BL2A

The Measurement of Soft X-Ray Excited Optical Luminescence of a Ge+
Implanted Silica Glass Sample
T. Yoshida1, S. Muto2 and H.Yoshida3
1
EcoTopia Science Institute, Nagoya University, Nagoya 464-8603, Japan
Department of Materials, Physics and Energy Engineering, Nagoya University, Nagoya 464-8603, Japan3
Department of Applied Chemistry, Graduate School of Engineering, Nagoya University, Nagoya 464-8603,
Japan
2

109

soft X-ray, and show the spectra in Fig. 2 together
with Ge L2,3-edge XANES of a Ge (0) powder
sample recorded in the TEY mode. It is interesting tosee good correspondence among them, and this result
may indicate that the Ge (0) species relate to thegeneration of 3.1 eV PL.

Intensity (arb. unit)[1] L. Rebohle, J. von Borany, R. A. Yankov, W.
Skorypa, I. E. Tyschenko, H. Froeb and K. Leo, Appl.
Phys. Lett. 71 (19) (1997) 2809.
[2] M. L. Brongersma, A. Polman, K. S. Min, E.
Boer, T. Tambo, H. A. Atwater, Appl. Phys. Lett. 72(1998) 2577.
[3] Skuja L. N. and Etzian W., Phys. Status Solidi A
96 (1986) 191.

X1/2

Ge+ 2x1016/cm2

Ge+ 2x1015/cm2

unirrad.

2.0

2.5 3.0 3.5 4.0
Photon energy /eVGe(0)Gepowder
metal

1180

1200

1220

1240

1260

X-ray energy (eV)Photon yield (arb. unit)

Fig. 1. PL spectra excited by 5 eV light for silica
glass before and after Ge+ implantation.

Intensity (arb. unit)

The
discoveryofroom
temperature
photoluminescence (PL) from indirect band gap
materials has generated important interest in this field.
In particular, Ge doped silica glass is expected toapply to a light emission source for communication ina LSI chip as well as to single electron devices, since
it shows blue and violet electroluminescence around
3.1 eV [1, 2]. Although the Ge doped silica glass has
been studied extensively, the origin of the 3.1 eV PL
and the emission mechanism still remain
controversial. The purpose of this paper is to discuss
the relation between the optical properties and thechemical state of Ge atoms in a silica glass, applying
X-ray absorption fine structures (XAFS).
The sample used in this study was a synthesized
silica glass (T-4040, OH content: 800ppm) producedby Toshiba Ceramics, Japan, 13 mm in diameter and
2 mm in thickness. Mass analyzed Ge+ ions of 30 keVwere injected into the samples at room temperaturenormal to the sample surface. The Ge+ fluence ranged
from 2 x 1014 to 2 x 1016 cm-2. The measurement ofluminescence of the samples induced by soft X-ray
irradiation was carried out on the beamline BL1A andBL2A at UVSOR, Institute for Molecular Science.
The luminescence was focused by a lens in the UHVchamber to the monochromator (CP-200, JOBINYVON) and detected by a multi-channel analyser(OMA III, EG&G PRINCETON APPLIEDRESEARCH).
PL spectra before and after the Ge+ implantation
were measured at room temperature using a Hitachi
F-4500 spectrometer.
Fig. 1 shows PL spectra excited by 5 eV light
before and after Ge+ implantation. A weak and broadband is observed before implantation, while a newemission band around 2.7 eV appears after the
implantation of 2 x 1015 cm-2. This PL band was
ascribed to the B2Į oxygen deficient centers [3]associated with the displacement damages. When the
fluence is more than 2 x 1016 cm-2, a sharp andintense 3.1 eV PL band replaces the 2.7 eV band.
We tried to measure Ge L2,3-edge XANES spectraof the 2 x 1016 cm-2 implanted sample in the total
electron yield mode (TEY), but the concentration of
Ge atom were too low to measure. On the other hand,
we found that an intense soft X-ray induced
luminescence of this sample is observed at 3.1 eV.
Therefore, we measured the excitation spectra of the
3.1eV band by changing the incident energy of the1280 1180

16

Ge 2x10 /cm2

L3-edge

1200

1220

L2-edge

1240

1260

1280

X-ray energy (eV)Fig. 2. Comparison between the luminescence yield
spectrum of the 3.1 eV band for a silica glass
irradiated by Ge+ and Ge L2,3-edge XANES spectrumof a Ge powder.

BL2A㩷

Local-Structures of P of LiMnFePO4 Doped with Mn
A. Nakahira1, 2, S. Hayashi1, M. Sato2, Y. Takamatsu1, Y. Nishio1, Y. Kawabe1, S. Misu1,
K. Kumadani1, T. Shirai1 and H. Aritani3
1

Faculty of Engineering, Materials Science and Engineering Dept, Osaka Prefecture University, 1-1 Gakuen-cho,Naka-ku, Sakai 599-8531, Japan
2
Kansai Center, Institute of Materials Research (IMR), Tohoku University, 1-1 Gakuencho, Naka-ku, Sakai
599-8531, Japan
3
Faculty of Engineering, Saitama Institute of Technology, 1690 Fusai-ji, Fukaya 369-0293, Japan
LiFePO4. Figure 1 shows the results of P-K edge forthe powders of LiFePO4 added with Mn(LiMnFePO4). The P-K edge of (NH4)2HPO4 wasmeasured as a reference material. The peaks of
LiFePO4 and LiFePO4 added with Mn (LiMnFePO4)
were higher than (NH4)2HPO4 as a reference material.
Furthermore, LiFePO4 added with Mn (LiMnFePO4)
had the higher energy of P-K edge peak, compared tomonolithicLiFePO4. It is obvious that the peak shiftof P-K edge for LiFePO4 added with Mn(LiMnFePO4) was attributed to the addition of Mn
into LiFePO4.

Normarized absorption (a.u.)

LiFePO4 (lithium iron phosphate) is an interestingalternative positive electrode material for lithium andlithium-ion rechargeable batteries application becauseof due to its low cost and good specific capacity.
Although these LiFePO4 (lithium iron phosphate)materials have some advantages in terms ofenvironmental benignity, potential low-cost synthesis,
cycling stability, and high temperature capability, themajor drawback point for LiFePO4 is the poor lithium
ion conductivity, which have limited its
power-demanding applications such as hybrid electricvehicles and other power generating equipments.
A number of various efforts has been made in thepast decade to solve the conductivity problem for
LiFePO4. Nowadays, the main problems for LiFePO4
is still the poor lithium ions conductivity caused by
the low intrinsic electronic conductivity. Therefore,
the modification of LiFePO4 added with Mn havebeen attempted. In this study, LiFePO4 added withMn (LiMnFePO4) was synthesized through thechemical synthesis of LiFePO4 materials by thesolution processing process and the heat-treatments inorder to improve the specific capacity and cyclingstability etc. The purpose of this study is toinvestigate the P-K edge local structures for LiFePO4
added with Mn (LiMnFePO4).
The synthesis of LiFePO4 and LiFePO4 added withMn (LiMnFePO4) was carried out using LiOH,(NH4)2HPO4, MnSO4 and FeSO4, and the ascorbic
acid was used as a carbon source. These stock
solutions were prepared for 1 mol/L LiOH, 0.5 mol/L(NH4)2HPO4, 0.5 mol/L FeSO4 and 0.5 mol/L MnSO4.
The synthetic reaction was carried out through thehydrothermal treatments in autoclave bomb at various
temperatures. The products were aged for various
times and filtered with Buchner funnel. The products
were dried at 323K for 24 hours and ground withalumina mortar. The products were characterized by
XRD and FT-IR and their microstructures by SEM
and TEM. The local structures around P-K edge forthe powders of LiFePO4 added with Mn(LiMnFePO4) were evaluated and characterized bymeasuring X-ray adsorption near edge structure(XANES) at BL2A at UVSOR in Okazaki.
The components of products of LiFePO4 addedwith Mn (LiMnFePO4) were evaluated by XRD andthe products were identified to be olivine type(NH4)2HPO4SynthesisͲA1
SynthesisͲA2

2150

2160

Energy / keV
Fig. 1. XANES results of P-K edge for LiFePO4
added with Mn (LiMnFePO4).

110

BL2A

Evaluation of Al Local Structure of Cu-Al based Talcite Clay
A. Nakahira1, 2, S. Hayashi1, M. Sato2, Y. Takamatsu1, S. Misu1, K. Kumadani1 and H. Aritani3
1

Faculty of Engineering, Materials Science and Engineering Dept, Osaka Prefecture University, 1-1 Gakuen-cho,Naka-ku, Sakai 599-8531, Japan
2
Kansai Center, Institute of Materials Research (IMR), Tohoku University, 1-1 Gakuencho, Naka-ku, Sakai
599-8531, Japan
3
Faculty of Engineering, Saitama Institute of Technology, 1690 Fusai-ji, Fukaya 369-0293, Japan

Hydrotalcite is one of unique clay minerals among
the layered double hydroxide (LDH) materials, which
are composed of double layered structures, its
composition formula of [M2+1-X M3+X (OH)2࡮A
nX/n࡮nH2O]. These materials are expected to be
useful due to wide potentials as filers, catalysis, drugdelivery system (DDS) and adsorption materials.
“MgAl-talcie basically composed of Mg2+ and Al3+,
which is called hydrotalcite, is most popular among
LDH materials .
In order to improve various properties of LDH,
especially Mg-Al LDH (hydrotalcite), recently, thedevelopments of novel LDH materials modified withthe substitution of Mg2+ and Al3+ sites with anotherM2+ and M3+, are attempted for expanding the new
application fields. We also have developed the new
types of LDH materials modified with the addition of
Zn2+, Ni2+, Fe3+ and other cations of M2+ and M3+. In
this study, we successfully synthesized the novel
LDH materials composed of Cu2+-Al3+ basedhydrotalcite by the soft solution process(co-precipitation processing). Therefore, we, here,
attempted to investigate the Al local structure of
Cu2+-Al3+ based hydrotalcite with the X-ray
adsorption
near
edge
structure(XANES)
measurements to clarify of the fine structure.
The stock solutions of 0.2 mol/dm3 CuCl2 aqueous
solution. 0.1 mol/dm3 AlCl3 aqueous solution, and0.05 mol/dm3 NaHCO3 were prepared. The pH of
mixing solution was controlled with the NaHCO3
solution and aqueous NaOH solution. The products
were synthesized by the co-precipitation method as asoft chemical process. CuAl-talcite with M2+/M3+=2~4 were synthesized by adding mixing solution
into 0.05 mol/dm3 NaHCO3 at RT. 1 mol/dm3 NaOH
was simultaneously added into the aqueous solution
under pH 10. The precipitation was aged at room
temperature for a few hours for Cu2+-Al3+ basedhydrotalcite. The products were filtered with Buchner
funnel. They were separated and sufficiently washedby deionized water and dried at 50 ͠ for 24 hours.
The products of Cu2+-Al3+ based hydrotalcite werecharacterized with various methods. The local
structures around Al for Cu2+-Al3+ based hydrotalcite
were evaluated by measuring X-ray adsorption nearedge structure at BL2A in UVSOR with KTP.

For the comparison, Mg-Al hydrotalcite samples
prepared by the co-precipitation method weresimilarly evaluated. Cu2+-Al3+ based hydrotalcite
samples synthesized by the co-precipitation method
was powdered. the results of powder X-ray
diffraction analysis showed that all samples for
Cu2+-Al3+ based hydrotalcite with different Cu2+/Al3+
ratio of 2 to 4 were identified to be a layered doublehydroxide. Figure 1 shows the results of XANES of
Al-K edge for Cu2+-Al3+ based hydrotalcite and thecommercial Mg-Al hydrotalcite and Mg-Al
hydrotalcite samples prepared by the co-precipitation
method. The spectrum of Cu2+-Al3+ basedhydrotalcite products were similar to one ofcommercial Mg-Al hydrotalcites. The difference
between the peak positions of P-K edge of Cu2+-Al3+
based hydrotalcite with different Cu2+/Al3+ ratio of 2
to 4 was confirmed by XANES This differences wascaused by the content of Cu in hydrotalcite structures.

Normalized Intensity / arb. units

Al-K edge

Cu-Al LDH
(Cu/Al = 4)

Cu-Al LDH
(Cu/Al = 3)

Mg-Al LHD
(Commercial)

1550

1560
1570
Energy / eV

1580

Fig. 1. Results of XANES of Al-K edge of the
Cu2+-Al3+ based hydrotalcite.

111

BL2A

Mo LIII-Edge XANES Study of Silylated Mo/H-MFI Catalysts for Methane
Dehydroaromatization
H. Aritani1, F. Ohya1, K. Kuramochi1, T. Sugawara1, N. Naijo1, K. Takanashi1 and
A. Nakahira2
1

Department of Life Science & Green Chemistry, Saitama Institute of Technology, Fukaya 369-0293, Japan
2
Graduate School of Engineering, Osaka Prefecture University, Sakai 599-8531, JapanAs a typical GTL (Gas-To-Liquid) catalyst in a
direct conversion process of LNG, MoOx-modified
H-MFI zeolite (Mo/H-MFI) catalyst shows highactivity for methane dehydroaromatization to produce
benzene and naphthalene in the absence of oxygen.
The important problem on the Mo/H-MFI catalysts is
a rapid deactivation by coking and/or carbon
contamination on the catalyst surface, and thus,
highly active GTL catalysts with durable activity
have been called for. Many workers have been
revealed that reduction of Mo species is brought
about in contact with methane in initial step. Reduced
Mo ions tend to react methane to form carbide and/or
oxycarbide species in next step [1]. The carbide
species is active for methane dehydroaromatization,
however, it is difficult to avoid the deactivation due tocarbon deposition on the H-MFI surface. The carbon
deposition proceeds mainly on the acid sites of
H-MFI extraframework. On H-MFI interpores,
appropriate acid sites shows durable activity with low
deactivation rate. In the present study, silylated
Mo/H-MFI catalysts were prepared in order to cover
the acid sites of H-MFI extraframework. By silylation
using organosilane regents, coking sites inextraframework can be covered selectively. In fact,
the silylated Mo/H-MFI showed durable activity inthe reaction with CH4-H2 at 973 K. To characterize
the bare- and sylilated Mo-H-MFI catalysts, Mo
LIII-edge XANES studies were applied. Red-ox statesof the active Mo species on H-MFI with/without
silylation were evaluated before/after the methane
dehydroaromatization.
Catalysts were prepared by impregnation of H-MFI
support with MoO2(acac)2íCHCl3 solution, andfollowed by drying overnight and calcination at 773
K for 3 h. The amount of MoO3-loading is 5.0 wt% inthis study. H-MFI supports with Si/Al2 = 20-72 weresynthesized hydrothermally at 413 K for a week, andfollowed by ion-exchanging with NH4Cl andcalcination at 873 K. Silylation was carried out byimpregnating triethoxyphenylsilane-heptane solution(Mo/Si=5) with Mo/H-MFI, and dried at 353 K for
onernight. The catalytic activity of methane
dehydroaromatization was evaluated by means of
fixed bed flow reaction, as described in a separate
paper [2]. Mo LIII-edge XANES spectra were
measured in BL2A of UVSOR-IMS in a
total-electron yield mode using InSb double-crystal
monochromator. Photon energy was calibrated byusing Mo metal-foil at Mo LIII-edge, and normalized
XANES spectra and their second derivatives arepresented. REX-2000 (Rigaku) software was used bynormalization of each XANES spectrum.
Figure 1 shows the XANES spectra of bare(shown in top) and silylated (in bottom) Mo/H-MFI
catalysts (pretreated with CO(2%)-He) reacted with
CH4(20%)-He at various temperatures. Reduction of
Mo species reflects the shift of lower edge energy.
After reaction at 973-1023 K, Mo species on silylated
Mo/H-MFI are reduced easier than that on bare one.
At 1073 K, Mo2+ species due to Mo2C are formed
definitely in both catalysts. The silylated catalyst
showed a maximum activity at 973 K, and thus, thesilylation relates to the enhancement of Mo species toform partially carbonized (Mo-oxycarbonates, inmajor) ones. At the same time, carbon deposition was
inhibited. These results suggest the inhibition ofcarbon coking and formation of active Mo species atlower temperature at 973K by silylation (Mo/Si=5).

Fig. 1. Mo LIII-edge XANES of Mo/H-MFI (Si/Al2=40)
catalysts before and after silylation with phenylsilane(Mo/Si=5) reacted at various temperatures.

[1] H. Aritani, S. Shinohara, S. Koyama, K. Otsuki,
T. Kubo and A. Nakahira, Chem. Lett. 35 (2006) 416.
[2] H. Aritani, H. Shibasaki, H. Orihara and
A. Nakahira, J. Environm. Sci. 21 (2009) 736.

112

BL2A

Mg-K XANES Study of Mg-Doped Hydroxyapatite
H. Murata1, K. Shitara1, F. Oba1 and I. Tanaka1, 2
1

Department of Materials Science and Engineering, Kyoto University, Kyoto 606-8501, Japan
Nanostructures Research Laboratory, Japan Fine Ceramics Center, Nagoya 456-8587, Japan

2

Hydroxyapatite (HAp) is a main inorganic
component of human bones and teeth. HAp in humanbodies contains various kinds of defects such as Ca
vacancies, excess H and impurities. They affectbiological properties of HAp [1] and thus it isimportant to understand their electronic and atomistic
structures.
Mg is one of the essential elements for humanbodies and a major impurity in human bones. It isconsidered to have an important role for bone
re-modeling [1]. In this study, we investigated the
local environments of Mg in HAp by Mg-K X-ray
absorption near edge structure (XANES).
Undoped and Mg-doped HAp samples weresynthesized using the solution-precipitation method.
Starting materials were 0.1 mol/L solutions ofCa(NO3)2·4H2O, Mg(NO3)2·6H2O and (NH4)2HPO4.
Precipitations were matured at 353 K for 5 hours in
Ar atmosphere. Nominal Mg/(Ca+Mg) ratio was set
to 0, 2, 5 and 20 at%. Mg-K XANES was measured at
BL2A in UVSOR. The incident X-ray beam was
monochromatized by beryl double crystals. XANES
signals were collected with GaAsP photo-diodes(HAMAMATSU G-1127-02). Pellets of samples were
mounted on carbon tapes. For a detailed analysis of
Mg-K XANES, theoretical spectra were obtained by
the first-principles supercell method with a core-hole
effect [2].
Mg-doped HAp samples were characterized by
XRD. No secondary phases were observed. The XRD
patterns showed peak shifts and broadening by Mg
doping. This indicated that Mg was incorporated into
the HAp crystals.
The Mg-K XANES spectra of these samples were
measured at BL2A. Usually, XANES signals in a soft
X-ray region are collected by the total electron yield(TEY) method. However, our samples had a tendency
to be charged up and we could not obtain high-quality
spectra by the TEY method. Therefore, thefluorescent yield (FY) method with GaAsP
photodiodes was employed. Figure 1 shows Mg-K
XANES of the Mg-doped HAp samples measured by
the FY method. The FY method gives good spectra
for Mg with small concentrations.
It is found that the Mg-K XANES for 2, 5 and 20
at%Mg have the same features. In order to elucidate
the local environments of Mg in HAp via theinterpretation of Mg-K XANES, theoretical spectra
were obtained. HAp has two non-equivalent Ca site,i.e., the Ca-1 (columnar) site and Ca-2 (triangular)site. We considered models in which one Mg was
substituted for the Ca-1 or Ca-2 sites. However,neither model reproduced the experimental Mg-K

XANES spectra. This indicated that Mg in HAp is not
simply located at these sites. Our previous study
suggested that the XANES of dopants in HAp isaffected by association with the Ca vacancy andexcess H [3]. Mg in HAp is likely to be associatedwith other defects as well. A detailed analysis of thedefective structures is in progress.

Intensity / arb. units

Expt.

HAp:Mg20 at%

5 at%

2 at%

Mg at Ca-2 site
Calc.

Mg at Ca-1 site
Calc.

1300

1310 1320 1330 1340
Transition Energy / eV

Fig. 1. Experimental and calculated Mg-K XANES of
Mg-doped Hap.
[1] A. Boanini et al., Acra Biomater. 6 (2009) 1882.
[2] I. Tanaka et al., J. Am. Ceram. Soc. 88 (2005)2013.
[3] H. Murata et al., J. Phys.: Condens. Matter 22(2010) 384213.

113

BL3U

Hybridized Electronic States near EF of Potassium-Doped Picene Probed bySoft X-Ray Emission Spectroscopy
H. Yamane and N. Kosugi
Department of Photo-Molecular Science, Institute for Molecular Science, Okazaki 444-8585, Japan
School of Physical Sciences, The Graduate University for Advanced Studies, Okazaki 444-8585, Japan
Recently, a new organic superconductor has been
discovered for a Potassium (K)-doped picene [see,Fig. 1(a)], with high transition temperatures (< 18 K),
which depend on the K concentration [1]. In order toelucidate the nature of the superconductivity in the Kdoped picene, detailed information on the electronic
structure is required. It has been suggested fromtheoretical calculations that the dopant (K) orbitals
are higher in energy and weakly hybridized with themolecular states close to the Fermi energy (EF) [2];however, experimental proofs for the formation of
hybridized electronic states of the K-doped picene
have not yet been reported.
In this work, in order to understand the local
electronic structure of the K-doped picene, soft X-ray
spectroscopies were applied at BL3U. The K-doped
picene was prepared by the co-evaporation of picene(99.9% purity, NARD inst. Ltd) and K (SAES getters)on SiO2/Si(111) at the fixed deposition rates with thepost-annealing at 370 K. The K concentration and the
crystallinity of the KxPicene (x = 0, 3) were confirmed
by CuKĮ X-ray diffraction (XRD) and fluorescenceyield X-ray absorption spectroscopy (FY-XAS).
Figure 1(a) shows the XRD scans of the KxPicene(x = 0, 3), which agree well with the earlier study [1].
For the K-doped picene, the K atoms are intercalated
in the stacked picene molecules.
Figure 1(b) shows FY-XAS spectra of the KxPicene(x = 0, 3) at the normal incidence. The strong XAS
peaks, originating from the C 1s ĺ LUMO (ʌ*) and
LUMO+1 (ʌ*) transitions [3,4], appear at the incident
photon energy (hȞin) of 284–285 eV. These peaks are
shifted to the lower hȞin side and broadened upon theK doping due to the electron transfer from K 4s to
LUMO and LUMO+1 of picene.
Figure 2 shows the selected soft X-ray emission(XES) spectra of the KxPicene (x = 0, 3) at various
hȞin. The elastic peak shows linear dispersion to thehigher emission energy (hȞout) side with increasing
hȞin. At hȞin = 310.0 eV, the fluorescence feature in
XES (hȞout < 290 eV) for the K-doped picene shows adoping-induced peak near EF at hȞout = 288 eV very
weakly, as indicated by the downed arrow, which is
not observable for the undoped picene. From the hȞindependence of the XES spectra, we found that thedoping-induced peak near EF is pronounced after theK 2p excitation, hȞin  295 eV [see, Fig. 1(b) and 2].
These FY-XAS and XES results indicate that themolecular states of picene are hybridized with K 4s
near EF upon the K doping.

0

10

20

Intensity / cps(a)2T / °

40

CuKD XRD: K xPicene/SiO2/Si(111)×8undoped(x = 0)×8

0Intensity / normalized

30

K-doped(x = 3)
1

2

Wave-vector Transfer / Å-1(b)

3

FY-XAS (normal incidence)
KxPicene/SiO2/Si(111)
undoped(x = 0)K-doped(x = 3)

280

290

300

310

Incident Photon Energy (hvin) / eV

Fig. 1. (a) CuKĮ XRD and (b) FY-XAS spectra of
KxPicene (x = 0 and 3) films on SiO2/Si(111).
hvin-dep. XES
KxPicene/SiO2/Si(111)Intensity / normalized: undoped (x = 0)
: K-doped (x = 3)

hvin =
310.0 eV
310.0 eV295.8 eV292.2 eV

240

260

280

300

320

Emission Energy (hvout) / eV

Fig. 2. XES spectra of KxPicene (x = 0 and 3) films asa function of incident photon energy (hȞin).
[1] R. Mitsuhashi et al., Nature 464 (2010) 76.
[2] G. Giovannetti and M. Capone, Phys. Rev. B 83
(2011) 134508.
[3] H. Ågren et al., Chem. Phys. 196 (1995) 47.
[4] M. Obst et al., Environ. Sci. Tech. 45 (2011) 7314.

114

BL3B

Optical Spectroscopy of ZnO:Al Thin Films
N. Kashiwagura, T. Hashida and T. Takeya
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 of solar
cell and liquid crystal display is ITO (Indium Tin
Oxides) films. However, ITO films include so sparse
Indium that the substitutive materials are expected.
ZnO thin films are one of the hopeful materials 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 induced electric power, thesubstrate temperature were 12 W, 200 ͠, respectively,
for all samples in common. The sputtering gas (Ar)pressure was 0.125, 0.2 and 0.3 Torr. for sample 1,
sample 3 and sample 5, respectively.
Optical absorption, luminescence and excitation
spectra were measured using the BL3B beam in thetemperature range of 8-300 K.
Figure 1 shows the optical absorption spectra for
ZnO thin films. The sharp edge at 150 nm (8.3 eV)
for sample 1 is due to the absorption of fused silica
substrate. The several broad bands below 360 nm (3.3eV) are due to the band-to-band transition of ZnO. All
the samples are efficiently transparent as TCO thinfilm above 380 nm (visible light region).
Figure 2 shows the luminescence spectra withvarious excitation wavelengths in sample 1 at 300 K.
The band-to-band excitation below 360 nm produces
several broad luminescence bands at 405 nm, at 425
nm and at 440-490 nm. In these spectra, theadditional peaks around 540 and 590 nm are theinfluence from outer ceiling lights, which could not
be perfectly eliminated. As the excitation wavelength
increases (the excitation energy decreases), theintensity of luminescence bands decreases as a whole.
Figure 3 shows the luminescence spectra insample 5 at 300 K. The spectra exhibit several broad
luminescence bands at 405 nm, at 425 nm and at
440-490 nm. The spectra are very similar to that ofsample 1. With increasing the excitation wavelength,
the luminescence intensities decrease as a whole.
The band around 425 nm is estimated to be
intrinsic
luminescence (for example, self-trapped
exciton). The broader band around 470 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 Al doped ZnO thin films were
investigated. It seems to be important to investigate
the exact Al concentration in thin films which are
made by sputtering with the target including 2% Al
nominally.

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

Fig. 2. Luminescence spectra in sample 1 at 300 K.

Fig. 3. Luminescence spectra in sample 5 at 300 K.

115

BL3B

Optical Spectroscopy of Cd1-xMnxTe Crystals
M. Yamaga, Y. Ogoshi, A. Kanetake, S. Takano, T. Nohisa, S. Tsuda and H. Tsuzuki
Department of Material and Design Engineering, Gifu University, Gifu 501-1193 Japan
Cd1-xMnxTe
crystalshavemagneto-opticalproperties induced by Mn2+ ions. They are very
useful as applications to an optical isolator.
In a previous activity report [1], optical properties
of Cd1-xMnxTe crystals (x=0.5) were reported. In
this report, the composition x is extended from 0.1 to
0.9 in the whole range of the composition.
Cd1-xMnxTe crystals were grown by the vertical
Bridgeman method. The compositions, x, of thecrystals are 0.1, 0.25, 0.3, 0.5, 0.75, 0.9 being equal
to the melting compositions [2].
Figure 1 shows the absorption spectra observed in
Cd1-xMnxTe (x=0.1, 0.25, 0.3, 0.5, 0.75, 0.9) at 300 K.
The band gap energy related to the absorption edge ofthe crystal increases with an increase of x up to 0.5
and slightly decreases above 0.6.
This result
suggests that the phase transition of the crystal
Cd1-xMnxTe occurs around x=0.6 from zinc blende to
hexagonal phases.
The sample temperature isdecreased from 300 K, the band edge is shifted to theshort wavelength, resulting in the large band gap.
Figure 2 shows temperature dependence of theluminescence spectra for Cd0.9Mn0.1Te. The peak of
the luminescence is around 720 nm and just below theband edge energy observed at 20 K.
Theluminescence peak is shifted to the longer wavelengthand the intensity markedly decreases with an increaseof temperature. The intensity is negligibly small
above 140 K. In the same way, Fig. 3 shows theluminescence spectra of Cd0.25Mn0.75Te at various
temperatures. The luminescence peak is shifted tothe short wavelength compared with that for
Cd0.9Mn0.1Te. However, the line width for
Cd0.25Mn0.75Te is the largest for all crystals with thecompositions of x=0.1, 0.25, 0.3, 0.5, 0.75, 0.9.
We consider the origin of the luminescence from
Cd1-xMnxTe (x=0.1, 0.25, 0.3, 0.5, 0.75, 0.9) crystals.
The luminescence may be assigned to weakly
localized excitons. The line width has the maximum
for x=0.75. This result is consistent with the ESR
results of Mn2+ in Cd1-xMnxTe (x=0.1, 0.25, 0.3, 0.5,
0.75, 0.9) that the ESR line width has also the
maximum at x=0.75. The line broadening is due tolocal magnetic field created by Mn2+ cluster in the
Cd1-xMnxTe crystals.

Cd1-xMnxTe

X = 0.1
X = 0.25
X = 0.3
X = 0.5
X = 0.75
X = 0.9-1

Absorption coefficient [cm ]

10

6

2

-2
200

400

600

800

1000

1200

Wavelength [nm]Fig. 1. Absorption spectra of Cd1-xMnxTe.
Cd0.9Mn0.1Te
800
14K
20K
40K
60K
80K
100K
120K

Intensity [a.u.]

600

400

200

0
550

600

650

700

750

Wavelength [nm]Fig. 2. Luminescence spectra in Cd0.9Mn0.1Te.
Cd0.25Mn0.75Te

Intensity[a.u.]

3000[1] M. Yamaga, UVSOR Activity Report 38 (2010)
117.
[2] M. Prakasam, O. Viraphong, L. Teule-Gay, R.
Decourt, P. Veber, E. G. Villora and K. Shimamura, J.
Crystal Growth 318 (2011) 533.

14K
20K
40K
60K
80K
100K
120K
140K

2000

1000

0
550

600

650

700

750

Wavelength [nm]Fig. 3. Luminescence spectra in Cd0.25Mn0.75Te.

116

BL4B

Annealing Effect on Local Structure of Diamond-Like Carbon andTungsten-containing Diamond-Like Carbon Films
A. Fujimoto and K. Kanda
Laboratory of Advanced Science and Technology for Industry, University of Hyogo Kamigori, Ako, Hyogo
678-1205, Japan
Diamond-like carbon (DLC) film is expected to use
as various industrial fields, because of its some
excellent properties such as hard, low friction
coefficient, low surface energy and so on. Recently,
incorporation of Tungsten (W) into DLC films was
proposed because it was known that thisincorporation improved the film’s thermal durability.
In the present study, the effect of annealing on DLC
and tungsten containing DLC (W-DLC) films invacuum was investigated from the measurement of
the annealing temperature dependence on the carbon
K-edge near-edge x-ray absorption fine structure(NEXAFS) spectra of the commercial DLC and
W-DLC films fabricated by various methods forindustrial use.
Sample commercial DLC and W-DLC films wereprepared on the Si substrate with |200 nm film
thickness using several deposition methods: 1) DLC
and W-DLC films by unbalanced magnetron
sputtering (UBMS) (KOBELCO Co., Ltd.), 2) DLC
film fabricated by ion plating (IP) (NANOTEC Co.,
Ltd.), 3) DLC film fabricated by plasma-enhanced
CVD (PE-CVD) (Nippon-ITF Inc.). These samples
were annealed in a furnace (Thermo RIKO Co. Ltd;GFA430) under vacuum conditions (<10-4 Pa). The
annealing temperature was 673, 773, or 873 K, andthe annealing time was 32 h. C-K NEXAFS spectra
were measured at the BL4B of UVSOR and BL09A of
NewSUBARU. The NEXAFS spectra were measured
in the energy range 275-320 eV with 0.5 eV FWHM
resolution in the total electron yield mode.
In the C-K NEXAFS spectra of DLC films, sharp
peak, which is assignable to the transitisons from theC 1s level to the unoccupied S* levels of the sp2(C=C) and/or sp (C҂C) sites, is observed at 285.4 eV.
The intensity of this peak increased with theannealing temperature, which indicated the increaseof sp2 ratio in the film. The procedure used toestimate the sp2/(sp2+sp3) ratio is described elsewhere[1]. Figure 1 shows the dependence of thesp2/(sp2+sp3) ratios of the films on the annealing
temperature. The sp2/(sp2+sp3) ratios of thecommercial W-DLC and DLC films as depositedwere estimated to be 0.3-0.4. The sp2/(sp2+sp3) ratiosof all the W-DLC and DLC films increased with theannealing temperature indicating that a graphite
structure had formed[2]. Therefore, the onset ofincrease in sp2/(sp2+sp3) ratio was positioned at
higher temperature indicated that the film had strong
thermal durability. In short, thermal durability ofthese films was stronger in the order of IP DLC,117

UBMS W-DLC, UBMS DLC, and PE-CVD DLC
films
As a result, the present study confirmed following
facts from local structure level by the measurement of
NEXAFS spectra: 1) Thermal stability of DLC films
was varied on the deposition method of film. 2)
Incorporation of W improves the film’s thermal
durability.

Fig. 1. Dependence on annealing temperature ofsp2/(sp2+sp3) ratio:. a): UBMS W-DLC and : UBMSDLC films;. b): PE-CVD DLC and : IP DLC films.
Points at 273 K indicate sp2/(sp2+sp3) ratio asdeposited.
[1] 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.
[2] K. Kanda, J. Igaki, N. Yamada, R. Kometani and
S. Matsui, Diamond Relat. Mater. 18 (2009) 490.

BL4B

Electronic Structural Changes of LiNi0.5Mn1.5O4 during Lithium-Ion
Intercalation/Deintercalation
T. Okumura, M. Shikano and H. Kobayashi
Research Institute for Ubiquitous Energy, National Institute of Advanced Industrial Science and Technology(AIST), Midorigaoka 1-8-31, Ikeda, Osaka 563-8577, Japanand the only nickel ion affect the redox reaction at x >1. Thus, the electronic structural change of oxide ion
is also crucial for considering redox reaction at
charging/discharging batteries, and the contribution of
the oxide ion on redox reaction differs in various
redox cation species.

Lithium-ion batteries (LIB) with high energy
density, safety, and durability are required for electricvehicles (EV), hybrid electric vehicles (HEV), andplug-in hybrid electric vehicles (PHEV) as their powersources. Lithium nickel manganese spinel oxide
LiNi0.5Mn1.5O4 is one of the candidates for the positiveelectrode materials for LIBs as an alternative to theconventional LiCoO2 cathode because of its low costand low toxicity[1]. Moreover, its specific capacity
above 4.5 V is suitable for designing thenext-generation LIB with the high energy density. In
order to get the guide to improve the superior
electrochemical properties, understanding the detailed
electronic structural changes in LixNi0.5Mn1.5O4 duringlithium-ion intercalation/deintercalation process isessential as reported in other electrode materials[2].
The LixNi0.5Mn1.5O4 samples with various lithium
concentrations have been picked up from the batteries
after intercalation or deintercalation as shown in Fig.
1. Ni L3-edge, Mn L3-edge and O K-edge XANES
spectra were measured with total electron yield mode.
The Ni and Mn L3-edge spectra at various lithiumcontents (shown in Fig. 2 (a) and (b), respectively)
correspond to the electronic transitions from the 2p3/2
states to an unoccupied 3d state. At x < 1, theabsorption edge peak at the right shoulder of Ni
L3-edge peaks clearly appeared with the decrement of
the lithium-ion contents although the significant
change could not be observed in Mn L3-edge peaks.
This result indicates that the charge variation duringlithium-ion deintercalation was compensated by theredox reaction of the Ni2+/Ni4+ couple[3]. On the other
hand, at x > 1, the absorption edge energy of Mn
L3-edge peaks shifted towards lower energy levels
with an increase of lithium content while Ni L3-edge
peaks were mostly maintained. This result indicatesthat the charge variation during lithium-ion
intercalation was compensated by the redox reactionof the Mn3+/Mn4+ couple [3]. The large voltage gap
between x < 1 and x > 1 in Fig. 1 results from thedifference of the redox-ion species.
The O K-edge spectra were shown in Fig. 2 (c).
The absorption edge peaks appeared around 528-535
eV particularly indicate the transition from O 1s to
hybridized orbital of O 2p with transition-metal 3d[2].
At x > 1, the shape of the peaks changed with theincrease of the lithium content as well as the shape ofMn L3-edge peaks changes as written above, although
it maintained at x > 1.
These results indicated that both manganese andoxide ions contributed on the redox reaction at x < 1,

Fig. 1.
Lithium-ion intercalation/deintercalation
behavior in LiNi0.5Mn1.5O4 spinel oxide. Circles
indicate lithium content of samples measured
XANES spectra.

Fig. 2. XANES spectra of Ni L3-edges (a), Mn
L3-edges (b), and O K-edges (c) in LixNi0.5Mn1.5O4.
[1] Q. Zhong et al., J. Electrochem. Soc. 144 (1997)
205.
[2] for example, T. Okumura et al., Dalton Trans. 40(2011) 9752.
[3] Y. Terada et al., J. Solid State Chem. 156 (2001)
286.

118

BL5U

Valence-Band Electronic Structure of N-Implanted TiO2
K. Soda1, T. Owada2, K. Morita2, S. Harada1, E. Kuda1 and T. Yoshida3,1
1

2

Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan
Department of Physical Science and Engineering, Nagoya University, Nagoya 464-8603, Japan
3
EcoTopia Science Institute, Nagoya University, Nagoya 464-8603, Japanbonding valence band at EB ~ 7 eV are enhanced dueto the resonance photoemission of the Ti3d electrons[3], while the states at EB = 3 ~ 6 eV show no
prominent increase, because they are mainly derived
from O 2p and N 2p states. These are consistent witha band calculation for TiO2 and N-substituted
TiO2-xNx [1].
Similar results have already reported by Batzill etal., who concluded that the impurity states just above
the valence band without the band-gap narrowing
were formed by N doping [3]. However, the present
results together with the NEXAFS ones suggest thatthe implanted N occupies the O site and forms acommon valence band extending into the band gap inthe surface region; the top of the valence band islocated at EB ~ 2.5 eV, which is comparable with
reported threshold of the photoactivity in thedecomposition of methylene blue for the N-doped
TiO2 [1]. The Ti 3d defect states just below theconduction band also induce visible-light absorption
but are not responsible for the photoactivity.

TiO2 is well-known as a photocatalyst with a band
gap of ~3 eV [1]. Recently, Yoshida et al. have foundthat TiO2 gains the photoactivity for visible light by
N+ ion implantation and that N occupying at the Osite is responsible for the visible-light activity by the
N K-edge near-edge X-ray absorption fine structure(NEXAFS) measurement [2]. In this study, we haveinvestigated valence electronic structure of such N+
implanted TiO2 to clarify the origin of thevisible-light photoactivity.
Photoelectron measurement were carried out atroom temperature in the angle-integrated mode withtotal energy resolution of ~0.1 eV. Specimens wererutile TiO2 (100) plates, which were first in situ
cleaned by 5-keV Ar+ ion sputtering and thenimplanted with 5-keV N+ or N2+ ions. The origin of
the binding energy EB was set to the Fermi energy EF
of a gold plate on a sample holder.
Figure 1 shows typical photoelectron spectra
recorded before and after the N+ implantation withthe excitation photon energy hQ. In the figure, thebackground due to the secondary electrons are
subtracted from the measured spectra and theirintensity is normalized with the integrated intensityfrom EF up to EB = 12 eV. In the upper panel, spectra
after the Ar+ sputtering and subsequent N+
implantation of ~8 x 1020 ions m-2 are compared byshifting the spectrum after the implantation towards
the high binding energy side so that the valence band
bottom (EB = 8 ~ 9 eV) may coincide with each other.
The N K-edge NEXAFS measurement for this sample
confirms the formation of visible-light active N
species. After the sputtering, a broad hump appears in
the upper part of the gap around EF. This is due to the
reduction of Ti or defect formation, induced by the
sputtering, which also leads to the increase inelectrons in the Ti 3d conduction band. Observed
apparent shift of the Fermi edge towards the highbinding energy side after the implantation implies thedecrease in defects and conduction electrons, which
may be caused by the recovery of the chemical bond,i.e. new Ti-N bond formation. Spectral increase dueto the N 2p states derived from the Ti-N bond appears
in the upper part of the O 2p valence band and in thelower region of the gap (EB = 3 ~ 6 eV). The
assignment of these new bands after the implantationis also confirmed by the Ti 3p-3d resonance behavioraround hQ = 50 eV, which is shown in the lower panelof the figure for the spectra after the implantation. At
hQ = 50 eV, the hump around 0.5 eV, Ti 3d
conduction band or defect states, and the O 2p-Ti 3d

after implantation
after Ar sputtering

1

INTENSITY / arb. units

hQ/eV = 40

0
hQ/eV
40
50

0
10

Fig. 1.

50BINDING ENERGY / eV

Valence-band spectra of N-implanted TiO2.

[1] R. Asahi et al., Science 293 (2001) 269.
[2] T. Yoshida et al., Mater. Sci. Forum 561-565(2007) 567.
[3] M. Batzill et al., Phys. Rev. Lett. 96 (2006)
026103.

119

BL5U

Effects of Periodicity Imperfection on Mini-Bands in Semiconductor
Superlattice
K. Hashimoto1, D. Shimura1, M. Kuwahara2, S. Harada1, T. Ujihara1, M. Matsunami3 andS. Kimura3
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 highly-optimized intermediate bandsolar cell, of which a high conversion efficiency canreach 74.6% theoretically, is proposed [1]. The
concept is to utilize intermediate bands in the carrier
excitation in order that the band structure is adjusted
to the solar light spectrum. As intermediate bands, themini-band due to a periodic quantum structure(superlattice structure) is used. But there is a problem
that it is presently impossible to produce quantum
dots of equal size and perfect periodicity.
In this study, the effects of periodicity imperfection
on the mini-band structure, especially holemini-bands, were investigated. We examined asuperlattice structure which consists of multiple
quantum wells because the preparation method has
been established.
Figure 1 shows the detailed structures of eachsample: A GaAs bulk sample, a periodic superlattice
and a disordered superlattice of which second welllayer was made thinner. Amorphous arsenic layers
were deposited on the sample surfaces to suppress
oxidation, and they were removed just before themeasurements by heat-treatment. On the evaluation of
the mini-band structure, the synchrotron-radiation
photoemission measurements were performed at
BL5U of the UVSOR facility.
Figure 2 shows the electron density mappings in avalence band around * point. The red dots and green
dots indicate the band structure determined from thepeak values of the second order differential. They area heavy hole band and a light hole band, respectively.
In the case of bulk GaAs, the band structures arequadric
curves.
Inaddition,
theyaredoubly-degenerated at kz = 0. On the other hand, theband structures of the periodic superlattice and thedisordered superlattice are not degenerated at kz = 0.
Moreover, they are also folded symmetrically on thewhite arrows, which indicate the interval of theBrillouin zone due to the periodic structure ofsuperlattice. Therefore, their band structures aremini-bands.
The mini-band structures are characterized by thewidth of band splitting between the heavy holemini-band and the light hole mini-band. Figure 3
shows the width of band splitting as a function of thewavenumber. The width of the disordered superlattice
is wider than that of of the periodic superlattice. This
is attributed to the following discussion: Due to thethinner well, the heavy hole band and the light holeband of the disordered superlattice were lower thanthose of the periodic superlattice. The band shift
amount of the light hole band is larger than that of theheavy hole band by a difference in the effective mass.
As a result, the width of the disordered superlattice iswider.

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

Fig. 2. Electron density mappings of (a) GaAs, (b)periodic superlattice and (c) disordered superlattice.

Fig. 3. Width of band splitting between heavy holemini-band and light hole mini-band.
[1] T. Nozawa and Y. Arakawa, Appl. Phys. Lett. 98(2012) 171108.

120

BL5B

Photodegradation in Amorphous Chalcogenide Semiconductor Films by
Vacuum Ultraviolet Absorption Spectroscopy
K. Hayashi
Department of Electrical and Electronic Engineering, Gifu University, Gifu 501-1193, Japan
It is well known that amorphous chalcogenide
semiconductor materials show a variety ofphotoinduced effects [1-4]. The prominent
photoinduced phenomenon in these materials is atime-dependent decrease in photocurrent during andafter irradiation of bandgap (BG) light [3,4]. Thisphenomenon is usually called photodegradation andis explained in terms of light-induced metastabledefects (LIMD). In the device application of thesematerials, LIMD creation is a serious problem.
Although a number of models have been proposed for
LIMD creation [4], details of the mechanism
underlying LIMD creation in these materials are still
not clear. Understanding the physical mechanism
underlying metastability is one of the importantfundamental problems related to these materials.
Recently, I found a new phenomenon thatphotocurrent by irradiation of BG light increased inamorphous chalcogenide films. This phenomenon is
not explained in terms of LIMD creation. These
photoinduced 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 new phenomenon
was an irreversible change. 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. However, several previous studies ofphotoinduced phenomena have reported that there isno direct correlation among photodegradation,
photodarkening and photoinduced volume change.
Our recent studies have focused on thephotoinduced energy structure changes in the vacuum
ultraviolet (VUV) region by irradiation of BG light.
Few studies have examined whether the photoinducedchange of the chemical bonds contributes tophotodegradation and photodarkening. To understandthe dynamics and the correlation of photoinduced
structure changes in the VUV region and thephotoinduced change of photocurrent in amorphouschalcogenide semiconductors, real time in-situ
measurements are required. In previous report, thosedynamics and correlations are examined bymeasuring the total photoelectron yield (TPEY) thatcan measure the optical property in the VUV region.
However, the beam position, brightness, and theenergy stability were insufficient in the observation of
the dynamics of TPEY by irradiation of BG light in
UVSOR. In this report, I will discuss the stability oflight source.

Samples used for the measurement of thetransmission spectrum were amorphous chalcogenide(a-As2S3 and a-As2Se3) semiconductor thin films
prepared onto aluminum thin films. Typical thicknessof the sample and the aluminum films were around
200nm and 100nm respectively. The aluminum film
was also used in order to eliminate the higher order
light from the monochromator in the VUV region.
The measurements were performed at room
temperature. A pinhole of 1.5mm in a diameter wasinserted between the monochromator and sample toremove stray light. The intensity of the VUV light
was monitored by measuring the TPEY of a gold
mesh.
Figure 1 shows the spectra of the light source under
the same condition monitored by measuring the
TPEY of a gold mesh. As shown in the figure, thespectrum reproducibility in the top-up operation is not
good. We are examining the method to increase theaccuracy of measurement at present.

Fig.1. Spectra of the light source under the same
condition monitored by measuring the TPEY of agold mesh.
[1] K. Tanaka, Rev. Solid State Sci. 4 (1990) 641.
[2] S. R. Elliot, J. Non-Cryst. Sol. 81 (1986) 71.
[3] D. L.Staebler and C. R. Wronski, Appl. Phys. Lett.
31 (1977) 292.
[4] K. Shimakawa, A. Kolobov and S. R. Elliott,
Adv. Phys. 44 (1995) 475.

121

BL6U

Extended Electronic States in Crystalline Guanine-Cytosine
Watson-Crick Multilayers
R. Friedlein1, Y. Wang1, H. Yamane2 and N. Kosugi2
1

School of Materials Science, Japan Advanced Institute of Science and Technology, Ishikawa 923-1292, Japan
2
Dept. of Photo-Molecular Science, Institute for Molecular Science, Okazaki 444-8585, Japan

In order to elucidate the role of intermolecularinteractions in the pre-biotic assembly of RNA
precursor oligomers, it is crucial to understand, if andin which way one-dimensional, stacked assemblies of
Watson-Crick (WC) pairs may be formed without thepresence of a backbone. With scanning tunnelingmicroscopy studies limited to monolayer assemblies[1], the nature of the stacking between individual
nucleic acids as well as of WC pairs has so far only
been studied theoretically [2].
In recent home-based experiments we have foundthat the two nucleic acids guanine and cytosine
assemble into an ordered multilayer film based on thehydrogen bonding within WC pairs (Fig. 1). The
stability of WC pairs is high such that a macroscopic
phase separation does not occur. Thin mixed films,
about 2-3 nm thick, have been prepared underultra-high vacuum conditions by co-deposition of
guanine and cytosine molecules, at equal rates withrespect to the amount of molecules, on the surface ofhighly-oriented pyrolytic graphite, HOPG(0001), heldat 25 qC. The samples have then been exposed towater vapor in situ, and transferred in air to theexperimental set-up at BL6U of the UVSOR facility.
Photon-energy
dependent
valenceband
photoelectron spectra obtained in the normal emission
direction are shown in Fig. 1. The onset of the spectra
is related to the highest occupied molecular orbital(HOMO) of guanine. It shows a clear dispersion withthe photon energy which is consistent with theformation of extended electronic states related to the
guanine HOMO similar to those in guanine
multilayer films [3]. This demonstrates aquasi-columnar arrangement of the guanine
molecules oriented in the direction normal to thesurface and an anti-parallel stacking with two
molecules per repeat unit along the stacks.
Non-dispersing states at higher binding energy,
denoted “A” and “B”, are related to cytosine
molecules that have a significantly smaller
intermolecular S-S overlap.
In order to accomplish an anti-parallel stacking in alayered arrangement, subsequent WC pairs must be
flipped with respect to each other. As verified byquantum-chemical calculations [4], in thisconfiguration, the cytosine molecules stack with thosefrom neighboring WC pairs. This is in stark contrast
to the arrangement realized in RNA which indicates
an important role of the backbone sugars inpreventing otherwise preferred anti-parallel stacking
configurations.

Financial support from the SHIBUYA and
MARUBUN Foundations is acknowledged.

Fig. 1. The structure of a cytosine-guanine WC pair(top), and photon-energy-dependent photoelectron
spectra of the mixed film (bottom). The dispersion of
the guanine-derived HOMO is indicated by a solidline.
[1] S. Xu, M. Dong, E. Rauls, R. Otero, T. R. Linderothand F. Besenbacher, Nano Lett. 6 (2006) 1434.
[2] C. Acosta-Silva, V. Branchadell, J. Bertran and A.
Olivia, J. Phys. Chem. B 114 (2010) 10217.
[3] R. Friedlein, Y. Wang, A. Fleurence, F. Bussolotti,
Y. Ogata and Y. Yamada-Takamura, J. Am. Chem.
Soc. 132 (2010) 12808.
[4] Y. Wang, F. Negri, H. Yamane, N. Kosugi and R.
Friedlein, in preparation.

122

BL6U

Band Alignment at CNT/n-Type 4H-SiC Interface Formed by Surface
Decomposition of SiC
T. Maruyama1, S. Sakakibara1, Y. Ishiguro1, H. Yamane2 and N. Kosugi2
1

Department of Materials Science and Engineering, Meijo University, Nagoya 468-8502, Japan
2
Institute for Molecular Science, Okazaki 444-8585 Japan

Carbon nanotube (CNT) growth by surface
decomposition of SiC is a synthesis method for
self-organized CNT films only by heating SiC(000-1)
carbon-face in a high vacuum [1]. By this method,high-density, well-aligned zigzag-type CNTs can beproduced perpendicular to the substrate surfaces. In
addition, CNTs were atomically bonded to SiC
without any interlayers [2]. Previously, we reported
both electrical property and electronic structure for
CNT/n-type 6H-SiC and showed the presence ofSchottky barrier at the interface [3]. In this study, we
performed photoemission spectroscopy (PES)
experiments for CNT/n-type 4H-SiC interface anddetermined the band alignment.
CNT films formed by surface decomposition ofn-type 4H-SiC(000-1) were used as samples. For PES
measurements, we prepared three kinds of samples
with different CNT thicknesses; CNTs (~300 nm) on
4H-SiC, CNTs (several nm) on 4H-SiC and
4H-SiC(000-1) substrate. The PES measurements
were carried out in a high-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 at room temperature (~300K).
Figure 1 (a) shows PES spectrum near valenceband maximum (VBM) of CNTs (~300 nm) formedon 4H-SiC with that of Au. The Fermi edge isobserved in the spectrum of CNTs, confirming thatthey were metallic, as reported for CNT films on
6H-SiC [3]. The band alignment at the CNT/SiC
interface can be evaluated by measuring the binding
energy of C 1s core level at the CNT/4H-SiC
interface, the energy separation between the VBM
and the C 1s core level of 4H-SiC and that between
the Fermi edge and the C 1s level of CNTs [3]. Figure
1 (b) shows PES spectrum of the C 1s core level of
CNT/SiC interface. By peak fitting, the energy
separation in the C 1s core level between CNT and
4H-SiC was estimated to be 1.36 eV. From the PES
measurements, we also determined the binding
energy of C 1s peaks of CNTs (~300 nm) to be284.35 eV and the energy separation between the
VBM and C 1s core level of SiC to be 279.30 eV.
From these values, the band alignment at the
CNT/4H-SiC interface was obtained, which is shownin Fig. 2. It should be noted that, in contrast with the
CNT/n-type 6H-SiC interface, Schottky barrier was
not formed at the CNT/4H-SiC interface. This result
is consistent with recent experimental results about
the current-voltage (I-V) characteristics of CNT(~300nm)/n-type 4H-SiC heterojunction, which showed
Ohmic behavior [4].

(a)hQ=60 eVAu
CNT

0.4 0.2 0.0 -0.2 -0.4

Binding Energy [eV]Intensity (a.u.)

Intensity (a.u.)This work was supported in part by the Joint
Studies Program (2011) of the Institute for Molecular
Science (IMS).

(b) 2C(sp )(CNT)

Si4C4O4

C 1s

Si4C4-xO2
SiC

288 286 284 282 280

Binding Energy [eV]Fig. 1. PES spectra of (a) VBM of CNTs with the
Fermi edge of Au (hQ = 60 eV), and (b) C 1s core
spectrum of CNT(several nm)/4H-SiC (hQ = 350 eV).

Fig. 2. Band alignment of CNT/n-type 4H-SiC
heterojunction.
[1] M. Kusunoki et al., Appl. Phys. Lett. 77 (2000)
531.
[2] M. Kusunoki et al., Chem. Phys. Lett. 366 (2002)
458.
[3] T. Maruyama, S. Sakakibara, H. Ito, H. Yamane, E.
Shigemasa and N. Kosugi, UVSOR Activity Report38 (2011) 127.
[4] S. Sakakibara, Unpublished master’s thesis. Meijo
University (2012).

123

BL6U

Valence-Band Dispersion of Metal Phthalocyanines (I): Intramolecular andIntermolecular Electron-Phonon Coupling in Crystalline Films of Zinc
Phthalocyanine (ZnPc)
H. Yamane and N. Kosugi
Department of Photo-Molecular Science, Institute for Molecular Science, Okazaki 444-8585, Japan
School of Physical Sciences, The Graduate University for Advanced Studies, Okazaki 444-8585, Japan
It has been predicted by theoretical studies that
local/nonlocal electron-phonon (e-ph) couplings playa crucial role in the charge transport mechanism inorganic molecular solids [1]. In order to examine theevidence for the e-ph coupling in organic molecularsolids experimentally, the temperature dependence ofthe electronic structure of well-crystallized samples
should be taken into account.
In this work at BL6U, we prepared a crystalline
film of zinc phthalocyanine (ZnPc), a promisingmaterial for organic electronics, with a flat-lie
orientation on the Au(111) surface, wherein theintermolecular ʌ-ʌ interaction exists along the surface
normal direction. The crystallinity of the ZnPc film
was confirmed by the incident-angle dependence of
X-ray absorption spectroscopy, the low-energy
electron diffraction, and the CuKĮ X-ray diffraction.
From the angle-resolved photoemission spectroscopy(ARPES) experiments on the crystalline ZnPc film on
Au(111), we observed an intermolecular valenceband dispersion [E(k)] and the clear evidence for thee-ph coupling as a function of the temperature.
Figure 1 (a) shows the E(k) relation of the ZnPc
crystalline film prepared on Au(111), which is
obtained from the incident photon energy dependenceof normal emission ARPES spectra at 15 K. The
highest occupied molecular orbital [HOMO (ʌ, a1u)],
Zn 3d (b1g), and ʌ’s derived peaks show a dispersive
shift with the same periodicity. This periodicity
corresponds well with the ī and Y points, estimatedfrom the lattice constant of 3.32 Å determined by the
X-ray diffraction [Fig. 1 (b)]. This correspondence
indicates that the observed E(k) is ascribed to theintermolecular interaction. The tight-binding fitting to
the E(k) relation gives a HOMO transfer integral (t)of 25 meV and hence a HOMO bandwidth (= 4t) of
100 meV, which is the same order as a reorganizationenergy for local phonons of CuPc [2].
In order to examine the e-ph coupling in crystalline
ZnPc film, we measured the temperature dependenceof ARPES at the ī and Y points with the 60 and 90
eV incident photons, respectively. As seen in Fig. 1(c), the ī-point HOMO energy shifts depending on
the temperature, while the Y-point HOMO energy
seems independent on the temperature. From thelineshape analysis, we found that both the HOMO
bandwidth (i.e., energy difference between the ī- andY-HOMO) and the ī-HOMO peakwidth are narrowed
with increasing the temperature, which is an opposite
trend in a remanent Au(111) Fermi edge. This is thedistinct indication of the local/nonlocal e-phinteraction; that is, band narrowing due to theinteraction with local phonons (intramolecularvibrations) and band widening due to interaction withnonlocal phonons (intermolecular lattice vibrations)[1]. Thus, it is considered that a transition between
incoherent hopping transport associated with the local
e-ph coupling and coherent band transport associatedwith the nonlocal e-ph coupling occurs within theenergy scale of 100 meV in the crystalline ZnPc.
From the E(k) relation, in the coherent band transport
at the low temperature, the effective mass and themobility of the HOMO hole are 13.8m0 and 28.9
cm2/Vs at 15 K, respectively.

1HOMO(ʌ, a1u)

2

0

Zn 3d
(b1g)

3

4

ʌ’s

5

ʌ’s+ı’s6
4

5

Wave Vector kA / Å-1

2T = 26.805°(aA = 3.32 Å)
×5
Au(111)
10

20

30

40

50

2T for CuKD / degree(c) ZnPc(250Å)/Au(111) @ T = 0°hv = 60 eV (~ *)hv = 90 eV (~ Y)

Intensity / arb. units

Binding Energy (vs. EF) / eV0(b) ZnPc(250Å)/Au(111) @300K

Intensity / cps(a) ZnPc(250Å)/Au(111) @15K

Remanent
Au peak

HOMO

2

1

T/K
300
180
120
80
15

0

Binding Energy / eV

Fig. 1. (a) Intermolecular E(k) relation of thecrystalline ZnPc film on Au(111) at 15 K, wherein thesecond derivative of the ARPES spectra were used formapping out. White curves for the HOMO and Zn 3d
derived bands are the best-fit results in the tightbinding model. (b) CuKĮ X-ray diffraction of thecrystalline ZnPc film on Au(111) at 300 K. (c)Temperature dependence of ARPES in the
HOMO-band region at the ī and Y points, measured
with 60 and 90 eV incident photons, respectively.
[1] C. A. Perroni, V. M. Ramaglia, and V. Cataudella,
Phys. Rev. B 84 (2011) 014303, and refs therein.
[2] S. Kera, H. Yamane, and N. Ueno, Prog. Surf. Sci.
84 (2009) 135.

124

BL6U

Valence-Band Dispersion of Metal Phthalocyanines (II): Phase Reversal of
the Intermolecular HOMO-Band Dispersion between Crystalline Films ofNon- and Per-Fluorinated Zinc Phthalocyanines
H. Yamane and N. Kosugi
Department of Photo-Molecular Science, Institute for Molecular Science, Okazaki 444-8585, Japan
School of Physical Sciences, The Graduate University for Advanced Studies, Okazaki 444-8585, JapanThe energy band dispersion [E(k)] is a fundamental
basis to understand electric properties of solids suchas hole mobility (ȝh). In the field of organic solids,
the study of the intermolecular E(k) shows a rapid
progress due to the needs of the interpretation of thecharge transport mechanism in organic electronic
devices. However, due to the weak intermolecularinteraction and the difficulty in preparing crystallized
films for the E(k) measurement, the intermolecularE(k) has been observed for high-ȝh materials. In order
to elucidate and control the functionality of organic
semiconductors, a systematic experiment on theintermolecular E(k) can play a crucial role.
Recently, we have observed the intermolecularE(k) in a crystalline zinc phthalocyanine (ZnPc) film
on Au(111) using the angle-resolved photoemission
spectroscopy (ARPES) at BL6U [1]. One of the key
points of this observation is that the E(k) relation of
ZnPc can be a benchmark for the systematic study on
the intermolecular interaction in terms of the intermolecular distance and/or the orbital symmetry bychanging terminal groups or central metals in themolecule. In this work, we have studied the effect of
the fluorination of ZnPc (i.e., F16ZnPc) on the E(k).
Figure 1 (a) and (b) show the ex situ CuKĮ X-ray
diffraction (XRD) and the in situ incident-angle (Į)
dependence of N-K X-ray absorption spectra (XAS),

R

10

×6
Au(111)20

30

2T for CuKD / degree

40

Intensity / normalized(b) XAS: F 16ZnPc(150Å)/Au(111) @15K

390

: D = 0°
: D = 20°
: D = 35°
: D = 50°
: D = 70°

D = 60°
T = 0±1°

Įhv

hv step= 4 eV(d) E(k) map @HOMO
0

eí

E

HOMO (ʌ)

hv
/eV
50
70

Binding Energy (vs. EF) / eV

2T = 27.565°(aA = 3.23 Å)

R
N
R
N N
N Zn N
N N
R
R
N
R
R
R R
R R

R

Intensity / normalized to photon flux

Intensity / cps
0

R

R

R

[1] H. Yamane and N. Kosugi, a separate page in thisvolume.

(c) ARPES: EBL-F16ZnPc@15K(a) CuKD XRD: F 16ZnPc(150Å)/Au(111) @300K
R

taken by the sample-current mode, measured for a
150-Å-thick F16ZnPc film on Au(111), respectively. A
F16ZnPc-derived XRD peak appears at 2ș = 27.565°,
which gives a lattice constant of 3.23 Å. This value isin good agreement with the molecular orientation
estimated from XAS, wherein the sharp ʌ* peaks at
397–400 eV are strongest at Į = 60°. From theseresults, we found that the molecules form a wellknown ȕ-bilayer (ȕBL)-crystalline structure, which issimilar to Į-crystalline ZnPc, with their molecularplane parallel to the Au(111) surface.
Figure 1 (c) shows the incident photon energy (hȞ)
dependence of ARPES at the normal emission for theȕBL-crystalline F16ZnPc film on Au(111). A highest
occupied molecular orbital (HOMO)-derived peakshows a clear periodic shift with hȞ. From the E(k)relation for the HOMO band of F16ZnPc and ZnPc in
Fig 1 (d), we found that the bandwidth of F16ZnPc(120 meV) is larger than that of ZnPc (100 meV),which is consistent with their intermolecular distance.
Moreover, the dispersion phase for the HOMO bandof F16ZnPc is flipped from that of ZnPc. This reversal
may originate from the ʌ···F 2p interaction, which isunder theoretical consideration.

1

2

F16ZnPc
4

410

420

Photon Energy / eV

430

6

0

1

90
110

2

ZnPc

130
400

5

6

5

4

3

2

1

0

Binding Energy (vs. EF) / eV

4

5

Wave Vector kA / Å-1

Fig. 1. (a) CuKĮ XRD of F16ZnPc films (150 Å) on Au(111). Inset is a molecular structure of ZnPc (R = H) and
F16ZnPc (R = F). (b) Incident-angle (Į) dependence of N-K X-ray absorption spectra for the crystalline film of
F16ZnPc on Au(111) at 15 K. (c) Photon energy (hȞ) dependence of ARPES for the crystalline film of F16ZnPc on
Au(111) at 15 K. (d) Comparison of the HOMO-band dispersion of ȕBL-F16ZnPc (top) and Į-ZnPc (bottom).

125

BL6U

Valence-Band Dispersion of Metal Phthalocyanines (III): Dimerization of
Mn-3d-Derived Band in Crystalline Films of ManganesePhthalocyanine (MnPc)
H. Yamane and N. Kosugi
Department of Photo-Molecular Science, Institute for Molecular Science, Okazaki 444-8585, Japan
School of Physical Sciences, The Graduate University for Advanced Studies, Okazaki 444-8585, Japan
Metal phthalocyanines (M Pc) have attracted muchattention due to their unique electronic and magnetic
properties in terms of organic electronics/spintronics.
In this work, we have examined the intermolecularinteraction in a crystalline film of MnPc, a promisingorganic spintronic material, by using angle-resolved
photoemission spectroscopy (ARPES) at BL6U.
Figure 1 (a) shows the CuKĮ X-ray diffraction(XRD) for a 200-Å-thick MnPc film on Au(111). A
MnPc-derived XRD peak appears at 2ș = 27.250°,
which gives a lattice constant of 3.27 Å. In the N-K
X-ray absorption spectra (XAS) for the crystalline
MnPc film on Au(111) [Fig. 1 (b)], sharp transition
peaks appear at 397–403 eV. Among these, peak X at
398 eV shows a broad incident angle (Į) dependence,
which may originate from the hybridization between
the N p and Mn d states. Other peaks are ascribed to
the N 1s ĺ ʌ* transitions. A sharp Į dependence of
these ʌ* peaks indicates the flat-lie orientation, whichis consistent with XRD.
Figure 1 (c) shows the incident photon energy (hȞ)
dependence of ARPES at the normal emission for the
crystalline MnPc film prepared on Au(111) at 15 K.
The highest occupied molecular orbital (HOMO) of
ZnPc is derived from the C 2p orbital; on the other
hand, the HOMO of MnPc is derived from the Mn 3d
orbital [1]. The HOMO (Mn 3d) and HOMO-1 (C 2p)

10

20

Au(111)30

2T for CuKD / degree

40

Intensity / normalized(b) XAS: MnPc(200Å)/Au(111) @15K

390

X

: D = 0°
: D = 20°
: D = 35°
: D = 50°
: D = 70°

2

0

D = 60°, T = 0±1°
hv step = 4 eV

Į

hv

eí

Mn 3d hv
/eV
50
70
90

410

420

Photon Energy / eV

430

4

1

C 2p
2

3

110
130

400

ī Y2a Y Y2a ī
Mn 3d

E
C 2p

2

(d) –d I(E)/dE of ARPES

Binding Energy (vs. EF) / eV

2T = 27.250°(a = 3.27 Å)

Intensity / normalized to photon flux

Intensity / cps
0[1] e.g., M.-S. Liao and S. Scheiner, J. Chem. Phys.
114 (2001) 9780; M.-S. Liao, J. D. Watts and M.-J.
Huang, Inorg. Chem. 44 (2005) 1941.

(c) ARPES: D-MnPc@15K(a) CuKD XRD: MnPc(200Å)/Au(111) @300K

×5

derived peaks of MnPc show a clear periodic shift
with hȞ. Moreover, we have found that the Mn 3d
derived peak consists of two components with thedifferent dispersion, indicating the relatively strong
interaction at the Mn site.
Figure 1 (d) shows the valence band dispersion
E(k) map for the crystalline MnPc film on Au(111),
wherein the symmetric ī and Y points are estimatedfrom the lattice constant of 3.27 Å from the XRD data.
The agreement between the periodicity of C 2p peak
and the symmetric points indicates that the observed
E(k) is ascribed to the intermolecular interaction. For
the Mn 3d peak, the peak split and its dispersion are
clearly seen in the E(k) map. Furthermore, theperiodicity of the Mn 3d dispersion is 1/2 times ofthat of the C 2p dispersion as labeled Y2a. A possible
origin of the Mn 3d dispersion is dimerization of theMn dʌ state.
Moreover, various band dispersions are observed in
the energy region of 3–4 eV. There is a possibility ofthe deconvolution of the broad photoemission peakinto individual molecular orbitals by the combination
with theoretical calculations.

3

2

1

0

Binding Energy / eV

4

4

5

-1

Wave Vector kA / ÅSUM
22–d I(E)/dE

Fig. 1. (a) CuKĮ XRD of MnPc films (200 Å) prepared on Au(111). (b) The incident angle (Į) dependence of N
K-edge XAS of crystalline films of MnPc on Au(111) at 15 K. (c) The photon energy (hȞ) dependence of ARPES
and (d) the intermolecular E(k) relation, which is obtained from the second derivative of the raw ARPES data, of
the crystalline film of MnPc prepared on Au(111) at 15 K.

126

BL6B

Far-Infrared Reflective Study of Alkali Niobate Ceramics
T. Nishi, T. Hayakawa, H. Yamada and K. Kakimoto
Graduate School of Engineering, Nagoya Institute of Technology Nagoya 466-8555, Japan
Alkali niobate is one of the promising candidatesof lead-free piezoelectric materials. Especially,
Li-doped (Na,K)NbO3 (LNKN) solid solution shows
excellent piezoelectric properties. Although it was
expected that the distortion of NbO6 unit would
contribute to the good ferroelectric properties, thedetails are still unclear. Thus, the knowledge of therelation between the vibration of NbO6 unit andferroelectric properties is strongly needed. The
vibration of the NbO6 is observed as a ferroelectric
soft mode in the infrared (IR) spectrum. So, this study
was aimed to the analysis of the ferroelectric softmode of Na0.5K0.5NbO3 which is a base material of
the LNKN.
The Na0.5K0.5NbO3 ceramic sample wassynthesized by a solid-state reaction method. The
surface of Na0.5K0.5NbO3 ceramic was polished andused for IR reflectivity study. The reflectivity
far-infared spectram was obtained in the range of 20 –8000 cm-1 at room temperature using a synchrotron
radiation source, then the spectra were corrected byusing a Michelson interferometer (Bruker, IFS66v).
Figure 1 shows the reflectance spectram of
Na0.5K0.5NbO3 ceramics at room temperature. The
spectrum was analyzed using the following formula,

Reflectivity

0.4

2

4 6

2

4 6

2

4 6

100
1000
-1
Wavenumber (cm )Fig. 1. Infrared reflectance spectra of Na0.5K0.5NbO3
ceramics at R.T..
0.9
Raw date
Fitting result

0.8n

H (Z ) H f 

0.6

0.2

Reflectivity

Z 2jLO  Z 2  iZJ jLO
22
j 1 Z jTO  Z  iZJ jTO

0.8

(1).

Figure 2 shows the infrared reflectance spectram
around the wavenumber of the soft mode and thefitting curve. Table 1 shows the fitting parameters of
the soft mode. The soft mode frequency obtained
from the IR spectroscopic analysis was 255 cm-1.
Relative permittivity Ťr calculated from the soft mode
wavenumber was Ťr of 520. This result was different
from the measured value (capacitance method, Ťr of
555 at 1kHz). The permittivity in the range of THz,which was the observed soft mode, is contributed aionic polarization. On the other hand, the permittivity
in the range of kHz was due to ionic and orientation
polarizations. This suggested that the ionicpolarization contributed the NbO6 unit accounted for
94 % of the total permittivity of (Na0.5K0.5)NbO3.
Thus, the relation between the distortion of the NbO6
units and a ferroelectric property can be evaluated byusing the IR spectroscopy analysis. We will
investigate the relation between a ferroelectric
property and the amount of a distortion of the NbO6
by analyzing the IR spectra of Li-doped (Na,K)NbO3
solid solution.

0.7

0.6

0.5
160

180

200 220 240-1Wavenumber (cm )

260

280

Fig. 2. Infrared reflectance spectra around the softmode.
Table 1. Fitting parameters of the soft mode.
Fitting parameterIRȦTO
254.51
ȦLO
434.34ȖTO
163.80ȖTO
568.14This research was supported by the Industrial
Technology Research Grant Program in 2007 fromthe New Energy and Industrial Technology
Development Organization (NEDO) of Japan.

127

BL7B

Luminescence Decay of Organic-Inorganic Perovskite Compounds
M. Koshimizu, N. Kawano and K. Asai
Department of Applied Chemistry, Graduate School of Engineering, Tohoku University, Sendai 980-8579, Japan128

the inorganic layer and the organic layer were excited.
The decay characteristics were also quite similar.
Hence, these results indicate that the excitation in theorganic layer has little influence on the luminescenceand scintillation behavior.

Intensity (a.u.)

20000
15000
10000
5000
0
350

400

450

500

550

600

Wavelength (nm)Fig. 1. Luminescence spectrum of the sample excited
at 300 nm.

Intensity (a.u.)

1000

Excitation: 300 nm
100

10

1
-20

0

20

40

60

80

100

Time (ns)100

Intensity (a.u.)

Quantum confinement in semiconductors has beena subject of great interest from the viewpoint of theapplications for optical devices, as well as the basic
science.
Amongvariousmaterials
havinglow-dimensional quantum confinement structure,Pb-basedorganic-inorganic
perovskite-type
compounds have attracted much attention due to theirunique optical properties. The crystal of thesecompounds composed of alternating inorganic andorganic layers. The inorganic and organic layers arecomposed of PbX62- octahedra and organic amine,
respectively. Due to the large difference in band-gap
energy and dielectric constant in the organic andinorganic layers, an exciton in the inorganic layer istightly confined, resulting in huge binding energy(several hundreds of meV) and large oscillator
strength.
We have been developing scintillation materials
using these compounds. In recent years, the demand
for fast scintillation materials has increased for
applying in radiation detectors with an excellent
timing property and the ability to operate at a highcounting rate.The large binding energy and oscillator
strength of the exciton enables us to fabricate
scintillation materials with excellent timing response[1]. During the course of the development, we foundthat the scintillation properties largely depend on theamine molecules used in the organic layer, in which
no luminescence or scintillation occurs. In order toreveal the origin of this dependence, theluminescence decay characteristics should beobtained under excitation (1) only in the inorganic
layers and (2) both in the inorganic and the organic
layers. In this report, we show that the excitation in
the organic layers has little influence on theluminescence behavior.
The chemical formula of the sample compound is(C6H5C2H4NH3)2PbBr4. The single crystals of thecompound were grown by solvent diffusion technique.
The photoluminescence spectra and time profiles
were measured using vacuum ultraviolet (VUV) lightas an excitation source at UVSOR BL7B operated insingle bunch mode.
Figure 1 shows the photoluminescence spectrum.
The prominent peak at 410 nm is ascribed to free
excitons. A small shoulder at 440 nm may be ascribed
to excitons trapped at some defects or impurities,
because the intensity of this shoulder has beendifferent in different crystals. The shape of the spectra
was quite similar for the excitation wavelength rangeof 50 – 300 nm.
Figure 2 shows the photoluminescence time
profiles at 410 nm under excitation at 300 nm and180 nm. In the case of the excitation at 180 nm, both

Excitation: 180 nm
10

1
-20

0

20

40

60

80

100

Time (ns)Fig. 2. Luminescence time profiles at 410 nm.
[1] K. Shibuya et al., Appl. Phys. Lett. 84 (2004)
4370.
[2] K. Shibuya et al., Jpn. J. Appl. Phys. Part2(Express Letter) 43 (2004) L1333.
[3] S. Kishimoto et al., Appl. Phys. Lett. 93 (2008)
261901.

BL7B

Photoluminescence upon Vacuum Ultraviolet Excitation in Pr3+ ion doped
Inorganic Materials
Y. Inaguma, H. Horiguchi, S. Sasaki and D. Mori
Department of Chemistry, Faculty of Science, Gakushuin University, Tokyo 171-8588, Japan
Pr3+ ion doped inorganic materials has been
extensively investigated in relation with various
applications [1]: laser materials utilizing visible
emission from the 3P0 and 1D2 states (blue-green andred emission, respectively), or infrared emission fromthe 1G4 state of Pr3+, field emission display (FED)
devices utilizing red 1D2 emission, scintillators or
tunable ultraviolet (UV) lasers utilizing UV emission
from the 4f5d state. The emission color or energy inPr3+ ion doped inorganic materials is strongly related
to the electronic structure, especially, the band gap
energy of host materials. Therefore the emission color
can be controlled by the choice of host materials [2].
Recently phosphor materials emitting UV lightupon vacuum ultra violet (VUV) excitation aredesired as the substitutes of mercury lamp due to thetoxicity of mercury. As mentioned above, since theUV emission from the 4f5d state in Pr3+ can beobserved, Pr3+ ion doped compounds are expected tobe candidates of the UV phosphors. The UV emission
from the 4f5d state would be observed when the band
gap energy of the host material is in the VUV region.
In this study we then chose Pr3+ ion doped
alkaline-earth phosphates with the band gap energy inthe VUV region and investigated their luminescence
properties.
Herein,wereportonthephotoluminescence upon VUV excitation of
Ca3(PO4)2:Pr among them.
The polycrystalline samples were synthesized by aconventional solid state reaction at elevatedtemperature. The mixture of starting materials:
alkaline-earth carbonate such as CaCO3, (NH4)H2PO4
and Pr nitrate solution were calcined at 1000qC in air.
The calcined powder was pressed into pellets andsintered at 1000-1200qC in air. The phase
identification for the samples was carried out by thelaboratory powder X-ray diffraction (XRD) using a
Rigaku RINT 2100 diffractometer with a Bragg
Brentanogeometry(graphite-monochromatized
CuKD radiation). The emission and excitation spectra
were recorded in the beamline BL7B at the UVSOR
facility.
The obtained sample was crystallized to be thehexagonal E-type form of tricalcium phosphates
composed of isolated PO4 tetrahedra and Ca2+ ions,which was confirmed by a powder XRD experiment.
Figure 1 shows the emission spectrum upon
excitation of 173 nm (right side) and excitation
spectrum emission monitored at 280 nm (left side) for
Ca3(PO4)2:Pr 2% at room temperature. Here, theemission spectrum was not corrected, while theexcitation spectrum was corrected for the spectral

distribution of excitation light source. As seen in
Fig. 1, the UV emission peaks at 240 and 280 nm
were observed, which correspond to the transitions
from the 4f5d state to the 4f2 (3HJ) state of Pr3+. In the
excitation spectrum, the peaks around 175 nm and200-220 nm were observed, which correspond to thehost absorption band, i.e. the intra charge transfer ofPO43-[3] and the transition from the ground state of
4f2(3H4) to the 4f5d state of Pr3+, respectively.
It was found that Ca3(PO4)2:Pr exhibits the UVemission corresponding to the transition from the
4f5d state of Pr3+ and the Pr3+ ion doped
alkaline–earth phosphates are candidates of UV
phosphors.
The luminescence properties upon VUV excitationin other Pr3+ ion doped alkaline phosphate will bereported in the near future.

Intensity / arb.unit

Excitation
Oem = 280nm

150

200

Emission
Oex = 173 nm

250

300

Wavelength / nm

Fig. 1. Emission spectrum upon excitation of 173 nm(right side) and excitation spectrum emission
monitored at 280 nm (left side) for Ca3(PO4)2:Pr 2%
at room temperature.
The authors thank Mr. M. Hasumoto and Prof. S.
Kimura for their experimental supports and helpfuladvices.
[1] E. Van der Kolk et al., Phys. Rev. B 64 (2001)
195129.
[2] Y. Inaguma et al., Inorg. Chem. 50 (2011) 5389.
[3] H. Ling et al., J. Solid State Chem. 177 (2004)
901.

129

BL8B

Photoelectron Spectrum of Zn-Phthalocyanine/C60 Interface under theLight Irradiation
S. Tanaka, K. Fukuzawa and I. Hiromitsu
Interdisciplinary Faculty of Science and Engineering, Shimane University, Matsue 690-8504 Japan130

It is known that the OSC with a BHJ layer showshigher energy conversion efficiency than that with a
PHJ OSC [3]. However, the photoelectron spectrumof ZnPc:C60 (volume ration = 1:1) BHJ sample
showed no significant photoinduced shift under thelight irradiation. The present results indicate that
semipermeable membranes for the electric charges
are required on both sides of the BHJ layer togenerate a charge current flowing in the OSC.

(a) C60 (3nm) on ZnPc (15nm)
hQ = 40 eV
Bias = -5 V

Intensity / Arb. units

Dark
Under irradiation~ 0.04 eV

4.8

5.0

5.2

5.4

5.6

Kinetic energy / eV(b)
Photoinduced shift / meV

The electronic structure at the organic/organic andorganic/metal interfaces plays an important role forthe photovoltaic properties of the organic solar cell
(OSC). The photoelectron spectroscopy is the one ofthe powerful methods to investigate the electronic
structure of the organic/organic and organic/metalinterfaces [1]. We have recently reported theelectronic structures of the organic thin films under
the light irradiation which were investigated by using
the photoelectron spectroscopy with an external lightsource. In the previous report, we reported thephotoinduced shift in the photoelectron spectra of theZinc-phthalocyanine (ZnPc)/C60 interface [2]. The
shift direction of the photoelectron spectrum wascorresponded to the polarity of the photovoltaic effectat the ZnPc/C60 interface. In this report, the progressof the experimental results of the photoelectron shift
under the light irradiation is reported.
The organic films were fabricated by the vacuum
deposition technique. Two kinds of sample were
fabricated: the planer hetero-junction (PHJ) sample
and the bulk hetero-junction (BHJ) sample. The PHJ
sample had a planer interface of ZnPc and C60 on theindium-tin oxide substrate. The BHJ sample wasfabricated by the co-deposition of ZnPc and C60. For
the observation of the photoinduced effect on theelectronic structure, the photoelectron spectra under
the light irradiation were compared with that underdark. A solar simulator was used as the light source
for the irradiation. The simulated airmass 1.5 solar
illumination was exposed to the sample in a vacuum
chamber through a viewport. All the photoelectron
measurements were carried out with a photon energyof 40 eV at room temperature.
The photoelectron spectrum of the ZnPc/C60
interface was shifted toward higher kinetic energies
by the light irradiation (Fig.1 (a)). The C60 thickness
dependence of the photoinduced shift on thephotoelectron spectrum is shown in Fig. 1 (b). The
shift value was saturated at around 10nm of the C60
layer thickness and was approximately 60 meV. The
photoinduced shift was caused by the photovoltaic
effect at the ZnPc/C60 interface. The photogenerated
electrons were accumulated in the C60 layer, causing
the acceleration of the photoelectrons (Fig. 1(b) inset).
The increase of the shift value in the initial stage of
the C60 deposition was probably related to theincrease of the interfacial area of the ZnPc/C60.
However, the saturated shift value of 60 meV wasapproximately one order smaller than the open circuit
voltage of the ZnPc/C60 PHJ OSC. The reason of thisdiscrepancy is still an open question.

100
80
60
40
䋭 䋭 䋭 C60 䋭 䋭 䋭

20

ZnPc

0

ITO

0

5

10

15

20

C60 thickness / nm

Fig. 1. (a) Photoelectron spectra of the C60 (3 nm)
on the ZnPc layer at the secondary electron cutoff
region. The red (black) curve indicates thephotoelectron spectrum under the light irradiation(dark). (b) C60 thickness dependence of thephotoinduced shift of the photoelectron spectrum.
Inset shows the schematic image of the charge
distribution at the sample surface under the lightirradiation.
[1] H. Ishii et al., Adv. Mater. 11(1999) 605.
[2] S. Tanaka et al., UVSOR Activity Report 38(2010) 136.
[3] T. Taima et al., Appl. Phys. Lett. 85 (2004) 6412.

BL8B

Influence of Growth Conditions on the Interface Electronic Structure at
Planar Organic Donor/Acceptor Heterojunctions of Sexithiopheneand C60
U. Hörmann1, Y. Nakayama2, Y. Ozawa3, W. Brütting1 and H. Ishii2, 3
1

Institute of Physics, University of Augsburg, Augsburg 86135, Germany
Center for Frontier Science, Chiba University, Chiba 263-8522, Japan
3
Faculty of Engineering, Chiba University, Chiba 263-8522, Japan2

On the way to high performance organic solar cells
tremendous effort has been put into increasing theshort circuit current by structural device optimization.
Little attention has, however, been paid to the influence of morphology on the open circuit voltage (Voc).
Here we investigate the electronic structure at sexithiophene (6T)/C60 hetero interfaces under different
preparation conditions. In solar cells this systemyields a Voc of 0.5 V for room temperature grownfilms. Accompanied by morphological changes, thisvoltage is reduced to 0.37 V if C60 is deposited onto ahigh temperature grown 6T film.
In high vacuum 6T layers of 10 nm and 5 nm
thickness have been deposited onto PEDOT:PSS(Clevios AI4083) covered indium tin oxide substrates
at substrate temperatures of 100 °C (HT) and 30 °C(RT), respectively. In both cases, the deposition rate
was set to 0.3 – 0.8 Å/s. C60 was deposited stepwise
onto the 6T films at a rate of 0.3 Å/s at room temperature. Ultraviolet photoelectron spectroscopy(UPS) was performed at BL8B at UVSOR. The photon energy, incident angle and photoelectron emission
angle were set to 30 eV, 45° and 0°, respectively
throughout this study. For valence-band region (VB)
spectra a bias of +5 V was applied to the sample inorder to avoid charging, whereas -13 V wasapplied
for secondary electron cutoff (SECO) measurements.
As expected from previous studies, heating in vacuum increases the work function of PEDOT:PSS,
from 5.0 eV to 5.3 eV. Figure 1 shows the recorded
UPS spectra of the 6T films and C60 overlayers on
them. For direct comparison RT (black) and HT (red)data are shown in one plot. Both the SECO and VB
spectra of the neat 6T film remain energetically unchanged by substrate heating. However, the features
become sharper for the HT grown film. This canprobably be attributed to the presence of two different
crystalline 6T phases in the RT film as indicated by
x-ray diffraction (XRD) measurements [1].
Interestingly the whole spectrum (VB and SECO)
is shifted by ' 0.12 eV towards higher binding energies, when 0.3 nm of C60 are deposited onto the HT
6T film. This indicates the occurrence of band bending at the C60/6T(HT) interface as illustrated in Figure
2. By contrast, no band bending is observed in the RT
case, where the energetic position of the overall spectrum remains unchanged by 0.3 nm of C60 (eventhough the HOMO onset shows a slight downshift due
to spectral broadening).
In both cases the SECOs shift towards higher vac-uum level energies with increasing C60 coverage. Thiscan possibly be ascribed to mutual polarization of themolecules.The RT shift of +0.15 eV is in accordancewith Ref. [2]. In the HT case the vacuum level recovers to the neat 6T value of 4.6 eV. Note, however, that
6T features are still present in the VB spectrum in thiscase. This suggests that the SECO shift probably has
not reached its maximum and might stem from an
incomplete C60 coverage on the extremely rough
6T(HT) surface observed by AFM.
Comparison of the 5 nm VB spectra reveals thatgrowing C60 on HT 6T shifts the C60 HOMO to larger
binding energies from 1.9 eV to 2.0 eV. This shift
results in a reduction of the 6T/C60 charge transfer
gap for the HT sample (indicated in Fig.2), which isin accordance with the reduced Voc observed for
6T(HT)/C60 solar cell devices.

Fig.1. SECO (left) and VB (red) spectra of C60 grownon the RT (black) and HT (red) grown 6T.

Fig.2. Schematic energy diagram of C60 grown on RT(left) and HT (right) grown 6T. Values are energy
relative to the PEDOT:PSS Fermi level (units: eV).
[1] A. Hinderhofer and F. Schreiber (UniTübingen),
unpublished.
[2] A. Wilke and N. Koch (HU Berlin), in prep.

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