J. Li
and D. B. Zhu, Chem. Comm. 46 (2010) 3256.
[2] J. Zhong et al., Appl. Phys. Lett. 93 (2008)
023107.

BL4B

Synchrotron Radiation Effect on the Local Structure ofHighly-Hydrogenated Diamond-Like Carbon Film
K. Kanda, M. Uemura and T. Hasegawa
Laboratory of Advanced Science and Technology for Industry, University of Hyogo,
Ako 678-1205, Japanof sp2 bonded carbon atoms can be extracted bynormalizing the area of the resonance corresponding to
1soS* transitions at 285.4 eV with the area of a large
section of the spectrum. The absolute sp2/(sp2+sp3)
ratio was determined by the comparison with that fromthe NEXAFS spectrum of graphite.
The sp2/(sp2+sp3) ratio in the DLC film, estimatedfrom NEXAFS study, was plotted to irradiated SR
dose in Fig.1. The sp2/(sp2+sp3) ratio increased
steeply from |0.5 before irradiation of SR to |0.6
after the irradiation of SR of 300 mA·h dose. Beyond
300 mA·h dose, the sp2/(sp2+sp3) ratio keeps |0.6. As
a result, carbon double bonding can be regarded toformed due to the desorption of hydrogen fromhighly-hydrogenated DLC film in the dose region
under 300 mA·h.

Diamond-like carbon (DLC) film is expected to use
as lubrication material in space, where oil cannot be
used, because its some excellent properties such as
hard, low friction coefficient, low surface energy andso on. It was known that the DLC films are durable
against X-ray exposure, because it was reported thatDLC films are not etched by exposure to soft X-ray(SX) in the absence of oxygen gas [1]. However, manykinds of novel DLC films have been fabricated by theremarkable advancement in DLC industry in last
decade. Recently, irradiation of SX in a vacuum into
DLC films has been reported to lead the desorption ofhydrogen and an increase in the film density, hardness,and refractive index [2]. Information on themechanism of the deterioration of DLC film by X-ray
exposure is undoubtedly important for use of DLC
films in space safely.
In the previous study [3], we reported that the SX
exposure in the vacuum with the absence of oxygen to
the DLC films formed using the amplitude-modulated
RF CVD method leaded to the desorption of hydrogenand etching. In the present study, we measured
NEXAFS spectra of the DLC films at BL8B1 ofUVSOR to discuss of the variation of local structure of
DLC film by the irradiation of SX.
DLC films were deposited with 200-nm thickness
on Si wafers by using the amplitude-modulated RF
plasma-enhanced CVD method. Three kinds of DLC
film, each with different hydrogen content, were usedas samples. The SX irradiation of the DLC film wascarried out at beamline 6 (BL-6) of NewSUBARU
synchrotron facility in the University of Hyogo. The
electron energy of the NewSUBARU ring was 1.0
GeV and the ring current was 220 mA on a top-up
mode. The synchrotron radiation (SR) at the BL-6
sample stage had a continuous spectrum from IR tosoft X-rays, which are lower than 1 keV. Sample DLC
films were exposured to SR at room temperature.
Local structure was evaluated from the Near edge
X-ray absorption fine structure (NEXAFS) study
performed at the BL4B of UVSOR. The NEXAFS
spectra were measured in the energy range 275-320 eVwith 0.5 eV FWHM resolution in the total electron
yield mode.
A pre-edge resonance at 285.4 eV is not visible inthe spectrum of diamond because diamond consists of
only carbon atoms in sp3 (C-C) sites. Therefore, thepeak intensity of this resonance is considered as a goodindex of sp2 content. The procedure for determinationof sp2 content from the NEXAFS measurements wasestablished in the previous studies [4, 5]. The amountsp2(sp2+sp3)

0.8

0.6

0.4

0.2

0

2000
4000
6000
SR dose / mA࡮h

8000

Fig. 1. Dose dependence of sp2/(sp2+sp3) ratio in the
DLC film[1] H. Kyuragi and T. Urisu, Appl. Phys. Lett. 50(1987) 1254.
[2] H. Matsuura, S. Ohkubo, K. Oda and T. Ushiro,
SEI Technical Review 171 (2007) 39 (Japanese).
[3] K. Kanda, K. Yokota, M. Tagawa, M. Tode, Y.
Teraoka and S. Matsui, Jpn. J. Appl. Phys. to be
published.
[4] K. Kanda, T. Kitagawa, Y. Shimizugawa, Y.
Haruyama, S. Matsui, M. Terasawa, H. Tsubakino, I.
Yamada, T. Gejo and M. Kamada, Jpn. J. Appl. Phys.
41 (2002) 4295.
[5] K. Kanda, Y. Shimizugawa, Y. Haruyama, I.
Yamda, S. Matsui, T. Kitagawa, H. Tsubakino and T.
Gejo, Nucl. Instrum. Methods Phys. Res. B 206 (2003)
880.

120

BL4B, BL1A

Study on the Electronic Structures of the Layered
Li2MO3-LiMO2 Materials in Li De-Intercalation Process

H. Kobayashi1 Y. Takenaka2 and M. Shikano1
1
Research Institute for Ubiquitous Energy Devices, AIST, Ikeda, Osaka 563-8577, Japan2
Department of Chemistry and Materials Engineering, Faculty of Chemistry, Materials andBioengineering, Kansai University, Suita, Osaka 564-8680, Japan[1] A. Ito et al., J. Power Sources 183 (2008) 344.
[2] A. Ito et al., J. Power Sources 195 (2010) 567.
[3] Y. Arachi et al., Solid State Ionics 176 (2005) 895.
[4] H. Kobayashi et al., J. Power Sources 146 (2005)
640.

121

5.0
4.5

Voltage / V

4.0
3.5
3.0
2.5

: Sample A: Sample B

2.0
1.5
0

50

100

150

200

250

300

Capacity / mAh/g

Fig. 1. 1st charge curves for the samples A and B.

a b

Intensity ( Arb. Unit )The layered Li2MO3-LiMO2 (M =transition metal)
materials are one of the promising positive electrodematerials of lithium secondary battery because of theirlarge capacity when operated above 4.6 V [1].
Especially, Li[Ni0.17Li0.2Co0.07Mn0.56]O2 displays a 1stdischarge capacity of c.a. 280 mAh/g in the voltage
range of 2.5 to 4.8 V and keep a reversible capacity ofc.a. 250 mAh/g after 50 cycles [2]. Several papers
have reported on the mechanism why these materials
show large reversible capacity. However, the initialcharge and discharge process are still ambiguous. We
have reported on the characteristic structural change
during Li de-intercalation for LiNi1/2Mn1/2O2 [3,4]using synchrotron radiation. Detailed information on
the crystal and electronic structures above 4.6 V isvery important in order to improve the calendar life
and thermal stability of these materials and, therefore,
the structural changes of Li1.01Ni0.49Co0.21Mn0.29O2(Sample A) and Li1.20Ni0.17Co0.10Mn0.53O2 (Sample B)
were investigated in Li de-intercalation process.
The de-lithiated samples from samples A and B
were electrochemically prepared using coin-type cellswith Li/1M LiPF6 in EC:DMC(1:2)/samples.
Electronic structures were investigated by the total
electron yield mode at BL4B and BL1A in UVSOR.
The data were analyzed using the program Athena.
Figure 1 shows the 1st charge curves of thesamples A and B in the voltage range of 2.0 and 4.8 V.
The Li/Sample A cell showed the 1st charge capacityof 282 mAh/g with the characteristic plateau region
around 4.5 V. On the other hands, the Li/Sample B
cell showed the 1st charge capacity of 251 mAh/g
without plateau region.
Figure 2 shows the Ni L-edge XANES spectra forthe de-lithiated samples from the samples A and B(BL4B). The spectrum measured contained peaks aand b, corresponding to the valence state of 2+ and
3~4+, respectively. The sample A showed that theintensity of peak b increased to y=0.8 with Li
de-intercalation. On the other hand, the sample B
showed that the maximum of peak b was observed at
y=0.4. These results indicate that Li de-intercalation
mechanism is different between these samples.

y=0.8
y=0.4
Sample B

y=0
y=0.8
y=0.4

Sample A
840

850

y=0
860

870

880

Energy ( eV )Fig. 2. Ni L-edge XANES spectra for the de-lithiatedsamples from the samples A and B (BL4B).

BL5U

Effects of Periodicity Imperfection on Mini-Bands in
Semiconductor Superlattice

K. Hashimoto1, T. Ujihara1, Y. Maeda1, M. Kuwahara2, H. Miyazaki3 and Y. Takeda1
1
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan
2
EcoTopia Science Institute, Nagoya University, Nagoya 464-8603, Japan
3
UVSOR Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
Recently, a quantum dot solar cell is proposed toachieve a high conversion efficiency that can reachover 50%, theoretically. The concept of the quantum
dot solar cell is to utilize mini-bands in the periodicalignment of quantum dots as a middle band in the
carrier excitation in order that the band structure isadjusted to the solar light spectrum. The theoretical
calculation is based on an assumption of "theperfectly periodic potential." However, it is presently
impossible to produce quantum dots of equal size andperfect periodicity. In this study, the effects ofperiodicity imperfection on the mini-band structure,especially hole mini-bands, were investigated. In
place of aligned quantum dots, we examined asuperlattice structure which consists of multiple
quantum wells because the preparation method has
been established.
We fabricated three specimens: a GaAs bulk
sample, an AlGaAs/GaAs superlattice, and adisordered AlGaAs/GaAs superlattice. Figure 1
shows the detailed structures of each sample. The
second well layer in the disordered superlattice wasmade thicker. Amorphous arsenic layers were
deposited on the sample surfaces to suppress
oxidation. The arsenic layer was thermally removed
just before measurements. On the evaluation of themini-band structure, the synchrotron-radiation
photoemission measurements at BL5U of the
UVSOR facility were performed.
Figure 2 shows the energy spectra in a valence bandaround * point. These mappings were processed by
the second order differential with respect to thebinding energy. There are mainly two bands that aresymmetrical at the red arrows. The lower and higher
energy bands are thought to be a heavy hole band anda light hole band, respectively. Theoretically, heavy
hole band and light hole band in GaAs aredoubly-degenerated at the * point. But, in the present
result, these bands are split and its width of the GaAs
sample is 180 meV corresponding to the width ofband splitting at the wavenumber k// = 0.0669 Å.
Thus, these mappings show the band structures at thiswavenumber.
The bands of (a) GaAs and (b) periodic superlattice
are widely dispersive, while those of (c) disordered
superlattice are flatter. The width of dispersion
depends on the degree of confinement effect. This
flatter band is considered to be due to the quantum
level in the thicker well layer.
Figure 3 shows widths of band splitting between theheavy hole band and the light hole band as a functionof wavenumber. We found that the band splitting
width of superlattice is larger than that of GaAs bulk.
As far as we know, this result is the first direct
observation of the band splitting due to themini-bands of the ordered superlattice. The splittingenergy of the disordered superlattice is wider
comparing to that of GaAs. However, thewavenumber dependence is small because the band
splitting is due to the narrow quantum level in thequantum well. These results imply that a disorder of
only one layer in superlattice disturbs the ideal
mini-band structure.

Fig. 1. Schematic images of the sample structure. (a)GaAs bulk sample, (b) periodic superlattice and (c)disordered superlattice.

Fig. 2. Photoemmision spectra of (a) GaAs bulk
sample, (b) periodic superlattice and (c) disordered
superlattice.

Fig. 3. Width of band splitting between heavy holeband and light hole band.

122

BL5U

Angle-Resolved Photoemission Study on Epitaxial Graphene
Grown on SiC(0001)

A. Ruammaitree, H. Hu, H. Nakahara, K. Akimoto, S. Harada, K. Soda and Y. Saito
Department of Quantum Engineering, Graduate School of Engineering, Nagoya University,
Nagoya 464-8603, Japan
Graphene, which has many exotic properties such as
high mobility and linear dispersion (Dirac cone) at theK-point in the Brillouin zone, can be synthesized byvarious methods. Annealing SiC is one of the efficientapproaches for a large scale production of graphene[1], where it is important to control the layer numberof grown graphene. To determine its layer number,there are several techniques such as Auger electron
spectroscopy, x-ray photoelectron spectroscopy andellipsometry. The accuracy of their techniques
depends on their models. In this report, not only theelectronic structure but also the estimation of layer
number of graphene epitaxially grown on SiC isstudied by angle-resolved photoemission spectroscopy(ARPES). The result can give us the accurate numberof the epitaxial graphene layers.
The epitaxial graphene on an n-type Si-terminated
6H-SiC(0001) substrate of 12u3u0.25 mm3(CRYSTAL BASE co. ltd.) was prepared byannealing it in an atmosphere of argon, as follows.
After ultrasonic cleaning with acetone, the substrate
was immediately mounted on a sample holder and put
in a chamber with the base pressure of ~10-8 Pa. In
order to remove oxides on the substrate surface, wedeposited silicon around 2 monolayers on it. Si
molecules deposited are expected to crash andeliminate the oxide molecules from the surface. Thenthe sample was transferred, without exposure to air, toan argon chamber where it was annealed by resistive
heating in the atmosphere of argon. The annealing
temperature was measured by an optical pyrometer.
The epitaxial growth of graphene was confirmed byreflection high energy electron diffraction (RHEED).
The transverse and longitudinal orientations of thesubstrate are [-1, 0, 1, 0] and [-1, 2, -1, 0], respectively,
and the unit cell of the graphene rotates by 30 degrees
to the SiC cell.
The ARPES measurement was conducted at room
temperature with a hemispherical electron analyzer(MBS A-1) and a circularly polarized photon beam as
an excitation source. To obtain a clear Dirac coneimage at the K point, the excitation photon energy was
set to 80 eV and the angle between the photon beamand the surface normal was 23 degrees. Total energy
resolution including the thermal broadening wasestimated to be 0.12 eV by measuring the Fermi edgeof an Au film.
ARPES is one way for estimating the layer numberof epitaxial graphene [2]. Figure 1 shows theelectronic structure around the K point (zero wave
number in the figure) for the sample prepared by123

annealing at 1948 K under the argon pressure of 0.03
MPa. The green dot curves represent ʌ and ʌ* bands
which originate from double layered graphene, while
the red lines stand for a linear dispersion of a graphenemonolayer. The presence of a brighter curve in thevalence band and a bright area in the band gap
suggests that this sample may be composed of themonolayer and bilayer graphenes. This is consistent
with our x-ray diffraction observation of the averaged
layer number of ~1.6, which will be reported
elsewhere.

S bandS bandFig. 1. Electronic band dispersion around the K point(zero wave number) of the graphene of the SiC(0001),
which was synthesized by annealing at 1948 K under
the argon pressure of 0.03 MPa.
[1] J. Hass et al., J. Phys.: Condens. Matter 20 (2008)
323202.
[2] T. Ohta et al., Phys. Rev. Lett. 98 (2007) 206802.

BL5U

Valence-Band Structure of Mg2Si in High-Pressure Phase
K. Soda, T. Shimada, H. Kondo, S. Harada, K. Nishida and M. Hasegawa
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan124

a small but definite intensity is observed around EF,
which might be ascribed to defect states within theband gap and might even imply the conduction
bottom near EF as expected by the band structure
calculation and the n-type thermoelectric power. On
the other hand, the intensity at EF is clearly increased
and the valence band top moves to EF for the highpressure phase. This is consistent with both the highelectric conductivity and the p-type thermoelectric
power, although it is difficult to study the unoccupied
conduction band by photoelectron spectroscopy.
Further study is intended for specimens
synthesized in Ar atmosphere.

2

0.2
1

Mg2Si

0.1

DOS
High pressure phase
Normal pressure phase
Au

0
2

1.5

1
0.5
0
BINDING ENERGY / eV0

DENSITY OF STATES / eV-1cell-1

0.3

INTENSITY / arb. units

Valence-band electronic structure has been studied
for Mg2Si in a metastable hexagonal high pressure
phase by photoelectron spectroscopy. Semiconducting
Mg2Si and its related compounds have received muchattention because of its high thermoelectric properties
with light constituent elements [1-3]. Mg2Si in anormal pressure phase show n-type thermoelectric
power, probably due to Mg-derived defects, and its
related p-type compounds are desired to develop.
Recently we have found high p-type thermoelectric
power and large electric conductivity for themetastable high pressure phase. According to a band
structure calculation [3], the fundamental band gap is
expected to collapse to a pseudogap in so far reported
high pressure phases, orthorhombic anti-cotunnite (>7.5 GPa at room temperature) and hexagonal
Ni2In-type (> 21.3 GPa) structures. However, theelectronic structure has not been clarified for thepresent high pressure phase yet.
Specimens of metastable hexagonal high pressure
phase Mg2Si were prepared under 7 GPa at 1173 K bymulti-anvil cell. Mg2Si in a normal pressure phase
was also obtained by pressing Mg2Si powder without
annealing. Valence-band spectra were recorded at 10
K with a total energy resolution of 0.12 eV at theexcitation photon energy hQ of 70 eV. The origin of
the binding energy EB was set to the Fermi energy EF
of a reference Au film electrically connected to thespecimen. Although the specimens were in situ
fractured or filed to clean their surfaces for thephotoelectron measurement, contamination by oxide
species, unfortunately, was not completely removed.
However, it seems not to affect the valence-bandstructure near EF so much, since the bands of theoxide species are located at EB > ~ 4 eV [6, 7].
Figure 1 shows typical valence-band spectra near
EF of Mg2Si in the normal and high pressure phases
together with an Au spectrum and density of states(DOS) for the normal pressure phase, which was
calculated by an all-electron full-potential linearized
augmented plane wave (FLAPW) method with the
WIEN2k code [4] at the experimental lattice constanta = 0.6354 nm of the anti-fluorite structure.
Calculated indirect band gap of ~ 0.2 eV is narrower
than experimentally reported (0.63 eV [3] and 0.77
eV [5]), which may be partly attributed to thewell-known tendency in the density functional
calculation. The spectral intensity is normalized at EB~ 2 eV and the origin of EB for the calculated DOS isset to the conduction bottom because of the n-type
thermoelectric properties of the normal pressure
phase. For the normal pressure phase, the spectral
intensity decreases towards EF, as predicted. However,-0.5

Fig. 1. Valence-band photoelectron spectra of Mg2Si
in normal and high pressure phases. Fermi edge of areference Au film and calculated density of states(DOS) of normal pressure phase Mg2Si are also
shown for comparison.
[1] Y. Noda et al., J. Jpn. Inst. Mtals 53 (1989) 487.
[2] M. Akasawa et al., J. Appl. Phys. 104 (2008)
013703.
[3] F. Yu et al., Solid State Commun. 150 (2010) 620.
[4] K. Schwarz et al., Comp. Phys. Commun. 147(2002) 71.
[5] U. Winler, Helv. Phys. Acta. 28 (1955) 633.
[6] M. Cardona et al., Phys. Stat. Sol. 58 (1973) 483.
[7] M. Brause et al., Surf. Sci. 398 (1998) 184.

BL5B

In-situ Observations of Photoinduced Effects on Amorphous Chalcogenide
Semiconductors by Synchrotron Orbital Radiation

K. Hayashi
Department of Electrical and Electronic Engineering, Gifu University, Gifu 501-1193, Japan
Introduction
It is well known that amorphous chalcogenide
semiconductor materials show a variety ofphotoinduced
effects [1-3].
Photodarkening(photoinduced optical absorption edge shift),photodegradation (photoinduced defect creation) andphotoinduced volume change are representative
photoinduced phenomena. These phenomena arecaused by irradiation of light (bandgap (BG) light)
having photon energy comparable to the optical
bandgap. These phenomena can in general be either
irreversible, i.e. the photoinduced changes arepermanent after irradiation, or reversible, in whichcase the changes can be removed by annealing to theglass-transition temperature. The photoinducedchanges by BG light of the x-ray diffraction and thevolume have directly shown that these phenomena are
due to a change of the local structure of the amorphous
network. Several previous studies of photoinduced
effects have reported that there is no direct correlation
among photodegradation, photodarkening andphotoinduced volume change. The detail is notclarified though the dynamics of those photoinduced
effects certainly seems to be different.
Our recent studies have focused on thephotoinduced energy structure changes in the vacuum
ultraviolet (VUV) region by irradiation of BG light.
To obtain a wide knowledge of the photoinduced
phenomena, we have investigated photoinduced
effects in the VUV region by the transmission spectraof amorphous thin films prepared onto polymer
membrane substrates. In previous work, we reported
the VUV transmission spectra of substrate-free
evaporated amorphous thin films that removed thesubstrate [4]. Few studies have examined whether thephotoinduced change of the chemical bonds
contributes to photodarkening and photodegradation.
To understand the dynamics and the correlation ofphotoinduced structure changes in the VUV region
and the photodegradation in amorphous chalcogenide
semiconductors, real time in-situ measurements are
required. However, it is difficult to measure thetransmission change while irradiating the BG light.
Therefore, those dynamics and correlations are
examined by measuring the total photoelectron yield(TPEY) that can measure the optical property in the
VUV region. In this report, we will compare thetransmission spectrum and the TPEY spectrum in theamorphous thin films.

Experimental
Samples used for the measurement of thetransmission spectrum were substrate-free evaporated
amorphous chalcogenide (a-As2S3 and a-As2Se3)
semiconductor thin films. Samples used for themeasurement of the TPEY spectrum were preparedonto quartz substrates which fabricated gold
electrodes by conventional evaporation technique.
The thickness of the amorphous films was fromaround 160 nm to 300 nm. The measurements were
performed at room temperature at the BL5B beamline. A pinhole of 1.5 mm in a diameter was inserted
between the monochromator and sample to remove
stray light. The intensity of the VUV light wasmonitored by measuring the TPEY of a gold mesh.
Results and Discussion
Figure 1 shows the VUV transmission spectrum
and the TPEY spectrum of a-As2S3 films at room
temperature. It is necessary to note it because theintensity correction by incident light strength is not
performed as for both spectra. As shown in the figure,the absorption peak corresponds very well in both
spectra. The escape depth of the photoelectrons andthe creation of the secondary electrons might
influence the difference of both spectral line shapes.
It will be necessary to discuss this point in the future.

a-As2S3
Intensity [a.u.]Transmission

TPEY

38

40

42
44
46
48
Photon Energy [eV]50

52Fig. 1. VUV transmission and total photoelectron
yield (TPEY) spectra of a-As2S3 film.
[1] Ke. Tanaka, Rev. Solid State Sci. 4 (1990) 641.
[2] D. L. Staebler and C.R. Wronski, Appl. Phys. Lett.
31 (1977) 292.
[3] S. R. Elliot, J. Non-Cryst. Sol. 81 (1986) 71.
[4] K. Hayashi, UVSOR Activity Report 34 (2006)
79.

125

BL5B

Characterization of 10-Bilayer TiO2/ZnO Mirrors at 2.74 nm
H. Kumagai1, Y.Tanaka1, M.Murata1, M.Sanjo1, T. Shinagawa2 and M. Chigane2
1
Graduate School of Engineering, Osaka City University, Osaka 558-8585, Japan
2
Inorganic Materials Lab., Osaka Municipal Technical Research Institute,
Osaka 536-8553, Japan
x-ray mirrors once further studies are conducted interms of the growth rate, interface, surface roughness,and reactant vapors. Further reduction of the surface
and interface roughness will lead to the fabrication of
normal-incidence reflectors at the water-window
wavelengths.
㩷The authors have proposed the use of a novel metaloxide multilayer for soft-x-ray reflectors atwater-window wavelengths [1], because an oxide
multilayer can prevent the formation of an alloy at theinterface without any buffer layer, and the absorptionof oxygen in oxides is negligible at water-window
wavelengths; moreover, the metal oxide multilayer
can be fabricated by atomic layer deposition or
atomic layer epitaxy (ALE). These techniques can beused to control surfaces on an atomic scale bysequentially dosing the surface with appropriate
chemical precursors and then promoting surface
chemical reactions that are inherently self-limiting.
We have found that the self-limiting adsorption
mechanism is effective in the fabrication of thin oxide
films such as aluminum oxide and titanium oxide.
Moreover,wereportedtheexperimental
demonstration of a high reflectivity of over 30% at awavelength of 2.74 nm and an incident angle of 71.8°
from the normal incidence using novel metal oxide
multilayers consisting of titanium oxide andaluminum oxide fabricated by controlled growth byatomic layer deposition with sequential surface
chemical reactions [1]. For high-power x-ray
processing, crystalline multilayer mirrors might bemore useful than those with amorphous layers.
Therefore, the authors recently demonstrated thefabrication of novel oxide superlattice structures ofcrystalline TiO2/ZnO on sapphire substrates for
multilayer mirrors with high reflectivity at 2.74 nm[2, 3].
In the fabrication study, first of all, (001)-oriented
wurtzite ZnO and (100)-oriented rutile TiO2 films
were grown on a (001)-oriented sapphire substrate byatomic layer epitaxy at a substrate temperature of
450 °C [4, 5]. These films were grown layer by layer
owing to the self-limiting nature of surface chemical
reaction. The ZnO films were grown using DEZ andH2O vapors at a constant growth rate of nearly
0.26 nm/cycle. The TiO2 films were deposited using
TCT and H2O at a constant growth rate of nearly
0.076 nm/cycle.
We fabricated a 10-bilayer TiO2/ZnO multilayer
mirror by atomic layer epitaxy at a substratetemperature of 450 °C [2, 3]. Figure 1 shows that themultilayer mirror achieved a high reflectivity of
29.4% at a wavelength of approximately 2.74 nm anda grazing angle of 2T=10°. The FWHM of theexperimentally obtained reflectivity at 2.74 nm was
0.041 nm, which corresponds to 'O/O=1.5%. The
TiO2/ZnO multilayer is expected to be useful for soft㩷Fig. 1. Experimental reflectances of the ALE-grown
10-bilayer TiO2/ZnO structure on a sapphire (0001)
substrate at a wavelength of approximately 2.74 nm
and a grazing angle of 2T=10°.
[1] H. Kumagai, K. Toyoda, K. Kobayashi, M. Obara
and Y. Iimura, Appl. Phys. Lett. 70 (1997) 2338.
[2] H. Kumagai, Y. Tanaka, M. Murata, Y. Masuda
and T. Shinagawa, Journal of Physics: Condensed
Matter 22 (2010) 474008 pp.1-7.
[3] H. Kumagai and Y. Yanagihara, Review of Laser
Engineering 38 (2010) 976.
[4] H. Kumagai, Y. Masuda and T. Shinagawa,
Journal of Crystal Growth 314 (2011) 146.
[5] H. Kumagai, Y. Masuda and T. Shinagawa,Materials Science and Engineering (2011) in press.

126

BL6U

Band Alignment at CNT/SiC Interface Formed by Surface Decomposition
T. Maruyama1, S. Sakakibara1, H. Ito1, H. Yamane2, E. Shigemasa2 and N. Kosugi2
1
Department of Materials Science and Engineering, Meijo University,
Nagoya 468-8502, Japan
2
Institute for Molecular Science, Okazaki 444-8585, JapanE SB

result indicates that the band bending occurred at theinterface, which might be induced by deep level traps.
This work was supported in part by the Joint
Studies Program (2010) of the Institute for Molecular
Science (IMS).

(a)hQ=60 eV
Au
CNT

0.4

0.2

0.0 -0.2 -0.4

Binding Energy

Intensity [a. u.]

Fig. 1. Schematics of CNTs on SiC formed by surface
decomposition.

Intensity [a. u.]

Carbon nanotube (CNT) growth by surface
decomposition of SiC is a synthesis method for
self-organized CNT films only by heating SiC single
crystal at high temperature in a vacuum [1]. By thismethod, high density aligned zigzag-type CNTs canbe produced perpendicular to the substrate surfaces(Fig. 1). Moreover, it has been demonstrated thatgrown CNTs are atomically bonded to SiC substrate[2]. This is a distinguished characteristic of thismaterial, because there are some interlayers at theinterface in other CNT/semiconductor heterojunctions.
In this study, to clarify the electronic structure at the
CNT/SiC interface, we performed the photoemission
spectroscopy (PES) experiments.
CNT films formed by surface decomposition of
6H-SiC(000-1) face were used as samples. The
current-voltage (I-V) measurement for CNT (~400nm)/SiC heterojunction showed distinct rectifying
behavior with the forward direction occurring at
positively biased CNTs. For PES measurement, thesamples with the CNT thickness of about several nm
were also prepared to investigate the CNT/SiC
interface. The PES measurement was carried out in ahigh-resolution angle-resolved PES system at BL6U.
All PES spectra were measured at normal emission
with the incident photon energy of 60 and 350 eV.
The overall energy resolution was below 100 meV atroom temperature (~300 K).
Figure 2 (a) shows photoemission spectrum near
valence band maximum (VBM) of CNTs (~400 nm)
with that of Au. The Fermi edge is observed in the
spectrum of CNTs, indicating that they were metallic
tubes.
The value of Schottky barrier height at theinterface between metallic CNTs and SiC wasevaluated by the energy difference between the C 1s
core levels at the CNT/SiC interface (shown in Fig. 2(b)) and the VBM to the core level separations in theCNT(~400nm)/SiC and SiC substrate samples, as
described by

(b)

C 1s

CNT
SiC

288 286 284 282 280

Binding Energy [eV]Fig. 2. Photoemission spectra of CNT (~400 nm)/SiC
heterojunction: (a) valence band near the Fermi level(hQ = 60 eV), (b) C 1s core level (hQ = 350 eV).

CNT ( CNT / SiC )(CNT / SiC )SiC
E gSiC  E CSiC E CSiC E CCNT
1s  EVBM  E C1s
1s
1s

3.0  280.87  1.76  284.22 1.41 [eV ]Fig. 3. Band alignment of CNT/SiC heterojunction.
where EiS denotes the energy of feature i in sample s,[1] M. Kusunoki et al., Appl. Phys. Lett. 77 (2000)
531.
[2] M. Kusunoki et al., Chem. Phys. Lett. 366 (2002)
458.
[3] O. Madelung, Semiconductors: Data Handbook(Springer-Verlag) (2004) 61.

and the band gap of 6H-SiC, E gSiC , is 3.0 eV [3]. The
obtained band alignment at the CNT/SiC interface isshown in Fig. 3. The Schottky barrier height at the
CNT/SiC interface, 1.4 eV, is fairly high, comparedwith conventional metal/SiC Schottky barriers. Our

127

BL6B

Far-Infrared Reflective Study of Alkali Niobate Ceramics

H. Matsudo, S. Koide, T. Nishi, I. Kagomiya and K. Kakimoto
Graduate School of Engineering, Nagoya Institute of Technology Nagoya 466-8555, Japan
223cm-1

Introduction
Alkali niobate is one of the promising candidatesof lead-free piezoelectric material. Li-doped
Na0.5K0.5NbO3 (LNKN) solid solution shows an
excellent piezoelectric property of d33=235 pC/N
when the Li content is 6.0 mol% (L6). The phasetransition of LNKN varies with Li content, andchanges at room tempereature from orthorhombic totetragonal at 5-7 mol% of Li content. However,
LNKN undergoes the temperature-dependent
successive structural phase transition. We haveinvestigated such phase transition by using Raman
scattering to clarify its origin and the related goodpiezoelectric property. Especially, the ferroelectric
property recorded at room temperature was closely
related to the frozen-modes of optical lattice vibration
below room temperature. Therefore, the phonon
behavior of L6 has been investigated more. In thisstudy, IR reflectivity of LNKN ceramics wasmeasured, and the effect of Li incorporation against
the lattice vibration of LNKN was discussed.

150K

140K

Reflectance (a.u.)

130K

120K

110K

Experimental Procedure
The ceramic sample was synthesized by an ordinalsolid-state reaction method. The surface of
Li0.06(Na0.5K0.5)0.94NbO3 (L6) and Na0.5K0.5NbO3(NKN) ceramics were polished and used for IR
reflective study.
Their reflectivity far-infrared spectra were obtained
at 78 and 300 K using the beamline BL6B, then the
spectra were corrected by using Michelsoninterferometer (Bruker, IFS66v) and synchrotron
radiation source.

100

200

300

400

500

600

700

-1

Real part of permittivity (-)

Wave number (cm )Fig. 1. Far-infrared reflectance spectra of L6 ceramics
at 110, 120, 130, 140 and 150 K.

Results and Discussion
Dielectric permittivity of L6 ceramics wascalculated by Kramers-Kronig transform method.
Any dielectric anomalies originated from structuralphase transitions has not been detected below 273 K
in the previous impedance analysis. However, an
optical mode was disappeared during the heating step.
Figure 1 shows the measured reflectance spectra of
non-poled L6 ceramics at 110, 120, 130, 140 and
150 K. The optical mode was observed at 223 cm-1 attemperatures below 150 K. This mode was notdetected for NKN. Figure 2 shows the real part ofpermittivity of L6 ceramics and the fitting result. The
strong peak of the mode could not fit well to a
traditional oscillator model. A strong Li-O mode in
LNKN seems to be correlated with the result, and wewill investigate the further mode analysis for thepoled-samples and single crystal.

@ 110K
60

40

20

Real part of permittivity
0

Fitted
40

50

60

70

80

90

100

Omega (THz)
Fig. 2. Real part of permittivity of L6 ceramics at
110 K and the fitting result based on traditional
oscillator model at 110 K.
Acknowledgement
This research was supported by the Industrial
Technology Research Grant Program in 2007 fromthe New Energy and Industrial Technology
Development Organization (NEDO) of Japan.

128

BL7U

Effects of Temperature Dependent Chemical Potential
on Seebeck Coefficient of TiS2

T. Takeuchi1,2, A. Yamamoto2 and K. Ogawa2
EcoTopia Science Institute, Nagoya University, Nagoya 464-8603, Japan2
Department of Crystalline Materials Science, Nagoya University, Nagoya 464-8603, Japan
1

Temperature dependence of chemical potential wasmeasured for the n-type TiS2 thermoelectric material
by means of angle resolved photoemission
spectroscopy (ARPES), and compared with its
measured Seebeck coefficient. The ARPES
measurements were performed using 22 eV photon
source and the MBS A1 analyzer located at the BL7U
in the UVSOR at Okazaki, Japan.
Figures 1 (a) and (b) show typical example of
ARPES intensity images obtained for the present
sample at 15 K and 140 K along the momentum line
crossing the M point. The intensity-peak varying with
momentum is observable in the ARPES images. This
experimentally revealed fact indicates that the qualityof samples was good enough, and also that theelectronic states in this sample are described with the
Bloch states k. The presence of Bloch states k is
responsible for the metallic electrical conduction
observed for this sample. The energy shift of ARPES
spectra with varying temperature is also observablefor the present sample. This energy shift in ARPES
spectra must be caused by the temperaturedependence of chemical potential.
The chemical potential of n-type materials
generally moves toward lower energy with increasing
temperature. By considering that (a) the chemical
potential of apparatus is kept constant regardless ofthe sample temperature and (b) the chemical potential
observable in the photoemission spectra is pinned at
the chemical potential of the apparatus, it is naturally
understood that the photoemission spectrum at given
temperature shifts towards the opposite energy
direction of the chemical potential shift of samples.
Indeed, the ARPES spectrum of the n-type TiS2, in
which the chemical potential is supposed to movetowards lower energy, was moved towards higher
energy with increasing temperature.

In order to quantitatively extract the temperaturedependence of chemical potential from the measured
ARPES spectra, the ridge in ARPES intensity, thatindicates the energy–momentum dispersion, wasplotted in Fig. 1 (c). From the determinedenergy-momentum dispersion, we identified that thechemical potential shift of г6 meV took place during
the temperature increase from 15 K to 140 K [1].
The contribution of chemical potential to Seebeckcoefficient was calculated from the measured value ofchemical potential shift. We assumed here that thecontribution of chemical potential to Seebeckcoefficient is linearly varying with absolutetemperature as it is described in the Sommerfeld
theory, and mathematically transformed the equation
SP (T ) ; ( P (T )  H F ) / ( e T ) into the following formula.
SP (T ) | T ( P (T )  P (T1 )) / ( e (T 2  T12 ))(1)

The resulting value was SP (T ) 43 ǴV/K at 140 K,
which exceeds 30 % of the measured value, S(140 K)= 135 ȝV/K [1].
The present analysis clearly proved that thecontribution of chemical potential to Seebeckcoefficient is not negligibly small but considerablylarge. We have confirmed that this fact is also true for
the Bi2Te3 and Bi2Se3 thermoelectric materials [2]. It
is confidently argued, therefore, that the proper
understanding of the relation between chemical
potential and Seebeck coefficient should be required
to develop a practical thermoelectric material
possessing a large magnitude of Seebeck coefficient.
[1] T. Takeuchi, A. Yamamoto, K. Ogawa, Mater. Res.
Soc. Symp. Proc. (2011), in press.
[2] A. Yamamoto, K. Ogawa, T. Takeuchi, Materials
Transactions (2011), in press.

Fig. 1. ARPES intensity I(k,e) of TiS2 measured at (a) 15 K and (b) 140 K. In order to observe the states above P, the 140 K
spectrum was divided by the Fermi-Dirac distribution function. The peak energy is plotted as a function of momentum in (c).
The chemical potential shift of 6 meV was caused by the temperature increase from 15 K to 140 K [1].

129

BL7B

Reflection Spectroscopy of 6FDA-Based Polyimide Films
H. Ito, N. Fujita, T. Yabumoto, S. Matsumoto and H. Matsumoto
School of Science & Technology, Meiji University, Kawasaki 214-8571, Japan

Introduction
Photo-alignment is attracting attention as a methodof fabricating the alignment layer in liquid crystal(LC) displays. In this method, a polyimide (PI) film isirradiated by linearly polarized ultraviolet light(LPUVL). In an alignment layer produced by thismethod, the LC molecules are oriented by amechanism that is not understood.
We investigated the optical properties of PI films
irradiated by LPUVL in the ultraviolet (UV) andvacuum ultraviolet (VUV) regions to examine themechanism.
We used fluorine-containing PI that is expected tofind wide application.

Reflection measurements
UV and VUV reflectance spectra of the films were
measured at up to 17 eV with a 3-m normal incident
monochromator (grating: G2 and G3) at BL-7B of
UVSOR. A silicon photodiode sensor was used todetect the reflected light.
Results and Discussion
Figure 2 shows the UV and VUV reflectancespectra of the 6FDA-based PI films from 4.5 to 17 eV.
The solid line indicates the reflectance before LPUVL
irradiation, and the dotted line indicates thereflectance after 8 J of LPUVL irradiation.

Experimental
Materials
We investigated the optical properties of PIs based
on 2,2ƍ-bis(3,4-dicarboxyphenyl)hexafluoro propane
dianhydride (6FDA). Four types of phenylene wereused as the diamine: 1,3-phenylene diamine (mPD),
4-methyl-1,3-phenylene
diamine(MPD),
2,4,6-trimethyl-1,3-phenylene diamine (TMPD), and2,3,5,6-tetramethyl-1,4-phenylene (TeMPD). They
were synthesized at the Nagai lab, Meiji University[1]. Figure 1 shows their chemical structures.

㪘Before LPUVLirradiation
8J irradiation㫄㪧㪛㪍㪝㪛㪘 㪄㪘㪧㪛㪝㪛㪘 㪄㪤

Reflectance (a.u.)

㪙
㪍

㪙
㪫㪤 㪧㪛
㪍㪝㪛㪘 㪄

㪘㪙㪍㪝㪛㪘 㪄

㪘㪫㪼㪤 㪧 㪛

㪙

5

7

9

11

13

PhotonEnergy (eV)1517Fig. 2. Reflectance of 6FDA-based PIs.
The spectra of the PI films are fairly similar andconsist of a prominent band group around 5.5 eV (theA band) and another broad band group (the B band)near 17 eV.
The appearance of the A band varied in the four
spectra taken before LPUVL irradiation. The reason
may be differences in the diamines.
A significant spectral change was observed after
LPUVL irradiation of the 6FDA-TeMPD film. But
the spectrum of the 6FDA-mPD film showed the leastchange after irradiation. This difference in the degreeof spectral change implies that the methyl groups
added to phenylene in the diamine affect thephoto-alignment of PI films.

Fig. 1. Chemical structures of 6FDA-based PIs.
Sample preparation
We dissolved the PIs with N,N-dimethylacetamide(DMAc). The PI films were prepared by spin coating
on quartz substrates. The PI-coated substrates were
baked at 100°C for 1 h. The film thicknesses were
controlled to about 400 nm.
These PI films were irradiated by LPUVL from a
200 W Hg–Xe lamp. Linearly polarized light was
obtained by reflection from a quartz plate at
Brewster’s angle. The degree of polarization was 99%
at 254 nm. The irradiation intensity was about
700 ȝW/cm2 at 254 nm. We irradiated the films with
8 J of LPUVL.

[1] S. Miyata et al., J. of Appl. Polym. Sci. 107(2008) 3933.

130

BL7B

Auger-Free Luminescence Emitted from A2ZnCl4 (A = Cs, Rb) Crystals
A. Ohnishi1, M. Saito1, M. Kitaura1, M. Sasaki1, T. Kajitani2 and M. Itoh2
1
Depeartment of Physics, Yamagata University, Yamagata 990-8560, Japan2
Department of Electrical and Electronic Engineering, Shinshu University,
Nagano 380-8553, Japan

The Auger-free luminescence (AFL) is a peculiar
type of intrinsic luminescence in ionic crystals inwhich the valence excitation through a nonradiative
Auger process is energetically forbidden. The AFL ischaracterized by a high quantum yield and a short
lifetime of the order of ns. Therefore, it is very usefulas the fast scintillator for high-energy physics,
positron emission tomography and so on.
In the present work, we have investigated theluminescence of A2ZnCl4 (A = Cs, Rb) crystals under
the outermost core excitation, in order to find theexistence of AFL.
A2ZnCl4 (A = Cs, Rb) crystals were prepared byevaporating a stoichiometric mixture of solutions of
CsCl or RbCl and ZnCl2. Experiments were
performed at the BL7B beamline of UVSOR. The
emission spectra in this measurement were notcorrected for the spectral dispersion. The excitation
spectra were corrected for the spectral distribution of
the light source.
When Cs2ZnCl4 was excited at 10 K at 21.4 eV,
two luminescence bands appeared at 3.2 and 4.2 eV.
The blue line of Fig. 1 shows the excitation spectrum
for the 4.2 eV band. The excitation spectrum for the
3.2 eV band was practically the same as that for the
4.2 eV band. For reference, the reflection spectrum
measured for the cleaved surface of Cs2ZnCl4 is also
shown by the red line. Sharp peak at 7.3 eV is
assigned to the n = 1 exciton absorption, which is dueto the transition from the Cl 3p valence band to the
conduction band. Sharp peaks at 13.2 and 13.9 eV
may be core exciton bands due to the transitions fromthe outermost core band to the conduction band.
The excitation spectrum rises at around 13.5 eV as
the photon energy is increased. This threshold energy
corresponds almost to the transition energy between
the outermost core band and the conduction band, andthus both of the 4.15 eV and 3.31 eV bands aresupposed to be the AFL bands due to the radiative
recombination between a valence electron and anoutermost core hole. It is noteworthy that thethreshold energy of the AFL is slightly higher thanthe lowest-energy core exciton band at 13.2 eV. This
fact suggests that core excitons decay nonradiatively
through Auger-electron emission.
Luminescence decay measurements were also
performed using a time-correlated single-photon
counting technique under the single bunch operation.
It was proved that, at 10 K, the 4.15 eV and 3.31 eV
bands are composed of a fast decay component withthe lifetime of about 1.8 ns.

In Rb2ZnCl4 crystal, two AFL bands peaking at 2.8
and 4.3 eV were observed at 10 K under excitation at
21.4 eV. In Fig. 2, excitation (blue) and reflection(red) spectra of Rb2ZnCl4 are shown. The excitation
spectra for the 2.8 and 4.3 eV bands were the same.
Sharp peaks at 7.3 and 16.5 eV in the reflection
spectrum is assigned to the n = 1 exciton and core
exciton absorption bands, respectively. From theexcitation spectrum, it turns out that both bands arestimulated with photons above about 14 eV. This
result suggests that the 2.8 and 4.3 eV bands are the
AFL bands in Rb2ZnCl4. However, it should benoticed that the threshold energy is located at theenergy lower than the core exciton peaks. The reason
for such a disagreement is under investigation now.

Fig. 1. Excitation spectrum at 10 K for the 4.2 eVband in Cs2ZnCl4 (blue line). Reflection spectrum(red line) is also shown for reference.

Fig. 2. Excitation spectrum at 10 K for the 4.3 eVband in Rb2ZnCl4 (blue line). For reference, reflection
spectrum is also shown by a red line.

131

BL8B

Electronic Structure of Delocalized Singlet Biradical Ph2-IDPL Solid Film
1

K. Kanai1, Y. Noda2, K. Kato2, T. Kubo3, K. Nakasuji4 and K. Seki2
Department of Physics, Faculty of Science and Technology, Tokyo University of Science,
Noda 278-8510, Japan2
Department of Chemistry, Graduate School of Science, Nagoya University
Nagoya 464-8602, Japan3
Department of Chemistry, Graduate School of Science, Osaka University,
Osaka 560-0043, Japan
4 Fukui University of Technology, Fukui 910-8505, Japan

Recently, Kubo et al. have successfully isolated abiradical hydrocarbon, diphenyl derivative ofs-indacenodiphenalene (Ph2-IDPL) which possesses a
relatively small HOMO-LUMO gap Eg and strong
intermolecular interactions [1]. The partial occupationof the LUMO is thought to contribute to stabilizationof the intermolecular interaction, leading tointermolecularcovalency.
Thus,astrongintermolecular interaction is characteristic of singlet
biradicals.
In this work, the electronic structure of a Ph2-IDPL
film has been investigated. A small Eg compared withthat of typical Ǹ-conjugated small molecules wasobserved even for the amorphous film of Ph2-IDPL.
This result indicates that the small Eg is an importantcharacteristic of the singlet biradical electronic
structure. Moreover, the photon energy dependence ofultraviolet photoemission spectra shows that thestacked Ph2-IDPL molecular chain in thepolycrystalline film develops an energy band
structure in the direction of the surface normal of thefilm. The intermolecular covalency therefore evolves
into the quasi one-dimensional energy band along themolecular stacking direction.
The UPS measurements were performed at theBL8B of the UVSOR.
Figure 1 shows UPS and IPES spectra of Ph2-IDPL
amorphous film on graphite substrate. The simulated
spectra for the monomer well explain the observed
spectra over the entire energy range. Since strong
intermolecular coupling among the molecules is
expected to significantly alter its electronic structure,these results clearly show that this Ph2-IDPL film ongraphite is composed of the monomers. Notably, theobserved Eg is rather small. From the results shown in
Fig. 1, Eg is about 1.2 eV, which is smaller than thoseof many other typical Ǹ-conjugated molecules.
Figure 2 shows the UPS and IPES spectra of
Ph2-IDPL crystal and the polycrystalline film. The
film spectra differ from the spectra for a monomer. In
particular, the electronic structure of the film aroundEF produces broader spectral features compared to themonomer spectra. The UPS spectrum of the film
closely resembles that of the crystal. The Eg for thepolycrystalline film was estimated to be 0.84 eV.
These results well explains the previously reported
ambipolar field effects of amorphous Ph2-IDPL film

Fig. 1. UPS and IPES spectra of Ph2-IDPL film ongraphite. Simulated spectra labeled “a” were
calculated by the DFT method for monomer.

Fig. 2. UPS and IPES spectra of Ph2-IDPL film ongraphite and simulated spectra for the dimer.
by Chikamatsu et al. [2]. The small Eg would result ina small injection barrier both for electrons and holes
at the electrode interface. The extremely small Eg
observed in the polycrystalline Ph2-IDPL film ispossibly caused by strong intermolecular coupling.
[1] T. Kubo et al., Angew. Chem. Int. Ed. 44 (2005)
6564.
[2] M. Chikamatsu et al., Appl. Phys. Lett. 91 (2007)
043506.

132

BL8B

Interface Electronic Structuresof a High Performance Organic Planar Hetero-junction Solar Cell

Y. Nakayama1, J. Wagner2, Y. Tanaka3, S. Machida3,
T. Nishi1,4, A. Opitz2, W. Brütting2 and H. Ishii1,3
1Center for Frontier Science, Chiba University, Chiba 263-8522, Japan
2
Institute of Physics, University of Augsburg, Augsburg 86135, Germany
3
Graduate School of Advanced Integration Science, Chiba University, Chiba 263-8522, Japan
4
UVSOR Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
Organic photovoltaics (OPVs) are expected as next
generation power conversion devices complementary
to conventional Si-based solar cells. For OPVs, a bulk
hetero-junction concept is generally adopted toredeem small exciton diffusion length, neverthelessthis concept also has disadvantages (e.g. spoiled
effective mobility Peff) that depress the power
conversion efficiency. A planar hetero-junction (PHJ)
OPV based on diindenoperylene (DIP) and C60 wasrecently reported to reveal high performance [1]. DIP
has to be deposited on a heated hole injection layer(HIL) to achieve high efficiency. The origin(s) of the
improved performance, which may be attributed to a
reduced interface energy barrier and/or enhanced bulk
Peff of DIP, have not been clarified yet. In this study,
we determined the interface electronic structures of
the DIP-C60 based high performance PHJ-OPV.
EDOT:PSS substrates were prepared by a method
described in Ref. [1] in nitrogen atmosphere. The
work function of the substrate varied in the range of
4.9 r 0.2 eV. DIP was deposited (ca. 0.05 nm/s) onto
the polymer substrates at ca. 100 oC in high vacuum.
Note that heating of the polymer up to the deposition
temperature significantly increased the work function[Fig. 1 (a)]. C60 was deposited (ca. 0.02 nm/s) onto a
DIP thin film at RT. All PES spectra were taken withthe photon energy of 40 eV at BL8B in UVSOR. For
the organic films of 1.5 nm or thicker, sample bias of+5 V was applied in order to avoid sample charging.
Figure 1 shows the UPS secondary electron cutoff(SECO) and spectra of a PEDOT:PSS substrate and
DIP overlayers. The onset of the highest occupied
molecular orbital (HOMO) of DIP ()h) appeared at
0.2 eV from the Fermi level (EF). This value isapparently smaller than the literature (0.45 eV [1]) of
DIP films deposited at RT. In addition, a negative
vacuum level shift (') is evinced by SECO, which isalso strikingly in contrast to the RT deposition case.
The HOMO level offset ('E) at the C60-DIP
interface is estimated to be 1.0 eV as shown in Fig. 2.
Slightly positive ' was revealed at this interface. Theenergy level diagram of the topical OPV cell is shownin Fig. 2 (d). Despite aforementioned HOMO energy
difference, 'E is identical to that of the interface
formed on DIP deposited at RT (1.05 eV [1]). The
reduced )h by high temperature DIP deposition
should be a reason of improved OPV performance.

Fig. 1. (a) SECO and (b) UPS of DIP deposited on aheated PEDOT:PSS substrate. A simulated density-ofstate curve of DIP is shown in the upper panel of (b).

Fig. 2. (a) UPS spectra of C60 deposited on DIP(9 nm)/PEDOT:PSS. (b) Magnified spectra of (a). Thin redlines indicate hypothetical DIP contribution. (c)SECO spectra of C60 on DIP/PEDOT:PSS. (d) An
energy level diagram of the hetero-interfaces in the
DIP-C60 based high performance PHJ-OPV [1]. The
numbers represent the energy with respect to EF.
[1] J. Wagner et al., Adv. Funct. Mater. 20 (2010)
4295.

133

BL8B

Electronic Structures of Room Temperature Ionic Liquids
Studied by UPS and IPES
T. Sato1, T. Iwahashi1,2, Y. Sakai1, K. Kanai3, T. Nishi4 and Y. Ouchi1
1

Department of Chemistry, Nagoya University, Nagoya 464-8602, Japan
Venture Business Laboratory (VBL), Nagoya University, Nagoya 464-8601, Japan
3
Faculty of Science and Technology, Department of Physics, Tokyo University of Science,
Chiba 278-8510, Japan
4
UVSOR Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
2

Intensity / arb. units

Room temperature ionic liquids (RTILs) are salts
RTILs other than [DEME]TFSA even if the calculated
in a liquid phase at ambient temperature and pressure,energy gap has some differences for each cation.
and have attracted much attention due to their unique
Considering that the calculated energy gap of thenature of nonvolatility, high conductivity, and wide
TFSA anion is much narrower than that of theelectrochemical window. Recently, RTILs are utilized
quaternary ammonium-based cations except theas a potential gate material for organic field effect[DEME]+ cation, we can propose that both the HOMO
transistors as well as capacitors, fuel cells, organicand LUMO levels of the ammonium-based RTILs
photovoltaic cells, and Li-ion batteries. In suchexamined in this study, except [DEME]TFSA, isapplications, electronic structure, such as HOMO andattributed to the TFSA anion. The observed energy
LUMO levels, of RTIL strongly affects thegaps of the quaternary ammonium-based RTILs
performance of the devices, and hence understanding
without methoxy group are also ~2 eV wider than theof the electronic structures of RTILs should be a key
calculated gap of the TFSA anion similarly to the caseissue for further development of such potentialof [emim]TFSA. The energy gap broadening of ~2 eV
applications.
seems to be a condensed state effect for the TFSAIn this brief report, we examined the HOMO and
RTILs.
LUMO levels and the energy gaps of the TFSA
RTILs with cations of 1-ethyl-3-methylimidazolium(g)([emim]+), 1-propyl-1-methyl pyrrolidinium ([P13]+),[DEME]
1-butyl- 1-methyl pyrrolidinium ([P14]+), 1-propylBF4+
1-methylpiperidinium ([PP14] ), N,N,N-trimethyl(f)[DEME]
propylammonium ([TMPA]+), and [DEME]+ cations
TFSA
by UPS and IPES. UPS measurements were carried(e)out at the beamline BL8B2 of UVSOR facility at[TMPA]TFSA
Institute for Molecular Science. IPES measurements
were conducted at Nagoya University. The samples(d)[PP13]
were prepared by putting one droplet of RTIL on
TFSA
scratched Au substrates. All measurements were
(c)
[P14]
performed under ultra-high vacuum condition of 
í8
í7
TFSA
2 × 10 ~ 1× 10 Pa.
(b)
Figure 1 shows the UPS and IPES spectra of the[P13]
TFSA RTILs with respect to the Fermi level of Au.
TFSA
The results for [DEME]BF4 are also shown as a(a)[emim]
reference. As is seen in Fig. 1, the HOMO and LUMO
TFSA
levels of each RTIL differ from those of the isolated
1)ions estimated from the ab initio calculations, and㪈㪇㪌㪇㪄㪌 㪄㪈㪇
these energy shifts are toward the lower and higher
energies for the cation and the anion, respectively. This
Binding energy / eV vs. Fermi level
energy shift is understood as a contribution ofFig. 1. UPS and IPES spectra of (a) [emim]TFSA,
Madelung energy which stabilizes and destabilizes the(b) [P13]TFSA, (c) [P14]TFSA, (c) [PP13]TFSA, (e)orbital energies of the anion and the cation,[TMPA]TFSA, (f) [DEME]TFSA, and (g)respectively. On the other hand, we found that the[DEME]BF4. The abscissa is set the Fermi level ofenergy gaps of the quaternary ammonium-based
the Au substrate as zero. Inset shows the MO pattern
RTILs are wider with shifting the HOMO and LUMO
of HOMO of [DEME]+ cation. The triangles on thelevels higher and lower binding energy side,
UPS spectra represent the peaks originating from O
respectively, relative to those of [emim]TFSA as
2p of [DEME]+ cation.
shown in Fig. 1. It is noteworthy that not only theenergy gap but also the HOMO and LUMO levels are[1] K. Kanai et al., J. Electron Spectrosc. Relat.
almost the same for the quaternary ammonium-based
Phenom. 174 (2009) 110.

134

BL8B

Experimental Evidence of Acceptor-Like Levels
in SrGa2S4:Eu2+ Phosphor Thin-Films
Y. Suzuki1, M. Kitaura1, S. Tanaka2, H. Kominami3, K. Hara3, A. Ohnishi1 and M. Sasaki1
1Faculty of Science, Yamagata University, Yamagata 990-8560, Japan
2
Interdiscriplinary Faculty of Science and Engineering, Shimane University, Matsue 690-8504, Japan
3
Research Institute of Electronics, Shizuoka University, Hamamatsu 432-8011, Japan

135

6

3

(a)

㩷

1

2
SrGa2S4:Eu
hQ =18eV
0

㩷

(b)

2

4
㩷

INTENSITY (arb. units)

The type of charge carrier in narrow-gap
semiconductors can be artificially controlled by theaddition of impurities. The exchange of carrier type
between electron and holes is challenge in wide-gap
semiconductors with ionic character, because thep-type conduction due to holes is hard to occur due tothe large effective mass and strong phonon coupling.
Therefore, there have been few reports that succeeded
the carrier type exchange in the wide-gap
semiconductors. Recently, Tanaka et al. have reported
that the hot carrier type in inorganic
electroluminescence (EL) thin-film devices is
exchangeable for different rare-earth ion doping [1].
Strontium thiogallate (SrGa2S4) is also one of the hostmaterials showing such carrier type exchange. This
material has the feature of n-type in Ce3+-doping. The
carrier type is changed to p-type in Eu2+-doping.
Judging from the practical sense mentioned above,these results are to be very surprising. On the other
hand, there are ambiguous points on the fundamental
physic of such p-type conduction, e.g., the presenceof acceptor-like levels and the contact type at theinterface between phosphor and metal electrodethin-films. The elucidation of these points is our main
purpose:hence,
ultraviolet
photoelectron
spectroscopy (UPS) experiment was carried out.
Eu2+-doped SrGa2S4 thin-films were deposited onquartz substrates heated at 200 oC using electron
beam sources. The concentration of Eu2+ ions was set
to be 2 mol% in the preparation. The thickness was
adjusted to be 100 nm using quartz oscillators. The
as-deposited thin-films were annealed at 850 oC in the
H2S+Ar gas mixture for 30 minutes, in order tosupply with sulfur. The composition and crystalinityof SrGa2S4:Eu2+ thin-films were checked by X-ray
diffraction measurement. In UPS experiment, thesurface was cleaned in vacuum with Ar+ ion gun. An
appearance of the clean surface was judged from thepeak energy and spectral shape of the Ga-3d peak.
Figure 1 (a) shows the typical UPS spectrum of
SrGa2S4;Eu2+ thin-films at room temperature under
excitation with photons at 18 eV. The absolutebinding energy was determined by referring to the
Fermi level of tantalum plates for the fixing ofsamples. In Fig. 1 (a), a prominent band appears
around 5.0 eV. This band is ascribed to the valenceband dominated by S-3p orbitals [2]. As shown in Fig.
1 (b), this band starts at Eth=0.86 eV. The fundamental
absorption edge of SrGa2S4 is located at 4.0 eV,

10

Eth=0.86eV
EF

2+

0

5
0
4
2
BINDING ENERGY (eV) 㩷

0Fig. 1. UPS spectrum of SrGa2S4:Eu2+ thin-films atroom temperature under excitation with photons at
18 eV: (a) Valence band region and (b) near-Fermi
level region.
according to the absorption spectrum at room
temperature [3]. We suppose that the fundamentaledge is due to the band-to-band transition, because no
exciton absorption peaks were found in the reflection
spectrum at room temperature. In this case, thethreshold energy for the valence band is expected tobe 2.0 eV, higher than the value of Eth. This result canbe explained by considering that acceptor-like levels
are on the valence band. Therefore, our experimental
results are adopted as one of the evidence supportingthat holes are dominant charge carrier in the
SrGa2S4:Eu2+ thin film.
UPS experiment allows us to know not only theelectronic structure but also the physical constants
characteristics of materials. We determined theionization potential of SrGa2S4:Eu2+ to be 4.96 eV.
This value is higher than the work functions of metals
except gold [4]. From this fact, it is suggested that theinterface between SrGa2S4:Eu2+ and metal electrodethin-films shows a tendency to form the Shottky-typebarrier for hole conduction. The Shottky-type barrier
may also involve the p-type conduction in the
SrGa2S4:Eu2+ thin-film.
[1] K. Tanaka and S. Okamoto, Appl. Phys. Lett. 89(2006) 203508.
[2] M. Kitaura et al., UVSOR Activity Report 37(2010) 137.
[3] C. Chartier et al., Electrochem. Solid-State Lett. 9(2006) G53.
[4] H. B. Michaelson, J. Appl. Phys. 48 (1977) 4729.

BL8B

Photoinduced Shift of the Photoelectron Spectra at the
Zn-Phthalocyanine/C60 Interface
S. Tanaka, K. Fukuzawa and I. Hiromitsu
Interdisciplinary Faculty of Science and Engineering, Shimane University,
Matsue 690-8504, JapanThe electronic structure of the organic/organic andthe organic/electrode interface of the organic
electronics devices, such as organic light emitting
diode, organic solar cell, and organic thin filmtransistor, has attracted much interest since it affects
the performance of these devices [1]. For example,
the authors have recently reported that thelithium-phthalocyanine thin film on indium-tin oxide(ITO) improves the device performance of theorganic solar cell. This improvement was caused by
the formation of the advantageous electronic structure
for the hole extraction at the ITO/organic layer [2, 3].
It is widely accepted that the electronic structure of
the organic layer in the devices can be inferred fromthe photoelectron spectra of the organic thin films.
The photoelectron spectroscopy is a powerful tool forthe estimation of the electronic structure of thesamples. Thus, the photoelectron spectroscopy under
the solar light irradiation can be useful to understandthe electronic structure of the organic solar cells
under the working condition. In the present study, we
focused on the electronic structure of adonor/acceptor (DA) heterojunction of the organic
solar cells. As a typical DA interface of the organic
solar cell, the Zn-phthalocyanine (ZnPc)/C60 interface
was studied.
The photoelectron spectra were measured as afunction of the thickness of the organic layer on an
ITO substrate in situ. For the observation of thephotoinduced effect on the electronic structure, thephotoelectron spectra under the light irradiation werecompared with that under dark. A solar simulator wasused as the light source for the irradiation. The
simulated airmass 1.5 solar illumination was exposedto the sample in a vacuum chamber through a
viewport. The distance from the viewport to thesample was approximately 30 cm. All thephotoelectron measurements were carried out with aphoton energy of 40 eV at room temperature. The
overall energy resolution was approximately 0.2 eV.
Figure 1 shows the photoelectron spectra of the C60(1 nm) layer on the ZnPc (15 nm) layer. The abscissa
indicates the kinetic energy of photoelectron. The red
dots and the black dots represent the photoelectron
signals with and without the light irradiation,
respectively. The peak at around 33.8 eV is thehighest occupied molecular orbital (HOMO) of ZnPc.
The HOMO peak position was shifted toward higher
kinetic energies approximately 0.06 eV under thelight irradiation compared with the HOMO peakunder dark. In contrast, no significant shift wasobserved on the HOMO peak position of the ZnPc

film without the C60 over layer (not shown). The
photoinduced shift showed dependences both on thelight intensity and on the thickness of the C60 layer.
In general, there are several possibilities for theorigin of the spectrum shift: the degradation of themolecules, the charging up of the sample, the thermal
effect, and the photoinduced charge redistribution.
We have considered that the shift was caused by thephotoinduced charge redistribution at the ZnPc/C60
interface from the following reasons. The degradationof the molecules should cause nonreversible changes
on the photoelectron spectrum, however, the present
shift was observed reversibly. The charging up is also
excluded, since the charging up would cause a shift
toward the lower kinetic energies and deform thespectral shape. Although the slight increase of thetemperature of the substrate was observed while thelight irradiation, the effect of the temperature on thephotoelectron spectra is empirically negligible in thepresent temperature range. Hence, the shift of the
HOMO peak can be attributed to the photoinduced
effects on the charge redistribution at the DA
interface. These preliminary results indicate that thephotoelectron spectroscopy under the light irradiation
is a valuable method to investigate the electronic
structure of the organic solar cell system under theworking conditions.
C60 (1nm) / ZnPc (15nm) / ITO

Intensity / Arb. units

hQ = 40 eV
Dark
Under light irradiation

0.06eV

32

33

34

35

36

Kinetic energy / eV
Fig. 1. Photoelectron spectra of C60 (1 nm) / ZnPc
(15 nm) on the ITO substrate with (red) and without(black) the light irradiation.
[1] H. Ishii et al., Adv. Mater. 11 (1999) 605.
[2] S. Tanaka et al., UVSOR Activity Report 37(2009) 80.
[3] S. Tanaka et al., Appl. Phys. Lett. 97 (2010)
253306.

136

BL8B, BL4B

Identification of Catalytic Active and Inactive Species in Nitrogen Doped
TiO2 Photocatalyst by N K-edge XAFS Analysis
T. Yoshida1, S. Muto2 and E. Kuda2
1
EcoTopia Science Institute, Nagoya University, Nagoya 464-8603, Japan2
Department of Materials, Physics and Energy Engineering, Nagoya University,
Nagoya 464-8603, Japanchemical environment similar to that in TiN. Morethorough observation suggested that double-peakaround 400 eV in Fig. 1 (c) shifted to the lower
energy side compared with that of TiN, which waswell reproduced by the theoretical prediction using
FEFF code when N occupies one of the O sites ofTiO2 [Fig. 1 (d)]. On the other hand, the XANES
spectrum of the sample implanted with the N+ fluenceof 5 x 1021 m-2 followed by heat-treatment (almost
inactive to visible-light : I-cat) shows a distinct single
peak around 401 eV [Fig. 1 (e)]. This feature of the
XANES spectrum was succesffuly reproduced bytheoretical simulations based on the model where an
O atom in TiO2 was replaced by NO2 [Fig. 1 (f)].

Introduction
In the field of catalytic chemistry, a specificfunction in a solid catalyst crucially depends on thechemical state of a specific active component (activesite). In this context X-ray absorption fine structure(XAFS) is expected as one of the most powerful
techniques for chemical state analysis of the active
site, because it provides information on the local
structure and electronic states around the specificelement of interest. In the present study, asophisticated combination of XAFS and a theoretical
calculation method allows us identification ofphotocatalytic active and inactive species in nitrogen
doped TiO2 catalysts.
Experimental
The samples used in this study were TiO2 (1 0 0)
single crystals (5 x 5 x 0.5 mm3), supplied by
Furuuchi Kagaku, Japan. Mass analyzed 100 keV N2+
ions (50 keV/ N+ ion) were injected into the samples
at room temperature, perpendicular to the sample
surface. The N+ fluence ranged from 1 to 5 x 1021 m-2.
After the ion implantation, parts of the samples were
heat-treated at 573 K for 2 hours in air.
A typical photocatalytic experiment consisted of
placing the N+-implanted sample in 0.5 ml of aqueous
methylene-blue (MB) solution (9.8 ȝmol/L) andsubsequent exposure to visible-light using a 15 W Xe
lamp with a cut filter for O > 430 nm.
N K-edge XANES spectra of the N+-implanted
TiO2 samples were measured at the BL8B&4B stationof UVSOR-II at the Institute for Molecular Science,
Okazaki, Japan. Data were recorded at room
temperature in total electron yield mode, and the
X-ray energy dependence of the N Auger electron
yield was monitored.

i) (NO)ih) (NO)sg) (N)iXANESf) (NO 2)se) I-catd) (N)sc) H-cat.

b) TiN calc.

Results and Discussion
The photocatalytic activity reached its maximum ata fluence of 3 x 1021 m-2 and then decreased with thefluence. The sample implanted at a fluence of
5 x 1021 m-2 followed by heat-treatment at 573 K was
almost photocatalytically-inactive under visible-lightirradiation.
Figure 1 shows N K-edge XANES spectra of the
N+-implanted TiO2 samples, a TiN powder andtheoretical spectra of several structure models.
Common XANES features in Fig. 1 (a) and (c)suggest that N in the sample implanted by 3 x 1021
m-2 (highest active photocatalyst: H-cat) is in aa) TiN exp.

395

400

405

410

X-ray energy (eV)415

Fig. 1. N K-edge XANES spectra of (a) TiN crystal,(c) the sample N+-implanted with 3 x 1021 m-2, (e)that with 5 x 1021 m-2 followed by heating at 573 K
for 2 h, and theoretical spectra of structure modelsfor (b) TiN, (d) substitutional N-doping, (e)substitutional NO2-doping, (g) interstitial N-doping,(h) substitutional NO-doping and (i) interstitial
NO-doping.

137