Ⅲ-2

Materials Sciences

461BL7U 出田真一郎.docx

Ⅲ -2. Materials Sciences

BL7U

Electronic Structure of Triple-- Layer Cuprate Superconductor
Bi2Sr2Ca2Cu3O10+δ Observed by ARPES
S. Ideta1, T. Yoshida2, S. Ishida3, K. Takashima4, S. Uchida3, A. Fujimori4 and K. Tanaka1
National Institutes of Natural Science, Institute for Molecular Science, Okazaki, 444-8585, Japan
Graduate School of Human and Environmental Studies, Kyoto University, Kyoto 606-8501, Japan
3
National Institute of Advanced Industrial Science and Technology (AIST), Tsukuba 305-8568, Japan
4
Department of Physics, University of Tokyo, Tokyo 113-0033, Japan1

2

Effects of intra-multilayer interaction between the
CuO2 planes, which may contribute to theenhancement of the critical temperature (Tc) in themulti-layer high-Tc cuprate superconductors (HTSCs)[1], have been investigated to elucidate the mechanismof superconductivity. The Bi-based HTSC is classified
by the number of the neighboring CuO2 layers (n):single-layer (n = 1) Bi2Sr2CuO6+δ (Bi2201), doublelayer (n = 2) Bi2Sr2CaCu2O8+δ (Bi2212), and triplelayer (n = 3) Bi2Cr2Ca2Cu3O10+δ (Bi2223). The Tc of
optimally doped Bi2223, 110 K, is the highest among
the Bi-family cuprates. In spite of both extensive
experimental and theoretical studies [2-6], themicroscopic origin of the Tc enhancement in the triplelayer HTSCs still remains unclear.
In recent ARPES studies of Bi2223, we havereported two bands originating from the outer andinner CuO2 planes [7-9]. Here, we report an ARPES
study of Bi2223 using low-energy photons with linear
polarization along the Cu-O bond direction. We
observed three bands and three Fermi surfaces (FSs)
resulting from interaction between the three CuO2
planes.
High-quality single crystals of optimally doped
Bi2223 (Tc = 110 K) were grown by the travelling
solvent floating zone method. ARPES experiments
were carried out at BL 7U of UVSOR using linearlypolarized light of hν = 7 eV - 8.5 eV. The total energy
resolution and angular resolution were set at 7 - 8 meV
and 0.05 deg, respectively. Temperature was set at T =12 K and clean sample surfaces were obtained bycleaving the single crystals in-situ in an ultrahighvacuum better than 7×10-9 Pa.
A schematic picture of FSs is shown in Fig. 1(a).
The FS from the outer plane (previously denoted asOP) is split into two FSs due to the bonding andantibonding combinations of the two OPs, as denotedby OP-B and OP-A, respectively. Their superstructuresdue to the structural modulation of the BiO layer are
also shown. Energy-momentum (E-k) intensity takenin the SC state along the momentum cuts indicated byblack curves in panel (a) are shown in Figs. 1(b)-1(e).
While the OP band is not clearly split near the nodal
region [Fig.1(b)], away from the nodal region, thisseemingly single feature splits into two dispersions as
indicated by arrows for OP-B and OP-A bands.
In conclusion, we have observed the splitting of thetriple-layer bands by ARPES. This result may give a

key to elucidate the mechanism of the triple-layer
cuprate superconductivity with the highest Tc of 110 K.

Ⅲ-2

Fig. 1. Energy-momentum intensity plots in the Cu-O
bond direction taken with hν = 8 eV at T = 12 K.
ARPES intensity plots [(b)-(e)] correspond to cuts inpanel (a). IP, OP-B, and OP-A denote the inner-plane,bonding outer-plane, and antibonding outer-planebands, and SS is the superstructures.
[1] J. M. Wheatley et al., Nature 333 (1988) 121.
[2] E. Pavarini et al., Phys. Rev. Lett. 87 (2001)
047003.
[3] S. Chakcraverty et al., Nature 428 (2004) 53.
[4] S. Okamoto et al., Phys. Rev. Lett. 101 (2008)
156401.
[5] D. L. Feng et al., Phys. Rev. Lett. 88 (2002) 107001.
[6] T. Sato et al., Phys. Rev. Lett. 91 (2003) 157003.
[7] S. Ideta et al., Phys. Rev. Lett. 104 (2010) 227001.
[8] S. Ideta et al., Phys. Rev. B 85 (2012) 104515.
[9] S. Ideta et al., Physica C 470 (2010) S14-S16.

45

121BL1B
淡野照義.docx
Ⅲ -2. Materials
Sciences

BL1B

Terahertz Spectroscopy of Solid Acid Proton Conductor
CsHSO4 - CsH2PO4 Single Crystals
T. Awano

Faculty of Engineering, Tohoku Gakuin University, Tagajo 985-8537, Japan

0.6

Cs3(HSO4)2(H2PO4)

0.5

REFLECTIVITY

0.4

E // b

300K
78K
10K

0.3
0.2
0.1
0
0.8

300K
78K
10K

E⊥b

0.6
0.4
0.2
0

20

40
60
80
WAVE NUMBER (cm -1)

100

Fig. 1. Reflectivity spectra of Cs3(HSO4)2(H2PO4)
crystal from CsCO3, H2SO4 and H3PO4.
0.5
0.4

Cs2(HSO4)(H2PO4)(a) as grown

0.3
0.2
REFLECTIVITY

Ⅲ-2

Proton conductors are interested because of the use
in fuel cells. Those in medium temperature areimportant for eco-system in house with high efficiency.
Inorganic solid acids such as Cs2(HSO4)(H2PO4)
crystal are newly discovered proton conductors in suchmedium temperature. Especially, this crystal keeps
high ionic conductivity once after heating above 380
K. It is interesting to make clear the mechanism of suchconductivity dependence on thermal history.
Meanwhile, terahertz spectral region is locatedbetween vibration and translation motion of the mobileion. The “attempt mode” of conduction ion is also
observed in superionic conducting crystals of silver,
cuprous and alkali ions at the frequency as their atomic
mass. This vibrational mode is directly connected withtranslational movement of conduction ion. Namely, theconduction ions surrounded by tetrahedral cage try tomove out at this frequency.
Three methods have been reported for crystal
growth of Cs2(HSO4)(H2PO4). Haile who discovered
this compound obtained it from stoichiometricaqueous solutions of cesium carbonate, sulfuric acidand phosphate acid [1]. By this method, both ofCs2(HSO4)(H2PO4) and Cs3(HSO4)2(H2PO4) are grown.
Other two methods were used by some groups. The
one is growth from stoichiometric aqueous solutions of
CsHSO4 and CsH2PO4 [2]. The other is growth fromstoichiometric sulfuric acid solution of CsH2PO4 [3].
Cs3(HSO4)2(H2PO4) also has superionic conducting
phase at high temperatures [4].
The grownCs2(HSO4)(H2PO4) and Cs3(HSO4)2(H2PO4) single
crystals were 0.2 mm × 2 mm × 6 mm and 5 mm × 8
mm × 15 mm in maximum size after one week growth,
respectively.
Figure 1 shows reflectivity spectra ofCs3(HSO4)2(H2PO4). Anisotropic peaks were observed
as shown in the figure. The proton conduction assistedby hydrogen bond with rotating of SO4 or PO4
tetrahedron occur along b-axis of the crystal.
Reflectivity increase at low wave number region at theE//b spectrum at 300 K seems to be due to the proton
conduction.
Figure 2 shows reflectivity spectra ofCs2(HSO4)(H2PO4). Anisotropic peaks were observed
also in these spectra, however, the difference was not
so clear. Crystal structures of Cs3(HSO4)2(H2PO4) andCs2(HSO4)(H2PO4) are similar and corresponding
peaks were observed. The high temperature phase is
cubic and there was no difference with direction ofelectric field. Increase at low wave number was also
observed at the 300 K spectrum.

E // b 300K
E⊥b 300K
E // b 78KE⊥b 78K

0.1
00.4(b) 440 K annealed

0.3
0.2

300K
78K

0.1
0

20

40

60
80
100
WAVE NUMBER (cm -1)

120

Fig. 2. Reflectivity spectra of Cs2(HSO4)(H2PO4)
crystal from CsHSO4 and CsH2PO4; (a) as grown (b)after heating at 440 K.
[1] C. R. I. Chisholm and S. M. Haile, Acta Cryst. B
55 (1999) 937.
[2] A. Matsuda, T. Kikuchi, K. Katagiri, H. Muto and
M. Sakai, Solid State Ionics 177 (2006) 2421.
[3] Y. Yamane, K. Yamada and K. Inoue, Solid State
Ionics 179 (2008) 483.
[4] S. M. Haile, G. Lentz, K.-D. Kreuer and M. Maier,
Solid State Ionics 77 (1995) 128.

46

131BL2B 田中仙君.docx

Ⅲ -2. Materials Sciences

BL2B

Degradation of CH3NH3PbI3 to PbI2 Driven by Annealing in Vacuum
S. Tanaka1, T. Kajikawa1 and K. R. Koswattage2

Graduate School of Science and Engineering, Kindai University, Osaka 577-0818, Japan
2
Center for Frontier Science, Chiba University, Chiba 263-8522, Japan

Organolead halide perovskites have recently
emerged as a photovoltaic material for high- efficiency
organic photovoltaic [1-3]. The organolead halide
perovskites have the advantageous characteristics forthe photovoltaics such as the direct bandgap near thevisible region, the small electron/hole effective masses,
the large absorption coefficients, and long diffusion
length of excited carriers, etc. The power conversion
efficiency (PCE) of the organic photovoltaic usedperovskites is approaching the PCE of silicon solar
cells. However, there are several problems for practicaluse of the perovskite solar cells. One of the problems
is its stability. A prototype organolead halide
CH3NH3PbI3 (MAPbI3) shows degradation to PbI2, amaterial has different electronic properties from thoseof the starting perovskite. As the degradation
mechanisms of MAPbI3, the chemical degradation
through the interaction with water and thethermodynamic degradation through the creation ofvolatile molecular defects has been proposed [4]. In
this report, we studied effect of the thermal annealingof MAPbI3 on the electronic structure by using
photoelectron spectroscopy to understand an impact ofthe decomposition of the perovskite structure.
Thin films of organolead halide perovskite MAPbI3
were fabricated by the 2-step vapor deposition method
at the Kindai University. As the first step, lead chloride(PbCl2) was deposited on a pre-cleaned indium-tinoxide substrate. Methylammonium iodide (CH3NH3I)
was subsequently deposited on the PbCl2 film as thesecond step. Upon CH3NH3I deposition, PbCl2 reactedwith CH3NH3I in situ and formed MAPbI3 films. The
fabricated samples were stored in dark, and transferredto the UVSOR BL2B. The samples were exposed in air
before the photoelectron measurements. Hence thesurface of samples would be partially degraded.
Figure 1 shows the photograph of the MAPbI3
sample before and after annealing. The annealing
treatment was performed in vacuum with a basepressure of 3×10−7 Pa. The annealing temperature andtime were 140 C and 10 min, respectively. During theannealing, the pressure was increased to 3×10−5 Pa.
The red-brownish color of MAPbI3 film was turned
yellowish after the annealing treatment. This indicatedthe annealing treatment induced the degradation of
MAPbI3 to PbI2. The photoelectron spectra of before
and after the annealing treatment were measured withthe photon energy of 100 eV. The results are shown in
Fig. 2. The secondary cutoff of the sample was shifted
toward higher kinetic energies after annealing. On theother hand, the Pb 5d core-levels were shifted toward
lower kinetic energies after annealing. The

photoelectron signal from I 4d core-levels were also
shifted toward lower kinetic energies with same valuesby annealing. (not shown here) The valence band
region showed slight changes of shape but nosignificant shift of the valence band edge position.
However, the signal-to-noise ratio of the spectrum was
not adequate to detailed discussions about the valenceband region.
The coordination number of Pb ions in the MAPbI3
lattice is 6 and is also maintained in the PbI2 lattice.
Therefore, it is expected that the chemical environmentfor Pb ions in MAPbI3 are similar that in PbI2. Hence,the variation of kinetic energy of Pb 5d photoelectrons
implies differences of the interaction from thesurrounding molecules during the escape process. The
shift toward the same direction of I 4d core-levels may
also support this assumption. Further analysis ofphotoelectron spectra to understanding the effects of
MAPbI3 degradation is in progress.
(a)(b)

Fig. 1. Photograph of the CH3NH3PbI3 thin film (a)before and (b) after annealing.
(a)

(b)

Intensity / Arb. unitsh = 100 eVBias =  5 V

45Intensity / Arb. units

1

6

7

Kinetic energy / eV

72

(c)

Pb 5d

After
annealing

Beforeannealing
74

76

78

85

90

95

Kinetic energy / eV

Fig. 2. Photoelectron spectra of (a) secondary
electron cutoff, (b) Pb 5d and (c) valence band regions.
Bias voltage of −5 V was applied for the secondary
electron measurements. Black and red line are thephotoelectron spectrum of before and after annealing
treatment, respectively.
[1] A. Kojima et al., J. Am. Chem. Soc. 131 (2009)6050.
[2] M. M. Lee et al., Science 338 (2012) 643.
[3] P. P. Boix et al., J. Phys. Chem. Lett. 6 (2015) 898.
[4] I. Deretzis et al., Appl. Phys. Lett. 106 (2015)
131904.

47

Ⅲ-2

141BL2A
佐藤充孝.docx
Ⅲ -2. Materials
Sciences

BL2A

Trial of Line Analysis for Si-Doped Porous Hydroxyapatite Prepared on
Mg-Based Alloy
M. Sato1, M. Togo2 and A. Nakahira2, 3Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan
Graduate school of Engineering, Osaka Prefecture University, Sakai 599-8531, Japan
3
Kansai Center for Industrial Materials Research, Tohoku University, Sakai 599-8531, Japan
1

2

Ⅲ-2

Hydroxyapatite (HAp) has been widely used asalternate materials our hard tissue due to its excellent
biocompatibility and osteo-conductivity. In addition,
HAp also has protein absorption ability and ionexchange ability. On the basis of these properties, wehave prepared the functional element-doped HAp suchas Mg, Si, Fe, Cu, and Ag by conventional wet process
so far and investigated the effect of addition of thoseelements on generation phase, microstructure, andother properties. Since HAp powders are usuallycoating on the surface of implant in a practical use,
structural investigation of these interface is importantin order to perform of its full potential as artificial hardtissue. Hydrotermal Hot Press (HHP) is unique
technique to obtain HAp bulk sample having a porous
structure [1]. This porous structure has many
advantages for bone formation. In this study, Si-doped
HAp was grafted with AZ-31 (Mg-3Al-1Zn) alloy by
HHP method, and elution behavior of Mg ion from
AZ-31 substrate toward a HAp phase was investigated
by XAFS measurement.
First of all, Si-doped amorphous calcium phosphate(ACP) powder was prepared by sol-gel method using
calcium metal, tetraethyl orthosilicate (TEOS),
phosphoric acid and ethanol at 273 K. Since the Ca/P
ratio of obtained ACP powder is usually around 1.5,
the obtained Si-doped ACP and Ca(OH)2 powders
were mixed so that their (Ca+Si)/P ratio is 1.67. And
then, Si-doped HAp bulk sample was prepared on AZ31 alloy by hydrothermal hot press method at 423 K
for 2h. In order to obtain the interface information,
obtained bulk sample was cut and polished, andsynchrotron beam was irradiated respectively at thesubstrate, the interface and the HAp phase by changing
a measurement position of sample. The Si K-edge
and Mg K-edge XANES spectra for obtained bulk
sample were corrected by fluorescence method using
InSb and Beryl double crystal monochrometer andsilicon drift detector (SDD) at BL2A in UVSOR,
respectively. Figure 1 shows Si K-edge XANES
spectra of Si-HAp prepared by HHP process obtained
from interface and HAp phase. The obtained XANES
spectra did not change by the obtaining position andresembled to that of calculated spectrum. This result
indicate that the added Si ion exist into HAp structureand Si doped HAp bulk sample had been obtained by
HHP technique. Moreover the shape of spectrum did
not change by correcting position.
Figure 2 shows Mg K-edge XANES spectra ofsample obtained at different position. Obtained

XANES spectra changed by changing the position, andthat obtained from HAp phase showed similar
spectrum to calculated spectra. This result indicatesthat the Mg ion eluted from substrate during HHP
process forms a solid solution with HAp structure. In
addition, spectrum obtained from near-interface has
similar trend to that of Mg(OH)2. From this result, it issuggested that the local structure around Mg atomchanges gradually toward the HAp phase fromsubstrate.

Fig. 1. Si K-edge XANES spectra of samples.

Fig. 2. Mg K-edge XANES spectra of samples.
[1] T. Onoki, S. Yamamoto, H. Onodera and A.
Nakahira: Mater. Sci. Eng. C 31 (2011) 499.

48

151BL2A 有谷博文.docx

Ⅲ -2. Materials Sciences

BL2A

Mo L-Edge XANES Study on the Formation of Catalytically Active Mo
Carbide Species on H-MFI for MTB Catalysts
H. Aritani1, S. Mogi1, N. Naijo1, T. Sugawara1, M. Akutsu1, K. Kawashima1 and A. Nakahira321Graduate School of Engineering, Saitama Institute of Technology, Fukaya 369-0293 Japan
Advanced Science Research Laboratory, Saitama Institute of Technology, Fukaya 369-0293 Japan
3
Graduate School of Engineering, Osaka Prefecture University, Sakai 599-8531 Japanis remarkable during the MTB reaction after 60
minutes. These results indicate that formation of αMo2C species relates to the deactivation, and inhibitionof the carbonization is effective for high and durable
MTB reactivity.

MTB (Methane To Benzene) reaction is one of arevolutionary process for direct conversion of natural
gas to petroleum compounds, i.e., from methane tobenzene. Mo-modified H-MFI is a typical catalyst withhigh activity for MTB. During the methane
dehydroaromatizaiton, Mo2C-like active carbide
species are formed on H-MFI catalysts. However,deactivation of MTB reactivity can not be avoided
because of carbon contamination due to excess
dehydrogenation of methane. In this case, structuralchange of active Mo species should be occur during
the deactivation [1]. For highly active and durable
reactivity for MTB reaction, detail of local structure ofactive Mo species and their structural changes needs tocharacterize. In the present study, Mo LIII-edge
XANES study has been applied to characterize theactive Mo species on H-MFI during the MTB reaction.
Catalyst samples were prepared by impregnation ofeach H-GaAlMFI (Si/Al2=40) support with
MoO2(acac)2−VO(C5H7O2)3−CHCl3 solution (5wt%
as MoO3), and followed by dried and calcined at 773
K. The H-GaAlMFI (Al/Ga=100) supports weresynthesized hydrothermally at 413 K for a week, andfollowed by ion-exchanging with NH4Cl and calcinedat 873 K. The catalytic MTB reactivity was evaluatedby means of fixed-bed flow reaction at 1023 K, as
described in a separate paper [1]. 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. REX-2000 (Rigaku Co,) software
was used for normalization of each XANES spectrum.
In our previous study, Mo-V (Mo/V=40) comodified H-GaAlMFI (Al/Ga=100) catalyst showshigher activity and lower deactivation rate than Mo/HMFI. In the case of Mo/H-MFI, partially carbonized
Mo (Mo2+ and Mo4+) species can be seen after CO
pretreatment. Figure 1 shows the LIII-XANES spectraof reduced MoO2 by CH4(10-20%)-H2 treatment at 973
K. Metallic Mo are formed, as well as the reduction
with H2. In this case, formation of Mo-carbide species
such as Mo2C can not be seen. In contrast, Mo2C-likecarbide species seems to be formed on the Mo-V/HGaAlMFI catalyst by CO(2%)-H2 treatment at 1023 K,
as shown in Fig. 2. It suggests that Mo2C-like carbide
species are formed before the MTB reaction by CO
pretreatment. After the MTB reaction, metallic Mo
species are also seen during the reaction for 10-200
minutes. For Mo/H-MFI, formation of α-Mo2C species

Ⅲ-2

Fig. 1. Mo LIII-edge XANES of reference Mo
compounds (Mo metal and α-Mo2C) and reduced
MoO2 with CH4-H2 atmosphere at 973 K.

Fig. 2. Mo LIII-edge XANES of Mo-V/H-GaAlMFI
catalysts after MTB reaction for 0-200 minutes at 1023
K.
[1] H. Aritani, H. Shibasaki, H. Orihara and A.
Nakahira, J. Environm. Sci. 21 (2009) 736.

49

161BL2A
山本知之_CAP.docx
Ⅲ -2. Materials
Sciences

BL2A

X-Ray Absorption Near-Edge Structure Analysis of Yb Ions in Er-Doped
CaZrO3
A. Kanemoto1 and T. Yamamoto1,2

Faculty of Science and Engineering, Waseda University, Tokyo 169-8555, Japan
Institute of Condensed-Matter Science, Waseda University, Tokyo 169-8555, Japan1

2

Ca1 - xErxZr1-yYbyO3
y = 0.01
Intensity (arb. units)

Ⅲ-2

Phosphor materials have been extensively studied
because of their importance in industrial applications.
Among them, rare-earth doped oxides with perovskite
structures are very attractive inorganic phosphors,
because of their luminescence efficiency and stability.
In addition to such conventional phosphors,
upconversion type phosphors, which can emit photons
with higher energy (shorter wavelength) by irradiationof lower energy (longer wavelength) photons, havebeen also actively investigated. It was reported that Er
doped CaZrO3 shows such upconversion luminescenceby irradiation of approx. 1000 nm photons to produce
green emissions and additional Yb doping enhances
luminescence intensity of Er-doped CaZrO3 [1]. In
order to understand the upconversion mechanism andthe reason why Yb doping could enhance emission
intensity, it is essential to know local environment of
Yb ions in Er-doped CaZrO3. In the current study, we
tried to analyze electronic structure of dilute Yb ionsin Er-doped CaZrO3 by Yb-M X-ray absorption nearedge structure (XANES) measurements.
Samples were synthesized by the solid state reaction
method. Reagent grade high purity powders of CaCO3,
ZrO2, Er2O3 and Yb2O3 were weighed changing theconcentration of Yb ions. These powders were mixed
and ground in an agate mortar for 30 min, which werepressed into pellet form. Resultant pellets weresintered in air at 1673 K for 6 hrs.
Crystal structure of the synthesized specimens was
examined by powder X-ray diffraction. Although tiny
amount of secondary phases appears, majority of thesample specimens could be determined to be
crystallized in orthorhombic perovskite structured
CaZrO3.
Yb-M XANES measurements were performed at
BL2A of UVSOR. Incident beams weremonochromatized
with
InSb
double-crystal
monochrometer. Sample powders were placed upon
the carbon adhesive tape, which are attached on first
Cu-Be dinode of the electron multiplier.
Observed Yb-M5 XANES spectra of Yb codoped
Er-CaZrO3 were compared with that of Yb2O3, which
are shown in Fig. 1. As shown in this figure, finestructures of Yb-M5 XANES spectra of the samples
here synthesized, i.e., Yb doped CaZrO3:Er, are almost
identical
withoutanydependencies
upon
concentration of Yb ions in Er-doped CaZrO3. In
addition, those show similar profiles as that of Yb2O3.
Hence it can be concluded that the doped Yb ions aretrivalent state from the current XANES analysis.

y = 0.03
3+

Yb2O3(Yb )1496 1498 1500 1502 1504
Energy (eV)Fig. 1. Observed Yb-M5 XANES spectra of Ybdoped CaZrO3:Er and that of Yb2O3.
[1] V. Singh et al., J. Appl. Phys. 112 (2012) 063105.

50

Ⅲ -2. Materials Sciences

171BL2A 中平敦.docx

BL2A

Structural Evaluation of AlPO-5 Synthesized by Hydrothermal Processwith XAFS
A. Nakahira1, 2, M. Togo1, H. Takahashi1, K. Yoshikawa1, M. Sato3 and H. Aritani4
Faculty of Engineering, Osaka Prefecture University, Sakai 599-8531, Japan
2
Kansai Center, IMR, Tohoku University, Sakai 599-8531, Japan
3
Institute of Materials Research, Tohoku University, Sendai 980-8577, Japan
4
Saitama Institute of Technology, Fukaya 369-0293, Japan

1

ALPO-5 is one of the most popular aluminum
phosphates (AlPO4), although there are many ofaluminum phosphate, generically known as "AlPO-n".
These aluminum phosphates, AlPO-n, have the uniqueframework structures similar to zeolites. Also, they
have unique microporous structures and are expected
to be used as a catalyst and a molecular sieve [1-2].
Recently these microporous aluminum phosphates
have attracted much attention, because AlPO-nframework can be incorporated with cationic species
from monovalent to pentavalent. Also, the syntheses ofheteroatom-doped AlPO-n were attempted to modify
the microporous structures and enhance theirproperties through a variety of processing. For
example, molecular-sieve catalysts of AlPO aresignificantly useful for the selective oxidation of linear
alkanes, such as the catalytic oxidation in air ofcyclohexane [3]. In this study, AlPO-5 were focusedon among AlPO-n families. In special, the synthesis for
AlPO-5 was attempted by hydrothermal process andthe structural evaluations were carried out for samples
by hydrothermal process.
AlPO-5 were synthesized by the hydrothermal
method. Al(OH)3 and aluminum alkoxide as analuminum source and phosphoric acid as a phosphorus
sourcewereused.
Furthermore,
TEAOH(tetraethylammonium hydroxide) were used as anorganic template. The molar composition of thehydrothermal
synthesismaterials
was
Al:P:TEAOH:H2O = 2:2:X:250 (X=3~6).
These starting reagents were mixed under magnetic
stirring and pH was maintained to be pH3 with H2SO4
solution. After aging for 1h, starting materials
solutions were transferred into teflon-lined stainlesssteel autoclave and carried out by the hydrothermal
treatment at 423 K to 463 K for 0 h to 48 h. After thehydrothermal treatments, obtained samples werecollected with Buchner filter and washed with
sufficient deionized water. And after drying at 323 K
overnight, samples were crushed with agate mortar.
The dried samples were characterized by X-ray
diffraction analysis (XRD) and FT-IR and TG-DTA.
Microstructures were observed with FE-SEM. The
local structures around Al-K for the obtained products
were characterized by measuring X-ray adsorption
near edge structure (XANES) at BL2A in UVSOR.
Various aluminum phosphates were synthesizedthough hydrothermal treatment under a variety ofhydrothermal conditions (hydrothermal temperatures,
synthetic time, various pressures, and various starting
materials).
The evaluation of X-ray diffraction analysis
revealed that products obtained by the hydrothermal
process at 443 K for 12 h to 48 h were identified to be
crystalline AlPO-5 without another phase. SEM
observation showed that products were rectangularwith approximately 5 μm in length and 500 nm indiameter. Figure 1 shows the results of Al-K XANES
spectra of AlPO-5 samples synthesized by thehydrothermal treatment for 0 h to 12 h. From theseresults, the local structure of Al-K XANES forAlPO-5
synthesized by the hydrothermal treatment for 12 h
was different from ones for AlPO-5 samples
synthesized by the hydrothermal treatment for 0 h to 8
h, suggesting that AlPO-5 local-structure was
successfully synthesized by the hydrothermal
treatment for 12 h.

Normalized Intensity (a.u.)

Al K-edge

1540

(a)
(b)
(c)
(d)

1560

1580

1600

1620

1640

Photon Energy / eV

Fig. 1. Al-K XANES spectra of AlPO-5 synthesized
by the hydrothermal treatment for 0 h to 48 h. (a) 0h,(b) 4h, (c) 8h, and (d) 12h.
[1] D. R. Pyke, P. Whitney, and H. Houghton, Applied
Catalysis, 18 (1985) 173.
[2] D. Li, J. Yao and H. Wang, Progress in Natural
Science: Materials International 22 (2012) 684.
[3] J. M. Thomas, R. Raja, G. Sankar and R. G. Bell,
Nature 398 (1999) 27.

51

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181BL2A
小林英一.docx
Ⅲ -2. Materials
Sciences

BL2A

In Situ Observation of the Dielectric Breakdown of Insulating Material
Using Soft X-Ray Absorption Spectroscopy
E. Kobayashi1, K. K. Bando2, K. K. Okudaira3 and T. Okajima1

Kyushu Synchrotron Light Research Center, Tosu, 841-0005, Japan
National Institute of Advanced Industrial Science and Technology, Tsukuba, 305-8565, Japan
3
Chiba University, Chiba, 263-8522, Japan
1

Fig. 1. Experimental setup.
4.5

1.0

4.0
3.5
Intensity (arb. units)

Ⅲ-2

Insulating materials are used in various parts ofelectric device. In recent years, it has become
necessary to develop excellent insulating material withimprovement in performance. In this study, theelectronic state on aluminum oxide thin film undervoltage applied conditions was studied by using nearedge X-ray absorption fine structure (NEXAFS).
Aluminum oxide is recognized as an importantmaterial because it is the prototype of wide-gap
insulator. The elementary processes of dielectric
breakdown of insulating materials are not understoodvery well from a microscopic point of view.
An alumina film was made on a polished stainlesssteel plate by dip-coating with a diluted alumina sol(Kawaken Fine Chemicals, F1000). The coated plate
was dried at 333 K, 6 hours and calcined at 573 K, 3
hours in an oven.
Al K-edge NEXAFS spectra of aluminum oxide thinfilm were measured at the beamline 2A of the UVSOR.
The experimental setup for NEXAFS measurement isshown in Fig. 1. The soft x-ray was passed through ametal mesh and was irradiated to the sample. The
fluorescence from sample was detected by silicon drift
detectors. Potential of the sample was applied between
the metal mesh and metal plate. All experiments were
performed at room temperature.
Figure 2 shows the Al K-edge NEXAFS spectra ofaluminum oxide obtained from PFY mode at various
applied voltage. The spectra were measured while
applying a voltage. This sample was breakdown at ~14 V. After breakdown, the potential to the sample
could not be applied. These spectral features are almost
same. And the spectra show two strong peaks at around1569 eV and 1572 eV, respectively. These spectralfeatures are also different from those of Al2O3 powder[1]. This indicates the local structure of Al inaluminum oxide thin film using in this study isdifferent from those in Al2O3 crystal.
Inset of the Fig. 2 shows the expanded spectra at
rising part of the spectra. The spectra with underapplied voltages of 8 V show little energy shift. But the
spectra at over 9.9 V are shifted slightly to the low
energy side. Then, the spectrum after the breakdown
was shifted significantly to the low energy side. This
trend was also observed for an aluminum oxide thinfilm calcined at 473 K.

Intensity (arb. units)

2

0.5

0.0
1562

3.0

1563

1564

1565

1566

1567

Photon Energy (eV)after breakdown
V=0V
V = 12 ~ 8.5 V

2.5

V = 9.9 ~ 8.6 V
V=8 V

2.0

V=6V
V=4V

1.5

V=2V
V=0V

1.0
0.5
0.0
1560

1570

1580

1590

1600

Fig. 2. Al K-edge NEXAFS spectra of aluminumoxide thin film obtained from PFY mode at various
applied voltage.
[1] M. Mogi et al., Mat. Trans. 45 (2004) 2031.

52

Ⅲ -2. Materials Sciences

191BL3B 小玉展宏.docx

BL3B

Material Dependence of Multiple Luminescence from Self-Trapped
Excitons in Orthoborates with Complex Oxoabions upon Vacuum UV
Excitation
N. Kodama, K.Yanagidaira, M. Kudo and T. Takahashi

Graduate School of Engineering and Resource Science, Akita University, Akita 010-8502, Japan53

decay times of the STE luminescence, as follows. The
fast decay components of STE(I) and STE(II) in LSB
and CLOB and the fast component of STE(I) in YSB
originate from singlet STE states in the BO33- groupwith C1 or Cs symmetry, respectively. In contrast, theslow components of STE(II) and STE(III) in LSB havebeen ascribed to a triplet STE with C1 symmetry andthose of CLOB have been ascribed to a triplet STEwith Cs symmetry. Furthermore, the slow componentof STE(II) in YSB is thought to originate from a triplet
STE with C1 symmetry.
The excitation spectra of LSB, CLOB, and YSB
shown in Figs. 3(a), (b), and (c), measured at 293 K, 8
K, and 293 K, respectively, were detected at various
fixed wavelengths. The excitation spectra of LSB
increased stepwise in intensity with increasing
excitation energy; that is, with decreasing excitation
wavelength. The features of the excitation spectra
below the band edge are very similar to those of CLOB
and YSB. The excitation intensities at 90 nm for LSB,LCOB, and YSB were two to five times greater than at
180 nm, near the band edge. These results suggestthat two or more electron–hole pairs are produced fromone absorbed photon when excited by photons withenergies greater than or equal totwice the band gap energy; that is,
it may be possible to generate
double- or multiple-self-trapped
excitons (double- or multipleSTEs) in LSB, CLOB, and YSB.

10
5
0
70 100

50

40
30
20

10

150 200
250
Wavelength / nm
Fig.3(a)0200 25070 100 150
Wavelength / nm
Fig.3(b)30

Emission Intensity /a.u.(×10-3)

20

15

Emission Intensity /a.u.(×10-3)

Emission Intensity /a.u.(×10-3)

We report the temperature and material dependenceof multiple intrinsic luminescence spectra from selftapped excitons(STEs) in the three low-symmetry
orthoborates with complex oxoanions : LaSc3(BO3)4(LSB), Ca4LaO(BO3)3 (CLOB), and YSc(BO3)2(YSB).
Furthermore, we examine the possibility of generating“multiple self-trapped excitons”(MSTE) in LSB,
CLOB, and YSB crystals upon excitation by photons
with more than twice the band gap energy, on the basisof STE excitation spectra in the VUV region.
Absorption spectrum of an LSB single crystal isshown in Figure 1. The absorption increased rapidly
below about 200 nm. This result indicates that the bandedge of LSB is around 200 nm. Moreover, the rise
wavelength (rising) of absorption was in agreementwith that of excitation. On the basis of this result, therises around 200 nm were assigned to the band edges
of LCOB and YSB, similar to that of LSB.
Figures 2(a), 2(b), and 2(c) show the temperaturedependence of the intrinsic luminescence spectra for
LSB, LCOB, and YSB, respectively, at 11 or 8 K, The
luminescence spectra exhibited multiple bands of
STE(I) and (II) (and/or (III)) for LSB, LCOB, and YSB,
respectively, located at energies corresponding to the
STE bands. The luminescence spectra for LSB under
VUV excitation contained three distinct broad bands
associated with STE(I), (II), and (III), with peaks at
245–272 nm, 311–320 nm, and 386–389 nm,
respectively, in the range of 11–293 K, as shown in Fig.
1(a). For CLOB and LSB, three distinct broad bands(two weak bands and one intense band) appeared.
These bands are associated with STE(I), (II), and (III),with peaks at 245–273 nm, 326–328 nm, and 410–413
nm, respectively, in the range of 8–293 K, as shown in
Fig. 2(b). YSB exhibited one intense broad emission
band, with the shortest wavelength of the three samples,
associated with STE(I) with a peak at 247–254 nm, andanother broad band associated with STE(II) with apeak at 290–296 nm in the range of 8–293 K
As the STE(I) emission intensity decreased in LSB
and YSB, the STE(II) and STE(III) emission
intensities increased. This suggests that there are
multiple (three in LSB, two in YSB) local minima on
the adiabatic potential energy surfaces, i.e., barriers for
exciton self-trapping in LSB and YSB.
The decay times for the luminescence in LSB at 78
K were 20 ns for STE(I), 230 ns and 4.5 μs for STE(II),and 7.4 μs for STE(III). The decay times for the
luminescence in CLOB at 78 K were 24 ns for STE(I),540 ns and 2.0μs for STE(II), and 6.0 μs for STE(III).
For YSB, the decay times for STE(I) and (II)luminescence at 293 K were 23 and 643 ns,
respectively. Taking into account the above results,
the site symmetry of the BO33- oxoanion group, and the
B-O bond length, we considered the structures and spinstates of the STEs in LSB, CLOB, and YSB crystals.
The existence of multiple stable structures and spinstates for STEs in these crystals can be inferred fromthe temperature dependence of the intensities and

20

10
0
200
70 100 150
Wavelength / nm
Fig.3(c)250

Fig. 1. The absorption spectrum of an LSB single
crystal at 293 K.
Fig. 2. Temperaturedependenceof
intrinsic
luminescence spectra in (a) LSB, (b) CLOB, and (c)YSB in the range of 11(8)– 293 K.
Fig. 3. Excitation spectra obtained by monitoring theintensity of the intrinsic luminescence at different
fixed wavelengths in (a) LSB at 293 K, (b) CLOB at 8
K, and (c) YSB at 293 K.

Ⅲ-2

201BL3B
黒澤俊介_CAP.docx
Ⅲ -2. Materials
Sciences

BL3B

Transmittance Spectra and Other Properties of Metal Elements Doped
LiCaAlF6 Single Crystals
C. Tanaka1, S. Kurosawa2, Y. Yokota2, A. Yamaji1, R. Murakami1, T. Horiai1, H. Chiba1,
Y. Ohashi1, K. Kamada2,3 and A. Yoshikawa1,2,3
Institute for Materials Research, Tohoku University, Sendai 980-8577, JapanNew Industry Creation Hatchery Center (NICHe), Tohoku University, Sendai 980-0845, Japan3
C&A corporation, Sendai 980-8579, Japan1

2

crystal, and were ascribed to self-trapped and trapped
exciton [2]. In this wavelength region, an emission
peak around 205-210 nm was observed and ascribed to
Pb2+ center consistently with the previous report about
the Pb doped LiCAF crystal [3]. In addition, we
revealed that there was another emission peak around
830 nm in the long wavelength region. This emission
peak can originate from defect in the structure due tononstoichiometry of LiCAF lattice caused by Pb
evaporation, or Pb+ center like in the LiBaF3 latticewhich has the emission around 880 nm in [4].
10080Transmittance / %Ⅲ-2

Neutron detectors using neutron scintillators havebeen investigated to find their application in thehomeland security scanning. A 3He-gas proportionalcounter has been used widely as a conventional sensor
for the thermal neutron detection due to the highcapture cross-section for the thermal neutrons and thelow sensitivity to γ-rays. However, the importance ofalternative neutron detectors such as scintillators has
increased due to the supply crisis of 3He gas by theexcessive demand. Therefore, the neutron scintillators
have become gradually of interest as an alternative forthe 3He gas.
In the search for novel neutron scintillators, we havedeveloped a LiCaAlF6 (LiCAF) scintillator crystal
containing 6Li with the high capture cross-section for
the thermal neutrons. In addition, the effective atomic
number and density of the LiCAF are relatively low
which results in the desired small detection efficiency
for γ-ray.
The ns2-type cations such as In+ and Pb2+ are
expected to show luminescence. In this study, we grew
LiCAF crystals doped by the ns2-type elements andinvestigated their properties to obtain a new candidate
for the doping ion in the LiCAF crystal.
Non-doped and In, Pb-doped LiCAF crystals withthe diameter of 2 mm were grown by a micro-pullingdown (μ-PD) method. We measured transmittancespectra of the polished specimens of non-doped and In,
Pb doped LiCAF by the SI-Photodetector (AXUV 100(IRD)) in the wavelength range from 80 to 300 nm at
liquid helium temperature at BL3 of UVSOR.
Transmittance spectra of the Non-doped and In, Pb
doped LiCAF are shown in Fig. 1. The spectrum of Indoped LiCAF has no remarkable difference from nondoped one except for the onset of the transmittancedrop near the band-edge which is shifted towards
longer wavelengths. This may be caused by some
defects created by the doping. In the transmittancespectra of Pb doped LiCAF specimen an intense
absorption peak was observed at 196 nm followed byother absorption peaks at shorter wavelengths and theband edge of LiCAF host at about 115 nm. Absorption
peaks at 196 and 135 nm could be ascribed to the socalled A and C absorption bands of the Pb2+ center [1].
Radioluminescence spectra of non-doped and Pb
doped LiCAF crystals under X-ray irradiation were
measured at RT. The emission peaks around 290, 380
and 510 nm were observed for non-doped LiCAF

60
40

In 0.5%
Pb 0.5%Non-doped

20
0

100

150
200
Wavelength / nm

250

300

Fig. 1. Transmittance spectra of non-doped and In, Pb
doped LiCAF in 80 – 300 nm at liquid heliumtemperature.
[1] P. W. M. Jacobs, J. Phys. Chem. Solids 52 (1991)
35.
[2] M. True, Y. Chen, M. Kirm, S. Vielhauer and G.
Zimmerer, J. Lumin. 124 (2007) 279.
[3] A. Novoselov, J. Pejchal, M. Nikl and A.
Yoshikawa, Nucl. Sci. Symp. Conf. Rec. N45-7 (2007)
2184.
[4] L. Prado, N. D. Vieira Jr, S. L. Baldochi, S. P.
Morato, J. P. Denis, N. Tercier and B. Blanzat, J. Phys.
Chem Solids 57 (1996) 413

54

211BL3B 黒澤俊介_CAP.docx

Ⅲ -2. Materials Sciences

BL3B

Charge Transfer Luminescence of Ce-Doped and Undoped Ytterbium
Pyrosilicate
T. Horiai1, S. Kurosawa2, R. Murakami1, A. Yamaji1, Y. Shoji1, 3, J. Pejchal4, Y. Ohashi1, 3,
K. Kamada2, 3, Y. Yokota2 and A. Yoshikawa1, 2, 3
Institute for Materials Research, Tohoku University, Sendai 980-8577, JapanNew Industry Creation Hatchery Center (NICHe), Tohoku University, Sendai 980-8579, Japan
3
C&A Corporation, Sendai 980-8579, Japan
4
Institute of Physics CAS, Cukrovamicka 10, Prague, 162 00, Czech Republic
1

2

Charge transfer (CT) luminescence of Yb3+ can havefast scintillation decay time, and such scintillationmaterials are expected to be applied to medical
imaging. In previous study, the charge transfer
luminescence of Ce-doepd Yb2SiO5 was evaluated,
and the intensity became highest at near 40 K [1,2]. In
this study, we search the CT luminescence forpyrosilicate group as the first step.
We prepared Ce-doped and undoped Yb2Si2O7
single crystals grown by the micro-pulling-down
method, and photoluminescence spectra for thesesamples were evaluated within 7.5-300 K at BL3 of
UVSOR. The emission spectra were examined with aspectrometer (spectoropro-300i, Acton research).
Figures 1 and 2 shown the temperaturedependences of the PL emission spectra for the
undoped and Ce-doped Yb2Si2O7 single crystals under
210 nm excitation in the temperature range of 7.5300K, respectively. Since absorption edge was around230 nm, the 210 nm photons were selected as anexaction source. Moreover, variations of the excitation
wavelengths demonstrated that the strongest emission
was observed at 210 nm excitation.
The undoped Yb2Si2O7 had emission peaks ataround 420 and 580 nm. The intensity of the 420 nm
emission band was degraded when temperature of thespecimen exceeded 100K, while that of the 580 nm
band remained almost constant and independent on
temperature. Thus, 420 nm band would be originatedby CT system to 2F5/2 of Yb3+, and thermal quenching
was observed. On the other hand, the emission at 580
nm may originate from defects.
Oppositely, the intensity of the 420 nm emission
band of the Ce-doped Yb2Si2O7 increased whentemperature of the specimen became higher. Mainly
this peak was assigned to Ce3+ 5d-4f emission, andCTS to 2F5/2 of Yb3+components can be included. 580
nm emission peak was also observed in the PL spectraof Ce-doped Yb2Si2O7, while this emission was
suppressed at RT. Moreover, from 200 to 300 K, Ce3+
emission intensity increased. Thus, the electron can bere-trapped into Ce3+ absorption band, and this result
suggested energy transfer from this defect site to Ce3+.

Ⅲ-2

Fig. 1. The photoluminescence of undoped YbPS
excited by 210 nm within 7.5-300 K.

Fig. 2. The photoluminescence of Ce-doped YbPS
excited by 210 nm within 7.5-300 K.
[1] L. V. Pieterson et al., J. Luminescnece 91 (2000)
177.
[2] D. W. Cook et al., J. Luminescence 79 (1998) 185.

55

221BL3B
黒澤俊介_CAP.docx
Ⅲ -2. Materials
Sciences

BL3B

Evaluation of the Bandgap Energies for Garnet-Type Infra-Red
Scintillators
A. Yamaji1, S. Kurosawa1, 2, R. Murakami1, T. Horiai1, M. Kitaura3, Y. Ohashi1, K. Kamada2,
Y. Yokota2, A. Ohnishi3 and A. Yoshikawa1,2Institute for Materials Research, Tohoku University, Sendai 980-8577, JapanNew Industry Creation Hatchery Center, Tohoku University, Sendai 980-8579, Japan
3
Yamagata University, Yamagata 990-8560, Japan
1

2

Ⅲ-2

Inorganic scintillator crystals have been widely used
for radiation detection such as high energy physics,
security applications and medical imaging. Recently,
we have studied the real-time dose monitor system
during the radiation therapy using scintillator with
infrared emission. Radiation therapy is one of the
primary treatments of tumor; however, radiation
overdoses which cause medical accidents or underdoses which cause incomplete treatments have often
been reported. Thus, real-time dose monitoring
systems are required.
One of the concepts for the monitoring is the use of
a 2-3 mm size scintillator material in the irradiated area(tumor) in the patient’s body [1]. When the scintillator
is irradiated with the incident X-rays, the intensity ofthe scintillation light give us information related to thedose in real time. Since human body had a low
absorption region from approximately 700-1400 nm[2], scintillator with near infrared emission could beapplied in the real time dose monitor system. Thus, wehave investigated the optical and scintillation
properties of some materials in the near infra-red
region.
In this work, we grew garnet-type oxide crystals for
scintillators with near infrared emission and evaluatedtheir optical properties. Since the light yield depends
on the band gap energy of a host crystal, we estimatedthe band-gap energies of the studied crystals. Band gap
measurement is useful for the material design ofscintillator.
The crystal growth of garnet-type oxide crystals
were performed from the melt by the micro pulling
down method using an Ir crucible with a radio
frequency heating system. Nominal compositions ofgrown crystals were Gd3Ga5O12 (GGG), Lu3Ga5O12(LGG), Y3Ga5O12 (YGG) and (Gd0.9M0.1)3
Ga4.7M’0.3O12 (M = Mg, Ca, Sr and Ba, M’ = Zr, Hf),
respectively. The crystals were mechanically polishedto the 1 mm thickness.
Transmittance spectra were measured at the beamline BL3B at the UVSOR facility; The specimens wereglued on holes (1.5 mm diameter) of a copper platewith a silver paste (Fujikura Kasei, DOTITE D-550),
and their temperature was kept at 8 K with liquid
helium in order to suppress lattice vibrations.
Transmitted light was detected with a silicon
photodetector (IRD, AXUV 100).
The transmittance spectra were obtained as shownin Fig.1. Band gap energies were estimated from theabsorption edges of these transmittance spectra. The
calculated bandgaps are listed in table 1. These results
contribute to the material design of scintillator withnear infrared emission

Fig. 1. Transmittance spectra of YGG, LGG and
GGG crystals at 8 K.
Table 1. Bandgaps of garnet-type oxide crystals.

[1] E. Nakata et al., Presentation at RadiologicalSociety of North America (2008).
[2] Y. Y. Huang et al., Dose-Response (Prepress)
Formerly Nonlinearity in Biology, Toxicology, andMedicine (2009).

56

231BL3B 黒澤俊介.docx

Ⅲ -2. Materials Sciences

BL3B

Determination of the Band-Gap Energy of (Gd, La)2Si2O7 Scintillator
Crystals Using the UVSOR Facility
R. Murakami1, S. Kurosawa2, A. Yamaji1, M. Kitaura4, T. Horiai1, Y. Shoji1,3,
Y. Ohashi1, Y. Yokota2, K. Kamada2,3, A. Ohnishi4 and A. Yoshikawa1,2,3
Institute for Materials Research, Tohoku University, Sendai, JapanNew Industry Creation Hatchery Center, Tohoku University, Sendai, Japan
3
C&A Corp., Sendai, Japan
4
Yamagata University, Yamagata, Japan
1

2

Inorganic scintillating materials are widely used forthe medical imaging, oil well logging, astronomy, etc.
Silicate scintillators such as Ce:Gd2SiO5 (Ce:GSO)
and Ce:(Lu,Y)2SiO5 (Ce:LYSO) are applied in thesefields, due to their high light outputs, short decay times,
and so on. For example, Ce:GSO, has a good light
output of ~12,500 photons/MeV and short scintillation
decay time of ~56 ns [1]. Recent studies demonstrated
that pyrosilicate-type crystals including Ce:Gd2Si2O7(Ce:GPS) and Ce:Lu2Si2O7 (Ce:LPS) are also very
promising scintillating materials [2,3]. Their light
output exceeds that of Ce:GSO over two times [4].
Heavily Ce-doped (~10%) GPS crystals can be
successfully grown from melt, and have the diameterover 1 inch. On the other hand, undoped GPS crystals
cannot be grown from melt, because this material melts
incongruently. However, heavily Ce-doped GPS
crystals are not favorable, because high Ceconcentration would lead to the degradation of lightyield due to concentration quenching and selfabsorption. In contrast, Ce:(Gd,La)2Si2O7 (Ce:LaGPS) does not demonstrate the quenching problems.
In Ce:La-GPS, part of Ce3+ is easily substituted withinactive La3+, because both Ce3+ and La3+ have similar
ionic radius [5].
Scintillation properties of Ce:La-GPS have been
investigated using the crystals grown by the floatingzone (FZ) method [5,6]. La 9% substituted crystal
showed high light output (~36,000 photons/MeV) andhigh energy-resolution (~5.0%, 662 keV, FWHM) [5].
The FZ-grown crystals were a few mm in dimensions.
Alternatively, the Ce:La-GPS crystals grown by the
Czochralski (Cz) process were over 1 inch in diameter.
It was also demonstrated that 25 - 50% La-substituted
GPS crystals can be repeatedly grown by the Cz
method [7].
Such pyrosilicate-type scintillators show high light
output at high temperature. The light output of Tl:NaI,
which have been used for oil well logging degrade toroughly 80% at 150C [8], while the light output of theCe:La-GPS does not almost degrade at temperature
over 150C [9]. One of the important keys is thetolerance for the photoionization, and it depends on theenergy difference between the top of 5d level andconduction band. However, band-gap energy of LaGPS has not been reported; therefore, there are few
guides for the material design.

We have succeeded the growth of undoped La-GPS
with the chemical composition of (Gd,La)2Si2O7
single crystals by the micro-pulling-down (μ-PD)
method. The crystals with 1 mm thickness were cut
from the ingots and were polished. The band-gap
energies were determined at the BL3B beamline in the
UVSOR facility. The transmittance spectra were
measured at low temperature (~5 K) with the liquid He
under ultra-high vacuum (~10-7 Pa). The band-gap
energies were estimated to be 7.1 -7.2 eV.
[1] K. Takagi and T. Fukuzawa, Appl. Phys. Lett., 42(1983) 43.
[2] C. L. Melcher et al., Delft Unviersity Press(SCINT95), ISBN 90-407-1215-8, (1996) 309.
[3] S. Kawamura et al., IEEE Nuclear Sci. Symp. Conf.
Rec. (2006) 1160.
[4] S. Kawamura et al., IEEE TNS 54 No.4 (2007)
1383.
[5] A. Suzuki, A. Yoshikawa et al., Appl. Phys. Express
5 (2012) 102601.
[6] S. Kurosawa et al., NIMA 744 (2014) 30.
[7] A. Yoshikawa et al., Cryst. Growth Des., 15 (2015)
1642.
[8] Saint Gobain, NaI(Tl) Data Sheet,
http://www.crystals.saintgobain.com/uploadedFiles/SGCrystals/Documents/NaI(Tl)%20Data%20Sheet.pdf[9] S. Kurosawa et al., NIMA 772 (2015) 72.

57

Ⅲ-2

241BL3B
黒澤俊介.docx
Ⅲ -2. Materials
Sciences

BL3B

Evaluation of Ce:(Gd, La)2Si2O7 Scintillator Using an UVSOR Beam Line
S. Kurosawa1, R. Murakami2, Y. Shoji2,3, A. Yamaji2, T. Horiai2, Y. Ohashi2,
Y. Yokota1, K. Kamada1,3 and A. Yoshikawa1,2,3
Institute for Materials Research, Tohoku University, Sendai, JapanNew Industry Creation Hatchery Center, Tohoku University, Sendai, Japan
3
C&A Corp., Sendai, Japan
1

2

Ⅲ-2

Scintillators are used in various fields such as
medical imaging, astronomy and so on. Evaluation oftemperature dependences of some scintillation
parameters such as light output and decay time isimportant to understand if the scintillation material issuitable for applications in low or high temperature
conditions.
In oil well logging, scintillator is exposed totemperatures above 350 K measuring the subsurface
geologic formation properties. Therefore, scintillator
with a high light output even at high temperatures isrequired.
Recently, scintillation properties of (Ce0.01, Gd0.90,
La0.09)2Si2O7 (Ce:La-GPS) have been reported. The
Ce:La-GPS crystal had a light output of approximately
35,000 photons/MeV, FWHM energy resolution of
5.0% at 662 keV [1,2]. Moreover, the light output over
35,000 photons/MeV was found constant in thetemperature range from 0 to 423 K. In addition,
FWHM energy resolution of Ce:La-GPS (roughly 7–8%) at 662 keV remained constant up to 373 K [3].
Thus, this crystal can be applied to oil well logging orother radiation detection application at hightemperature conditions.
In this study, we evaluated luminesce properties of
Ce:La-GPS in low temperature in BL3B of UVSOR.
The luminescence was measured with a spectrometer(spectoropro-300i, Acton research) consisting of a
CCD camera from approximately 5 – 300 K.
Figure 1 shows the emission intensity around 390
nm originated from 5d-4f transition of Ce3+ as functionof temperature, where excitation wavelength wasapproximately 240 nm originated from 4f-5d4 translonof Ce3+. This result showed that the luminescenceintensity remained constant, not only up to 400 K, also
down to 5 K.
Compared with other scintillation materials, wefound the invariable properties is rare case; in general,
intensities of several scintillators doped with Ce3+
increase as temperature fall or disappear at low
temperature (5-100 K). Thus, we found Ce:La-GPS
has “strange” and “unique” temperature dependence.
We also measured thermo-luminesce for thismaterial in UVSOR (BL3B), and we found that thismaterial had small number of traps between valence
and conduction band compared with conventional
scintillators. Some traps would suppress the light
output (luminance intensity), and the suppression fromthe traps are related to temperature. Thus, the small
number of traps would be related to be “strange” and“unique” temperature dependence.
Moreover, we also measured the bandgap energy of
La-GPS in UVSOR, and this material had wider gap
energy than conventional scintillators [4]. This widebandgap can be related to the “strange” and “unique”
temperature dependence in high temperature region(up to approximately 400 K), because, in this case, theenergy difference between the bottom of the
conduction band and excitation band of Ce3+ (5d1level) can be wider than conventional scintillators. Dueto the wide energy, even in high temperature, electron
located in excitation band of Ce3+ can de-excite to
grand state of Ce3+ (4f), and a scintillation photon is
generated. On the other hand, in the case of the low
difference, electrons tend to move from the excitation
band of Ce3+ to conduction band, and light output isquenched.

Fig. 1. Intensity of Ce3+ 5d-4f transition as a functionof temperature.
[1] A. Suzuki, S. Kurosawa, A. Yoshikawa et al.,Applied Physics Express 5 (2012) 102601.
[2] S. Kurosawa, A. Yoshikawa et al., Nucl. Inst. andMeth. in Phys. Res. A 774 (2014) 30.
[3] S. Kurosawa, A. Yoshikawa et al., Nucl. Inst. andMeth. in Phys. Res. A 772 (2015) 72.
[4] R. Murakami and S. Kurosawa, UVSOR Activity
Report, submitted.

58

251BL3B 黒澤俊介.docx

Ⅲ -2. Materials Sciences

BL3B

Thermoluminescence of SrHfO3 Ceramics
1,

H. Chiba S. Kurosawa1,2, R. Murakami1, A. Yamaji1, Y. Ohashi1, J. Pejchal1,3, K. Kamada2,4,
Y. Yokota1,2 and A. Yoshikawa1,2,4
Institute for Materials Research, Tohoku University, Sendai 980-8577,Japan2
New Industry Creation Hatchery Center(NICHe), Sendai 980-8579,Japan
3
Instiute of Physics, AS CR, Cukrovarnická 10,162 53 Prague ,Czech Republic
4
C&A Corporation, Sendai 980-8577,Japan
1

In Positron Emission Tomography (PET), Lutetium
based scintillators are typically used because of their
good performance and high density. However, Lucontaining scintillators have high-rate intrinsic
background of roughly 300 counts/s/cm3 because
natural abundance of 176Lu radio isotope (beta emitter)is 2.59%. This is why Lu-containing scintillators canhave worse sensitivity. Hafnium has a high atomic
number of 72, and the abundance of its only
radioisotope is less than 0.2% (174Hf, half-life is morethan 1015 year). Thus, a new scintillator containing
Hafnium can have low intrinsic background and highdetection efficiency.
Recently, SrHfO3 doped with Ce (Ce:SHO) has
attracted attention as a scintillating material due to ahigh density of 7.65 g/cm3, a high effective atomic
number of 64 and negligible intrinsic-background [1].
Al/Ce:SrHfO3 ceramics can be prepared by the Spark
Plasma Sintering(SPS) method due to the high melting
point of Hafnium [2]. It has improved opticalproperties and radiation response when compared to
the Al-free material.
Considering the above, we fabricated Al/Ce:SrHfO3
by the SPS method; For starting materials, HfO2
(99.95% pure), SrCO3 (99.99% pure), α-Al2O3(99.99% pure), CeO2 (99.99% pure) powders wereused. Ce:SHO ceramic samples were pre-sintered in achamber filled with air. SHO ceramic samples wereprepared from the SHO powder using a 10-mmdiameter carbon punch and 30-mm-diameter carbon
dies, SPS chamber (Fuji Denpa, DR.SINTER)
We used carbon punch and dies were used in thesintering process under vacuum, strongly reduction
environment is established. This leads to creation of
the oxygen vacancies in the material. These defects
decrease the light output of the samples, they need to
be investigated to optimize the material performance.
Thus, thermoluminescence method was employed.
In BL3B of UVSOR, we used thermoluminescencemeasurement from 4 K to room temperature for
SrHfO3 samples; The samples were irradiated at 220
nm for 10 minutes at very low temperature 4~6 K.
After the irradiation, we raised temperature slowly andmeasured the emission intensity with a CCD.
The thermoluminescence glow curve is shown in
Fig.1hawse can observe some peaks from 100 to 200K,
and these peaks may originate from Oxygen defects.
As for future plans, we will investigate how theannealing can diminish Oxygen defects concentration

using the same method.

Ⅲ-2

Fig. 1. Thermo Luminescence spectra of AlCeSrHfO3.
[1] E. Mihóková, V. Jary´, L. S. Schulman and M. Nikl,
Chemical Physics Letters (in press).
[2] S. Kurosawa et al., Radiat. Meas.56 (2013) 155.
[3] S. Blahuta et al., IEEE Transaction on Nuclear
Science NS 60 (2013) 3134.

59

261BL3B
小南裕子.docx
Ⅲ -2. Materials
Sciences

BL3B

Study of Emitting Mechanism on UV Emitting ZnAl2O4 Phosphor Using
Vacuum Ultra Violet Photoluminescence System
H. Kominami1, T. Ito1 and S. Kubota2

Graduate School of Integrated Science and Technology, Shizuoka University, Hamamatsu 432-8651 Japan
2
Faculty of Engineering, Shizuoka University, Hamamatsu 432-8651 Japan

CL Intensity [a.u.]

4 kV

200

400
600
800
E, 200 h annealed
D, 100 h annealed
B, 2 h annealed
A, as-depo

200

400
600
800
Wavelength [nm]Fig. 1. Cathodoluminescence of ZnAl2O4 thin films
prepared by thermal diffusion of ZnO and sapphire
substrate, excited by anode voltages of 4 and 10 kV
and sample current density of 30 μA/cm2.

1
PL Intensity (a.u.)

Ⅲ-2

In our previous work, it was found that ZnO-basedoxide materials showed UV emission under excitation
by electron beam. Especially, ZnAl2O4 showed strong
UV emission peaked around 250 nm. However, thebasic property and the emission mechanism of
ZnAl2O4 were not cleared. For the evaluation of thephysical properties, it is difficult to evaluate because itis the influence of surface scattering, such as theabsorption coefficient and refractive index reflectance.
By thinning the ZnAl2O4, we thought that it is possible
to evaluate these fundamental properties and themechanism of UV emission from ZnAl2O4. In thisstudy, ZnAl2O4 thin firm was prepared by sputteringand thermal diffusion to the sapphire substrate.
ZnAl2O4 thin films was prepared by thermal
diffusion of ZnO film to sapphire substrate. After
deposition, the films were annealed in air. From crystal
analysis, ZnAl2O4 polycrystalline was obtained.
Figure 1 shows CL spectra of the ZnAl2O4 thin films
before and after annealing for 2, 100 and 200 hours,
respectively, excited at anode voltage of 10 and 4 kV
and sample current density of 30 μA/cm2. Both films
showed UV emission from ZnAl2O4. From thespectrum of 2 hours annealing, in spite of ZnO thickfilm was still remained, the visible emission around
500 nm was not appeared. CL intensity of 100 hours
annealed film showed 25 times stronger than that of 2hours annealing. In other words, it suggested that thenot only increasing of thickness but also improvementof crystallinity of ZnAl2O4 film by longer annealing
time.
Figure 2 shows PL excitation spectra of ZnAl2O4
powders prepared by different ZnO/Al2O3 ratio. The
excitation band was shifted to higher energy when
Zn/Al ratio was decreased. It indicates that Zn/Al ratio
was influenced the optical properties. The red line in
the figure is the transmittance spectrum of ZnAl2O4
thin film. The absorption edge was closed to lower
ZnO/Al2O3 ratio. It seems that the film was notstoichiometry. Moreover, the refraction index anddielectric constant were not obtained from the analysisof transmittance and reflection spectra. It was thoughtthat the ZnAl2O4 film was not uniform and clear theboundary between ZnAl2O4 and substrate. To obtain
the optical parameters such as refraction index anddielectric constant, it is necessary to prepare the film
which is clear boundary at the each layer.

0.5
ZnAl 2O4
Thin film
Zn0.90Transmittance (a.u.)1

Zn0.95
0
140

160

180
200
220
Wavelength (nm)240

Fig. 2. Photoluminescent excitation spectra of
ZnAl2O4 powders and transmittance spectrum of
ZnAl2O4 thin film measured at BL3B.

60

Ⅲ -2. Materials Sciences

271BL3B 福井一俊.docx

BL3B, BL7B

PL Spectra of Laser Irradiated a-CNx Thin Films
1

K. Nakamura1, K. Ikeda1, K. Fukui1 and K. Yamamoto2

Department of Electrical and Electronics Engineering, University of Fukui, Fukui 910-8507, Japan2
Far-infrared region Development Research Center, University of Fukui, Fukui 910-8507, Japan
Upon band to band excitation, a-CNx thin film show awide visible range emission band. This band consistsof many emission bands which reflect the variation of
the a-CNx components and/or structures. However, allemission bands are commonly excited same two
photon energy range around 4.5 eV and 5.5 eV.

Amorphous carbon nitride (a-CNx) is known to havean interesting mechanical, electrical and opticalproperties. Therefore, it has been actively carried out
research on synthesis, structure and properties. The aCNx has also been expected as an inexpensive whitelight emitting diode material. However, since theoptical properties of a-CNx are sensitive to the
preparation conditions, photoluminescence (PL)spectra of the samples evaporated under same
conditions are, in some cases, different with each other.
Recently, it is found that laser irradiated a-CNx thinfilms after depositing show much higher PL intensity
than those of as-deposited a-CNx thin films, and their
PL spectra have better reproducibility [1,2]. In thisreport, we show the RF power dependence of the
combined excitation emission spectra (CEES) by using
laser irradiated a-CNx thin films.
All samples on Si substrates are deposited about 1
μm by the RF-sputtering method at Ryukoku
University [1]. The sputtering conditions with graphitetarget in the low pressure nitrogen gas environment of
every samples are same without RF power of 40 W, 60
W and 100 W. Beam diameter of laser-irradiation isabout 1.2 mm, and irradiation time is 240 minutes byusing He-Cd laser (325 nm) with 20 mW. All CEES
measurements have been performed at BL3B and
BL7B at about 20 K.
Figure 1 shows CEES of RF power 60 W a-CNx thinfilm. High intensity regions are represented by warm
colors. It is found that the PL peak is located around
2.8 eV, and peak energy has no excitation energy
dependence. It means that this emission process is
unique process upon the band to band excitation of aCNx. Figure 1 also suggest that this PL is strongly
excited around 4.5 eV and 5.5 eV. Both the excitationenergy independence of the PL peak and the existenceof two peaks (4.5 eV and 5.5 eV) in the excitation
spectra are observed in both 40 W and 100 W a-CNx.
Figure 2 shows the emission spectra at the excitationenergy 4.6 eV for all samples. All spectra normalized
by the maximum intensity. Although it is clear that theemission peak is shifted to lower photon energy sidewith an increase of the RF power, these three samples
are emitted almost same photon energy region from 1.6
eV to 4.0 eV. In addition to the asymmetric spectral
shape of these emission bands, it probably suggeststhat these PL bands consists of many emission bands
which are continuously distributed from 1.6 eV to 4.0
eV, and RF power affects the distribution of the a-CNx
component which corresponds to each emission band.
In summary, we measured the combined excitation
emission spectra of laser irradiated a-CNx thin films.

Ⅲ-2

Fig. 1. CEES of a-CNRF-60W at 9K.

Fig. 2. Emission Spectrum at 9K.
[1] M. Satake et al., 2013 JSAP Autumn Meeting
28p-PB3-11.
[2] D. Itoh et al., UVSOR Activity Reports 41 (2014)
68.

61

281BL3B
河相武利.docx
Ⅲ -2. Materials
Sciences

BL3B

Energy Transfer from I− Centers to In+ Centersin Co- Doped NaCl:I−, In+ Crystals
A. Iguchi and T. Kawai

Graduate School of Science, Osaka Prefecture University, Sakai 599-8531, Japan

260

Optical Density

0.3 RT
0.2

Wavelength (nm)
240
220
-

NaCl:In+C abs.

Lumi. Int. (arb. unit)

280

NaCl:I
Lumi.

0.1

0
4.4 4.6 4.8 5.0 5.2 5.4 5.6 5.8 6.0
Photon Energy (eV)Fig. 1. Luminescence spectrum (blue) of NaCl:I− andabsorption spectrum (red) of NaCl:In+ at RT.

Luminescence Intensity (arb. units)

Ⅲ-2

Alkali halide crystals have the wide band-gap up to
the vacuum ultraviolet energy region and are a suitable
candidate host for doping of impurity ions. Though
optical studies for impurity centers doped in the alkali
halide crystals have been extensively performed,
energy transfer between two kinds of impurity centersin co-doped alkali halide crystals is comparatively lessstudied [1,2]. In this study, we have investigated theenergy transfer from I− centers to In+ centers in NaCl
crystals co-doped with the I− and In+ ions at the BL-3B
line of UVSOR.
Figure 1 shows the absorption spectrum of NaCl:In+
and the luminescence spectrum of NaCl: I− in theenergy range from 4.4 to 6.0 eV. In NaCl:In+, theabsorption band observed around 5.2 eV is called the
C absorption band and is attributed to the intra-ionictransition from the 1A1g to 1T1u states [3,4]. In NaCl:I−,
the luminescence band appearing around 5.3 eV iscalled the NE luminescence band and comes from theone-center type localized exciton, which is localized
mainly on the central site of a substituted I− impurity
with small lattice relaxation [5]. The NE luminescenceband in NaCl:I− has an large overlap with the Cabsorption band in NaCl:In+. The fact indicates thepotential of the energy transfer from the I− centers to
the In+ centers in co-doped NaCl:I−, In+ crystals.
According to the Förster model [6,7], the criticaldistance between both centers for resonant energy
transfer is estimated to be about 3.0 nm from theoverlap with the luminescence and absorption bands.
Figure 2 shows the excitation spectra in NaCl:I−,
NaCl:In+, and co-doped NaCl:I−,In+ crystals at room
temperature (RT). The NE luminescence band in
NaCl:I− is efficiently excited in the energy region
around 6.4 eV and 7.3 eV, where the absorption banddue to the I− centers is located. In NaCl:In+, theluminescence band at 3.0 eV, which comes from theIn+ centers, is efficiently excited at the C absorption
bands of the In+ centers.
In the co-doped NaCl:I−,In+ crystal, the excitation
spectrum for the 3.0 eV luminescence band of the In+
centers has the broad band around 5.7 eV in addition
to the C band of the In+ centers. Since the 5.7 eV bandcorresponds to the excitation band for the NE
luminescence band in NaCl:I−, the fact implies theexistence of the energy transfer from the I− centers to
the In+ centers in the co-doped NaCl:I−, In+ crystal. In
order to clarify the energy transfer mechanism fromthe I− centers to the In+ centers in the co-doped NaCl:I−,
In+ crystals, the measurements of the decay kinetics ofthe luminescence are needed.

300

4.0

250

RT

Wavelength (nm)200

NaCl:In+for 3.0 eV

NaCl:I for 5.3 eVC
NaCl:I -,In+for 3.0 eV

4.5

5.0 5.5 6.0 6.5
Photon Energy (eV)

7.0

7.5

Fig. 2. Excitation spectra of NaCl:In+ (red), NaCl:I−(blue), and co-doped NaCl:I−, In+ (purple).
[1] A. F. Muñoz and J. O. Rubio, Phys. Rev. B 38
(1988) 9980.
[2] A. Méndez et al., J. Lumin. 79 (1998) 269.
[3] A. Ranfagni et al., Advances in Physics 32 (1983)
823.
[4] P. W. M. Jacobs, J. Phys. Chem. Solids 52 (1991)
35.
[5] I. Akimoto et al., Phys. Status Solidi C 6 (2009)342.
[6] T. Förster, Disc. Faraday Soc. 27 (1959) 7.
[7] D. L. Dexter, J. Chem. Phys. 21 (1953) 836.

62

Ⅲ -2. Materials Sciences

291BL3B 河相武利.docx

BL3B

Luminescence of Copper-Heavily-Doped Cesium Iodide Crystals
1

T. Hirai1 and T. Kawai2

College of Science and Engineering, Ritsumeikan University, Kusatsu 525-8577, Japan
2
Graduate School of Science, Osaka Prefecture University, Sakai 599-8531, Japan

Impurity ions doped into insulating materials, suchas Tl: CsI and Tb: Gd2O2S, have been widely used as
scintillators in γ- and X-ray detectors. In order toremarkably improve performance of those detectors,
attempts have been recently made to develop a new
type of scintillator utilizing semiconductor
microcrystals instead of the impurity ions, since theradiative recombination rates of excitons insemiconductor microcrystals should be much higher
than those in the impurity ions [1]. In the present study,
we have focused on microcrystals of semiconducting
Cu compounds that can be formed in a CsI crystal [2].
We have investigated photoluminescence (PL) andphotoluminescence excitation (PLE) spectra of Cu
heavily doped into CsI bulk crystals, which wereprepared by the Bridgman method, by use ofsynchrotron radiation at the BL3B line in UVSOR.
Figure 1 shows a PL spectrum (black) of a Cu (10mol%): CsI crystal at 5 K under excitation at 200 nm,
which corresponds to the band-to-band excitation of
CsI. A strong PL band is observed at 440 nm. A weakPL band can be also observed at around 405 nm as ashoulder of the PL band at 440 nm. In addition, there
appear two PL bands at 290 and 340 nm, which
originate from radiative recombination of self-trapped
excitons (STEs) in CsI [3]. Three PLE spectra detected
at 340 (violet), 440 (red) and 405 (blue) nm are also
shown in Fig. 1. The PLE spectrum for the 340-nm PL
band corresponds to that for the PL bands of STEs in
CsI. In the PLE spectrum for the 405-nm PL band, two
PLE bands are observed at 260 and 280 nm. For the
440-nm PL band, there appear three PLE bands at 245,
270 and 290 nm.
In Fig. 2, a PL spectrum (black) of a Cu (40 mol%):
CsI crystal under excitation at 200 nm at 5 K is shown.
According to the phase diagram of the CuI-CsI system[4], the mixed melts of 60 mol% CsI and 40 mol% CuI
form Cs3Cu2I5 crystals. In the PL spectrum, there
appear two PL bands at 440 and 405 nm, which arequite similar to the 440- and 405-nm PL bands in Cu(10 mol%): CsI (Fig. 1). Futhermore, two PLE spectra
detected at 440 (red) and 405 (blue) nm shown in Fig.
2 are likely to correspond to those for the 440- and 405nm PL bands shown in Fig. 1. These PL and PLE
spectra suggest that those PL bands in the two samples
have the same origin. The 440-nm PL band is certainly
identical to the 440-nm PL band arising from radiative
recombination of excitons in Cs3Cu2I5, which has beenreported by Hirao et al. [5]. Although the origin of the405-nm PL band is unclear, the PL band perhaps comesfrom Cu impurities in CsI crystals.
From the above results, we can conclude that the Cu(10 mol%): CsI crystal contains Cs3Cu2I5
microcrystals. Additionally, we notice that theexcitation of CsI host crystals in the Cu (10 mol%): CsI
crystal gives rise to the photoluminescence originatingfrom the Cs3Cu2I5 microcrystals.

Ⅲ-2

Fig. 1. PL spectrum of a Cu (10 mol%) : CsI crystal
under excitation at 200 nm (black), and its PLE spectra
detected at 340 (violet), 405 (blue) and 440 (red) nm.

Fig. 2. PL spectrum of Cs3Cu2I5 excited at 200 nm(black) , and its PLE spectra detected at 405 (blue) and440 (red) nm.
[1] K. Shibuya et al., Rev. Sci. Instrum. 78 (2007)
083303.
[2] M. Nakayama et al., Japanese Unexamined Patent
Application No. 2003-147343.
[3] H. Nishimura et al., Phys. Rev. B 51 (1995) 2167.
[4] A. Wojakowska et al., J. Chem. Eng. Data 48(2003) 468.
[5] N. Hirao et al., Meet. Abs. Phys. Soc. Jpn. 58(2003) 742.

63

301BL3B
吉野正人.docx
Ⅲ -2. Materials
Sciences

BL3B

Excitation and Emission Spectra of Pr3+ in Ga- Doped SrY2O4
M. Yoshino, S. Nishiki and S. Watanabe

Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan

Ⅲ-2

The trivalent lanthanide ions (e.g., Ce3+, Nd3+, Er3+)
in Oxide crystals have drawn attentions due to theirapplication for luminescent materials in NIR to UV
regions such as solid-state lasers or phosphors. The
trivalent cerium ion, Pr3+, has also attracted attentions
as luminescence centers. In this work, the excitation
spectra and emission spectra for Pr3+ in Ga-doped
SrY2O4 crystal have been measured. For thecomparison, the spectra for host oxide, Ga-doped
SrY2O4 have been also measured. The Ga-doped
SrY2O4 sample and Pr,Ga-doped SrY2O4 sample areproduced by solid state reactions. Pr-doped samples
are annealed in N2-H2 atmosphere. The concentrationof Ga and Pr in the samples are 3 mol% and 1 mol%,respectively. The emission spectrum of Ga-doped
SrY2O4 at 200 nm excitation is shown in Fig. 1. The
broad peaks from interband optical transition exist
around 400 nm. The excitation spectrum monitored at
400 nm emission is shown in Fig. 2. The peak near 200
nm relates the absorption from the optical transition

around the band edge in Ga-doped SrY2O4. This
corresponds with that in the spectrum monitored at 423
nm in SrY2O4 [1]. The emission spectrum for Pr3+ in
Ga-doped SrY2O4 is shown in Fig. 3. The spectrum
shows sharp peaks originated from Pr3+ 4f-4f transition
and the broad peaks of interband optical transition
from the host are weak while the excitation wavelengthis 200 nm. It can be said that excitation energy could
transferred from host to Pr3+ ion. Figure 4 shows theexcitation spectrum monitored 520 nm for Pr3+ in Gadoped SrY2O4. The peaks around 320 nm and 280nm
originate in the absorption of the 4f-5d transitions of
Pr3+. Then, in the excitation at around 320 nm and 280
nm, Pr3+ ions are excited directly, the emission
spectrum mainly shows the sharp peaks originatedfrom Pr3+ 4f-4f transition as shown in Fig. 3.

Fig. 1 Emission spectrum of Ga-doped SrY2O4.

Fig. 2 Excitation spectrum of Ga-doped SrY2O4.

Fig. 3. Emission spectrum of Pr, Ga-doped SrY2O4.

Fig. 4. Excitation spectrum of Pr, Ga-doped SrY2O4.

[1] M. Yoshino, S. Watanabe and Y. Ichikawa,UVSOR Activity Report 36 (2009) 122.

64

Ⅲ -2. Materials Sciences

311BL3B 大野宣人.docx

BL3B

Photoluminescence of TlBr-TlI Mixed Crystals
1

N. Ohno1 and K. Matsui2

Fundamental Electronics Research Institute, Osaka Electro-Communication University,
Neyagawa 572-8530, Japan
2
Graduate School of Engineering, Osaka Electro-Communication University,
Neyagawa 572-8530, Japan
self-tapped excitons decreases due to the smaller
exciton-phone interaction, resulting in decreasing of
Stokes shift of the luminescence band width.

Thallous halide crystals are compound semiconductors attractive for fabrication of γ ray detectors.
They exhibit high stopping power originated fromhigh atomic number and high density [1]. TlCl and
TlBr crystalize in cubic CsCl-type structure, while thestable form of TlI is orthorhombic structure below
170ºC [2] and is cubic CsCl-type one above it [3,4].
Therefore, it is difficult to grow a large-sized single
crystal of TlI under atmospheric conditions. In thepresent study, mixed singles of TlBr-TlI in cubic
CsCl-type structure have been grown and examined
to reveal the optical properties such as the opticalabsorption, photoluminescence and luminescence
excitation spectra.
The crystals TlBrxI1-x were grown by the
Bridgeman method after mixing appropriate amountsof purified TlBr and TlI. A plate of the sample cut
from as-grown crystals was polished with alumina
abrasive plastic film, then etched in hot distilled water.
The optical measurements were made at 10 K and attemperatures up to 300 K at BL3B station.
Figure 1 shows and absorption spectra (black
curves) of TlBrxI1-x mixed crystals for different
compositions measured at 10 K. In pure TlBr (x=1.0),
the fundamental absorption edge is of the indirectforbidden exciton transition [5]. As clearly seen in
Fig.1, the indirect absorption edge moves to lower
energy with increasing iodine concentration. At x=0.4,
the absorption-edge energy is located at around 2.0
eV. Below x≤0.3, the absorption spectra showed noclear absorption edge which was ascribed to the phasetransition in TlBrxI1-x mixed crystals. In I-rich mixed
crystal, the cubic phase becomes unstable below theroom temperature.
On exciting TlBrxI1-x mixed crystals with light at
the energy above the indirect absorption edge, a
Stokes-shifted broad luminescence band is observed
as clearly seen in Fig. 1 (blue curves). In TlBr (x=1.0),
the shape of the spectra is almost Gaussian aspreviously observed in lightly I-doped TlBr [6]. The
broad band was identified to be originated fromannihilation of excitons self-trapped at an iodine ion
due to the impurity-induced self-trapping. It is worth
noting in Fig. 1 that the amount of the Stokes shift
and the full width at the half maximum of theobserved luminescence band become continuously
smaller as the iodine concentration is increased.
Therefore, with increasing I-ion concentration in
TlBrxI1-x mixed crystals, the lattice relaxation of the

Ⅲ-2

Fig. 1. Photoluminescence (blue) and absorption
coefficient (black) spectra of TlBrxI1-x at 10 K.
[1] K. Hitomi, T. Shoji and K. Ishii, J. Crystal
Growwth 379 (2013) 93.
[2] N. Ohno and M. Itoh, J. Phys. Soc. Jpn. 62 (1993)
2966.
[3] K. Heidrich, W. Staude, J. Treusch and H.
Overhof, Solid State Commun. 16 (1975) 1043.
[4] A. Fujii and I. Ishikawa, J. Phys. Soc. Jpn. 56(1987) 4581.
[5] J. Nakahara, K. Kobayashi and A. Fujii, J. Phys.
Soc. Jpn. 37 (1974) 1312.
[6] J. Nakahara and K. Kobayashi, J. Phys. Soc. Jpn.
40 (1976) 180.

65

321BL3B
北浦守.docx
Ⅲ -2. Materials
SciencesBL3B

VUV Absorption Spectra of Racemic Alanine Crystals
S. Tanaka, M. Kitaura and A. Ohnishi

Department of Physics, Faculty of Science, Yamagata University, Yamagata 990-8560, Japan6-1

absorption coefficient (cm )1.5x10

1x10

6

5x10

5

0

6

10

14
18
Photon Energy (eV)22

Fig. 1. VUV absorption spectra of a racemic alanine
crystal, derived from the Kramers-Kronig analysis ofreflectivity spectra in 5-25 eV range. The data weremeasrued at room temperature.
5

3x10
-1

absorption coefficient (cm )Ⅲ-2

Amino acids are simple chiral molecules with Land D-forms. Protein molecules are composed of only
L-amino acids. This is called “homochirality”, which
has attracted much attention in relation to the mysteryof life. It has been generally accepted that thehomochirarity of proteins comes from the creation of
L-amino acids by circularly polarized VUV light inouter space region. So far, VUV absorption spectra ofamino acids have been studied using aqueous solutions[2] and thin films [3]. To our knowledge, there is noreport on VUV absorption spectra of amino acidcrystals. We have measured VUV reflectivity spectraof L-alanine single crystals to obtain the spectra ofvarious optical constants. In the present study, we havemeasured reflectivity spectra of racemic alanine
crystals, and investigated the change of VUVabsorption spectra on crystal structure, which gives us
information on intermolecular interaction among
alanine molecules.
Racemic alanine single crystals were obtained fromsaturated aqueous solution of commercial alanine
powder. The crystals are of needle shape. The crystal
axes were distinguished by two polarizers arranged ina cross Nicole alignment. To obtain VUV absorption
spectra, we performed the Kramers-Kronig (K-K)
analysis of reflectivity spectra. In the K-K analysis, itis necessary to determine absolute reflectivity over theenergy range as wide as possible. We obtained absolutereflectivity by the correction using refractive indexes
in the transparent region below 2 eV.
Figure 1 shows VUV absorption spectra of aracemic alanine crystal for different configurations oflinearly polarized light. A few peaks appear below 15
eV. No remarkable peak is found above 15 eV. Thisfeature is similar to VUV absorption spectra of Lalanine crystals [4]. On the other hand, there is a linear
dichroism between them. Figure 2 shows VUVabsorption spectra extended in the 5-12 eV range.
Absorption peaks for a racemic alanine crystal areapparently broader than those for L-alanine crystals [4].
The reason for such broadening of absorption peaks is
not clear. To clarify it, theoretical approaches based onelectronic state calculations are effective. Theoretical
calculations were applied for alanine molecules [5,6].
However, there is no instance applied for alanine bulk
crystals. Therefore, theoretical calculations for alanine
bulk crystals are desired.

5

2x10

5

1x10

0

6

8
10
Photon Energy (eV)12Fig. 2. VUV absorption spectra of a racemic alanine
crystal, extended in the 5-12 eV range.
[1] P. A. Snyder et al., Biopolymers 12 (1973) 975.
[2] T. Inagaki et al., Biopolymers 12 (1973) 1353.
[3] M. Tanaka et al., J. Electron Spectrosc. Relat.
Phenom. 181 (2001) 177.
[4] S. Tanaka et al., UVSOR ACTIVITY REPORT 42(2015) 70.
[5] F. Kaneko et al., J. Phys. Soc. Jpn. 78 (2009)013001.
[6] T. Kaneko et al., CHIRALITY 23 (2011) E52.

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331BL3B 北浦守.docx

BL3B

Construction of Vacuum Referred Binding Energy Diagram in
Gd3(Al,Ga)5O12:Ce3+ Phosphors
1

K. Asami1, J. Ueda1, M. Kitaura2 and S. Tanabe1

Graduated School of Human and Environmental Studies, Kyoto University,
Kyoto 606-8501, Japan2
Department of Physics, Faculty of Science, Yamagata University,
Yamagata 990-8560, Japan

Ce3+- doped Gd3Al5-xGaxO12 (GAGG) is apromising optical material for the scintillator [1]. It is
known that the quantum efficiency (QE) of GAGG:
Ce3+ becomes lower with increasing Ga content, x [2].
To understand optical properties of a phosphor, anelectronic structure of host material should be
investigated. Dorenbos suggested the construction
method of the vacuum referred binding energy(VRBE) diagram [3]. This diagram can be constructedfrom the spectroscopy data of host exciton energy andcharge transfer transition (CTS) of Eu3+. To determine
thosetransitionenergies,
wemeasured
photoluminescence excitation (PLE) spectra of bothCe3+- and Eu3+- doped GAGG at UVSOR BL3B.
Figure 1 (a) shows the PLE spectra of Ce3+- doped
GAGG phosphors with various Ga content. In PLE
spectra, the host exciton peaks are observed at around200 nm and three bands derived from Ce3+ are
observed at 220 nm (4f-5d3, 4, 5), 350 nm (4f-5d2) and
450 nm (4f-5d1). Several peaks at 270 nm and 310 nm
are due to the intra – 4f transitions of Gd3+ in the host[1]. From the host exciton peaks, the bandgap energies
of GAGG (x = 1, 2, 3 and 4) are estimated to be 6.84,
6.81, 6.63 and 6.55 eV, respectively. With increasing
Ga content, the bandgap becomes narrower. Figure 1(b) shows the PLE spectra of Eu3+: GAGG phosphors.
The strong broad band is located at around 250 nm,
which is attributed to CTS of O2--Eu3+ →O--Eu2+. With
increasing Ga content, the peak wavelength of CTS
shifts to longer. The top of valence band is calculated
by subtracting the CTS energy from the determined 4f
state energy of Eu2+.
Based on the method of VRBE diagram construction
by Dorenbos [3], the VRBE diagram of Ce3+-doped
GAGG are constructed by using measurement dataobtained in Fig.2 . The energy Ev which is top ofvalence band shifts higher with increasing Ga content
while the energy Ec which is the bottom of conduction
band slightly changes. When the energy of Ce3+:5d1 isclose to Ec, the excited electrons located at Ce3+:5d1
state tend to move to the conduction band by thermal
ionization. From this diagram, with increasing Ga
content, the energy gap between Ec and Ce3+:5d1
becomes smaller. Therefore QE of Ce3+ in GAGG
system becomes decrease by thermal ionizationquenching.

Ⅲ-2

Fig. 1. PLE spectra of (a). Ce3+-doped GAGG (Left)and (b). Eu3+-doped GAGG (Right).

Fig. 2. Constructed VRBE diagram of Ce3+:GAGG.
[1] M. Kitaura, A. Sato, K. Kamada, A. Ohnishi and M.
Sasaki, J. Appl. Phys. 115 (2014) 083517.
[2] J. M. Ogiegło, A. Katelnikovas, A. Zych, T. Jüstel,A. Meijerink and C. R. Ronda, J. Phys. Chem. A 117(2013) 2479.
[3] P. Dorenbos, J. Lumin. 134 (2013) 310.

67

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Ⅲ -2. Materials
Sciences 北浦守.docx

BL3B, BL6B

Local Structure of Gd3+ and Ce3+ Ions in Gd3Al5(1-x)Ga5xO12 Crystals
Studied by UV and IR Absorption Spectroscopy
C. Oyama1, M. Kitaura1, M. Ishizaki1, K. Kamada2, S. Kurosawa2, A. Ohnishi1 and K. Hara3
1

Ⅲ-2

Faculty of Science, Yamagata University, Yamagata 990-8560, Japan
2
NICHe, Tohoku University, Sendai 980-0845, Japan
3
RIE, Shizuoka University, Hamamatsu 432-8011, JapanThe Ce:Gd3Al2Ga3O12 (Ce:GAGG) garnet crystals
has been known as an inorganic scintillators with highlight output, high energy resolution, and no
hygroscopic nature [1]. The scintillation properties
strongly depend, especially on the composition ratio of
Al3+ and Ga3+ ions. The control of the composition
ratio of Ga3+ and Al3+ ions is of great importance tobring out the high performance of Ce:GAGG. The
crystals of garnet type have three cation sites named A,
B, and C, in which cations are coordinated by eight,
six, and four oxygen ions, respectively. The chemical
formula of oxide garnets is expressed as A3B2C3O12.
Al3+ and Ga3+ ions can regularly occupy octahedral (B)and tetrahedral (C) sites. So far, the distribution of Al3+and Ga3+ ions has been investigated in garnet solid
solutions. Although there are a number of reports on
this subject, it still remains obscure which sites Ga3+
ions occupy preferentially [2,3]. From the viewpoint ofbasic science and industrial application, it isworthwhile to investigate the site preference of Ga3+
ions in Ce:Gd3Al5(1-x)Ga5xO12 crystals. In the present
study, Gd3+ and Ce3+ 4f-4f absorption spectra ofundoped and cerium-doped Gd3Al5(1-x)Ga5xO12 (0.4 x1) crystals have been measured at low temperatures
below 10 K in ultraviolet and infrared ranges [4].
Generally, the 4f-4f transitions of rare- earth ions giverise to narrow absorption peaks, and they are expected
to be sensitive to cation distribution. Thus, Gd3+ andCe3+ 4f-4f absorption peaks can utilize as a probe toclarify the site preference of Ga3+ ions.
Figure 1 shows absorption spectra of undoped
Gd3Al5(1-x)Ga5xO12 (a) and cerium-doped Gd3Al5(1-x)
Ga5xO12 (b), which are due to 4f-4f transitions of Gd3+
and Ce3+ ions, respectively. In Fig. 1(a), the peak
positions and spectral widths of absorption peaks for x= 0.4 were almost the same as those for x = 1.0. This
result suggests that, for x = 0.4, Ga3+ and Al3+ ionsoccupy octahedral and tetrahedral sites, respectively.
This opinion is consistent with that by Ogiegro et al.
[2], but is inconsistent with that by Kanai et al. [3] The
peak positions for x = 0.6 and 0.8 are shifted to lower
energy side. The spectral widths were clearly broader,compared to that for x = 0.4 and 1.0. The peak shift andspectral broadening indicate that, for x = 0.6 and 0.8,
Ga3+ and Al3+ ions occupy both of octahedral andtetrahedral sites. In Fig. 1(b), the position of theabsorption peak shifts to lower wavenumber side withincreasing x. This is explained by the reduction of
crystal field strength at Ce3+ ions, which is due to thereplacement of Ga3+ ions for Al3+ ions [5]. The spectral
width for x = 0.4 is largest of all x, suggesting thevariety of crystal field at Ce3+ ion sites. In oxide
garnets, oxygen vacancies and antisite ions are known
as primary imperfections [6]. It is thus likely thatvarious Ce3+ ions coupled with such imperfections areformed. The present study reveals that the variety of
Ce3+ ions coupled with imperfections is the largest for
x = 0.4. The reason for this is unclear, so that it is an
open subject to be solved in near future.

Fig. 1. Absorption spectra due to 4f-4f transitions of
Gd3+ ions in undoped Gd3Al5(1-x)Ga5xO12 (a) and Ce3+
ions in cerium-doped Gd3Al5(1-x)Ga5xO12 (b). These
data were observed below 10 K.
[1] K. Kamada et al., J. Cryst. Growth 352 (2012) 88.
[2] J. M. Ogiegro et al., J. Phys. Chem. C 35 (2013)322.
[3] T. Kanai et al., J. Am. Ceram. Soc. 462 (2008) 456.
[4] For x = 0 and 0.2, it is not possible to obtain single
crystals of garnet phase.
[5] P. Dorenbos, J. Lumi. 151 (2014) 224.
[6] M. Kitaura et al., J. Appl. Phys. 115 (2014) 083517.

68

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351BL4U 大東琢治.docx

BL4U

Observation of Morphology of a Fuel Cell by Using a Humidity Control
Sample Cell for STXM
2

T. Ohigashi1, 2, Y. Inagaki1, T. Horigome1 and N. Kosugi1, 2

UVSOR Synchrotron, Institute for Molecular Science, Okazaki 444-8585, Japan
School of Physical Sciences, The Graduate University for Advanced Studies (SOKENDAI), Okazaki 444-8585,
Japan1

One of the advantages of a scanning transmission Xray microscopy (STXM) is tolerance for ambientconditions of samples. Due to high transmittance of the
X-ray, the STXM does not require a vacuum condition
but enables to observe electronic states of various
samples in atmospheric pressure or in water. Moreover,a focusing optical element of the STXM, namely, aFresnel zone plate, has characteristics such as long
working distance and long focal depth. These features
enable to perform in-situ observation, which is
impossible for an electron microscopy; therefore,various sample cells have been developed widely [1].
Control of humid atmosphere around the sample issometimes needed as a promising in-situ observation
method [2]; for example, a fuel cell working underhigh humid atmosphere and high temperature. In thisstudy, we have developed a humidity control sample
cell and performed a test measurement.
The present humid control sample cell (shown in Fig.
1) is a small chamber consisted of two silicon nitride
membranes (thickness of 100 nm) as windows [3]. A
small sensor to measure humidity and temperature(SHT7x, Sensirion AG) is set inside of the chamber.
The cell has three ports for inlet/outlet of gas flowing
and two of them were used in this study. The inlet andoutlet were connected to a feedthrough of the STXM
chamber with stainless tubes and dry/humid helium
gas is flowed into the cell to control the humidity.
Humidity in the gas is added by bubbling pure water ina bottle and its flow rate is controlled by using a needle
valve manually. The inside of the STXM chamber wastransferred by helium gas until air pressure. As aperformance test, we could change the humidity from
16.1 to 80.4% at 29C (temperature inside of the
STXM chamber was higher than RT).
By using the cell, morphological change of thinsections of a fuel cell fixed on Formvar membrane wasused as a test sample. 2-dimensional distributions offluorine in the polymer were obtained by using the Kabsorption edge of fluorine (687 eV). X-ray
transmission images below and above of the edge (682
and 692 eV) were obtained. After conversion to opticaldensity images, the distribution of fluorine wasobtained by subtracting these images. The distributionof fluorine of the humidity at 16% and 8% are shownin Fig. 2. In these images, bright areas show highconcentration of fluorine and dark areas like cracks areseen on the center. By changing the humidity from
16% to 80%, 11% of the dark area decreased. Thischange of the area was caused by swelling of water
vapor of the polymer.
As a future plan, temperature control and autocontrol (i.e. feedback of the humidity and temperature)system are under consideration. As further applicationof this humid cell, observation of living biologicalsamples, such as cells and bacteria, will be performed.

Fig. 1. The inside and the base plate of the humidity
control sample cell.

(a)

(b)

500 nm

500 nmFig. 2. Distribution of fluorine in the fuel cell (bright
pixels show high concentration). The humidity of thecell were at (a) 16% and (b) 80%.
[1] X. H. Zhu et al., UVSOR Activity Report 41 (2014)72.
[2] V. Berejnov et al., ECS Transactions 41 (2011)395-402.
[3] T. Ohigashi et al., accepted.

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353BL4U_Way-Fang2012-2028.docx
Ⅲ -2. Materials
Sciences

BL4U

Influence of the Oxygen Vacancy at Ta/TaOx Interface on Resistive
Switching Memories
Y. F. Wang1, W. F. Pong1, J. S. Chen2, J. W. Chiou3, T. Ohigashi4 and N. Kosugi4

Department of Physics, Tamkang University, Tamsui 251, Taiwan
Department of Materials Science and Engineering, National Cheng Kung University, Tainan 701, Taiwan
3
Department of Applied Physics, National University of Kaohsiung, Kaohsiung 811, Taiwan
4
UVSOR Synchrotron, Institute for Molecular Science, Okazaki 444-8585, Japan1

2

Fig. 1. The STXM stack mapping of the crosssectional views of the Ta/TaOx/Pt thin film.

O K-edge

Normalized Absorption (arb. units)

Ⅲ-2

In this work, we have studied the electronic structure
and the location of oxygen vacancy in the interface of
the Ta/TaOx/Pt thin film and identify the type of
conducting path by scanning transmission x-ray
microscopy (STXM). The amorphous TaOx thin filmof thickness of ~150 nm was deposited by electron
beam evaporation on Pt/SiO2/Si substrates. Active
electrodes of Ta were thermally evaporated
respectively on the TaOx thin film with an equivalent
thickness of ~150 nm. As shown in Fig. 1, the STXM
stack mapping displays the cross-sectional views of
the Ta/TaOx/Pt thin film. Obviously, due to a uniform
diffusion occurs in the vicinity of the Ta/TaO2 interface,
the image of the Ta/TaOx/Pt thin film revealed that theinterfacial TaOx was ~180 nm in thickness. The
experiments were performed at the 4U beamline.
Figure 2 presents the corresponding O K-edge x-ray
absorption near-edge structure (XANES) spectra of the
Ta/TaOx/Pt thin film. The O K-edge STXM-XANES
spectra were recorded at three selected regions, cyan,
red and yellow, in the Ta/TaOx/Pt thin film,
respectively. According to the dipole-transition
selection rule, the features at ~530-545 eV are
attributed to the electron excitations from O 1s-derived
states to 2p-derived states, which are approximately
proportional to the density of the unoccupied O 2pderived states.1 The intensities of the O K-edge STXMXANES spectra of the red region are significantly
higher than that of the cyan region, which reflects theincrease in the number of unoccupied O 2p-derived
states. In other words, the STXM-XANES results
demonstrate that the population of defects at the O sites
in the Ta/TaOx interface and confirming the enhanced
density of states of O 2p-derived states, as the
population of defects and dangling bonds at/above
ECBM or EF in the Ta/TaOx interface. The STXMXANES results intensely support the phenomena thatthe occurrence of O 2p vacancy at Ta/TaOx interface
affects the resistive switching effect of the Ta/TaOx/Pt
thin film.

Ta

Ta+TaOx

TaOx

530

540

Photon Energy (eV)550

Fig. 2. The corresponding O K-edge STXM-XANES
spectra were recorded at three selected regions, cyan,
red and yellow, in the Ta/TaOx/Pt thin film,
respectively.
[1] J. W. Chiou et al., Appl. Phys. Lett. 81 (2002) 3389.

70

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Ⅲ -2. Materials Sciences

BL4U

Resolving the Electronic Structure of TiO2 Core-Shell Nanostructures
Using Scanning Transmission X-Ray Microscope
Y. F. Wang1, S. H. Hsieh1, J. W. Chiou2, W. F. Pong1, T. Ohigashi3 and N. Kosugi3
Department of Physics, Tamkang University, Tamsui 251, Taiwan
Department of Applied Physics, National University of Kaohsiung, Kaohsiung 811, Taiwan
3
UVSOR Synchrotron, Institute for Molecular Science, Okazaki 444-8585, Japan1

2

In this work, a three-dimensional (3D) hierarchicalTiO2 is successfully constructed using scanningtransmission x-ray microscopy (STXM). The core
portion of the sample contains rutile TiO2 nanowire(NW) while the rutile nanoparticles (RNP) and anatase
TiO2 are sequentially located on the shell region. The
TiO2 NW array provides a fast electron transport
pathway due to its quasi-single-crystalline structure
and the 3D configuration with NPs in the shell portion
provides a large surface area for more efficient photocharge separation without significantly sacrificing theelectron collection efficiency [1] which attracted ourinterest to probe this nano-scale interface phenomenon.
Figure 1 shows the (a) scanning electron microscope(SEM) and (b) STXM images of focused ion beam
(FIB) milled ANP/RNP/NW sample (above the reddashed line) on fluorine doped tin oxide (FTO)
substrate (below the red dashed line). The STXM
image was recorded at 460 eV. Fig. 2(a) shows themagnified STXM image in middle of sample of Fig. 1(b), where Ti L2,3-edge x-ray absorption near edge
spectra (XANES) had been measured and analyzed byprinciple component analysis (PCA) method. Fig. 2(b)was generated by PCA method analyzed in the same
area shown in Fig. 2(a) and indicates similar spectrawith the same color which defines the spatial
distribution of background (blue), Pt (yellow), carbon
film (purple) and surface (red), semi-surface (green),shell (brown), semi-core (orange), core (cyan) regionsof ANP/RNP/NW.
Figure 2(c) depicts the Ti L2,3-edge XANES ofdifferent regions of ANP/RNP/NW. Line shapes of egstates in L3-edge around 458 eV show the anatase
phase TiO2 in the surface (red) and semi-surface(green) regions [2], the rutile phase TiO2 in core (cyan)
region and mixed spectra of anatase and rutile phases
in shell (brown) and semi-core (orange) regions. These
results reveal that even the size of ANP is smaller than
the spatial resolution of STXM [3], the spatial
distribution of different phases of TiO2 core-shellstructures can still be successfully resolved under
STXM after FIB milling.

Fig. 1. (a) SEM image of ANP/RNP/NW on FTO. (b)STXM image of ANP/RNP/NW on FTO measured at
460 eV.

Fig. 2. (a) Magnified STXM image of ANP/RNP/NW.
(b) Spatial distribution of background (blue), Pt(yellow), carbon film (purple) and surface (red), semisurface (green), shell (brown), semi-core (orange),core (cyan) regions of ANP/RNP/NW. (c)Corresponding Ti L2,3-edge XANES of different
regions of ANP/RNP/NW.
[1] J. Y. Shin et al., Adv. Funct. Mater. 21 (2011) 3464.
[2] J. S. Yang, Appl. Mater. Interfaces 5 (2013) 7425.
[3] G. S. Henderson et al., Phys. Chem. Minerals 29(2002) 32.

71

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Ⅲ -2. Materials
Sciences

BL4U

Effect of Fe2O3 Coating on ZnO Nanorod Probed by
Scanning Transmission X-Ray Microscopy
Y. X. Chen1, Y. F. Wang1, J. W. Chiou2, C. L. Dong1, W. F. Pong1, T. Ohigashi3and N. Kosugi3
Department of Physics, Tamkang University, Tamsui 251, Taiwan
Department of Applied Physics, National University of Kaohsiung, Kaohsiung 811, Taiwan
3
UVSOR Synchrotron, Institute for Molecular Science, Okazaki 444-8585, Japan
1

2

Ⅲ-2

A novel Fe2O3/ZnO core-shell nanorods forphotoelectrochemical (PEC) water splitting isdeveloped recently [1]. The PEC performances, suchas photocurrent response and incident photons toelectrons (IPCE), vary with the thickness of Fe2O3
coating layer. However, the mechanism of the
improved PEC activity remains unclear. In this work,
pristine and Fe2O3 coated-ZnO core-shell nanorods
have been investigated by using x-ray spectroscopic
coupled with microscopic approaches (i.e., x-ray
absorption near-edge structure (XANES) and scanningtransmission X-ray microscopy (STXM)).
Figures 1(a)-(d) present the O K-edge STXM
images and its stack mappings of selected single
nanorod of Fe2O3/ZnO and pristine ZnO. The stackmappings display blue, yellow, red and green areas,
corresponding to the different regions that areassociated with different thickness and chemical
properties of the nanorods. Figures 1 (e)-(h) show the
XANES spectra, which correspond to Figs. 1 (a)-(d).
The different colored spectra reflect the various regionof the nanorods. The O K-edge probes the electrontransitions from O 1s core level to the O 2p-Zn 3d/4sp
hybridized states. The area under the spectrum reflect
the amount of density of unoccupied O 2p-derived
states. Thus, based on the spectroscopic results, theelectronic structures of the layer region and the corearea differ in both samples. Besides, strong anisotropic
effects are observed for both samples based onpolarized-XANES. Notably, the density of states in thesurface of the coated- and pristine nanorods aredifferent and the results suggest that the oxygenrelated deficiency is formed in the core-shell nanorods
which can affect the photocatalytic activity.
The enhanced PEC performance of Fe2O3 coatedZnO compared with pristine ZnO is likely to be owing
to the presence of interface that causes varied electron
density of states in the core-shell structured nanorods.

Fig. 1. Upper panel: Polarized O K-edge scanningtransmission X-ray microscopy image and its
corresponding stack mappings of selected single Fe2O3
coated- ((a)and (b)) and pristine ZnO nanorods ((c) and(d)). Lower panel: Polarized O K-edge STXMXANES spectra of Fe2O3 coated- ((e) and (g)) andpristine ZnO nanorods ((f) and (h)). The color
corresponds to the different region of the nanorods.
[1] Y. K. Hsu et al., ACS Appl. Interfaces 7 (2015)
14157.

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Ⅲ -2. Materials Sciences

BL4U

Mapping Oxygen Activation on Different Surface Facet of Palladium
Nanocrystals for Organic Catalysis
M. W. Lin1, Y. L. Lai1, M. H. Cao2, T. Ohigashi3, N. Kosugi3, Q. Zhang2 and Y. J. Hsu1*
National Synchrotron Radiation Research Center, Hsinchu 300, Taiwan
Institute of Functional Nano- & Soft Materials (FUNSOM) Soochow University, 215123, P. R. China
3
Institute for Molecular Science, Okazaki 444-8585, Japan
Electronic mail: yjhsu@nsrrc.org.tw
1

The O2 activation process where inert ground triplet
O2 is excited to produce highly reactive singlet O2
involves a key step in many organic oxidation andcancer treatment. However, it puzzles scientist what
factor induces the change of electron spin state of O2
molecules, although it has demonstrated that thepresence of noble metal nanoparticles can promote thegeneration of singlet O2. In previous works, Q. Zhang
et al., first demonstrated that surface facet of metal
nanocrystals is a key parameter for tuning the activities
of generating singlet oxygen. The experiments andsimulations elucidated that a spontaneous spin−flipprocess may occur during the molecular absorption
when appropriate surface facets are selected. On a
selected surface facet, the charge state of the Pd surface
is a critical parameter to activate O2 to generate a
species that behaves like singlet O2 both chemicallyand physically [1]. Based on this finding, we prepareda metal–semiconductor hybrid system in which thenanocrystals of Pd{100} facets are supported on TiO2
surface. Through illumination of appropriate light, theelectrons are anticipated be transferred from TiO2 tometallic Pd catalysts and thus enhance the oxygen
activation.
To confirm the charge transfer at the heterojunction
between TiO2 and Pd heterojunction, chemical maps ofoxygen distributions within microaggregates of
SiN/TiO2-Pd were studied by scanning transmission
X-ray microscopy (STXM). The TiO2-Pd hybrid
structures were obtained according to following
procedures. The Pd nanocubes with 20 mg TiO2-P25
were synthesized according to a previously described
protocol [2]. In a typical synthesis, 8.0 mL of anaqueous solution containing poly(vinyl pyrrolidone)(PVP, MW = 55 000, Sigma– Aldrich, 856568-100g;
105 mg), l-ascorbic acid (AA, Sigma–Aldrich, A027825g; 60 mg), and KBr (300 mg) was heated in a 50 mL
three- necked flask for 10 min at 80 oC with magnetic
stirring. Subsequently, 3.0 mL of an aqueous solutionof potassium palladium(II) chloride ([K2PdCl4],
Aladdin, 1098844-1g ; 63 mg) was added with apipette, and the reaction mixture was stirred at 80 oC
for 3 h. The powders were collected by centrifugation
and washed with deionized water to remove excess
PVP. The final product was dissolved into thedeionized water and then dipped onto the SiN substrate.
Figure 1 displays the STXM images of (a) SiN/
TiO2-Pd (cubes {100} facets) and (b) SiN/TiO2-Pd(octahedrons {111} facets), which are optical density(OD) maps obtained at 530.9 eV, the major
characteristic absorption peaks of TiO2. The brightregion represents elemental distributions of O in TiO2.
The corresponding micro NEXAFS spectra extracted
from different color circles as shown in Figs. 1(a) and(b) are displayed in Figs. 2(a) and 2(b), respectively.
The resonance at 530.9 eV for both surface facet is
assigned to 1σu → 1πg* transition for TiO2 contactedwith Pd surface. Noticeably, the absorption peak athigher photon energy for TiO2-Pd{100} (at 533.3 eV)shifts to lower photon energy in comparison with TiO2Pd{111} (at 533.6 eV), indicating that significant
electrons are transferred from TiO2 to the Pd{100}surface. The preliminary results indicate that thechemical mapping of oxygen activation on different
surface facet of Pd nanocrystals will shed light ondesigning high efficient noble metal nanocatalysts fororganic oxidation and cancer treatments.
(a)

(b)

Fig. 1. The OD images of (a) SiN/TiO2-Pd (cubes{100} facets) and (b) SiN/TiO2-Pd (octahedrons {111}facets) obtained at 530.9 eV.
(a)

(b)

2.6

2.2

2.4

2.0
1.8

Intensity

2.0

A
B
C
D
E

2.2

1.6
1.4
1.2

1.8

1.4
1.2
1.0
0.8

1.0

0.6

0.8

0.4
0.2

0.6
0.4

520

A
B
C
D

1.6

Intensity

2

0.0

530

540

550

Photon energy (eV)

560

570

520

530

540

550

560

570

Photon energy (eV)Fig. 2. Micro-NEXAFS spectra of the O K-edge
extracted from the Figs. 1(a), and 1(b), respectively.
[1] L. Ran, M. Keke, Y. Xiaodong, Y. Wensheng, H.
Yaobing, W. Jianyong, F. Yao, W. Xisheng, W. Xiaojun,
X. Yi and X. Yujie, J. Am. Chem. Soc. 135 (2013) 3200.
[2] B. Li, R. Long, X. Zhong, Y. Bai, Z. Zhu, X. Zhang,
M. Zhi, J. He, C. Wang, Z.-Y. Li and Y. Xiong, Small
8 (2012) 1710.

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山本知之.docx
Ⅲ -2. Materials
Sciences

BL4B

Valence State Analysis of Mn Ions Doped in CaTiO3
2

Multiferroic materials, which have both
ferroelectric and ferromagnetic properties, have been
extensively studied for their potential applications inindustry. It was reported that the ABO3 type oxides,
e.g., BiFeO3 and YMnO3, show multiferroic properties,
and, on the other hand, dilute doping of the magnetic
element into the ferroelectric material gives
ferromagnetic property, which means such materials
also have a multiferroic property. One of the examplesof the later ones is Fe doped Bi4Ti3O12 [1]. To
understand the mechanism of newly appeared
ferromagnetism by dilute doping of magnetic elementin such ferroelectric materials, it is mandatory to know
the local environment of doped ions. However, it is not
so easy to determine the local environment of suchdilute dopants that one often skips such importantanalysis. In the case of the valence state analysis of 3dtransition metals, analysis with L2,3-edge X-ray
absorption spectrum (XAS) is quite powerful. We
could successfully determine local environment of Mn
and Fe ions doped in Bi4Ti3O12 by such XAS analysis[2, 3]. In the current study, valence state of Mn ions in
Mn-doped CaTiO3 is investigated by the Mn-L3 XAS
measurements.
Samples were prepared by the solid state reaction
method. High purity powders of CaCO3, TiO2 andMn2O3 were used as starting materials. These powders
were weighed by changing the concentration of Mnions, which were mixed and ground in an agate mortar
for 30 min. with ethanol. After drying these mixtures,
they were calcined in air for 2 hrs at 1423 K. Resultant
powders were mixed and ground again, which werepressed into pellet form. These pellets were sintered in
air at 1523 K for 6 hrs.
Crystal structures of the sintered samples were
examined by the powder X-ray diffraction with Cu-Kα
X-rays. Resultant XRD patterns show no extra peaks
except for those from orthorhombic perovskite
structured CaTiO3. In addition, lattice parameters
deduced from the observed XRD pattern decrease dueto the doping of Mn ions in CaTiO3. These results
suggest doped Mn ions are substituted at one of thecation sites, i.e., Ca2+ or Ti4+ sites, in CaTiO3.
Mn-L3 X-ray absorption spectrum of Mn doped
CaTiO3 were measured at BL4B in UVSOR. Incident
beam was monochromatized by varied-spacing planegrating (800 l/mm) and energy resolution was tuned bycontrolling the heights of the slits settled at upper andlower reaches of grating. Samples were placed on the

carbon adhesive tape, which was attached on the first
Cu-Be dinode of the electron multiplier. All the XAS
measurements were carried out in a total electron yield(TEY) mode by collecting the sample drain current.
Observed Mn-L3 XAS spectrum of Mn-doped
CaTiO3 is shown in Fig. 1 together with those of MnO(Mn2+) and MnO2 (Mn4+), in which that of MnO2
shows much better agreement with experimental
profile of Mn-doped CaTiO3. From these results, i.e.,
reduction of cell parameter by XRD and XANES
analysis, it can be concluded that Mn ions aresubstituted as Mn4+ ions in CaTiO3.

Intensity (arb. units)

Ⅲ-2

T. Murai1, K. Mori1, D. Yamada1, H. Murata2 and T. Yamamoto1,3

Faculty of Science and Engineering, Waseda University, Tokyo 169-8555, Japan
Association of International Arts and Science Institute of Natural Science, Yokohama City University, Yokohama
236-0027, Japan
3
Institute of Condensed-Matter Science, Waseda University, Tokyo 169-8555, Japan1

MnO
CaTi0.95Mn0.05O3

MnO2
640

645

Energy (eV)650

Fig. 1. Observed Mn-L3 X-ray absorption spectra of
Mn-doped CaTiO3 (middle) and Mn oxides (MnO: top
and MnO2: bottom).
[1] X. Q. Chen et al., Solid State Comm. 150 (2010)1221.
[2] K. Nishimura et al., IEEE Trans. Mag. 50 (2014)
2502306.
[3] K. Nishimura et al., UVSOR Activity Report 2013.

74

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Ⅲ -2. Materials Sciences

BL4B

Effect of Carbon Substitution on the Magnetic Properties of NanoArchitectural ZnO
Y. F. Wang1, W. F. Pong1, J. W. Chiou2, Y. Takagi3, T. Ohigashi3, T. Yokoyama3
and N. Kosugi3
Department of Physics, Tamkang University, Tamsui 251, Taiwan
Department of Applied Physics, National University of Kaohsiung, Kaohsiung 811, Taiwan
3
UVSOR Synchrotron, Institute for Molecular Science, Okazaki 444-8585, Japan
1

2

Based on our previous publication in Nanoscale 6,
9166-9176 (2014) entitled "Observation of the originof d0 magnetism in ZnO nanostructure using x-raybased microscopic and spectroscopic techniques", weobserved zinc vacancies (VZn) induced orbital
moments in oxygen sites near the surface of nanoarchitectural ZnO via combined synchrotron based
techniques and first principle calculations [1].
Moreover, since the carbon-doped ZnO have been
proposed and observed to be ferromagnetic with Curie
temperature well above room temperature, anddifferent carbon-related defects can result entirely
different magnetism in carbon-doped ZnO as
theoretically predicted. The correlation between
carbon-related defects and magnetism has become an
important issue experimentally for non-magnetic
impurities induced dilute magnetic semiconductors.
Due to this novel but controversial issue, nanoarchitectural ZnO is an ideal system to investigate. The
defects-related phenomenon due to the large surface tovolume ratio and intrinsic VZn on the surface ofnanostructures provide an opportunity to study how
carbon implantation affects the electronic structuresand magnetism of ZnO.
Figure 1 shows the C K-edge XANES and XMCD
spectrum of (a) carbon-implanted ZnO nano-wires(C:NWs) and (b) carbon-implanted ZnO nano-cactus(C:NCs). The XMCD signals appear around 287-292
eV suggest the origin of enhancement in room
temperature ferromagnetism (RTFM) of carbonimplanted nano-architectural ZnO can be correlated to
π-like hybridization states between carbon and oxygen[2], or zinc atoms. A weaker XMCD signal shows in
C:NCs with respect to NWs suggests VZn may suppress
ferromagnetic ordering between magnetic-impurities
since C:NCs contain higher density of VZn which
provides higher possibility for carbon substitution into
VZn sites and hybridize with oxygen to form non-spinpolarized CZn [3]. Carbon substituted oxygen (CO) can
introduce holes into system which satisfies Stonercriterion easily and results stronger localization spinsin carbon with respect to VZn [4]. Spin-polarized CZnCI (interstitial carbon) complex in VZn sites should beconsidered since highly dense CI can form carbonrelated complexes easily, but relatively weakmagnetization and XMCD signal in C:NCs withrespect to C:NWs suggests dominant spin-polarized
carbon impurity is CO [5].

(a)

(b)

Ⅲ-2

Fig. 1. C K-edge XANES and XMCD spectrum of (a)C:NWs and (b) C:NCs.
[1] Shashi B. Singh Nanoscale 6 (2014) 9166.
[2] Y. F. Wang et al., Sci. Rep. 5 (2015) 15439.
[3] S. K. Nayak et al., Phys. Rev. B 86 (2012) 054441.
[4] H. Peng et al., Phys. Rev. Lett. 102 (2009) 017201.
[5] S. Sung et al., Semicond. Sci. Technol. 26 (2011)014038.

75

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北浦守.docx
Ⅲ -2. Materials
Sciences

BL6B

UV-Induced Infrared Absorption Change in SrAl2O4:Eu Crystals: Influenceof Eu2+ Concentration on Trap Filling Process
Y. Tsurumi1, M. Kitaura1, A. Ohnishi1, T. Ishibashi2, S. Furukawa2, H. Oda2 and A. Yamanaka2
Facalty of Science, Yamagata University, Yamagata 990-8560, Japan
Chitose Institute of Science and Technology, Chitose 444-8585, Japan

1

2

Ⅲ-2

The control of shallow electron traps is mostsignificant in the development of high luminance longpersistent phosphors. So far, the energies of trap depths
have been evaluated from both sides of experimental
and theoretical approaches. The so-called “Dorenbos
model” [1] provides us fruitful information on the
optimization of co-dopants working as photocarrier
traps. On the other hand, the trap filling process is also
important; however, it has not been discussed so much,and still remains obscure. It is thus necessary to clarify
the factors dominating the trap filling process, and to
determine the optimal condition which gives rise tohigh luminance long persistent phosphorescence. For
these purposes, we have studied the influence ofeuropium concentration on trap filling process in
SrAl2O4:Eu2+ crystals grown by the Floating zone
method.
Figure 1 shows absorption spectra of SrAl2O4:Eu
crystals with europium concentrations of 0.1 (a), 0.3(b), and 1 mol% (c). These data were observed at 9 K.
Blue and red lines indicate absorption spectra observed
before and after UV-irradiation at 375 nm, which cancause long persistent phosphorescence [2,3]. An
absorption band appears around 2000 cm-1 under UVirradiation. The UV-induced band around 2000 cm-1
could not be observed for undoped SrAl2O4 crystals.
This result suggests that the Eu2+ 4f-5d absorption
triggers an appearance of the UV-induced band.
According to the Dorenbos model, the lowest-energy
5d level in the Eu2+ 4f 5 5d 1 configuration locates justbelow the bottom of the conduction band of host
SrAl2O4. Excited electrons at the lowest-energy 5dlevel can easily migrate to be trapped by latticeimperfections. Therefore, it is more likely that the UVinduced band is caused by the formation of electron
traps.
The UV-induced band was weakened when theeuropium concentration is increased from 0.1 to 1.0
mol%. Apparently, the efficiency of trap filling is
influenced by the concentration of Eu2+ ions. As theeuropium concentration becomes higher, thepenetration depth becomes smaller. If excited electrons
cannot migrate a long distance, they will be trapped bylattice imperfection in the vicinity of Eu2+ ions exited
by UV-light. Since the amount of such imperfections
becomes low with increasing europium concentration,
the UV-induced band is weakened as we observed.
The present study demonstrated that the trap filling
occurs most efficiently in the case of low Eu2+
concentration. This information gives a guiding
principle for the material design of high luminancelong-persistent phosphors. Unfortunately, we could
not exhibited UV-induced absorption change for
SrAl2O4:Eu2+ crystals with Eu2+ concentrations lessthan 0.1mol%, because of difficulties for adjustingsuch extremely low europium concentration. Thissubject is now under investigation.

Fig. 1. Absorption spectra of SrAl2O4:Eu crystals
with europium concentrations of 0.1 (a), 0.3 (b), and
1.0 mol% (c). These data were observed at 9 K. Blueand red lines indicate absorption spectra observed
before and after UV-irradiation at 375 nm.
[1] P. Dorenbos and J. J. Bos, Rad. Meas. 43 (2008)
139.
[2] Y. Kamiyanagi et al., J. Lumi. 122-123 (2007) 509.
[3] M. Kitaura, J. Ceram. Proc. Res. 12 (2011) s33.

76

Ⅲ -2. Materials Sciences

381BL6B 北浦守.docx

BL6B

Suppression of Electron Traps in Ce:GAGG Crystals by Mg2+ Co-Doping
1

R. Inaba1, M. Kitaura2, K. Kamada3, S. Kurosawa3, A. Ohnishi2 and K. Hara4

Graduate School of Science and Engineering, Yamagata University, Yamagata 990-8560, Japan2
Department of Physics, Faculty of Science, Yamagata University, Yamagata 990-8560, Japan3
New Industry Creation Hatchery Center, Tohoku University, Sendai 980-8579, Japan
4
Reserch Institute of Electronics, Shizuoka University, Hamamatsu 442-8529, Japan

Cerium doped Gd3Al2Ga3O12 (Ce:GAGG) crystals
have been developed as an inorganic scintillator withexcellent features such as high density, high lightyield, and short decay time [1]. The scintillation
properties are influenced by the existence of latticeimperfections, because they are responsible for theoccurrence of delayed luminescence [2]. In order toimprove scintillation properties, it is necessary to findout how to suppress lattice imperfections in theprocess of crystal growth. Therefore, the informationof crystal defects is indispensable. Recently, it wasreported that co-doping of alkali-earth ion is effective
for the improvement of scintillation properties [3].
However, the mechanism of such alkali-earth ion
co-doing has not yet been clarified. So far, we haveinvestigated infrared absorption spectra of Ce:GAGG
crystals under UV-irradiation at low temperatures [4].
The data were limited in the energy range of 0.05-1.2
eV. In the present study, we have extended themeasurement of absorption spectra to higher energy
side. Furthermore, we have investigated UV-induced
absorption change for Mg co-doped Ce:GAGG(Ce,Mg:GAGG) crystals.
Ce:GAGG and Ce,Mg:GAGG were grown by theCz method. The concentration of cerium andmagnesium ion was set to 1.0 and 0.1 mol%,respectively. Experiment was performed at thebeamline BL6B of UVSOR. The UV-irradiation was
carried out at 3.31 eV with a picosecond light pulser.
The sample temperature was set to 9K.
Figure 1 shows absorption spectra of Ce:GAGG
crystals. A black line represents the absorption
spectrum obtained before UV-irradiation. When the
UV-irradiation is turned on, the absorption spectrum
exhibits a broad band with the peak at 1.5 eV, as
indicated by a blue line. The UV-light can excite
electrons to a Ce3+ 5d level in the conduction band,and triggers the formation of electron traps [4]. In
addition, the Ce3+ 4f-4f peak around 0.5 eV is
weakened by 3 % under UV-irradiation at 9K. As theformation of electron traps is responsible for the
weakening of the Ce3+ 4f-4f peak, the concentrationof electron traps is estimated to be 1018 cm-3. Thisvalue is two orders of magnitude higher than theconcentration of lattice imperfections in typicalintrinsic semiconductors.
Figure 2 shows absorption spectra ofCe,Mg:GAGG
crystals
beforeandafter
UV-irradiation. The UV-induced band around 1.5 eV

disappears by Mg ion co-doping. This fact indicatesthat the Mg co-doping hinder the formation ofelectron traps. There are two hypotheses concerning
the rope of Mg co-doping: One is that Mg2+ ions play
as charge compensators to suppress latticeimperfections, and the other is that electron traps aredestabilized by controlling the energy band of host
crystals. Further experimental works are needed tosolve the role of Mg co-doping.

Fig. 1. UV-induced change of absorption spectra in
Ce:GAGG crystals measured at 9 K.

Fig. 2. UV-induced change of absorption spectra in
Ce,Mg:GAGG crystals measured at 9 K.
[1] K. Kamada et al., J.Cryst. Growth 352 (2012) 82.
[2] M. Kitaura et al., J. Appl. Phys. 115 (2014)083517.
[3] K. Kamada et al., Opt. Mater. 41 (2015) 63.
[4] R. Inaba et al., UVSOR Activity Report 42 (2015)
83.

77

Ⅲ-2

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原明人.docx
Ⅲ -2. Materials
Sciences

BL6B

Measurement Temperature Dependency of Ultrashallow Thermal Donors
and Observation of Shallow Donors in Carbon- and Hydrogen-Doped
Czochralski Silicon Crystals
A. Hara and T. Awano

Department of Electronic Engineering, Tohoku Gakuin University, Tagajo 985-8537, Japan
2.8

USTDs

STD(H)s

2.6

70 K

2.4
2.2

60 K

A
B

Peak intensity

2.0

C

50 K

1.8

45 K

1.6

40 K

1.4

35 K

1.2

30 K

1.0

25 K

0.8

21 K

0.6

18 K

0.4

15 K

0.2
0.0

50

100

150

200

Wave number (cm-1)

250

300

Fig. 1. Temperature dependency of optical absorption
peaks of USTDs and STD(H)s.
1.2
STD(H)s
USTDs

1.0

Normalized intensity

Ⅲ-2

We previously reported on ultrashallow thermal
donors (USTDs) in carbon- and hydrogen-doped
Czochralski silicon (CZ Si) crystals [1, 2]. USTDs arehydrogen-like donors with slightly different energy
levels, some of which show negative central-cellcorrection. To the best of our knowledge, these are theshallowest energy levels as compared with those ofpreviously reported donors in Si crystals. However, we
identified USTDs as shallow donors by a comparison
with effective-mass theory. In this study, we measured
the intensity of USTDs as a function of measurement
temperature to identify USTDs as shallow donors.
Carbon-doped CZ Si samples were doped withhydrogen by annealing in wet oxygen at 1300°C for 60
min. The samples were then cooled to room
temperature by rapidly exposing to air. For carbon
doping, the Si melt for preparing the ingot was dopedwith carbon powder during CZ Si crystal growth.
Transmission spectra were obtained with a far-IR
spectrometer of the BL6B beamline at different
temperatures using a flowing cryostat.
Figure 1 shows the temperature dependency of
USTDs and STD(H)s at different measurement
temperatures. Both spectra decrease as themeasurement temperature increases up to 70 K. Figure
2 shows variations in the peak intensity of USTDs and
STD(H)s against the measurement temperature. The
intensities of both defects show a similar dependency
on the measurement temperature, which indicates thatthe energy levels of both are very close. Since
STD(H)s are identified as shallow donors [3], it isreasonable that the energy levels of USTDs are very
close to those of STD(H)s. This is consistent with ourprevious proposal that USTDs are donors with energy
levels close to the conduction band.
In addition, we observed broad peaks at highmeasurement temperatures of around 50 K. They are
indicated as A, B, and C in Fig. 1. The differences inwavenumber between A and B, and B and C, areconsistent with the predictions of effective-mass
theory. Thus, the peaks A, B, and C are identified as
1s-2p0, 1s-2p+-, and 1s-3p+- transition of new shallow
donor, respectively.
In summary, we observed the measurement
temperature dependency of the spectra of USTDs. Thisbehavior is consistent with our previous results, whichwere derived from effective-mass theory. Moreover,
we found new broad optical absorption peaks, which isconsistent with the predictions of effective-mass
theory.

0.8
0.6
0.4
0.2
0.0

0

20

40

60

Temperature (1000/T (K))

80

Fig. 2. Temperature dependency of USTDs and
STD(H)s.
[1] A. Hara, T. Awano, Y. Ohno, and I. Yonenaga, Jpn.
J. Appl. Phys. 49 (2010) 050203.
[2] A. Hara and T. Awano, Jpn. J. Appl. Phys. 54 (2015)
101302.
[3] R. C. Newman, J. H. Tucker, N. G. Semaltianos, E.
C. Lightowlers, T. Gregorkiewicz, I. S. Zevenbergen,
and C. A. J. Ammerlaan, Phys. Rev. B 54 (1996)
R6803.

78

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401BL6B 柿本健一.docx

BL6B

Evaluation of Domain Wall in Alkali Niobate Piezoelectrics
by Far-Infrared Reflective Method
Y. Taniguchi, N. Matsubara, K. Yoshida, T. Fuchigami and K. Kakimoto

Graduate School of Engineering, Nagoya Institute of Technology, Nagoya 466-8555, Japan
It is generally accepted that domain structure is amajor factor for the origin of piezoelectric property,
and specifically that domain walls contributesignificantly to the properties. We evaluated domain
walls both for (Na,K)NbO3 crystal and ceramics byRaman spectroscopy and suggested that ferroelectricand piezoelectric properties at the walls might bedifferent far from those of adjacent domain regions [1].
In the previous work, we observed various permittivity
depending on the focused regions of (Na,K)NbO3
crystal by using mid-infrared as a light source [2].
However, there were some problems of the shortdistance in mapping measurement and of the low
intensity of Infrared(IR) reflectivity. Therefore, longdistanceandhigh-intensity
line
mappingmeasurements were performed in this work usingsynchrotron radiation source, and we tried to evaluate
domain walls especially for the permittivity in a stripe
domain structure.
(Na0.55K0.45)(Nb0.995Mn0.005)O3(NKN) crystal withstripe domains(20-80 μm width) was measured by a
FT-IR spectrometer (Bruker, VERTEX 70v). IR
measurements were performed vertically to thedomain wall every 10 μm in a total range of 100 μm.
The spectra were fitted to obtain permittivity
according to the following equation:ε (ω) = ε∞ + Σn ω pn / ω on - ω - nγnω222permittivity ε’ showed values between 65 and 70
except at the measurement positions of 30, 50 and 100
μm. Permittivity ε’ increased periodically in 20-50
measurement range, and the values at 30, 50 and 100
μm positions were approximately 80, which isapproximately 20% higher compared to the otherpositions. These results suggested that the permittivity
increases at the domain walls. Therefore, domain walls
could be evaluated on the basis of the increase inpermittivity by IR reflectivity measurement. In
addition, we think that ferroelectric and piezoelectric
properties remarkably improve with increasing thedensity of domain wall in the materials.

Fig. 1. IR reflectivity spectra of NKN crystal obtainedusing synchrotron radiation and mid-infrared as lightsource.

(1)where, ε∞ is the high-frequency dielectric constant, ωpand ωo the plasma and longitudinal frequencies, and γ
the damping constant. The complex dielectric function
is related to the reflectivity spectrum by the equation
R = |√ε (ω) – 1 / √ε (ω) + 1|2(2)

Figure 1 shows IR reflectivity spectra of NKN
crystal. The reflectivity was 0.6 when using midinfrared. On the other hand, the reflectivity
significantly increased by using synchrotron radiation,and it was 0.8 at a maximum even in 0.25 aperture
diameter. The reason is that the synchrotron radiation
had a more than 6 times larger reflection intensity in Si
bolometer than that of mid-infrared. Hence, IR
reflectivity measurement by synchrotron radiation
should give reliable data of permittivity.
Considering the width of the stripe domains in the
NKN crystal (20-80 μm), permittivity ε’ should change
periodically in 20-80 μm measurement range owing to
the effect of domain wall. Figure 2 shows thepermittivity of NKN crystal at each measurement
position with an aperture diameter of 0.25 mm. The

Fig. 2. Permittivity ε’ of NKN crystal at measurement
positions obtained using synchrotron radiation as lightsource.
[1] Y. Taniguchi and K. Kakimoto, Japanese Journal ofApplied Physics 54 (2015) 10ND09.
[2] Y. Taniguchi, T. Yamazaki, M. Kato and K.
Kakimoto, UVSOR Activity Report 42 (2015) 85.

79

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高橋隆.docx
Ⅲ -2. Materials
Sciences

BL7U

High-Resolution ARPES Study of Noncentrosymmetric
Weyl Semimetal Candidate NbP
1

Ⅲ-2

S. Souma1, Z. Wang2, H. Kotaka2, T. Sato3, K. Nakayama3, Y. Tanaka3, H. Kimizuka3,
K. Yamauchi2, T. Oguchi2, K. Segawa4, Y. Ando2,5 and T. Takahashi1,3

WPI Research Center, Advanced Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan
2
Institute of Scientific and Industrial Research, Osaka University, Ibaraki 567-0047, Japan3
Department of Physics, Tohoku University, Sendai 980-8578, Japan
4
Department of Physics, Kyoto Sangyo University, Kyoto 603-8555, Japan
5
Institute of Physics II, University of Cologne, Köln 50937, Germany

Weyl semimetals (WSMs) manifest a novel
quantum state of matter where the bulk conduction andvalence bands cross at discrete points, defined as Weylnodes, with linear dispersion in all the momentum (k)directions in three-dimensional Brillouin zone, which
can be viewed as a three-dimensional analogue ofgraphene breaking time-reversal or space-inversion
symmetry [1,2]. The WSMs can host many exotic
physical phenomena such as anomalous Hall effects,chiral anomalies, and magnetoelectric effects. The
most intriguing prediction for WSMs is the emergenceof Fermi arcs on their surfaces. Unlike twodimensional metals showing closed Fermi surface (FS),the Fermi arcs in WSMs are disjoint, open curves.
Recently, density functional theory predicted that
noncentrosymmetric transition-metal monopnictide
family TaAs, TaP, NbAs, and NbP are WSMs [3,4]. To
firmly establish the WSM nature of monopnictides andto build a basis for the proposed exotic phenomena, it
is of particular importance to experimentally establish
the fermiology of transition-metal monopnictide
family.
In this study, we have performed high-resolution
angle-resolved photoemission spectroscopy (ARPES)
of NbP. By utilizing the low photon energy andvariable polarization characteristics of the beamline
BL-7U, we have succeeded in observing the fineelectronic structure that supports the WSM nature of
NbP [5].
Figure 1(a) displays a schematic summary of theexperimentally observed FS of NbP aroundX. Theenergy bands labeled here, S1-S4, all of which aresurface states, obey different selection rules ofphotoelectron intensity. For example, band S1 isdominantly seen with vertical polarization [Fig. 1(b)].
On the other hand, the intensity of outer band S2, whichforms a tadpole FS, is greatly enhanced with circularly
polarized photons [Fig. 1(c)]. There exist another dogbone-shaped FSs, S3 and S4, elongated alongXM,
which are better resolved with horizontally polarized
photons [Fig. 1(d)]. We found that overall crossshaped FS in the experiment is reasonably reproducedby our calculation and the FS (S2) which connects
Weyl nodes W2± can be viewed as a combination of
two Fermi arcs. Namely, one arc starts and ends at W2,
and the other is connected to W1. It is thus likely that

only S2 forms a Fermi arc and all the others (S1, S3, andS4) form trivial FS.
To examine the WSM nature of NbP, we choose
closed k-loop (Γ-X-M-Γ) surrounding the odd
number of (three) Weyl nodes and counted the total
number of FS crossings. Since only an open Fermi arc
can cross this loop an odd number of times, the total
odd number of FS crossings would be a hallmark of
the existence of Fermi arcs in WSM. As seen from Fig.
1(a), the Fermi arc (S2) crosses this loop only once, andthe trivial FSs (S1, S3, and S4) cross six times.
Therefore the observed FSs cross the k-loop seventimes in total, supporting the WSM nature of NbP. The
present result opens a pathway for the Fermi-arc
engineering of Wey semimetals.

Fig. 1. (a) Experimental FS aroundX of NbP. (b)-(d)Second-derivative intensity of the ARPES spectra
measured with vertically-, circularly-, andhorizontally-polarized
photons,
respectively.
Polarization vector of lights is indicated by arrow.
[1] X. Wan et al., Phys. Rev. B 83 (2011) 205101.
[2] L. Balents, Physics 4 (2011) 36.
[3] H. Weng et al., Phys. Rev. X 5 (2015) 011029.
[4] S.-M. Huang et al., Nature Commun. 6 (2015) 7373.
[5] S. Souma et al., arXiv:1510.01503. (Phys. Rev. B
Rapid Commun., in press)

80421BL7U 伊藤孝寛.docx

Ⅲ -2. Materials Sciences

BL7U

Angle-Resolved Photoemission Study of the Quasi-One-Dimensional
Organic Conductors (TMTTF)2X (X = AsF6, SbF6)
T. Ito1,2, A. Kawakita1, T. Hajiri1, K. Tanaka3,4, S. Kimura5,6 and T. Nakamura3,4

1
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan
Nagoya University Synchrotron Radiation Research Center, Nagoya University, Nagoya 464-8603, Japan
3
Institute for Molecular Science, Okazaki 444-8585, Japan
4
School of Physical Sciences, SOKENDAI (The Graduate School of Advanced Studies), Okazaki 444-8585, Japan
5
Graduate School of Frontier Biosciences, Osaka University, Osaka 565-0871, Japan
6
Department of Physics, Osaka University, Osaka 565-0871, Japan
2

The Fabre salts (TMTTF)2X and Bechgaard salts(TMTSF)2X (X = PF6, AsF6, etc.) belong to a familyof quasi-one-dimensional organic conductors. As thechemical pressure of the system tuned with the
combination of TMTTF/TMTSF and anion X, thesematerials allow for a variety of ground states, such as
antiferro magnetic insulator, spin-Peierl state, spindensity wave state and superconductor; furthermore
one observes a crossover from a Luttinger liquid
toward a Fermi-liquid metal, charge-ordred (CO) andcharge-localized (loc.) insulator [1, 2].
In this study, we have performed temperaturedependent
angle-resolved
photoemission
spectroscopy(ARPES)onsingle-crystalline(TMTTF)2X (X = AsF6, SbF6) to clarify the relation
between the electric structure and the thermodynamic
properties. ARPES measurement were performed at
UVSOR-Ⅲ BL7U by utilizing the bulk-sensitive low
photon energy (hν = 8 eV) as well as the micro focus
beam (15×100 μm2).
Figures 1 (A1-A8) and (S1-S6) show thetemperature dependence of ARPES image along theΓX direction of (TMTTF)2AsF6 and (TMTTF)2SbF6,
respectively. The schematic illustrations of thecorresponding electronic structure have been shownin Fig. 1 (M1) and (M2) for comparison. At low
temperature CO states (T = 30 K < TCO = 102 and 157
K for X = AsF6 and SbF6, respectively [3]), we foundhighly dispersive band away from the Fermi level(EF) which is accompanied with the lower energy
shoulder around X point, and that the band foldedaround the top of the dispersive features (kF ~ 0.25
and 0.2 Å-1 for X = AsF6 and SbF6, respectively).
From the comparison with the previous ARPES
results of metallic (TMTSF)2X [4, 5], we haveexpected that the observed dispersive features can be
ascribed as the spinon and holon bands, which ischaracterized by the spin-charge separation in the
one-dimensional system (Figs. 1 (M1) and (M2)).
With increasing temperature, the main dispersive
feature of X = AsF6 and SbF6 gradually becomes
weak around T = 50 – 150 K and T = 100 – 200 K,
respectively. On the other hand, at T = 200 – 300 K,
the broad non-dispersive feature (shaded area in
Figs.1 (M1) and (M2)) around 1 eV dominates theelectronic structure of both compounds, though theslight dispersive feature remains in X = AsF6. The
observed temperature dependence seems to roughly
scale with the charge-order transition temperature ofeach compound. Therefore, we expect that theobserved spectral weight transfer from the dispersive
spinon-holon bands near EF to the localized boundstate is caused by the transition from charge-orderedinsulator to charge-localized state. It should be notedthat the metallic states of (TMTTF)2AsF6 at T = 300
K have been found as a sizable spectral weight at EF
originating from the tail of the localized bound states.

Fig. 1. Temperature dependence of ARPES imagesof (TMTTF)2AsF6 (A1-A8) and (TMTTF)2SbF6(S1-S6) along ΓX direction. (M1, M2) The schematic
picture of spinon-holon dispersions (solid and dashedlines) and localized bound state (shaded area) of(TMTTF)2AsF6 (M1) and (TMTTF)2SbF6 (M2),respectively.
[1] T. Ishiguro et al., Organic Superconductors(Springer-Verlag, Berlin, 1998).
[2] E. Rose et al., J. Phys.: Condens. Mat. 25 (2013)
014006.
[3] M. Dressel et al., Adv. Cond. Mat. Phys. 2012(2012) 398721.
[4] S. Hirate et al., UVSOR Activity Report 39(2012) 97.
[5] M. Mitamura et al., UVSOR Activity Report 40(2013) 110.

81

Ⅲ-2

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BL7U

Systematic Change of Fermi Surfaces with As Doping in Iron Pnictide
Superconductor NdFeP1-xAsx(O,F)
S. Miyasaka1, A. Takemori1, T. Adachi1, M. Uekubo1, T. Kobayashi1, S. Tajima1, T. Hajiri2,3,
M. Matsunami3,4 and K. Tanaka3,4

Department of Physics, Graduate School of Science, Osaka University, Toyonaka 560-0043, Japan
Department of Crystalline Materials Science, Graduate School of Engineering, Nagoya University, Nagoya
464-8603, Japan
3
UVSOR Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
4
School of Physical Sciences, The Graduate University for Advanced Studies (SOKENDAI), Okazaki 444-8585,
Japan
1

2

Ⅲ-2

x = 1.0, however, the quietly different Fermi surfaces
exist. As shown in Figs. 1(a) and (d), this sample has
the electron Fermi surface with circle shape and thepropeller-like one. These clear changes of Fermi
surfaces around zone center and boundary at x ~ 0.8
are related with the change of transport and otherproperties, and supports the scenario of two Tc-risingmechanisms and two Fermi surface states in RFeP1xAsx(O,F).

The iron-based superconductors have various Fermi
surface states, which are closely related with the local
crystallographic structure around Fe ion. Recently, we
studied the physical properties in 1111-type P/As solid
solution system, RFeP1-xAsxO0.9F0.1 (R = La, Pr andNd). In these systems, the local structural parameters
coupled with Fermi surface conditions are
systematically varied by P/As solid solution. Our
results have revealed that these systems have twodifferent Fermi surface states and nesting conditions
below and above x = 0.6~0.8 [1, 2]. To observe thechange of Fermi surfaces by P/As solid solution
directly, we have investigated the angle resolved
photoemission spectroscopy using single crystals of
NdFeP1-xAsx(O,F) with x = 0.4, 0.8 and 1.0 for F
concentration y ~ 0.1.
The single crystals of NdFeP1-xAsx(O,F) weregrown by high pressure technique using cubic anvilpress and self-flux method [3]. Tc of the single crystals
in the present work were 12 K for x = 0.4, 23 K for x =0.8 and 44 K for x = 1.0, respectively. The angle
resolved photoemission spectroscopy were measured
at BL7U of UVSOR facility in Institute for Molecular
Science using incident photons with hν = 36 eV
linearly polarized.
As shown in Figs. 1(a)-(f), the results of angle
resolved photoemission spectroscopy indicate theexistence of three hole and two electron Fermi surfaces
in x = 0.4, and two hole and two electron ones in x =0.8 and 1.0 samples, i.e. the change of FS state aroundx = 0.8. The orbital characters of these Femi surfaces
and electronic bands have been assigned bypolarization-dependent angle resolved photoemission
spectroscopy, and the dxz hole band around zone center
particularly shows the systematic and important x
dependence. As shown in Fig. 1(c), the dxz hole Fermi
surface exists around zone center at x = 0.4. The energy
level of dxz band decreases with increasing x, and thisband almost sinks below Fermi energy in the x = 0.8
sample. With further increasing x, this dxz Fermi
surface completely disappears at x=1.0 (See Fig. 1(a).).
Other distinct change of Fermi surfaces has been
observed around zone boundary. In the x = 0.4 and 0.8
samples, we observed two electron Fermi surfaces
with circle or elliptic shapes around zone boundary. In

Fig. 1. Fermi surface mapping on kx-ky plane of
NdFeP1-xAsx(O,F) with x = 1.0 (As100) ((a) and (d)), x= 0.8 (As80) ((b) and (e)) and x = 0.4 (As40) ((c) and(f)), respectively. The upper and lower panels show theresults by using P- and S-polarization configurations,
respectively. The insets of panel (d)-(f) show the detailof Fermi surface mapping around zone boundary in the
S-polarization configuration. The red lines indicate theobserved Fermi surfaces.
[1] S. Miyasaka et al., J. Phys. Soc. Jpn. 82 (2013)
124706.
[2] K. T. Lai et al., Phys. Rev. B 90 (2014) 064504.
[3] A. Takemori et al., JPS Conf. Proc. 1 (2014)
012111.

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Ⅲ -2. Materials Sciences

BL7U

Electronic Structure of Sr1-xCaxFe2(As1-yPy)2 (x = 0.08, y = 0.25)
Revealed by Angle Resolved Photoemission Spectroscopy
T. Adachi1, S. Ideta2, 3, K. Tanaka2, 3, S. Miyasaka1 and S. Tajima1

Graduate School of Science, Department of Physics, Osaka University, Toyonaka 560-0043, Japan
2
UVSOR Facility, Institute for Molecular Science, Okazaki 444-8585, Japan3
School of Physical Sciences, The Graduate University for Advanced Studies (SOKENDAI), Okazaki 444-8585,
Japan1

It is well known that there is a remarkablecorrelation between a crystal structure and Tc in ironbased superconductor. Thus, it is essential forunraveling the superconducting (SC) mechanism inthis system to clarify how the electronic structurechanges with the crystal structure. Recent ARPES
study has reported that in SrFe2(As1-yPy)2 whose c/a(the ratio of a- and c-axes lattice constants), which isindex of structural anisotropy, is smaller than that ofBaFe2(As1-yPy)2, innermost dxy hole Fermi surface
shrinks and splits into two pockets, while in
BaFe2(As1-yPy)2, all the hole Fermi surfaces areconnected in the entire kz region [1, 2]. The reason isconsidered to be the stronger three-dimensionality(3D) in SrFe2(As1-yPy)2 originates from enhanced
interlayer hopping matrix elements due to the smaller
c-axis lattice constant and thus structural anisotropy(c/a).
In this measurement, we observed the electronic
structure of Sr1-xCaxFe2(As1-yPy)2 (x = 0.08, y = 0.25,
Tcmax=32K) whose c/a is smaller than that of
SrFe2(As1-yPy)2 and whose Fermi surfaces are
considered to be more three-dimensional. A
momentum distribution curve (MDC), a E-k intensity
plot and a second derivative plot of hole Fermi
surfaces around Γ point measured by using linearizeds-polarized light are shown in Fig. 1. As seen in Fig. 1,
around Γ point, the hole band whose orbital characteris assigned to dxy sinks below EF, consistent with the
ARPES result of SrFe2(As1-yPy)2. Figure 2 shows SC
gaps around Γ and X points as a function of θFS in
Sr1-xCaxFe2(As1-yPy)2 (x = 0.08, y = 0.25). SC gaps ofhole Fermi surface around Γ point are isotropic within
the error bar. On the other hand, SC gaps of electron
Fermi surface around X point show a different
behavior. The SC gap for outer αelectron Fermi
surface decreases rapidly toward θFS ~ 40 deg and at
θFS ~ 40 deg, node of SC gap is observed. Moreover,
the SC gap for inner β electron Fermi surface
decreases gently with θFS increasing. These SC gap
anisotropy are different from those of BaFe2(As1-yPy)2(y=0.30) [3] and it suggests that there is a relationshipbetween SC gap anisotropy and structural anisotropy(c/a).

Ⅲ-2

Fig. 1. The
electronic
structureof
Sr1-xCaxFe2(As1-yPy)2 (x = 0.08, y = 0.25) around Γpoint. (a) MDC at EF, (b) E-k intensity plot, and (c)second derivative plot.

(a)

(b)

hν = 26 eV
T = 12 K(c)(d)

Fig. 2. (a), (c) SC gaps around Γ and X points as afunction of θFS in Sr1-xCaxFe2(As1-yPy)2 (x = 0.08, y =0.25). (b), (d) kF around Γ and X point where SC gaps
were measured. The Fermi angle is defined so that theΓ-X line is θFS = 0 deg.
[1] T. Yoshida et al., Phys. Rev. B 106 (2011) 117001.
[2] H. Suzuki et al., Phys. Rev. B 89 (2014) 184513.
[3] T. Yoshida et al., Sci. Rep. 4 (2014) 7292.

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Sciences

BL7U

Superconducting Gap of the Electron-Doped Cuprate Superconductor
Pr1.3-xLa0.7CexCuO4 Studied by Angle-Resolved Photoemission Spectroscopy
M. Horio1, K. Koshiishi1, T. Ohgi2, S. Ideta3, K. Tanaka3,
T. Adachi4, Y. Koike2 and A. Fujimori1

1

Department of Physics, University of Tokyo, Tokyo 113-0033, Japan.
Department of Applied Physics, Tohoku University, Sendai 980-8579, Japan.
3
UVSOR Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
4
Department of Engineering and Applied Sciences, Sophia University,
Tokyo 102-8554, Japan.

2

Ⅲ-2

The momentum dependence of the superconducting(SC) gap carries rich information about the mechanismof superconductivity. While it is generally acceptedthat the SC gap of hole-doped cuprates has d-wave
symmetry, the symmetry for electron-doped cuprates
is still controversial.
The most direct probe of the symmetry of the SC
gap is angle-resolved photoemission spectroscopy(ARPES). Although rapid surface degradation and thesmall size of the SC gap have prevented thoroughinvestigation of the momentum dependence [1,2], an
ARPES study on optimally doped electron-doped
cuprates reported that the symmetry was basically dwave but the gap took a maximum around the hot spot
where the Fermi surface and antiferromagnetic (AF)Brillouin zone boundary cross, suggesting large
contribution of AF spin fluctuations to thesuperconductivity [3]. However, some studies haveraised the possibility that the gap symmetry changes
from d-wave to s-wave by electron doping [4].
Recently, a new annealing method, which is called
protect annealing and conducted to remove impurity
oxygen that stabilizes AF order and destroys
superconductivity, has enabled superconductivity inelectron-doped cuprates with less Ce concentration [5].
An ARPES study has revealed that sufficient protect
annealing strongly suppresses the quasi-particle
scattering at the hot spot while the SC transition
temperature (Tc) remains high [6]. In order to reveal
the momentum dependence of the SC gap of protectannealed samples, we have performed ARPES
measurements with higher energy resolution.
Pr1.3-xLa0.7CexCuO4 (x = 0.10, PLCCO) single
crystals were grown by the traveling-solvent floatingzone method. We prepared two PLCCO samples whichshowed Tc’s of 26.3 K (sample #1) and 26.9 K (sample#2) after annealing at 800 °C for 24 h. ARPES
measurements were performed at UVSOR BL7U.
Linearly polarized light with hν = 16.5 eV was used.
The total energy resolution was set at 8 meV. Sample#1 was measured once and Sample #2 for three times.
Prior to each measurement, the samples were cleaved
in situ under the pressure ~ 5×10-11 Torr.
Figure 1 shows energy distribution curves (EDCs)
near (0, π) taken below and above Tc. Below Tc, theleading edge was shifted by 1.2-1.5 meV towardhigher binding energy with good reproducibility,
suggesting the opening of SC gap, although theamount of shift was a little smaller than 1.9 meV which
was reported in previous studies [2,3]. Unfortunately,
we could not reveal further detailed momentum
dependence during this beam time, and additional
measurements are planned in the near future.

Fig. 1. SC gap of PLCCO. EDCs taken below andabove Tc for (a) sample #1 and (b)-(d) sample #2. (b)to (d) are results of the first to third measurements,
respectively. All the EDCs were taken at the
momentum indicated by a green point in the inset of(a).
[1] T. Sato et al., Science 291 (2001) 1517.
[2] N. P. Armitage et al., Phys. Rev. Lett. 86 (2001)
1126.
[3] H. Matsui et al., Phys. Rev. Lett. 95 (2005) 017003.
[4] J. A. Skinta et al., Phys. Rev. Lett. 88 (2002)
207005.
[5] T. Adachi et al., J. Phys. Soc. Jpn. 82 (2013)
063713.
[6] M. Horio et al., Nat. Commun. 7 (2016) 10567.

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BL7U

Electron-Phonon Scattering in Graphene: Probing of the PhononDispersion by the Angle-Resolved Photoelectron Spectroscopy
S. Tanaka

The Institute of Scientific and Industrial Research, Osaka University, Ibaraki 567-0047, JapanThe electron-phonon coupling rules many
properties of solids, and have been extensively studied.
Recently, we proposed a new experimental method
which enables us to probe the electron-phonon
scattering, which is the elemental process of theelectron-phono coupling, in HOPG graphite by using
the angle-resolved photoelectron spectroscopy(ARPES). Here, we report an investigation using thesame method for the epitaxial graphene on SiC.
The sample used was the single-layer graphene on
the SiC crystal manufactured by Graphene Platform
Inc. (Tokyo, Japan). ARPES measurements werecarried out at the BL7U at the UVSOR facility in
Institute for Molecular Science. Figure 1(a) shows thephoton-energy dependence of the surface normal
photoelectron spectra of the graphene. There are steps
near 160meV of the binding energy at the photon
energy of near 10.8eV. These steps are converted to thepeak shapes by differentiating the spectra with thebinding energies as shown in Fig. 1(b). The heights ofthe steps are plotted as a function of the photon energy[Fig. 1(c)]. There are no electronic band near the Γpoint in graphene, and in fact these steps are the Fermiedge which are shifted as much as the energy of thephonon which is scatters the electron from the Diraccone at the K-point to the Γ-point of the unoccupied
band located at 10.3eV above the Fermi level. Theenergy shift is due to the energy conservation ruleduring the phonon emission process of the electron.
Possible phonon modes are LO- and TO-branches
whose calculated dispersions are shown in Fig. 2(b).
The ARPES map whose intensities are differentiatedwith the binding energy are shown in Fig. 2(c) for two
directions Γ-K and Γ-M). Obviously, the energies
shift as a function of the electron-momentum. This isdue to the momentum conservation rule, and these
relate to the phonon dispersion. It should be noted thatthe electron momentum at the Fermi level of the Diraccone is somewhat ambiguous since Fermi circle isformed in the case of the graphene/SiC due to thesubstrate effect. Therefore, a convolution with aphonon momentum should be taken into account. Fig.
2(d) shows simulation for the scattering processincluding the broadening induced by the experimental
resolution. The TO-phonon is assumed in thesimulation, and the agreement between theexperiments and the calculation are good. The
simulation assuming LO-phonon (not shown here)cannot reproduce the experiments at all, indicating that
the TO-phonon is responsible for the electron-phonon
scattering in this case. Although this suggests a highefficiency of the TO-mode for the electron-phonon
coupling, what observed here is the electron-phonon
scattering between specific wave functions of thegraphene, and detailed comparison with the theoretical
calculation including the matrix elements both of theelectron-phonon scattering and the optical transition
will be necessary for more detailed discussion.

Ⅲ-2

Fig. 1. (a) Surface normal ARPES spectra, (b) theirdifferentiation, and (c) step heights in the spectra as afunction of the photon energy.

Fig 2. (a) Schematic drawing of the electron
emission process, (b) phonon dispersions of graphene,(c) diffrentciated ARPES maps along the -K (upper)and -M (lower) directions at hν= 10.8eV (d)Simulation of the electron-phonon scattering ingraphene.

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BL7B

Electronic Structure and Decay Profiles of Auger-Free Luminescencefrom α- and β- BaLu2F8 Crystals
H. Tsuchida1, S. Tanaka2, M. Kitaura1, A. Ohnishi1, J. Pejchal3, A. Yamaji4 and S. Kurosawa3
Faculty of Science, Yamagata University, Yamagata 990-8560, Japan
Faculty of Science and Engineering, Kinki University, Higashiosaka 577-8502, Japan3
New Industry Creation Hachery Center, Tohoku University, Sendai 980-8579, Japan
4
Institute for Materals Research, Tohoku University, Sendai 980-8679, Japan
1

2

Ⅲ-2

Auger-Free Luminescence (AFL) is caused byradiative transition of core holes to the valence band[1]. The spectral shape and radiative lifetime of AFLs
mainly reflect electronic structure such as density ofstates (DOS) and wave functions of valence andoutermost core bands, which are dependent on local
ionic or covalent bonds. AFLs of compounds withdifferent modifications are very much interesting incontrolling AFLs. From this viewpoint, we have muchattention to the AFL of BaLu2F8. This compound has
two modifications of α- and β-phases [2]. The former
and latter belong to orthorhombic and monoclinic
systems, respectively. The α- and β-phases are stable
above and below 900 oC, respectively. In the present
study, we have studied electronic structures of α- andβ-BaLu2F8 crystals by measurements of x-ray
photoelectron spectroscopy (XPS), and reflectancespectra. Furthermore, we have measured emission
spectra and decay profiles of Auger-free luminescence(AFL) from α- and β-BaLu2F8 crystals under the single
bunch operation of UVSOR.
Figure 1 shows reflectance (red) and absorbance(blue) spectra of α- and β-BaLu2F8 crystals, whichwere measured at room temperature. The absorption
edge of α-BaLu2F8 is higher than that of β-BaLu2F8.
Furthermore, the reflectance spectrum of α-BaLu2F8 isclearly different from that of β-BaLu2F8. On the other
hand, it was confirmed that the XPS spectra are in good
agreement with each other, indicating that theelectronic structures of valence and outermost core
bands are almost same between them. Therefore, thedifference between reflectance and absorbance spectra
is attributed to that in the structure of the conduction
band.
From the analysis of XPS spectra, the energy
difference between the top of the valence band and thatof the outermost core band was determined to be 7.9
eV, the value of which is smaller than the band-gap
energy estimated from the higher energy side of thelowest reflectance peak. Therefore, it is evident that αand β-BaLu2F8 crystals accompany radiative core holetransitions between the valence band and the top of theoutermost core band, resulting in AFLs. Actually,
emission spectra exhibited an AFL band around 5 eV.
The spectral shape is same between α- and β-BaLu2F8
crystals. Decay profiles of the AFLs were also same,as shown in Fig. 2. The lifetimes of AFLs were
determined to be 610 ps (α-BaLu2F8) and 560 ps (β-BaLu2F8), respectively. Apparently, similarities inemission spectra and decay curves of AFLs reflect theresults of XPS spectra.

Fig. 1. Reflectance (red) and absorbance (blue)
spectra of α- and β-BaLu2F8 crystals, measured at room
temperature.

Fig. 2. Decay profiles of AFLs from α- and βBaLu2F8 crystals, measured at room temperature under
excitation at 21.4 eV. A red line indicates the pulse
profile of an excitation light pulse.
[1] M. Itoh, Recent Res. Devel. Physics 4 (2003) 205.
[2] A. A. Kaminskii et al., J. Alloys Compd. 275-277(1998) 442.

86Ⅲ -2. Materials Sciences

491BL7B 越水正典.docx

BL7B

Analysis of the Luminescence Properties of Gd2Si2O7:Ceand Gd2Si2O7:La, Ce with Vacuum Ultraviolet Excitation
M. Koshimizu1, T. Yanagida2, Y. Fujimoto1 and K. Asai1

Department of Applied Chemistry, Graduate School of Engineering, Tohoku University, Sendai 980-8579,
Japan
2
Graduate School of Materials Science, Nara Institute of Science and Technology, Ikoma 630-0192, Japan1

Ce-doped Gd2Si2O7 (GPS) has been attracting
considerable attention owing to its excellent
scintillation properties [1]. Recently, it is proved that
La-doing is an effective method to enhance theperformance [2]. In this study, we analyzed theluminescence properties of Ce-doped GPS with andwithout La with vacuum ultraviolet (VUV) excitation
to elucidate the effect of La incorporation on theenergy transfer process from the host matrix to Ce3+
ions.
The single crystals of Ce-doped GPS with andwithout La incorporation were grown via floating-zone
method. The luminescence properties werecharacterized at BL7B of UVSOR. Luminescence
spectra were obtained using an optical multichannel
analyzer equipped with a liquid-nitrogen-cooled CCD
at different excitation wavelengths. For themeasurements of luminescence time profiles, single
bunch operation of UVSOR was used. The
luminescence photons were monochromatized anddetected
withamultichannel-plate-mounted
photomultiplier tube.
Figure 1 shows the luminescence spectra of the Cedoped GPS crystals with and without La incorporation.
A prominent band was observed at 370–400 nm for all
the crystals, and is attributed to the 5d-4f transition of
Ce3+.
Figures 2 and 3, respectively, show theluminescence time profiles of the Ce-doped GPS
crystals with and without La incorporation withexcitation at 60 and 300 nm. The decay behavior was
expressed satisfactorily with single exponentialfunctions for all the time profiles. For GPS doped with
Ce and La at 1 and 10 mol% respectively, the decay
time constants were 52 and 60 ns with excitation at 300
and 60 nm, respectively. On the other hand, the decay
time constants were 41 and 52 ns with excitation at 300
and 60 nm, respectively. The excitation at 300 nm
corresponds to the direct excitation of Ce3+, whereas
the excitation wavelength at 60 nm corresponds to thehost excitation. The difference in the luminescenceprocess is that energy transfer is involved in theluminescence process for the excitation wavelength of
60 nm. It is clearly seen that the GPS crystal with La
incorporation has a similar decay behavior for
excitation wavelengths of 60 and 300 nm, whereas aconsiderable difference was observed for the GPS
crystal without La. This result indicates that the La
incorporation has an influence on the energy transfer
process.

Ⅲ-2

Fig. 1. Luminescence spectra of the Ce-doped GPS
crystals with and without La incorporation.

Fig. 2. Luminescence time profiles of GPS dopedwith Ce and La at 1 and 10 mol%, respectively.

Fig. 3. Luminescence time profiles of GPS dopedwith Ce at 10 mol%.
[1] S. Kawamura, J. H. Kaneko, M. Higuchi, F. Fujita,
A. Homma, J. Haruna, S. Saeki, K. Kurahige, H.
Ishibashi and M. Furusaka, Nucl. Instrum. Methods
Phys. Res. A 583 (2007) 356.
[2] S. Kurosawa, T. Shishido, A. Suzuki, J. Pejchal, Y.
Yokota and A. Yoshikawa, Nucl. Instrum. Methods
Phys. Res. A 744 (2014) 30.

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BL7B

Influence of Barium Concentration on the Reflection Spectra for Ternary
Niobium Bismuth Phosphate Glasses
1

National Institute of Advanced Industrial Science and Technology, Osaka 563-8577, Japan
2
Kansai University, Osaka 564-8680, Japan

BaO-Nb2O5-P2O5 glass is one of niobium phosphateglass systems having a wide glass forming region withhigh Nb2O5 contents [1]. In this system, NbOx
polyhedra construct glass network structure with PO4
tetrahedra. Since the glass shows high refractive index
due to high niobium concentration. It is of interest for
the application to optical glasses in precision opticaldevices. However, the glass with high niobium content
gives yellowish and blue colours and the relationshipbetween optical properties, such as refractive index,
dispersion, absorption edge, etc., and structure aroundniobium ions or connection of niobium-oxygen
polyhedra in the network structure are not clear yet. In
the present study, reflectivity of the bismuth phosphateglasses has been measured to investigate
compositional dependence of the band structure in thevacuum ultraviolet region.
Ternary xBaO-10Nb2O5-(90-x)P2O5 (x = 10, 20, 30,
40, 50 and 55 mol%) glasses and xBaO-20Nb2O5-(80x)P2O5 glasses (x = 10, 20, 30, 40, 50 and 55 mol%)were prepared by using a conventional liquid-quench
method. Reflectivity of optically polished samples wasmeasured in the photon energy region of 4-25 eV withan incident angle of 10 at the BL7B. Figure 1 showsreflection spectra of the xBaO-10Nb2O5-(90-x)P2O5
glasses. Strong reflection peak was observed at around
9.5eV and 22 eV for all the glasses. The peak areassigned to the transition from O 2p and P 3p states to
Nb 4d and Ba 5d states from a report of DFT
calculation for barium niobium phosphate compounds[2]. The origin of the higher band at 22 eV is not clear
yet. The unknown peak at 22 eV increased withincreasing BaO content. Therefore, the peak could be
due to the state of Ba2+. Contrary, intensity of the 9.5
eV peak did not show remarkable change with the BaO
content. However, the peak position shifted toward
lower energy by about 0.3 eV, while the band width(HWHM) decreased by about 0.6 eV from an analysis
of multi-Gaussian deconvolution for the K-K
transformed absorption spectra. These changes
predicted a higher energy shift of absorption edgeenergy and the prediction was consistent with theexperimental result of absorption spectra in visible
region [3].
On the other hand, reflection spectra of ternaryxBaO-20Nb2O5-(80-x)P2O5 glasses are shown in Fig.
2. As shown in the figure, the similar peaks wereobserved. The peak around 22 eV increased withincreasing BaO content. The intensity of the 9.5 eVband was almost constant against the BaO content too.

This work was carried out as a study of CSTI, SIP
Project (Development of Advanced Glass Processing
Technologies).

0.30

xBaO-10Nb2O5-(90-x)P2O5 glass

Reflectivity

0.25

55

0.20

50
40

0.15

30

0.10

20

0.05

x=10

10

15

20

25

Photon energy/ eVFig. 1. Reflection spectra of ternary xBaO10Nb2O5- (90-x)P2O5 glasses.

Reflectivity

Ⅲ-2

N. Kitamura1, K. Fukumi1, Y. Hisano2, A. Uchiyama2 and H. Kozuka2

0.30

xBaO-20Nb2O5-(80-x)P2O5 glass

0.25

55

0.20

50
40

0.15

30

0.10

20

0.05

x=10

10

15

20

Photon energy/ eV

25Fig. 2. Reflection spectra of ternary xBaO20Nb2O5- (80-x)P2O5 glasses.
[1] N. Shcheglova and T. V. Avlas, Fiz. Khim. Stekla
16 (1990) 879.
[2] I. Cho et al., Eur. J. Inorg. Chem. 2011(2011) 2206.
[3] N. Kitamura et al., Extended abstract 1st Int. Nat’l.
Conf. Phosphate glasses, Pardvice, 2014.

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511BL7B 藤岡淳.docx

BL7B

Optical Conductivity Spectra in the Course of Bandwidth-Control MetalInsulator Transitions in Nd1-xPrxIr2O7
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J. Fujioka1, K. Ueda1 and Y. Tokura1,2

Department of Applied Physics, University of Tokyo, Tokyo 113-8656, JapanRIKEN Center for Emergent Matter Science (CEMS), Wako 351-0198, Japan

Strongly correlated electron system has been one of
the central topics in the condensed matter physics
where the interplay among charge, spin, orbital degreeof freedom frequently induces a variety of intriguingstates. Among them, interplay between electron
correlation and relativistic spin-orbit coupling in 5dtransition metal oxides has recently attracted muchattention for the realization of the correlated
topological insulator/semimetals [1]. The Weyl
semimetal (WSM) is an example, which ischaracterized by a Weyl dispersion in bulk andnontrivial surface states with Fermi-arc.
The pyrochlore iridates R2Ir2O7 (R = rare-earth,Y,and Bi ions) is one of representative materials to
explore these phenomena. The metal-insulatortransition (MIT) can be tuned by changing R-ions via
the modulation of 5d electron hopping interaction andaccordingly of effective electron correlation (U). The
R = Pr compound is metallic down to the lowest
temperature, while others with smaller R-ionic radius(r) thermally turns into an insulator concomitantly withor separately from the antiferromagnetic (all-in all-out)order at low temperatures. So far, there have been
many theoretical studies on the evolution ofelectronic/magnetic state upon the MIT. For example,
it is proposed that the Weyl semimetal can be realized
in the vicinity of MIT. Despite extensive research, asystematic study regarding the evolution of electronic
structure upon this enigmatic MIT is still lacking.
In this study, we investigate the charge dynamics
upon the MIT on series of pyrochlore iridates(Nd1-xPrx)2Ir2O7 by means of optical spectroscopy. For
the measurements, high-quality (hard, dense, and wellstoichiometric) polycrystalline samples are prepared.
We measured reflectivity spectra for polycrystals of
Nd2Ir2O7 between 12 meV and 5 eV using Fouriertransform spectrometer and grating-type spectrometerin temperature range between 5 K and 290 K. We
obtained the reflectivity between 5 and 30 eV at room
temperature by using synchrotron radiation at BL7B in
UVSOR. We determined complex optical constants
through the Kramers-Kronig transformation withproper extrapolations of reflectivity in the low andhigh energy regions.
In Fig. 1(a), we show the optical conductivity
spectra for the mixed compounds of R = Nd1-xPrx (x =0, 0.2, 0.5, and 1) at 5 K (< TN), which are located in
vicinity to the quantum phase transition of the all-in
all-out magnetic order [2]. An optical gap is observedfor x = 0.5 as well as for x = 0, while the magnitude of
the gap systematically decreases as x increases (0.035

eV for x = 0.2 and 0.012 eV for x = 0.5). Figure 1 (b)exhibits the T variation of optical conductivity spectra
for x = 0.5. Upon cooling below TN, the spectral weight(SW) rapidly piles up around 0.04 eV at the expense ofa strong depletion of conductivity in lower energy,
resulting in the formation of the charge gap as small as
0.012 eV. It is quite different T evolution of spectrafrom other compounds indicative of the emergence ofunconventional ground states.
We discuss the possible electronic states of x = 0.5,
which is located near the quantum MIT point. At firstglance, the observed T evolution of the spectra isreminiscent of the band reconstruction due to theformation of charge/spin density wave states, in which
the SW transfers to right above the optical gap. This is
however not likely since the band reconstruction by thefolding of Brillouin zone may not occur in the case of
all-in all-out type magnetic order. Another possible
scenario is emergence of Weyl Mott insulator. It istheoretically proposed that the electronic band can be
gapped while maintaining topological aspects byelectron-electron correlation. In this state,characteristic absorption peaks due to interband
interactions are predicted to show up in opticalconductivity, which is not inconsistent with thepresently observed results.

Fig. 1. (a) Optical conductivity spectra for x = 0, 0.2,
0.5 and 1 of R = Nd1-xPrx at 5 K. (b) Opticalconductivity spectra at several temperatures for x = 0.5.
[1] X. Wan et al., Phys. Rev. B 83 (2011) 205101.
[2] K. Ueda et al., Phys. Rev. B 92 (2015) 121110(R).

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521BL8B
宮崎隆文.docx
Ⅲ -2. Materials
Sciences

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Ultraviolet Photoelectron Spectra of Sc3C2@C80
T. Miyazaki1, *, H.Yagi1, T. Zaima1, R. Sumii2,
H. Okimoto3, N. Izumi3, Y. Nakanishi3, H. Shinohara3 and S. Hino1

Graduate School of Science and Engineering, Ehime University, Matsuyama 790-8577, Japan
2
Institutes for Molecular Science, Okazaki 444-858, Japan
3
Graduate School of Science, Nagoya University, Nagoya 464-8602, Japan(*Present address: Faculty of Science, Okayama University, Okayama 700-8530, Japan)

1

the Eonset. The small onset energy of Sc3C2@C80 isqualitatively consistent with the previous DFT
calculation by Sato et al. [1] in which the band gap of
Sc3C2@C80, La2@C80 and Sc3N@C80 is estimated to be
0.07, 0.41 and 1.23 eV, respectively.

Sc3C2@C82

Intensity / a.u.

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Fullerene cages often encapsulate metal atoms. C80
endohedral fullerenes attract attention because of thereasons such as empty Ih-C80 is not stable and cannot
be isolated but it becomes stable upon encapsulation of
metal atoms. Intramolecular charge transfer fromencapsulated species to the carbon cage gives themunique and different characteristics from empty
fullerenes. Sc3C2@C80 is one of such endohedralcluster fullerenes and well studied because its highly
symmetric Ih cage structure is expected to result in
unique properties. Recently, Sato et al. [1] reported
considerably high carrier mobility (the highest among
the polycrystalline endohedral fullerene films) andhigh electrical conductivity of polycrystalline
Sc3C2@C80 film compared to those of La2@C80 and
Sc3N@C80. Their DFT calculation suggests that highconductivity of Sc3C2@C80 is attributed not only to thehigh carrier mobility but also to its small band gap(0.07 eV), and that Sc3C2@C80 can be considered as asemimetal at room temperature. To examine whether
Sc3C2@C80 really has such a small bandgap, we
measured the Ultraviolet Photoelectron Spectra (UPS).
Sc3C2@C80 was synthesized according to theprevious procedures. The UPS were measured by thephotoelectron spectrometer at the beamline 8B2 at the
Ultraviolet Synchrotron Orbital Radiation Facility(UVSOR), and by our laboratory equipment which
consists of He-I, (hv = 21.218 eV) discharge lamp
(combination of MB Scientific L-1 and M-1) and a
SCIENTA SES-100 electron energy analyzer. All the
spectra are referenced to the Fermi level of thesubstrate, which was determined by measuring the
Fermi edge of the gold. The Sc3C2@C80 specimen for
the UPS measurements were deposited by sublimationfrom a resistively heated quartz crucible.
Figure 1 shows the UPS of Sc3C2@C80 obtainedwith hv = 20 ~ 60eV. The spectra exhibit seven peaks
denoted as A-G in the figure. The intensity ratios of
those peaks oscillate with the incident photon energy,
which is due to the spherical shape of the molecule andis the characteristics of the UPS of fullerenes. The
onset energy of the UPS (Eonset), which is the energy
separation between the Fermi level and the valenceband maximum, is 0.2 eV for Sc3C2@C80. This value
is much smaller than the Eonset of La2@C80 (0.77 eV)and Sc3N@C80 (1.1 eV). The Fermi level of the n-typesemiconducting fullerene thin film usually does not
situate at the mid gap but closer to the conduction band
minimum, so the band gap of Sc3C2@C80, La2@C80
and Sc3N@C80 is probably between Eonset and twice of

h=
60 eV
55 eV

GF

50 eV

E

45 eV

D

40 eV
35 eV

C B

30 eV

A

25 eV
21.22 eV
20 eV

10

8

6

4

2

Binding Energy / eV0Fig. 1. Ultraviolet
photoelectron
spectraof
Sc3C2@C80 with incident photon energy at hν = 20 ~60 eV.
[1] S. Sato, S. Seki, G. Luo, M. Suzuki, J. Lu, S.
Nagase and T. Akasaka, J. Am. Chem. Soc. 134 (2012)
11681.

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