BL7B

Scintillation and Luminescence Properties of Sn2+-doped Oxide Glass[1]H. Masai1, T. Yanagida2, Y. Fujimoto2, M. Koshimizu3, T. Yoko1 and K. Asai3
1

Institute for Chemical Research, Kyoto University, Kyoto 611-0011, Japan
2
Kyushu Institute of Technology, Kitakyushu 808-0196, Japan3
Department of Applied Chemistry, Graduate School of Engineering, Tohoku University, Sendai 980-8579, Japan
RE (rare-earth)-free phosphors have long beendeveloped in terms of abundance as natural resource.
Among them, RE-free glasses have advantage thatvarious kinds of elements can be incorporated tocontrol the emission properties. One application of the
RE-free glasses may be scintillators, because most of
conventional scintillators include RE ions asluminescent centers and as the constituent of the hostmatrix. In this study, we demonstrated theapplicability of RE-free glasses for scintillators.
The SnO-SrO-B2O3 and SnO-ZnO-P2O5 glass
samples were fabricated by a conventional
melt-quenching technique. The raw materials weremelted in platinum crucibles, and subsequently werequenched on a stainless steel at 473 K. The glass
samples were annealed at the glass transition
temperatures.
The pulse height spectra of the D-ray from 241Am
were obtained for the scintillation detectors quippedwith sample scintillators. The emission and excitation
spectra of samples were measured at room
temperature under the irradiation of synchrotron
radiation having energies of 6–20 eV at the UVSOR
facility (BL7B).
Figure 1 shows the pulse height spectra of thesamples. Because the experimental set up and theexcitation source was the same for each sample, themultichannel analyzer (MCA) channel is roughly
proportional to the number of photoelectrons. A clear
peak was observed for each sample. This result
indicates that the light yield was sufficient for
detection of ionizing radiation whose energy is on theorder of MeV, although the light yield wassignificantly smaller than that of Li-glass.
Figure 2 shows contour plot of emission andexcitation spectra of the 0.5SnO-24.5SrO-75B2O3 and
5.0SnO-60ZnO-40P2O5 glass. Each contour plotshows the photon energy of excitation (ordinate) andemission (abscissa), and the intensity axes are shownon an identical linear scale. We observed two
excitation bands; one is the band at ~6 eV
corresponding to the S0-S2 transition of Sn2+. We
observed another peak at above 20 eV. Assuming that
this band corresponds to the excitation in the host glass,the emission occurs via the energy transfer from thehost to the Sn2+ center. The observation of thehigh-energy band suggests that effective scintillation
occurs, because the energy transfer is essential in
obtaining effective scintillation.

Counts (arb. units)

(a)

(b)

X 10

X 0.1

䎓

䎔䎓䎓

䎕䎓䎓

(c)

Li-glass䎖䎓䎓䎗䎓䎓

MCA channel

Excitation energy / eV

Fig. 1. Pulse-height spectra of the samples for 241Am
D-ray; (a): 5SnO-60ZnO-40P2O5,(b): 0.5SnO-24.5SrO-75B2O3,(c): 0.1SnO-24.9SrO-75B2O3.

Excitation energy / eV

Emission energy / eV

Emission energy / eV

Fig. 2. Contour plot of emission and excitation
spectra of the 5.0SnO-60ZnO-40P2O5 (lower) and
0.5SnO-24.5SrO-75B2O3 (upper).
[1] Hirokazu Masai, et al., Appl. Phys. Lett., 101(2012) 191906.

123

BL2A

Mo LIII-Edge XANES Study of MoO3-Modified H-GaAlMFI Catalysts for
Methane Dehydroaromatization
H. Aritani1, T. Sugawara1, N. Naijo1, K. Takanashi1, S. Mogi1 and A. Nakahira2
1

Department of Life Science & Green Chemistry, Saitama Institute of Technology, Fukaya 369-0293, Japan
2
Graduate School of Engineering, Osaka Prefecture University, Sakai 599-8531, Japan

MoO3-modified H-MFI (Mo/H-MFI) is a typicalcatalyst for methane dehydroaromatization, which isan important reaction for direct GTL (Gas to Liquid)
processes. The H-MFI zeolite (SiO2/Al2O3=28-40)
shows unique properties for dehydroaromatization oflight alkanes such as propane and ethane. The
catalytic activity is based on strong acidity andselective sieving effects. By MoO3 modification, it
shows high activity for methane dehydroaromatiozation (MTB reaction). In this reaction,
MoOx species are reduced and carbonized to form
Mo2C-like carbide species. The carbide species plays
a key role in methane dehydrogenation. As a sametime, carbon contamination on Mo/H-MFI proceeds.
It is a major source for catalytic deactivation. Because
the deactivation relates to the strong acidity on
H-MFI, an improvement of H-MFI support is called
for. In the present study, Ga-containing H-MFI(GaAl-MFI) supports have been synthesizedhydrothermally [1], and MoOx-modified GaAl-MFI
catalysts have been applied to employ the methane
dehydroaromatization. By substitution of Ga ions
onto H-MFI framework, the acidity weakens slightly.
The effect of Ga ion onto H-MFI, highly active anddurable catalysts can be expected. In the present
report, Mo species on GaAl-MFI catalysts have beencharacterized by Mo LIII-edge XANES to evaluate thecarbide species after the MTB reaction.
Catalysts were prepared by impregnation of
H-GaAlMFI support with MoO2(acac)2íCHCl3
solution, and followed by drying overnight andcalcination at 773 K for 3 h. The amount of
MoO3-loading is 5.0 wt% in this study. H-GaALMFI
supports were synthesized hydrothermally at 413 K
for a week, and followed by ion-exchanging with
NH4Cl and calcination at 873 K. The catalytic activityof methane dehydroaromatization was evaluated bymeans of fixed bed flow reaction, as described in a
separate paper [2]. The catalyst samples are denoted
as Mo/GamMFIn, in which m=Ga/Al and n=Si/Al2
atomic ratios. Mo LIII-edge XANES spectra were
measured in BL2A of UVSOR-IMS in a
total-electron yield mode using InSb double-crystal
monochromator. Photon energy was calibrated byusing Mo metal-foil at Mo LIII-edge, and normalized
XANES spectra and their second derivatives arepresented. REX-2000 (Rigaku) software was used bynormalization of each XANES spectrum.
Figure 1 shows the Mo/H-MFI28(Ga-free),
Mo/GaAl-MFI28 (intraframework Ga onto MFI with
Ga/Al=10-50), and Mo-Ga/H-MFI28 (Ga2O3-MoO3

Fig. 1. Mo LIII-edge XANES spectra (top) and theirsecond derivatives (bottom) over MoOx- modified
GaAl-MFI catalysts after MTB reaction.
co-impregnation) catalysts after MTB reaction. For
all the catalysts, the edge energy values of L-XANES
are lower than bare Mo2C, suggesting the formationof micro-particles which consist of metallic andcarbonized Mo. Mo/H-Ga50MFI28 shows the
maximum reactivity for MTB reaction, and thus, Ga
ions do not affect the Mo carnbonization. Mo-Ga
interaction can not be seen in both the structural
evaluation and catalytic properties. It is concluded
that formation of active Mo carbide species proceeds
independently of Ga modification.
[1] K. Nagashima, S. Nakamura, K. Okada, A.
Nakahira and H. Aritani, Bull. Chem. Soc. Jpn. 82(2009) 1203.
[2] H. Aritani, H. Shibasaki, H. Orihara and A.
Nakahira, J. Environm. Sci. 21 (2009) 736.

124

BL2A

Local Structure of Na-K and K-K edge of Artificial Zeolites Prepared from
Blast Furnace Slag
M. Sato1, K. Kumadani2, T. Shirai2 and A. Nakahira1, 2
1

Kansai Center for Industrial Materials Research, Tohoku University, Osaka 599-8531, Japan
Department of Material Science and Engineering, Osaka Prefecture University, Osaka 599-8531, Japan2

result shows that exchanged K+ ions in zeolite
prepared from BF slag exist in similar environment
with zeolite 3A.

The depletion of natural resources is one of the ourserious problems. Therefore, the development of thenovel recycling technology has been performedenergetically. Blast furnace slag (BF slag) is one of
the waste materials that the development andestablishment of the effective recycling technology
and system become the urgent assignment. The BF
slag has been produced approximately 25 million tons
per year, and it is recycled as Portland cement, basecoarse materials, fine aggregate for concrete andground improvement materials, and so on. However,recycling rate of BF slag is insufficient by comparing
to the amount of generation. In this study, A and Y
type zeolite were prepared using BF slag as a starting
material to reduce an emission of waste materials, andlocal structure of Na-K and K-K edge of products
were measured in order to investigate the condition of
exchanged cation in zeolite.
BF slag and sodium aluminate were used as
starting materials. Firstly, BF slag was ground to
become their particle size less than 70 Pm. Groundslag powder was undergone an acid treatment toremove Ca component. Then, Ca removed slag
powder was hydrothermally treated in NaOH solution
at 368 K. In this study, the ratio of Si to Al waschanged in the range from 0.5 to 2.4 by addition of
sodium aluminate. Obtained samples were ion
exchanged in 1000 ppm KNO3 solution at room
temperature for several periods. K-K edge XANES
spectra of obtained samples were measured by a total
electron yield mode at room temperature using InSb
double-crystal monochrometer, and Na-K edge
XANES spectra was also measured by a total electron
yield mode at room temperature using Beryl
double-crystal monochrometer at BL-2A station of
UVSOR.
Figure 1 shows Na-K edge XANES spectra of A
type zeolite prepared from BF slag before and after
ion exchange. Commercial A type zeolite and A typezeolite prepared using reagents were used as
references. The position of second peak around 1080eV was different between references and sample. This
result shows that Na environment in zeolite preparedfrom BF slag more distorts than references [1].
In the case of K-K edge XANES spectra of A typezeolite prepared from BF slag after ion exchange.
Although clear spectra did not obtain for the sample
ion exchanged for 5 and 30 min, obtained spectra
were similar spectra with zeolite 3A, which is thepotassium form of the type A crystal structure. This

Fig. 1. Na-K edge XANES spectra of A type zeolite
prepared from BF slag and references. (a) Commercial
A type zeolite, (b) A type zeolite prepared using
reagents, (c) A type zeolite prepared from BF slag, (d)after ion exchange for 30 min.
[1] D. R. Neuville, L. Cormire, A-M. Flank, R. J.
Prado and P. Lagarde, Eur. J. Mineral 16 (2004)
809-816.

125

BL2A

Investigation of Local Structure of P-K edge of Fe doped Amorphous
Calcium Phosphate Powders
M. Sato1, T. Nagayasu2, S. Hayashi2 and A. Nakahira1, 2
1

Kansai Center for Industrial Materials Research, Tohoku University, Osaka 599-8531, Japan
Department of Material Science and Engineering, Osaka Prefecture University, Osaka 599-8531, Japan

2

Hydroxyapatite (Ca10(PO4)6(OH)2, HAp) has been
well understood as an alternate materials of our hardtissue because of their excellent osteoconductivity
and biocompatibility. Besides, HAp also has good
protein adsorption ability and ion exchange ability.
Therefore, many studies for bioactivity anddegradable behavior of Mg, Mn, Fe, Si, Zn doped
HAp have been performed.
In a previous study, we prepared Fe doped HAp
powders by co-precipitation and hydrothermal
method, and investigated the effect of Fe addition on
the crystal phase, microstructure, local structurearound Fe, Ca and P atoms and other properties [1, 2].
From the result of XAFS measurement, it wassuggested that the added Fe mainly existed as
D-FeOOH and D-Fe2O3 phase, and small amount of
Fe was substituted into the HAp structure. Since thegeneration of D-FeOOH phase is caused by the alkali
condition of preparation method, sol-gel method,which does not require the alkali condition andtypical method to obtain amorphous calcium
phosphate, was used as a preparation method.
Metal Ca, phosphoric acid and Fe acetylacetonate
were used as starting materials. These were mixed up
in ethanol solution at 353 K for 4 hours becoming the
Ca/P ratio of 1.67 and the ratio of Fe from 1 to 10
mol% for Ca. After that, they were filtered and driedat 323 K for overnight. P-K edge XANES spectra ofobtained HAp powders were measured in a total
electron yield mode at room temperature using InSb
double-crystal monochrometer at BL-2A station of
UVSOR.
Figure 1 shows P-K edge XANES spectra of Fe
doped amorphous calcium phosphate powders
prepared by sol-gel method. Commercial HAp wasused as a reference. The obtained spectra of Fe doped
sample showed similar spectra independent from theratio of Fe. This result indicates that the local
structure around P atom of Fe doped samples does notchange by Fe addition. Besides, clear peak shift tohigh energy site and spectrum broadening wereobserved by comparison with the commercial HAp.
From these results, it is suggested that relatively
higher amount of Fe can be doped than our previous
studies [1, 2], and this data is important to clarify thecondition and solid solubility limit of Fe in the HApstructure.

Fig. 1. P-K edge XANES spectra of Fe doped
amorphous calcium phosphate powder prepared bysol-gel method.
[1] M. Sato, S. Yamamoto, Y. Nishio, Y. Takamatsu
and A. Nakahira, UVSOR Activity Report 38 (2011)115.
[2] M. Sato, Y. Kawabe, S. Misu, S. Hayashi and A.
Nakahira, UVSOR Activity Report 39 (2012) 81.

126

BL2A

Study of the Local Strucures of LTA-type Zeolite Nanopowder by XANES
Evaluations
A. Nakahira1, 2, T. Tagami1, T. Shirai1, M. Sato2, T. Nagayasu1 and H. Aritani3
1

Faculty of Engineering, Osaka Prefecture University, Gakuencho, Sakai 599-8531, Japan
2
Osaka Center, IMR, Tohoku University, Gakuencho, Sakai 599-8531, Japan
3
Saitama Institute of Technology, Fukaya 369-0293, Japan

indicated that the obtained products were LTA typezeolites. XRD results showed that after themechanical mixing processing for 2 hs the groundproducts were also identified to be LTA type.
Products obtained after the mechanical mixing
processing had the average particle of a few hundred
nanometer in diameter. Suggesting the nanosized LTA
zeolite powders were successfully obtained by after
the mechanical mixing processing for 2 hs.
Fig.1 shows the XANES spectra of zeolites andreference materials. As shown in Fig. 1, XANES
spectra of sample obtained after the mechanicalmixing processing for 2h the same spectra as the LTA
without mechanical mixing. These results suggest
nanosized LTA zeolite powders possess the good
crystallinity and also the same local structure of Al
after the mechanical mixing processing.

It is well-known that zeolites are microporous,
aluminosilicate minerals commonly used as various
commercial adsorbents. In fact, zeolites are widelyused in industry for water purification, as catalysts,
for the preparation of advanced materials and Cs
purification in nuclear reprocessing. In addition, they
purification are used to extract nitrogen from air toincrease oxygen content for both industrial andmedical purposes. However, the modification of
microstructures and morphology is needed forzeolites to expand their applications in various
industrial fields. [1-3]Recently, as one approach for the modification of
microstructures and morphology of zeolites, muchattention are paid to the synthesis of the finely
nanosized zeolite powders by the novel processing.
Various processing are attempted to synthesize thenanosized zeolite powders.
In this study, the finely nanosized zeolite powders
were synthesized by the mechanical mixing
processing with high energy grinding forces and the
local structures were investigated for finely nanosizedzeolite powders prepared by the mechanical mixing
processing.
Zeolite was synthesized by the following method.
LTA type zeolites were synthesized at 60 ͠ to 90 ͠by the soft chemical solution processing using
sodium silicate as s silicon source and aluminumsilicate as an aluminum source. The mixture of
sodium silicate and aluminum silicate was aged for
12h and then treated under hydrothermal conditionsfor 12 h to 48 h without stirring. The solid productwas filtered, frequently washed, and air-dried at room
temperature. Subsequently, the powder of this LTA
type zeolites powder was milled by the mechanicalmilling method with the high energy forces for 1h to
24 h. After the milled and ground nanosized LTA
type zeolite powder were characterized by Xrd
method and TG/DTA. The microstructures ofnanosized LTA type zeolite powder were observed byscanning electron microscopy. The local structures
around Al for the products of nanosized LTA typezeolite powder were characterized by measuring
X-ray adsorption near edge structure (XANES) at
BL2A in UVSOR.

Fig. 1. Al XANES spectra of zeolites and referencematerials.
[1] M. E. Davis and R. F. Lobo, Chem. Mater., 4(192) 756.
[2] A. Nakahira, S. Nishimura, H. Aritani, T.
Yamamoto and S. Ueda, J. Mater. Sci. 36 (2001)
1885.
[3] A. Nakahira, S. Yamamoto, M. Sato, Y. Takamatsu,
S. Misu, Y. Kawabe, Y. Nishio and H. Aritani,
UVSOR Activity Report 38 (2011) 112.

The results of XRD for synthesized zeolite

127

BL2A

XANES Evaluation of Zeolites obtained from Waste Glass
A. Nakahira1, 2, T. Tagami1, T. Shirai1, K. Kumadani1, M. Sato2 and H. Aritani3
1

Faculty of Engineering, Osaka Prefecture University, Gakuencho, Sakai 599-8531, Japan
2
Osaka Center, IMR, Tohoku University, Gakuencho, Sakai 599-8531, Japan
3
Saitama Institute of Technology, Fukaya 369-0293, Japansolid product powder were characterized bymeasuring X-ray adsorption near edge structure(XANES) at BL02A in UVSOR.
The results of XRD indicated that the products
synthesized under the optimum synthetic conditions
were identified to be a LTA type zeolite without otherphase, although some products contained LTA andother phases. These results suggest that the fine
LTA zeolite powders were successfully obtained fromwaste glass by eco-soft chemical processing with low
environmental load.
The XANES spectra of products and referencematerials are shown in Fig.1. XANES spectra ofsample obtained the same spectra as the commercial
LTA. These results suggest the fine LTA zeolite
powders possess the good crystallinity and also thesame local structure of Si as the commercial LTA(com LTA).

Normalized (a.u.)

Glass is widely used in many fields and an ideal
material for recycle, whereas it is used for new glass
container manufacture it is virtually infinitely
recyclable. The glass component in municipal waste
is usually made up of bottles, broken glassware, lightbulbs and other items. Glass makes up a large
component of household and industrial waste. The
use of the recycled glass as aggregate in concrete andothers has become popular in modern times, althoughlarge scale researches have been carried out long year.
It is expected that glass recycling is the process ofturning waste glass into usable products because theuse of recycled glass in new containers helps saveenergy. In fact, recycling of waste glass and othersources help in brick and ceramic manufacture, and it
conserves raw materials, reduces energy consumption[1-4].
As one approach of the recycle of waste glass,much attention are paid to the synthesis of zeolites for
waste glass by soft chemical processing, that is,
eco-processing withlow environmental load.
Especially, the finely zeolite powders are expected tosynthesize by the novel soft chemical eco-processing.
Various eco-soft chemical processing with low
environmental load are attempted to synthesize thenovel zeolites.
In this study, the finely LTA zeolite powders weresynthesized
by thenovel
soft
chemical
eco-processing with low environmental load and the
local structures were investigated for zeolite powders
prepared by the novel eco-soft chemical processing
with low environmental load.
Glass was prepared from the waste glass, such as awindow sheet glass. Zeolite was synthesized by thefollowing eco-soft chemical processing with low
environmental load. LTA type zeolites weresynthesized at 90 °C by the eco-recycle soft chemical
processing using sodium silicate as s silicon source
and aluminum silicate as an aluminum source. The
mixture of both sodium silicate and aluminum silicate
was aged for 24h and then heat-treated underhydrothermal conditions for 48 h with magnetic
stirring. The solid product was filtered, frequently
washed, and air-dried at room temperature.
Subsequently, the powder of the zeolite was milled by
the mortar. The solid products were evaluated by
XRD method and TG/DTA equipments. The
microstructures of the solid product powder wereobserved by scanning electron microscopy (SEM).
The local structures around Si for the products of theas-syn LTA1
as-syn LTA2
com LTA

glass

1800

1820

1840

1860

1880

Energy / eV

Fig. 1. Si XANES spectra of zeolites and referencematerials.
[1] Y. Shaoa, T. Leforta, S. Morasa and D.
Rodriguezb, Cement and Concrete Research. 30(2000) 91.
[2] A. R. Boccaccini, M. Bücker, J. Bossert and K.
Marszalek, Waste Management. 17 (1997) 39.
[3] A. Nakahira, S. Yamamoto, M. Sato, Y. Takamatsu,
S. Misu, Y. Kawabe, Y. Nishio and H. Aritani,
UVSOR Activity Report. 38 (2010) 112.
[4] M. Sato, S. Yamamoto, Y. Nishio, Y. Takamatsu,
K. Kumadani and A. Nakahira, UVSOR Activity
Report . 38 (2010) 114.

128

BL3B

Photoluminescence of Si-O-C(-H) Ceramics Derived from Polysiloxanes
T. Kawai1 and M. Narisawa2
1

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

129

transfer from the S1 to T1 states. In order to obtain
information on the PL centers of the Si-O-C(-H)
ceramics, the shape of the PL bands is carefully
analyzed.
Luminescence Intensity (arb. units)

The photoluminescence (PL) properties of Si-O-C
materials derived from organosilicon polymers havebeen attracting widespread attention in recent years [1,2]. Simple chemical composition of such Si-O-C
without special activator was promising for use inlight-emitting diode (LED) systems. Though some
Si-O-C transparent materials were synthesized bysol-gel methods or CVD methods, the precise
adjustment of resulting chemical composition was noteasy in such sol-gel or CVD methods. Recently, we
succeeded in synthesizing Si-O-C(-H) ceramics witha white appearance by using hydrogen in thepyrolysis process for silicone resin particles [3, 4]. In
this study, we have investigated the PL properties of
the Si-O-C(-H) ceramics decarbonized at different
temperatures.
Figure 1 shows the PL spectra of the Si-O-C(-H)
ceramic decarbonized at 750qC. Under excitation on
the shorter wavelength than 240 nm, the broad PL
bands are observed around 320 and 450 nm. The 450
nm bands have a long tail on the longer wavelength
region. On the other hand, the excitation on the longer
wavelength than 250 nm induces the broad PL bands
peaking at 420 nm and having a shoulder structurearound 540 nm. The shape of the PL bands indicatesthat the Si-O-C(-H) ceramic decarbonized at 750 {C
has several PL centers.
Figure 2 shows the PL spectra of the Si-O-C(-H)
ceramics decarbonized at 1100 qC under excitation at
various wavelength. The broad PL bands with a peakat 500 nm are observed. The peak energy andband-shape of the PL bands are independent of theexcitation wavelength. The fact indicates that the
decarbonization at 1100 qC leads to thetransformation to the more stable PL centers.
In the simple chemical composition of Si-O-C(-H),
the oxygen deficiency centers called ODC (II) areoften proposed as possible PL centers. The excitationand emission due to the transition between S0 (singlet
ground state) and S1 (singlet excited state) in ODC(II) occur around 5.0 eV (245 nm) and 4.3 eV (290nm), respectively [5]. Since the 320 nm PL band of
the Si-O-C(-H) ceramic decarbonized at 750 qC
appears only under excitation on the shorter
wavelength than 240 nm, the 320 nm PL band would
be attributed to the S1 ψS0 radiative transition in
ODC(II). On the other hand, the 450 nm PL bandmight come from the transition between S0 and T1(triplet excited state) in ODC (II). The appearance ofthe 450 nm PL band under excitation on the shorter
wavelength than 240 nm implies the existence of

4.0

3.0

Photon Energy (eV)2.0
Si-O-C(-H) (750°C)

RT

exc. 180 nm
220 nm
260 nm
300 nm
300

400

500
600
Wavelength (nm)700

Fig. 1. Luminescence spectra of the Si-O-C(-H)
ceramics decarbonized at 800 qC.
Luminescence Intensity (arb. units)2

3.0

Photon Energy (eV)2.0
Si-O-C(-H) (1100°C)

400

500
600
Wavelength (nm)4.0
RTexc. 180 nm
220 nm
260 nm
300 nm300

700

Fig. 2. Luminescence spectra of the Si-O-C(-H)
ceramics decarbonized at 1100 qC.
[1] A. Karakuscu et al., J. Am. Ceram. Soc. 92 (2009)2969.
[2] S. Gallis et al., Appl. Phys. Lett. 97 (2010)
081905.
[3] M. Narisawa et al., Bull. Chem. Soc. Jpn. 85(2012) 724726.
[4] M. Narisawa et al., J. Am. Ceram. Soc. 95 (2012)
3935.
[5] L. Skuja: J. Non-Cryst. Solids 239 (1998) 16.

BL3B

Intrinsic Luminescence and Host Sensitization of Rare-earth Ions inLanthanum Scandium Borate upon VUV Excitation
N. Kodama, Y. Morisawa, A. Abe, T. Sakashita and T. Takahashi
Faculty of Engineering and Resource Science, Akita University, Akita 010-8502, Japan

concentration and then decreases monotonically athigher concentrations, which indicates concentrationquenching by energy transfer between two Sm3+ ions.
For LSB:Tb, the STE emission decreases withincrease in the Tb3+ concentration (1-40 at.%), as
shown in Fig. 3(c). An increase in the 5D4ĺ7FJ
transition of Tb3+ accompanies the decrease of the
STE emission, which also indicates efficient energy
transfer from the STE to Tb3+ ions. The two decay
times of the intrinsic emissions, II (85 ns) and III (2.5ȝs), become shorter in the Tb3+-doped samples, whichis consistent with the decrease in the time-averaged
STE emission intensity with increasing Tb3+
concentration.
The Sm3+/STE or Tb3+/STE ratios gradually
increase in intensity with temperatures from 7.8 to
293 K, which suggests that the energy transfer fromthe STE to Sm3+ or Tb3+ is thermally activated.
Fig. 1(a).

Fig. 1(b).

3500

2500
Ȝex = 70 , 160 nm

2500

1500

2000

323 nm

1500

380 nm1000

1000

379 nm

253 nm0200

251 nm324 nm

500

500

0

250

300 350 400 450
Wavelength / nm

500

200

250

300 350 400
Wavelength / nm

     Fig. 2.

1.2

ὼ 3at.%Sm

0.41000 ὼ 5at.%Sm
ὼ 8at.%Sm

800 ὼ 10at.%Sm
600 ὼ 30at.%Sm

400
200

0.0
50

100

150
200
Wavelength / nm

0
200

250

250

300 350 400
Wavelength / nm

Fig. 3(b).

130

8000

3+ 4
6
Ȝex = 70 , 160 nm Sm ( G5/2ĺ HJ )J=7/2
ὼ undoped
J=9/2
ὼ 1at.%Sm

6000
4000

ὼ 3at.%Sm
ὼ 5at.%Sm
ὼ 8at.%Sm
ὼ 10at.%Sm
ὼ 30at.%Sm

J=5/2

2000
0
200

300

400 500
600
Wavelength / nm

450

500

Fig. 3(c)J=11/2

700

14000
12000
Intensity / a.u.

10000

500

Ȝex = 70 , 160 nm

1200 ὼ 1at.%Sm

Ȝem = 311 nm , 380 nm
ὼ polycrystal
ὼ single crystal

0.8

12000

450

Fig. 3(a)1400 ὼ undoped

Band gap transition of BO3-group

Intensity / a.u.

Intensity / a.u.

2.0
1.4

T = 7.8 K

2000

Intensity / a.u.

Intensity / a.u.

Ȝex = 70 , 160 nm

T = 293K

3000

Intensity / a.u.

Host sensitization of rare-earth ions (RE3+ = Gd3+
and Ce3+) has been observed in some phosphates andborates[1-3].
The
intrinsic
luminescencecharacteristics of lanthanum scandium borate,
LaSc3(BO3)4 (LSB) were investigated, in which theabsorption of vacuum ultraviolet (VUV) light by thehost may facilitate efficient energy transfer to Re3+
ions. Here, energy transfer was observed from theself-trapped exciton (STE) to Sm3+ and Tb3+ ions.
Emission and excitation spectra were measured attemperatures in the range of 7.8-293 K using theUVSOR facility. Figures 1(a) and (b) show emission
spectra observed for an undoped LSB polycrystal
with various excitation wavelengths at 293 and 7.8 K.
The emission spectrum with VUV excitation at 70
and 160 nm consists of three intrinsic broad bands at
253, 323, and 379 nm. The feature of the emission
spectrum for a LSB single crystal is the same as thatfor the polycrystal. Figure 2 shows excitation spectra
at 293 K for both the LSB polycrystal and single
crystal, which consist of fairly broad bands between
50 and 215 nm. The band edge of the single crystal isobserved around 215 nm and the broad bands are
assigned as bandgap excitations or moleculartransitions of the BO33- group. Both the single crystal
and polycrystal spectra exhibit the same features witha sharp rise below 215 nm.
The three intrinsic emissions can be assigned torecombination of the distinct STEs (I, II, and III) that
could be associated with bandgap excitations ormolecular transitions of the BO33- group. The decay
times of the emissions, I (253 nm), II (323 nm) andIII (379 nm) are less than 10 ns, 115 ns, and 3.4 ȝs,
respectively. By analogy with the STE states in alkali
halides [4], the I (fast) and II (fast) STE emissions
originate from singlet transitions on center andoff-center configurations, respectively, whereas the
III (slow) STE emission is suggested to originate
from a triplet transition on the off-center
configuration. The emission spectra of undoped LSB,
Sm3+-doped LSB polycrystals (LSB:Sm), andTb3+-doped LSB polycrystals (LSB:TB) excited at 70
and 160 nm at 296 K are shown in Figs. 3(a)-(c).
For LSB:Sm, the intrinsic emissions (STE
emission) decrease with increase in the Sm3+
concentration (1-30 at.%) (Fig. 3(a)) and this isaccompanied by the appearance of intense emissions
from the 4G5/2ĺ6HJ transitions of Sm3+ (Fig. 3(b)).
This suggests efficient energy transfer from the STE
to Sm3+ ions. The Sm3+ emission intensity passes
through a maximum at 3-5 at.% with increasing Sm3+

10000
8000
6000
4000

Ȝex = 70 , 160nm
䋭 undoped
䋭 1at.%Tb䋭 5at.%Tb䋭 10at.%Tb䋭 20at.%Tb䋭 30at.%Tb䋭 40at.%Tb

J=5

T = 293K

J=6
J=4

2000
0
200

Tb3+ (5 D4ĺ7FJ)

300

J=3

600
700
400 500
Wavelength / nm
Fig. 1.

Emission spectra observed in undoped LSB excited at 70 nm and
160 nm at 293 K (a) and 7.8 K (b).
Fig. 2. Excitation spectra of undoped LSB poly- and single-crystal.
Fig. 3. Emission spectra of undoped LSB (a), LSB:Sm (b) and
LSB:Tb excited at 70 nm and 160 nm.
[1] S.P. Feofilov et al., Phys.Rev.B. 74 (2006) 085101.
[2] S.P. Feofilov et al., J. Lumin. 125 (2007) 80.
[3] N. Kodama et al., Abstract of 15th Int. Conf. on Defects in
Insulating Materials (2010) A90.
[4] K.Kan’no et al., Pure & Appl. Chem. 69 (1997) 1227.

BL3B

Effect of Ion Implantation on the Crystallinity of YAlO3 Single Crystal
Y. Horii, T. Morimoto and Y. Ohki
Waseda University, 3-4-1 Ohkubo, Shinjuku-ku, Tokyo 169-8555, Japan
500
400

Absorbance(10 cm )

200
100
0
30

32

34
36
2T [deg]38

40

PL Intensity [arb. unit]

Fig. 2. In-plane XRD patterns observed before andafter the implantation of P+ or B+ ions.

A

A: Before

B
B: P+ implantedC
C: B+ implanted

1.5

1.6
1.7
Detected Photon Energy [eV]1.8

PL Intensity [arb. unit](a)A
B and C
(overlapped)

2

3
4
5
Detected Photon Energy [eV]6

PL Intensity [arb. unit](b)

Beforebeforeimplanted
PP+ ionirradiatedB
implanted
B+ ionirradiatedA
B and C
(overlapped)

4

2

4.5
5
5.5
6
6.5
Detected Photon Energy [eV]7(c)Fig. 3. PL spectra observed before and after P+ or B+ ion
implantation, induced by 6.4 eV (a), 7.7 eV (b) and
8.2 eV (c) photons.

10
3(10 cm )+ion irradiatedimplanted
BB
15
-2

300

4
3

Beforebeforeion irradiated
P+Pimplanted
15-2

Intensity [arb. unit]

Yttrium aluminate (YAlO3) is a promising
candidate for a gate insulator in advanced MOS
devices. However, ion implantation, which is usuallyrequired after the deposition of a gate insulator,
would cause damage to the gate insulator. Therefore,
the effects induced by implantation of ions were
examined by measuring optical absorption, photoluminescence (PL) and X-ray diffraction (XRD)
spectra.
The samples examined are commercially available
YAlO3 (100) single crystals in the shape of a platewith a thickness of 0.5 mm. Positive phosphorus (P +)or boron (B+) ions were implanted to the samples at afluence of 1.0×1015 cm-2. Before and after the ion
implantation, optical absorption, in-plane XRD and
PL were measured.
Optical absorption increases significantly atenergies slightly lower than the gap energy as shownin Fig. 1. The XRD peak shown in Fig. 2 shifts after
the implantation of P+ and B+ ions. Furthermore, its
peak height becomes smaller. Figure 3 shows PL
spectra induced by 6.4, 7.7 and 8.2 eV photons,
where curve A represents PL spectra observed before
ion implantation, while curves B and C represent thoseobserved after the implantation of P+ and B+ ions,
respectively. Among several photoluminescencebands observable in YAlO3, two bands due to Cr3+ ataround 1.70 eV induced by 6.4 eV photons [1],another one due to self-trapped exciton at 4.2 eV
induced by 7.7 eV photons [2], and another one due
to an antisite defect at 5.5 eV induced by 8.2 eV
photons [3] become smaller or disappear by the ion
implantation. These three PLs appear only when thesample is crystalline [1, 3, 4]. All these facts indicatethat the appearance of localized electronic states in theband gap and the degradation of crystallinity are
induced by the ion implantation.

4

5
6
7
Photon Energy [eV]8[1] E. Hirata et al., Jpn. J. Appl. Phys. 49 (2010) 091102.
[2] Ch. Lushchik et al., J. Phys. Condens. Matter 6(1994) 11177.
[3] Y. V. Zorenko et al., Opt. Spectrosc. 96 (2004) 536.
[4] M. Murakami et al., J. Appl. Phys. Lett. 78 (2001) 2664.

Fig. 1. Absorption spectra of YAlO3 single crystals
implanted with P+ or B+ ions.

131

BL3B

Optical Spectroscopy of Eu3+ in GaN Codoped with Mg and Eu
M. Yamaga, S. Tsuda and K. Takagi
Department of Mathematical and Design Engineering, Gifu University,Gifu 501-1193, Japan

Europium doped GaN is a candidate of red
light-emitting diodes (red LEDs). Recent optical
spectroscopy of Eu3+ in GaN(Mg) has revealed theelectronic structure of Eu3+ associated with defects
and temperature-dependence hysteresis of the Eu3+
luminescence [1]. However, the quantum efficiencyof the Eu3+ luminescence was less than 1 %. In
order to improve the quantum efficiency, it is very
important to examine the relaxation process from theband-to-band excitation to the Eu3+ excited state.
GaN films, grown by hydride vapor phase epitaxy(HVPE) on (0001) sapphire substrates and dopedwith Mg (p-type) were implanted with 300 keV Euions. The GaN(Mg):Eu sample was supplied by
Prof. K. P. O’Donnell at University of Strathclyde,
UK.
Figure 1 shows the absorption spectrum of
GaN(Mg):Eu at 300 K. The sharp step around 360 nm
is corresponding to the band-edge energy of GaN.
The interference fringes due to the thin film structure
were observed in the spectrum.
The band-to-band excitation below 360 nm
produces fairly strong Eu3+ luminescence lines around
620 nm, and three broad bands with peaks at 440, 540
and 600 nm as shown in Fig. 2. The Eu3+
luminescence lines decrease drastically beyond 360
nm and the two broad bands at 440 and 540 nm
appeared clearly. These broad bands may be assigned
to localized centers, for example, donors, acceptors ordefects.
Relaxation to Eu3+ and emitting processes of Eu3+
are strongly associated with the rise and fallingcomponents of the decay curves of the Eu 3+
luminescence, respectively. The decay curve of theEu 3+ luminescence in GaN:Eu can fit a single
exponential function with the decay time of 0.23 ms[2]. However, the decay curve from GaN(Mg):Eu
d o e s n o t fi t t o a s i n g l e e x p o n e n t i a l , b u t
multi-exponential, at least three components.
Otherwise, the decay curve fits to a sum of a single
exponential function and power function of t n (n=-1),
assuming that there are two different relaxation
processes to Eu 3+; one is rapid relaxation fromband-to-band excitation immediately to the Eu3+
excited states; the other is relaxation to Eu3+ via
donor-acceptor recombination process. Namely, thelong decay time (larger than 2 ms) of Eu3+ may be
due to fairly long feeding time. The possible model isshallowly trapped centers, for example, donors or
acceptors trapping weakly electrons or holes,
respectively. These bands with the double peaks at
440 nm and 540 nm in Fig. 2 may be associated withsplitting of the acceptor bands consisting of Mg andN wavefunctions.

GaN:Eu, Mg
5Absorbance

4
T=300 K

3
2
1
0
200

300
400
500
Wavelength (nm)600

Fig. 1. The absorption spectrum of GaN(Mg):Eu.

T=8 K

3

Intensity (arb. units)

Oex=70 nm

2
Oex=280 nm

1
Oex=360 nm

0
400

500
600
700
Wavelength (nm)800Fig. 2. Photoluminescence of Eu3+ in GaN(Mg):Eu
excited with various wavelengths in the band-to-band
transitions.
[1] K. P. O’Donnell, P. R. Edwards, R. W. Martin, K.
Lorenz, V. Darakchieva, A. Alves and M. Bockowski,
to be published (2013).
[2] I. S. Roqan, K. P. O’Donnell, R. W. Martin, P. R.
Edwards, S. F. Song, A. Alves and M. Bockowski,
Phys. Rev. B 81 (2010) 085209.

132

BL3B

Photoluminescence due to Intraionic 4f-4f Transitions of Gd3+ Ions in
Gd3(Al,Ga)5O12 Single Crystals
S. Ishii1, M. Kitaura1 K. Kamada2, A. Ohnishi1 and M. Sasaki1
1Dept. of Physics, Faculty of Science, Yamagata University, 1-4-12 Kojirakawa, Yamagata 990-8560, Japan
2
Materials laboratory of Furukawa Co. Ltd, 1-25-13 Kannondai, Tsukuba 305-0856, Japan

Resonant energy transfer between impurity ionshas been studied in the viewpoint of fundamentals
and applications in inorganic phosphor materials [1].
The control of this process is one of conventional
methods to enhance the luminescence yield in thedevelopment of phosphor materials. The main work is
to hunt for better combinations of donor and acceptor
ions on the basis of an overlap of absorption andluminescence spectra. The pair of Ce3+ and Gd3+ ionsis known to be one of such combinations. In mostionic crystals, the absorption due to 4f-5d transitionsof Ce3+ ions and the emission due to 4f-4f transitionsof Gd3+ ions are located in the near ultraviolet range[2]. Since they have a strong overlap with eachother, resonant energy transfer from Gd3+ to Ce3+ ionsare possible in crystals including them.
Gd3Ga3Al2O12 (GAGG) belongs to the cubic
system with garnet structure. Ce3+ doped GAGG
(Ce:GAGG) attracts much attention as an inorganic
scintillator for the detection of x- and Ȗ–rays, because
it has scintillation properties superior to other opticalcrystals [3]. As mentioned above, since the
Ce:GAGG crystal contains Gd3+ and Ce3+ ions,
there is the possibility that host sensitization of the
Ce3+ 5d-4f emission occurs by the energy transferfrom Gd3+ to Ce3+ ions. An investigation of the hostsensitization process is of great significance for therealization of higher luminescence yield and shorter
decay time, so it is worthwhile to investigate
photoluminescence (PL) properties of GAGG crystals
with vacuum ultraviolet (VUV) photons. In thepresent study, we have measured PL and
PL-excitation (PLE) spectra of GAGG crystals withuse of synchrotron radiation. Experiment was carriedout at the BL3B beam line of UVSOR. Here, wereport a part of results in our experiment.
Figure 1 shows the absorption spectrum of GAGG
at 6K. Narrow peaks are observed around 4 –5 eV.
Similar peaks have been observed for Gd3+ activatedphosphors [2]. We thus assign them to 4f-4ftransitions of Gd3+ ions. The strong absorption occurs
at 5.96 eV. This corresponds to the fundamental
absorption edge of GAGG.
Figure 2 shows PL and PLE spectra of a GAGG
crystal. The PL spectrum was obtained at 6 K under
excitation at 6.52 eV. A sharp peak appears at 3.92
eV in the PL spectrum. The PLE spectrum for the
3.92 eV peak exhibits narrow peaks, due to 4f-4ftransitions of Gd3+ ions, around 4 – 5 eV. The 3.92
eV PL peak is also assigned to 4f-4f transitions of
Gd3+ ions, because this peak is efficiently excited in

the same position as narrow peaks in Fig. 1. As thephoton energy is increased, the excitation for the 3.92
eV peak occurs at the fundamental absorption edge at
5.96 eV again, and it continues for the high-energy
side above 5.96 eV. The intensity of the 3.92 eV peakis gradually increased around 20 eV, and reaches
twice at 30 eV. This is interpreted as an appearance ofthe multiplication of electron-hole pairs in GAGG.

Fig. 1. Absorption spectra of a GAGG single crystal
observed at 6 K.

Fig. 2. Emission (red line) and excitation (black line)
spectra of a GAGG single crystal observed at 6 K.
The emission spectrum was observed under excitation
at 6.52 eV. The excitation spectrum was measured forthe 3.92 eV emission peak.
[1] D. L. Dexter, J. Chem. Phys. 21 (1953) 836.
[2] T. Kano, Phosphor Handbook, CRC Press (Boca
Raton) 2nd ed., Chap.3, Sect.3 .
[3] K. Kamada et al., J. Cryst. Growth 352 (2012) 88.

133

BL4B

Nitrogen Doped TiO2 Photocatalyst Prepared by Low Energy N+
Implantation
T. Yoshida1, E. Kuda2 and H.Yoshida2

1

EcoTopia Science Institute, Nagoya University, Nagoya 464-8603, Japan
Department of Applied Chemistry, Graduate School of Engineering, Nagoya University, Nagoya 464-8603,
Japan

2

Figure 2 shows N K-edge XANES spectra of the
N+-implanted TiO2 and TiN samples. Common
XANES features in (b) and (d) suggest that N in thesample implanted by 3 x 1021 m-2 (TiO2-N(3)) is in achemical environment similar to that in TiN.
Careful observation led us to notice that two peaks
around 400 eV for the catalyst sample shift to lower
energy side than those for the TiN sample, which waswell reproduced by the theoretical prediction using
FEFF code when N occupies one of the O sites ofTiO2.
On the other hand, the XANES spectra of thesample implanted by 1 x 1021 m-2 (TiO2-N(1)) andthat implanted by 3 x 1021 m-2 followed by heating at
573 K (TiO2-N(3)-H) showed a distinct single peakaround 401 eV. This peak could be empiricallyattributed to formation of the species such as N
bonds near the surface [3], and which also indicatesthat the inactive N-O species dominates in TiO2-N(1)and the active nitrogen formed in TiO2-N(3) changed
to the inactive N-O species by oxidation at 573 K.
These results are in conflict with our previous study,
in which the photocatalytically active nitrogen waspreferentially produced by 50keV N+ implantation
with the lower nitrogen concentration.

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. Asahi et al.
reported that the substitutional doping of N into TiO2
contributes to band gap narrowing to provide
visible-light response [1].
In our previous study, we prepared visible-light
response TiO2 photocatalysts by the implantation of
50 keV N+ ions [2]. We found that the photocatalytic
active nitrogen species, i.e., nitrogen substituting for
the O sites in TiO2, are preferentially formed in thesamples implanted with a low N+ fluence (< 3 x 1021
m-2). In the samples prepared with a higher N+
fluence, the inactive N-O species generated to reduce
photocatalytic activity. In the present study, a low
energy, 5 keV N+ implantation was performed inorder to inject nitrogen atoms near the surface ofTiO2.
The samples used in this study were TiO2 (1 0 0)
single crystals (5 x 5 x 0.5 mm3). Mass analyzed 10
keV N2+ (5 keV N+) ions were injected into thesamples at room temperature, perpendicular to thesample surface. As calculated by a Monte Carlosimulation using SRIM code (Fig. 1), the implanted
nitrogen distributes up to ca. 25 nm, peaking around
10 nm in depth from the surface, and they would
increase monotonously with N+ fluence. Thus, the5keV N+ implantation enables to inject nitrogenatoms in the shallower region compared with the 50
keV N+ implantation in our previous study.
N K-edge XANES spectra of the samples were
measured at the BL-4B. Data were recorded at room
temperature in total electron yield mode.

Absorbance (arb. unit)-2

5keV N
6

c)

b)

+

50keV N

+

a)

-1

-1

Atoms ( nm ion )

8x10

d)

4

395

400

405

410

X-ray energy (eV)415

Fig. 2. N K-edge XANES spectra of (a) N+-implanted
at 1 x 1021 m-2, (b) N+-implanted at 3 x 1021 m-2, (c)N+-implanted at 3 x 1021 m-2 followed by heating at
573 K for 2 h, and (d) TiN.

2
0
0

50

100
Depth (nm)

150

200

[1] R Asahi et al., Science 293 (2001) 269.
[2] T. Yoshida et al., Mater. Trans. 48 (2007) 2580.
[3] J-H. Wang et al., Anal. Chim. Acta, 476 (2003)
93.

Fig. 1. Depth distributions of implanted nitrogen per
incident N+ ion calculated by SRIM code for 5 and 50
keV N+ implantations.

134

BL4B

Surface-Structure Change in Lithium-Ion Battery Using Anti-Fluorite Type
Positive ElectrodeT. Okumura, M. Shikano and H. Kobayashi
National Institute of Advanced Industrial Science and Technology (AIST), Ikeda, Osaka 563-8577, Japan
battery (#1 to #4). Thus, the surface Li2O also reacts
as well as the oxidation of internal Li5FeO4 phase.
For example, Li2O2 is one of the candidates of
product compounds. After that, the Li2O re-formed
after discharging battery (#5). Thus, theelectrochemical reaction of the surface Li2O phasewould also contributed on the charge / discharge
properties in the battery using anti-fluorite Li5FeO4positive electrode.

Fe-based materials as positive electrodes oflithium-ion batteries have been paid much attention todue to their advantages of being low cost andenvironmentally benign. For example, LiFePO4 isused as the positive electrode material of commercial
lithium-ion batteries in spite of relatively lower
capacity. To improve battery performance, thepositive electrode materials of high capacity are
required. Defect anti-fluorite-type structure Li5FeO4
is one of the candidate because its theoretical capacity
reaches ~694 mAh g-1 (four Li in a formula unit) [1,2]. In previous work, the relatively rechargeable
capacity of Li5FeO4 could be obtained after firstcharging until two Li per Li5FeO4 formula unit [1].
However, the changes in electronic and crystal
structure of Li5FeO4 during the electrochemical
reaction have not been clarified yet. Recently, we
have analyzed the structural changes of Li5FeO4
during charging process using XRD, hard-XAFS and
7
Li MAS NMR. In this study, the surface-structurechange during electrochemical reaction was analyzed
by O K-edge XAFS spectra using total-electron-yield(TEY) mode.
Li5FeO4 was synthesized by solid-state reaction.
The positive electrodes were made from a mixture
containing 60 wt% Li5FeO4, 32 wt% acetylene blackand 8 wt% polytetrafluoroethylene, which were
casted and pressed on nickel mesh. Coin cells wereassembled using metallic lithium foil as a negativeelectrode. A 1 mol dm-3 ethylene carbonate/dimethylcarbonate solution of LiPF6 was used as electrolyte.
Electrochemical test was carried out at a constantcurrent density of 1/30 C. For measuring XAFS
spectra, samples (#1~5) were obtained from the cells
after charging or discharging, as shown in Fig. 1.
The O K-edge spectra were shown in Fig. 2+.
3d-transition metal (TM) oxides generally have thetwo peaks around 532-538 eV since O 2p orbital
hybridized with TM 3d t2g and eg as shown in the
spectrum of Fe2O3. However, this specific shape
could not be observed in the all samples in this study(#1~5). This result indicates that the surface of thesamples is covered with another phase containing
lithium salts, and iron oxide phase could not be
detected by TEY mode. To identify the surface phase,the spectra of various lithium salts were measured as
partially shown in Fig. 2. The spectrum of pristine
sample (#1) is similar to that of Li2O, that is, thesurface of #1 would be covered with Li2O since thelattice parameters of anti-fluorite-type structure Li2O
is close to those of Li5FeO4. The shapes of the spectra
around 535-538 eV were changed during charging

Voltage / V vs. Li/Li+

4.0
3.5

#4
#3

#2

3.0
2.5 #1
2.0

ChargeDischarge

#5

1.5
0

50

100

150

200

250

300

350

Capacity / mAh g-1

Fig. 1. First cycle of lithium-metal half cell using
Li5FeO4 positive electrode. Circles indicate thesamples measured XANES spectra.

#5
Normalized Intensity / a.u.

#4

㩷

#3
#2
#1
Li2O2
Li2O
Fe2O3

㪌㪊㪇

㪌㪊㪌

㪌㪋㪇

㪌㪋㪌

㪌㪌㪇

㪌㪌㪌

㪌㪍㪇

㪌㪍㪌

Energy / eV

Fig. 2. O K-edge XANES spectra measured bytotal-electron-yield mode.
[1] A. Hirano et al., Solid State Ionics 176 (2005)
2777.
[2] C.S. Johnson et al., Chem. Mater 22 (2010) 1263.

135

BL5U

Ultrahigh-Resolution Direct Observation of Mini-bands Formed in
InGaAs/AlGaAs Superlattice
D. Shimura1, F. Ichihashi1, K. Nishitani1, S. Harada1, M. Kuwahara2, T. Ito1,
M. Matsunami3, S. Kimura3, M. Tagawa1 and T. Ujihara1
1

Department of Materials Science and Engineering, Nagoya University, Nagoya 464-8603, Japan2
Department of Applied Physics, Nagoya University, Nagoya 464-8603, Japan
3
UVSOR Facility, Institute for Molecular Science, Okazaki 444-8585, Japan

Intermediate-band solar cells have attracted great
attention for their potential to achieve 74.6 %efficiency over the Shockley-Queisser limit of asingle gap solar cell [1-3]. The concept of theintermediate-band solar cell is to increase
photocurrent by using electrons excited to a
conduction band via the intermediate-bands.
Mini-bands formed in superlattice structures are often
utilized for the formation of the intermediate-bands.
Indeed, recently, it has been demonstrated thatquantum-well superlattice solar cells induced two
step photon absorption and photocurrent generation atroom temperature [4]. In addition, these reports
indicated that the transport of the carriers in themini-bands was essential in order to the photocurrent
generation based on two step absorption. Therefore it
is very important to figure out the mini-bandstructures which influence the transport properties
and the conversion efficiency. In this study, we
directly observed the dispersions of mini-bands
formed in InGaAs/AlGaAs quantum-well superlatticeby angle-resolved photoemission spectroscopy(ARPES) with synchrotron radiation.
The structure of the specimen is shown in Fig. 1.
The superlattice was grown by a low-pressure
metalorganic vapor phase epitaxy. On the evaluationof the mini-band structure, normal-emission andin-plane
ARPES
measurementsinthevacuum-ultraviolet region were performed at BL5U
of the UVSOR-III storage ring. The measurements
were carried out at 10 K under the ultrahigh vacuum
atmosphere (less than 1×10-8 Pa).
The ARPES intensity mapping image of the
X-valley of InGaAs which was originated from thesurface of the superlattice specimen is shown in Fig.
2. The zero energy indicates the valence-band
maximum and the zero wave vector indicates the
X-valley. The valence band of InGaAs and As 3d corelevel-derived peaks were observed. In addition, it wasnoticed that the peaks which were not existed in the
InGaAs bulk were observed ranging from -2.0 to 0
eV of energy. The yellow circles indicate themini-band-derived peaks of the second orderdifferential, which indicate the mini-band structures
formed in the superlattice structure. These peaks weregood consistent with the theoretical values calculated
by Kronig-Penney model, especially the heavy holemini-bands.
By ARPES measurement at high resolution, we canexperimentally evaluate the mini-band structures.

Fig. 1. Details of the structures of the specimen.

Fig. 2. The ARPES intensity mapping image of
InGaAs/AlGaAs superlattice. The yellow circles
indicate the mini-band-derived peaks of the second
order differential.
[1] A. Luque and A. Martí, Phys. Rev. Lett. 78 (1997)
5014.
[2] T. Nozawa and Y. Arakawa, Appl. Phys. Lett. 98(2011) 171108.
[3] W. Shockley and H. J. Queisser, J. Appl. Phys.
32 (1961) 510.
[4] M. Sugiyama, Y. Wang, K. Watanabe, T. Morioka,
Y. Okada and Y. Nakano, IEEE J. Photovoltaics 2(2012) 298.

136

BL5B

Characterization of Amorphous Chalcogenide Films by Vacuum Ultraviolet
Absorption Spectroscopy
K. Hayashi
Department of Electrical and Electronic Engineering, Gifu University, Gifu 501-1193, Japan

amorphous chalcogenide films. As shown in the
figure, the As 3d core level absorption peaks of the
as-deposited film show significant differences inshapes and positions in comparison with thoseobtained for the well-annealed film. It can explain the
difference by the reports [5] that many dangling
bonds and homo bonds (As-As and S-S bonds etc.)
exist in the as-deposited amorphous films. In the next
step, the structural changes by the irradiation of the
bandgap light will be investigated by the VUVabsorption spectroscopy.

Amorphous semiconductor materials are very
expected as a potential material for optoelectronic
devices because these materials are very sensitive tothe light and show a variety of photoinduced
phenomena [1-3]. For the device application, it isnecessary to sufficiently understand the physicalproperty of these materials. Although a large numberof studies have been done on the photoinduced
phenomena of these materials, little is known about
the details of these mechanisms. These phenomena
were studied by exciting outer core electrons with theirradiation of light with the energy corresponding tothe optical bandgap or sub-bandgap. The interest has
been attracted for the change of the optical properties
in the energy region of the visible light. We areinteresting for the changes of the optical properties in
the higher energy region. To obtain a wide knowledgeof the photoinduced phenomena, it is necessary toinvestigate to the photoinduced effects on wide
energy region. In previous reports, we reported thephotoinduced change at the VUV reflection and total
photoelectron yield spectra of the amorphouschalcogenide thin films induced by bandgap light.
Recently, we became able to measure directly the
VUV absorption spectra of amorphous thin films [4].
In this paper we report results on the study of thedifference of the structure between the as-deposited
and well-annealed films by the VUV absorption
spectroscopy.
Samples used for the measurement of the VUVabsorption spectrum were amorphous chalcogenide(a-As2S3 and a-As2Se3) semiconductor thin films
prepared onto aluminum thin films by conventional
evaporation technique. Typical thickness of thesamples and the aluminum films were around 180nm
and 100nm respectively. The aluminum film was also
used in order to eliminate the higher order light fromthe monochromator in the VUV region. The
measurements were carried out at room temperature
at the BL5B beam line of the UVSOR facility of theInstitute for Molecular Science. And the spectrum
was measured by using the silicon photodiode as adetector. A pinhole of 1.5mm in a diameter wasinserted between the monochromator and sample toremove stray light. The intensity of the VUV light
was monitored by measuring the TPEY of a gold
mesh. The positions of the core levels for the samples
were calibrated by referencing to the 2p core levelabsorption peak of the aluminum film.
Figure 1 shows VUV absorption spectra of As 3dcore level for the as-deposited and the well-annealed

Fig. 1. VUV absorption spectra of As 3d core level
for as-deposited and well-annealed amorphouschalcogenide films.

[1] Ke. Tanaka, Rev. Solid State Sci. 4 (1990) 641.
[2] K. Shimakawa, A.Kolobov, and S. R. Elliott, Adv.
Phys. 44 (1995) 475.
[3] Ke. Tanaka, Encyclopedia of Nanoscience andNanotechnology, 7 (2004) 629.
[4] K. Hayashi, UVSOR Activity Report 34 (2007)
79.
[5] Ka. Tanaka, Researches of the Electrical
Laboratory, 779 (1977) 35.

137

BL6U

Highly Ordered Multilayer Films of Fragile Bio-Molecules - The Band
Dispersion in Adenosine and Guanosine Assemblies
R. Friedlein1, S. Z. N. Demon1, H. Yamane2 and N. Kosugi2
1

School of Materials Science, Japan Advanced Institute of Science and Technology, Nomi, Ishikawa 923-1292,
Japan
2
Dept. of Photo-Molecular Science, Institute for Molecular Science, Okazaki 444-8585, Japanthe formation of extended electronic states as requiredforaband-likechargetransport
alongquasi-one-dimensional assemblies.
This work was supported by the Joint Studies
Program (No. 608, 2012-2013) of the Institute for
Molecular Science.

Guanosine and adenosine are purine nucleosides
that are composed of guanine or adenine units,
respectively, attached to a ribose sugar. Their
molecular structure is depicted in Figure 1 a and b,
respectively. As the basic units of RNA and DNA, as
well as in the case of adenosine-triphosphate (ATP)as the primary energy source for most bio-molecular
processes, they are in the center of life itself.
Since these fragile bio-molecules decompose atelevated temperatures, however, it has been
intrinsically difficult to investigate the electronic
properties of their assemblies using surface science
techniques. In the present study, thin multilayer filmsof adenosine, guanosine as well as of the Na salt ofguanosine triphosphate (Fig. 1c) have been processedat about 100 °C onto graphite (0001) surfaces byspray deposition from solution inside a home-based
ultra-high vacuum system. Films were then exposedin situ to water vapor before being transferred on air
to the experimental set-up at BL 6U.
Photon-energy dependent photoelectron spectra ofabout an about 2-4 nm thick adenosine film wereobtained at normal emission with photon energies in
between 43 and 130 eV. As shown in Fig. 1d), a
pronounced feature related to the highest occupied
molecular orbital (HOMO), centered at a binding
energy of about 2.8 eV, is composed of at least twobands, denoted “A” and “B”. While the dispersion of
the one at lower energy, “A”, can be traced over
several periods, “B” appears to be non-dispersing. As
shown recently for guanine and adenine thin films [1,2], the presence of dispersing S electronic statesindicates the formation of homeotropically aligned,
highly-ordered columnar assemblies. A more precise
evaluation of the band dispersion of “A” with a
one-dimensional tight-binding fit reveals the band
width of 384 r 20 meV and the length of the repeat
unit of 6.78 r 0.03 Å which is consistent with two
inequivalent molecules per unit cell and the size ofthe overlap integral of 96 r 5 meV. Note that thestacking is distinguishably different from that inadenine assemblies with its four molecules per unit
cell [2] and also from DNA and RNA spiralconfigurations.
Guanosine triphosphate within films of its Na salt
and guanosine have been found to form thin films
which are structurally and electronically similar to
those of adenosine. Our results show that the presenceof water molecules within the crystal does not prevent

Fig. 1. Photon-energy dependent photoelectron
valence band spectra of the hydrated adenosine
multilayer films. “A” and “B” denote features related
to HOMO-derived electronic states.
[1] R. Friedlein, Y. Wang, A. Fleurence, F. Bussolotti,
Y. Ogata and Y. Yamada-Takamura, J. Am. Chem.
Soc. 132 (2010) 12808.
[2] Y. Wang, A. Fleurence, Y. Yamada-Takamura and
R. Friedlein, Chem. Commun. 47 (2011) 12349.

138

BL6B

Far-Infrared Reflective Study of Alkali Niobate Ceramics
T. Nishi, K. Tsuchida, K. Kato, N. Kato, M. Watanabe and K. Kakimoto
Graduate School of Engineering, Nagoya Institute of Technology, Nagoya 466-8555, Japan
1

The Lix(Na0.5K0.5)1-xNbO3 ceramic samples were
produced by a solid-state reaction method. The
surfaces of the LNKN ceramics were polished for IR
reflectivity measurement. The reflectivity spectra wererecorded in the range of 20 – 8000 cm-1 at room
temperature using a synchrotron radiation source. The
spectra were corrected by using a Michelsoninterferometer (Bruker, IFS66v). The IR reflectivity
spectra were fitted by :(1)

Reflectivity

0.6
0.4
0.2
0

1

100

8000

1000

Wavenumber (cm-1)
Fig.
1.
Infraredreflectancespectraof
Li0.02(Na0.5K0.5)0.98NbO3 ceramics at R.T. and red line
represented fitting result using equations (1) and (2).
The Pm-derived peak60

İ// (Li = 2 mol%)
Bmm2
Bmm2 + Pm

40

20

0

where, İ is the high-frequency dielectric constant,Ȧp and Ȧo are the plasma and longitudinal frequencies,and Ȗ is the damping constant. The complex dielectric
function is related to the reflectivity spectrum by theequation,
R = | ¥䓔 (Ȧ) – 1 / ¥䓔 (Ȧ) + 1|2

Fitting result(using eq. 1 and 2)

0.8

0

100
300
200
Wavenumber (cm-1)

400

Fig. 2. Variation of imaginary part of permittivity withwavenumber for Li0.02(Na0.5K0.5)0.98NbO3 ceramic.

150

The Pm-derived peak
Li = 6 mol%
Li = 4 mol%
Li = 2 mol%
Li = 0 mol%(2)

Figure 1 shows the reflectivity spectra of LNKN
ceramics (x = 2 mol %) at room temperature. Figure 2
shows the complex permittivity (imaginary part)
calculated from the fitting results (using eq. (1) and(2)). A weak phonon peak was observed around 100
cm-1. This peak cannot be assigned with the only
orthorhombic Bmm2 which represents NKN crystal
structure (x = 0 mol %). Therefore, it can be found that
Li substitution led to another phase in the NKN crystal
structure. From XRD analysis, it has been obtainedthat LNKN ceramics possess the distorted monoclinic
Pm phase. Therefore, this peak could be assigned withthe mixed phase structure including monoclinic Pm
phase. Figure 3 shows the phonon analysis results of
LNKN ceramics by using FT-IR data. The phonon
intensity around 100 cm-1 are increased with Liconcentration. This result suggests that the proportionof Pm phase in LNKN ceramics increases withincreasing the amount of Li content.

139

100
Bmm2-derivedİ//䓔 (Ȧ) = 䓔  + n Ȧ2pn /Ȧ2on – Ȧ2 - nȖnȦExperimental(Li = 2 mol%)İ//Alkali niobate ceramics are one of the promising
candidates for lead-free piezoelectric materials.
Especially, Li-doped (Na,K)NbO3 (LNKN) solid
solution shows excellent piezoelectric properties.
Although it has been expected that the distortion oflattice structure due to the small ionic diameter ofsubstited-Li ion would contribute to the good
ferroelectric properties, the particular reason for this is
still unclear. In our previous work, we evaluated the
local lattice structure of alkali niobate ceramics using
FT-IR analysis. The present work reports the
spectroscopic characterization of LNKN ceramics
using infrared (IR) technique.

50

0

0

100 200 300 400
Wavenumber (cm-1)

Fig. 3. Dependence of the imaginary part of
permittivity on Li content in LNKN ceramics.
In the LNKN phonon spectra, more phonon modes
are observed due to the existence of Li in the lattice.
It is concluded that Li substation distorted the unit cellof the NKN ceramics and introduced the monoclinic
Pm phase.

BL7B

VUV Reflection Spectroscopy of Mg-Ni Alloy Thin Films
S. Sato, R. Matsushima, K. Hirasawa, M. Anzai and H. Matsumoto
School of Science & Technology, Meiji University, Kawasaki 214-8571, Japanchanges were observed in the Mg and Ni films before
and after the films were exposed to H2. However,some strong peaks were observed in the region
beyond 10 eV in the spectrum of the MgNi alloy
thin film after it was exposed to H2. The positionsof these peaks were 11, 14, 21, and 32 eV. The
hydrogenation reaction of the Mg and MgNi alloys
is represented as follows [3]:Mg2Ni + 2H2 ĺ Mg2NiH4
Mg + H2 ĺ MgH2
The peaks beyond 10 eV were assigned Mg2NiH4,
which is the hydride form of the MgNi alloy,
because no changes were observed in the spectra of
the Mg film. The transparent hydride of the MgNialloy film exhibited a large energy gap andfundamental absorption in the VUV region.

Magnesium nickel (MgNi) alloy are well-known
metallic alloys for hydrogen storage, and they exhibit
switching mirror behavior from metal to transparent
films. This behavior is due to the fact that thehydrides of the MgNi alloys are transparent. The
alternant hydrogenation and dehydrogenation
reactions depend on their switching behavior.
Therefore, the optical properties of these alloys are
significantly changed by hydrogenation. However,there have been no reports about the opticalproperties of these MgNi alloy materials in thevacuum ultraviolet (VUV) region. The transparent
hydride forms of the MgNi alloy films exhibit alarge energy gap and fundamental absorption in the
VUV region. VUV reflection spectroscopy is theone of the most important analysis methods forverifying the fundamental absorption properties ofsuch transparent films [1, 2]. On the basis of VUVreflection spectroscopy, the refraction spectra of thehydride and dehydride forms of MgNi alloy thinfilms prepared by the radio-frequency (RF)magnetron sputtering method were systematicallyinvestigated.
A switching MgNi alloy film was prepared by the
RF magnetron sputtering method according toReference [3]. Metallic Mg (99.9%) and Ni plates(99.9%) were used as targets. The incident RF
power was 100 W for Ni and 3050 W for Mg. The
sputtering was performed under an argon gas
atmosphere at a total gas pressure of 0.8 Pa, and theargon flow was fixed at 20 sccm. The films weresputtered onto silica substrates at 100qC. The film
thicknesses were approximately 100 nm.
Inaddition, a 5-nm-thick palladium layer was deposited
onto the films at room temperature. These films
were maintained under a hydrogen (H2) atmospherefor 24 h to prepare the hydride forms of the films.
Reflection spectra of the films were measured in
the vacuum ultraviolet region up to 40 eV with a 3 m
normal incidence monochromator (gratings: G1, G2,
and G3) using the BL7B beamline at UVSOR-III. A
silicon photodiode sensor was used as the detector forthe reflected light.
Figure 1 shows the refraction spectra of Mg, Ni,
and Mg-Ni alloy thin films before and after they were
exposed to H2. A change in the refraction spectra
before and after H2 exposure was observed only inthe case of the MgNi alloy thin film. The change
in the spectra was more pronounced in the case of the
MgNi alloy thin film deposited with a Mg RF powerof 40 W. As evident in the Fig. 1, no spectral

Reflectance

Mg-Ni alloy (after H 2 purge)

Mg-Ni alloy (before H2 purge)

Mg (before and after H purge)2

Ni (before and after H purge)
2

5

10

15
20
25
30
Photon energy (eV)3540Fig. 1. Reflection spectra of Mg, Ni, and Mg-Ni alloy
thin films before and after H2 purge.
[1] K. Nakagawa et al., UVSOR Activity Report 32(2005) 88.
[2] Y. Iwai et al., UVSOR Activity Report 35 (2008)
96.
[3] K. Yoshimura, J. Surf. Finish. Soc. Jpn 56 (2005)
882.

140

BL7B

Photoluminescence upon Vacuum Ultraviolet Excitation in Gd3+ Ion Doped
Inorganic Materials
Y. Inaguma, R. Horiguchi, S. Sasaki and D. Mori
Department of Chemistry, Faculty of Science, Gakushuin University, Tokyo 171-8588, Japan

Recently phosphor materials emitting UV lightupon vacuum ultra violet (VUV) excitation aredesired as the substitutes of mercury lamp due to thetoxicity of mercury. Among them, Gd3+ ion doped
compounds are expected to be candidates of the UV
phosphors since the f-f transition from 6PJ to 8S7/2 in
Gd3+ corresponds to UV region [1-4]. The UVemission would be observed when the band gap
energy of the host material is in the VUV region. In
this study we then chose alkaline-earth phosphates
with the band gap energy in the VUV region. Herein,
we synthesized A3(PO4)2:Gd3+(A = Ca, Sr, Ba) andinvestigated the luminescence properties upon VUVexcitation.
The polycrystalline samples of A3(PO4)2:Gd3+(A =Ca, Sr, Ba) were synthesized by a conventional solidstate reaction at elevated temperature. The mixture ofstarting materials: alkaline-earth carbonate such as
ACO3 (A = Ca, Sr, Ba), (NH4)H2PO4 and Gd2O3 or Gd
nitrate solution were calcined at 1000qC in air. The
calcined powder was pressed into pellets and sinteredat 1000-1200qC in air. The phase identification for the
samples was carried out by the powder X-ray
diffraction (XRD) using a Rigaku RINT 2100
Brentano geometry
diffractometer with a(graphite-monochromatized CuKD radiation). The
emission and excitation spectra were recorded by a
JASCO FP-6500 fluorometer upon the photoexcitation using a Xe excimer lamp ( O = 172 nm) orin the beamline BL7B at the UVSOR facility.
The obtained samples were crystallized to be thehexagonal E-type form of tricalcium phosphates for A= Ca, and hexagonal form with a space group R 3 m
for A = Sr and Ba. Figure 1 shows the emission
spectra upon photo excitation of 172 nm for A3(PO4)2:
Gd3+ 5% ( A = Ca, Sr, Ba) at room temperature. As
seen in Fig.1, the UV emission peaks at 307 nm, 310
nm were observed for all compounds, whichcorrespond to the f-f transitions, from 6P5/2 to 8S7/2,
from 6P7/2 to 8S7/2 states of Gd3+, respectively, andSr3(PO4)2: Gd3+ exhibited the highest emission
intensity. Figure 2 shows the excitation spectrum
emission-monitored at 310 nm for Sr3(PO4)2: Gd3+
1% at room temperature. Here, the excitation
spectrum was corrected for the spectral distribution of
excitation light source. In the excitation spectrum, thepeaks around 142, 275, and 155 nm correspond to the
f-f transitions from the ground state of 8S7/2 to 2Q23/2
state, 8S7/2 to 2Q23/2 of Gd3+, and the host absorption
band, i.e. the intra charge transfer of PO43-[5, 6],
respectively. While another peak at 130 nm has notbeen assigned at the present time.
Consequently, it was found that A3(PO4)2: Gd3+ (A= Ca, Sr, Ba) exhibits UV emission corresponding to
the f-f transition and they are candidates of UV
phosphors upon VUV excitation.
P7/2 ḵᴾ8S7/2

Intensity / arb.unit

6

Ca
Sr
Ba

6

P5/2 ḵᴾ8S7/2

300

305

310 315 320
Wavelength / nm

325

330

Intensity / arb.unit

Fig. 1. Emission spectrum upon excitation of 172 nm
for A3(PO4)2:Gd3+5% (A = Ca, Sr, Ba).

Host
(155 nm)

Excitation
Oem = 310 nm

2

ᴾ8S7/2㩷℉ Q23/2(142 nm)

100

?

8

6

S7/2㩷℉‒ I11/2(275 nm)

150

200
250
Wavelength / nm

300

Fig. 2. Excitation spectrum monitored at 310 nm for
Sr3(PO4)2: Gd3+ 1%.
The authors thank Mr. M. Hasumoto, Prof. S. Kimura,and Prof. K. Fukui for their experimental supports
and helpful advices.
[1] H. Yoshida, R. Yoshimatsu, S. Watanabe and K.
Ogasawara, Jpn. J. Appl. Phys. Part .1 45 (2006) 146.
[2] N. Yokosawa, G. Sato and E. Nakazawa, J.
Electrochem. Soc. 150 (2003) H197.
[3] T. Hirai, H. Yshida, S. Sakuragi and N. Ohno, Jpn.
J. Appl. Phys. Part .1 46 (2007) 146.
[4] S. Okamoto, R. Uchino, K. Kobayashi and H.
Yamamoto, J. Appl. Phys.,106 (2009) 013522.
[5] E. Nakazawa and F. Shiga, J. Lumin. 15 (1977)
255.
[6] H. Liang, Y. Tao, J. Xu, H. He, H. Wu, W. Chen, S.
Wang and S. Su, J. Solid State Chem. 177 (2004) 901.

141

BL7B

Ascription of the Luminescent Components of Cs2ZnCl4 Single Crystal
N. Yahaba1, M. Koshimizu1, T. Yanagida2, Y. Fujimoto3, R. Haruki4, F. Nishikido5,
S. Kishimoto6 and K. Asai1
1

Department of Applied Chemistry, Graduate School of Engineering, Tohoku University, Sendai 980-8579, Japan
2
Graduate School of Life Science and Systems Engineering, Kyushu Institute of Technology, Kitakyushu
808-0196, Japan
3
Institute of Materials Research, Tohoku University, Sendai 980-8579, Japan
4
Advanced Science Research Center, Japan Atomic Energy Agency, Tokai-mura 319-1195, Japan
5
Molecular Imaging Center, National Institute for Radiological Sciences, Chiba 263-8555, Japan
6
Institute of Materials Structure Science, High Energy Accelerator Research Ogranization, Tsukuba 305-0801,
Japan

Scintillation materials exhibiting short decay arehighly required for detection of high-energy X-ray
photons at high counting rate. Recently, we havefound that Cs2ZnCl4 crystals have excellent
scintillation characteristics. In this report, we present
the luminescence and scintillation dynamics of
Cs2ZnCl4.
Figure 1 shows the scintillation time profile excitedwith 50 keV X-ray. The profile was expressed as asum of two exponential components. The fast andslow ones have lifetimes of 1.8 and 10 ns,
respectively. The relative intensity of each component
was different for different crystals. The relative
intensity hardly depended on the raw materials. In
contrast, the contribution of the slow component
increased after annealing the crystal.
Figure 2 shows the luminescence spectra. Under
excitation at 88 nm, two luminescence bands wereobserved at 290 and 380 nm. This result is quite
similar to a previous report [1, 2], and these bands are
ascribed to the Auger-free luminescence. In the caseof excitation at 198 nm, luminescence was observed
at 320 nm. In this case, the excitation photon energy
was lower than the bandgap energy of Cs2ZnCl4, andthus the observed luminescence band was ascribed to
defects or impurities.
Figure 3 shows the luminescence time profile
under VUV excitation. The lifetime of the decay at
290 and 380 nm was estimated to be 2 ns, which isconsistent with the fast component in the scintillation
time profile. In contrast, the lifetime of the decay at
320 nm was estimated to be approximately 10 ns.
These results indicate that the slow component isascribed to defects would be eliminated by treating
the crystal after growth.

Without annealing
After annealing

Fig. 1. Scintillation time profiles of Cs2ZnCl4
crystals before and after annealing.

Fig. 2. Luminescence spectra of Cs2ZnCl4 under
VUV excitation.

[1] A. Ohnishi, M. Kitaura, T. Otomo and M. Sasaki,
J. Phys. Soc. Jpn. 72 (2003) 2400.
[2] A. Ohnishi, M. Kitaura, M.Itoh and M. Sasaki, J.
Phys. Soc. Jpn. 81 (2012) 114704.

Fig. 3. Luminescence time profile at different
wavelength.

142

BL7B

Electronic Status of Pr3+ in Glass Host by Characterizing
Photoluminescence of Pr3+ Doped APLF Glass Scintillator
T. Nakazato1, M. Tsuboi1, K. Takeda1, Y. Minami1, T. Hori1, T. Shimizu1,
R. Arita1, K. Yamanoi1, R. Nishi1, Y. Shinzato1, R. Sogi1, N. Sarukura1, H. Azechi1,
M. Cadatal-Raduban2, T. Murata3, S. Fujino4, H. Yoshida5, T. Suyama6, K. Fukuda6,
A. Yoshikawa7, N. Sato8, H. Kan8 and K. Kamata8
1

Institute of Laser Engineering, Osaka University, 2-6 Yamadaoka, Suita, Osaka 565-0871,Japan
Centre for Theoretical Chemistry and Physics, Institute of Natural and Mathematical Sciences, Massey
University, Albany, Auckland, 0632 New Zealand
3
Kumamoto University, 2-39-1 Kurokami, Kumamoto 860-8555, Japan
4
Kyushu University, 6-10-1 Hakozaki, Higashi-ku, Fukuoka 812-8581, Japan
5
Ceramic Research Center of Nagasaki, 605-2 Hasami-cho, Higashisonogi-gun, Nagasaki 859-37, Japan
6
Hamamatsu Photonics K.K., 325-6 Sunayama-cho, Naka-ku, Hamamatsu City, Shizuoka 430-8587, Japan
7
Institute for Materials Research, Tohoku University, 1-1-2 Katahira, Aoba-ku, Sendai, Miyagi 980-8577, Japan
8
Furukawa Co. Ltd., 2-6-1 Marunouchi, Chiyoda-ku, Tokyo 100-8370, Japan2

The realization of nuclear fusion is highly expectedto satisfy the energy demand in the future. Fusion
plasma diagnostics at higher plasma areal density
than 3 g/cm2 is one of the biggest breakthrough innuclear fusion research. Scattered-neutron diagnostics[1]
utilizing
Pr3+-doped
20Al(PO3)3-80LiF(APLF+Pr) glass scintillator [2] is expected toachieve plasma areal density measurement in fusion
core. APLF+Pr, which have high sensitivity for
scattered neutrons and fast response time about 5.4 ns,
is the key factor of the measure.
In this work, we report the luminescence properties
of APLF+Pr at different Pr3+ concentrations(0.5mol%㨪3mol%).
We have measured Photoluminescence (PL)spectra and Photoluminescence Excitation (PLE)
spectra of APLF+Pr at 30K, 100K, 200K, and 300K.
The experiment was conducted at BL7B utilizing G2
grating. PL spectra are measured with excitation
wavelength at 180nm, 200nm and 217 nm. The
APLF+Pr samples, having Pr3+ density at 1mol%, andcut into 10 mm  10 mm  2.5 mm, wereprepared. PL and PLE spectra of Pr3+-doped
LiCaAlF6 (Pr:LiCAF) was measured as well forcomparison.
The PL spectra of APLF+Pr excited by 180 nm
radiation at 27 K, 100 K, 200 K and 300 K are shownin Fig. 1. 2 sharp peaks and 3 broad peaks have beenfound from 200 nm to 300 nm. The 2 peaks strongly
depend on temperature and intensities decrease withincreasing temperature. However the 3 broad peaks
do not strongly depend on the temperature ascompared with the 2 sharp peaks. The broad peaks of
227 nm, 236 nm and 261 nm are ascribe to 4f5d-5H4,
4f5d-5H5, and 4f5d-5H6, 5F2 and 5F3 respectively
based on the results of Pr:LiCAF PL spectra. The
sharp peaks of 253 nm and 273 nm are also ascribedto 1S0-5F4 and 1S0-1G4 respectively.

143

Fig. 1. PL spectra of APLF at various
temperature.
[1] M. Moran et al., Rev. Sci. Instrum. 75 (2004)
3592.
[2] T. Murata et al., IEEE Trans. Nucl. Sci., 57 (2010)
1426.

BL8B

Photoelectric Emission from Ag/zinc-phthalocyanine Surface
Under Near-UV Light Irradiation
S. Tanaka, K. Fukuzawa, T. Otani and I. Hiromitsu
Interdisciplinary Graduate School of Science and Engineering, Shimane University,
1060 Nishikawatsu, Matsue, Shimane 690-8504, Japan

3.3 eV. The shift of vacuum level is frequently
observed at the organic/metal interface [3], although
no anomalous photoelectric emission has beenreported. The surface structure of the Ag/ZnPc
sample was observed by atomic force microscope(AFM). The AFM images showed that Ag was
deposited on the ZnPc surface with grain sizes of few
nano meters. The structure of the Ag particles could
be related to the reduction of the surface barrier
potential. The fact that the anomalous photoelectric
emission was suppressed at the thicker Ag region (Fig.
2. inset) could support this hypothesis.

The organic semiconductors have gained
significant importance in the optoelectronic devices
in the last decade. The electronic structure of theorganic semiconductor has attracted interest from theview point of not only the basic science but also thepractical application. For example, the electronic
structure of the organic semiconductor/metal interface
has an important role on the device performance ofthe organic solar cells [1]. Recently, we have studied
the light irradiation effects on the electronic structureof the organic/metal and organic/organic interfaces toinvestigate the electronic structure under the working
condition of the solar cells [2]. During the research,
we found an anomalous photoelectric emission fromthe organic thin films with a surface treatment bymetal deposition.
Figure 1 shows the photoelectric emission from thezinc phthalocyanine (ZnPc) film with Ag deposition
under the 360 nm (= 3.44eV) laser irradiation. The
nominal thickness of the ZnPc and Ag layer were 10
nm and 0.05 nm, respectively. The work function of
Ag (~4.5 eV) and the ionization potential of ZnPc(~5.0 eV) are larger than the energy of the laser light.
Thus, the present photoelectric emission could be an
anomalous phenomenon at a particular condition of
the organic/metal interface. The threshold of lightenergy for the photoelectron emission wasapproximately 3.3 eV.
To investigate the effect of the Ag deposition on
the electronic structure of the sample surface,photoelectron spectroscopy was performed. The ZnPc
layer was deposited on the indium tin-oxide (ITO)
substrate by vacuum deposition at the preparationchamber. Ag was also deposited on the ZnPc layer byvacuum deposition. The thickness of the eachmaterial was monitored by quartz microbalance. The
sample was transferred to the measurement chamber
without breaking the vacuum. All the photoelectron
measurements were carried out at room temperature.
Figure 2 shows the variation of the photoelectron
spectra of Ag deposited on the ZnPc substrate withincreasing thickness. The highest occupied molecularorbital of ZnPc was shifted approximately 0.5 eV
toward lower kinetic energies by the deposition of 0.5
nm of Ag. The secondary electron cutoff was also
showed a corresponding shift. These results infer thatthe potential barrier for the photoelectric emission at
the sample surface was decreased by the Ag
deposition. However, the estimated value of the work
function of the sample surface is much larger than the

Ag (0.05nm) on ZnPc

Intensity / Arb. unit

Bias = -10 V

ehQ = 3.44eV

Ag(0.05nm)

ZnPc

9

10

11

12

Kinetic energy / eV

Photoelectric emission
intentisy / Arb. unit

Fig. 1. Photoelectric emission from the Ag (0.05nm)/ZnPc (10nm) /ITO emitted by near-UV lightirradiation (hQ = 3.44 eV). The sample was biased by
í10 V to measure the electron with small kinetic
energies.

Intensity / Arb. units

hQ = 40 eV

20000

10000

0
0

5

10

15

20

Ag thickness / nm

Ag(20.0nm)/ ZnPc(15nm)/ ITO
Ag(10.0nm)/ ZnPc(15nm)/ ITO
Ag(3.0nm)/ ZnPc(15nm)/ ITO
Ag(1.0nm)/ ZnPc(15nm)/ ITO
Ag(0.5nm)/ ZnPc(15nm)/ ITO
ZnPc(15nm)/ ITO

30 31 32 33 34 35 36

Kinetic energy / eV

Fig. 2. Photoelectron spectra of ZnPc and Ag/ZnPc
sample near the Fermi level. Inset: The Ag thickness
dependence of the photoelectric emission intensity.
[1] S. Tanaka et al., Appl. Phys. Lett. 97 (2010)
253306.
[2] S. Tanaka et al., UVSOR Activity Report 39(2012) 130.
[3] H. Ishii et al., Adv. Mater., 11 (1999) 605.

144

BL8B

Does Polarity of the Giant Surface Potential Impacts Charge Transfer
at Organic-Organic Interfaces?
H. S. Lim1, Y. Ozawa1, Y. Uragami1, Y. Nakayama2, Y. Noguchi1,2 and H. Ishii1,2
1Graduate School of Advanced Integration Science, Chiba 263-8522, Japan
2
Center for Frontier Science, Chiba University, Chiba 263-8522, Japan145

The present results implies occurrence of chargetransfer from D-NPD to the overlayers irrespective of
the polarity of GSP, which is contradicting to ourprevious proposition [2]. In order to figure out theelectronic phenomena at the topical organic-organicinterface, further systematic studies by altering theoverlayer characteristics are highly anticipated.
䎹䏖䎠䎃䎐䎘䎹

䏋Q䎃䎠䎃䎗䎓䎃䏈䎹

䎐䎓䎑䎕

䎋䏙䎌

䎫䎲䎰䎲䎑䎃䎶䎨䎦䎲䎃䎒䎃䏈䎹䎱䏒䏕䏐䏄䏏䏌䏝䏈䏇䎃䏌䏑䏗䏈䏑䏖䏌䏗䏜䎃䎒䎃䏄䏕䏅䎑䎃䏘䎑

䎐䎓䎑䎔

䎖䎖䎑䎜䎗
䎓䎑䎔

䎋䏌䏙䎌

䎖䎗䎑䎖䎙

䎋䏌䏌䏌䎌

䎐䎕䎗䎑䎓䏐䏈䎹䎒䏑䏐

䎐䎓䎑䎖

䎐䎓䎑䎗

䎖䎗䎑䎔
䎖䎗䎑䎓

䎐䎕䎜䎑䎖䏐䏈䎹䎒䏑䏐

䎖䎖䎑䎜
䎋䏌䏌䎌

䎖䎖䎑䎛
䎓
䎕
䎗
䎙
䎛 䎔䎓
䎷䏋䏌䏆䏎䏑䏈䏖䏖䎃䏒䏉䎃䎤䏏䏔䎖䎃䎒䎃䏑䏐

䎖䎗䎑䎗䎙

䎋䏌䎌

䎘

䎙 䎖䎕
䎖䎖
䎖䎗
䎮䏌䏑䏈䏗䏌䏆䎃䏈䏑䏈䏕䏊䏜䎃䎒䎃䏈䎹䎖䎘

Fig. 1. UPS spectra of the Alq3 [(i) : 0nm, (ii) : 0.1nm,(iii) : 0.4nm, (iv) : 1.6nm, (v) : 6.4nm] on D-NPD(5nm) interfaces. The gray line shows the spectra
contribution of D-NPD, which is derived from aprocedure described in Ref. [2].
䎹䏖䎃䎠䎃䎐䎘䎃䎹

䏋Q䎃䎠䎃䎗䎓䎃䏈䎹
䎓䎑䎗

䎓䎑䎖
䎓䎑䎕

䎖䎗䎑䎓䎛

䎋䏙䎌

䎫䎲䎰䎲䎃䎑䎃䎶䎨䎦䎲䎃䎒䎃䏈䎹䎱䏒䏕䏐䏄䏏䏌䏝䏈䏇䎃䏌䏑䏗䏈䏑䏖䏌䏗䏜䎃䎒䎃䏄䏕䏅䎑䎃䏘䎑

Ito et al. reported spontaneous building up of thesurface electric potential (giant surface potential;
GSP)
invacuumdepositedfilmsoftris(8-hydroxyquinolinato) aluminum(III) (Alq3) [1].
The origin of GSP can be attributed to a slightalignment of dipole moment of the molecules inaverage. The polarity of GSP in the case of Alq3 is
positive in increasing the film thickness.
We have reported how the existence/absence of GSP
in an overlayer impacts the electronic nature atorganic hetero-interfaces [2]. These works, however,are conducted on overlayers of non-GSP or positive
GSP materials. In this context, an Alq3-derivativetris(7-propyl-8-hydroxyquinolinato)
aluminum(III)(Al(7-prq)3) is an interesting material because thisreveals a negative GSP [3].
In this experiment, we investigated the influence ofdifferent GSP direction at organic interfaces of a
N,N’-bis(1-naphthyl)-N,N’-diphenyl-1,10-biphenyl-4,40-diamin (D-NPD) film capped with the Alq3 orAl(7-prq)3 overlayer.
Ito is used as substrate and is cleaned byultrasonication in neutral detergent, isopropanol andacetone. nm-thick D-NPD layers were thendeposited onto the ITO substrate. Alq3 or Al(7-prq)3
was subsequently deposited onto D-NPD from 0.1nm
to 10.0nm in a step-by-step manner. Ultraviolet
photoelectron spectroscopy (UPS) was performed at
BL8B at UVSOR. 40 eV of photon energy was used.
-5 V bias was applied during secondary electron
cutoff (SECO) measurement. The experiment iscarried out at room temperature.
At the D-NPD/Alq3 interface, the highest occupied
molecular orbital (HOMO) peak of D-NPD (as wellas the SECO position) moved with growing the
thickness of Alq3 by 0.1 eV to low kinetic energy side.
This implies that deposition of Alq3 induces
downward band-bending in the D-NPD layers,
namely charge transfer from D-NPD to Alq3. As
growing the Alq3 overlayer further thicker, the SECO
position shifted down proportionally to the Alq3
thickness representing GSP.
D-NPD HOMO peak moved to low kinetic energy
side also at the D-NPD/Al(7-prq)3 interface, yet theshifting value (~0.06 eV) was a little less than thecase of the unsubstitued Alq3. While SECO did notexhibit apparent shift in the thin coverage region, it
moved to the high kinetic energy side in increasing
the Al(7-prq)3 thickness, which can be attributed to
GSP of the opposite polarity to the Alq3 layers.

䎋䏌䏙䎌
䎓䎑䎓䎙

䎋䏌䏌䏌䎌

䎖䎗䎑䎖䎕

䎚䎕䎑䎕䏐䏈䎹䎒䏑䏐

䎓䎑䎔
䎓䎑䎓

䎐䎓䎑䎔

䎖䎗䎑䎖
䎖䎗䎑䎕
䎖䎗䎑䎔

䎋䏌䏌䎌

䎐䎖䎙䎑䎗䏐䏈䎹䎒䏑䏐

䎖䎗䎑䎓

䎋䏌䎌

䎓
䎕
䎗
䎙
䎛 䎔䎓
䎷䏋䏌䏆䏎䏑䏈䏖䏖䎃䏒䏉䎃䎤䏏䎋䎚䎐䏓䏕䏔䎌䎖䎃䎒䎃䏑䏐

䎖䎗䎑䎖䎛

䎘

䎙 䎖䎕
䎖䎖
䎖䎗
䎮䏌䏑䏈䏗䏌䏆䎃䏈䏑䏈䏕䏊䏜䎃䎒䎃䏈䎹䎖䎘

Fig. 2. UPS spectra of the Al(7-prq)3 [(i) : 0nm, (ii) :0.1nm, (iii) : 0.4nm, (iv) : 1.6nm, (v) : 10.0nm] on
D-NPD (5nm) interfaces.
[1] E. Ito et al., J. Appl. Phys. 92 (2002) 7306.
[2] Y. Nakayama et al., Organic Electrons 13 (2012)2850.
[3] T. Isoshima et al., submitted.