Research

How Can We Decrease Defect Density in Si Multicrystals
to Realize High-Efficiency Solar Cells?
To manipulate crystal defects and impurities in Si
multicrystals is a promising route to improve conversion
efficiency of solar cells. Therefore, development of a novel
crystal growth technique to permit artificial control of defects
has been pursued. It would be ideal if such a technique isbased on the practical directional solidification. For thispurpose, fundamental understanding of mechanisms on
generation of defects during crystal growth is of crucial
importance so that one can answer an open question, “Howcan we decrease defect density in Si multicrystals?”
In order to answer the question, we started with model
crystal growth experiments using a seed crystal to consist ofa couple of single crystals with an artificially controlled grainboundary configuration, namely <110> tilt grain boundary.
Spatially resolved X-ray rocking curve analysis and evaluationof etch-pit density clarified that “sub-grain boundaries”, which
consist of a cluster of dislocations and give detrimental effecton solar cell performance, are generated from the grainboundary as crystal growth proceeds. Furthermore, theamount of dislocations was found to be strongly dependent
on the grain boundary configuration. In a particularconfiguration, dislocations were mostly localized in one of thegrains to consist of a grain boundary [1].
Since Si expands when crystallized from the melt,
compressive stress is supposed to be applied owing to thecontact with a rigid crucible. This led us to assume that local
shear stress around a grain boundary, which would dependon the configuration of the grain boundary, might explain theexperimentally observed difference in defect density. To see
if this hypothesis is reasonable, we carried out threedimensional finite element analysis of shear stress for amodel bi-crystal with a single grain boundary. By taking
anisotropic elastic constants into account, calculations were
performed for various grain boundary configurations.
Interestingly, dislocation density was found to be wellcorrelated with the calculated local shear stress around thegrain boundary. The shear stress is likely to be reduced when

0.8
0.6
0.4
0.2
0.0

6

grain boundaryx10

Growth direction

1.0

Dislocation density 106/cm2

generation of
dislocations

generation of
dislocations

Contact angle ：2°

Contact angle：63°

Grain boundary

50μm

Fig. 2 Comparison of dislocation density around grainboundaries with different contact angles of adjacent dendrite
crystals. It is seen that parallel contact is favorable in termsof reduction of dislocations.
the rotational symmetry is improved along the direction
perpendicular to the plane where external stress from the
crucible is introduced.
This knowledge inspired us to suppress dislocations
during directional solidification by controlling the coherencyof the grain boundary without any seeds so that one could
decrease shear stress around the grain boundary.
The attempt to change the coherency of the grainboundary was made by utilizing nucleation of plural dendrite
crystals. Since the upper plane of the dendrite crystal can belimited to either {110} or {112}, the coherency of the grainboundary can be controlled by the contact angle between
adjacent dendrite crystals. Crystal growth was started withdendritic nucleation by establishing the amount ofsupercooling larger than 10 K. The contact angle wasmeasured by using electron back scattering pattern analysis.
Dislocation density measured by counting etch-pits revealed
that parallel contact of adjacent dendrite crystals is useful tosuppress generation of dislocations. On the other hand,increase of the contact angle resulted in increase of thedislocation density. This confirms that control of the coherencyof grain boundaries at the initial stage is very important for
decreasing dislocations in the whole ingot [2].
References[1] I. Takahashi, N. Usami, K. Kutsukake, G. Stokkan, K.
Morishita, and K. Nakajima, J. Cryst. Growth 312, 897-901(2010).
[2] N. Usami, R. Yokoyama, I. Takahashi, K. Kutsukake, K.
Fujiwara, and K. Nakajima, J. Appl. Phys. 107, 013511 (2010).

5mm

shear stress small

large

large

Key Words
Dislocation, Grain Boundary, Dendrite
small

Fig. 1 Spatial distribution of dislocations in model crystals
grown using seeds with controlled grain boundaryFig.1
configuration, and shear stress distribution calculated byfinite element analysis. It is clearly seen that dislocations are
generated from grain boundaries and localized at one of thecrystal grains with larger shear stress.

KINKEN Research Highlights 2010 IMR

Contact toNoritaka Usami (Crystal Physics Division)
E-mail: usa@imr.tohoku.ac.jp
URL: http://www.xtalphys.imr.tohoku.ac.jp

Energy-related Materials

11Research

Phase Stability of Metallic Nano-Clusters Studied by Positron 2D-ACAR :Fermi Surface - Brillouin Zone Interactions

(b)

[010]

[100]

Al (fcc)[a.u.]

[010]

Al‐Ag

2.5

[a.u.]

Al3Ag (L12)

Al‐Zn

2

[010]

[010]

[100]

AlAg (L10)[100]

Zn (fcc)[100][a.u.][a.u.]

1.5[100][a.u.]

AlAg3 (L12)1
Ag (fcc)[100][a.u.]

Fig. 2 3D plots of calculated 2D-ACAR anisotropies projectedalong the [100] direction and Fermi surfaces by FLAPW
calculations for pure Al, Al3Ag, AlAg, AlAg3, pure Ag, andpure Zn (fcc).

References[1] Y. Nagai, M. Hasegawa, Z. Tang, A. Hempel, K. Yubuta,
T. Shimamura, Y. Kawazoe, A. Kawai, and F. Kano, Phys.
Rev. B 61, 6574 (2000).
[2] Y. Nagai, T. Chiba, Z. Tang, T. Akahane, T. Kanai, M.
Hasegawa, M. Takenaka, and E. Kuramoto, Phys. Rev. Lett.
87, 176402 (2001).
[3] Y. Nagai, T. Toyama, Z.Tang, K.Inoue, T.Chiba, M.
Hasegawa, S.Hirosawa and T.Sato: Phys. Rev. B 79,
201405(R) (2009).
Key Words
Embedded Nanoparticle,
Surface

Al‐Zn

Positron

Annihilation,

Fermi

Contact to

[010]

Al‐Ag

[010]

(a)

3

[010]

It is well known that a positron, an antiparticle of anelectron, is a sensitive probe of open volume defects because
it can be trapped in the defects and then annihilates therewith the surrounding electron conveying significant informationon the electronic environment around the defect site with theemitted two gamma-rays. However, we recently found that apositron can be also trapped in solute nano-clusters withpositron affinities higher than that of matrix even if the nanoclusters are free from the open volume defects. We named
this phenomenon quantum positron confinement. [1]. Using
the quantum positron confinement, we can reveal electronic
structure of the solute nano-clusters by two-dimensional
angular correlation of positron annihilation radiation(2D-ACAR) method [2].
As an example, Fig. 1 shows the anisotropies of electron
momentum distributions in Ag and Zn clusters in Al
experimentally obtained by 2D-ACAR method [3]. We candiscuss the Fermi surface topology of the solute nanoclusters from the 2D-ACAR anisotropies. By comparing withthe theoretical calculations shown in Fig. 2, the chemicalcompositions (electron/atom ratios) in the solute clusters are
determined. The Ag cluster is close to Al:Ag=1:3 and the Zncluster is almost pure Zn.
Based on these results, we proposed a general rule tosuggest preferential chemical composition of solute nanoclusters coherent to matrix with nearly free electrons. The
Fermi surface - Brillouin zone interaction is a key tounderstand the chemical composition realized in the solutenano-clusters because the presence of necks on the Fermi
surface contributes to the reduction in the total electronenergy due to the band gap effect.

e/a

[010]

A general approach is outlined to predict the chemicalcompositions of solute nano-clusters embedded in materials
with nearly free electrons. Based on the experimental resultsof the two-dimensional angular correlation of positronannihilation radiation (2D-ACAR) and the corresponding
theoretical calculations, we show that the Fermi surface Brillouin zone interaction is a key to understand the chemicalcomposition realized in solute nano-clusters. This concept isalso useful to understand the stability of impurity clusters
formed in nuclear materials by irradiation.

Yasuyoshi Nagai (Irradiation Effects in Nuclear and Their RelatedMaterials Division)
E-mail: nagai@imr.tohoku.ac.jp
URL: http://wani.imr.tohoku.ac.jp/[100][a.u.][100][a.u.]

Fig. 1 3D plots of experimental 2D-ACAR anisotropies
projected along the [100] direction for (a) Ag cluster in Al and(b) Zn cluster in Al.

12

Energy-related Materials

IMR KINKEN Research Highlights 2010

Research

Hydrogen Incorporation and
Optical Absorption of Tungsten Oxide Films

WO2/WO3Hx/SiO2 sample
MeV He++analyzing beamkeV proton
implantation beam

before implantation

0.2
after implantation)

0.1
0.0
400

Recoil H yield (counts/channel)

light source

0.3
Transmittancetransmitted light

2000
1500
1000
500
0
0

600
800
Wavelength (nm)1000

recoil hydrogen
before implantation
after implantation
SiO2
substrate
WO3 layer (500nm)

0.8

5
-1

4

0.6

3
0.4
2
0.2

1
0
0

17

17

1x10
2x10
-2
Hydrogen ion fluence [cm ]

0.0
17
3x10

H retained
In WO3 layer

H
W

WO3

SiO2

Fig. 2
Absorption coefficient and retained hydrogenconcentration in the sandwiched WO3 layer, plotted as afunction of incident 3 keV H+ ions. The photo-images of Pd/
WO3/SiO2 sample as prepared and exposed by 1%H2-Ar
mixture gas are inserted in upper part.
the film increased with the coloring process. The evidence ofthe coloration accompanied by the hydrogen incorporation
supports the double injection model, in which protons andelectrons are simultaneously inserted into a WO3 film to
reduce WO6+ to WO5+. The relation between hydrogen andcoloration has been further analyzed under ion implantation,and structural and compositional effects on the opticalproperties were systematically analyzed to optimize thepreparation condition of the films [2].
Figure 2 shows changes of optical absorption in the
wavelengths between 600 and 900 nm of W/WO3/SiO2
sample, and hydrogen retention in the sandwiched WO3layer under 3 keV H+ ion implantation at a room temperature.
Because the incident H ion totally loses its energy in the W
layer, no damage was created in the sandwiched WO3 layer.
At the beginning of the implantation, nearly 100 % hydrogenatoms implanted in the W layer diffused into the WO3 layer
and were retained, accompanied with coloration of the WO3layer. The coloration was completed at the fluence of about
1x1017 cm-2, corresponding to the increase of he Hconcentration about 0.2 H/W. The evaluated coloration
efficiency about 40 cm2C-1 was fairly consistent with theprevious experiment under Ar-H2 gas exposure [1], indicatingthat mechanism of the ion implantation induced coloration is
the same as the gaschromic phenomenon.
References[1] S. Nagata, A. Inouye, S. Yamamoto, B. Tsuchiya, K.
Takano, K. Toh and T. Shikama, J. Alloys and Comp. 446447, 558 (2007).
[2] A. Inouye, S. Yamamoto, S. Nagata, M. Yoshikawa and T.
Shikama, Nucl. Instr. and Meth. B 267, 1480 (2009).
Key Words
Tungsten Oxides, Ion-beam, Optical Absorption

200
400
Channel number

Fig. 1 Schematics of the simultaneous measurement system
for hydrogen concentration and optical absorption of tungsten
oxide films. During the keV H+ ion implantation into surface
layer, energies of hydrogen recoiled by MeV He ions areanalyzed, and a transmitted light spectrum is obtained.

KINKEN Research Highlights 2010 IMR

1% H2-Ar gas exposure on Pd/WO3

Hydrogen concentration [H/W]

In near-term fusion devices such as ITER (International
Thermonuclear Experimental Reactor), tungsten is acandidate material for plasma facing components because of
its superior physical and mechanical properties at hightemperature. Although the pure tungsten metal shows very
little solubility of hydrogen, re-deposited layer on the first wall
surface can contain carbides and oxides, retaining hydrogenisotopes up to considerably high concentration. It is essentially
important to clarify the transport behavior of hydrogenisotopes in the first wall surface to control fuel balance andtritium inventory. On the other hand, it is known that tungsten
tri-oxide films covered with a thin catalyst layer show areversible coloration phenomenon under hydrogen exposure.
Because the hydrogen gas has a relatively low explosive limit
in the atmosphere, development of a sensor of hydrogen gas
is also important to handle hydrogen safely. The mechanismof this gasochromic phenomenon is not fully understood,
while the optical characteristic of the tungsten oxide film wasinvestigated by electrochemical techniques and by infraredabsorption
Recently, we demonstrated that hydrogen atoms wereincorporated in the WO3 film during the gasochromiccoloration by direct measurement of hydrogen concentrationin the film using the ion beam analysis technique combinedwith optical measurements as shown in Fig. 1 [1]. Tungsten
oxide thin films were prepared on a SiO2 glass substrates bya reactive RF magnetron sputtering using a metal W target.
The concentration depth profiles of hydrogen in the nearsurface layer were measured by the Elastic Recoil Detection
Analysis (ERDA) technique, performed in a scattering
vacuum chamber which was connected to a 1.7 MV tandem
accelerator. The optical transmittance was simultaneously
measured by a CCD camera equipped with a monochrometer.
It was clearly observed that the hydrogen concentration in

3 keV H+ ion implnataion into W /WO3

Change of absorbance [µm ]Ion-beam analysis techniques combined with opticaldetection can provide useful non-destructive and in-site
observation of retention and transport behavior of hydrogenin nuclear materials. We demonstrate simultaneous
measurements of optical absorption and hydrogenconcentration in tungsten oxide films under gas exposure
and/or ion implantation, to reveal the coloration phenomena
accompanied with an insertion of hydrogen in the tungsten
tri-oxide film.

Contact toShinji Nagata (Nuclear Materials Science Division)
E-mail: nagata@imr.tohoku.ac.jp

Energy-related Materials

13ResearchIn-situ TEM Observation and First Principle Calculation Study on
Hydride Formation in Zirconium Alloy

The mechanical degradations of zircaloy claddings in
LWRs strongly depend on distribution of hydrides because of
their characteristics in mechanical properties. Therefore, it isimportant to understand the formation mechanism of Zr
hydride to control the behavior of embrittlement of cladding
materials [1].
In this study, we investigated the formation of Zr hydrides
associated with dislocations both by electron microscopy,
molecular dynamics and first principle calculations.
1. In-situ TEM observations of Zr hydride under hydrogenion implantation
In-situ observations during hydrogen ion implantation
were performed with a TEM interfaced with an ion accelerator.
In electron transparent thin foils, low density of hydrides wereintroduced during electrochemical polishing. Growth
processes of Zr hydrides accompanied by the formation of
dislocations were observed simultaneously under irradiation
with 150-keV H2+ ions at room temperature as shown in Fig.
1. The crystallographic relationship between the Zr matrix
and Zr hydrides was identical with the previous studies. The
direction of growth was in the direction. Dislocations wereformed at the tip of hydrides. The contrast implied they wereextended into partial dislocations. The distance between
them were less than 5 nm and 20nm, respectively, in Zr
matrix and at a tip of hydrides.

Dislocation

Zr(hcp)+2H-16Zr(hcp)+2H(固溶体)
Solid solution-16.2

Zr(fcc)+2H(固溶体)Z(f ) 2H(固溶体)
Solid solution-16.4

γ-ZrH+H

-16.6

δ
δ-ZrH
2-x

ε-ZrH2-16.8

Volume increase byV
y 17.4%
%

-17
21

22

23

24

25

26

27

28

29

30

31

Volume per Zr-atom (Å3)

Fig. 2 First principle calculation of zirconium hydrides
compared with pure zirconium.

2. Modeling for hydride formation
Molecular dynamics simulation was performed torepresent the formation of partial dislocations which areextended from a basal dislocation and roughly 4 nm apart,well agree with the experimental results, and suggest fairly
low stacking fault energy in hydride. The first principle
calculations employing VASP 4.6 code was applied to theZr-H system to clarify (1) stability of Zr-H solid solution, (2)preferable crystal plane of H atoms and (3) stable Zr hydride
crystal. As shown in Fig. 2, it is concluded that lower energy
was achieved in Zr (fcc)-H solid solution than Zr (hcp)-H, andε or δ-ZrH2 is the stable intermetallic compound among
possible allotropes. Accumulation of hydrogen at partial
dislocations is presumably attributable for the hydride
formation.

Key Words
Zirconium, Hydrogen Implantation, Transmission Electron
Microscope

Hydride
100nm(a) 6.1×1019 H/m2

Zr(fcc)+2H

-8.2
-15.8

References[1] H. Abe, Japan/Korea Joint Seminar, Annual Meeting ofAtomic Energy Society of Japan Sept 16, 2009,Tohoku
University

Zr(Fe,Cr)2

g

-15.6
-8.0

Total energy p
per Zr atoom (eV))The growth process of Zr hydride in Zircaloy-4 wasinvestigated using in-situ TEM observations underimplantation with hydrogen ions followed by molecular
dynamics and first principle calculations so as to clarify theformation mechanism. Stacking fault associated with partial
dislocations on the hydride is presumed to have an importantrole for the formation.

(b) 1.0×1020 H/m2

Contact toHiroaki Abe (Nuclear Materials Engineering Division)
E-mail: abe.hiroaki@imr.tohoku.ac.jp
URL: http://www-lab.imr.tohoku.ac.jp/~abeken/(c) 1.2×1020 H/m2(d) 1.5×1020 H/m2

Fig. 1 The preferential growth of zirconium hydride along
under 150-keV H2+ ions at room temperature.

14

Energy-related Materials

IMR KINKEN Research Highlights 2010

Research

Complex Hydrides with (BH4)– and (NH2)– Anions
as New Lithium Fast-Ion Conductors
We have studied lithium ion conductivities of complex
hydrides Li2(BH4)(NH2) and Li4(BH4)(NH2)3, both withcombinations of the (BH4)– and (NH2)– anions. Ac impedance
measurements demonstrate that both Li2(BH4)(NH2) andLi4(BH4)(NH2)3 are novel lithium fast-ion conductors (2×10−4
S/cm at room temperature) due to being provided new
occupation sites for Li+ ions.
We have recently reported that a complex hydride
Li(BH4), a potential candidate for advanced hydrogen storage
materials [1], exhibits lithium fast-ion conduction (10–3 S/cm)
due to its structural transition at 390 K [2]. Research anddevelopment of lithium (fast-)ion conductors is significantly
important because they can be potentially used as solid
electrolytes in all-solid-state lithium ion batteries. From theapplication point of view, it is highly desirable to enhance theconductivity of Li(BH4) at room temperature (RT).
Figure 1 shows the temperature dependences of theelectrical conductivities of Li2(BH4)(NH2) and Li4(BH4)(NH2)3[3]. Li2(BH4)(NH2) exhibits fast-ion conductivity of 2×10–4
S/cm at RT, which is four and five orders of magnitude higher
than those of the host hydrides Li(BH4) and Li(NH2),
respectively, and the conductivity monotonically increases
upon heating. The activation energy for conduction
significantly decreases at around 368 K from 0.56 eV (303348 K) to 0.24 eV (above 368 K). This significant change in
the activation energy is attributed to the melting ofLi2(BH4)(NH2). The total ion conductivity reaches up to 6×10–2
S/cm after melting at the highest temperature measured (378
K). Li4(BH4)(NH2)3 also exhibits fast ion conductivities of
2×10–4 S/cm even at RT. The activation energy for conduction
is evaluated to be 0.26 eV, which is less than half that inLi2(BH4)(NH2) before melting and Li(BH4), indicating thehigher mobility of Li+ ions.

0Electrical conductivity (S/cm)

10425 400

Temperature (K)
375
350
325

300

-1

10

-2

10

-4

Li4(BH4)(NH2)3

10

Li(1)

H

B

N
H

Li(3)

Li(1)

Li(2)

N
Li(2)

Fig. 2
Local atomic structures of Li2(BH4)(NH2) andLi4(BH4)(NH2)3. Red, green, orange and blue solid circles
show Li, B, N and H, respectively.
References[1] S. Orimo, Y. Nakamori, J.R. Eliseo, A. Züttel and C.M.
Jensen, Chem. Rev., 107, 4111 (2007).
[2] M. Matsuo, Y. Nakamori, S. Orimo, H. Maekawa and H.
Takamura, Appl. Phys. Lett., 91, 224103 (2007).
[3] M. Matsuo, A. Remhof, P. Martelli, R. Caputo, M. Ernst, Y.
Miura, T. Sato, H. Oguchi, H. Maekawa, H. Takamura, A.
Borgschulte, A. Züttel and S. Orimo, J. Am. Chem. Soc., 131,
16389 (2009).

E-mail: orimo@imr.tohoku.ac.jp-610

URL: http://www.hydrogen.imr.tohoku.ac.jp

Li(BH)4

-7

10

2.2

H

B

Shin-ichi Orimo (Hydrogen Functional Materials Division)

10

-8

H

Li4(BH4)(NH2)3

Contact to-5

10

Li2(BH4)(NH2)Key Words
Complex Hydride, Lithium Fast-Ion Conductor, Solid
Electrolyte

Li2(BH4)(NH2)-3

10As shown in Fig. 2, both Li2(BH4)(NH2) and Li4(BH4)(NH2)3
have plural occupation sites for Li+ ions with different
tetrahedral coordination consisting of (BH4)– and (NH2)–
anions. Li2(BH4)(NH2) has eighteen Li(1)s and eighteen
Li(2)s per unit cell while Li4(BH4)(NH2)3 has twelve Li(1)s,twelve Li(2)s and eight Li(3)s. Calculation of the size of theeach occupation site (one good factor to estimate the size ofthe bottleneck for Li+ ion diffusion) demonstrates that 50% Li+
ions occupy the rather small Li(2) (approximately 56% inradius and 18 % in volume of the Li(1)) in Li2(BH4)(NH2). The
larger activation energy for conduction for Li2(BH4)(NH2) thanfor Li4(BH4)(NH2)3 can be explained by the Li+ ions with lower
mobility in the smaller Li(2).

Li(NH2)

2.4

2.6

2.8

3.0

-1

3.2

3.4

1000/T (K )Fig. 1 Temperature dependences of the electrical conductivityof Li2(BH4)(NH2) and Li4(BH4)(NH2)3. The melting temperatureof Li2(BH4)(NH2), 365 K, is indicated as Tm. For reference, thedata of Li(BH4) and Li(NH2) as host hydrides are also
shown.

KINKEN Research Highlights 2010 IMR

Energy-related Materials

15Research

Hydrothermal Synthesis of Actinide Dioxides
for Fabrication and Reprocessing of MOX Nuclear Fuel

Spent Fuel

0.2

0.3

(b1) size 1

Cubic

Frequency

Frequency

0.1
0.0
0.3

(b2) size 2

0.2

(c)

0.2

0.1

0.1
2

3

4 5 6 78(2) EtOH + Hexylaldehyde

3

0.0

4 5 6

0.1
0.0
0.3

(b2) size 2

0.2

1

0.8

(b1) size 1

0.2

Spherical

2

1
size /µm

Frequency

0.0
0.1

(a)

Frequency

23

4 5 673

4 5 67

10
Aspect ratio

2

3

4

2

3

4

2

3

4

(c)

0.6
0.4
0.2

0.1
0.0
0.1

(a)

(3) EtOH + Guanidinium carbonate

Rectangularparallelpiped

0.2

2

3

4 5 6 781
size /µm

2

3

0.0

4 5 6

1

2

(b1) size 1

0.2

(c)

0.1
0.0
0.2

10

Aspect ratio

Frequency

The small emission of carbon dioxide from the nuclear
power generation has made it attractive energy sources. In
the commercial nuclear power plants, uranium dioxides or
the solid solution with plutonium dioxides, i.e. so-called the"mixed oxides" (MOX), with fluorite-structure have been used
exclusively because of the high melting point (>2800ºC).
However, the fuel fabrication consists of complicated solidchemical processes including a precipitation from solution,
its calcination followed by H2 reduction, and crush andgranulation to control the stoichiometry and the particle size.
Once the fuel is used, its reprocessing to extract burnableuranium and plutonium includes troublesome solution
processes using flammable organic solvents. We havefocused on hydrothermal synthesis, i.e. solid-solution systemat temperature of 100-450ºC and pressure of 0.1-22 MPa, inorder to dispense with flammable organic solvents and to
reduce the amounts of radioactive wastes.
The lower dielectric constant of the high-temperaturewater leads to equilibrium shifts from soluble metallic ions tosolid metallic oxides. The supercritical hydrothermal synthesis
was widely reported for metallic oxides with fluorite structuresuch as CeO2, ZrO2, HfO2 and ThO2. In spite of the lack ofexperimental reports for AnO2 (An = U, Np, Pu), the fluoritetype AnO2 is expected to be produced under the hydrothermal
condition.
Figure 1 shows a schematic flow-sheet of the proposedhydrothermal process [1]. The first step is the reprocessing,i.e. a removal of fission products from actinides owing to thenature of actinide ions which exclusively form stable
coordination compounds with carbonate anion (CO32-). The(1) EtOH + Hexylamine

Frequency

We have recently focused on application of hydrothermal
synthesis of actinide dioxides for fabrication and reprocessingof MOX fuel in order to improve the nuclear fuel cycle much
simpler and safer. This process enables us to reduceamounts of radioactive wastes and possible risks of fire and/or explosion, and also to control various specifications of the
MOX fuel.

(b2) size 2

0.1

0.1

(a)

0.0
0.1

2

3

4 5 6 781
size /µm

2

3

4 5 6

0.0

1

2

3

4 5 67

10
Aspect ratio

Fig. 2 Shape (a), size (b) and aspect ratio (c) of UO2 particles
obtained by hydrothermal process with two additives.

second step is the combination of the reprocessing and thefabrication of MOX fuel, i.e. the hydrothermal process at
450ºC for 30 minutes with a reductant. This process separates
actinide oxides (MOX fuel) from cesium ion down to 40 ppm.
When only ethanol is added as a reductant, the product of
UO2 crystal is indefinite, whereas the second additive suchas hexylamine, hexylaldehyde, or guanidinium ammonium
alters a fashion how the crystal grows of UO2, i.e. the shapeand the size (Fig. 2). Simultaneously, the additives govern
the reductive atmosphere and determines the nonstoichiometry x of UO2+x in the range of 0.05-0.2. The detailed
chemistry in the control of the specifications of AnO2 is now
underway.
References[1] T. Yamamura, et al., Preparation method of crystals ofoxides of actinides, PCT/JP2009/055458 (2009).
Key Words
Stoichiometry, Actinide Dioxides, Crystal Growth

1st step:
Dissolution to Carbonate Solution
Solid Phase
Liquid Phase

Actinides (U, Np 0.7, Pu 11.0, Am 0.5)
Values:Alkaline Metals (Cs 3.7, Rb 0.5) Ratios
toU /103 ppm

Rare Earths (La, ..., Lu, Sc, Y)
Alkaline Earths (Ca, Sr, Ba)
mainly Solid Phase

Others (Zr, Mo, Re)
Noble Metals (Ru, Pd)

Contact toTomoo Yamamura (Radiochemistry of Metals Division)
E-mail: yamamura@imr.tohoku.ac.jp
URL: http://www.shiokawa.imr.tohoku.ac.jp/

2nd step:
Hydrothermal process
Solid Phase

Actinide Oxides

Liquid Phase

Alkaline Metals
Cs: 40 ppm

Fig. 1 Hydrothermal process for MOX fuel fabrication andreprocessing.

16

Energy-related Materials

IMR KINKEN Research Highlights 2010

Research

Production of Ni-Cr-P-B Metallic Glass-coated Bipolar Plate forFuel Cell by HVOF Spray-coating Method
In this study, the newly designed bipolar plate for proton
exchange membrane fuel cells (PEMFC) was produced byspray-coating the Ni65Cr15P16B4 metallic glass on Al plate
having a bipolar plate flow field and the electricity generation
tests with the single fuel cell having the glass-coated biplar
plates were conducted. As a result, the single cell with themetallic glass-coated bipolar plates showed very high I-V
performance as well as the cell with the carbon bipolar
plates.
Proton exchange membrane (PEM) fuel cell has been
widely recognized as one of the most powerful candidates fornear future power generating devices of the automobiles andthe home appliances. Bipolar plates are one of the mostimportant components of PEM fuel cells and are multifunctionalas they conduct electricity from cell to cell, they separate thefuel gas from the oxidant gas, and their flow field supplies thegases to the electrodes.
Metallic glassy alloys are also potential materials forbipolar plates because they have many advantageous
characteristics over crystalline alloys. For example, glassy
alloys exhibit higher mechanical strengths and higher
corrosion resistance. We have developed the Ni-based
metallic glassy alloys suitable for bipolar plate application [1,2].
In this work, we adopted the high-corrosion-resistant NiCr-P-B metallic glass for bipolar plate production andproduced a new metallic glass-coated bipolar plate by thehigh velocity oxy-fuel (HVOF) spray-coating method andexamined the electricity generation properties of the single
fuel cell with those metallic glass-coated bipolar plates.
Figure 1 shows the outer view of the Ni65Cr15P16B4
metallic glass-coated Al bipolar plate produced by HVOF
spray-coating and subsequent hot-pressing. The glass film
was deposited on both sides of the plates.
Figure 2 shows the I-V curves of the single fuel cell with

Fig. 2 I-V curves of the single cell with carbon, Al and
Ni65Cr15P16B4 metallic glass bipolar plates.

carbon graphite, Al and Ni65Cr15P16B4 glassy alloy bipolar
plates measured at the cell temperature of 353 K at the gas
flow rate of 0.1 L•min-1. The results measured after 50 times
repetition for aging the membranes were indicated in the
figure. At the cell voltage of 0.5V, the single cells employing
graphite, Al, and Ni65Cr15P16B4 glassy alloy bipolar plates
generated the current density of 700, 200, and 750 mA•cm-2,
respectively. As clearly seen in the figure, the single cell with
the Ni65Cr15P16B4 glass-coated bipolar plates showed better
I-V performance than that with Al bipolar plates. This isbecause the Al bipolar plates are easily subjected tocorrosion, resulting in the increase in contact electricalresistance. The I-V performance of the single cell with the
Ni65Cr15P16B4 glass-coated bipolar plates is as high as thatwith the carbon graphite bipolar plates. This means that thecorrosion resistance of the Ni65Cr15P16B4 metallic glass is sohigh that it can be applied to bipolar plates in practice.
References[1] S. Yamaura, M. Yokoyama, H.M. Kimura and A. Inoue, J.
Phys.:Conf. Ser. 144, 012001 (2009).
[2] M. Yokoyama, S. Yamaura, H.M. Kimura and A. Inoue,
Int. J. Hydrogen Energy 33, 5678 (2008).
Key Words
Metallic Glass, Bipolar Plate, Fuel Cell
Contact toShin-ichi Yamaura (NEDO Laboratory for Metallic Glass Research
Collaboration and Promotion)
E-mail: yamaura@imr.tohoku.ac.jp

Fig. 1 Outer view of the Ni65Cr15P16B4 metallic glass-coated
Al bipolar plate.

KINKEN Research Highlights 2010 IMR

Energy-related Materials

17Research

A New Method for Size Estimation of Cu Nano-precipitates in Fe
Based on Positron Quantum-dot Confinement

The main origin of the embrittlement of the nuclear reactor
pressure vessel (RPV) steels due to long-term in-service
neutron irradiation is considered to be the (sub) nanoprecipitates of Cu contained in the old RPV enhanced by theneutron irradiation. In the early stage of Cu subnanoprecipitate formation, it has been difficult to obtain the size ofthe precipitates even by high-resolution transmission electron
microscope or three-dimensional atom probe (3D-AP).
Positrons are confined in Cu subnano-precipitates in Fe
by positron affinitive trapping. The positron quantum-dot
confinement leads to smearing of positron-electron
momentum density distribution around the Fermi cut-off
momentum (pF) as shown in Fig. 1. Due to the uncertainly
principle between position and momentum, the momentum
smearing is inversely proportional to the size of the quantumdot. Therefore, we have observed the momentum smearing
by two-dimensional angular correlation of positron annihilation
radiation (2D-ACAR) and estimated the average size of the
Cu nano-precipitates where positrons are confined.

Occupation Number

(a)

(c)

Momentum distribution
sampled by positron (NACAR)

Momentum

Difference between NACAR for
Embedded nano-precipitate and bulk

(b)

Bulk
Embeddednano-precipitate

pF

pF

MomentumpF

3D-AP observation

Difference of NACAR in the aged Fe-Cu alloy
between bulk Cu and Cu nano-precipitates

A new method for size estimation of Cu subnanoprecipitates in Fe based on a positron quantum-dot-state isdeveloped. The size of Cu subnano-precipitates which are
too small to be detected even by the other technique, suchas 3D-AP, can be estimated

0.1h

0.1h
0.5h
2h

0.1h
0.5h
2h

0.5h
2h

0
Δp ∝ 1/r (r : Radius of Cu

pF

nano-precipitate)r can be estimated by Δp

0

5

10

15

Momentum [ x10-3 m0c]

Fig. 2 Difference of NACAR in the aged Fe-Cu alloy between
bulk Cu and Cu nano-precipitate. The smearing effect clearly
depends on the aging time, i.e., size of Cu nano-precipitates.
The smearing effect (Δp) is assessed quantitatively using theGaussian function convolution for NACAR of bulk Cu and thensize of Cu nano-precipitate is estimated.

Fe-Cu alloy was aged at 475°C for 0.1 to 20 h to form theCu nano-precipitates with various sizes. Figure 2 shows theaging time dependency of momentum smearing effectobserved by 2D-ACAR. The average sizes of the Cu nanoprecipitates were estimated from the smearing effect andthen compared with those by 3D-AP as also shown in Fig. 2.
Both the results agree with each other for samples aged for
1-20 hours. Furthermore, the present method gives the sizesof the Cu subnano-precipitates which are too small to be
detected even by the 3D-AP for samples aged for 0.1-2 hours[1]. Thus it is demonstrated that positron annihilation
technique is a powerful tool to estimate the size of Cu
subnano-precipitates in the early stage of formation.
Reference[1] T. Toyama, Y. Nagai, Z. Tang, K. Inoue, T. Chiba, M.
Hasegawa, T. Ohkubo and K. Hono, Tetsu to Hagane, 95,
118, (2009).
Key Words
Embedded Nano-precipitates, Positron Annihilation, Threedimensional Atom Probe0
Smearing effect (Δp )Contact toTakeshi Toyama (International Research Center for Nuclear
Materials Science)

MomentumE-mail: ttoyama@imr.tohoku.ac.jp
URL: http://www.imr-oarai.jp/

Fig. 1 (a) Electron occupation number around the Fermi
momentum, (b) Angular correlation of positron annihilation
radiation spectrum (NACAR), (c) Difference of NACAR between
bulk and embedded nano-precipitate.

18

Energy-related Materials

IMR KINKEN Research Highlights 2010

Research

Intergrown alpha-Type Phase in beta-Type Phase TmAlB4 Crystal
Studied by High-Resolution Electron Microscopy
A TmAlB4 crystal with a ThMoB4-type (beta-type)structure phase was studied by electron diffraction (ED) andhigh-resolution electron microscopy (HREM). A HREM
image clearly exhibits an intergrown lamellar structure of an
YCrB4-type (alpha-type) phase in the matrix of the beta-typephase in TmAlB4 crystal. The lamellar structure can becharacterized by a tiling of deformed hexagons, which are acommon structural unit in the alpha-type and beta-type
structures.
It is well known that the rare-earth metal borides haveyielded intriguing systems to study fundamental problems inphysics and chemistry [1]. The rare-earth (RE) metal
aluminoboride system REAlB4 has been paid much attention.
Recently, multiple magnetic transitions were reported tooccur in TmAlB4 at low temperatures below anantiferromagnetic transition temperature TN [2]. REAlB4
crystals have two different structure types; alpha- and betatype structures. Both phases are built by a unique structuralunit, which consists of six heptagonal atomic columns
containing edge-shared two pentagons (Fig. 1).
Fig. 2 (a) HREM image and (b) corresponding ED pattern, of
an intergrowth nanostructure taken with the incident electron
beam parallel to the [001] direction. Deformed hexagons in(a) correspond to the structural unit in Fig. 1.

Fig. 1 Structural unit of the deformed hexagon consisting ofedge-shared two pentagonal and six heptagonal atomic
columns.
In Fig. 2, it is unequivocally demonstrated that anintergrowth nanostructure of the alpha-type phase exists inthe beta-type matrix [3]. Based on ED and HREM observations,slivers of the alpha-type “tiling” were clearly observed to beinsinuated into the beta-type TmAlB4 structure. Characteristic
diffuse streaks in ED are consistent with the lamellar structure.
As can be seen from a tiling, there are coherent interfaces
between both phases. The crystallographic relationshipbetween the beta- and the alpha-type phases can be
expressed as (1-10)beta || (010)alpha and [001]beta || [001]alpha.
The present direct observation of the intrinsic nanostructure
indicates that the intergrowth is the origin of low temperature
anomalies in the physical properties.

KINKEN Research Highlights 2010 IMR

References[1] e.g. T. Mori, "Higher Borides", in: Handbook on the
Physics and Chemistry of Rare Earths, ed. K. A. Gschneidner
Jr., et al., 38, 105 (2008).
[2] T. Mori, H. Borrmann, S. Okada, K. Kudou, A. LeitheJasper, U. Burkhardt, and Y. Grin, Phys. Rev. B 76, 064404(2007).
[3] K. Yubuta, T. Mori, A. Leithe-Jasper, Y. Grin, S. Okada
and T. Shishido, Mater. Res. Bull., 44, 1743 (2009).
Key Words
Rare Earth Aluminoboride, Intergrowth Structure, Electron
Microscopy
Contact toKunio Yubuta (Advanced Research Center of Metallic Glasses)
E-mail: yubuta@imr.tohoku.ac.jp

Energy-related Materials

19Research

Cryogenfree 23 T Superconducting Magnet with a 7.5 T
YBa2Cu3O7 Insert
In order to develop a cryogenfree high-field
superconducting magnet generating over 20 T, a Y123coated conductor insert is used to enhance a cryogenfree
18 T superconducting magnet. The Y123 tape does not
require stainless steel reinforcement. The Y123 insert, whichis almost the same size as the current Bi2223 insert,
is designed to generate 7.5 T at 187 A. It is able to achieve
this since without stainless steel reinforcement the
number of turns can be greatly increased. A cryogenfree
superconducting magnet can achieve a static magnetic
field of 23 T, which is higher than that generated by practicalsuperconducting magnets using liquid helium.
Static high-field facilities around the world require hugeelectrical powers in the range 20 to 30 MW to power a watercooled resistive magnet of a hybrid magnet system. It isdifficult for university laboratories to afford such high electricity
costs. Superconducting technology is essential for constructing
high static field facilities that do not consume huge amounts ofpower. We introduce an improvement to the cryogenfree
superconducting magnet that involves incorporating an
YBa2Cu3O7 (Y123) coated conductor tape and enables themagnet to generate a high static field of 23 T.
We focus on the larger critical current density of
YBa2Cu3O7 (Y123) coated conductors than that of Bi2223
tape at 4.2 K. When we employ the new insert fabricated withhigh performance Y123 conductors, higher magnetic fields
over 20 T can be targeted. Although the splendid Jc properties
in strong magnetic fields over 20 T for B//c is difficult even for
the Y123 superconductor, the magnetic field applied to an
insert in a background magnet is almost parallel to thesuperconducting tape surface. This means that the c-axisparallel component is quite small in the insert in backgroundfields. We understand that the critical current characteristic in
the c-axis-perpendicular field component for Y123 is far
excellent, when Bi2223 and Y123 are compared each other.
Figure 1 shows the Ic-B characteristics for Bi2223 and Y123
at 4.2 K. Y123 coated conductors developed as practicallong-length 4 mm wide and 0.1 mm thick tapes have the Ic

Table I Comparison of the Bi2223 insert for 18T-CSM andthe Y123 insert for 23T-CSM
Superconductor

Bi2223

Supercond. tape size

mm2 4.3 × 0.21

Cowinding reinforcement mm

SUS 0.3

Y123
4.0 × 0.1
Hastelloy 0.05

Inner diameter

mm

90

90

Outer diameter

mm

176

176

Coil height

mm

252

252

No. of layers

50 (25 DP)

50 (25 DP)

Total turns

3548

9350

Operation current

A

161

187

Central field(15.6 T backup)

T

18.1

23.3

MPa

98

439

Hoop stressvalue of 80 A at 77 K in a self-field. It was found that the Y123
has splendidly high Ic potentials in fields for both B//c andB┴c directions to design the insert with high current densities.
The coil load lines are well determined to generate 23 T for
B┴c. In this case, we understand that the c-axis-parallel fieldcomponent is 3.1 T.
Table I lists the design parameters for the Y123 insert coilof the 23 T cryogenfree superconducting magnet (23T-CSM).
The coil current density was increased by a factor of overthree from 50 to 160 A/mm2 by replacing the Bi2223 coil with
the Y123 one. Reinforcement is reduced from 0.3-mm-thickstainless steel tape to 0.1-mm-thick Hastelloy tape including
the coated conductor substrate. The Y123-coated conductor
tape with Hastelloy reinforcement can overcome hugeelectromagnetic stresses of up to 1000 MPa. Therefore, it isexpected that the Y123 tape can withstand the electromagnetic
force of 440 MPa in the 23T-CSM. This means that stainlesssteel reinforcement is no longer required for the insert coil.
As a result, a Y123 insert coil of almost the same size as theBi2223 insert coil will generate 7.5 T at 187 A, because thenumber of turns without stainless steel reinforcement can begreatly increased. A cryogenfree superconducting magnet
can achieve a 23 T static magnetic field in a 52 mm
experimental bore at RT [1].
References[1] Superconducting Magnet with a 7.5 T YBa2Cu3O7 Insert
Coil, K. Watanabe, S. Awaji, G. Nishijima, S. Hanai and M.
Ono, Appl. Phys. Express 2, 113001 1-3 (2009).
Key Words
Y123 Coated Conductor, High Field, Cryogenfree
Superconducting Magnet
Contact toKazuo Watanabe (High Field Laboratory for Superconducting
Materials)

Fig. 1 Comparison of the Ic properties at 4.2 K in fields for B┴cand B//c between Ag-sheathed Bi2223 tape (Sumitomo ElectricCo.) and Y123 coated conductor tape (SuperPower Co.).

20

Energy-related Materials

E-mail: kwata@imr.tohoku.ac.jp
URL: http://hflsm.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2010

Research

Promotional Effect of Au/Pd Bimetallic Nanoparticles
for Photocatalytic Activities
Au/Pd bimetallic nanoparticles were successfully
prepared from aqueous solution of the corresponding noble
metal salts and immobilized on TiO2 by the effect of highpower ultrasonic irradiation. These nanoparticles exhibited
unique promotional effect for TiO2 photocatalysis depending
upon the inner nanostructure even if they had same content
of Au/Pd ratio.
Designing nano-morphologies is expected to be one ofthe promising approaches to develop the novel highperformance catalysts. The author reported the preparationof bimetallic nanoparticles consisting of Au-core and Pd-shellusing unique methodology utilizing the high energy reaction
sites induced by high power ultrasonic irradiation [1]. The
inner structures of alloy nanoparticles such as core/shellstructure have been receiving increasing attention to enhance
the catalytic activities of the particles with high performanceas compared with that of the constituent metals. In this study,
the photocatalytic activities of TiO2 supporting Au/Pd
bimetallic nanoparticles prepared by ultrasonic irradiation
were evaluated.
Au-core/Pd-shell bimetallic nanoparticles dispersions areprepared by the previously reported sonochemical method[1]. TiO2 (P-25) was added to the dispersion of
the nanoparticles, and followed by the sonication (200 kHz,
6 W/cm2) under air atmosphere to immobilize noble metal
nanoparticles on the surface of TiO2 (sample A) [2].
According to the HAADF-STEM image, nanosized Aucore/Pd-shell nanoparticles with multiple twined structure
were confirmed to be immobilized on TiO2 surface (Fig. 1(a)).
The mixtures of Au and Pd monometallic nanoparticles
individually prepared were also immobilized by the sameprocedure (sample B). On the other hand, by annealing
sample A, the core/shell structure was broken and changed
into ordered state where Au and Pd atoms were
homogeneously located in each particle (sample C, shown in
Fig.1 (b)) [3]. Neither sintering of the supported nanoparticles
nor growth in their sizes was not observed after the annealing.
ICP analyses showed that the contents were similar between
sample A, B, and C.
Photocatalytic H2 evolution was employed for evaluation

Table 1 H2 evolution by sonochemically prepared photocatalysts
SampleH2 / mmol (h∙g-cat (mW∙cm-2))-1
UV illuminationa

Vis illuminationb

A (Au-core/Pd-shell)

140.0

17.4

B (Au-Pd mixture)

74.7

10.2

C (Annealed)

213.1

13.0

a 250 – 385 nm. b 385 – 740 nm.

for the samples. Ar-purged ethanol aqueous solutions
containing the prepared photocatalysts were illuminated by a
Xe lamp in a closed glass vessel with a quartz window. The
amounts of evolved H2 depended on the wavelength of theilluminated light and the nanostructure of the supported
nanoparticles (Table 1). These results indicate that thebimetallic nanoparticles (samples A and C) effectively trap
the photo-excited electrons compared with the mixture(sample B). Moreover, XPS studies suggested the change inthe band structures depending on the inner structures of theimmobilized bimetallic nanoparticles [4].
These results suggest that designing inner structure of
the bimetallic nanoparticles provides new functions differingfrom intrinsic features of constituting elements.
References[1] Y. Mizukoshi, T. Fujimoto, Y. Nagata, R. Oshima, Y.
Maeda, J. Phys. Chem. B, 104, 6028-6032 (2000).
[2] Y. Mizukoshi, Y. Makise, T. Shuto, J. Hu, A. Tominaga, S.
Shironita and S. Tanabe, Ultrason. Sonochem., 14, 387-392(2007).
[3] T. Nakagawa, H. Nitani, S. Tanabe, K. Okitsu, S. Seino,
Y. Mizukoshi and T. A. Yamamoto, Ultrason. Sonochem., 12,249-254 (2005).
[4] Y. Mizukoshi, K. Sato, T. J. Konno and N. Masahashi.,Appl. Catal. B Environ., 94, 248–253 (2010).
Key Words
Catalyst, Bimetallic Nanoparticle, UltrasoundContact toYoshiteru Mizukoshi (Osaka Center for Industrial Materials
Research)
E-mail: mizukosi@imr.tohoku.ac.jp

(a)

(b)

Pd shell

URL: http://polar.imr.tohoku.ac.jp

TiO2
Au-Pd

Au core

TiO2

2nm

2nm

Fig. 1.1 HAADF-STEM
images ofimages(a) as preparedand (b)
Au/Pd andbimetallic
Fig.
HAADF-STEMof (a)
asannealed
prepared(b)
nanoparticles immobilized on TiO2. Schematic views of the product are also shown.
annealed Au/Pd bimetallic nanoparticles immobilized on
TiO2.

KINKEN Research Highlights 2010 IMR

Energy-related Materials

21