Ⅲ-2

Materials SciencesⅢ -2. Materials Sciences
BL5U
BL5U

Detwinned Electronic Structure of Ba0.82K0.18Fe2As2 Observed by ARPES
S. Ideta1,2, M. Nakajima3 and K. Tanaka1,2

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

1(b) and will analyze the temperature dependent
ARPES spectra and compare with the previous
ARPES study [8]. In the electronic structure of the ZX and Z-Y directions, we expect that the occupation of
the dyz and dxz orbitals might be the same above Ts,N.
In summary, we have performed an ARPES studyof detwinned K-Ba122 (x = 0.18) and measured theelectronic structure in the Z-X, Z-Y, and Z-X/Y
directions. We found that the electronic structure in theZ-X and Z-Y directions shows different band
dispersions, reflecting the electronic structure from thesingle domain due to the uniaxial pressure. We will
analyze the detailed temperature dependence of the
ARPES spectrum of dyz and dxz.

Iron-based superconductors have an interesting phase
diagram with the antiferromagnetic (AFM) transition
and the structural phase transition. Recently, nematicity,
defined as broken rotational symmetry [a trigonal(C4)to-orthorhombic (C2) structural transition], has shedlight on the understanding of the mechanism on theiron-based superconductivity [1-4]. In hole-doped
BaFe2As2 (Ba122) system, Ba1-xNaxFe2As2 and Ba1-x
KxFe2As2 (K-Ba122) have shown the magnetic orderwithout C4 symmetry breaking [5, 6] and the moment’s
direction would be swapped from in-plane to out-ofplane [7]. Besides, the superconductivity is suppressed
between the superconductivity and the C4-magnetic
phase. The electronic structure at the C4 magnetic phasehas been unclear yet, and it would give us a greatinterest to elucidate the mechanism of the hole-doped
Ba122.
Here, in order to compare with results obtained fromthe previous study [8], we have demonstrated atemperature dependent angle-resolved photoemission
spectroscopy (ARPES) experiment to elucidate theelectronic structure of detwinned underdoped KBa122 (x = 0.18) with the orthorhombic (C2) phasetransition below the structural and AFM transition
temperatures (TN,S ~ 100 K). In hole-doped Fe-SCs, it
has been known that the anisotropic resistivity
disappears in contrast with the electron-doped Fe-SCs[9]. However, we have clearly observed the orbital
anisotropy corresponding to the two-fold symmetry at
the C2 phase in the hole-doped Fe-SC.
High-quality single crystals of Ba0.82K0.18Fe2As2 (Tc~ 10 K) were grown by self-flux technique. ARPES
experiments were carried out at BL5U of UVSOR-III
Synchrotron using the linearly s polarized light of hν= 60 eV. Temperature was set at T = 6 K-150 K andclean sample surfaces were obtained for the ARPES
measurements by cleaving single crystals in-situ in an
ultrahigh vacuum better than 1×10-8 Pa.
Figures 1(a) and 1(b) show the second-derivative
ARPES intensity plots along the Z-X and Z-Y
directions, respectively, corresponding to the Z points.
In Figs. 1(a) and 1(b), the direction of the uniaxialpressure is perpendicular or parallel to the spolarization vector respectively [4] as shown by
photos, and both band dispersions are considered to be
the electronic structure from the strain free single
domain. In Fig.1(c), we also measured twinned single
crystals and found that two dispersions simultaneously
corresponding to the dxz and dyz orbitals shown in Figs.
1(a) and 1(b), respectively. In future work, we focus
on the dyz and dxz orbitals as shown in Figs. 1(a) and

Fig. 1. Electronic structure of underdoped K-doped
Ba122. (a), (b), (c) Second-derivative ARPES
intensity plots for detwinned (Z-X and Z-Y directions)and twinned Ba0.82K0.18Fe2As2 taken at hν = 60 eV. The
dyz and dxz bands are clearly observed differently.
Photos of the sample setup and schematic Brillouin
Zone are shown on the right side. Red arrows indicate
the direction of uniaxial pressure.
[1] R. M. Fernandes et al., Nat. Phys. 10 (2014) 97.
[2] R. M. Fernandes et al., Phys. Rev. Lett. 111 (2013)
127001.
[3] T. Shimojima et al., Phys. Rev. B 90 (2014) 12111(R).
[4] Y. Ming et al., PNAS 108 (2011) 6878.
[5] L. Wang et al., Phys. Rev. B 93 (2016) 014514.
[6] A. E. Böhmer et al., Nat. Commun. 6 (2015) 7911.
[7] D. D. Khalyavin et al., Phys. Rev. B 90 (2014)
174511.
[8] S. Ideta et al., Activity Report 45 (2017) 77.
[9] E.C. Blomberg et al., Nat. Commun. 4 (2014) 1914.

43

Ⅲ02- 1 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL1U
BL1U

Photon Induced Positron Annihilation Lifetime Spectra of Ce:GAGG Crystals
Measured Using Ultrashort Laser-Compton-Scattered Gamma Ray Pulses
K. Fujimori1, M. Kitaura1, Y. Taira2, M. Fujimoto3, Y. Okano4, H. Zen5, M. Katoh3,8,
M. Hosaka6, J. Yamazaki3, K. Kamada7 and A. Ohnishi1

1
Faculty of Science, Yamagata University, Yamagata 990-8560, Japan
National Institute of Advanced Industrial Science and Technology, Tsukuba 305-8568, Japan
3
UVSOR Synchrotron Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
4
Center for Mesoscopic Sciences, Institute for Molecular Science, Okazaki 444-8585, Japan
5
Institute of Advanced Energy, Kyoto University, Uji 611-0011, Japan
6
Synchrotron radiation Research Center, Nagoya University, Nagoya 464-8603, Japan
7
New Industry Creation Hatchery Center, Tohoku University, Sendai 980-8579, Japan
8
Hiroshima Synchrotron Radiation Center, Hiroshima University, Higashi-hiroshima, 739-0046, Japan2

Ⅲ-2

Ce3+-doped Gd3Al2Ga3O12 (Ce:GAGG) is known as
an inorganic scintillator with high light yield and goodenergy resolution. It has been demonstrated that, in thismaterial, antisite Gd2+ ions adjacent to oxygen
vacancies are responsible for shallow electron traps [1].
Such electron traps cause the occurrence ofphosphorescence with long lifetime and the lowering oflight yield, and thus they have to be suppressed in theprocess of crystal growth. It was pointed out that theantisite Gd3+ ions and oxygen vacancies are introducedin Al/Ga defficient compositions. In order to improve
scintillation properties of Ce:GAGG crystals, it isimportant to find out the best way to suppress Al/Ga
deficiency.
The existence of cation vacancies has been
investigated by positron annihilation lifetime
spectroscopy, because positrons are attractive fornegative charged cation vacancies. The position
annihilation lifetime spectroscopy (PALS) generally
requires radioisotope such as 22Na as a positron source.
In this case, positrons are injected from the outside of acrystal, so that pair annihilation gamma rays generate
not only from the inside of the crystal but also from theoutside. The latter has to be removed to analyze thepositron annihilation lifetime at the cation vacancy site,
because it is responsible for the uncertainty of the
lifetime. This problem is to be solved to perform thelifetime analysis more accurately.
Photon induced PALS (PIPALS) is the only solutionthat can solve the above-mentioned problem [2]. Thismethod utilizes ultrafast laser Compton-scattered (LCS)
gamma ray pulses generated by the normal collision ofelectron beam and ultrashort laser pulses. The advantage
is that positrons can be created inside the crystal throughthe pair creation process. PIPALS allows us to clarify theexistence of cation vacancies in Ce:GAGG crystals. For
this purpose, we performed PIPALS experiment using
ultrashort LCS gamma ray pulses at the BL1U beamlineof UVSOR synchrotron facility.
The size of the Ce:GAGG crystal used in our
experiment was 10×10×10 mm3. The details of PIPALS
experiment have been reported elsewhere. The photon
energy, pulse width, and repetition rate of the LCS
gamma ray pulses were 6.6 MeV, 2 ps, and 1 kHz,

respectively. The pair annihilation gamma ray from theCe:GAGG crystal was detected with two
photomultiplier tubes attached BaF2 crystals, whichwere faced each other. The waveforms were stored in
the memory of a digital storage oscilloscope, when the
output signals from two photomultiplier tubes coincidedwith each other. The PIPALS spectrum of a Ce:GAGG
crystal was obtained from the distribution of the time
half the maximum intensity.
The PIPALS spectrum of a Ce:GAGG crystal isshown in Fig. 1. Solid circles indicate experimental data.
Apparently, the PIPALS spectrum is of a single decay
component. A solid line indicates the result of curve
fitting analysis, and it reproduces the experimental data.
From the fit, the lifetime is determined 178 ± 14 ps. Thisvalue is almost agreement with the lifetime of the pairannihilation of a positron with an electron in the bulk. A
longer lifetime component is slightly observed in the
tailof the main component. The higher statisticalaccuracy for experimental data is needed for detail
lifetime analysis.

Fig. 1. PIPALS spectrum of a Ce:GAGG crystal
measured at room temperature.
[1] M. Kitaura et al., Appl. Phys. Lett. 113 (2018)
041906.
[2] Y. Taira et al., Rev. Sci. Instrum. 84 (2013) 053305.

44

Ⅲ02- 2 -Ⅲ -2. Materials Sciences
BL1B
BL1B

Polarized Terahertz Reflectivity Spectra of Ca3TaGa3Si2O14 Piezoelectric
Crystals with Ordered Langasite Structure
M. Kitaura1, K, Kamada2, Y. Yokota2, Y. Ohashi2, S. Watanabe3 and A. Ohnishi1
Faculty of Science, Yamagata University, Yamagata 990-8560, JapanNew Industry Creation Hatchery Center, Tohoku University, Sendai 444-8585, Japan
3
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan1

2

formed in CTGS. The phonon polariton dispersion
curves obtained are drawn in Fig. 2.

Langasites are expressed by the chemical formulaof A3BC3D2O14, and are composed of four cations AD. These materials belong to the trigonal system of thespace group P321, and are known piezoelectricmaterials with high electromechanical coupling
factors. They are classified into two categories: ordertype and disorder-type. In order-type langasites, eachcation site is occupied by different element. In
disorder-type, one element occupies multi cation sites.
Order-type langasites are preferred, because theypresent better thermal stability of piezoelectric
properties. Ca3TaGa3Si2O14 (CTGS) is one of suchorder-type langasites. CTGS exhibits a higher
electrical resistivity which is advantage to hightemperature use of sensor applications [1]. Despite
such an industrial key material, the fundamental
optical functions such as dielectric constant andrefractive index in the terahertz region have not yet
been studied so far. In order to determine terahertz
optical functions of CTGS, we have measured
polarized terahertz reflectivity spectra of CTGS
crystals at 9 K.
A typical result is shown in Fig. 1 by using dottedcurves. One can see remarkable dichroism between the
spectra for Ec and E//c configurations, which reflects
the crystal structure of CTGS. According to theconsideration based on group theory [2], the E and A2
modes are IR-active for Ec and E//c configurations,
respectively. Thus, it is likely that the two bands are
assigned to E and A2 modes for Ec and E//c
configurations, respectively.
We determined the complex dielectric function fromthe fit of Eq. (1) with experiment data.

,

Ⅲ-2

Fig. 1. Polarized terahertz reflectivity spectra of
CTGS for Ec and E//c configurations, measured at
9K. Dotted and broken curves indicate experimental
and calculated results, respectively.

(1)Fig. 2. Phonon polariton dispersion curves of CTGS
for Ec and E//c configurations, calculated on thebasis of Fig. 1.

where ε∞ is the high-frequency dielectric constant,  isthe frequency of terahertz wave, and TO(LO) andTO(LO) are frequencies and damping factors of TO(LO) phonons, respectively. The parameters
determined from the fit of Eq. (1) are listed in Table 1.
The reflectivity spectra calculated using the complex
dielectric function are indicated by broken curves. In
Fig. 1, the calculated reflectivity spectra reproduce theexperimental ones. The lowest frequency bands
exhibit the high-reflectivity plateau in the highfrequency side of TO and the dip at the position of LO.
These features imply that phonon polaritons are

Table1: Parameters obtained from the fit of Eq. (1)with experimental data.

Ec
E//c

∞

TO
(THz)

TO
(MHz)

LO
(THz)

LO
(MHz)

16.3
9.6

1.58
2.57

15.5
68.6

2.32
3.34

53.9
26.9

[1] X. Fu et al., J. Ceram. Soc. Jpn. 124 (2016) 523.
[2] J. Lan et al., Phys. Stat. Sol. 242 (2005) 1996.

45

Ⅲ02- 3 -Ⅲ -2. Materials Sciences
BL1B, BL7B
BL1B, BL7B

Extremely Sharp Structure of the Vibration Modes Derived from Hydrogen
Bonds of Single-crystalline L-alanine
2Ⅲ-2

Z. Mita1, H. Watanabe1,2 and S. Kimura1,2

Graduate School of Frontier Biosciences, Osaka University, Suita 565-0871, Japan
Department of Physics, Graduate School of Science, Osaka University, Toyonaka 560-0043, Japan
1

Amino acids are the most basic molecules, which
compose proteins in vivo. Since almost all amino acids
have chiral molecule structure, the crystal structures aswell as the optical properties are expected to beanisotropic. One of amino acids with optical isomers,
alanine [CH3CH(COOH)NH2] has orthorhombic
molecular structure [1]. Many spectroscopic
measurements have been performed for alanine so far[2,3,4]. To clarify the thermal effect of molecularvibrations of different origins, we have measured
polarization and temperature dependences of opticalconductivity spectra using single crystalline alanine.
We have grown the single crystalline L-alanine byusing the solvent evaporation method and confirmed
to be orthorhombic structure by taking X-ray Lauediffraction pattern. Anisotropic optical reflectivity
spectra from the terahertz (THz) to vacuum-ultraviolet(VUV) regions have been measured by using BL1B
and BL7B of UVSOR-III and laboratorial equipments.
The obtained spectra were converted to opticalconductivity [σ(ω)] spectra by the Kramers-Kronig
analysis. Temperature dependence of anisotropic σ(ω)
spectrum in the THz region is shown in Fig. 1. To
reveal its temperature dependence, each peak was
fitted with the function of the Lorentz model [5]. As
shown in Fig. 1, there are 18 peaks originating frommolecular vibrations. The spectra are strong
anisotropy and all peaks have very large temperaturedependence. Specifically, the peak positions shift tothe high energy side and the widths become narrowsignificantly with decreasing temperature. However,the peak areas of all peaks were approximately
constant. In comparison of the obtained peak energy
with a previous investigation [2], peaks of number 1 to
12 originate from the vibration of hydrogen bond.
Figure 2 shows the peak widths of all peaks as afunction of peak energy at 10 K. Red (blue) bars
indicate the vibration with (without) the contributionof hydrogen bond. The peaks originating fromhydrogen bond have extremely sharp structure, sohydrogen bond is a vibration where thermal
fluctuation effects mainly appear. Therefore, tocompare the peak width with one another at a very low
temperature may become one of useful methods toidentify the origin of the peaks, hydrogen bonds or not.
In addition, at 10 K, the peak widths originating fromhydrogen bonds become almost constant near the
resolution limit of the spectrometer, suggesting thepossibility of the quantum fluctuations.

(a) E⊥c⑬⑰②-⑤① ⑥⑦⑧ ⑩(b)E // c

10K

⑫

⑮

⑨

⑱
⑪
⑭⑯

300K

Fig. 1. Anisotropic optical conductivity [σ(ω)]
spectra at temperatures from 10 to 300 K in the THz
region perpendicular to the c-axis (a) and parallel tothe c-axis (b).
⑯

10 K
⑭

⑰
②-④

⑤-⑧

①

⑬
⑩
⑨ ⑪⑫

⑮

⑱

Fig. 2. Peak widths of all peaks in Fig. 1 at 10 K in
THz region. Red (blue) bars show the vibrations with
(without) the contribution of hydrogen bonds.
[1] M. Fleck and A. M. Petrosyan, Salts of Amino AcidsCrystallization Structure and Properties (Springer,2014) p.26.
[2] J. Bandekar et al., Spectrochimica Acta Part A:
Molecular Spectroscopy 39 (1983) 357.
[3] M. Yamaguchi et al., Appl. Phys. Lett. 86 (2005)
053903.
[4] V. S. Minkov et al., J. Struct. Chem. 51 (2010) 1052.
[5] F. Wooten, Optical Properties of Solids 82 (Academic
Press, 1972) 42.

46

Ⅲ02- 4 -Ⅲ -2. Materials Sciences
BL1B
BL1B

Ground State of Ultrashallow Thermal Donors in Silicon
A. Hara and T. Awano

Faculity of Engineering, Tohoku Gakuin University, Tagajo 985-8537, Japan
We discovered ultrashallow thermal donors(USTDs) in carbon- and hydrogen-doped Czochralski
silicon (CZ Si) crystals [1, 2]. To the best of ourknowledge, these are the shallowest energy levels
among those of currently reported donors in Si crystals.
In addition, the ground (1S) state of USTDs was foundto be split into two states [3, 4]. This is direct evidence
that the ground state of USTDs consists of a linear
combination of the wave functions of the conduction
band minimum. However, electron spin resonancemeasurements have shown nearly isotropic spectrawith slight anisotropy, thus indicating Td symmetry. If
Td symmetry is correct, the ground state must split into
three levels [5, 6]. To evaluate another ground state,
we measured the optical absorption spectra at a regionof low wavenumbers using BL1B.
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 rapid exposure to air. For carbon
doping, the Si melt used to prepare the ingot wasdoped with carbon powder during CZ Si crystal
growth. Transmittance spectra were obtained using a
BL1B beamline with a flowing cryostat.
Figure 1 shows Trans. (40 K)/Trans. (8 K), where
Trans. (T K) represents the transmittance at T K [3, 4].
According to our previous research using BL6B [3, 4],
the two broad dips observed at approximately 100 and
150 cm-1 at 40 K originate in the transition from theupper ground state to the 2P0 and 2P+− states. The
upper ground states are occupied by thermal excitationof electrons from the ground state.
Figure 2 shows Trans. (40 K)/Trans. (5 K) at aregion of 20–100 cm-1. No clear difference wasobserved between the two transmittance spectra.
These indicate that the ground state of USTDs weresplit into two levels but not three. Thus, symmetry of
USTDs was concluded not to be Td.
In summary, the ground-state splitting of USTDs
was evaluated at a low-wavenumber region. However,new peaks were not observed at 20–100 cm-1. This
indicates possibility that the ground state of USTDs are
split into two and that symmetry of USTDs is not Td.

2.0

Ground state – 2P+- USTD-3

Trans.(40 K) / Trans. (8 K)

1.8

USTD-1

1.6

Ground state – 2P0
USTD-1 USTD-3

1.4
1.2

a

b

a* b*

1.0
c*d*

0.8
0.6

c d

0.4
0.2
0.0

50

100

150
200
Wave number (cm-1)

250

300

Fig. 1. Trans. (40 K)/Trans. (8 K) of USTDs [4].
2.0Trans.(40 K) / Trans. (5 K)

1.8
1.6
1.4
1.2
1.0
0.8
0.6
0.4
0.2
0.0

20

30

40

50
60
70
Wavenumber (cm-1)

80

90

100

Fig. 2. Trans. (40 K)/Trans. (5 K) of USTDs.
[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] A. Hara and T. Awano, Jpn. J. Appl. Phys. 56(2017) 068001.
[4] A. Hara and T. Awano, Jpn. J. Appl. Phys. 57(2018) 101301.
[5] W. Kohn, in Solid State Physics, ed. F. Seitz and D.
Turnbull (Academic Press, New York, 1957) Vol. 5, p.
257.
[6] R. L. Aggarwal and A. K. Ramdas, Phys. Rev. 140(1965) A1246.

47

Ⅲ02- 5 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL2A
BL2A

XAFS Measurement for Alloying Element in Tempered Nitrogen Martensite
M. Sato and T. Murata

Institute for Materials Reseach, Tohoku University, Sendai 980-8522, Japan

Ⅲ-2

It is well known that the martensite phase obtainedby rapid quenching is hard but brittle. Therefore thebalance between hardness and toughness is adjusted by
the subsequent tempering treatment. At that time,
cementite is precipitated in case carbon steel, andFe16N2 or Fe4N is precipitated in case of nitrogen steel[1, 2].
From the experimental results, the Fe16N2 precipitatesin the temperature range between 100 °C and 250 °C,
and transition to Fe4N occurs at tempering temperaturearound 300 °C. In addition, it has been found thatprecipitation of alloy nitride also occurs because thesubstitutional element can be diffused [3].
Although the alloying elements are known to partiallydissolve as solid solutions in these precipitates, there isno data of Si and Mo in iron nitrides. The solid solutionof alloying elements into iron nitrides greatly affect thestability of iron nitrides and also affect the subsequent
precipitation of alloy nitride. Therefore, in order tounderstand the hardness change associated withtempering, it is important to clarify the solid solution
state of such substitutional elements.
In this study, changes in the chemical state of thesubstitutional elements such as Si and Mo in the
precipitates formed during tempering were investigated
by XAFS measurement.
Fe-1mass%Si and Fe-1mass%Mo alloys were usedas starting materials. They were homogenized at 1523
K for 24h and furnace-cooled until room temperature.
The Fe-1Si-0.3N and Fe-1Mo-0.3N alloys wereprepared by nitriding and quenching (N-Q) processusing NH3/H2 mixed gas at 1273 K for 1 h. Thenobtained these alloys were tempered at 773 K, andprecipitated Fe4N generated during tempering wereextracted using Iodine-alcohol procedure. The Si Kedge and Mo L3-edge XANES spectra were correctedby fluorescence method using InSb double crystal
monochrometer and silicon drift detector (SDD) at
BL2A in UVSOR, respectively. Obtained data wereanalyzed using Athena software.
Figure 1 shows Si K-edge XANES spectra ofextracted Fe4N powders. As shown in this figure, pure
Si, SiO2 and Si3N4 were used as references.
The obtained XANES spectra are different for eachsamples, and the extracted residue from the tempered
sample at 200 °C had a peak around 1847 eV, while thethat obtained from the sample tempered at 500 °C
showed double peaks around 1844 eV and 1847 eV,
respectively. From the comparison with the reference
sample, the peak at 1847 eV is presumed to be due tothe oxidation that occurred during extraction or storage.
On the other hand, it can be seen that the peak around1844 eV shows similar shape to Si3N4. From the TEM
observation, since the precipitation of the Si-nitrides
has not been observed in this sample, it is consideredthat the Si makes a solid solution into Fe4N and the
local structure of Si in Fe4N has similar to that of Si3N4.
Similar results were obtained in the case of Fe-1Mo0.3N alloy. In the future, detailed examination of solid
solution state of alloying element in iron nitrides
should be carried out by combination of N K-edge andO K-edge measurement of these samples at BL4B andprediction of XANES spectra from calculation.

Fig. 1. Si K-edge XANES spectra of references andextracted precipitates. (a) Pure Si, (b) SiO2, (c) Si3N4,(d) extracted residue from tempered sample at 200 °C
for 1 h, and (e) extracted residue from tempered
sample at 500 °C for 1h.
[1] L. Cheng and E. J. Mittemeijer, Met. Trans. A 21A(1990) 13.
[2] L. Cheng, A. Bottger and E. J. Mittemeijer, Met.
Trans. A 23A (1992) 1129.

48

Ⅲ02- 6 -Ⅲ -2. Materials Sciences
BL2A
BL2A

Mo L3-edge XANES Study of Active and Durable Mo-Carbide Species on
H-MFI for Methane Dehydroaromatization Catalysts
H. Aritani1, K. Kuramochi2, R. Yamazaki2, H. Miyanaga1, S. Sato1 and A. Nakahira3
Graduate 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 Japan1

2

the catalyst showed high MTB reactivity, it issuggested that the deeply carbonized species relate to
the highly active one. For Mo/H-MFI (Si/Al2=30), theMo species are different from those of Mo/H-MFI in
Si/Al2=40 even in H2 co-feed and/or V comodification. Deeply reduced Mo species on H-MFI(Si/Al2=30) support have been invisible, and α-Mo2C
like species were formed in any cases. Because the
Mo/H-MFI (Si/Al2=30) catalysts showed durable
reactivity for MTB, the correlation between thereactivity and the active Mo-carbide species is still
unclear. It is now in progress.

As an innovative GTL (Gas-To-Liquid) process,methane dehydroaromatization (methane-to-benzene
reaction; MTB) has been studied for its potential in the
direct conversion of methane into liquid fuels. It has
been investigated intensively that Mo-modified HMFI (Mo/H-MFI) zeolite catalysts show high activity
for methane dehydroaromatization. Since the MTB
reaction is revolutionary for direct conversion ofnatural gas, deactivation due to coking over thecatalyst cannot be avoided during the methane
conversion. It is a serious problem for development of
MTB reaction, and thus, clarification of thedeactivation process over the catalysts is one of theimportant point for develop of highly active anddurable MTB catalysts. Not only carbon deposition onacid sites of H-MFI but also excess carbonization ofactive Mo species causes the deactivation of MTB. It
is widely accepted that addition of hydrogen tomethane reactant is effective for inhibiting thedeactivation. In the present study, Mo L3-edge
XANES study is introduced to characterize the active
Mo-carbide species on H-MFI after the MTB reaction.
Mo/H-MFI and Mo-V/H-MFI catalysts wereprepared by impregnation of H-MFI (Si/Al2=30-50)
support with MoO2(acac)2−CHCl3 or MoO2(acac)2−
VO(acac)2−CHCl3 solution (5.0wt% as MoO3;
Mo/V=10), and followed by drying overnight andcalcination at 773 K. H-MFI supports weresynthesized hydrothermally at 413 K for a week, andfollowed by ion-exchanging with NH4Cl andcalcination at 873 K. The catalytic activity of MTB
was evaluated by means of fixed bed flow reaction, as
described in a separate paper.[1] The reactant gas isCH4(20%)+H2 (0-3%)+He(base) at the flow rate of 30
mL/min, and the reactivity was evaluated at 1023 K byusing the 0.25 g of each Mo/H-MFI catalyst. Mo LIIIedge XANES spectra were obtained in BL2A of
UVSOR-IMS in a total-electron yield mode using
InSb double-crystal monochromator. Photon energy
was calibrated by using Mo metal-foil at Mo LIII-edge,
and normalized XANES spectra are presented byusing REX-2000 (Rigaku) software.
Figure 1 shows the Mo L3-edge XANES spectra of
Mo/H-MFI and Mo-V/H-MFI catalysts (Si/Al2=30
and 40) after MTB reaction at 1023 K. For Mo/H-MFI
in Si/Al2=40, partially carbonized Mo species(MoCxOy) were possibly formed. By H2 co-feed,
reduction of Mo species advanced. In case of Mo-V/HMFI (Si/Al2=40), the Mo species were deeply
carbonized and formed MoCx (x>0.5) species . Since

H-MFI (Si/Al2=30)

TEY

1.0

Mo/H-MFI (CH4)
Mo/H-MFI (CH4+H2)
Mo-V/H-MFI (CH4)
Mo-V/H-MFI (CH4+H2)
Mo-Metal (ref.)-Mo2C (ref.)

0.5

0.0
2520

2

1.0

TEY

2530
2540
Energy / eV H-MFI (Si/Al =40)

Mo-V/H-MFI (CH4)
Mo/H-MFI (CH4+H2)
Mo-V/H-MFI (CH4)
Mo-V/H-MFI (CH4+H2)
Mo-Metal (ref.)-Mo2C (ref.)

0.5

0.0
2520

2530
Energy / eV

2540

Fig. 1. Mo L3-edge XANES of Mo/H-MFI and MoV/H-MFI catalysts (Si/Al2=30 and 40) after MTB
reaction at 1023 K.
[1] H. Aritani et al., J. Environm. Sci. 21 (2009) 736.
[2] H. Aritani et al., UVSOR Activity Report 2016 44(2017) 52.
[3] H. Aritani et al., UVSOR Activity Report 2017 45(2018) 44.

49

Ⅲ02- 7 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL2A
BL2A

Local Environment Analysis of Sr Ions in Octacalcium Phosphate
K. Arakawa 1 and T. Yamamoto1,2,3

Faculty of Science and Engineering, Waseda University, Tokyo 169-8555, Japan
Insitute of Condensed-Matter Science, Waseda University, Tokyo 169-8555, Japan
3
Kagami Memorial Research Institute for Materials Science and Technology, Waseda University, Tokyo 169-0051,
Japan
1

2

Considering the above experimental and theoretical
XANES and theoretical solution energy the local
environment of Sr ions in OCP could be determined.

Intensity (arb. units)

Ⅲ-2

Calcium phosphate basis bioceramics have been
widely studied because of their efficient medical
applications. Among these bioceramics, octacalcium
phosphate (OCP) has a great attention, since it is aprecursor of the hydroxyapatite (HAp), which canchange into the natural bone in human body. It wasreported that the existence of the trace elements such
as Zn and Mg can enhance the speed of bonereconstruction [1]. However, the mechanism of suchenhancement has not yet been understood. To
understand the mechanism, it is essential to know the
local environment of doped ions. For such purpose, wehave studied wide variety of the functional materials
by using the X-ray absorption near edge structure(XANES) and the first-principles calculations within adensity functional theory (DFT), which could
successfully determine the local environment of the
trace elements in β-tricalcium phosphate [2-6]. In thecurrent study, local environment of Sr ions in OCP has
been investigated by the above strategy, i.e., combined
use of XANES measurements and DFT calculations.
Sr-doped OCPs were synthesized with precipitation
method changing the concentration of doped Sr ions.
Crystal structures of these Sr-doped OCPs were
examined by the powder X-ray diffraction, whichshowed all the samples are single phased and nosecondary phase could be observed. Elemental
analysis was also carried out with Energy Dispersive
X-ray (EDX) analyzer equipped with the scanningelectron microscope (SEM) to confirm theconcentration of Sr ions in OCP. For the local
environment analysis of Sr ions in OCP, Sr-L3 XANES
spectra were observed at BL2A, UVSOR, in a total
electron mode. Sample powders were settled on the
carbon adhesive tape, which was put on the first dinodeof the electron multiplier. The synchrotron radiations
from the storage ring was monochromatized with the
InSb(111) double-crystal monochrometer. Observed
Sr-L3 XANES spectra are shown in Fig. 1. Although it
can be determined that Sr is divalent from thiscomparison, we cannot get any further information on
local environment of Sr ions in OCP only from theseexperiments. To get more information, theoretical
XANES spectra for the models of Sr-doped OCP wereconstructed by the all-electron linearized augmented
wave plus local orbital package, WIEN2k [7], withcore-hole effect. Furthermore solution energy for thesemodels were also calculated with same package.

1940

SrCO3

OCP:Sr

1950

1960

Energy (eV)

1970

1980

Fig. 1. Observed Sr-L3 XANES spectra of SrCO3 andSr-doped OCP.
[1] H. Kawamura, A. Ito, S. Miyakawa, P. Layrolle, L.
Ojima, H. Naito, N. Ichinose and T. Tateishi, J. Biomed.
Mater. Res. 50 (2000) 184.
[2] K. Kawabata, H. Sato and T. Yamamoto, J. Ceram.
Soc. Jpn. 116 (2008) 108.
[3] K. Kawabata, T. Yamamoto and A. Kitada, Physica
B406 (2011) 890.
[4] K. Kawabata, T. Yamamoto and A. Kitada, Mater.
Trans. 56 (2015) 1457.
[5] K. Kawabata, H. Sato and T. Yamamoto, UVSOR
Activity Report 2007 35 (2008) 108.
[6] K. Kawabata and T. Yamamoto, UVSOR Activity
Report 2010 38 (2011) 110.
[7]P. Blaha, K. Schwarz, G.K.H. Madse, D. Lvasnicka,
J. Luiz, WIEN2k (Wien, Austria, Univ. Vienna, 2001).

50

Ⅲ02- 8 -Ⅲ -2. Materials Sciences
BL2A
BL2A

P-K XANES Study for Hydration Reaction of α-Tricalcium Phosphate in
Aqueous Solutions with Various Ions
H. Murata, R. Hashii and A. Nakahira

Department of Materials Science, Osaka Prefecture University, Sakai 599-8531, Japansoaked in the solution for 12 h remained while samples
were transformed to HAp in the aqueous solution
without Zn2+. Powder X-ray diffraction analyses gavethe same tendencies to that of P-K XANES. These
results means that hydration of α-TCP to HAp does notoccur even on the surface. Therefore, It is implied thathydration reaction of α-TCP to HAp do not simply
occur on the surface but relate solution-reprecipitation
process.

A number of biomaterials which are used in humanbodies have been developed in order to improve
quality of life (QOL) for a long time. Calcium
phosphates are promising biomaterials which is usedas alternative materials of human bones and teeth. αtricalcium phosphate (Ca3(PO4)2, α-TCP) is one ofmain components of bone cements which are hydrated
and hardened in human bodies.[1] Hydration of α-TCP
usually forms hydroxyapatite (Ca5(PO4)3OH, HAp) of
which human bones and teeth are composed. It is
known that coexistent ions and molecules in aqueous
solutions affects hydration reaction of α-TCP. Since it
is considered that hydration reaction starts on surfaceof α-TCP, it is essential to investigate surface and bulk
states of α-TCP hydrated with various kinds ofenvironments.
X-ray absorption near edge structure (XANES) is a
powerful tool to investigate local structures. In soft Xray region, XANES analyses can reveal both of
surface and bulk states using methods with different
attenuation length, electron yield and fluorescent yield.
[2] In detail, electron yield method collects
information around several nano-meters below thesurface while fluorescent yield done around several
micro-meters below it. In this study, we investigatedsurface and bulk structures of hydrated α-TCP in
aqueous solutions which contains various kinds of ionswith P-K XANES.
α-TCP powders were purchased from Taihei
Chemical Industrial Co., Ltd. They were soaked inaqueous solutions with various kinds of ions.
P-K
XANES spectra were collected at BL2A in UVSOR by
a total electron yield (TEY) mode with a drain current
method and a partial fluorescent yield (PFY) modewith a silicon drift detector (SDD). Samples were
mounted using carbon tapes on Cu plates. X-ray beam
was monochromated using InSb double crystals.
Figure 1 shows typical P-K XANES spectra of αTCP samples hydrated in 0.1 mol/L NH3 aqueous
solution at different temperature for 24 h and referencematerials. P-K XANES of α-TCP has different featuresfrom those of HAp and we can distinguish from eachother. P-K XANES of samples shows that α-TCP
hydrated at 70 °C and 80 °C were transformed to HAp
while α-TCP hydrated at 37 °C still remained.
Co-existent ions, such as Zn2+, disturb hydration
reaction of α-TCP. We investigated α-TCP samples
hydrated in aqueous solutions which contains Zn2+.
Both of TEY and PFY results indicated that samples

Fig. 1. Typical P-K XANES spectra of α-TCP
hydrated in different temperature for 24 h andreferences materials.
[1] R.G. Carrodeguas and S. De Aza, Acta Biomater. 7(2011) 3536.
[2] M. Kasrai et al., Appl. Surf. Sci. 99 (1996) 303.

51

Ⅲ02- 9 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL2A
BL2A

Ga L3-edge XAFS Analysis of Ga2O3 Loaded Al2O3 Photocatalysts
T. Yoshida1, R. Ito2, M. Akatsuka2, Y. Kato2, A. Ozawa2, M. Yamamoto1 and T. Tanabe1
1

Ⅲ-2

Advanced Research Institute for Natural Science, Osaka City University, Osaka 558-8585, Japan
2
Graduate School of Engineering, Osaka City University, Osaka 558-8585, Japan

Recently Ga2O3 has attracted a lot of interests as aphotocatalyst for water splitting and CO2 reduction withwater, and various efforts have been paid to improve its
photocatalytic activity. Yamamoto et al. have reported
improvement of photocatalytic activity for CO2
reduction with water under UV irradiation using Ag asa co-catalyst. In the present study, we have focused togeometrical or morphological effects of Ga2O3 particles
supported on Al2O3 (referred as Ga2O3/Al2O3, hereafter)
for water splitting and CO2 reduction under UV lightirradiation without using the Ag cocatalyst. The reasonsof utilization of Al2O3 as the support are twofold, (1) toincrease surface area of Ga2O3, as evidenced by theobservation that Ga2O3/Al2O3 was used for the removalof NOx, and (2) that Al2O3 hardly show thephotocatalytic activity for both water splitting and CO2
reduction.
Ga2O3/Al2O3 photocatalyst samples were preparedby an impregnation method. Ga(NO3)3･8H2O (Kishida
Chemical Co. Ltd. purity 99.0%) and 1.0 g γ-Al2O3(Sumitomo Chemical Co. Ltd. purity 99.99%) wereadded to 200 mL of distilled water and stirred with amagnetic stirrer in air and dried up, followed bycalcination at 823 K for 4h, resulting the Ga2O3/Al2O3
samples. The loaded amounts of Ga2O3 were 5, 10, 20,
40 and 60 wt%. Pure Al2O3 and unsupported Ga2O3
samples (referred as 0 and 100 wt%, respectively) werealso prepared in the similar procedure. Ga L3-edge
XANES were measured with the beam line of BL5S1
at Aichi Synchrotron Radiation Center and BL2A at
UVSOR, Institute for Molecular Science in Japan,
respectively.
Photocatalytic CO2 reduction with H2O under UVlight irradiation was tested for 0.1 g of one of thesamples set in a fixed-bed flow reactor cell under CO2
gas flow. The UV light intensity was 35 mW/cm2 in therange of 254 ± 10 nm. Before the test, the sample wasirradiated with a 300 W Xe lamp for 1h under CO2 gas
flow with the flow rate of 20.0 mL/min. Then the
reduction test was started introducing a NaHCO3
aqueous solution (1.0M) of 10.0 mL and CO2 gas witha flow rate at 3.0 mL/min under the UV lightirradiation. The reaction products (CO, H2 and O2)
were analyzed by a gas chromatograph equipped witha thermal conductivity detector (GC-TCD).
All prepared Ga2O3/Al2O3 samples have shown
significantly higher photocatalytic activity comparedto that of pure Ga2O3 (non-supported Ga2O3). The
cause of the improvement is discussed considering
detailed characterizations by Ga L3-edge XAFS

measurements.
Figure 1 shows Ga L3-edge XANES spectra. A peak
appeared near 1123 eV is caused by a Ga-O-Al bond.
[1-3] For lower Ga2O3 loaded samples, this peak was
relatively high compared to other peaks. This indicatesthat some interaction between Ga2O3 and Al2O3
occurred in lower loaded samples. For higher Ga2O3
loaded samples, the peak intensity at 1123 eV becameless and the spectrum transformed to that of Ga2O3
indicating loaded Ga2O3 well crystallized. The result of
XANES analysis was quite consistent with that of ourseparate EXAFS and XPS analyses.

Fig. 1. Ga L3-edge XANES spectra of all preparedsamples. The spectrum of β-Ga2O3 is also given forcomparison. (a) 0.15 vol% Ga2O3/Al2O3 (b) 0.32 vol%Ga2O3/Al2O3 (c) 0.73 vol% Ga2O3/Al2O3 (d) 2.6 vol%Ga2O3/Al2O3 (e) 4.2 vol% Ga2O3/Al2O3 (f) 100 vol%Ga2O3/Al2O3.Peaks appeared at 1123 eV is assigned tobe caused by Ga-O-Al bond formation.
[1] K. Shimizu, M. Takamatsu, K. Nishi, H. Yoshida,
A. Satsuma, T. Tanaka, S. Yoshida and T. Hattori, J.
Phys. Chem. B 103 (1999) 1542.
[2] X. T. Zhou, F. Heigl, J. Y. P. Ko, M. W. Murphy,
J. G. Zhou, T. Regier, R. I. R. Blyth and T. K. Sham,
Phys. Rev. B 75 (2007) 125303.
[3] N. H. Tran, R. N. Lamb, L. J. Lai and Y. W. Yang,
J. Phys. Chem. B 109 (2005) 18348.

52

Ⅲ02- 10 -Ⅲ -2. Materials Sciences
BL2A
BL2A

Removal of Elastic Scattering for Partial Fluorescence Yield XANES
Measurements
H. Murata

Department of Materials Science, Osaka Prefecture University, Sakai 599-8531, Japan
X-ray absorption near edge structure (XANES) is a
powerful tool for materials science since its targets arewidely varied. For example, ultra-dilute dopants (~ppm), thin films, amorphous materials can be analyzed.
In soft X-ray region, it is difficult to use transmission
mode due to low transparency. Instead, electron yieldand fluorescent yield method were often adapted for
XANES measurements. Although there are someproblems such as “self-absorption effect”, a number ofbeamlines for soft X-ray absorption introduced
fluorescence yield method. In the case of UVSOR, asilicon drift detector (SDD) for partial fluorescentyield (PFY) method are available at BL2A and BL4B.
[1, 2] Using SDD, elastic scattering of incident X-ray
was sometimes observed. Since energy of incident Xray scans from low to high, usual “region of interest”(ROI) sampling cannot remove elastic scattering. In
this study, energy variable fitting method wasdeveloped to removal elastic scattering for PFY
XANES measurements.
PFY XANES measurements were performed at
BL2A in UVSOR. X-ray beam was monochromatedby beryl, KTP and InSb double crystals. Conversion of
SDD channel to X-ray energy were estimated using
linear function. Its factors were determined using
several energy of incident X-ray. It was noted thatrather strong elastic scattering was observed in several
cases. Amorphous materials and samples containing
nano-particles gave strong elastic scattering due totheir broad X-ray diffraction pattern, halo pattern. In
addition to this, ultra-dilute dopants also done since
SDD are moved close to samples to improve detection
efficiency of fluorescent X-ray. Therefore SiO2 wasdiluted by TiO2 nanopowder for model samples.
Figure 1 shows typical X-ray fluorescence spectraof SiO2 diluted by TiO2 nano-powders. Elastic
scattering of incident X-ray was also detected. As
energy of incident X-ray increase, peaks of elastic
scattering were shifted. Figure 2 shows XANES
spectra estimated by Si-Kα and elastic scattering ofincident X-ray. It was note that the elastic scattering
have intensities below Si-K absorption edge. Elastic
scattering have small peak at the white line of SiO2 due
to overlapping of Si-Kα and the elastic scattering.
Therefore it behave as background. We also measured
more diluted samples, the background caused by theelastic scattering were well reduced using this method.
Therefore, this method were useful for samples ofwhich fluorescent X-ray was overlapped.

Ⅲ-2

Fig. 1. Typical fluorescent X-ray spectra of SiO2
diluted by TiO2 nano-powders with different inciden
X-ray energy. Arrows indicates elastic scattering ofincident X-ray.

Fig. 2. Si-K XANES of SiO2 diluted by TiO2 nanopowders. Red and blue lines indicates spectra
measured by Si-Kα X-ray and elastic scattering ofincident X-ray.
[1] H. Murata et al., UVSOR Activity Reports 2013 41(2014) 52.
[2] H. Murata et al., UVSOR Activity Reports 2013 41(2014) 56.

53

Ⅲ02- 11 -Ⅲ -2. Materials Sciences
BL2A
BL2A

Local Environment of Multivalent Cations in Ion-Exchanged Layered
Titanate and Titanate Nanotubes
A. Fujimoto, H. Murata and A. Nakahira

Department of Materials Science, Osaka Prefecture University, Sakai 599-8531, Japan

Ⅲ-2

These results implied that Al3+ were incorporated intotitanate nanotubes and layered titanates. Detailanalyses of local environments for Al3+ in titanate
nanotubes and layered titanates have been in progressusing XANES simulation by first-principles manner.

Titanate nanotubes have attracted many researchers
due to their unique structures.[1] They consists of TiO6octahedra linked by sharing edges, which isconsidered to be based on structures of layered titanate.
Titanate nanotubes and layered titanates have cationexchange ability. Since they have variable-lengthinterlayers, various kinds of cations can beincorporated. Generally, univalent cations, M+, areeasily ion-exchanged into them. On the other hand,
multivalent cations are slightly difficult to introduce
into them. In this study, we tried to synthesize ionexchanged titanate nanotubes and layered titanateswith multivalent cations, and investigated their local
environments by X-ray absorption near edge structure(XANES).
Layered titanate samples were prepared with asolid-state reaction method. Na2CO3 and TiO2 wereused as starting materials. They were mixed andground in an agate mortar. Then, they were calcined at
1000°C for 8 h. Titanate nanotubes samples weresynthesized by the hydrothermal method based on
Kasuga et al. [2]. Anatase-type TiO2 nano-powders
and 10 mol/L of NaOH aqueous solution were sealed
into Teflon-lined autoclave and they were heated at
150°C for 24 h. Obtained samples were ion-exchanged
in aqueous solutions of HCl, MgCl2 or AlCl3.
XANES spectra were measured at BL2A in UVSOR.
Partial fluorescent yield method with a silicon drift
detector (SDD) were adapted. X-ray beam was
monochromated using beryl double crystals for Na-K
and Mg-K edge, and KTP double crystals for Al-K
edge. Samples were mounted by a carbon tape on Ti
plates. It was noted that X-ray fluorescence from ionexchanged titanate nanotubes gave rather week signals.
For example, only 10 cps for Al-K were observed at
the peak top that the incident photon energy was 1570
eV. Therefore, we set maximum 15 second/points fordwell time.
As an example of results, fig.1 shows typical Al-K
XANES spectra of Al3+-exchanged samples andreferences. Al3+-exchanged titanate nanotubes andNa2Ti3O7 have similar spectra to each other. Thismeans local environments of Al3+ in them were alsosimilar. When titanate nanotubes and Na2Ti3O7 aresoaked in aqueous solutions, pH usually increase
because of H+ incorporation into them, which often
forms hydroxides. But our samples have different
spectra from those of hydroxide and starting materials.

Fig. 1. Typical Al-K XANES spectra of titanate
nanotubes and layered titanate which were ionexchanged in AlCl3 aqueous solutions. Spectra ofreference samples were also shown.
[1] D. V. Bavykin et al., Adv. Mater. 18 (2006) 2807.
[2] T. Kasuga et al., Langmuir 14 (1998) 3160.

54

Ⅲ02- 12 -Ⅲ -2. Materials Sciences
BL2A
BL2A

Local Structure Investigations of Y3Al5O12 Induced by Swift Heavy Ions
Irradiations
S. Yoshioka1, K. Yasuda1, S. Matsumura1, N. Ishikawa2, K. Kobayashi3 and K. Okudaira4

Department of Applied Quantum Physics and Nuclear Engineering, Kyushu University, Fukuoka 819-0395, Japan
Nuclear Science and Research Center, Japan Atomic Energy Agency (JAEA), Tokai 319-1184, Japan
3
Kyushu Synchrotron Light Research Center (SAGA-LS), Tosu 841-0005, Japan
4
Graduate School of Advanced Integration Science, Chiba University, Chiba, 263-8522, Japan

1

Yttrium aluminum garnet (Y3Al5O12 or YAG) is
considered as an important material for laser opticalapplication. Ion irradiations are currently used to tailormaterials properties. In particular, for garnet structurematerials, swift heavy ion irradiations, produce
amorphous tracks that can change the microstructure
and the magnetic or magneto-optical properties[1]. In
the present study, we made a combined study of
NEXAFS measurements and first principles
calculations on Y3Al5O12 with special interests on thelocal environment of Y and Al after irradiated withswift heavy ions (SHI).
Synthetic single crystals of Y3Al5O12 were used inthis study. The crystals were cut into sheets with a(111) plane surface and were polished to a mirror
finish. The specimens were irradiated with 340 MeV
Au ions to fluences of 3 × 1011 and 1 × 1012 cm-2 at theH1 beamline of the tandem ion accelerator facility at
the Japan Atomic Energy Agency (JAEA) in Tokai. Al
K-edge and Y L3-edge XANES measurements were
performed at the BL2A beamline of UVSOR Okazaki,
Japan, using the total electron yield method (TEY). A
Beryl and an InSb double crystal monochromators
defined Al K and Y L3 absorption edges, respectively.
The samples were set with their surface perpendicular
to the incident X-ray beam. All measurements of
XANES spectra were carried out in vacuum of 1 × 10−5
Pa at room temperature. To interpret the local structurefrom the experimental spectra, we used theoretical
spectra by the full-potential linearized augmented
plane wave (APW) plus local orbitals technique as
implemented in WIEN2k code.
Figure 1 shows the XANES spectra with the Al Kedge acquired from the Y3Al5O12 samples before andafter irradiation with 340 MeV Au ions with fluences of
3 × 1011 and 1 × 1012. The Al K-edge spectrum featuresof the irradiated sample show significant change withincreasing SHI irradiation fluences. On the other hand,
the Y L3-edge spectrum features of the irradiatedsamples are very similar to that the pristine sample, as
shown in Fig. 2. These results suggest that the SHI
irradiations to the Y3Al5O12 sample preferentiallyinduced the local structure changes of Al, compared toY.

Ⅲ-2

Fig. 1. Al K-edge XANES spectra of pristine andirradiated Y3Al5O12 with 340 MeV Au ions to fluencesof 3×1011 and 1×1012 cm-2.

Fig. 2. Y L3-edge XANES spectra of pristine andirradiated Y3Al5O12 with 340 MeV Au ions to fluencesof 3×1011 and 1×1012 cm-2.
[1] J. M. Costantini, S. Miro, F. Beuneu and M.
Toulmonde, J. Phys.: Condens. Matter 27 (2015) 496001.

55

Ⅲ02- 13 -Ⅲ -2. Materials Sciences
BL2B
BL2B

Excitation Energy Dependent UPS of Solution-processed Perovskite Film1

Graduate School of Advanced Integration Science, Chiba University, Chiba 263-8522, Japan
2
Graduate School of Engineering, Chiba University, Chiba 263-8522, Japan3
Molecular Chirality Research Center, Chiba University, Chiba 263-8522, Japan

In the perovskite solar cell, the interfaces play
crucial roles in the charge extraction, transfer andrecombination. The general chemical formula for theperovskite is ABX3. The most commonly usedworking layer is the solution-processed film of
methylammonium lead iodide (MAPbI3) where A, B,
and X are methylammonium (MA), lead (Pb) andiodine (I), respectively. The electronic properties of
the interface involving the perovskite ABX3 should bestrongly affected by the elements of the surface
termination layer being either A, B, or X3.
Previously, we have examined the uppermost
surface elements of the solution-processed MAPbI3
perovskite film by comparing the valence band spectra
measured with different probing depths using
ultraviolet photoelectron (UPS) and metastable-atomelectron (MAES) spectroscopies [1]. In this work, wehave measured the excitation energy dependent UPS
in the range between 24.3 eV and 110 eV to confirm
our peak assignment.
Figure 1 shows the energy dependent UPS. Basedon the reported DFT calculations [2, 3], we tentatively
assigned the peak at 3.3~3.6 eV to the I 5p derivedstates, and the peak at 4.7 eV to the Pb 6s and 6p
derived-states. The features at 8.8 and 10.8 eV are
dominated by MA molecular states, with additional Pb
6s states at 8.8 eV.
The intensities of the peaks vary with the excitationenergy as shown in Fig. 1. The features of I 5p shows
strong intensity on the low energy and decreased as theincreases of the photon energy. Pb 6p and 6s, however,always shows lower intensity compared with I 5p. The
features of MA (N 2p and C 2p) show strong intensityand decreased on the high energy.
The observations are fully consistent with thephotoionization cross-section [4] (Fig. 2) confirming
our peak assignment.

Pb

MA

I

110 eV
100 eV
95 eV
90 eV
85 eV
80 eV
75 eV
70 eV
65 eV
60 eV
55 eV
50 eV
40 eV
38 eV
36 eV
34 eV
32 eV
30 eV
28 eV
26 eV
24.3 eV

Energy Dependence

MA

12

10

8

6

4

2

0

Binding Energy (eV)Fig. 1. Energy dependent UPS spectra of MAPbI3 film.

Photoionization cross section (Mb/atom)

Ⅲ-2

A. Mirzehmet1 and H. Yoshida2,3

10
1
0.1

N 2p
C 2p
I 5p
Pb 6p
Pb 6s

0.01
0.001
1E-4

10

20

30
40
50
Photon energy (eV)

60

70

80

Fig. 2. Photoionization cross-section of MAPbI3
elements.

[1] M. Abduheber, et al., The 79th JSAP Autumn
Meeting (2018).
[2] W-J. Yin et al., Appl. Phys. Lett. 104 (2014) 063903.
[3] J. Endres et al., J. Phys. Chem. Lett. 14 (2016) 2722.
[4] J-J. Yeh, Atomic calculation of photoionization crosssections and asymmetry parameters (Gordon and Breach
Science, Publishers, 1993).

56

Ⅲ02- 14 -Ⅲ -2. Materials Sciences
BL3B
BL3B

Correlation between Intrinsic Luminescence Intensity and Structure in(Ca1-xSrx)2Al2SiO7 and (Ca1-xSrx)2MgSi2O7 Melilite Crystals
N. Kodama, H. Kubota, T. Uematsu and T. Takahashi

Graduate School of Engineering Science, Akita University, Akita 010-8502, Japan

Crystals with a layered structure are expected to
exhibit strong luminescence from peculiar twodimensional excitons resulting from their large binding
energy. (Ca1-xSrx)2Al2SiO7 and (Ca1-xSrx)2MgSi2O7
melilite crystals have a tetragonal sheet structure, which
consists of five-membered rings of TO4 (T = Al, Si, Mg)
AlO45-, SiO44-, or MgO46- oxoanions, resulting in selftrapped excitons (STEs) under band-to-band (interband)
excitation or molecular transitions in the VUV region.
The formation energy and the luminescence properties of
excitons are considered to depend on the nature of theoxoanion-linked structure, such as the dimensionality(intra-layer distance) or two-dimensional modulation. To
date, very little research has focused on the formation ofself-trapped excitons in the dimensionality andmodulation of the structure. We examined the correlation
between the intrinsic luminescence intensity and the twodimensionality and structure modulation (superstructure)in (Ca1-xSrx)2Al2SiO7 and (Ca1-xSrx)2MgSi2O7 (x = 0, 0.25,
0.5) melilite crystals. We report intrinsic luminescencefrom self-trapped excitons (STEs). In addition, we
discuss the structural correlation of the intrinsic
luminescence intensity from STEs.
Luminescence and excitation spectra of STEs in (Ca1-x
Srx)2Al2SiO7 and (Ca1-xSrx)2MgSi2O7 (x = 0, 0.25, 0.5)
crystals were measured in the temperature range of 14293 K using the undulator beamline BL3B at the UVSOR
facility. Under VUV excitation at 160 and 70 nm,
intrinsic luminescence from self-trapped exitons (STE)
was observed in these crystals.
The dependence of the intrinsic luminescenceintensity on the Sr fraction x was examined for thesecrystals. In (Ca1-xSrx)2Al2SiO7, one luminescence bandwith a peak at 271-300 nm associated with STE appears
at temperatures of 14-293 K. The luminescence spectraof (Ca1-xSrx)2MgSi2O7 with x = 0.0, 0.25, and 0.5
consisted of two broad bands with peaks at 238-246 nm
and 340-350 nm associated with STEs in the range of
14-293 K.
Figures 1(a) and 1(b) show the luminescence spectrafor (Ca1-xSrx)2Al2SiO7 and (Ca1-xSrx)2MgSi2O7 (x = 0.0,
0.25, 0.5) excited at 160 nm at 14 K.
In (Ca1-xSrx)2Al2SiO7, the crystal with x = 0.5
exhibited one intense STE luminescence band, and theSTE luminescence intensity decreases rapidly at x =0.25. No significant luminescence was observed at x =0.0. On the other hand, in (Ca1-xSrx)2MgSi2O7, theintensity of STEs passes through a maximum atintermediate fraction x = 0.25, as x increases. For the
dependence of luminescence intensity on the Sr fraction,
x, two possibilities are plausible: one is a two-

dimensional effect and the other is a structuralmodulation effect (superstructure).
First, the intra-layer distance increases withincreasing x, because of the larger ionic radius involvedin substituting Sr2+. One possible explanation for suchan intense luminescence at x = 0.5 in (Ca1-xSrx)2Al2SiO7and (Ca1-xSrx)2MgSi2O7 (x = 0.0, 0.25, 0.5) is that alarger two-dimensional effect, i.e., the intra-layer
distance increases with increasing Sr fraction x, whichleads to larger two-dimensionality. However, theenlargement of intra-layer distance is not very large
even at x = 0.5 (a few percent). The second possible
explanation is that a modulation effect of commensurate
or incommensurate structure leads to an increase inlocal sites with larger distortion potentials. The
formation of STEs and the luminescence intensity may
be strongly affected by whether the crystals form a
modulated structure (commensurate or incommensurate
phase) and by the amplitude of modulation. We foundthat a (Ca1-xSrx)2Al2SiO7 crystal with 0.5, which shows
the most intense luminescence among the three x
fractions (0.0, 0.25, and 0.5), exhibited commensurate
modulation with two-fold periodicity along the c-axisby X-ray diffraction. Moreover, incommensurately andtwo-dimensionally modulated structures with Sr
fractions, x, ranging from 0.04 to 0.32 have been
observed in (Ca1-xSrx)2MgSi2O7 [1]. Based on thesestructural changes with Sr composition, the intense
luminescence at x = 0.5 is considered to be due mainly
to a structural modulation effect.
λex = 160 nm
Ca2Al2SiO7 (T=8K)(Ca0.75Sr0.25)2Al2SiO7 (T=10K)(Ca0.5Sr0.5)2Al2SiO7 (T=14K)

0.2
0.15

0.006

0.1

0.004
0.002

0.05
0

0

200

300

230

280

320

3
Intensity / a.u.

Intensity / a.u.

0.25

380

400 500 600
Wavelength / nm

700

2

λex =

160nm

Ca2MgSi2O7(Ca0.75Sr0.25)2MgSi2O7(Ca0.50Sr0.50)2MgSi2O7

1
0
200

300 400 500
Wavelength / nm

600

(a)
(b)
Fig. 1. Dependence of intrinsic luminescence on Sr
fraction x in (a) (Ca1-xSrx)2Al2SiO7 and (b) (Ca1-x
Srx)2MgSi2O7 (x = 0.0, 0.25, 0.5) at 14 K under
excitation at 160 nm. The inset shows the enlarged
luminescence spectra in the range from 230 to 390 nm.
[1] J. C. Jiand, M. Schosnig, A. K. Schapaer, K. Ganster,H. Ragaer and L. Toth, Phys. Chem. Minerals 26 (1998)
128.

57

Ⅲ02- 15 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL3B
BL3B

Evaluation of Fluorescence Lifetimes of Scintillation Fibers
H. Oikawa, M. Kitaura, A. Ohnishi, Y. Tajima and H. Y. Yoshida
Yamagata University, yamagata 990-8560, Japan

Ⅲ-2

In high-energy physics experiments, particle
detectors with fast time response are absolutely
imperative. Especially, plastic scintillators are the first
choice for radiation converter, because they havefluorescence with short decay time of nanosecond
range. This feature is necessary for successive
detection of particles at high count rates. In order todevelop plastic scintillators, it is indispensable toevaluate fluorescence lifetimes with the use of a shortpulse light source in the wide range from visible tovacuum ultraviolet. In the present study, we havemeasured fluorescence spectra and decay curves forsome polystyrene-based plastic scintillators.
Samples used in our experiment can be available
commercially, and they are labelled as Y-8, Y-11, B-3,
O-2, R-3, and SCSF-78. As an example, thefluorescence spectrum of SCSF-78 is shown in Fig. 1.
This spectrum was obtained under excitation at 200
nm. Two bands peaking at 335 nm and 435 nm appear
in Fig. 1. Fluorescence decay curves were measured
under single bunch operation by adopting timecorrelated single photon counting method. The decay
curve for the 335 nm band is shown in Fig. 2. The blueline indicates experimental data.
Lifetimes were analyzed by using the data analysis
framework named “ROOT” [1]. For the lifetime
analysis, we adopted the equation�I(𝑡𝑡) � � P(𝑡𝑡 � ) G(𝑡𝑡 𝑡 𝑡𝑡 � )𝜃𝜃(𝑡𝑡 � − 𝑇𝑇� )d𝑡𝑡 � ,(1)

P(𝑡𝑡) � � ���(−𝑡𝑡𝑡𝑡𝑡� ) ,

(2)

where

and

��
�

curves suggest the existence of plural relaxation
pathway under excitation in the fundamental absorption
region of host polystyrene.

Fig. 1. Emission spectrum of SCSF-78, measured atroom temperature under excitation at 200 nm.

Fig. 2. Fluorescence decay curve for the 335 nm bandof SCSF-78, measured at room temperature under
excitation at 200 nm.

(𝑡𝑡 𝑡 𝑡𝑡 � − 𝜇𝜇� )�
�
2𝜎𝜎��
�
(𝑡𝑡 𝑡 𝑡𝑡 − 𝜇𝜇� )�
� � (�)� ���� �𝑡
2𝜎𝜎��Table. 1. Fluorescence peak-wavelengths and lifetimes
determined in the present study.

G(𝑡𝑡 𝑡 𝑡𝑡 � ) � ���� �𝑡sampleλEX (nm)λEM (nm)τ1 (ns)τ2 (ns)τ3 (ns)

200

335

0.934

4.68

17.4

200

435

3.33

14.7

277

335

0.694

277

435

2.10

351

435

2.01

Y11

430

495

6.92

Y8

455

526

7.24

B3

351

442

2.04

O2

535

577

5.32

R3

550

606

6.70

SCSF78I(𝑡𝑡) represents the fluorescence decay curve observed.
P(𝑡𝑡) represents the superposition of fluorescence decay
curve functions. G(𝑡𝑡 𝑡 𝑡𝑡 � ) represents the resolution of
the experimental system, the time structure of the beam,and so on. A sum of two Gaussians (3) successfully
describes the measurement of the G(𝑡𝑡 𝑡 𝑡𝑡 � ) in thisanalysis. 𝜃𝜃(𝑡𝑡 � − 𝑇𝑇� ) is assumed a step function to
reproduce the origin of time 𝑇𝑇� . Table 1 shows
fluorescence lifetimes of samples determined in thepresent study, together with the excitation and emission
wavelengths. Decay curves for the 355 nm and 435 nm
bands of SCSF-78 are reproduced by two or three
exponential decay functions. Such complicated decay[1] R. Brun and F. Rademakers, Nucl. Instrum. Methods
Phys. Res. A 389 (1997) 81.

58

Ⅲ02- 16 -Ⅲ -2. Materials Sciences
BL3B, BL7B
BL3B, BL7B

Effect of Ce and Mg Concentration Ratio on the Properties of
Gd3Al2Ga3O12 Single Crystal Scintillators
K. Bartosiewicz1, A. Yoshikawa1,2,3, S. Kurosawa2,4, A. Yamaji1 and M. Nikl5

Institute for Material Research, Tohoku University, Sendai 980-8577, JapanNew Industry Creation Hatchery Center, Tohoku University, Sendai 980-8579, Japan3
C&A corporation, T-Biz, Sendai 980-8579, Japan
4
Facility of Science, Yamagata University, Yamagata 990-8560, Japan
5
Institute of Physics, Academy of Sciences of the Czech Republic, Na Slovance 1999/2, Prague 8, 18221,
Czech Republic
1

2

Normalized intensity

1.4
1.2

Mg=0ppm
Mg=500ppm
Mg=1000ppm
Mg=2000ppma) GAGG:Ce0.25%at

1.0
0.8
0.6
0.4
0.2
0.0
50

100

150

200

T (K)

250

1.4

Normalized intensity

In gadolinium gallium aluminum garnets(Gd3Al2Ga3O12, GAGG), efficient luminescence isdependent on dopants to create a radiative transition
within the forbidden band [1]. In medical imaging, Ce3+
is commonly used due to fast and bright 5d→4ftransition with the emission wavelength that is suitablealso for semiconductor photodetectors [2]. The Ce3+
activated GAGG is considered as a promising candidate
for the next generation Positron Emission Tomographymaterial due to its high light yield value (~58,000
ph/MeV), fast scintillation response (~100 ns), and highdensity (~6.2 g/cm3). However, this material suffersfrom the contribution of the undesired slow component
in the scintillation response and long rise time [3].
Recently, the divalent (Me2+) codoping strategy has
been used in order to mitigate the degraded scintillation
properties that result from charge carrier traps and havebeen shown to improve light yield and decay times insome scintillators [4].
The Mg2+ codoping in GAGG:Ce single crystal
significantly accelerates the scintillation decay time,
but at the same time the light yield become strongly
reduced [4]. To deeper understand the effect of Mg2+
codoping content on the scintillation characteristics of
GAGG:Ce, we recorded thermally stimulated
luminescence (TSL), see Fig.1. The results show that
increasing Mg2+ concentration creates new peaks
within the whole recorded temperature range. Thismeans that Mg2+ codoping increases the number ofmetastable levels, which act as electron traps. Due tothe transformation of Ce3+ to Ce4+ crystal lattice
become perturbed. Consequently, many defects arise,
which can significantly diminish scintillation
performances. Moreover, a rich pattern of TSL peaks
over RT is due to the presence of the deep electron
traps, which are associated with oxygen vacancies. To
conclude, TSL measurements revealed that Mg2+
codoping reduces the quality of the crystal lattice dueto the creation of many defects acting as electron traps.

1.2

300

350

Mg=0ppm
Mg=500ppm
Mg=1000ppm
Mg=2000ppmb) GAGG:Ce0.75%at

1.0
0.8
0.6
0.4
0.2
0.0
50

100

150

200

T (K)

250

300

350

Fig. 1. Normalized TSL glow curves (measured in the
10-3800 K temperature range for the for Mg2+ free andMg2+ codoped GAGG:Ce single crystals. Nominal
content of Ce3+ was (b) 0.25% and (b) 0.75% at,
irradiated at 10 K by UV radiation (12.4 eV)[1] M. Nikl and A. Yoshikawa, Adv. Optical Mater. 3(2015) 463.
[2] C. Ronda, J. Solid State Sci. Technol. 5 (2016)
R3121.
[3] M. Lucchini et al., Instr. Meth. Phys, Research A
852 (2017) 1.
[4] C. Foster et al., J. Cryst. Growth, in press(10.1016/j.jcrysgro.2018.01.028).

59
Ⅲ02- 17 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL3B
BL3B

Estimation of Optical Properties of RE2Si2O7 Sintered Compacts Using theUVSOR Synchrotron Facility
T. Horiai1, S. Kurosawa2, A. Yamaji1, S. Kodama1, S. Yamato1, Y. Shoji1,4, M. Yoshino1,
H. Sato2, Y. Ohashi2, K. Kamada2,3, Y. Yokota2 and A. Yoshikawa1,2,3
2

Ⅲ-2

1
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, Aramaki, Sendai 980-8579, Japan

Scintillation crystals convert energy of ionizing
radiation such as gamma-ray and X-ray into multiple
photons of energy 2~8 eV, and are used in many fields,
as radiation detectors [1-3]. The Ce-doped RE2Si2O7(RE= rare-earth ions) scintillators have been reported
to have a high light yield, short decay time and goodthermal stability. Here, the crystal system of RE2Si2O7
are changed by rare-earth ions [4]. In addition, theemission and excitation spectra of Ce3+ 5d-4ftransition are strongly depend on the host crystal
structure of RE2Si2O7 due to the involved Ce3+ 5dorbital are not shielded by outer electrons. Therefore,
it is important to estimate the influence of crystal
structure of RE2Si2O7 and the ionic radii of rare-earth
ions on optical properties of Ce-doped RE2Si2O7. In
addition, the bandgap energy of RE2Si2O7 is also
affected when the rare-earth ion changes. In this study,
we estimated the bandgap energy of RE2Si2O7 (RE=La,
Gd, Y and Lu) to investigate the change in bandgap
energy due to the ionic radii of rare-earth ions.
We prepared (Ce0.01 RE0.99)2Si2O7 (RE=La, Gd, Y
and Lu) sintered compacts.; As starting materials, weused La2O3, CeO2, Gd2O3, Y2O3, Lu2O3 and SiO2
powders the purities of which were at least 99.99%.
After weighing, the powders were well wet-blended,and they were molded into pellets using a hydraulicpress, and the pellets were pre-sintered at 1550 ºC for
12 hours under air. Finally, we sintered the pellets at
1600 ºC for over 24 hours under air. The emission
spectra of these sintered compacts were measured witha spectrometer (spectoropro-300i, Acton research),and the excitation spectra were measured with a photo
diode (IRD, AXUV 100) at BL3B of UVSOR.
Figure 1 shows the emission spectra of Ce-doped
RE2Si2O7 (RE=La, Gd, Y, Lu). From this figure, theemission peak due to Ce3+ 5d1-2f5/2 was changed withchanges in the ion radii of rare-earth ions. The
excitation spectra of Ce-doped RE2Si2O7 (RE=La, Gd,
Y, Lu) is shown in Fig. 2. The result of estimating thehost excitation band peak showed that the bandgap
energies of La2Si2O7, Gd2Si2O7, Y2Si2O7 and Lu2Si2O7
were approximately 7.0, 6.9, 7.2 and 6.9 eV,
respectively. All samples were found to have highbandgap energy of over 7.0 eV.

(Ce0.01 La0.99)2Si2O7(Ce0.01 Gd0.99)2Si2O7(Ce0.01 Y0.99)2Si2O7(Ce0.01 Lu0.99)2Si2O7

1.0
0.8
0.6
0.4
0.2
0.0

350

400
450
500
Wavelength [nm]550

Fig. 1. Emission spectra of Ce-doped RE2Si2O7 (RE=La,
Gd, Y and Lu) at 6 K

Fig. 2. Excitation spectra of Ce-doped RE2Si2O7(RE=La, Gd, Y and Lu) at 6 K[1] C. Rozsa, R. Dayton, P. Raby, M. Kusner and R.
Schreiner, IEEE Trans. Nucl. Sci. 37 (1996) 966.
[2] A. Baberdin, A. Dutova, A. Fedorov, M. Korzhik,
V. Ligoun, O. Missevitch, V. Kazak, A. Vinokurov andS. Zagumenov, IEEE Trans. Nucl. Sci. 55 (2008) 1170.
[3] C.L. Melcher, J.S. Schweitzer, R.A. Manente andC.A. Peterson, J. Crystal Growth 109 (1991) 37.
[4] J. Felsche, Journal of the Less Common Metals 21(1970) 1.

60
Ⅲ02- 18 -Ⅲ -2. Materials Sciences
BL3B
BL3B

Temperature Dependence of PL Spectra of p-terphenyl and1,4-bis(2-methylstyryl)benzene (bis-MSB) at 6-300 K
S. Yamato1, S. Kurosawa2,3, A. Yamaji1, T. Horiai1, S. Kodama1, M. Abe3, M. Yoshino1,
Y. Ohashi2, K. Kamada2,4, Y. Yokota2 and A. Yoshikawa 1,2,4
1
Institute for Materials Research (IMR), Tohoku University, Sendai 980-8577, JapanNew Industry Creation Hatchery Center (NICHe), Tohoku University, Sendai 980-8579, Japan
3
Fuculty of science, Yamagata University, Yamagata 990-8560, Japan
4
C&A Corporation, Sendai 980-8577, Japan.

2

Neutron detectors can be used in several fields suchas neutron diffraction for crystal structure analysis,
social infrastructure maintenance and nuclear fusion
physics [1-3]. Here, nuclear fusion technology is
expected as next effective energy generation technology.
In fusion reactors, neutrons are emitted from nuclear
fusion reactions, and the monitoring these particles is a
important technology to control fusion reactors andgenerate energy effectively.
Among several nuclei, the neutron cross section ofhydrogen nucleus is the largest in higher energy range.
1
H has relatively higher cross section for the fast neutron
than other materials. We focused on organic scintillators
which contains many 1H and have low detection
efficiency for gamma rays as background noise.
Although the neutron detectors in fusion reactors canbe exposed on high temperature of about from 400 to
600 K, conventional organic scintillators cannot be usedat high temperature. Plastic scintillators melt at around230 K. Liquids evaporate at high temperature and are
difficult to handle. Even crystals, only low-meltingpoint materials were developed. As a first step, we
aimed to develop the organic scintillators which can beused at over 270 K s. Also, I considered high-meltingpoint crystal have a potential to meet this demand.
We grew pure p-terphenyl crystal and pure 1,4-bis(2methylstyryl)benzene (bis-MSB) crystal by the selfseeding vertical Bridgman method using an enclosed
chamber [4]. Pulling down rate for growth was ~1 mm/h.
The melting point of p-terphenyl and bis-MSB are 485
and 352 K, respectively.
We measured the photoluminescence (PL) emission
spectra of the samples at 6-300 K with the beamline
BL3B at UVSOR facility. The spectra of p-terphenyl
excited by 245 nm photons and bis-MSB excited by 260
nm photons were shown in Fig. 1 and Fig. 2, respectively.
At 300 K the emission peaks of 375 and 390 nm wereobserved for p-terphenyl. As temperature decreased, theratio of peak of 375 nm for 390 nm became smaller andnew peak of 350 nm appeared. For bis-MSB, theemission peak of 484 nm became weaker and several
peaks were appeared in shorter wavelength range as
temperature decreased.
Since the molecular vibration was suppressed atlower temperature, emission peaks were split clearly. As
future works, we study on the relationship between

these results and temperature dependence of photoluminesce decay time.

Ⅲ-2

Fig. 1. Emission spectra of p-terphenyl crystal excitedby 245 nm photons at 6-300 K.

Fig. 2. Emission spectra of bis-MSB crystal excitedby 260 nm photons at 6-300 K.
[1] P. Zhang et al, Cement and Concrete Research 108(2018) 152.
[2] H. K. Lee et al., J. Supercond. Novel Magn. 31(2018) 1677.
[3] Y. K. Kim et al., J. Instrum. 7 (2012) C06013.
[4] A. Arulchakkaravarthi et al., J. Crystal Growth 234(2002) 159.

61
Ⅲ02- 19 -Ⅲ -2. Materials Sciences
BL3B
BL3B

Low-temperature Photoluminescence Spectra of Cs2HfI6
S. Kodama1, S. Kurosawa2,3, M. Ohno1, A. Yamaji1, M. Yoshino1, H. Sato2,
Y. Ohashi1, K. Kamada2,4, Y. Yokota2 and A. Yoshikawa1,2,4

Institute for Materials Research (IMR), Tohoku University,Sendai 980-8577, JapanNew Industry Creation Hatchery Center (NICHe), Tohoku University, Sendai 980-8579, Japan
3
Faculty of Science, Yamagata University, Yamagata 990-8560, Japan
4
C&A Corporation, Sendai 980-8577, Japan.
12

Intensity (a.u.)

(a)

12000

300K
150K
8K

10000
8000
6000
4000
2000
0

(b)
Intensity (a.u.)

Ⅲ-2

A2HfX6 (A: monovalent cation, X: halogen ion)compounds are host-emitting materials originatedfrom self-trapped exciton (STE) [1]. Since the crystal
lattice of A2HfX6 compounds can be a well-ordered
and high symmetry cubic structure (space group: Fm3m), producing a transparent single crystal is easy.
Actually, Our research group has developed a Cs2HfI6(CHI) single crystal as a noble red-emitting (around
650 nm) scintillator with a high light output of ~64,000
photons/MeV [2,3]. As mentioned above, theluminescence mechanism of CHI could be STE around[HfI6]2-, however, the detail of emission mechanism
was still unclear. In this year, we evaluated thetemperature dependence of CHI photoluminescence
properties to study the luminescent origin.
A single-crystalline CHI specimen was synthesizedfrom 99%-pure HfI4 and 99.999%-pure CsI using thevertical Bridgman growth method in our laboratory.
We measured photoluminescence emission spectra of
CHI at from 10 K to 300 K cooled by liquid Helium ina cryostat chamber of UVSOR BL3B. The light source
was a synchrotron light.
Figure 1 indicate the photoluminescence emission
spectra of CHI. When the sample was excited by 400
nm at 300 K in Fig. 1 (a), the typical broad emission
spectrum around 650 nm was observed. At 150 K one
additional peak appeared at 550 nm, and as
temperature decrease, the 550-nm-peak became very
intense, while the intensity of 650-nm-peak decreased.
The 550-nm-peak intensity leached the maximum at
10 K. CHI room temperature (300 K) emission excited
460 nm showed the different temperature dependence.
To 8 K, the peak intensity increased, and wavelength
shifted towards 700 nm.
Based on above temperature dependence, CHI was
considered to have several excitation-emission bands.
Moreover, the major band could change as temperaturechanges. Since the typical STE emission intensity
increases at low temperature, the low temperatureemission peak at 550 nm could be one STE band which
thermally quenches above 150 K. Further study shouldbe needed to understand the energy transfer between
room temperature emission band and low temperatureband. For the direct evaluation of STE center, electron
paramagnetic resonance technique would be available.

400

2500

500 600 700 800
Wavelength [nm]

900

300K
150K
8K

2000
1500
1000
500
0

400

500 600 700 800
Wavelength [nm]900Fig. 1. Photoluminescence emission spectra of CHI
crystal. (a) and (b) are spectra excited by 400 nm and 460
nm, respectively.
[1] B. Kang and K. Biswas, J. Phys. Chem. C. 120 (2016)
12187. doi:10.1021/acs.jpcc.6b02496.
[2] S. Kodama, S. Kurosawa, M. Ohno, A. Yamaji, M.
Yoshino, J. Pejchal, R. Král, Y. Ohashi, K. Kamada, Y.
Yokota, M. Nikl and A. Yoshikawa, Radiat. Meas. inprint (2019).
[3] S. Kodama, S. Kurosawa, A. Yamaji, J. Pejchal, R.
Král, Y. Ohashi, K. Kamada, Y. Yokota, M. Nikl and A.
Yoshikawa, J. Cryst. Growth. 492 (2018) 1.

62
Ⅲ02- 20 -Ⅲ -2. Materials Sciences
BL3B
BL3B

Photoluminescence upon Vacuum Ultraviolet Excitationin Pr3+ Ion Doped Alkaline–Earth Phosphates
Y. Inaguma1, R. Horiguchi1, S. Sasaki1, M. Yaname1, D. Mori1, K. Ueda1 and T. Katsumata2
Department of Chemistry, Facility of Science, Gakushuin University, Tokyo 171-8588, Japan2
Department of Chemistry, School of Science, Tokai University, Hiratsuka 259-1292, Japan1

Pr3+ ion doped inorganic materials have been
extensively investigated in connection with various
applications [1]: laser materials utilizing visible
emission corresponding to 3P0 → 3H4 and 1D2 → 3H4
transitions(blue-green and red emission, respectively),or infrared emission from the 1G4 state of Pr3+, fieldemission display (FED) devices utilizing red 1D2emission, scintillators or tunable ultraviolet (UV)
lasers utilizing UV emission from the 4f5d state. Since
the emission color or energy in Pr3+ ion doped
phosphors is closely related to the electronic structureof host materials, the choice of host materials is amatter of vital importance to control the emission
color/energy [2].
Recently phosphor materials emitting UV lightupon vacuum ultra violet (VUV) excitation are desiredas the substitutes of mercury lamp due to the toxicity
of mercury. Pr3+ ion doped compounds are candidatesof the UV phosphors because the UV emission fromthe 4f5d state in Pr3+ can be expected. The UVemission from the 4f5d state would be observed if theband gap energy of the host material corresponds tothe VUV region. We then chose Pr3+ ion doped
alkaline-earth phosphates with the band gap energy inthe VUV region and investigated their luminescence
properties. In our previous report[3], we presented theresults of UV photoluminescence upon VUVexcitation of Ca3(PO4)2:Pr. Herein, we report on the
UV photoluminescence of Pr-doped alkaline–earth
phosphates A3(PO4)2:Pr(A = Sr, Ba) in addition to
Ca3(PO4)2:Pr.
The polycrystalline samples were synthesized by asolid state reaction. The mixture of starting materials:
alkaline-earth carbonate ACO3(A = Ca, Sr, Ba),
(NH4)H2PO4 and Pr nitrate solution were calcined at1000 C in air. The calcined powder was pressed into
pellets and sintered at 1000-1200 C in air. The phase
identification for the samples was carried out by thelaboratory powder X-ray diffraction (XRD) using a
Rigaku RINT 2100 diffractometer or a PANalytical
X’Pert3 Powder diffractometer with a Bragg Brentano
geometry (CuKα radiation). The emission andexcitation spectra were recorded in the beamlines
BL3B and BL7B at the UVSOR facility.
Figure 1 displays the emission(right side) andexcitation(left side) spectra for A3(PO4)2:Pr5%(A = Ca,
Sr, Ba) at room temperature. As seen in Fig.1, the UVemission peaks corresponding to the transitions fromthe 4f5d state to the 4f2 (3HJ, 3FJ) states of Pr3+ wereobserved. In the excitation spectra, the peaks in therange of 150-220 nm were observed. The shorter andlonger wavelength regions correspond to the hostabsorption band, i.e. the intra charge transfer of PO43[4] and the transition from the ground state of 4f2(3H4)
to the 4f5d state of Pr3+, respectively. The emission andexcitation peaks for A3(PO4)2:Pr(A = Sr, Ba) appears at
the shorter wavelength than those for Ca3(PO4)2:Pr,
which is attributable to the difference in covalency of
Pr-O bonding.
Consequently, we found that A3(PO4)2:Pr(A = Sr, Ba)as well as Ca3(PO4)2:Pr exhibits the UV emission
corresponding to the transition from the 4f5d state of Pr3+
upon the VUV excitation and the Pr3+ ion doped alkaline–
earth phosphates are candidates of UV phosphors.
The authors thank Mr. M. Hasumoto, Mr. T. Yanofor their experimental supports.

Fig. 1. Emission spectra (right side) and excitation
spectra emission (left side) for A3(PO4)2:Pr5% (A = Ca,
Sr, Ba) at room temperature.
[1] E. Van der Kolk et al., Phys. Rev. B. 64 (2001)
195129.
[2] Y. Inaguma et al., Inorg. Chem. 50 (2011) 5389.
[3] Y. Inaguma et al., UVSOR activity report 2011 39(2012) 129.
[4] H. Liang et al., J. Solid State Chem. 177 (2004)
901.

63
Ⅲ02- 21 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL3B
BL3B

Emission Spectra of Pr-doped BaZrO3
M. Yoshino, J. Sugiyama, T. Kuyama and S. Watanabe

300

BaZrO 3:Pr
BaZrO 3ex

= 265 nm

300

500
600
Wavelength (nm)Intensity (arb. units)

Y15%-doped
BaZrO3:Pr= 265 nm700

= 490 nmem

before hydration
after hydration

250
300
Wavelength (nm)350

Fig. 3. Change in excitation spectra of Y-doped
BaZrO3:Pr with hydration. (λem= 490 nm).

Y15%-doped
BaZrO3:Pr= 616 nmem

before hydration
after hydration

200
700

ex

400
500
600
Wavelength (nm)200

x 2000400

BaZrO3:Pr
Y5%-doped
BaZrO3:Pr
Y15%-doped
BaZrO3:Pr

Fig. 2. Emission spectra of BaZrO3:Pr and Y-doped
BaZrO3:Pr. (λex= 265 nm).

Intensity (arb. units)

Intensity (arb. units)

Ⅲ-2

In this study, the excitation spectra and emission
spectra of Pr3+ in BaZrO3 crystal have been measured
and changes in the spectra with doping of Yttrium andhydration are examined. The samples are synthesizedby solid state reactions followed by annealing in N2H2 atmosphere in order to reduce Pr4+ to Pr3+. The
oxygen vacancies are created in Y-doped BaZrO3 andprotons are incorporated by hydration of the vacancies.
The concentration of Pr is 0.5 mol% and that of Y are
5 mol% and 15 mol%. The Y-doped samples are
hydrated at 523 K in humidified N2.
The emission spectra of BaZrO3:Pr and BaZrO3 hostat 265 nm excitation are shown in Fig. 1. The sharp
peaks of BaZrO3:Pr around 500 nm and 600-650 nm
originate from 4f-4f transitions of Pr3+. Since the broad
emission peaks around 350-550 nm appear in bothBaZrO3:Pr and BaZrO3 samples, these originate fromelectron transitions in host rather than from 4f-5dtransitions of Pr3+. The emission spectra of BaZrO3:Pr
and Y-doped BaZrO3:Pr shown in Fig. 2. The
intensities of Pr3+ emissions increase with 5 mol%doping of Y, but they do not increase monotonicallyand become smaller in 15 mol%. It is supposed that Y
substituted for Zr and oxygen vacancy promote theradiative transition in Pr3+ by decreasing the site
symmetry of the Pr3+, but oxygen vacancy prevent thetransition with creation of defect level in the bandgap
in high doping concentration. The excitation spectrumof Y15mol%-doped samples monitored at 490 nm and
616 nm are shown in Fig. 3 and Fig. 4. It is found thatthe changes in the emission intensity with hydration
around 265 nm excitation is larger in 490 nm than in
616 nm as is the case in the change with increasing Y
concentration in Fig. 2. The intensities around 220-250
nm decrease in both 490 nm and 616 nm. These two
characteristic changes may be due to the defect levelsof oxygen vacancy and hydrogen, respectively.

Intensity (arb. units)

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

250
300
Wavelength (nm)350

Fig. 4. Change in excitation spectra of Y-doped
BaZrO3:Pr with hydration. (λem= 616 nm).

Fig. 1. Emission spectra of BaZrO3:Pr and BaZrO3.
(λex= 265 nm).

64
Ⅲ02- 22 -Ⅲ -2. Materials Sciences
BL3B
BL3B

Temperature Dependence of Dielectric Function in Widegap Oxide Epilayer
T. Makino, T. Takeuchi and T. Asai

Depeartment of Electrical and Electronics Engineering, University of Fukui, Fukui 910-8585, Japan
Oxide-based strongly electron correlated materials
are important from both fundamental and applied
perspectives, due to their potential in the
optoelectronic and electronic devices.
A detailed understanding of the bulk and surface
optical properties of these materials is essential inorder to make the best selection. However, in spite ofits significance, very little work has been done on the
optical properties of oxide-based strongly electron
correlated materials, especially in the case of rareearth compounds [1].
ε-Ga2O3 has various advantages as a wide band-gap
semiconductor, and it is expected to be applied to
power devices such as high frequency devices bytaking advantage of this. In this way, it is essential forthe detailed understanding of the bulk and surface
optical properties of this material for the best choice.
In recent years, its crystal symmetry has been
identified; therefore little has been reported on the
optical properties of ε-Ga2O3 at present. In thisresearch, we aim to elucidate the interaction between
electrons and phonons in ε-Ga2O3 and report on thetemperature dependence of its dielectric function.
Figure 1 shows temperature dependence of thedielectric functions taken for Ga2O3 thin film. Only thereal part is shown. The situation in the case of double
layer (thin films grown on substrates) becomes very
complicated. In case of bulk crystals free from theirsubstrates, the conversion is rather simple. Althoughboth the transmissivity and reflectivity can beanalytically represented by using Fresnel’s effective
coefficient method with dielectric functions andthickness of the thin film, the dielectric functions
cannot be represented with transmissivity andreflectivity. In our case, some model dielectric
functions were adopted for this purpose.

Ⅲ-2

Fig. 1. Temperature dependence of real parts of dielectric
function for ε-Ga2O3 thin film.
[1] R. Roy, V. G. Hill, E. F. Osborn, J. Am. Chem. Soc.
74 (1952) 719.
[2] H. Playford et al., Chem. Eur. J. 19 (2013) 2803.
[3] F. Mezzadri et al., Inorgan. Chem. 55 (2016) 12079.
[4] H. Nishinaka et al., Jpn. J. Appl. Phys. 55 (2016)
1202BC.
[5] T. Uchida et al., Appl. Phys. Express 7 (2014) 021303.
[6] J. C. Lu et al., J. Cryst. Growth 299 (2007) 1.
[7] Y. P. Varshni, Physica (Utrecht) 34 (1967) 149.
[8] H. Qi et al., J. Quant. Spectros. Rad. Transfer 112(2011) 2507.
[9] M. A. Ordal et al., Appl. Opt. 22 (1983) 1099.

65
Ⅲ02- 23 -Ⅲ -2. Materials Sciences
BL3B

BL3B

Appearance of the Luminescence Band due to Ag− Centers by VUVIrradiation at Room Temperature in Co-doped NaCl:I−,Ag+ Crystals
T. Kawai and O.Yagi

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

VK centers created under photo-excitation at the I−
center would play an important role on the conversion
from the Ag+ to Ag− ions. Such valence change
processes might lead to the potential for producing theuncommon anions such as Au− and Pt2−.
500

400

Intensity (arb. unit)

RT

Wavelength (nm)
300

200

Excitation

Irradiation

A' lumi.
A+ lumi.
C
NaCl:Ag

NaCl:I -,Ag+-2.5 3.0 3.5 4.0 4.5 5.0 5.5 6.0 6.5 7.0
Photon Energy (eV)Fig. 1. Typical luminescence (solid curves) andabsorption (broken curves) spectra of NaCl:Ag− (blue)
and NaCl:I−,Ag+ (red) crystals at RT.
500-

Luminescence Intensity (arb. units)

Ⅲ-2

Since silver (Ag) atoms are normally present ascation impurities in ionic crystals, Ag− centers are notnaturally existed in alkali halide crystals. The Ag−
centers in alkali halides are created by the electrolytic
coloration, additive coloration, and/or X- and γ-ray
irradiation [1, 2]. In general, the Ag− centers are notcreated by irradiation of light in the vacuum ultra
violet (VUV) region. We found that the irradiation of
a VUV light on co-doped NaCl:I−,Ag+ crystals leads tothe appearance of the A’ luminescence due to the Ag−
centers through the experiment with the BL3B beamline.
Figure 1 shows the typical luminescence andabsorption spectra of NaCl:Ag− and NaCl:I−,Ag+crystals at room temperature (RT). The A’
luminescence and the C absorption bands due to theAg− center in NaCl:Ag− are observed at 2.81 and 4.46
eV, respectively. In co-doped NaCl:I−,Ag+ crystals, theabsorption and luminescence bands are observed at
6.78 and 5.02 eV, respectively. The absorption andluminescence bands due to the I− center are located at
the same energy positions as those due to the Ag+
center [2, 3]. Thus, the absorption and luminescencebands observed in the co-doped NaCl:I−,Ag+ crystals
consist of the sum of both bands due to respective
impurity centers.
Figure 2 shows the appearance and enhancement of
the A’ luminescence band obtained under excitation at
4.5 eV by prolonged irradiation of 6.74 eV light at RT.
The 6.74 eV light corresponds to the absorption bands
due to the I− and Ag+ centers. The enhancement of theA’ luminescence intensity is plotted as a function ofirradiation time of VUV light at 6.74 eV in the insert
of Fig. 2.
The appearance and enhancement of the A’
luminescence imply that the Ag− centers are created by
the irradiation of VUV light at 6.74 eV; that is to say,the conversion from the Ag+ to Ag− ions occurs in
NaCl host lattice at RT. When the VUV lightcorresponding to the absorption band due to the Ag+
centers is irradiated to single-doped NaCl:Ag+ crystals,
the A’ luminescence does not appear. Thus, theappearance of the A’ luminescence in the co-doped
NaCl:I−,Ag+ crystals would be related with theexistence of the I− center band.
In our recent study [4], the photoexcitation on the I−
center band at RT induces the creation of the electrons
and the VK centers in NaCl host crystals. The electrons
and the VK centers migrate through the host lattice at
RT and are trapped at impurity centers. Though thedetailed mechanism is unknown, the electrons and the+

Wavelength (nm)
400

300

NaCl:I ,Ag Exc. 4.5 eV
A' lumi.
RT
NaCl:I-, Ag +RT

120 min60 min30 min10 min3 min
0 min

Irradiation: 6.74 eV
0

2.5

50
100
Irradiation Time (min)3.0
3.5
Photon Energy (eV)

4.0

Fig. 2. Appearance and enhancement of the A’
luminescence in NaCl:I−,Ag+ by prolonged irradiationof 6.74 eV light at RT.
[1] K. Kojima et al., J. Phys. Soc. Jpn. 28 (1970) 1227.
[2] T. Kawai et al., Opt. Mater. 77 (2018) 30.
[3] I. Akimoto et al., Phys. Status Solidi C 6 (2009) 342.
[4] T. Kawai and A. Iguchi, J. Lumin. 207 (2019) 58.

66
Ⅲ02- 24 -Ⅲ -2. Materials Sciences
BL3B
BL3B

Optical Properties of Tl+-doped CsCaCl3 Single Crystals
K. Kubota and T. Kawai

Graduate School of Science, Osaka Prefecture University, Sakai 599-8531, Japan
Though optical properties of alkali halide crystals
doped with metal ions of ns2 electron configuration suchas Tl+, In+, and Ag− have been studied for a long time,
there have been few reports on the ns2 centers in othercrystals [1]. In this study, we focused on Tl+ centers in
ternary compound CsCaCl3. The CsCaCl3 crystals attract
attention as a host crystal for scintillator materials and theoptical properties of the CsCaCl3 crystals doped withvarious kind of impurity ions have been investigated [2].
However, the optical properties of the CsCaCl3 crystal
doped with ns2 ions are considerably less studied. Thus,
we have investigated the optical properties of the
CsCaCl3 crystals doped with the Tl+ ions.
CsCaCl3:Tl+ crystals were grown by the Bridgman
method from CsCl, CaCl2 and TlCl powders. The
optical measurements were performed at various
temperatures from 8 K to 300 K at the BL3B line of
UVSOR.
The absorption spectrum of a CsCaCl3:Tl+ crystal at
8 K is shown in Fig. 1. The A, B, and C absorption
bands due to the Tl+ ion were observed around 5.6, 6.7,
and 7.1 eV, respectively. These absorption bands are
attributed to the intra-ionic transitions from the ground
1
S0 state to the excited 3P1, 3P2, and 1P1 states in the Tl+
ion. The intensity ratio among the A, B, and C
absorption bands in CsCaCl3:Tl+ is similar to that of
the Tl+ centers in alkali halide crystals, while theseenergy positions are located at the higher energy
region than that of the Tl+ centers in alkali chloride
crystals. The difference of the energy positions mightcome from that of the lattice arrangement around theTl+ ion between CsCaCl3 and alkali halide crystals.
Figure 2 shows the temperature change of the Aabsorption band between 8 K and 300 K. As thetemperature rises, the bandshape of the A absorption
band broadens and the peak position shifts to thehigher energy side up to around 150 K. On the other
hand, above 150 K, the peak position shifts to the low
energy side. The change of the shift of the peak energy
would come from a slight change of the latticearrangement related with the tetragonal to cubic phasetransition in CsCaCl3.
Figure 3 shows the temperature dependence of thehalf-width of the A absorption band. The temperaturedependence of the half-width can be fitted by thefollowing theoretical formula considering the electron
lattice interaction,
where H(0) and hν are the half-width at 0 K and theeffective phonon energy, respectively. From the fitting
result, the values of 0.17 eV and 29.7 meV wereobtained as H(0) and hν, respectively.

Ⅲ-2
Fig. 1. Absorption spectrum of CsCaCl3:Tl+ at 8 K.

Fig. 2. Temperature change of the A absorption bandin the CsCaCl3 crystal at 8 K-300 K.

Fig. 3. Temperature dependence of the half-width of theA absorption band in the CsCaCl3 crystal.
[1] A. Ranfagni et al., Advance in Physics 32 (1983)
823.
[2] M. Tyagi et al., J. Appl. Phys. 113 (2013) 203504.

67
Ⅲ02- 25 -Ⅲ -2. Materials Sciences
BL3B
BL3B

UV-visible Spectroscopy of γ-Si3N4
L. Museur1, E. Feldbach2, A. Zerr3, M. Kitaura4 and A. Kanaev3

Laboratoire de Physique des Lasers, CNRS, Université Paris 13, 93430 Villetaneuse, France
2
Institute of Physics, University of Tartu, 1 W. Ostwald str., 50411, Tartu, Estonia
3
Laboratoire des Sciences des Procédés et des Matériaux, CNRS, Université Paris 13, 93430 Villetaneuse, France
4
Faculty of Science, Yamagata University, Yamagata 990-8560, Japan
related radiation-induced structural defects havedifferent origins. The excitonic origin may be
proposed for the near-IR PL, which has no allowed
sub-band gap transitions. The shifted visible PL band
apparently belongs to modified structural defects SiXSi.
This work has been carried out within the frameworkof the EUROfusion Consortium and French ResearchFederation for Fusion Studies and has received funding
from the Euratom research and training programme
2014–2018 under grant agreement No. 633053. The
views and opinions expressed herein do not necessarilyreflect those of the European Commission.
b
750

Intensity, a.u.

Ⅲ-2

Compounds having spinel structure have been
predicted to exhibit a very high tolerance to irradiation
with fast neutrons resulting in low swelling even athigh doses [1]. In this connection, the understandingof their electronic band structure and defects inducedby ionized radiation is of paramount importance forthe development of novel optical materials stable in
the environments of a fusion reactor. We performed thestudy of a polycrystalline -Si3N4 with spinel structure
by means of photoluminescence (PL) spectroscopy.
The sample of -Si3N4 was synthesized according to
the earlier reported procedure [2]. The starting powderof α-Si3N4 with oxygen content <2 wt.%, compacted
into a Pt capsule, was compressed to 14 GPa andheated to 1700 ºC for ~30 minutes. The sample
irradiation with He-ions of total dose 1017 cm-² was
performed in the Ion Beam Lab of Helsinki University.
The PL experiments were carried out at the BL3B
beamline of UVSOR facility with synchrotronradiation (SR) excitation in the UV-visible spectral
range. The samples were mounted on a LHe flow-type
cryostat, which permitted stabilizing the sample
temperature in the range between 300 and 8K. The
vacuum in the samples chamber was ~310-7 Pa. The
measured spectra were corrected for the detector
sensitivity and optical line transmission.
The PL and PL-excitation (PLE) spectra of thepristine and irradiated -Si3N4 are respectively shownin Fig. 1 and Fig. 2. According to Museur et al. [3], thedominant PL band in Fig. 1a excited with photons ofenergy above band gap Eg=5.05 eV belongs (inKröger-Vink notation) to X*SiXSi exciton bound toneutral defects. No significant PL of excitedmetastable paramagnetic nitrogen anion radical at 450
nm was observed. The absence of the self-trapped
exciton emission at 312 nm (4.0 eV) may be explainedby a structural quenching. After He+ irradiation andexcitation with photons of energy h≥Eg, The PL
strongly weakened and exhibited red shift indicating
structural modifications; simultaneously, a new nearIR PL band appeared at ~700 nm. In contrast, no
modification was observed for the PL spectra excitedwith h<Eg. The PLE spectra of the bound exciton (Fig.
2) evidenced red shift after the He+ irradiation.
Moreover, interband transitions above 5.05 eV (<250nm) did not appreciably contribute to PL and a new
band with the maximum at 530 nm was intensified
with the lower-energy excitation. Apparently, the

1x10

4

a

0
300

400

500

600

700

800

Wavelength, nm

Fig. 1. PL spectra of fresh (a) and irradiated (b) Si3N4 excited at 230 nm.
b

Wavelength, a.u.

1

a

100

150

200

250

300

350

400

450

Wavelength, nm

Fig. 2. Excitation spectra of PL at 500 nm of fresh(a) and irradiated (b) -Si3N4.
[1] K. E. Sickafus et al., Science 289 (2000) 748.
[2] N. Nishiyama et al., Sci. Rep. 7 (2017) 44755.
[3] L. Museur et al., Sci. Rep. 6 (2016) 18523.

68
Ⅲ02- 26 -Ⅲ -2. Materials Sciences
BL3B

BL3B

UV-visible Spectroscopy of MgAl2O4 Spinel Single Crystal
L. Museur1, E. Feldbach2, M. Kitaura3 and A. Kanaev4

Laboratoire de Physique des Lasers, CNRS, Université Paris 13, 93430 Villetaneuse, France
2
Institute of Physics, University of Tartu, 1 W. Ostwald str., 50411, Tartu, Estonia
3
Faculty of Science, Yamagata University, Yamagata 990-8560, Japan
4
Laboratoire des Sciences des Procédés et des Matériaux, CNRS, Université Paris 13, 93430 Villetaneuse, France

7.2 eV may be connected to antisite defects. The work
is in progress to interpret radiation-induced damage ofthe spinel structure.
This work has been carried out within the frameworkof the EUROfusion Consortium and French ResearchFederation for Fusion Studies and has received funding
from the Euratom research and training programme
2014–2018 under grant agreement No. 633053. The
views and opinions expressed herein do not necessarilyreflect those of the European Commission.
hexc, eV
9

8

7

6

5
PL=235 nm

Eg = 8.3 eV

PL intensity, a.u.

Spinel or MgAl2O4 is well known material, whichelectronic band and defect structures are not yet
understood. In the same time, this material has a strong
practical interest because of the predicted hightolerance to ionized radiation [1]. The knowledge of
its electronic and optical properties in connection withradiation-induced defects will help to develop novel
optical materials for extreme environments. In thisstudy, we performed the spectroscopic analysis of
MgAl2O4 single crystals by means of energy resolved
photoluminescence (PL) spectroscopy.
The experiments were carried out at the BL3B
beamline [2] of UVSOR facility with synchrotron
radiation (SR) excitation in the UV-visible spectral
range. The samples were mounted on the cold finger
of a LHe flow type cryostat, which permitted
stabilizing the sample temperature in the range
between 300 and 8K. The vacuum in the samples
chamber was ~310-7 Pa. In the complement to PL and
PL-excitation (PLE) spectra, transmission andreflection spectra were recorded. The measured
spectra were corrected on the detector sensitivity andoptical line transmission. The single crystal (S1) from
Alineason was (100) oriented and both sides polishedof size 5x5x0.5 mm3. Another crystal (S2) with (100)oriented surface, was available from the Tartu group.
The irradiation of the spinels with He+ ions of total
dose 1017 cm-² was performed in the Accelerator Lab
of Helsinki University.
The PL spectra of the pristine crystals S1 and S2 arequite similar. The most long-wavelength structured
near-IR emission at ~700 nm (1.8 eV) belongs to thewell-known R-lines of Cr3+ impurity. According toprevious interpretations, the strongest PL band at 5.4
eV belongs to the recombination of conduction-band
electrons with holes captured at the nearby oxygenions (e/h), and bands at 3.8 eV and 3.0 eV wererespectively assigned to F (two electrons trapped at anoxide-ion vacancy) and F+ (one electron trapped at anoxide-ion vacancy) centers (see in [3]). The PLE
spectra of different PL bands for pristine samples areshown in Fig. 1. All PL bands (except for that of Cr3+
impurity) show up upon the band gap excitation,
which energy Eg8.3 eV can be suggested. The He+
irradiation weakened PL and strongly modified PLE
spectra of all bands, except for that of Cr3+ impurity(Fig. 2). In particular, the interband transitions do notsignificantly contribute to these spectra and theintraband transitions shift to longer wavelengths. The
common excitation band of e/h, F and F+ emissions at

PL=320 nm

PL=420 nm

PL=520 nm

PL=680 nm

150

200

exc, nm

250

300

Fig. 1. PLE spectra of S1 crystal (T=8 K).
hexc, eV
8

7

6

5

PL = 240 nm

PL intensity, a.u.

1PL = 680 nm

PL = 330 nm

150PL = 520 nm

200

exc, nmPL = 440 nm

250

Fig. 2. PLE spectra of irradiated S2 crystal (T=8 K).
[1] K. E. Sickafus et al., Science 289 (2000) 748.
[2] M. Kitaura et al., J. Lum. 172 (2016) 243.
[3] A. Pille et al., Ceram. Int. 45 (2019) 8305.

69
Ⅲ02- 27 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL3B

BL3B

Construction of Vacuum Referred Binding Energy Diagram of YSiO2N
Doped with Lanthanoid Ions for Persistent Luminescence
Y. Kitagawa, J. Ueda and S. Tanabe

Graduate School of Human and Environment Studies, Kyoto University, Kyoto 606-8501, Japan

Ⅲ-2

Lanthanoid ions doped in inorganic materials form
localized energy levels in the host bandgap. The relative
positions of these levels play a very important role forthe carrier trapping process of persistent luminescence.
From the view point of developing novel persistentphosphors, it is necessary to investigate the electronic
structure of Ln3+/2+-doped compounds. Dorenbos
proposed the method to construct the vacuum referredbinding energy (VRBE) diagram of Ln3+/2+ incompounds based on spectroscopy [1,2]. In this work,
the VRBE diagram of YSiO2N, which has characteristic
Y3+ sites coordinated by six oxide and two nitride ions[3], was constructed. Since Ce3+ ions in oxynitrides
show 5d-4f luminescence with low energy because of
the stronger nephelauxetic effect by N3- than O2- [4],oxynitrides activated with Ce3+ have the potential to bea characteristic persistent phosphor with longer
luminescent wavelength. To obtain some parameters forconstruction of the VRBE diagram, photoluminescence
excitation (PLE) spectra of Ce3+- or Eu3+-doped
YSiO2N samples were measured at UVSOR BL3B.
Figure 1 shows the PLE spectra of Ce3+- or Eu3+doped YSiO2N samples. In both spectra, a small band
was observed at around 210 nm. This band is assigned
to the host exciton creation band. Considering thebinding energy between an electron and a hole, theoptical bandgap of YSiO2N was estimated to be 6.34 eV,
which is smaller than that of pure oxide materials withthe same Y/Si ratio. This is because for oxynitride
compounds, energy of the valence band top increases
due to N 2p orbitals. For the YSiO2N:Ce3+ sample,
several overlapped excitation bands, which are derived
from the 4f-5d transition of Ce3+, were observed ataround 260 nm (4f-5d4,5) and around 360 nm (4f-5d1,2,3).
In the case of the YSiO2N:Eu3+ sample, another broad
excitation band was observed in the range of 250-350
nm. This band is attributed to the charge transfer (CT)transition from the coordinating anions to Eu3+ ions.
Although for many Eu3+-activated oxides the CT bandis located at around 250 nm, the CT band of the Eu3+activated oxynitride sample has lower energy. Therefore,
this band is due to the CT transition from not O2- but N3to Eu3+.
Based on the measured spectroscopic data, the VRBE
diagram of YSiO2N was constructed and shown in Fig.
2. EC and EV mean the bottom of the conduction bandand the top of the valence band, respectively. The redand blue zigzag curves indicate the ground level ofdivalent and trivalent lanthanoid ions, respectively.
Lanthanoid ions with energy levels below EC have thepotential to be an electron trap. For YSiO2N:Ce3+, Sm3+

and Tm3+ is predicted to be the suitable co-dopant ionsfor persistent luminescence at room temperature
because they have appropriate activation energy, thedifference between EC and levels of divalent lanthanoidions.

Fig. 1. PLE spectra of YSiO2N samples doped with
Ln3+ (Ln = Ce or Eu). Both spectra were measured at
10 K.

Fig. 2. The constructed VRBE diagram of YSiO2N.
[1] P. Dorenbos, Phys. Rev. B 85 (2012) 1.
[2] P. Dorenbos, J. Solid State Sci. Technol. 2 (2012)
R3001.
[3] L. Ouyang et al., Phys. Rev. B 69 (2004) 094112.
[4] J. W. H. van Krevel et al., J. Alloys Compd. 268(1998) 272.

70
Ⅲ02- 28 -Ⅲ -2. Materials Sciences
BL4U
BL4U

STXM Study of Nanowire Cathode Materials for Li-Ion Batteries
D. Asakura1, 2, K. Akada3, Y. Harada2,3,4, H. Yuzawa5, T. Ohigashi5 and E. Hosono1,2

Research Institute for Energy Conservation, National Institute of Advanced Industrial Science and Technology(AIST), Tsukuba 305-8568, Japan
2
AIST-UTokyo Advanced Operando-Measurement Technology Open Innovation Laboratory (OPERANDO-OIL),
AIST, Kashiwa 277-8565, Japan
3
Institute for Solid State Physics, The University of Tokyo, Kashiwa 277-8581, Japan
4
Synchrotron Radiation Research Organization, The University of Tokyo, Sayo 679-5148, Japan
5
UVSOR Synchrotron Facility, Institute for Molecular Science, Okazaki 444-8585, Japan1

Li-ion battery (LIB) is a key energy-storage device torealize a low-carbon society. Particularly, for electricvehicles, LIBs with high energy density and high power
density are necessary. To develop such high-performance
LIBs, improving the performances of the cathodes anddesigning novel cathode materials are inevitable.
Nanowire materials are promising to enhance the rate
performance because of the short Li-diffusion path,
which should contribute to the power density of LIB.
For nanowire cathode materials, it is important toinvestigate how the Li diffusion proceeds in the
nanowire and to visualize the distribution of electronic
structure upon the charge/discharge. Thus, X-ray
microscopy should be powerful to clarify those issues.
In this study, we demonstrate ex situ scanningtransmission X-ray microscopy (STXM) for LiMn2O4(LMO) nanowire.
The LMO nanowire made by our group [1] wasmixed with carbon additive and a binder. The slurry was
dropped on Au(5 nm)/Ti(5 nm)/Si3N4(150 nm)
membrane window and dried. The thickness of thecoated film was thin enough to transmit soft X-rays.
This film as a working electrode was assembled as athree-electrode cell with Pt counter electrode, Ag/AgCl
reference electrode and 1 M LiNO3/H2O electrolyte
solution. Cyclic voltammetry experiments with a scan
speed of 0.5 mV/s were performed before the soft X-ray
measurements. The voltage range was set to 0.01 V (vs.
Ag/AgCl) and 1.0 V. Ex situ Mn L2,3-edge X-ray
absorption spectroscopy (XAS) and STXM were carriedout for the initial and charged LMO samples at BL4U of
UVSOR.
Figure 1(a) shows the Mn L2,3-edge XAS for the total
area of the charged LMO nanowire. The profile is almost
attributed to a Mn4+ state according to a multiplet
calculation. Figures 1(b-1) and 1(b-2) are the STXM
images for the nanowire measured with the photon
energies of A and B shown in Fig. 1(a). Figure 1(b-1) for
which the Mn3+ character is enhanced are not uniform
compared to Fig. 1(b-2). A difference between the opticaldensity image of Fig. 1(b-1) and that of Fig. 1(b-2) isdisplayed in Fig. 1(c), showing inhomogeneous
distribution of Mn3+ states remaining in the nanowire
after partial and insufficient oxidation, which is hardly
found by the XAS without imaging.
For the initial LMO, mixed Mn3+ and Mn4+ states
were confirmed (not shown here). Thus, the utility of

STXM to nanostructured electrode materials for LIBs isextremely high. In the near future, we will develop an
operando STXM measurement system to monitor
charge distribution dynamics during charge-discharge
reaction of LIBs.

Fig. 1. (a) Mn L2,3-edge transmission XAS for thewhole are of a nanowire. Charge-transfer multiplet
calculations for Mn3+ and Mn4+ are also shown [2]. (b1) and (b-2) STXM images for the photon energies A
and B in panel (a). (c) Difference between the opticaldensity of (b-2) and that of (b-1).
[1] E. Hosono et al., Nano Lett. 9 (2009) 1045.
[2] E. Stavitski and F. M. F. de Groot, Micron 41 (2010)
687.

71
Ⅲ02- 29 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL4U
BL4U

STXM Imaging of Water Surrounding Air Bubbles
2

Ⅲ-2

Y. Harada1,2, T. Ohdaira1, K. Akada1 and T. Ohigashi3

Institute for Solid State Physics, The University of Tokyo, Kashiwa 277-8581, Japan
Synchrotron Radiation Research Organization, The University of Tokyo, Sayo 679-5148, Japan
3
UVSOR Synchrotron Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
1

Hydrogen bond network of liquid water isspatiotemporally uniform on average, while the local
and instantaneous coordination of individual watermolecules is of course heterogeneous. Many discussions
have been made on the nature of the microscopic
heterogeneity. In the past 15 years we have studied
liquid water using synchrotron soft X-rays aiming at anessential understanding of this heterogeneity andobtained a lot of experimental results which supported
the existence of two characteristic moieties that
constitute liquid water [1, 2].
When the subject is extended to solid-liquid, liquidliquid and liquid-gas interfaces, we must consider aconnection between the bulk water having
heterogeneous hydrogen bond network and water at thesolid-liquid, liquid-liquid and liquid-gas interfaces
which we simply refer to as interfacial water [3]. In thecase of micro/nano bubbles liquid-gas interface isformed, and the presence of interfacial water
surrounding the micro/nano bubbles can be considered.
However, the thickness and hydrogen bond network of
the interfacial water has yet to be systematically
explored. Moreover, the electrostatic charge aroundmicro/nano bubbles and their stabilization mechanism
are still under debate [4]. In this work, we performed
chemical state imaging of water surrounding
micro/nano bubbles using O K-edge X-ray absorption
spectroscopy to elucidate the electronic state andnetwork structure of water molecules forming the
micro/nano bubbles by taking advantage of theexperience cultivated in the analysis of bulk water.
Water sample filled up with plus charged nanobubbles with the maximum ζ-potential of +23 mV, theaveraged diameter of 180 nm and the number density
around 1 x 108 particles/ml were prepared, all estimatedby using a ZetaView Particle Tracking Analyzer. A dropof the micro/nano bubble water was sandwiched in
between two 150 nm-thick SiN membranes. During thelimited beamtime we could find only micrometer sizedbubbles due to better possibility to find them than 100
nm order nano bubbles. Prior to the O K-edge imaging
we obtained N K-edge image to confirm that the particle
is indeed an air bubble that has stronger N K-edge
absorption inside than outside.
Figure 1 shows O K-edge STXM image of a micro
bubble obtained at (a) 532.04 eV and (b) 536.44 eV
excitations, respectively. At 532.04 eV, enhancement of
the signal inside the micro bubble occurred by theresonance to the π* state of an O2 gas (around 1 eV

above the typical absorption peak). At 536.44 eV
excitation which is close to the pre-edge resonance ofliquid water (around 1.5 eV above the typical absorptionpeak), we found surrounding water has specificchemical state in a micrometer region from the bubble
surface compared to the bulk region far from the bubble.
The origin of the appearance of different chemical stateof water is yet unknown. A possible explanation is thatwater in this region is under the electric dipole layer
around the bubble and influenced by the interaction withthe dipole moment of a water molecule, although the
thickness of the dipole layer could be the micrometersize only in a ultrapure water, which is not expected in
the present experimental protocol. An additional brightregion was found in Fig. 1(b) indicated by the blackarrow. This sphere was not observed in Fig. 1(a) and webelieve this is due to radiation induced damage of the
bubble interacting with both sides of the SiN
membranes. In the near future we need to trace the causeof the occurrence of the additional structure and realize
the imaging of real nanobubbles.

Fig. 1. O K-edge transmission XAS images of a
microbubble obtained (a) at 532.04 eV excitation(around 1 eV above the * state of typical O2 gas) and(b) at 536.62 eV excitation (around 1.5 eV above thepre-edge resonance of typical liquid water) .
[1] T. Tokushima et al., Chem. Phys. Lett. 460 (2008)
387.
[2] A. Nilsson and L. G. M. Pettersson, Nat. Commun.
6 (2015) 8998.
[3] K. Yamazoe et al., Langmuir 33 (2017) 3954.
[4] M. Alheshibri, J. Qian, M. Jehannin, and V. S. J.
Craig, Langmuir 32 (2016) 11086.

72
Ⅲ02- 30 -Ⅲ -2. Materials Sciences
BL4U
BL4U

STXM Studies on Vapor-assisted Quasi-2D Perovskite Solar Cells
H. W. Shiu1, M. S. Li2, L. C. Yu1, Y. L. Lai1, T. Ohigashi3, N. Kosugi3, P. Chen2 and Y. J. Hsu1
National Synchrotron Radiation Research Center, Hsinchu 300, Taiwan
2 Department of Photonics, National Cheng Kung University 701, Tainan, Taiwan
3
Institute for Molecular Science, Okazaki 444-8585, Japan
1

Organic hybrid perovskite is one of the most attractive
emerging photovoltaic devices due to its high opticalabsorption cofficient, long carrier diffusion lengths andlow manufacturing costs. Although the power conversion
efficiency (PCE) of 3D perovskite has been reported toreach a high record of 23 % for a small area device [1],how to improve the stability of organic hybrid perovskite
in high humid condition is an important issue. Recent
study indicated that a hybrid structured low dimensional
perovskites formed by replacing the cation with large
planar organic cation exhibits much more resistant tohumidity than the 3D perovskite devices.
In order to investigate the device performance andelectronic properties, we synthesized a bunch of
dimensional tunable quasi-2D perovskite by mixing
stoichiometric quantities of lead iodide (PbI2) and PEAI
(C8H9NH3) to yield a series of compounds with different
PEAI/PbI2 ratio on mp-TiO2/SiN substrate. With further
treatment of vapor-assisted MAI (CH3NH3I), a series ofcompact crystalline layer of (PEA)2(MA)n-1PbnI3n+1
were formed. As a results, a remarkable PCE of 19.1%
has been successfully achieved with PEAI/PbI2 = 0.05
(n = 40) which showed better performance than thepristine 3D perovskite of MAPbI3 (PCE =17.31%) [2].
However, at high PEAI/PbI2 ratio, for example
PEAI/PBI2 = 1 or 2, poor PCE of 5.85% and 0.02 %were observed. Besides, the SEM and KPFM results
suggest that the grain size and surface potential of PEAI
modified perovskite play an important role for thephotovoltaic performance. In order to understand thequsai-2D materials in a microscopic point of view, we
systematically studied the layer number dependent
behavior of (PEA)2(MA)n-1PbnI3n+1 perovskite byscanning transmission X-ray microscopy (STXM) and
Near-edge X-ray absorption fine structure (NEXAFS).
Two different substrates with low and regular mp-TiO2concentration were prepared for study the interplay
between PEAI modified-perovskite and mp-TiO2. Herein,
we mainly focus on the high PEAI/PbI2 ratio perovskite
on regular substrate to understand the correlation between
PEAI and the MAPbI3. Figure 1 shows the STXM carbon
map of (a) PEAI/PbI2 = 2 (n = 1) without MAI treatmentand (b) PEAI/PbI2 = 1 (n = 2) with MAI treatment. Only
weak contrasts were observed in Fig. 1 (a), however there
is no chemical difference between dark and bright regions
after detail examined by NEXAFS (not show here). The
results implied that for high PEAI/PbI2 ratio, the hybrid
perovskite performed morphologically and chemicallyuniform.
Figure 2 shows STXM carbon map and thecorresponding micro-spectroscopy of PEAI/PbI2 = 2 (n =1) extracted from 3D images after MAI treatment.

Compare with the result without MAI treatment (Fig.1(a)), clear bubble like structures with nearly 5 μm size
domain were observed. For pure (PEA)2PbI4 (n = 1), thecharacteristic carbon signal is C=C π* resonance at 285
eV, in contrast with pure MAPbI3 (n = ∞), thecharacteristic signal is C-H and C-N σ* resonance at
288.5 and 290 eV. As shown in Fig. 2 (b), the shape of red
and blue spectrum is similar, however the ratio of C=C π*
to C-N σ*resonance exhibits opposite behavior between
which indicated the phase separation behavior ispresented in PEAI/PbI2 mixtures. It is another evidence
for the formation of multiphase qusi-2D perovskite withn = 1, 2 ,3 and else.
(a)

(b)

Fig. 1. STXM carbon map of (a) PEAI/PbI2 = 2
without MAI vapor treatment and (b) PEAI/PbI2 = 1
with MAI vapor treatment on mp-TiO2/SiN substrate.
The mapping are obtained by the imaging subtraction
between 285 eV and the pre-edge.
(a)

(b)

(a)

Fig. 2. (a) STXM carbon map of PEAI/PbI2 = 2 with
MAI vapor treatment and (b) the corresponding microspectrum before (dark line) and after (red and blue line)
MAI treatment.
[1] Y. Rong, Y. Hu, A. Mei, H. Tan, M. I. Saidaminov,
S. Seok, M. D. McGehee, E. H. Sargent and H. Han,
Science 361 (2018) eaat8235.
[2] M.-H. Li, H.-H. Yeh, Y.-H. Chiang. U-S. Jeng, C.-J.
Su, H.-W. Shiu, Y.-J. Hsu, N. Kosugi, T. Ohigashi, Y.-A.
Chen, P.-S. Shen, P. Chen and T.-F. Guo, Adv. Mater. 30(2018) 1801401.

73
Ⅲ02- 31 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL4U

BL4U

In Depth Qualitative Structural Analysis of Nanocellulose Hybrid Materials
M.Y. Ismail1, M. Huttula2, M. Patanen2, H. Liimatainen1, T. Ohigashi3, H. Yuzawa3 and N. Kosugi3
Fibre and Particle engineering unit, University of Oulu, P.O. Box 3000, 90014 University of Oulu, Finland
2
Nano and Molecular Systems unit, University of Oulu, P.O. Box 3000, 90014 University of Oulu, Finland
3
UVSOR Synchrotron Facility, Institute for Molecular Science, Okazaki 444-8585, Japan1

Ⅲ-2

Hybrid materials are high strength and modulus
materials that consist of an organic-inorganic blend. We
have fabricated such kind of hybrid films using
cellulose nanofibrils (CNF) and nanosilica (5 - 30% wt)
embedded in a chitosan (Chi) biopolymer matrix using
a slow evaporation method, and characterized various
mechanical, chemical and structural aspects of thesefilms [1]. The combination of these constituents canresult in a sustainable packaging material with tailoredcharacteristics such as microbial resistance with thehelp of chitosan, high strength and stiffness andbiodegradability and with the aid of CNF and nanosilica.
The inorganic filler plays a major role in the mechanical,
thermal and UV absorption properties. The interaction
between the constitutes is a key factor in affecting the
hybrid’s properties.
Microscale structural properties of hybrids arecommonly investigated using field emission scanningelectron microscopy (FESEM) and transmission
electron microscopy. However, those techniques do notgive the chemical distribution of the structure.
Therefore, synchrotron radiation techniques are neededin order to complement the FESEM imaging. Scanning
Transmission soft X-ray Microscopy (STXM) was used
to investigate the spatial distribution of the chemicalcomponents of the studied hybrids, verifying theagglomerates of silica in the hybrid.
The STXM experiment was carried out at BL4U
beamline at UVSOR. The imaging was performed atenergies spanning over the C K-edge (280 and 320 eV)
and Si L2,3-edge (100 to 130 eV). The samples were“sandwiched” between resin blocks for ultramicrotome
sectioning, a reference spectrum of resin was collected
to ensure that the main material between the resin has
not been affected by the embedding. The reference Xray absorption spectra (XAS) at C K-edge of purecellulose nanofibers and chitosan were recorded andmatched with previous references [2].
Figure 1 shows examples of the obtained results for
10 wt% nanosilica hybrid film. Upper panels show
examples how the spectral information of STXM can beused to differentiate the silica-rich regions from the
CNF-chitosan film. The RGB images have been
constructed using a “stack fit” procedure of aXis2000
software [3]. In short, the XAS (shown in Figs. 1 b & d)of indicated regions (shown with circles in Figs. 1 a &c) are extracted and they are fitted to each pixel of theimage. The final color of a pixel depends on the weightof each RGB component spectra in that specific pixel.
In Fig. 1 b, the red XAS is a resin spectrum, extracted

from right lower corner of Fig. 1 a. The blue XAS
resembles both cellulose and chitosan spectra. Greenshows strong XAS with cellulose-type features. At Si
L2,3-edge (Figs. 1 c & d), the regions in red show atypical XAS for SiO2 [4], whereas blue and green
regions do not exhibit as strong resonance behavior. The
optical density images are shown in the lower panels(Figs. 1 e and f) and corresponding average C K-edge
XAS of the whole imaged region is shown in Fig. 1 g.
Figures 1 e and f represent the sum stacks below the C
K-edge (280-284 eV) and on the top of the C K-edge(288.2-300 eV). Figure 1 e shows strongly absorbing
white band and spots, which effectively disappear on
the top of the C K-edge resonance region (Fig. 1 f).
Instead, thin (width <150 nm) lines are clearlyhighlighted. This indicates that the white regions below
the C K-edge are SiO2 agglomerates, whereas thethinner stripes showing up at the C K-edge are bundlesof nanocellulose fibers aligned with the direction of
cutting or wrinkles in the CNF-chitosan film.

Fig. 1. Upper panels: RGB maps at a) C K-edge withb) corresponding XAS of indicated regions, and c) at
Si L2,3-edge with d) XAS of indicated regions. Lower
panels: Optical density images e) below and f) on the
top of the C K-edge, and g) the average XAS showing
the energy regions used to construct images e and f(grey shadowed areas). Scale bar 1 µm.
[1] M. Y. Ismail et al., submitted to Carbohydrate
Polymers.
[2] C. Karunakaran et al., PLOS ONE 10 (2015)
e0122959.
[3] http://unicorn. mcmaster.ca/aXis2000.html[4] D. Li et al., Am. Miner. 70 (1994) 622.

74
Ⅲ02- 32 -Ⅲ -2. Materials Sciences
BL4U

BL4U

Chemical and Spatial Identification for Gas-Dependent Nanobubbles
Sandwiched in Graphene Liquid Cells
C. H. Chuang1, W. H. Hsu2, I. S. Huang2, W. F. Pong1, T. Ohigashi3 and N. Kosugi3
Department of Physics, Tamkang University, Tamsui 251, Taiwan
2
Institude of Physics, Academia Sinica, Taipei 115, Taiwan
3
UVSOR Synchrotron Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
1

Since 2004, graphene retains the crown among themost studied two-dimensional materials for numerous
applications, not only due to the promising electric andthermal conductivity but also due to the true twodimensionality of single layer graphene. The idea of
replacing Si3N4 window by graphene in liquid cells isunder development, which will provide in-situ
conditions and real-time control [1]. X-ray source usedin scanning transmission X-ray microscopy (STXM)
can provide the X-ray absorption spectrum with theelement-specific determination and chemical bonding
information, even at the nanoscale resolution. Using the
combined ability of X-ray spectral feature in STXM
microscopes, in-situ detection is achieved in various
investigation [2]. Although the use of graphene sheet as
ultra-thin window has been proposed due to physicalenhancement, it is rarely seen in the utilization in STXM
owing to the limited knowledge in the solid-liquidinterface and electrochemical ability. In the energy
materials research, the catalytic activity on the externalsurface and active site is of prime importance which will
provide space-resolved information.
In this work, we observed structure formed by gas
molecules within the confinement of few-layers
graphene sheets by STXM experiment which was
carried out at the BL4U. The left inset in Fig. 1(a) shows
the C 1s mapping image of graphene by STXM whichrevel an elliptical dark-shade region at the center,corresponding to gas-formed structure surrounded bymulti-layer graphene. The transmission image througha variety of piece of graphene sheet demonstrates thevariation in mapping intensity. The absorption intensity
inside the elliptical shape at the center area is muchhigher than the signal of the graphene sheets, which isthe characteristic of the cross-section absorption of gas
region and was also identified as the gas-formed
structure area by transmission electron microscopy(TEM) experiments prior to this experiment. Individual
x-ray absorption spectrum (XAS) at the carbon nearedge corresponding to the selected area are shown in Fig.
1(a). The C1s spectra collected from outside of the gasformed structure exhibit a flat anisotropy of graphene
layer, while it is strongly based on the incident
polarization of X-ray beam and its directional dipoletransition selection rule. The inside area of ellipticalgas-formed structure is divided into three areas andfeatures owing to the spectral features around photon
energy 284 ~ 287 eV. It is indicative of the interface
bonding coordination and gas-filled contribution, andalso the relative ratios of π*/σ* state at 285.5 and 291.7

eV are one of the spectral evidence as for the position
dependence [3].
Figure 1(b) displays the N1s mapping image andrelative XAS. At the corresponding area as Fig. 1(a), thegraphene region outside the gas-formed structure has
the low edge jump at the N K-edge absorption due tothe unnoticed nitrogen content owing to the large-area
signal integration. Relatively, the intensity of N K-edge
absorption is verified for the higher absorption in thegas-formed region. The features at 398.9 and 400.5 eV
are mostly assigned to pyridinic/pyrrolic N-relatedgroups and –C ≡ N triplet state [3]. Both the featuresreflect the evidence of nitrogen bonding configuration
inside the structure, different from the outside surface
absorption. Thus, our STXM result for graphenematerials and its application show the interior phase ofgas as for the combination and absorption. It provides
valuable insights into the possible mechanism of gas
absorption at the graphene interface as well as theobservation of gas-condensed phase at solid-liquidinterface [4].

Fig. 1. (a) Individual XAS at carbon K-edge, selectedfrom the STXM mapping image (inset). (b) XAS at N
k-edge obtained from STXM mapping image (inset).
[1] S. Nemsak et al., J. Am. Chem. Soc. 139 (2017)
18138.
[2] T. Ohigashi et al., AIP Conference Proceedings 1741
(2016) 050002.
[3] C. H. Chuang et al., Scientific Report 7 (2017) 42235.
[4] C.-K. Fang et al., Scientific Reports 6 (2016) 24651.

75
Ⅲ02- 33 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL4U
BL4U

Effect of Defect on Perovskite Hydroxide for Oxygen Evolution Reaction
Y. R. Lu1, C. H. Chuang1, C. L. Dong1, W. F. Pong1, T. Ohigashi2, N. Kosugi2,
Y.-S. Liu3 and J. Guo3
Department of Physics, Tamkang University, Tamsui 251, Taiwan
UVSOR Synchrotron Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
3
Advanced Light Source, Lawrence Berkeley National Laboratory, Berkeley, CA 94720
1

2

Ⅲ-2

Perovskite-based materials have been largely
investigated for using as promising electrocatalysts foroxygen evolution reaction.
Control of oxygen
vacancies can modulate the electronic structure ofperovskite and is beneficial for enhancing OER
activity [1]. To further promote the oxygen evolution
reaction performance, metal cation vacancies areproposed in addition to oxygen vacancies [2]. However,
it is challenging to create the cation vacancies inperovskites owing to the fact that the formation energyof metal cation vacancies are too high. Herein, SnCoFe
perovskite hydroxide was prepared by a simple coprecipitation, and then treated with Ar plasma to obtain
the defective SnCoFe-Ar. An advanced SnCoFe
hydroxide for efficient oxygen evolution reaction isdeveloped recently [3]. The electrochemical properties
and the x-ray absorption spectroscopic studies haverevealed the correlation between the defective
structures and the catalytic activities. Yet, the detailsof these defective structures are not fully understood.
In this study, the SnCoFe perovskite hydroxide andthat treated with Ar have been investigated byscanning transmission X-ray microscopy (STXM)
which provides regional atomic/electronic structure
information.
Figures 1(a) and (b) present the Fe L-edge STXM
images and its stack mappings of selected region of
SnFeCo and SnFeCo-Ar perovskite hydroxides. The
stack mappings display red, yellow, and green areas,
corresponding to the different areas which areassociated with various thickness and the different
electronic structures are revealed for these two
samples. The spectral variations between SnFeCo and
SnFeCo-Ar suggest there are more defective site
formed after Ar treatment, and the chemical states of
Fe are quite different in the exterior and interior region
in the sample. Similarity, Figs 1(c) and (d) show theCo L-edge STXM. Interestingly, the Co is distributed
more homogeneously in SnFeCo-Ar than in the one
without Ar treatment. This suggest the Ar treatment
has greater effect on the Co site in SnFeCo.
Further O K-edge and Sn M-edge will be performed
in order to reveal the correlation between the OER
activity and the defective structures in these perovskitehydroxides.

Fig. 1. Fe L-edge ((a) and (b)) and Co L-edge ((c) and(d)) scanning transmission X-ray microscopy image
and its corresponding stack mappings of selected
region of SnFeCo and SnFeCo-Ar.
[1] J. A. Koza et al., Chem. Mater. 24 (2012) 3567.
[2] C. Yang et al., J. Phys. Chem. Lett. 8 (2017) 3466.
[3] D. Chen et al., Angew. Chem. Int. Ed. 57 (2018)
8691.

76
Ⅲ02- 34 -Ⅲ -2. Materials Sciences
BL4U
BL4U

Characterization of Two-phase Network Structure of Sulfur Cross-linked
Isoprene Rubber
K. Miyaji1, T. Nakajima1, Y. Sakaki1, J. Preeyanuch2 and Y. Ikeda3

Graduate School of Science and Technology, Kyoto Institute of Technology, Kyoto 606-8585, Japan
2
Research Strategy Promotion Center, Kyoto Institute of Technology, Kyoto 606-8585, Japan.
3
Faculty of Molecular Chemistry and Engineering, Kyoto Institute of Technology, Kyoto 606-8585, Japan1

The cross-linking reaction to form a threedimensional network structure in rubber is necessary togive a good rubber elasticity for rubber products.
Among the conventional cross-linking reactions, asulfur cross-linking (vulcanization) is the mostimportant cross-linking reaction which has been widelyused in rubber industries. However, the rubber network
formations are generally very complicate. The controlsof network structures are not easy because many crosslinking reactions proceed heterogeneously, dependingon various factors such as cross-linking reagents, fillers,
processing methods, and so on. Because the network
structures of rubber materials significantly influence
their mechanical properties, the evaluation of the rubber
network structures is an important for development ofrubber science and technology. In 2009, our researchgroup found the two-phase network structure of sulfur
cross-linked isoprene rubber composed of network
domain and mesh network using small-angle neutron
scattering. The network domains with high networkchain densities embedded in mesh network was
proposed. [1] This specific morphology was found to becontrolled by changing the amount of cross-linking
reagents. There are several analytical methods toinvestigate the morphology of rubber vulcanizates.
Scanning transmission X-ray microscopy (STXM) is
one of methods which provides chemical mapping in
nanometer scales. Also, both spatial and compositional
chemical information are provided at the same time. In
this study, STXM was used to confirm the two-phasenetwork morphology of isoprene rubber vulcanizate interms of the distribution of sulfur, and an existence of
network domain in the isoprene rubber vulcanizates was
confirmed.
The rubber compound was prepared byconventional mixing at room temperature on a two-rollmill with a water cooling system; isoprene rubber was
mixed with zinc oxide, N-(1,3-benzothiazol-2ylsulfanyl)cyclohexanamine (CBS), and sulfur.
Rubber vulcanizate was prepared by conventionalpressing method. This rubber vulcanizate was
extracted by tetrahydrofuran (THF) to remove the unreacted sulfur. The vulcanized rubber sheet after THF
extraction was cut by an ultra-microtome to prepare asmooth and flat surface with a thickness ofapproximately 100 nm for STXM measurement.
Figure 1 shows the STXM image of the vulcanizedisoprene rubber at 165.4 eV. The bright and dark regions

indicate the strong and weak absorptions by sulfur,
respectively. Thus, the brighter regions shown in Fig. 1
are possibly the clusters of sulfur in the vulcanizedisoprene rubber. As predicted by the SANS results, thenetwork domains were formed by the absorption of
sulfur and the accelerator, CBS on the ZnO clusters, andfollowed by sulfur cross-linking on the surface of ZnOclusters [1].
This STXM image successfully confirmed theinhomogeneous two-phase network morphology in thevulcanized isoprene rubber from the view point ofsulfur distribution. The localized sulfur cluster is one of
the evidences for the network domain formation.

Fig. 1. STXM image of the vulcanized isoprene rubber
at 164.5eV.
[1] Y. Ikeda et. al., Macromolecules 42 (2009) 2741.

77
Ⅲ02- 35 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL4B
BL4B

Chemical State Investigation of Nitrogen in Nitrided Steels by Soft X-ray
Absorption Spectroscopy
M. Sato and S-W. Young

Institute for Materials Reseach, Tohoku University, Sendai 980-8522 Japan

Ⅲ-2

Machinery structural materials such as gears, shafts
and molds are required to have high surface hardness
and high toughness because of the severe friction
and/or abrasion at the surface.
Nitriding is a surface hardening technique widelyused industrially and flows nitrogen from the surface
and combining with strong nitride forming elements(for example, Al, Cr, V, Ti, etc.) contained in steel toprecipitate clusters and alloy nitrides at near thesurface. Compared to other surface hardening methods
such as carburizing and shot peening, the nitriding has
advantages of less strain because it does not involve
phase transformation and deformation during
treatment [1, 2]. Therefore, the importance of nitriding
process is increasing in recent high-precision materials.
In the nitriding treatment, a sufficient amount of
nitrogen flows in the vicinity of the surface toprecipitate the alloy nitride at the initial stage of
nitriding, whereas precipitation is delayed at the insideof the sample because it takes long time for thenitrogen diffusion. As a result, the concentration
gradient of nitrogen is formed in a sample.
In this research, we aimed to investigate whether
such gradient change of nitrogen can be captured by
XAFS method. Since this measurement is our firstbeam time at BL4B, it is also an important object toacquire how to use the devices and to obtain the correctdata.
A sample with a nitrogen concentration gradientfrom the surface to the inside was prepared by plasma
nitriding treatment at 550 °C for 16 hours using the Fe0.5V-1.5Al alloy. In addition, nitrogen martensite steel
including 0.3, 0.5 and 2 mass% N prepared by nitriding
and quenching (N-Q) treatment using the pure iron,and the γ’-Fe4N and the BN powder were used as
reference materials.
XAFS measurement was carried out at the beam line
BL4B, and the N-K edge (380 ~ 450 eV) XANES
spectra were collected by the total electron yield
method, respectively. The measurement for the low
concentration sample was performed a plurality oftimes, and the obtained spectrum was merged.
Figure 1 shows the N-K edge XANES spectra ofnitrided Fe-0.5V-1.5Al samples at different depth. The
shape and intensity of the obtained spectrum weredifferent at the distance from the surface of the sample.
In the case of nitriding, the reaction proceeds bydiffusion of nitrogen from the surface to the inside of
the sample, so the amount of nitrogen becomes thinneras it becomes inside of the sample. Obtained spectra
showed such a concentration gradient of nitrogen well.

In addition, since the spectral shape is different
between the sample surface and the inside, it isunderstood that the chemical state of nitrogen isdifferent inside the sample. At the inside of the sample,
VN clusters are precipitated mainly. On the other hand,at the sample surface, AlN and Fe4N or Fe16N2 are
precipitated in addition to VN clusters, andsupersaturated nitrogen also exists. In the future,detailed analysis of the obtained spectrum will be
performed by combining simulations.

Fig. 1. N K-edge XANES spectra of nitride Fe-0.5V1.5Al samples obtained at different depth. (a) 50 μm(b) 200 μm (c) 400 μm and (d) 600 μm, respectively.
[1] Tekkou Zairyou, The Japan Institute of Metals andMaterials, 1985.
[2] K. Kawasaki, M. Okumiya and S. Uchida,
NETSUSHORI 50 (2010) 305.

78
Ⅲ02- 36 -Ⅲ -2. Materials Sciences
BL4B
BL4B

Charge State Analysis of Fe Ions in Cr-doped Y3Fe5O12
K. Shimazu1 and T. Yamamoto1,2,3

Faculty of Science and Engineering, Waseda University, Tokyo 169-8555, Japan
Insitute of Condensed-Matter Science, Waseda University, Tokyo 169-8555, Japan
3
Kagami Memorial Research Institute for Materials Science and Technology, Waseda University, Tokyo 169-0051,
Japan
1

2

Thermochromism is a phenomenon to change thecolor due to a change in temperature. There are wide
variety of organic materials, which show
thermochromism. However, few inorganic ones showsuch thermochromism. Among such inorganic
thermochromic materials, rare-earth (RE) iron garnet,
RE3Fe5O12, shows such color change [1]. Although it
was reported that red-shift occurs by Cr-doping in oneof the rare-earth iron garnet, Y3Fe5O12 [2], themechanism in such red-shift has not yet been clearly
understood. In the current study, we focused upon thechange in electronic structure of the Cr-doped Y3Fe5O12
garnet. Especially, change in charge state of Fe ions as afunction of doped Cr concentration was investigated by
the Fe-L3 X-ray absorption spectra (XAS)
measurements in this report.
Cr-doped Y3Fe5O12 garnet was synthesized withconventional solid state reaction method changing theconcentration of doped Cr ions. Crystal structures of thesynthesized powders were examined by the powder Xray diffraction, which showed all the samples are single
phased and no secondary phase could be observed.
From the observed XRD patterns, we obtained thevolumes of cells as a function of Cr concentration,
which are shown in Fig. 1. Volumes decrease as increaseof Cr concentration, which suggests Cr ions substitute at
Fe site as trivalent ions. These results can be explainedby the difference in ionic radii, i.e., ionic radius of Cr3+
ion is smaller than that of Fe3+ ion [3].
Fe-L3 XAS spectra were observed at BL4B, UVSOR,
in a total electron mode by collecting the drain current.
Sample powders were settled on the carbon adhesivetape, which was put on the Cu plate. Incident beams
were irradiated perpendicular onto the plate, which was
monochromatized with a varied-space grating (800l/mm). Observed Fe-L3 XANES spectra are shown in
Fig. 2. All the spectral profile of Fe-L3 XAS of Cr-doped
Y3Fe5O12 garnet is independent from the concentrationof doped Cr and shows similar profile as that of Fe3O4.
From these results we can safely conclude that valence
state of Fe ions do not change in all samples withdifferent Cr concentration from the mixture of 2+ and3+ as in Fe3O4. The red-shift is under investigation by
the calculated electronic density of states within adensity functional theory.

Ⅲ-2

Fig. 1. Relative volume (V/V0) of Cr-doped Y3Fe5O12(Y3Fe5-xCrxO12) garnet as a function of Cr concentration.

Fig. 2. Observed Fe-L3 XANES spectra of Cr-doped
Y3Fe5O12 (Y3Fe5-xCrxO12) with those of the standard
iron oxides.
[1] S. C. Parida et al., J. Phase Equil. 24 (2012) 5.
[2] H. Lu et al., J. Mater. Chem. C44 (2016) 15981.
[3] R. D. Shanon, Acta Cryst. A32 (1976) 751.

79
Ⅲ02- 37 -Ⅲ -2. Materials Sciences
BL4B

BL4B

Applications of Soft X-ray Spectroscopy to the Analysis of Nitrogen
Implanted TiO2 Photocatalysts
T. Yoshida1, M. Yamamoto1, Y. Kato2, A. Ozawa2 and T. Tanabe1

Advanced Research Institute for Natural Science, Osaka City University, Osaka 558-8585, Japan
2
Graduate School of Engineering, Osaka City University, Osaka 558-8585, Japan

1

(a)

(b)

4

10

around 400 eV was observed for the photocatalytic
active sample while the XANES spectrum of thephotocatalytic inactive sample showed a distinct single
peak around 401 eV. These results suggested two typesof chemical states of nitrogen, i.e., the photocatalytic
active N substituting the O sites and the inactive NO2
species. In the valence band photoelectron spectrum of
the photocatalytic active sample, the additional
electronic states were observed just above the valenceband edge of a TiO2. The electronic state would beoriginated from the substituting nitrogen and be
responsible for the band gap narrowing, i.e., visible
light response of TiO2 photocatalysts. On the other
hand, the resonant photoemission spectra of the active
sample under X-ray excitation around N K-edge were
almost the same regardless of X-ray energy, while
those of the inactive sample showed an additional
sharp peak in valence band under the X-ray excitation
around 401eV.
[1] A. Fujishima, X. Zhang and D. A. Tryk, Surf. Sci.
Rep. 63 (2008) 515.
[2] R. Asahi, T. Morikawa, T. Ohwaki, K. Aoki and Y.
Taga, Science 293 (2001) 269.
[3] R. Asahi and T. Morikawa, Chem. Phys. 339 (2007) 57.

(c)
5

unirrad.

8

3

4

2
17

5 x 10 /cm
1

6

unirrad.

2

4

1 x 10 /cm

17

2

0

17

2

5 x 10 /cm

2

540
550
560
X-ray energy (eV)

570

2

17

2

420

425

3
2

5 x 10 /cm

1
0

0
530

17

1 x 10 /cm
I / I01 x 10 /cm

2

I/ I0

17

I / I0

Ⅲ-2

Photocatalytic reactions at the surface of titanium
dioxide (TiO2) under UV light irradiation have been
attracting much attention in view of their practicalapplications to environmental cleaning such as self
cleaning of tiles, glasses, and windows.[1] It has beenreported that the doping of nitrogen into TiO2
contributes to band gap narrowing to provide visiblelight response.[2, 3] In the present study, nitrogen doped
TiO2 photocatalysts were prepared by nitrogen
implantation method, and investigated the chemical
states of the surface of the samples mainly by soft X-ray
absorption and photoelectron spectroscopies.
Nitrogen ions were implanted into TiO2 at 5 or 50 keV
at room temperature. Photocatalytic activity of the N+implanted sample was evaluated by the decomposition
reaction of methylene-blue solution under visible-lightirradiation. The photocatalytic activity changed with theamount of implanted nitrogen, and the photocatalytic
active and inactive N+-implanted samples were obtained.
As shown in Fig. 1, O K-edge and Ti L2,3-edge of
XANES analysis exhibited the formation of Ti3+
species in both photocatalytic active and inactive
samples (A-cat and I-cat). The Ti3+ species would beformed by the displacement effect of oxygen
atomsduring the N+ implantation. On the other hand,in N K-edge XANES spectra, a double-peak structure

455

460
465
470
X-ray energy (eV)

475

390

395

400 405 410 415
X-ray energy (eV)Fig. 1. (a) O K-edge XANES spectra of active sample (A-cat), inactive sample (I-cat) and a bare sample.
(b) Ti L2,3-edge XANES spectra of active sample (A-cat), inactive sample (I-cat) and a bare sample.
(c) N K-edge XANES spectra of active sample (A-cat) and inactive sample (I-cat).
.

80
Ⅲ02- 38 -Ⅲ -2. Materials Sciences
BL4B

BL4B

N K-edge XAS Spectroscopy on 3d Transition Metal Doped AlN Films
N. Tatemizo and S. Imada

Faculty of Electrical Engineering and Electronics, Kyoto Institute of Technology, Kyoto, 606-8585, JapanFor the last two decades, wide-band gap
semiconductors heavily doped with 3d transition
metals (3d-TM) have been extensively investigated inorder to realize devices with new functionality [1, 2].
To confer photoconversion of sun light as anintermediate band (IB) material [1], we have been
exploring the possibility of heavy 3d-TM doping to
AlN, which is in the family of III-nitrides [3-5]. For IB
material, Fe is thought not to be effective because an
Fe ion in an Al site of AlN have 5 3d electrons, that is,half-filled electronic structure. However, from theview point of chemical trend study of 3d-TM in AlN,
Fe is an interesting dopant; the additional band mainly
consisting of 3d states of Fe are formed very near tothe valence band (VB) of AlN according to ourtheoretical band structure calculation, while those ofTi, V and Cr are formed in the middle of the band gap[6]. To investigate the band structures of the Fe-doped
AlN (AlFeN), we synthesized AlFeN with various Feconcentration between 0-30% [7] and carried out N Kedge XAS measurements.
The films were deposited on SiO2 glass substratesby RF sputtering from an AlN target with Fe metal
chips on it. Results of XRD and TEM showed that thefilms have wurtzite structures [7]. N K-edge XAS
measurements for the films in fluorescence-detection
mode at room temperature. The electric field vectorsof the X-rays were parallel to the film plane.
Figure 1 shows the N K-edge XAS spectra of AlFeN(7.8%) and Ti-doped AlN (AlTiN, Ti: 8.3%) films [3].
The energies of the 3 main peaks indicated blackallows A, B, and C for the both films coincided withthose of un-doped AlN [3]. Regarding the relative
intensities of the peaks, there is a difference at the peak
C. This is thought to be due to the deference in theorientation axis of the films; the AlTiN film has a caxis orientation [3], while the AlFeN film has an a-axis
orientation [7]. In both spectra, pre-edge peaks wereclearly observed. Reflecting the large number of unoccupied d-states (9 or 10 states [3]) hybridized with
N p-states just below the bottom of the conduction
band, AlTiN film had a large pre-edge peak near themain peak. On the contrary, that of AlFeN was small
and far from the main peak A. According to the resultof Fe K-edge XANES analyses (not showing here), Fe
ions have an oxidation state of 3+ (3d5). It isreasonable that the spectrum of AlFeN had smaller
pre-edge peak than that of AlTiN. As intensities of preedge peaks reflect not only un-occupation of 3d states
but also symmetry around N, father study on thesymmetries around N (Fe/Ti) should be nessesary.

The difference of the peak energies of the pre-edge
peak and the main peak A is about 5 eV in the spectrumof AlFeN. This indicates that the un-occupied states
were formed near the VB of AlN. This is consistent
with our calculation of the band structure of AlFeN.

Ⅲ-2

Fig. 1. N K-edge XAS of AlTiN (Ti: 8.3%) and
AlFeN (Fe: 7.8%) films deposited on SiO2 glass
substrates.
[1] A. Luque et al., Phys. Rev. Lett. 78 (1997) 5014.
[2] T. Dietl et al., Science 287 (2000) 1019.
[3] N. Tatemizo et al., J. Mater. Chem. A 5 (2017)
20824.
[4] N. Tatemizo et al., J. Appl. Phys. 123 (2018)
161546.
[5] N. Tatemizo et al., J. Phys.: Condens Matter 29(2017) 085502.
[6] N. Tatemizo et al., Materials Science Forum 924(2018) 322.
[7] N. Tatemizo et al., AIP advances 8 (2018) 115117.

81
Ⅲ02- 39 -Ⅲ -2. Materials Sciences
BL4B
BL4B

Local Structure Investigation of Alloy Nitrides Precipitated by Interface
Precipitation during Ferrite Transformation
1

Ⅲ-2

M. Sato and T. Murata

Institute for Materials Reseach, Tohoku University, Sendai 980-8522, Japan

The pro-eutectoid ferrite formed from austenite is themost frequently appearing morphology in steel. In
general, it is known that grain boundary allotriomorph
ferrite are formed when the degree of supercooling fromthe Ae3 line is low, and the widmanstatten ferrite isformed as the degree of supercooling increases [1, 2]. In
addition, in carbon steel, when the strong carbideforming elements such as Ti, Nb, V, Cr and Mo areadded, the interface precipitation occurs and alloy
carbides are periodically generated on the γ / α interface
during ferrite transformation.
Although there are many reports on the ferrite
transformation behavior in carbon steel, there are very
few reports on nitrogen steel. Under such circumstances,
we have found that the interface precipitation of CrN
occurs in the ferrite transformation of Fe-1Cr-0.3N
austenite.
In this study, the change in chemical states of Cr andN during the ferrite transformation with interface
precipitation in Fe-1Cr-0.3N austenite.
Fe-1mass%Cr alloy were used as starting materials.
They were homogenized at 1523 K for 24h and furnacecooled until room temperature. The Fe-1Cr-0.3N alloy
were prepared by nitriding and quenching (N-Q)
process using NH3/H2 mixed gas at 1273 K for 1 h. The
obtained alloys were isothermally treated at 973 K forseveral time of periods and then quenched into iced
brine. The N K-edge and Cr L-edge XANES spectra
were corrected by fluorescence method using silicon
drift detector (SDD) at BL4B in UVSOR, respectively.
Obtained data were analyzed using Athena software.
Calculation of XANES spectrum was performed by
FDMNES software and the crystal structural data wascorrected from NIMS web site [3].
Figure 1 shows Cr L-edge XANES spectra of samples.
In the L3 edge peak, intensity around 588 eV increased
with increasing holding time due to the increasing thevolume fraction of CrN precipitate.
Figure 2 shows N K-edge XANES spectrum ofsample isothermally treated at 750 °C for 10 min. andcalculated XANES spectrum of Fe4N and CrN. It isthought that the obtained spectra are consisted from the
combination of Fe4N and CrN. Further detailed analysis
will be performed using this kind of calculations.

Fig. 1. Cr L-edge XANES spectra of samples (a) As
quenched, (b) heat treated at 750 °C for 10 min. and(c) heat treated at 750 °C for 1 h.

Fig. 2. Calculated and measured N K-edge XANES
spectra. (a) Fe4N (calc.), (b) CrN (calc.) and (c) heat
treated at 750 °C for 10 min.
[1] Dube’ C. A., H. I. Aaronson and R. F. Mehl,
Re’v.met. 55 (1958) 201.
[2] H. B. Aaron and H. I. Aaronson: Metall. Trans. 2(1971) 23.
[3] https://crystdb.nims.go.jp/

82Ⅲ02- 40 -Ⅲ -2. Materials Sciences
BL5U

BL5U

ARPES Study of Iron-chalcogenide High-Tc Superconductors
K. Nakayama1, G. N. Phan1, H. Oinuma1, D. Takane1, K. Sugawara1,2, T. Takahashi1,2,3,
F. Nabeshima4, T. Ishikawa4, T. Shikama4, A. Maeda4 and T. Sato1,3

Department of Physics, Tohoku University, Sendai 980-8578, Japan
Center for Spintronics Research Network, Tohoku University, Sendai 980-8577, Japan3
WPI Research Center, Advanced Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan
4
Department of Basic Science, the University of Tokyo, Tokyo 153-8902, Japan1

2

Iron chalcogenides FeCh (Ch: S, Se, Te) have long
been attracting considerable attention as the
structurally simplest iron-based superconductors.
These materials show rich physical properties such as
superconductivity, antiferromagnetism, structuraltransition, and electronic nematicity. Intriguingly,
these physical properties exhibit gigantic pressure
effects, as exemplified by a marked increase in asuperconducting transition temperature (Tc) from 8 K
to 37 K in pressurized FeSe [1]. Understanding theorigin of such pressure-induced Tc enhancement
would provide a key to resolving the high-Tc
mechanism in iron-based superconductors.
Recent successful fabrication of FeSe1-xTex thin films
allows us to control the lattice parameters by tuning thelattice mismatch with substrate. Such strain engineering
has great capability for manipulating various physicalproperties. For example, the application of compressive
strain enhances the Tc value up to 12 K in FeSe on CaF2
[2, 3] and further up to 23 K in FeSe0.8Te0.2 on CaF2 [4].
In addition, strain-induced drastic changes in the
structural, magnetic, and nematic properties have been
discovered in FeSe1-xTex [5]. These findings haveprovided a rare opportunity for investigating theinterplay between the physical properties and the lattice
parameters at ambient pressure.
In this study, we have performed high-resolution
angle-resolved photoemission spectroscopy (ARPES)of compressive-strained and strain-free FeSe1-xTex thinfilms at BL5U, and determined their band structure in
the three-dimensional momentum space. Our
systematic ARPES measurements revealed the straininduced evolution of Fermi surface and electronic
nematicity as well as their relationship with Tc
variation.
High-quality compressive-strained FeSe1-xTex thinfilms with the thicknesses of approximately 300
monolayers were grown on CaF2 substrate by pulsed
laser deposition. Figure 1 shows the representative
ARPES results in the normal state of FeSe0.8Te0.2 (Tc =23 K) measured with circularly-polarized 40-eV
photons. We observed several dispersive bands near
the Fermi level (EF), which mainly originate from Fe
3d orbitals. Fermi surfaces consist of hole-like andelectron-like small pockets at the Brillouin-zone
center (Γ point) and corner (M point), respectively (seeFig. 1). Similar size of these Fermi surfaces supports
the semimetallic nature of FeSe0.8Te0.2. We have also
performed ARPES measurements on several different

samples such as strain free and different compositions,and found strain- and Te-substitution-induced changes
in the semimetallic band overlap and the nematicity.
Our observations provide important insights into theinterplay between Tc enhancement and lattice
parameters.

Fig. 1. (a) and (b) ARPES intensity plot as a functionof binding energy and wave vector measured aroundthe Brillouin-zone center and corner, respectively, on
FeSe0.8Te0.2 (Tc = 23 K) at T = 30 K. The data wereobtained with circularly-polarized 40-eV photons.
[1] S. Medbedev et al., Nature Mater. 8 (2009) 630.
[2] F. Nabeshima et al., Appl. Phys. Lett. 103 (2013)
172602.
[3] G. N. Phan et al., Phys. Rev. B 95 (2017) 224507.
[4] Y. Imai et al., Proc. Natl. Acad. Sci. USA 112 (2015)
1937.
[5] Y. Imai et al., Sci Rep. 7 (2017) 46653.

83
Ⅲ02- 41 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL5U

BL5U

Observation of Fermi-arc Surface States in CoSi
D. Takane1, Z. Wang2, S. Souma3,4, K. Nakayama1, T. Nakamura1, H. Oinuma1, K. Hori1,
K. Sugawara1,3, T. Takahashi1,3,4, Y. Ando2 and T. Sato1,3

Department of Physics, Tohoku University, Sendai 980-8578, Japan2
Physics Institute II, University of Cologne, 50937 Köln, Germany
3
Center for Spintronics Research Network, Tohoku University, Sendai 980-8577, Japan
4
WPI Research Center, Advanced Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan1

Ⅲ-2

The search for new fermionic particles, which has
long been an important research target in elementaryparticle physics, is now becoming an exciting
challenge in condensed-matter physics to realize
exotic fermions, as highlighted by the discovery of
two-dimensional (2D) Dirac fermions in graphene andhelical Dirac fermions at the surface of threedimensional (3D) topological insulators. The recent
discovery of 3D Dirac semimetals and Weyl
semimetals hosting massless Dirac-Weyl fermionsfurther enriches the category of exotic fermions. While
all these fermions could manifest themselves inelementary-particle physics, recent theories havepredicted new types of massless fermions incondensed-matter systems that have no counterparts inelementary particles. These fermions appear in the
crystals with specific space groups, such as CoSi, as
exemplified by the spin-1 chiral fermion and the
double Weyl fermion showing multifold banddegeneracies at the band-crossing point (node)
protected by the crystal symmetry [1-3] [Fig. 1(a)].
While a well-known spin-1/2 Weyl fermion shows a
Weyl-cone energy band dispersion with twofold
degeneracy at the node that carries the topologicalcharge (Chern number) C of ±1, these novel fermionscarry topological charge of ±2, and hence produce
multiple (double) Fermi-arc surface states.
In this study, we performed angle-resolved
photoemission spectroscopy (ARPES) of CoSi at
BL5U. By utilizing energy-tunable photons fromsynchrotron radiation, we found first experimental
evidence for unconventional Fermi-arc surface statesthat strongly support the presence of spin-1 chiralfermion and double Weyl fermion in the bulk [4].
High-quality single crystals of CoSi were grown bya chemical vapor transport method. To search for thepredicted Fermi-arc surface states, we selected hν = 67
eV, which corresponds to kz ~ 0.75π at kx ~ 0 to avoid
a possible complication from the bulk bands. As
shown in the upper left panel of Fig. 1(b), the ARPES
intensity at the Fermi level(EF) is elongated toward＿＿
two of four adjacent M points, resulting in a C2symmetric “Z”-shaped intensity pattern. This feature
is commonly observed at other hν’s [Fig. 1(b)] despitea sizable change in the kz value [0.3- 0.95π]. We foundthat the energy dispersion of the EF-crossing band that
produces the Z-shaped Fermi contour seems to beunchanged against a variation of hν These results

strongly suggest the surface origin of the Fermi
contour in Fig. 1(b). This Fermi contour is attributed
to the predicted Fermi-arc surface states, since thecalculated surface states are
and connect＿＿C2 symmetric＿＿with the bulk FSs at the Γ and M points, consistent
with the present ARPES results. The present study
strongly supports the existence of chiral topologicalfermions beyond Dirac and Weyl fermions incondensed-matter systems.

Fig. 1. (a) Spin-1 chiral fermion (left), and double
Weyl fermion (right). (b) Fermi-surface mapping of
CoSi at hν= 67 eV which indicates the existence ofFermi-arc surface states.
[1] B. Bradlyn et al., Science 353 (2016) aaf5037.
[2] G. Chang et al., Phys. Rev. Lett. 119 (2017) 206401.
[3] P. Tang et al., Phys. Rev. Lett. 119 (2017) 206402.
[4] D. Takane et al., Phys. Rev. Lett. 122 (2019) 076402.

84
Ⅲ02- 42 -Ⅲ -2. Materials Sciences
BL5U
BL5U

Investigating Electronic Structures of the Strained Transition Metal
Dichalcogenide Monolayers
2

D. J. Oh1,2, B. Kim1,2, K. Tanaka3, S. R. Park4 and C. Kim1,2

Department of Physics and Astronomy, Seoul National University (SNU), Seou 08826l,Republic of Korea
Center for Correlated Electron Systems, Institute for Basic Science (IBS), Seoul 08826, Republic of Korea
3
UVSOR Synchrotron Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
4
Department of physics, Incheon National University, Incheon 406-772, Korea

1

Transition metal dichalcogenides (TMDs) monolayer(ML) such as MoS2 and NbSe2 have recently drawn alot of attentions because of their novel properties
distinguished from their bulk formed materials [1, 2]. It
is desired to manipulate the electronic structures andmaterials properties of TMDs ML by applying external
field or strain. Uniaxial strained MoS2 ML was recently
realized by a special method using a flexible substrateand exhibit a interesting property [3]. The electronic
structure of monolayer MoS2 can be modified byuniaxial strain. When the current was applied to thesample, strained monolayer MoS2 can has valley
polarization. In addition, the direction of the valley
polarization can be controlled by the direction of thecurrent [3].
We bought the MoSe2 single crystal from the HQgraphene. We exfoliated the thick MoSe2 flakes fromsingle crystal on the viscoelastic substrate [4]. Then weapplied the uniaxial strain using a mechanicaltranslation stage [3]. Finally, we transferred a MoSe2
flake onto the degeneratly doped Si substrate. Samples
were annealed at 700K for 2h in UHV to cleaning thesurface before ARPES measurements. ARPES spectra
were acquired by hemispherical photoelectron analyzer(MBS A-1) equipped with deflectors that enable
momentum space mapping. The base pressure is better
than 5 × 10-11 torr. We used linearly polarized light with
80eV photon energy.
We have tried to measure the electronic structure ofstrained MoSe2 by angle resolved photoemission in
UVSOR BL5U. Figure 1(a) is the constant energy cut
obtained at -2.4 eV binding energy. There are two hole
pockets at Γ and K-point. In Fig. 1(b), there are band
splitting induced by spin-orbit coupling and it wascoincident with previous work [5]. Due to the small
spin-orbit coupling in MoSe2, it is hard to measure thesize of splitting. And we couldn’t measure any modified
electronic structure induced by uniaxial strain. We
assume that the applied uniaxial strain is too small tomodify the electronic structure. Or the applied uniaxial
strain was released during the annealing process.
Furthermore, we have check the charging effect. We
measure the energy distribution curve (EDC) with two
different entrance slit size. In Fig. 2, measured EDC isshifted to higher binding energy when we use larger
entrance slit. It is general tendency when there is acharging effect. It seems to be came from poor electricalcontact between Si substrate and MoSe2 flake or thepresence of SiO2 layer on the Si substrate.
Based on these results, we were able to identify the
problems to be solved. We will try to make better
electrical contact or use another conducting substrate.
And we will try to apply uniaxial strain using anothermethod which can maintain the uniaxial strain during
ARPES measurement.

Fig. 1. (a) Constant energy cut obtained at the -2.4 eVbinding energy (b) Measured dispersion along the reddashed line in Fig. 1 (a).

Fig. 2. Measured energy distributrion curves using two
different slit size.
[1] D. Xiao et al., Phys. Rev. Lett 108 (2012) 0196820.
[2] M. Ugeda et al., Nature Physics 12 (2016) 92.
[3] J. Lee et al., Nat. Materials 16 (2017) 887.
[4] A. Castellanos-Gomez et al., 2D Mater. 1 (2014)
011002.
[5] B. S. Kim et al., Sci. Rep. 6 (2016) 36389.

85
Ⅲ02- 43 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL5U

BL5U

Angle-Resolved Photoemission Spectroscopy Study of Fermi Surface andSuperconducting Gap in NdFeAs0.4P0.6(O,F)
Z. H. Tin1, T. Adachi1, A.Takemori1, S. Ideta2,3 K. Tanaka2,3 S. Miyasaka1 and S. Tajima1

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

hv=60eVky(/a)

2.0

2.0

1.0

1.0

0.0

0.0

-1.0

hv=60eV-1.0-1.0
k x(/a)-1.0
k x (/a)

1.0

1.0

Fig. 1. FS mapping of NdFeAs0.4P0.6(O,F) in P and Spolarized configurations (left and right panels) at 60eV.
-0. 5

k ll
0. 0

0. 5

1. 0

-1. 0

k ll
0. 0

0. 5

1. 0

z

64eV

c)
k (2/

6.5

-0. 5

6.0

z

6.0

k (2/
c)

6.5

-1. 0

k ll
0

6.5
k (2/
c)

z

70e V

6.0

6.0

1. 0

70eV

5.5

z

k(2/
c)

k ll
0. 0

-1. 0

1. 0

6.5

-1. 0

5.5

5.5

64eV

Fig. 2. kz dependence of FSs of NdFeAs0.4P0.6(O,F) in
P and S-polarized configuration (left and right panels).
Top and bottom panels show the results of hole andelectron FSs, respectively.
45ﾟ

-45ﾟ
45ﾟ

28ﾟ

In te n s ity (a rb . u n its )Ⅲ-2

The discovery of IBSs has bestowed us an alternative
direction to investigate the origin of superconductivity.
Amongst the various type of families in Fe-based
superconductors, 1111-system such as NdFeAs(O,F)
achieves the highest Tc, which is 55K. In spite of highest
Tc, 1111-system is not well-researched. The reasons
behind are the difficulty of synthesizing single crystal
and no cleaved surface with electrical neutrality. In ourprevious study, we are able to observe the evolution of
Fermi surfaces and Tc accordance to isovalent doping ofAs site to P. [1] In addition, we had measured thesuperconducting gap of NdFeAs(O,F) by using ARPES
method. In this report, we show the ARPES result of
NdFeAs0.4P0.6(O,F).
The single crystals of NdFeAs0.4P0.6(O,F) weregrown by self-flux method under high pressure using a
cubic anvil press machine. [1, 2] Tc of the single crystal
was about 11 K. The ARPES data were measured at
BL5U of UVSOR facility.
The in-plane mapping of Fermi surfaces (FSs) in P
and S-polarized photon configurations at 60eV wereshown in Fig. 1. In these mappings, three hole FSs andtwo electron FSs can be clearly seen in Γ(0,0) and M(1,1) points, respectively.
The kz dependence of all FSs were measured bychanging the energy of the incident photon and theresults were shown in Fig. 2. All the FSs showed lessdispersive feature along the kz direction. The largest
hole FS (arrow mark) remains around zone center evenby 60% P-doping. This largest FS is treated as surface
state from polar-cleaved surface of the 1111-type
superconductors. Γ points determined to be around 64
eV with inner potential of 18 eV.
The superconducting gap measurement was carriedout on the FSs around the Γ and M points in P-polarized
photon configuration. Figure 3 shows the symmetrizedenergy-distribution curves (EDCs) at the various k pointon the inner and outer hole FSs around the Γ point, andon the inner electron FS around M point. In thesymmetrized EDCs at 5 K, the coherence peaks or theshoulder like structures can be observed around 10 meV,suggesting the superconducting gap ~ 10 meV. However,these peaks or the shoulder like structure do notdiminish at 19 K. This behavior suggests that the peak
and the shoulder like features are due to pseudogap,
which was often observed in the Fe-based
superconductors.

-76ﾟ
14ﾟ

14ﾟ

-99ﾟ

-6ﾟ

-18ﾟ

-20ﾟ
-40ﾟ

0
50
E-E F (meV)

-113ﾟ

-35ﾟ
19K
5K

19K
5K

-50

-108ﾟ

-50

0
50
E-E F (meV)

19K
5K

-50

0
50
E-E F (meV)Fig. 3. Symmetrized EDC of inner (left), outer(middle) hole bands and inner electron band (right) at
5 K (blue) and 19 K (red).
[1] A. Takemori et al., Phys. Rev. B 98 (2018) 100501(R).
[2] A. Takemori et al., JPS Conf. Proc. 1 (2014) 012111.

86
Ⅲ02- 44 -Ⅲ -2. Materials Sciences
BL5U, BL7U
BL5U, BL7U

Study on Interaction between Charge Density Wave Collective Modes andElectrons
Y. Lee, J. Hyun, C.-Y. Lim, G. Lee, S. Kim and Y. Kim

Korea Advanced Institute of Science and Technology (KAIST), Daejeon, 34141, Republic of Korea
Among atomic-layer materials, transition-metal
dichalcogenides (TMDs) have been an attractive
system to research due to their various physicalproperties such as superconductivity, charge-density
wave (CDW), magnetic ordering, etc. Amongst
various properties, competition between CDW phaseand superconductivity provides an important
playground to study the possible connection between
CDW fluctuation and superconductivity. The suitablesystem for such study is 2H-TaSe2. It shows
superconducting transition Tc ~ 0.15 K. Also, there
exists two CDW transitions: from normal phase toincommensurate CDW phase at 122 K and fromincommensurate CDW phase to commensurate CDW
phase at 90 K [1].
In our previous work and other early studies, kinknear Fermi level is observed along M-K highsymmetry direction in 2H-TaSe2 [2, 3] which shows
clear temperature dependence across the CDW
transition. It is not expected behavior of ordinary kinkproduced through electron-phonon interaction. Instead,such temperature dependence implies a possible roleof CDW fluctuation on the observed renormalization.
If we can clearly demonstrate that the kink develops
below CDW transition temperature apart from theband folding and disappears after CDW relaxation, we
can conclude that the coupling between collective
excitation mode of CDW and electron is responsible
for the kink. It will be the first step towards theunderstanding on the relationship between CDW andsuperconductivity.
In this study, we have performed high-resolution
angle-resolved photoemission spectroscopy (ARPES)
on 2H-TaSe2 single crystals grown by chemical vapor
transport (CVT) method [4]. ARPES experiments
were conducted on Beamline 5U and 7U at UVSOR,
using linearly horizontal polarized photon of hv = 50
eV. Total energy resolution was set to 20 meV. Samples
were cleaved in situ in the condition of pressure lower
than 1×10-8 Pa at temperature 10 K. In order to observe
temperature evolution of kink, we measured theelectronic band structure at 10 K, 70 K, 110 K, 150 K
respectively.
Figure 1(a) shows the band dispersion of 2H-TaSe2
along M-K direction taken at 10 K. Dispersion also MK is free from band folding due to symmetry lowering.
We observed the kink at binding energy around 10
meV near M point. Figure 1(b) shows the kink (upper)and corresponding peak positions of momentum
distribution curve (MDC) fitting at four different
temperatures (lower). We observed that kink

disappears above 70K. Therefore, the present resultclearly suggests that the kink is weakened as CDW
phase transition occurs at 90 K, and thus it is originatedfrom electron-CDW coupling.

Ⅲ-2

Fig. 1. (a) Band dispersion of 2H-TaSe2 along M-K
direction of BZ taken at 10 K. Kink structure is shownin the red box. (b) Temperature evolution of kinkstructure and corresponding peak positions of MDC
fitting of 2H-TaSe2 (in the red box of (a)).
[1] D. E. Moncton et al., Phys. Rev. B 16 (1977) 801.
[2] T. Valla et al., Phys. Rev. Lett. 85 (2000) 4759.
[3] K. Rossnagel et al., Phys. Rev. B 72 (2005)
121103(R).
[4] L. H. Brixner et al., J. Inorg. Nucl. Chem. 24 (1962)
257.

87
Ⅲ02- 45 -Ⅲ -2. Materials Sciences
BL5U

BL5U

Polarization-dependent Angle-resolved Photoemission Study of
MAX Phase compound Ti2SnC
1

Ⅲ-2

T. Ito1,2, M. Ikemoto2, D. Pinek3, M. Nakatake4, S. Ideta5,6, K. Tanaka5,6 and T. Ouisse3

Nagoya University Synchrotron radiation Research center (NUSR), Nagoya University, Nagoya 464-8603, Japan
2
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan
3
Grenoble Alpes, CNRS, Grenoble INP, LMGP, F-38000 Grenoble, France
4
Aichi Synchrotron Research Center, Seto 489-0965, Japan
5
UVSOR Synchrotron Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
6
The Graduate University for Advanced Studies, Okazaki 444-8585, JapanMAX phase compounds, i.e., Mn+1AXn where M is atransition metal, A belongs to groups 13-16 and X is the
C or N element, have recently been attracted muchattention due to their possible application for new classof two-dimensional systems called MXenes by removing
A atoms [1]. On the other hand, the bulk electronic
structure of MAX phase has been studied mostly by
calculations, mainly because of lack of well-establishedsingle crystalline samples. In this study, we have
performed angle-resolved photoemission spectroscopy(ARPES) on MAX phase compound Ti2SnC [2] todirectly investigate the electronic structure of this system.
ARPES measurements were performed at the
UVSOR-III BL5U. Data were acquired at T = 25 K withhν = 71.5 eV which enables us to trace around the ΓKM
plane with inner potential of V0=10.7 eV estimated fromthe photon energy dependent measurement (not shown).
With using linear horizontally (LH), vertically (LV) andcircularly right (CR), left (CL) polarized photons, wehave obtained the orbital and orbital-angularmomentum (OAM) dependent ARPES images. It
should be noted that each ARPES images were obtained
without changing the photon incident angle relative tothe sample surface by utilizing two-dimensional
mapping lens mode of MBS A-1 analyzer.
Figure 1 shows the obtained Fermi surface (FS)image on the ΓKM plane (a) compared with the FS
topologies calculated by DFT method (b, c). The band
structure along the ΓM and ΓK line is shown in Fig. 2(a) together with the DFT calculation [2]. Six-folded
symmetry of hole-pockets α and β has clearly been
observed in consistent with the DFT calculations. In
addition, three-fold symmetry around the M point of thesmall FS γ formed by the electron and hole dispersive
feature crossing below the Fermi level along the ΓK line
has also been elucidated in Fig.1 (a). Furthermore, wehave found that the γ branch has been formed by the
Dirac-cone-like dispersive feature along ky axis (Figs. 2(b), (c)). More interestingly, the γ branch shows opposite
OAM sign without changing orbital symmetry withrespect to the ΓM line. The results indicate the changeof chirality of OAM between the two γ branches in Fig.
2 (c). To elucidate the relation between the observed
anomalies and the spin-orbit coupling effect, further
studies are intended.

Fig. 1. Fermi surface (FS) image on ΓKM plane of
Ti2SnC (a). (b,c) FS calculated by DFT shown asprojection along the [001] axis (b) and enlarged FS
around the KH axis (c).

Fig. 2. (a) Band structure along the ΓM and ΓK line of
Ti2SnC obtained by LH polarized photons. Solid lines
are DFT calculation. (b, c) orbital (b) and OAM (c)dependent ARPES image of Ti2SnC along the cut#1 (kx= −0.7 Å-1) in Fig. 1., which is obtained by subtracting
the ARPES results of linearly (LH−LV) and circularly(CR−CL) polarized photons, respectively.
[1] M. Basoum, MAX phases (Wiley, Weinheim 2013).
[2] J. Y. Wu, Y. C. Zhou, J. Y. Wang, Mat. Sci. Eng. A
422 (2006) 266.

88
Ⅲ02- 46 -Ⅲ -2. Materials Sciences
BL5B
BL5B

Study on Optical Properties of Metallic Sodium by Vacuum Ultra Violet in
UVSOR
1M. Kawaguchi 1, J. Saito1, H. Daido2 and T. Suemoto3

Sodium Technology Development Group, Japan Atomic Energy Agency, Tsuruga 919-1279, Japan
2
Institute of Laser Engineering, Osaka University, Suita 565-0871, Japan3
Toyota Physical and Chemical Research Institute, Nagakute 480-1192, Japan

Metallic sodium (Na) is well-known as a typicalalkali metal and very reactive material with oxygen
and/or moisture. The physical properties of Na havebeen researched experimentally and theoretically for along time [1, 2]. Recently, the paper entitled on“Demonstration of partially transparent thick metallic
sodium in the vacuum ultraviolet spectral range [3]”was published by one of the authors (H. D.), and it
demonstrated vacuum ultraviolet light (Wave length:115-180 nm) penetrates in several millimeters of Nawith high transmittance for the first time. To elucidate
theoretically this mysterious and novel phenomenon,
the aim of this research is to obtain the accurate
spectrum using BL5B in UVSOR.
Figure 1 shows the outline of the Na sample which
was fabricated in glove box of argon gas atmosphere(O2 concentration: less than 1.0 ppm, Dew point: lessthan -76 °C) in Sodium Engineering Research Facility(SERF) [4] in Japan Atomic Energy Agency where the
R&D programs on the advanced sodium handlingtechnology have been conducted. The Na sample
which consisted of the inner and safety cells was
designed to improve the safety for the chemical
reactivity and to prevent from oxidizing Na surface.
The inner and safety cells functioned correctly duringtransportation from SERF to UVSOR, and Na sample
maintained metallic luster. The gas conditions in theboth cells were inert argon gas. The ER grade of Na(Na: 99.98%, Ca and K: less than 10 ppm) which wasprovided from Metaux speciaux was filled between
MgF2 windows in the inner cells by immersion in theliquid Na for high quality. Figure 2 shows the Na
sample set up in the chamber of BL5B in UVSOR. The
chamber was vacuumed by turbo-molecular pumps
during measuring the transmittance.
Now, we started the measurement of thetransmittance for the four Na samples using 100-200
nm in wave length: (1) clean MgF2 windows, (2) MgF2
windows after immersion in liquid Na, (3) MgF2
windows + Na (thickness: 0.2 mmt) and (4) MgF2
windows + Na (thickness: 5 mmt). From the differenceof each spectrum, we recognized possibility to occur
the intransparent layer for vacuum ultraviolet light on
the MgF2 windows. In the near future, we will improve
the Na sample to solve the problems and conduct themeasurement again.

< Safety cell : Ar gas >< Inner cell : Ar gas >Spacer: 0.2 - 5mmt

Vacuum seal

Si photodiodeUVSOR BL5B

(Wave length: 100 - 200nm)

Na

Ⅲ-2
MgF2 window

Fig. 1. Outline of the inner and safety cells to prevent
oxidizing sodium surface.

Fig. 2. Na sample (inner and safety cells) set up in thevacuum chamber of BL5B in UVSOR.
[1] J. C. Sutherland et al., Optical properties of sodiumin the vacuum ultraviolet (ORNL-TM-1776, Oak
Ridge National Laboratory, 1967).
[2] R. W. Wood, Phys. Rev. 44 (1933) 353.
[3] H. Daido et al., Optics express 21 (2013) 28182.
[4]
https://www.jaea.go.jp/04/turuga/center/fpc/index.htm.

89Ⅲ02- 47 -Ⅲ -2. Materials Sciences
BL6U

BL6U

Orbital Mapping of One-dimensionally Ordered Picene Film on Ag(110)
Y. Yamada 1, M. Iwasawa1, N. Sumi1, M. Meissner2,3, T. Yamaguchi3, S. Kera2,3 and F. Matsui2
Faculty of Pure and Applied Sciences, University of Tsukuba, Tsukuba 305-8573, JapanUVSOR Synchrotron Facility, Institute for Molecular Science, Okazaki 444-8585, Japan3
School of Physical Sciences, The Graduate University for Advanced Studies, Okazaki 444-8585, Japan
1

2

Ⅲ-2

For an establishment of the state-of-art orbital
mapping and orbital tomography, a well-definedmolecular layer is surely required for a model system.
Especially, the molecular layer with single-domain
structure is preferred for the detailed analysis of theangular distribution of the photoelectrons. However,such a film of the small organic semiconductor
molecules has been rare, because they usually forms
rotational domains even on the single crystalline
substrate.
We have recently shown that, by using small organic
semiconductor molecules with strong intermolecularinteraction such as picene, nicely ordered molecular films
can be realized [1, 2]. We have further demonstrated that,using anisotropic surface of Ag(110), a nearly singledomain film of picene can be fabricated [3]. Here, we
tried to measure angular dependence of the photoelectron
intensity of the single-domain, multilayer film of picene
on Ag(110).
The single-domain film of picene were formed bysupplying approximately 10 nm of picene on the clean
on Ag(110). The thickness of the film was monitoredwith the quartz microbalance and the film structure were
examined in situ by scanning tunneling microscopy(STM) and low-energy electron diffraction (LEED).
Photoemission experiments were done at BL6U of
UVSOR, equipped with A1 analyzer (MB Scientific)
which enables the mechanical deflector scan in the
direction perpendicular to the analyzer slit. The intensity
mapping was measured with increasing acceptanceangle of photoelectrons, by applying potential of 400eV
to the mesh electrode in front of the sample. Sample
temperature was 15 K.
Figure 1 (a) shows the STM image of the multilayer
picene film on Ag(110). It is seen that the multilayer
film film consists of one-dimensional molecular rows
only along [1 -10] direction of the substrate. It is foundthat the molecular long axis of picene in the 1D row
structure is also along [1-10] direction, from thedetailed STM image of the row, shown in Fig.1 (b).
Therefore, it is considered that the picene film on theAg(110) surface is suitable for orbital mapping
because of the fact that the molecular axis is aligned in
one direction. However, we also have to note that thedetailed molecular arrangement in the present structure
have yet been fully clarified.
Figure 2(a) shows the UPS spectrum of themultilayer picene film on Ag(110). It is seen that
HOMO, HOMO-1 and HOMO-2 are close to eachother and difficult to distinguish in the sspectrum,
consistent to the previous reports. However, in thephotoelectron intensity mapping shown in Fig. 2(b),HOMO and HOMO-1 can be well separated; HOMO
maxima comes in the direction perpendicular to themolecular axis, while that of HOMO-1 comes
direction along the molecular axis. This behavior isconsistent with the simple simulation of thephotoelectron distribution, by means of FFT of themolecular orbitals as shown in the inset, suggestingthat the film is indeed consist of the single domain of
picene and that the system is suitable for further
detailed measurements and analysis.
(a)

(b)

Fig. 1. (a) STM image of picene multilayer film on
Ag(110), 200 nm x 200 nm, and (b) a detailed scan.
(a)(b) HOMO
HOMO-2HOMO

HOMO-1HOMO-1

Angle (deg)Fig. 2. (a) UPS spectrum of the picene multilayer
film on Ag(110) (b) Intensity mapping of HOMO and
HOMO-1. FFT of the molecular orbitals are displayedin the inset.
[1] Y. Hasegawa, Y. Yamada, T. Hosokai, Y. Wakayama,
K. Rasika, M. Yano and M. Sasaki, J. Phys. Chem. C 120(2016) 21536.
[2]C. Zhang, H. Tsuboi, Y. Hasegawa, M. Iwasawa, M.
Sasaki, Y. Wakayama, H. Ishii and Y. Yamada, ACS
Omega, (in press).
[3] 長谷川友里，山田洋一，佐々木正洋，表面と真空，2018 年 61 巻 6 号 p. 366-371.

90
Ⅲ02- 48 -Ⅲ -2. Materials Sciences
BL6B

BL6B

Analogy of Trapped Electron Centers in Ce:GGG and Ce:LuGG Crystals
T. Yagihashi1, M. Kitaura1, K. Kamada2, S. Kurosawa2, S. Watanabe3, A. Ohnishi1 and K. Hara4
Faculty of Science, Yamagata University, Yamagata 990-8560, JapanNew Industry Creation Hatchery Center, Tohoku University, Sendau 980-0845, Japan
3
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan
4
Research Institute of Electronics, Shizuoka University, Hamamatsu 432-8011, Japan
1

2

The nature of shallow electron traps in Ce:Gd3Al5-x
GaxO12 crystals has been studied so far. Grigorjeva et
al. have performed luminescence spectroscopy, andtentatively attributed the origin of shallow traps toGa2+ ions perturbed by the nearest neighboring Ce3+
ions [1]. We have carried out infrared absorption
spectroscopy, and found two types of shallow electron
traps [2]. The one is the antisite Gd2+ ion adjacent tooxygen vacancies, which is formed in the range of x <3.5. The other is of an unknown origin, which isformed in the range of x > 3.5. The existence of two
types of electron traps is likely, because the main
component of the conduction band minimum (CBM)
is varied from the Gd3+ 5d state to the Ga3+ 3s statewith increasing x [3]. The Ga3+ 3s character in the
CBM is common among REGa5O12 (RE=Y, Gd, andLu). Therefore, it is expect that the electron trap in
Ce:GGG is of the same origin as that in Ce:LuGG.
In the present study, we have measured infraredabsorption spectra of undoped GGG, Ce:GGG, andCe:LuGG crystals under irradiation and unirradiation
with UV-light. The crystals of Ce:GGG and Ce:LuGG
were grown from high-temperature melt by the micro
pulling down method. The concentration of cerium
ions was set 0.5 mol%. Experiment was performed at
the beamline BL6B. The photon energy of the UVlight was 3.31 eV.
Figure 1 shows UV-induced infrared absorption
spectra of (a) Ce:GGG and (b) Ce:LuGG crystals,
measured at 12 K. The data were obtained bysubtracting unirradated spectra from irradiated specta.
In Fig. 1(a), a mid-infrared (MIR) band appears in theregion below 4000 cm-1. This band was enhanced byhigh-temperatureannealingunderhydrogenatmosphere. In undoped GGG, the MIR band
disappeared completely, in contrast with undoped
GAGG. These facts suggest that Ce3+ ions play animportant role on the occurrence of the MIR band. In
Fig. 1(b), one can see the MIR band below 4000 cm-1,
almost the same as that in Fig. 1(a). From this result, it
is evident that electron traps in Ce:GGG and Ce:LuGG
are associated with Ga3+ ions, but not with Gd3+ andLu3+ ions.
As mentioned above, the CBM is mainly of Ga3+ 3s
state. Generally, the electronic structure of electron
traps reflects the CBM character. On this basis, Ga2+
ions are the most plausible candidate for electron traps
in Ce:GGG and Ce:LuGG crystals. As suggested by
Grigorjeva et al., the formation of such electron traps

may require the perturbation by Ce3+ ions in thevicinity of them. This is supported by missing of the
MIR band in undoped GGG. The enhancement of the
MIR band by high-temperature annealing underhydrogen atmosphere is also explained by the increasein the concentration of Ce3+ ions. In order to clarify theorigin of shallow electron traps responsible for teh
MIR band, further data by spectroscopy experimentand theoretical calculations are needed.

Fig. 1. UV-induced infrared absorption spectra of (a)Ce:GGG and (b) Ce:LuGG crystals measured at 12K.
[1] L. Grigorjeva et al., Opt. Mater. 75 (2018) 331.
[2] M. Kitaura et al., Appl. Phys. Lett. 113 (2018) 041906.
[3] P. Dorenbos, J. Lumin. 134 (2013) 310.

91
Ⅲ02- 49 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL6B
BL6B

Far-Infrared Reflective Analysis for Unipolar Fatigued Piezoelectric Ceramics
H. Nishiyama, D. Xie, R. Hasegawa, S. Maeda, Y. Ito, R. Kimata, N. Nagamatsu,
H. Hashimoto and K. Kakimoto
Graduate School of Engineering, Nagoya Institute of Technology, Nagoya 466-8555, Japan

Ⅲ-2

Piezoceramics is widely used in the industry toconvert electric energy to mechanical energy. As
generally accepted, piezoelectric effect is caused bylattice expansion and domain orientation. The
piezoelectric property deteriorates through repeateduse, and some researchers attribute this cyclic fatigue
to pinning of the domain structure. However, there
have been no investigation for the change of the lattice
vibration. Although far-infrared reflective (FIR)
analysis reveals the lattice vibration and many
researchers investigated for single crystal, nobody
reports for ceramics due to its complexity of vibrationmode assignment. We have measured FIR spectra forlead-free piezoceramics including alkali niobate
perovskite. In this study, we next investigated for
BaTiO3 which has already reported in single crystal [1].
BaTiO3 ceramics was prepared by conventional
solid state sintering. For comparing as-sintered andfatigued samples, unipolar cyclic E-field of 100 Hz
with a maximum field of 3 kV/mm was applied up to
105 cycles. FIR measurement was conducted by a FTIR spectrometer (Bruker, VERTEX 70v). A beam
splitter and a detector used in this measurement werea Mylar 6 μm and a Si bolometer, respectively. Notethat aperture diameter was 6 mm and then these were
hundred thousands of grains in the spotted area.
Figure 1 shows lattice vibration modes of BaTiO3
ceramics. From the FIR spectra shown in Figs. 1a andd, imaginary part of permittivity (Figs. 1b and e) was
calculated with Drude-Lorentz model [2]. Infraredactive vibration modes of tetragonal perovskite are
explained as 4E + 3A1, but we cannot fit with sevenvibration modes for ceramics well because of not only
the ratio of the polarization on its surface but also the
distribution of the domain width, which can make adeviation of the locations due to internal stress causedby surrounding domain structure. This is why we used
plenty number (23 waves) of Lorentz functions forfitting. Imaginary part of permittivity calculated with
Drude-Lorentz model was fitted again with Voigt
functions to separate overlapping of each peaks (Figs.
1c and f). According to the fitting result for crystal [1],these three peaks were identified Last, Slater, and Axe
modes, respectively. Moreover, Slater mode in 105
cycles shifted higher by 32.2 cm-1 comparing to the assintered one. Here, Slater mode locates 34 cm-1 in aaxis and 280 cm-1 in c-axis oriented single crystal [1].
Considering a polarization ratio of as-sintered sample
should be 2/3 for a-axis and 1/3 for c-axis due to its
randomness, the amount of c-axis oriented domainshould be increased. In other words, 105 cycles appliedBaTiO3 ceramics became more like c-axis oriented
single crystal. In these ways, the change of lattice
vibration for unipolar fatigued piezoceramics can beevaluated with FIR analysis.

Fig. 1. Lattice vibration modes calculated from FIR
spectra of BaTiO3 ceramics.
[1] T. Hoshina et al., Jpn. J. Appl. Phys. 53 (2014)
09PD03.
[2] H. Nishiyama et al., UVSOR Activity Report 2016
44 (2017) 84.

92
Ⅲ02- 50 -Ⅲ -2. Materials Sciences
BL6B

BL6B

Material Control by High Intensity THz Wave and Physical Property
Evaluation by Infrared Spectroscopy
A. Irizawa

The Institute of Science and Industrial Research, Osaka University, Ibaraki 567-0047, Japan

Optical study is one of the most powerful techniques
for investigating electronic states on solids. We canobtain direct of information about the electronic states,
a band structure, a symmetry of crystal structure, anda dielectric response of materials. In view ofexperimental affinity, the optical study is extremelycompatible with extreme conditions such as low
temperature, high magnetic field, and high pressure,excited state by pumping light, and their combinations.
The observation of electronic states can be concluded
with only optical operations in contrast with e.g.
photoemission spectroscopy. Additionally small
sample area and/or convoluted mixed-phase have needof microscopic technique, in such view optical study
including infrared spectroscopy is one of the best
choice for determining electronic states of materials.
Beamline 6B in UVSOR is adjusted for theinvestigations for optical study in a low-energy regionof infrared, i.e. FIR and MIR. In this report, we have
performed the optical transmittance measurements inthe longest wavelength region. The introduced dry-air
system is investigated by check the absorption lines ofgas components in the air.
The experiments are performed by using
Michelson-type interferometer in MIR region with
infrared microscope. Before this experiment, thissystem was operated in the air without closed
atmospheric condition. Otherwise we can get the best
condition by using N2 gas flow with rather quantity,
one compressed gas cylinder for 2 days. This time,
active-dry air system has been installed for thisbeamline and the performance was checked by IR
spectroscopy. Figure 1 shows the spectra of black-bodylight source with the atmospheric conditions of open
in the air, flowing dried air, and flowing N2 gas fromcompressed cylinder with a fine resolution of Δ=1 cm-1.
Between 1000 and 2000 cm-1, 3000 and 4000 cm-1, and5000 and 6000 cm-1, there shows clustered fineabsorption lines of the vapor in the air. The
atmospheric condition of humidity was around 15 %in the beamline which was not wet condition, but the
MIR light was strongly absorbed by vapor in theseregions. In case of dried air flow and N2 gas flow are
better than this, except for the absorption line of CO2
around 2300 cm-1, the obtained light source spectra areapplicable for estimating the ratio of spectra, i.e.
absorptance, transmittance, and reflectance. This
improvement must be the advantage for microscopic
spectroscopy under extreme conditions including theexperiments using organic materials that haveabsorption lines in an infrared region.

Ⅲ-2

Fig. 1. Spectra of light source with different atmosphericconditions.

93
Ⅲ02- 51 -Ⅲ -2. Materials Sciences
BL7U

BL7U

High-resolution ARPES Study of Nodal Fermion Materials
T. Sato1,2, Z. Wang3, K. Nakayama1, S. Souma2,4, D. Takane1, Y. Nakata1, K. Hori1,
H. Oinuma1, T. Takahashi1,2,4 and Y. Ando2

Department of Physics, Tohoku University, Sendai 980-8578, Japan2
Physics Institute II, University of Cologne, 50937 Köln, Germany
3
Center for Spintronics Research Network, Tohoku University, Sendai 980-8577, Japan
4
WPI Research Center, Advanced Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan1

Ⅲ-2

The search for new types of topological materials
hosting nodal fermions is one of the emergent topics incondensed-matter physics. One of the effective
strategies to search for nodal fermions is to utilize thepoint-group symmetries of crystal, i.e., mirror reflection,
rotation, and inversion symmetries in addition to timereversal symmetry, as highlighted by the discovery of
topological crystalline insulators hosting the surface
nodal fermions protected by mirror reflection symmetry.
Recent theoretical studies also predicted that
nonsymmorphic space-group symmetry combining
point-group symmetry and the fractional translation
further enriches the category of nodal fermions. This isdemonstrated by the prediction/observation of nodal
loops protected by the glide mirror (mirror reflection
plus translational) symmetry [1-3], as well as the Weylnodes in trigonal Te and the nodal lines in ZrSiS, bothof which are protected by the screw (rotation plustranslational) symmetry [2, 4, 5].
Recently, it was theoretically proposed that thelayered ternary telluride Ta3SiTe6 hosts nodal fermionsprotected by nonsymmorphic glide mirror symmetry[6]. This material crystallizes in the orthorhombic
structure with the space group No. 62 (Pnma). As
shown in Fig. 1(a), the basic structural unit of Ta3SiTe6
is a Te trigonal prismatic slab with Ta atoms locatedaround the center of this prism. Each unit cell contains
two such slabs which are overlaid with each other by
the inversion operation. First-principles band-structure
calculations [6] show that, when the spin-orbit
coupling (SOC) is neglected, Ta3SiTe6 displays afourfold-degenerate nodal line on the SR line in thebulk Brillouin zone [see Fig. 1(b)] due to theprotection by glide mirror symmetry. It is alsosuggested that when the SOC is included, the fourfold
degeneracy is slightly lifted and as a result
hourglasslike dispersions appear. To examine such
intriguing predictions, it is highly desirable toexperimentally establish the electronic band structureof Ta3SiTe6.
In this study, we performed low-energy highresolution ARPES measurements of Ta3SiTe6 at BL7U,
and found a signature of Dirac-like band dispersion
protected by the glide mirror symmetry.
Figure 1(c) shows a representative ARPES intensity
in the valence-band region of Ta3SiTe6 measured at T =40 K with hν = 26 eV photons. We find several highly
dispersive bands within 1 eV of the Fermi level. Thesebands mainly originate from the Ta 5d orbitals
hybridized with the Te 4p orbitals. We have also
performed detailed three-dimensional k-space mappingby tuning photon energy and observed the banddegeneracy that which supports the existence of nodal
fermions predicted by the calculation.

Fig. 1. (a) Crystal structure of Ta3SiTe6. (b) Bulk
orthorhombic BZ (blue) and corresponding surface BZ
projected onto the (001) plane (green). (c) ARPES
intensity as a function of binding energy and wave
vector measured at T = 40 K with linearly polarized
26-eV photons.
[1] Q. Xu et al., Phys. Rev. B 92 (2015) 205310.
[2] L. M. Schoop et al., Nat. Commun. 7 (2016) 11696.
[3] D. Takane et al., Phys. Rev. B 94 (2016) 121108(R).
[4] C. Chen et al., Phys. Rev. B 95 (2017) 125126.
[5] K. Nakayama et al., Phys. Rev. B 95 (2017) 125204.
[6] S. Li et al., Phys. Rev. B 97 (2018) 045131.

94
Ⅲ02- 52 -Ⅲ -2. Materials Sciences
BL7U
BL7U

Electronic State of the Two-dimensional Transition Metal Chalcogenide
Ferromagnets Studied by ARPES
M. Suzuki1, B. Gao2, K. Koshiishi1, C. Lin1, Y. X. Wan1, M. Kobayashi3,4, S. Ideta5,
K. Tanaka5, S.-W. Cheong2 and A. Fujimori1

1
Department of Physics, The University of Tokyo, Tokyo 113-0033, Japan
Rutgers Center for Emergent Materials and Department of Physics and Astronomy, Rutgers University,
Piscataway, New Jersey 08854, USA
3
Center for Spintronics Research Network, The University of Tokyo, Tokyo 113-0033, Japan
4
Department of Electrical Engineering and Information Systems, The University of Tokyo, Tokyo 113-8656, Japan
5
UVSOR Synchrotron Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
2

Since the discovery of graphene, there has been
tremendous interest in the development of new twodimensional (2D) materials and their functionalities[1]. The materials have layered crystal structures andthe layers are bonded to each other through van der
Waals (vdW) forces. In addition to graphene andtransition-metal dichalcogenides, vdW ferromagnets
have recently attracted much attention as candidate
materials for new types of spintronic devices.
Cr2Ge2Te6 (CGT) and Cr2Si2Te6 (CST) arerepresentative vdW ferromagnets, and along with other
vdW ferromagnets, e.g., CrI3, their electronic andmagnetic properties have been studied theoretically andexperimentally [2, 3]. The magnetism of CGT nanosheets
shows soft behaviors and is well described by the 2D
Heisenberg model [4].
To unveil the origin of the ferromagnetism in 2D
materials, understanding the physical properties fromthe electronic structure points of view is important. For
both fundamental understanding and potential
applications of 2D vdW ferromagnets, the knowledgeof the electronic structure of CGT related to theferromagnetism is indispensable. Therefore, we haveconducted angle-resolved photoemission spectroscopy(ARPES) in order to clarify the electronic structure of
CGT and CST.
Figures 1(a) and 1(b) show ARPES spectra of CGT
obtained by using p- and s-polarizations. The center ofthe cuts corresponds to the Γ point. One can see that ahole-like band centered at Γ is located at ~0.2 eV below
the Fermi level and the valence-band maximum does
not reach the Fermi level, indicating a semiconducting
property, consistent with the electric transport
properties [2, 5]. In addition, a similar hole band exists
at ~0.5 eV below the upper hole band. Photoemission
spectra of CGT are shown in Fig. 1(c). Difference
between s- and p-polarizations has been clearlyobserved. Especially, the position of the main peakaround 2 eV is shifted between the two polarizations.
Since Cr 3d electrons are supposed to form partial
density of states (PDOS) around the binding energy of~2 eV, according to first-principle calculation, we
consider that the PDOSs of different Cr 3d orbitals havebeen observed.

On the other hand, for CST, we could not observe
any differences as shown in Fig. 1(d). Moreover, its
spectral line shape is completely different from that of
CGT. However, it is not natural that electronic statesof the Cr 3d electrons is dramatically changed by only
substituting Si for Ge. Therefore, we consider that theresults for CST maybe due to extrinsic effects.
In order to reveal the origin of the ferromagnetismin these materials, further studies are required.

Fig. 1. ARPES spectra of .van der Waals ferromagnets.
(a), (b) Energy-momentum plots of Cr2Ge2Te6 using pand s-polarizations. (c) Comparison of the valenceband photoemission spectra of Cr2Ge2Te6 between pand s-polarizations. (d) Photoemission spectra of thevalence band of Cr2Si2Te6.
[1] K. S. Novoselov et al., Science 306 (2004) 666.
[2] X. Zhang et al., Jpn. J. Appl. Phys. 55 (2016) 033001.
[3] G. T. Lin et al., Phys. Rev. B. 95 (2017) 245212.
[4] C. Gong et al., Nature 546 (2017) 265.
[5] H. Ji et al., J. Appl. Phys. 114 (2013) 114907.

95
Ⅲ02- 53 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL7U
BL7U

Angle-resolved Photoemission Study of
Ullmannite NiSbS with the Cubic Chiral Crystal Structure1

Ⅲ-2

T. Ito1,2, T. Hosoya2, M. Nakatake3, S. Ideta4,5, K. Tanaka4,5, M. Kakihana6, D. Aoki7,
A. Nakamura7, T. Takeuchi8, H. Harima9, M. Hedo10, T. Nakama10 and Y. Onuki10

Nagoya University Synchrotron radiation Research center (NUSR), Nagoya University, Nagoya 464-8603, Japan
2
Graduate School of Engineering, Nagoya University, Nagoya 464-8603, Japan
3
Aichi Synchrotron Research Center, Seto 489-0965, Japan
4
UVSOR Synchrotron Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
5
The Graduate University for Advanced Studies, Okazaki 444-8585, Japan
6
Graduate School of Engineering and Science, University of the Ryukyus, Okinawa 903-0213, Japan
7
Institute for Materials Research, Tohoku University, Ibaraki 311-1313, Japan8
Low Temperature Center, Osaka University, Osaka 560-0043, Japan
9
Graduate School of Science, Kobe University, Kobe 657-8501, Japan
10
Faculity of Science, University of the Ryukyus, Okinawa 903-0213, Japan

Non-centrosymmetric compounds have attracted
much attention due to their interesting phenomena
relating with inversion symmetry breaking. The
ullmannite type NiSbS with the non-centrosymmetriccubic chiral structure has recently been suggested to bea possible candidate of Wely semimetal [1]. Indeed, split
Fermi surface topologies reflecting the cubic chiral
crystal structure and the spin-orbit interactions havebeen reported [2]. In this study, we have performed
angle-resolved photoemission spectroscopy (ARPES)
on NiSbS to directly investigate the electronic band
structure of this system.
ARPES measurements were performed at the
UVSOR-III BL7U. Data were acquired at T = 25 K withhν = 32 eV which enables us to trace around the XMRplane with inner potential of V0=24.7 eV estimated fromthe photon energy dependent measurement (not shown).
Clean surfaces were obtained by in situ cleaving on(001) plane.
Figure 1(a) shows the obtained Fermi surface (FS)image on the XMR plane compared with the reported
FS topologies [2]. The band structure along the XM line
is shown in Fig. 2 together with the band calculation [2].
From the present ARPES study, it has been found thatthe electronic structure of NiSbS seems to be well
reproduced by the calculation along the XMR plane. For
example, an electron pocket around the X point which
touches the Fermi level and then disperse to 0.5 eVaround the R point has clearly been observed in Fig. 2.
On the other hand, we found non-negligible
contributions of the ΓXM plane and the surface states.
The bands due to electron (hole) -like FS branch α (ε)
appear and merge with the bands along the XR line. Thismight originate from a large kz-broadening effect [3]. In
addition, narrow dispersion around 0.5 eV should be
assigned as surface states. The results suggest theapplicability of ARPES to pursue the Weyl points andnode features originating in the chiral crystal structurewith great care about bulk sensitivity as well as spin
(orbital) –polarizations at the electronic structure.

Fig. 1. (a) Fermi surface (FS) image on the XMRplane of NiSbS. Solid and dashed green lines are FS
topologies on the XMR and ΓXM plane reported in ref.
[2], respectively. (b) Cubic Brillouin zone shown withthe momentum axis.

Fig. 2. Band structure image along the XM line of
NiSbS. Solid and dashed lines are band calculation
along the XR and ΓM lines, respectively [2].
[1] A. Furusaki, Sci. Bulletin 62 (2017) 788.
[2] M. Kakihana et al., J. Phys. Soc. Jpn. 84 (2015)084711.
[3] T. Mitsuhashi et al., Phys. Rev. B 94 (2016) 125148.

96
Ⅲ02- 54 -Ⅲ -2. Materials Sciences
BL7U, BL5U

BL7U, BL5U

Direct Observation of the Energy Gap Generation of
Ruthenium mono-arsenide RuAs
Y. Nakajima1, T. Nakamura1, Y. Ohtsubo2,1, H. Kotegawa3, H. Sugawara3, H. Tou3
and S. Kimura1,2
1
Department of Physics, Osaka University, Toyonaka 560-0043, Japan
Graduate School of Frontier Biosciences, Osaka University, Suita 565-0871, Japan3
Department of Physics, Kobe University, Kobe 657-8501, Japan

2

Ruthenium mono-pnictides RuPn (Pn = P, As, Sb)
have been recently noted in the relation to iron pnictide
superconductors. Actually, Rh-doped RuP and nondoped RuSb become superconductor at low
temperatures [1]. Other interesting properties are thecharge-density wave in RuP [2] and the metal-toinsulator transition (MIT) in RuAs [1]. Thesephenomena are believed to be related to the origin of thesuperconductivity, but it is under debate at present [3].
RuAs, which is one of the RuPn family, is being
well investigated. The MIT originates from theformation of a superlattice of 3×3×3 of the originalunit cell [4]. However, the origin of the superlattice
formation has not been clarified yet. One possibility is
the electronic instability, and the other the lattice
distortion owing to the zigzag-chain of Ru site. To
clarify the origin of MIT, we are investigating theelectronic and the phonon structure as a function oftemperature.
According to the electronic band calculations, thedensity of states (DOS) within the energy of about 0.2 eVfrom the Fermi level (EF) decreases in the lowtemperature insulating (LT) phase from the hightemperature metallic (HT) phase as shown in Fig. 1, eventhough the finite DOS at EF still remains in the LT phase,i.e., no energy gap is expected. However, the opticalconductivity spectra as well as the electrical resistivity
have suggested that a clear energy gap appears in the LT
phase [5]. Then, we have performed the temperaturedependent angle-integrated photoelectron spectroscopy(AIPES) of RuAs to investigate whether the energy gap
can be realized at low temperatures or not.
Figure 2 shows the temperature-dependent AIPES
spectra divided by the Fermi-Dirac distributionfunctions for each temperature, which can be comparedwith DOS. The intensity at EF at 280 K (in the HTphase) is almost flat suggesting the constant DOS nearEF, but it slightly decreases with decreasing temperature
to 13 K in the LT phase. The energy at which thedivergence starts with decreasing temperature is about
0.15 eV, which corresponds to the energy where DOSchanges from the HT phase to the LT phase. Therefore,
the change of electronic structure was clearly observedin the AIPES spectra.

Ⅲ-2

Fig. 1. Density of states (DOS) of RuAs as hightemperature (HT) and low-temperature (LT) phases.
The zero energy was set as the Fermi level (EF). The
intensity in the energy region within 0.15 eV below
and above EF decreases from the HT phase to the LT
phase, but finite DOS remains even at the LT phase.

Fig. 2.
Temperature-dependent angle-integrated photoelectron
spectra of RuAs divided by the Fermi-Dirac distributionfunctions (FDs) for each temperature convoluted with aGaussian function of σ of 10 meV. The excitation
photon energy was 20 eV. (Inset) Same as the main
figure but without the division of FDs.
[1] D. Hirai et al., Phys. Rev. B 85 (2012) 140509(R).
[2] R. Y. Chen et al., Phys. Rev. B 91 (2015) 125101.
[3] Y. Kuwata et al., J. Phys. Soc. Jpn. 87 (2018)073703.
[4] H. Kotegawa et al., Phys. Rev. Mater. 2 (2018)055001.
[5] Y. Nakajima et al., UVSOR Activity Report 201745 (2018) 84.

97
Ⅲ02- 55 -Ⅲ -2. Materials Sciences
BL7U
BL7U

Measurement for the Dispersion of the Excited States in NbSe2 by Using the
Photon-Energy-Dependent ARPES
2

Ⅲ-2

S. Tanaka1 and K. Ueno2

The Institute of Industrial and Scientific Research, Osaka University, Ibaraki 567-0046, Japan
Department of Chemistry, Graduate School of Science and Engineering Saitama University, Saitama 338-8570,
Japan
1

The electronic-excitation dynamics of the material
is a key issue for understanding and developing thefunctional
opticaldevice.
Transition-metaldichalcogenides (TMDC) attracts interests of many
researchers because of its unique 2D character. Here,
we report the photon-energy dependent angle-resolved
photoelectron spectroscopy (ARPES) of NbSe2(2H),
which is one of the most studied TMDC, and examine
the band-dispersion of the excited state.
All the experiments were carried out at BL7U of the
UVSOR-III. The NbSe2 sample was cleaved in UHV andthe photon-energy-dependent ARPES was measured at
10K. The photon intensity was calibrated by the use of
the photodiode.
Figure 1 show the two-dimensional photoelectron
intensity map of the normal emission as functions of thebinding energy and photon energy. Peaks at the binding
energy around -1eV clearly show enhancements in
intensity at the photon energy of about 9.6eV. This canbe attributed to the resonant enhancement when thephoton energy matches the energy-difference between
the initial (occupied) and final (empty) bands of thephotoexcitation. Therefore, the energy position of theexcited (empty) bands can be deduced from the photonenergy-dependent measurement since the initial(occupied) band can be determined from the ordinal(non-resonant) ARPES measurement. The Fig. 2(a)shows the ARPES spectrum along the Γ-K line taken athν=11.4eV. We focus the two bands indicated as “upper”
and “lower” bands in the figure, and measured theirintensity as a function of the photon energy. Then theenergy is converted as Ef(k)=hν+EB(k), where Ef(k)denotes the final-state-energy at a specific electron
momentum of k, hν the photon energy and EB(k) the
binding energy in the negative scale at k, which can bederived from the dispersion curves [solid lines in Fig.2(a)]. The results are shown for several momenta in Figs
2(b) and (c) for upper and lower bands, respectively.
The intensity distributions can be well represented by
the Lorentzian curves whose peak positions changewith the electron momenta. These peak positions should
indicates the position of the excited (empty) bands into
which the photo-excitation occurs from the upper andlower bands in Fig. 2(a). Figures 2(d, e) show intensity
maps of the photoelectron peaks as functions of the final
state energy with respect the Fermi level and theelectron momentum. The peak positions obtained in
Figs. 2(b, c) are shown as solid lines, and these aredispersions of the excited bands of NbSe2.
These experimental results are well explained by thecomparison to the 1st principles band calculation
including the analysis of the photoexcitation selection
rules based on the band symmetry.

Fig. 1. Photoelectron intensity map of NbSe2 asfunctions of photon energy (vertical axis) and binding
energy (horizontal axis). Lines indicate the position of
the same final state energies.

Fig. 2. (a) APRES map of the NbSe2 at hν=11.4eV; (b,c) the intensity of the bands indicated in Fig. 2(a) as asa function of the photon energy after the energy
conversion; (d, e) the dispersion of the excited stateobtained from (b, c).

98
Ⅲ02- 56 -Ⅲ -2. Materials Sciences
BL7U

BL7U

Temperature Dependence of the Kondo Resonance Peak in Photoemission
Spectra of Rare-earth Compound YbMgCu4
K. Morikawa1, H. Shiono1, H. Sato2, S. Ideta3, K. Tanaka3, T. Zhuang4,
K. T. Matsumoto4, K. Hiraoka4 and H. Anzai1

1
Graduate School of Engineering, Osaka Prefecture University, Sakai 599-8531, Japan
Hiroshima Synchrotron Radiation Center, Hiroshima University, Higashi-Hiroshima 739-0046, Japan
3
UVSOR Synchrotron Facility, Institute for Molecular Science, Okazaki 444-8585, Japan
4
Graduate School of Science and Engineering, Ehime University, Matsuyama 790-8577, Japanthe observed 4f7/2 peak can be assigned to the Kondo
resonance peak in YbMgCu4. This consistency implies
that the observed 4f7/2 state in this study is indeed
responsible for the heavy-fermion behavior of YbMgCu4.
(a)Intensity (arb.units)

4f5/2

4f7/2

YbMgCu4hν = 24 eV
T=8K
-1.5

-1.0
-0.5
Energy, ω (eV)

(b)
8K
57 K
80 K
122 K
187 K

-0.4

-0.2
0
Energy, ω (eV)

0

(c)

80

Peak Position (meV)

In rare-earth compounds, the interaction between
itinerant conduction and localized-4f electrons (c-fhybridization) leads to a renormalized quasiparticle.
Such a heavy-fermion behavior is characterized by aresonance peak appearing at the energy scales of the
Kondo temperature TK in electronic excitation spectra.
For Yb compounds, the spin-orbit split Yb2+ 4f7/2 stateis interpreted as the Kondo resonance peak [1]. The
Yb2+ 4f7/2 state near the Fermi level (EF) will provides
important clues to the possible mechanism of theheavy-fermion states.
YbMgCu4 with C15b-type structure is known as aheavy fermion compound with the electronic specificheat coefficient γ ~ 62 mJ/K2mol [2]. The Kondo
temperature is estimated to be TK ~ 855 K from themagnetic susceptibility measurements [2]. Previous
photoemission spectroscopy measurements presented
a broad Yb2+ 4f7/2 peak at ~240 meV, which deviateconsiderably from the value of Kondo temperature [3].
To detect the heavy-fermion states in YbMgCu4, it isessential to identify the Yb 4f state near EF.
Here, we report on a study of the Kondo resonancepeak in photoemission spectra of YbMgCu4. The
experiments were performed at BL7U of UVSOR.
The energy resolution was 16 meV. The samples werecleaved in situ and maintained under ultrahigh vacuum(8×10-9 Pa) during the measurements.
Figure 1(a) shows the valence-band spectra at T = 8
K. The spin-orbit splitting of the Yb2+ 4f7/2 and Yb2+
4f5/2 states are observed at |ω| ~ 0.07 eV and 1.36 eV,
respectively [3, 4]. The observability of the 4f7/2 statenear EF is substantially improved by slightly
increasing the photon energy from those used inprevious studies, hν ~ 21 eV [3].
Figure 1(b) shows the temperature dependence ofthe photoemission spectra near EF. We found that a
sharp peak of the 4f7/2 state survives at even higher
temperature, T ~ 190 K, and the peak energy remains
unchanged. We determined the peak energies andplotted them as a function of temperature in Fig. 1(c).
The peak energy slightly decreases with decreasing
temperature from T = 187 K to 122 K, and then tends
to saturate at low temperatures. The peak energies
extrapolated to zero temperature is estimated to be ~69 meV. This value corresponds to TK ~ 805 K, which
is approximately consistent with TK reported from themagnetic susceptibility measurements [2]. Therefore,

Intensity (arb.units)

2

60
40
20
0

0

100
200
Temperature (K)Fig. 1. (a) Valence-band spectra of YbMgCu4 takenwith photon energy hν = 24 eV at T = 8 K. (b)Temperature dependence of the valence-band spectra
near EF. Vertical bars indicate the peak positions. (c)Energy shift of the Yb2+ 4f7/2 peak in panel (b). The
error bars of energies derived from the uncertainty ofthe peak positions.
[1] R. I. R. Blyth et al., Phys. Rev. B 48 (1993) 9497.
[2] J. L. Sarrao et al., Phys. Rev. B 59 (1999) 6857.
[3] H. Sato et al., J. Synchrotron Rad. 9 (2002) 229.
[4] F. Reinert et al., Phys. Rev. B 58 (1998) 12808.

99
Ⅲ02- 57 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL7U

BL7U

ARPES Study of Novel Topological Material Candidates
H. Oinuma1, S. Souma2,3, D. Takane1, T. Nakamura1, K. Hori1, R. Tsubono1, K. Nakayama1,
K. Sugawara1,2, T. Takahashi1,2,3, A. Ochiai1 and T. Sato1,2

Department of Physics, Tohoku University, Sendai 980-8578, Japan
Center for Spintronics Research Network, Tohoku University, Sendai 980-8577, Japan3
WPI Research Center, Advanced Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan
1

2

Ⅲ-2

Topological insulators (TIs) are a novel quantum
state of matter where gapless edge or surface states(SSs) appear within a bulk-band gap inverted by strong
spin-orbit coupling. The SSs of three-dimensional (3D)
TIs are characterized by the linearly dispersive Diraccone band that can be viewed as a sea of Dirac fermions
whose massless nature is guaranteed by the timereversal symmetry. The search for new topologicalmaterials is currently a hot topic in condensed-matterphysics.
Recently, it was predicted from the first-principles
band-structure calculations [1] that lanthanum (La)monopnictide with rock-salt structure LaX (X = N, P, As,
Sb, and Bi) becomes either topological Dirac semimetal(for X = N) or 3D TI (for X = P, As, Sb, and Bi) with
Dirac fermions at the surface, due to the band inversion
at the X point of bulk fcc Brillouin zone (BZ) [see Fig.
1(a)]. While rare-earth monopnictide RX (R: rare earth)
was intensively studied in 1980-90’s in relation to
heavy-fermion physics, this theoretical proposal [1]renewed the interest for RX in topological aspects andtriggered intensive transport, spectroscopic, andtheoretical studies. The occurrence of magnetic order insome RX compounds, e.g., cerium (Ce) monopnictides
CeX, also provides a rare opportunity to investigate theinterplay between topological properties andmagnetism, which has largely been left unexploredbecause of the absence of suitable material platform.
Therefore, experimental elucidation of the topologicalproperties in RX compounds is highly desired.
In this study, we performed high-resolution angleresolved photoemission spectroscopy (ARPES) of RX
family. By utilizing low photon energy and energytunable characteristics of BL7U, we determined thebulk and surface band structures and investigatedtopological properties.
Figures 1(b) and 1(c) show ＿the ARPES spectra
measured along a cut crossing the  point of the surface
BZ and the corresponding intensity plot, respectively,
obtained in the paramagnetic state of CeBi. We find thatat least threehole-like bands cross the Fermi level (EF)＿near the  point. Two of them show parabolic
dispersions and attributed to the topmost bulk valencebands with the Bi 6p orbital character, consistent with
our recent soft-x-ray ARPES measurements on RX
family [5]. The rest hole-like band that shows a linear
dispersion is assigned as the topological Dirac-cone SSs
with the Dirac point in the close vicinity of EF. By
performing high-resolution measurements at low

temperatures, we searched for magnetism-inducedchanges in the topological SSs, and also investigated the
topological property of other RX compounds.

Fig. 1. (a) Bulk and surface BZs with the schematicbulk Fermi surfaces at the Γ point (green) and the X
point (red and blue). (b) ARPES spectra
measured＿along a momentum cut crossing the  point of thesurface BZ with hν = 22 eV. (c) ARPES intensity plotof (b) as a function of binding energy and wave vector.
[1] M. Z. Hasan and C. L. Kane, Rev. Mod. Phys. 82(2010) 3045.
[2] X.-L. Qi and S.-C. Zhang, Rev. Mod. Phys. 83(2011) 1057.
[3] Y. Ando, J. Phys. Soc. Jpn. 82 (2013) 102001.
[4] M. Zeng et al., arXiv:1504.03492.
[5] H. Oinuma et al., Phys. Rev. B 96 (2017) 041120(R).

100
Ⅲ02- 58 -Ⅲ -2. Materials Sciences
BL7U

BL7U

Electronic Structure of Underdoped Triple-layer Cuprate
Bi2Sr2Ca2Cu3O10+δ
S. Ideta1,2, S. Adachi3, N. Sasaki3, S. Yamaguchi3, T. Watanabe3 and K. Tanaka1,2

National Institutes of Natural Science, Institute for Molecular Science, Okazaki, 444-8585, Japan
2
The Graduate University for Advanced Studies (SOKENDAI), Okazaki, 444-8585, Japan
3
Graduate School of Science and Technology, Hirosaki University, Hirosaki 036-8561, Japan1

To understand the mechanism of high-Tc
superconductivity in cuprates is one of the crucial issues
and the energy gap seen in the superconducting (SC)and normal states is believed to be an important piece ofevidence for the non-BCS behavior of the SC transition
for the underdoped regime in the high-Tc cuprate
superconductors. Two different representative energy
gaps are present in the SC and normal states, namely,
the SC gap and pseudogap in cuprate superconductors[1-3]. Origin of the pseudogap has been considered that
it is related to the superconductivity or a phenomenon
distinct from superconductivity, and sheds light on its
nature to elucidate the mechanism of high-Tc
superconductivity.
Triple-layer Bi-based cuprate superconductor,
Bi2Sr2Ca2Cu3O10+δ (Bi2223), which shows the highest
Tc (110 K) among the Bi-based cuprates, has two
inequivalent CuO2 planes with different hole carrier [46]. In this study, we have investigated the energy gap
anisotropy and we found an as-yet-unknown energy gap
with a node which is neither the general pseudogap nor
the nodeless energy gap in underdoped regime. We
observed two bands corresponding to the outer andinner CuO2 planes and the hole carrier deduced from the
IP and OP FSs suggests that IP is almost zero carrier and
OP is underdoped.
High-quality single crystals of underdoped
Bi2Sr2Ca2Cu3O10+δ (UD Bi2223, Tc = 80 K) weregrown by the TSFZ method. Underdoped Bi2223
samples were successfully obtained in the two-stepannealing in order to control their doping levels [7].
ARPES experiments were carried out at BL7U of
UVSOR-III Synchrotron. Clean sample surfaces wereobtained for the ARPES measurements by cleaving
single crystals in-situ in an ultrahigh vacuum better
than 6×10-9 Pa. The measurements were performed at
12 K and 95 K.
Figure 1(a) shows the ARPES spectra of UD Bi2223.
ARPES intensity mapping integrated within ±5 meVcentered at EF is shown taken at hν = 18 eV. Two Fermi
surfaces (FSs) are designated as the inner (IP) and outer(OP) CuO2 planes reported as previous ARPES studies[4, 5]. The FS momentum (kF) positions for IP (bluedots) and OP (red dots) have been determined by thepeak position of momentum- distribution curves in thenormal state as shown in Fig. 1(b). The hole carrier for
IP and OP deduced from the FS area are 0-2% and 810%, respectively, and hence, surprisingly the inner

CuO2 plane is almost non-doping. To see themomentum dependence of the pseudogap and the SC
gap, the energy gap is estimated from the peak positionof the energy-distribution curves and plotted asfunctions of a simple d-wave order parameter |cos(kxa)- cos(kya)|/2 [Figs. 1(c) and 1(d)]. As clearly seen in Fig.
1(c), the energy gap for IP is relatively temperatureindependent across Tc and the FS only remains nearnode above Tc as well as that in the SC state. On theother hand, in Fig. 1(d), the energy gaps for OP show ad-wave-like evolution at 12 K and close the gap andhave a Fermi arc around the nodal region at 95 K. To
evaluate the validity of the present finding, it is highly
desired to demonstrate a further systematic
experimental study in the near future.

Fig. 1. Electronic structure of underdoped Bi2223.
(a) ARPES intensity plots taken at T = 95 K. (b) Fermi
momentum plots on the outer (OP) and inner (IP)
Fermi surfaces. (c), (d) Temperature dependence of theenergy gap for OP and IP plotted as functions of the dwave order parameter.
[1] D. S. Marshall et al., Phys. Rev. B 76 (1996) 4841.
[2] A. G. Loeser et al, Science 273 (1996) 325.
[3] H. Ding et al., Nature (London) 382 (1996) 51.
[4] S. Ideta et al., Phys. Rev. Lett. 104 (2010) 227001.
[5] S. Ideta et al., Phys. Rev. B 85 (2012) 104515.
[6] S. Ideta et al., Physica C 470 (2010) S14.
[7] S. Adachi et al., Physics Procedia 65 (2015) 53.

101
Ⅲ02- 59 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL7B
BL7B

Photoluminescence and Scintillation of Rb1-xCsxCaCl3
1

Ⅲ-2

K. Takahashi1, M. Koshimizu1, Y. Fujimoto1, T. Yanagida2 and K. Asai1

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

Auger-free luminescence (AFL) is caused by theradiative transition of an valence electron to a core hole
state. The transition is competitive with the Auger
process, which is a much faster process, efficient AFL
cannot be observed in many compounds. In other words,
efficient AFL has been observed in limited compounds
such as alkaline- or alkaline-earth fluorides or CsClbased compounds.
AFL has been applied to fast scintillators owing to its
fast decay. A representative one is BaF2; however, CsClbased compounds have an advantage of long emission
wavelength appropriate for detection with aphotomultiplier tube. Thus far, we have developed fast
scintillators based on ternary CsCl-based compounds[1, 2]. In this study, we developed fast scintillators basedon quaternary compounds, Rb1-xCsxCaCl3.
The crystals of Rb1-xCsxCaCl3 was grown in a simple
solidification method after heating the mixed raw
powders at 400 K to remove the adsorbed water.
Figure 1 shows the X-ray-induced radioluminescence
spectra. A prominent band is observed at approximately
450 nm. The excitation spectra for this band exhibited apeak within the band-gap energy of the crystals.
Therefore, the band at 450 nm is attributed to thetransition at localized levels within the band-gap. Figure
2 shows the X-ray-induced radioluminescence spectra in
the UV region. Several emission bands were observed at
250 and 300–400 nm. We obtained the excitation spectraof these bands in the VUV region to discuss the origin of
the emission bands. The excitation spectra exhibited a
step-like structure at approximately 84 nm, whichcorresponds to the energy difference between theoutermost core level and the bottom of the conduction
band. On the basis of this result, the emission bands are
attributed to AFL.
Figure 3 shows the pulse-height spectra of thescintillation detectors equipped with the Rb1-xCsxCaCl3
crystals for 662-keV gamma-rays from 137Cs. Note thatthe coarse gain (CG) of the main amplifier was different
for the Rb1-xCsxCaCl3 crystals and GSO as a reference.
Based on the comparison of the full-energy peakchannel, the light yields of Rb1-xCsxCaCl3 crystals were
estimated to be 520–1200 photons/MeV.

Fig. 1. X-ray-induced radioluminescence spectra of
Rb1-xCsxCaCl3 crystals.

Fig. 2. X-ray-induced radioluminescence spectra of
Rb1-xCsxCaCl3 crystals in the UV region.

Fig. 3. Pulse-height spectra of the scintillation detectors
equipped with the Rb1-xCsxCaCl3 crystals for 662-keV
gamma-rays from 137Cs.
[1] N. Yahaba, M. Koshimizu, Y. Sun, T. Yanagida, Y.
Fujimoto, R. Haruki, F. Nishikido, S. Kishimoto and
K. Asai, Appl. Phys. Express 7 (2014) 062602.
[2] M. Koshimizu, N. Yahaba, R. Haruki, F. Nishikido,
S. Kishimoto and K. Asai, Opt. Mater. 36 (2014) 1930.

102
Ⅲ02- 60 -Ⅲ -2. Materials Sciences
BL7B
BL7B

Impurity Luminescences of Bulk and Thin Film AlN
W. Kamihigoshi1, Y. Maegawa1, K. Fukui 1, K. Yamamoto2, N. Tatemizo3 and S. Imada3

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
3
Faculty of Electrical Engineering and Electronics, Kyoto Institute of Technology, Kyoto 606-8585, Japan1

excited above the band gap energy and the excitation
spectra of these two bands are similar to the absorption
spectrum of AlN. As described above, the absorption
coefficient of AlN(film) from 4.2 eV to 5.2 eV is 100
times higher than that of AlN(bulk). This means that in
the case of AlN(film), the band-edge tail is able tooverlap with this 4.2-5.2 eV absorption band, and theinterband transition transfers its energy to both the UVband and the BH band emissions.

Fig. 1. CEES contour plots of AlN(bulk) andAlN(film).

Intensity (arb.units)

AlN(bulk)
peak fit
BH
BL

2

3

4

5

Emission Energy (eV)Fig. 2. Emission spectra of AlN(bulk).
BL

BH

UV

Intensity (arb.units)

Since AlN has a wide energy band gap of about 6
eV at room temperature, many unintentionallyimpurity absorption bands are possible in the bandgap.
The information in these bands is important for bothelectrical and optical applications of AlN.
Photoluminescence measurement is useful for the
study of impurity bands, as the absorption process ofthese bands often causes a luminescence process. In
this report, we focused on the blue - ultraviolet
absorption band that makes AlN yellow and the
combined excitation emission spectra (CEES) of bothbulk crystal and sputtered AlN(film) were measured.
A 0.5 mm thick c-plane AlN bulk crystalline
substrate (AlN(bulk)) studied in this report was grownin PVT (physical vapor transport) method by
CrystaAl-N GmbH. AlN substrate is transparent withyellow dots and has characteristic impurity absorption
band of ~10+2 cm-1 from 4.2 eV to 5.2 eV region. The
c-plane AlN thin film (AlN thin film) was depositedby RF (radio frequency) sputtering on SiO2 glass, andshows slightly pale yellow color. Film thickness isabout 1 μm and there is an absorption band of ~10+4
cm-1 at around same energy region [1]. The CEES
measurements were carried out at 10 K, in the order of
10-6 Pa, and from 2.2 eV to 10.0 eV.
Figure 1 shows the CEES of AlN(bulk) (left) andAlN(film) (right). Each contour plot is normalized at
the maximum intensity. Both AlN(bulk) and AlN(film)
have similar emission spectra with two peaks at 2.8 eV
and around 3.9 eV. However, excitation spectra of thetwo peaks are different for AlN(bulk) and AlN(film).
The CEES of AlN(bulk) is very similar to the results
for the AlN(bulk) substrates previously reported by
Alden et al. [2]. Figure 2 shows the emission spectrumof AlN(bulk) at an excitation energy 4.5 eV. The 3.9
eV emission peak is represented by a single Gaussianfit (UV band), while the 2.8 eV peak is represented by
two Gaussian fit (BH and BL bands). Figure 3 shows
integrated excitation spectra of UV, BH and BL bands.
The UV band and the BH bands have the same
excitation process which is excited only between 4.2
eV and 5.2 eV, but the excitation process of the BL
band is different and excited up to band gap energy of
AlN. On the other hand, both the 2.8 eV and the 3.9
eV peaks of AlN(film) are represented by a single
Gaussian fit, and their bandwidth reflecting crystal
quality is wider than that of AlN(bulk). Figure 1
suggests that the 2.8 eV peak of AlN(film) is consists
only of the BH band, since the 2.8 eV excitation
spectrum is similar to that of the 3.9 eV peak. Figure 1
also suggests both the UV band and the BH bands are

4.0

4.5

5.0

5.5

6.0

Excitation Energy (eV)Fig. 3. Excitation spectra of AlN(bulk).
[1] N. Tatemizo et al., J. Mater. Chem. A 5 (2017) 20824.
[2] D. Alden et al. Phys. Rev. Applied 9 (2018) 054036.

103
Ⅲ02- 61 -Ⅲ-2

Ⅲ -2. Materials Sciences
BL7B
BL7B

Observation of Mott Gap in Pyrochlore Ruthenates R2Ru2O7
R. Kaneko1, M. Masuko1, R. Yamada1, J. Fujioka2, K. Ueda1 and Y. Tokura1

Department of Applied Physics, University of Tokyo, Tokyo 113-8656, Japan
Graduate School of Pure and Applied Sciences, University of Tsukuba, Tsukuba 305-8577, Japan

Ⅲ-2

The electronic correlation plays a crucial role on theirelectronic property in transition metal oxides. A typicalexample is the Mott transition. Coulomb interaction
localizes electrons at each atomic site in a Mott insulator.
The effective control of the electron correlation, i.e.
bandwidth control, gives rise to a number of exotic
phenomena in the vicinity of Mott transitions [1].
The pyrochlore ruthenate R2Ru2O7 (RRO, R=Pr, Nd,… Lu) is an insulator which undergoes anantiferromagnetic transition at a low temperature,reminiscent of Mott systems. The effective electronic
correlation varies as a change of R ionic radius throughthe modulation of Ru-O bond angles which is related tothe electron hopping integral (the larger R gives thesmaller Ueff). Actually, a specific heat measurement
reveals that the Neel temperature TN, which is roughly
proportional to 1/Ueff in a simple localized-electron
picture, increases with an increase of the ion radius of
R3+[2].
What is missing so far is the insight into the electronic
band structure. In this experiment, we have aimed tomeasure optical reflectivity to obtain the opticalconductivity which encompasses an important
information about the electronic state including themagnitude of the charge gap, interband electronic
excitations, and so on.
The optical conductivity can be obtained by KramersKronig (KK) transformation of reflectivity in the wholeenergy range. Therefore, we measured the reflectivity ina range from 0.01 eV to 5 eV at our laboratory, and from
2 eV to 40 eV at UVSOR. This measurement enables us
to discuss the charge dynamics in the energy region ofinterest. We employed a high-pressure technique for
growth of samples which is hard and dense enough to bewell polished for the optical measurement.
Figure 1 shows the optical reflectivity for R=Pr, Nd,
Sm, and Y compounds. By the KK transformation of thereflectivity, the optical conductivity is obtained asshown in Fig. 2. A charge gap is clearly observed for all
samples. We defined the gap size as a value at thecrossing point between an energy axis and theextrapolated line from the low-energy slope of each
spectrum (dashed lines). The estimated gap size isplotted as a function of the ionic radius of R3+ in Fig. 3.
The gap size Δ becomes larger as the ionic radius of R3+
decreases, which is consistent with the existing theory.
It indicates that the electronic states in the 4d electron
system, which is characterized by a weak electron
correlation and large spin-orbit coupling compared tothose in 3d systems, are mainly governed by theelectronic correlation.

Reflectiv

1

2

Photon Energy (eV)Fig. 1. Optical reflectivity of R2Ru2O7 (R=Pr, Nd, Sm,Y) at the room temperature.

Photon Energy (eV)Fig. 2. Optical conductivity of R2Ru2O7 (R=Pr, Nd,
Sm, Y) at 10 K.

Fig. 3. The magnitude of the charge gap as a functionof R ionic radius.
[1] M. Imada, A. Fujimori and Y. Tokura, Rev. Mod.
Phys. 70 (1998) 1039.
[2] M. Ito et al., J. Phys. Chem. Solids 62 (2001) 337.

104
Ⅲ02- 62 -Ⅲ -2. Materials Sciences
BL7B
BL7B

Evaluation of Electrical Property of CaF2-metal Interface and Developmentof Vacuum Ultraviolet Detector
S. Kato, J. Otani, K. Suzuki and S. Ono

Nagoya Institute of Technology, Nagoya 466-8555, Japan
Vacuum ultraviolet light sources are applied invarious fields such as surface modification ofmaterials and cleaning of semiconductor substrates.
For monitoring such a light source, a detector forstably monitoring vacuum ultraviolet light is required.
Therefore, we have focused on the fluoride material
which is a wide gap material and has high durability
against vacuum ultraviolet light irradiation [1-4].
Using such fluoride material made it possible todevelop filterless and durable photodetectors. In thisstudy, we aimed to develop a detector capable ofoperating at zero bias by selectively bringing thefluoride compound CaF2 single crystal into contact
with metal.
To fabricate the device, a thin film of Au and Al was
deposited on both surfaces of a CaF2 single crystal
plate by vacuum deposition method. In order to allow
the Au electrode to function as a transparent electrode,
the Au electrode was deposited with a thickness
allowing vacuum ultraviolet light to pass through (Fig.
1).
In order to investigate a performance of thefabricated device as photodetector, we measured thecurrent value when the device was irradiated withultraviolet light (photocurrent) and not irradiated (darkcurrent). It was confirmed that the dark current value
was about 0.18 pA, the photocurrent value was about
602 pA, and the current value increased more than 3
digits by light irradiation at an applied bias of 100 V.
Even at zero bias, the current value increased morethan 3 digits by light irradiation. Furthermore, it was
confirmed the rectifying property depending on bias
direction as shown in Fig. 2.
We measured spectral sensitivity characteristics ofthe device at zero bias and transmittance of CaF2
single crystal as shown in Fig. 3. It was confirmed thatthe wavelength response region of the detector is lessthan 120 nm and coincide with the transmission end of
the CaF2 single crystal.
In summary, we achieved filterless VUV detector
meeting practical performance by applying CaF2
single crystal and metal interface.

Ⅲ-2
Fig. 1. Transmission spectra of Au electrode.

Fig. 2. The rectifying property depending on bias
direction.

Fig.3. Transmission spectrum and spectral response of
the detector.
[1] M. Ieda, T. Ishimaru, S. Ono, N. Kawaguchi, K.
Fukuda, T. Suyama, Y. Yokota, T. Yanagida and A.
Yoshikawa, Jpn. J. Appl. Phys. 51 (2012) 062202.
[2] T. Ishimaru, M. Ieda, S. Ono, Y. Yokota, T. Yanagida
and A. Yoshikawa, Thin Solid Films 534 (2013) 12.
[3] M. Yanagihara, H. Ishikawa, S. Ono and H. Ohtake,
Testing and Measurement: Techniques and Applications(2015) 151.
[4] K. Suzuki, M. C. Raduban, M. Kase and S. Ono,
Optical Materials 88 (2019) 576.

105
Ⅲ02- 63 -Ⅲ -2. Materials Sciences
BL7B

BL7B

Optical Study of FeSe Thin Films on LaAlO3 Substrate
M. Nakajima1, K. Yanase1, M. Kawai2, T. Ishikawa2, N. Shikama2, F. Nabeshima2,
A. Maeda2 and S. Tajima1
Department of Physics, Osaka University, Osaka 560-0043, Japan
Department of Basic Science, University of Tokyo, Tokyo 153-8902, Japan
1(a)

LaAlO30.01

0.1
1Energy (eV)

10
(b)

0

1

400
200
0
-200
-400
0

500

2

1.0
0.8
0.6
0.4
0.2
0

0.02

0.04 0.06
Energy (eV)

0.08

0.10

Fig. 1. (a) Reflectivity spectrum of LaAlO3 at room
temperature up to 40 eV. (b) Real and imaginary partsof dielectric function of LaAlO3 below 0.1 eV.
1.0

R

0.9

FeSe film on LaAlO3

(a)

0.8
0.7
0.6

1 (103 -1cm-1)

Ⅲ-2

In iron-based superconductors, the electronic
structure is strongly affected by in-plane strain. In
particular, the band structure of FeSe thin films differsfrom that of bulk FeSe due to the epitaxial strain [1].
The correlation between a superconducting transition
temperature (Tc) and the magnitude of strain indicatesthat the band-structure modification definitely alteredphysical properties [2], but how the electronic statechanges with application of strain is still unclear. In
this study, we performed optical spectroscopy on thinfilms of FeSe on LaAlO3 substrates with a tensile strain
and investigated the effect of in-plane strain.
Thin films of FeSe with the thickness of 150 nm were
fabricated on LaAlO3 by the pulsed-laser deposition
method [2]. Optical reflectivity measurement was
carried out on the FeSe film and the bare LaAlO3
substrate for the energy range of 0.05–2 eV using theFourier-transform infrared spectrometer (Bruker Vertex80v). High-energy reflectivity measurement of LaAlO3
for 2–40 eV at room temperature was performed at
BL7B of the UVSOR synchrotron facility. Dielectric
function of LaAlO3 was derived via the Kramers-Kronigtransformation, and the thin-film reflectivity spectrum
was fitted using a number of Lorentz oscillators, which
gives the dielectric function of FeSe [3].
Figure 1(a) shows the reflectivity spectrum of
LaAlO3 for a wide range of energy up to 40 eV. Three
optical phonon modes dominate the low-energy region
below 0.1 eV, whereas several peaks associated withinterband transitions are present in the high-energy
region above 6eV. In Fig. 1(b), we plot the dielectric
function of LaAlO3 derived from the reflectivity data.
The phonon modes are observed at 22.6, 53.0, and 80.8
meV.
The reflectivity spectrum of the FeSe film on
LaAlO3 at room temperature is shown in Fig. 2(a). Due
to multiple reflections within the film, the phononmodes of substrate LaAlO3 makes an impact on thereflectivity spectrum, giving rise to three peaks atcorresponding energies. Using the dielectric functionof LaAlO3, we can remove the substrate contribution
and extract the optical conductivity spectrum of FeSe.
The result is shown in Fig. 2(b). The peak at 29.6 meVcorresponds to an optical phonon mode of FeSe.
Compared with the spectrum of FeSe with acompressive strain [3], the low-energy spectral weightis smaller in the present case. The smaller carrier
density is likely related to the suppression of Tc.

R

2

3

FeSe

2
1
0
0

(b)
0.1
Energy (eV)0.2

Fig. 2. (a) Reflectivity spectrum of the FeSe thin filmon the LaAlO3 substrate at room temperature. (b)Optical conductivity spectrum of FeSe.
[1] G. N. Phan et al., Phys. Rev. B 95 (2017) 224507.
[2] F. Nabeshima et al., Jpn. J. Appl. Phys. 57 (2018)120314.
[3] M. Nakajima et al., Phys. Rev. B 95 (2017) 184502.

106
Ⅲ02- 64 -Ⅲ -2. Materials Sciences
BL7B
BL7B

Optical Investigations of Pr/Ce-doped APLF Glasses
as Potential Neutron Scintillators
T. Shimizu, Y. Minami, M. J. F. Empizo, K. Kawano, K. Kuroda, K. Yamanoi,
N. Sarukura and T. Murata
Institute of Laser Engineering, Osaka University, Suita 565-0871, Japanand Ce for improvement of decay times and conversion
efficiencies. In this study, we investigated the opticalproperties of Pr/Ce-doped APLF glasses with different
concentrations using the spectroscopy beamline (BL7B)
of the Institute of Molecular Science UVSOR
Synchrotron Facility. Figure 2 shows the PL spectra of
the Pr/Ce-doped APLF glasses. This results are presently
under analysis.

Fig. 1. Decay profiles of the (4f5d → 3H6) emissionsof (a) , (b) , (c) , and (d) Pr3+-doped APLF glasses atdifferent temperatures under excitation into the 4f5d
level [1]1.0

PrCe_co-doped_APLF PLE specrum
1Pr_0.1Ce
1Pr_0.2Ce
1Pr_0.5Ce
1Pr_1Ce_1
1Pr_1Ce_2
0.6
2Pr_1Ce
1Pr_2Ce
0.5Pr_1Ce
0.4
0.2Pr_1Ce
0.1Pr_1Ce0Pr_1Ce

0.8
Intensity (arb.unit)

Recently, neutron detection is required in inertial
confinement fusion research and infrastructure
inspection. The ratio of primary neutron and scatteredneutron corresponds to the density and radiusparameter of fusion plasma. Additionally, neutron
detection is considered as one of the nondestructive
methods for inspecting the form of infrastructure.
Scintillator materials which are able to detect anddiscriminate neutrons should then be developed to beable to satisfy the technological demand.
Previously, the optical properties of praseodymium(Pr3+)-doped and cerium (Ce3+)-doped APLF
[20Al(PO3)3-80LiF] glasses was studied as fast
response scintillators for neutron detection [1]. The
optical properties were characterized using
photoluminescence and photoluminescence excitationof the Pr3+- and Ce3+-doped APLF glass samples for
doping concentrations ranging from 0.1 to 3.0 mol%.
The APLF glass doped with Pr3+ revealed richemission bands under 180, 217, and 430-440 nm
excitations. The broad emission from 228 to 371 nm
for both excitations were assigned to theinterconfigurational 4f5d and intraconfigurational 4f
transitions of Pr3+ ions. The intensity of the emission
peaks was observed to increase as a result of
increasing Pr3+ concentration. In contrast, the APLF
glass doped with Ce3+ exhibited intense emission at
340 nm under 240 to 400 nm excitation which
originated from the 4f5d transition of Ce3+ ions. There
was no fluorescence from 4f transitions, but the strong
emission peak at 340 nm was shifted to shorter
wavelengths with decreasing Ce3+ concentration. The
decay times of APLF80+Pr3+ under 217 nm excitation
were constant at different temperatures from 0 K to300 K and became faster with increased doping
concentration from 19 ns (0.5% Pr3+) to 16 ns (3.0%
Pr3+). The decay times of APLF80+Ce3+ glasses were
the same in the range from 38 to 41 ns with neutron
excitation from a 252Cf source regardless of Ce
concentration. These results highlight that the
scintillation decay times from both Pr3+- and Ce3+doped APLF glasses are significantly faster thanconventional glass scintillators and therefore anadvantage for fast-response scintillator applications.
Investigation is underway to improve the scintillation
decay times and conversion efficiencies by co-doping
APLF glass with Pr3+ and Ce3+.
Now, we have developed co-dope APLF glass with Pr

0.2

0
100

150

200

250

300

350

Wavelength (nm)Fig. 2. PLE spectrum of Pr3+ and Ce3+ co-doped
APLF glass by using UVSOR.
[1] M. J. F.Empizo et al., J. Lumin. 193 (2018) 13.

107
Ⅲ02- 65 -Ⅲ-2

UVSOR User 3