KINKEN

Research Highlights
2013

1986. February
Courtesy of Tohoku University Historical Museum.

Institute for Materials Research,Tohoku University

KINKEN Research Highlights 2013
Contents
Preface���������������������������������������������������������������������������������������������������������������������������������������������������������� 1

Research Highlights

1. Infrastructural Materials
Microstructure Design of Structural Metallic Materials Division�������������������������������������������������������������������� 3
Chemical Physics of Non-Crystalline Materials Division������������������������������������������������������������������������������ 5
Biomaterials Science Division���������������������������������������������������������������������������������������������������������������������� 7
Crystal Chemistry Division��������������������������������������������������������������������������������������������������������������������������� 9
Deformation Processing Division����������������������������������������������������������������������������������������������������������������11
Advanced Analysis of Materials Division���������������������������������������������������������������������������������������������������� 13
Computational Materials Research Initiative���������������������������������������������������������������������������������������������� 15

2. Energy-related Materials
Surface and Interface Research Division��������������������������������������������������������������������������������������������������� 17
Metal Physics with Quantum Beam Spectroscopy Division����������������������������������������������������������������������� 19
Physics of Crystal Defects Division������������������������������������������������������������������������������������������������������������ 21
Materials Design by Computer Simulation Division������������������������������������������������������������������������������������ 23
Nuclear Materials Science Division������������������������������������������������������������������������������������������������������������ 25
Nuclear Materials Engineering Division������������������������������������������������������������������������������������������������������ 27
Non-Equilibrium Materials Division������������������������������������������������������������������������������������������������������������� 29
Hydrogen Functional Materials Division����������������������������������������������������������������������������������������������������� 31
Advanced Crystal Engineering Laboratory������������������������������������������������������������������������������������������������ 33
Analytical Science Division������������������������������������������������������������������������������������������������������������������������ 35

3. Electronic Materials
Theory of Solid State Physics Division������������������������������������������������������������������������������������������������������� 37
Magnetism Division������������������������������������������������������������������������������������������������������������������������������������ 39
Low Temperature Condensed State Physics Division�������������������������������������������������������������������������������� 41
Irradiation Effects in Nuclear and Their Related Materials Division������������������������������������������������������������ 43
Physics of Electronic Materials Division����������������������������������������������������������������������������������������������������� 45
Magnetic Materials Division������������������������������������������������������������������������������������������������������������������������ 47
Multi-Functional Materials Science Division����������������������������������������������������������������������������������������������� 49
Radiochemistry of Metals Division������������������������������������������������������������������������������������������������������������� 51
Materials Processing and Characterization Division���������������������������������������������������������������������������������� 53
Advanced Materials Development and Integration of Novel Structured Metallic and Inorganic Materials���� 55

Research Centers
International Research Center for Nuclear Materials Science�������������������������������������������������������������������� 57
Cooperative Research and Development Center for Advanced Materials(Formerly Advanced Research Center of Metallic Glasses)������������������������������������������������������������������� 59
High Field Laboratory for Superconducting Materials�������������������������������������������������������������������������������� 61
Kansai Center for Industrial Materials Research���������������������������������������������������������������������������������������� 63
Center for Computational Materials Science���������������������������������������������������������������������������������������������� 65
International Collaboration Center (ICC-IMR)��������������������������������������������������������������������������������������������� 67
Integrated Materials Research Center for a Low-Carbon Society(LC-IMR)����������������������������������������������� 69
Center of Neutron Science for Advanced Materials����������������������������������������������������������������������������������� 71
Laboratory of Low Temperature Materials Science������������������������������������������������������������������������������������ 73
Laboratory of alpha-Ray Emitters��������������������������������������������������������������������������������������������������������������� 75
Analytical Research Core for Advanced Materials������������������������������������������������������������������������������������� 77
Research and Development Center for Ultra High Efficiency Nano-crystalline Soft Magnetic Material���� 79

Author Index����������������������������������������������������������������������������������������������������������������������������������������������� 82

Keyword Index������������������������������������������������������������������������������������������������������������������������������������������� 83

Preface

Dear colleagues,

Professor Kotaro Honda established a steel research organization officially known as “the 2nd
Division of the Provisional Institute of Physical and Chemical Research” at Tohoku University in 1916. In
1987, the organization was restructured into its present form as a national collaborative research instituteaffiliated to Tohoku University and was consequently renamed “the Institute for Materials Research”(IMR), also known as “KINKEN” (which is the short form for “kinzoku zairyo kenkyujo,” the Japanesename for IMR).

Since its establishment, IMR has become known for its excellence in both fundamental and appliedresearch on metals and on a wide range of new materials. Numerous functional materials have beenstudied and developed at IMR, including Kichizaemon Sumitomo (KS) magnetic steel, new KS magnetic
steel, SiC fibers, compound alloys, metallic glasses, and others.

IMR has greatly contributed to the advancement of materials science. As one example, our study ofmagnetism resulted in the invention of the world’s strongest permanent magnet, the KS magnet, which
dramatically improves the performance of electrical machines. IMR has historically focused on
fundamental research to develop applications for a wide range of advanced technologies. More recently,
IMR has developed a wide array of materials including high-performance, high-quality, and multifunctional
materials such as amorphous alloys, bulk metallic glasses, intermetallic compounds, quasicrystals,oxides, ceramics, nanostructured Si and Ge, optical and electronic materials, solar cell crystals,
biomaterials, organic materials, hydrogen storage alloys, shaped crystals, and III-V, II-VI, and oxide
semiconductors.

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IMR KINKEN Research Highlights 2013

Director

IMR has completely recovered from the Great East Japan Earthquake that occurred on March 11,
2011, and we are working hard to achieve innovative research and development in the fields of materials
science and engineering. KINKEN will celebrate its 100th anniversary on May 21, 2016. We would like tointroduce the logo for KINKEN’s 100th anniversary, which has been selected from among 726 candidates.
We have redefined our missions and are establishing new roles for KINKEN in the post 3/11 era.

We are pleased to introduce KINKEN Research Highlights 2013, the seventh annual report containing
a collection of the research output during the previous year at KINKEN. We decided in 2007 to annually
prepare this report so that our colleagues around the world can be kept informed about our research
activities. The studies conducted at our individual research laboratories and research centers and those
jointly conducted with other academic institutions are provided, especially those in the key research
fields of infrastructural, energy-related, and electronic materials.

The world faces serious issues such as the deterioration of the environment and the depletion ofnatural resources and nonrenewable energy sources. IMR will promote research into materials to further
address the problems associated with these issues with the aims of sustaining human development andensuring a high standard of living for all people. In this way, we can offer a brighter future to subsequent
generations.

We hope that the KINKEN Research Highlights will enable you to better understand and support our
research activities and will aid the promotion of worldwide collaboration with our institute in the field ofmaterials research. We ask for your continued support and welcome any suggestions.

IMR KINKEN Research Highlights 20132

Infrastructural Materials

Mechanical Properties of Ferrite-martensite Dual Phase
Steels with Interphase PrecipitationF

errite-martensite dual phase steels with interphase precipitation of nano-sized vanadium carbides wereprepared and mechanical properties were systematically investigated. It was found that nano-sized vanadium
carbides are effective for improving the post-uniform elongation as well as increasing the strength of dual
phase steels.
Introduction
It is well known that ferrite-martensite dual
phase (DP) steels exhibit high strength and good
formability and has widely been in practical use [1].
Stress-strain curves of ferrite-martensite DP steels
are characterized by large work-hardening anduniform elongation but small post-uniform elongation.
On the other hand, interphase precipitated steels,composed of single phase ferrite and dispersion of
nano-sized alloy carbides in ferrite, are also known
as commercial steels having a good balance of highstrength and formability [2]. Stress-strain curves ofinterphase precipitated steels are such that work
hardening and uniform elongation are relatively low
but post-uniform elongation is quite large [3]. Thatis, ferrite-martensite DP steels and interphase
precipitated steels have both good and bad points.
It may be possible to significantly improve thestrength and ductility if nano-sized alloy carbides
are dispersed in ferrite phase of ferrite-martensite
DP steels, combining advantages of both steels,
which is a motivation for the present study.
In this study, ferrite-martensite DP samples withdispersion of nano-sized vanadium carbides (VC) in
ferrite phase were produced using ferrite
transformation accompanied with interphase
precipitation followed by water quenching in aV-added low carbon steel, and the mechanicalproperties of such DP samples were systematicallyinvestigated to understand the effect of nano-sized
precipitates on the strength and ductility of DP
steels.

Experimental
A low carbon steel with a chemical compositionof Fe-0.095%C-1.49%Mn-0.43%V (in mass%)(V-added steel) was used in the present study. After
austenitizing at 1100oC for 600 s, samples were
isothermally transformed at 690oC for different
holding periods followed by quenching, leading toferrite-martensite DP samples accompanied withinterphase precipitation of VC in ferrite. Conventional
ferrite-martensite DP samples with no VC particles
were also prepared using Fe-0.098%-1.5%Mn(V-free steel) for comparison. Microstructures werecharacterized by optical microscopy, transmission
electron microscopy (TEM) and electron backscatter
diffraction (EBSD). Tensile tests were carried out atroom temperature at an initial strain rate of 8.3 x 10 -4
s-1.
Results and Discussion
By using isothermal ferrite transformation
followed by water quenching, ferrite-martensite DP
samples with different volume fraction of ferrite has
successfully been obtained for both V-added andV-free steels. TEM observation has confirmed that
VC precipitates with an average diameter ofapproximately 10 nm were distributed in ferrite
phase of V-added DP samples. Tensile tests werecarried out for the obtained DP samples, andstrength and elongation were determined fromstress-strain curves. Fig. 1 shows the strength andductility of the samples as a function of volume
fraction of ferrite. It is seen that the strength ofFig. 1 Strength and ductility of V-free and V-added steels. (a) 0.2% proof stress and ultimate tensile strength, (b) uniform
elongation and (c) post-uniform elongation.

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IMR KINKEN Research Highlights 2013

IMR KINKEN Research Highlights 2013

Fig. 2 Strength-ductility balance of V-free and V-added steels. (a) Ultimate tensile strength vs uniform elongation and (b)ultimate tensile strength vs post-uniform elongation.

V-added DP samples is higher than that of V-free
DP samples, which should be due to precipitation
strengthening of ferrite by nano-sized VC particles.
Uniform elongation of the samples decreases byaddition of vanadium, but post-uniform elongation isslightly larger in the V-added DP samples than thatin the V-free samples.
Strength-ductility balance is compared between
V-added and V-free samples in Fig. 2. The balanceof ultimate tensile strength and uniform elongation in
Fig. 2a shows that the balance curve is shifted
toward high-strength and low-ductility region by V
addition, but the balance of both V-added and V-freesamples are located in the single curve. This result
indicates that V addition increases the strength of
DP samples with keeping a good balance of ultimate
tensile strength and uniform elongation. On theother hand, it is clearly seen that V addition improves
the balance of strength and post-uniform elongationin the DP samples.
To understand the difference in the mechanicalproperties of the DP samples between V-free andV-added steels, the work hardening behavior wascompared. It was found that the work hardening
rate during tensile test for ferrite-martensite DP
samples decreases by addition of VC. This can
explain the reason why the uniform elongation
decreased by addition of V. However, this result issomewhat surprising since nano-sized alloy
carbides are reported to have an effect to increase
the working hardening in the tensile test of single
phase ferrite steels [3]. To further understand themechanisms behind, digital image correlation (DIC)

analysis was applied to investigate the strain
distribution introduced in the DP samples during
tensile tests. The DIC analysis revealed the following
results. In the V-free steel, strain distribution wasquite heterogeneous, and in particular the strain
was more concentrated on soft ferrite phase.
However, it was clearly shown that the strain
distribution became more uniform in the V-added
steel. These tendencies suggest that precipitation
strengthening of ferrite by VC particles reduces thedifference in strength between ferrite and martensite
phase in DP samples, which would suppress
inhomogeneous deformation during the tensile test.
This corresponds well with the decreased work
hardening and smaller uniform elongation in the
V-added DP samples.
More homogeneous
deformation caused by VC dispersion may also
suppress formation of voids or cracks in the laterstage of tensile deformation. This would be thereason why the V addition led to an increase in thebalance of strength and post-uniform elongation. It
can be suggested that nano-sized alloy carbides
are effective for improving the post-uniform
elongation as well as increasing the strength of DP
steels.
References[1] “Fundamentals of Dual Phase Steels”, edited by R. A. Kot and
B. L. Bramfitt, AIME, (1981).
[2] Y. Funakawa, T. Shiozaki, K. Tomita, T. Yamamoto, and E.
Maeda: ISIJ International, 44, 1945-1951(2004).
[3] N. Kamikawa, Y. Abe, G. Miyamoto, Y. Funakawa, and T.
Furuhara: Tetsu-to-Hagané, 99, 352-361(2013).

Keywords: steel, mechanical properties, deformation
Tadashi Furuhara, Goro Miyamoto, and Naoya Kamikawa (Microstructure Design of Structural Metallic Materials Division)
E-mail: furuhara@imr.tohoku.ac.jp
URL: http://www.st-mat.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 20134

Infrastructural Materials

Structure of Hydrated Copper silicate gelsH

ydrated copper silicate gel samples were prepared using the conventional sol-gel method and stock
solutions whose Cu/Si molar ratios were different. The structures of the obtained gel samples were analyzed
using X-ray diffraction and extended X-ray absorption fine structure (EXAFS) spectroscopy. A common mediumrange structure similar to that of natural chrysocolla was produced in each gel sample. The local structure ofchrysocolla is readily produced in during sedimentation from any solutions containing CuO and SiO2
components.
Chrysocolla, Cu2H2(Si2O5)(OH)4∙nH2O, is asecondary mineral that forms in the oxidized zones
of copper deposits (Fig. 1). Because of the unique
Cu/Si molar ratio for natural chrysocolla, it isreasonable to assume that the production of natural
chrysocolla is associated with a regular chemical
reaction during sedimentation from solutions
containing copper and silica components, andunderstanding the mechanism of the mineralizationof chrysocolla is expected to provide useful
information on its cryptocrystalline structure. This
prompted us to study the disordered structure of asynthetic analogue of chrysocolla prepared using
the sol-gel method [1].
The sol-gel method and reagents CuSO4∙5H2O
and Si(OC2H5)4 (TEOS) were used to prepare
hydrated copper silicate gels with three different Cu/
Si molar compositions. The Cu/Si molar ratios forthe three obtained gel samples were 0.32, 0.83, and1.61 (hereafter named Cu1Si2, Cu1Si1, and Cu2Si1,
respectively).
The thermogravimetric (TG) curves for Cu1Si2,
Cu1Si1, Cu2Si1, and natural chrysocolla show
relatively smooth weight losses of 19.0, 22.0, 24.3,
and 20.2%, respectively. The differential thermal
analysis (DTA) curves for Cu1Si1 and natural
chrysocolla show exothermic peaks at ~680°C,
which correspond to thermal decomposition during
the formation of crystalline CuO. Although exothermic
peaks are not clearly shown in of the DTA curves forFig. 1 Chrysocolla from Chuquicamata, Chile.

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IMR KINKEN Research Highlights 2013

Fig. 2 X-ray diffraction patterns of three gel samples. The
pattern of chrysocolla from Chuquicamata, Chile and
SiO2 glass are also shown.

Cu1Si2 and Cu2Si1, evidence of a similar reaction
associated with the formation of CuO was found byheating Cu1Si2 and Cu2Si1 at 1000°C. This thermal
behavior is different from that observed for Cu(OH)2(spertiniite). Therefore, each of the synthetic gel
samples is not a simple assemblage of spertiniite
and amorphous or crystalline silica but it also
includes a substance similar to natural chrysocolla.
The X-ray diffraction (XRD) patterns for thethree gel samples and for natural chrysocolla areshown in Fig. 2. Each pattern consists of a few
diffuse maxima, and several broad peaks aresuperimposed on them. The profiles for Cu1Si2 andCu1Si1 indicate broad peaks at Q = 3.8, 4.2, and 4.6
Å-1, which correspond well with those for natural
chrysocolla. In particular, the XRD pattern for Cu1Si1
is quite similar to that for natural chrysocolla,suggesting the Cu1Si1 gel is isostructural withnatural chrysocolla. The first diffuse maximum in the
XRD pattern for Cu1Si2, on the other hand, isattributed to the structure of ordinary amorphous
SiO2. This feature enables us to consider that part of
the SiO2 in Cu1Si2 is encouraged to form amorphous
SiO2.
Fig. 3 shows the reduced distribution functions(RDFs) for the Cu1Si2, Cu1Si1, and Cu2Si1 samples
and for SiO2 glass and natural chrysocolla. The
atomic correlation peaks for every pair of constituentsIMR KINKEN Research Highlights 2013

Fig. 3 Reduced atomic pair distribution functions of thegels, chrysocolla from Chuquicamata, Chile, and SiO2
glass

Fig. 4 Cu K-edge k2-weighted X(k) functions of the gelsand chrysocolla from Chuquicamata, Chile.

are clearly observed up to ~15 Å for Cu2Si1, Cu1Si1,
Cu1Si2 and natural chrysocolla. However, the atomic
correlation for SiO2 glass becomes less distinct withincreasing radial distance and disappears over 10
Å. Since the oscillations in the reduced RDFs well
represent the degree of structural ordering, theobtained Cu1Si2, Cu1Si1, and Cu2Si1 gel samples
show structures whose middle-range orderings aredifferent from those for common glasses. The
dashed lines in Fig. 3 indicate the atomic distances
for the Si-O bonds in the tetrahedral unit and the
Cu-O bonds in the square planar unit, respectively.
The first nearest-neighbor correlation for each gel
represents a harmonization of the Si-O and Cu-O
correlations and suggests the existence of the SiO4
unit.
Fig. 4 shows Cu K-edge extended X-ray
absorption fine structure EXAFS χ(k) functions forthe synthetic gel samples and for natural chrysocolla.
The spectra for Cu1Si2, Cu1Si1, and natural
chrysocolla feature oscillation peaks at k = 7.5 and10 Å-1, while those peaks are absent in the spectrum
for Cu2Si1. This result indicates that the local
structures around Cu in Cu1Si2 and Cu1Si1 aresimilar to that of natural chrysocolla and that Cu2Si1
has a unique Cu environment.
Copper silicate gels were synthesized throughsol-gel precipitation from starting materials whose
Cu/Si molar ratios were different. The TG-DTA, XRD,
and EXAFS measurements indicated that middle-

range ordering similar to that of natural chrysocolla
was formed in each of the gel samples despite thevariation in the Cu/Si molar ratio of the starting stock
solutions. In particular, the structural feature of the
Cu1Si1 gel sample is equivalent to that of natural
chrysocolla. Therefore, the nearly stoichiometriccomposition of cryptocrystalline chrysocolla strongly
associates with its short- and medium-range order,
and this unique ordering is formed as a result of achemical reaction during precipitation from solutions
containing SiO2 and CuO components. The local
structural information about the CuO4 units aroundCu together with the common SiO4 units observedin the Cu1Si1 gel and natural chrysocolla enables us
to consider middle-range ordering such as thatfound in crystalline dioptase (CuSiO3∙H2O) andplanchéite. However, the linkages of the CuO4 andSiO4 units in chrysocolla are still unknown. Further
systematic investigation with methods such as highenergy X-ray diffraction is strongly required in orderto obtain high-resolution pair distribution function
data.
References[1] T. Hariu, H. Arima, and K. Sugiyama, J. Mineral. Petrol. Sci.,
108, 111-115 (2013).

Keywords: cryptocrystalline, pair distribution function, XAFS
Hiroshi Arima, Kazumasa Sugiyama (Chemical Physics of Non-Crystalline Materials Division)
E-mail: arimah@imr.tohoku.ac.jp, kazumasa@imr.tohoku.ac.jp
URL: http://www.xraylab.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 2013

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Infrastructural Materials

Improvement in Hard-Tissue Compatibility of β-type Titanium
Alloy by Formation of Hydroxyapatite through Alkali Treatment
and Metal Organic Chemical Vapor Deposition

The hard-tissue compatibility of the β-type Ti alloy Ti-29Nb-13Ta-4.6Zr (TNTZ) was improved using two
methods. The effects of the alloying elements Ta, Nb, and Zr on the hydroxyapatite (HAp) formability of Ti alloys
soaked in simulated body fluid after alkali treatment were investigated to improve the formability of HAp on
TNTZ-based implants for the human body. Furthermore, the adhesive strengths of HAp films deposited throughmetal organic chemical vapor deposition onto TNTZ substrates were improved by controlling surface
morphology of TNTZ substrates.

In biomedical applications, bone absorption
and degradation of bone quality are induced bystress shielding resulting from mismatch between
the Young’s moduli of living bone tissue and implant
devices. Therefore, a β-type Ti alloy, Ti-29Nb-13Ta4.6Zr (mass%, TNTZ) [1], has been developed witha low Young’s modulus to overcome this significant
problem. The Young’s modulus of TNTZ isapproximately 60 GPa [1], which is much closer tothat of cortical bone (10–30 GPa) than to those ofexisting Ti alloys. However, Ti and its alloys fuse withbone through fibrous tissues, and it takes a longtime for them to do so[2]. Therefore, it is necessary
to improve the hard-tissue compatibility of TNTZ.
A coating of hydroxyapatite (Ca10(PO4)6OH2;
HAp) is very effective in improving the hard-tissue

Fig. 1 SEM images of Ti-xM (nominal content x = 10, 20,
30, or 40 mass%; M = Nb, Ta, or Zr) alloy and each puremetal soaked in SBF for 1 week after alkali treatment in
1 M NaOH solution at 363 K for 259.2 ks.

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IMR KINKEN Research Highlights 2013

compatibility of Ti and its alloys. HAp directly bonds
with bone tissue without any fibrous cells forming. In
the present study, two methods were used to form
HAp on TNTZ in order to improve the hard-tissuecompatibility of TNTZ.
The first method of forming HAp on the TNTZ
surface is alkali treatment. HAp forms on the surfacesof pure Ti and its alloys immersed in a NaOH solution
and then in simulated body fluid (SBF) whose ion
concentrations are almost the same as those in thehuman body. However, the HAp formability of TNTZ
soaked in SBF after alkali treatment is lower thanthat of pure Ti that has been subjected to the same
treatment conditions. Therefore, it is important tounderstand the relation between the alloying
elements Nb, Ta, and Zr and HAp formability inorder to improve the HAp formability of TNTZ. In thisstudy, the HAp formability of Ti-xM (nominal content
x = 10, 20, 30, or 40 mass%; M = Nb, Ta, or Zr) andof each pure metal soaked in SBF after alkali
treatment was investigated [3].
Fig. 1 shows scanning electron microscope(SEM) images of the surfaces of (a) Ti-xNb and pureNb, (b) Ti-xTa and pure Ta, and (c) Ti-xZr and pureZr soaked in SBF for 1 week after alkali treatment in
1 M NaOH solution at 363 K for 259.2 ks [3]. The
HAp coverage on the Ti-xTa surfaces is almost
constant at ~95%; however, the coverage on the TixNb and Ti-xZr surfaces decreased from 90–95%
to 0% as the Nb and Zr contents increased from 10
to 100 mass% (i.e., pure Nb and Zr metals).
The X-ray diffraction (XRD) profiles for the alkalitreated Ti-xTa showed that Na2Ti5O11 and Na5TaO5
had formed on the surfaces (data not shown here)[3]. The XRD profiles for the alkali-treated Ti-xNb and
Ti-xZr, on the other hand, showed that only Na2Ti5O11
had formed on the surfaces (data not shown here)[3]. The Ti-Nb and Ti-Zr alloys showed good HApformability when the Nb and Zr contents were lower
than 10 mass% because enough sodium titanate waspresent on the surfaces. However, when the Nb or Zr
content was increased to 20 mass%, the number of
Ti atoms on the surfaces decreased, resulting in a

IMR KINKEN Research Highlights 2013

Fig. 2 AFM images of (a) TNTZ substrate surface withmirror-like finish and TNTZ substrates polished with (b)#2400-, (c) #1500-, and (d) #1000-grit abrasive papers.

decrease in the sodium titanate content on thesurfaces. These results suggest that the sodium
titanate and sodium tantalate on the surfaces were
responsible for HAp nucleation. Therefore, it isimportant to maximize the formation of sodium titanateand sodium tantalate on the TNTZ surface in order toimprove the hard-tissue compatibility of TNTZ.
The second method of forming HAp on the
TNTZ surface is dry coating; namely, metal organic
chemical vapor deposition (MOCVD), which is a
useful method of coating HAp films onto metal
surfaces. This study was undertaken with the aim of
improving the adhesive strength of HAp films
deposited onto TNTZ substrates through MOCVD
by optimizing the surface morphology; i.e., thesurface roughness and size of asperities of the
TNTZ substrates [4].
Fig. 2 shows atomic force microscopy (AFM)
images and surface roughness (Ra) values for TNTZ
substrates with a mirror-like surface finish and for
those polished with #1000-, #1500-, and #2400-gritabrasive papers [4]. The surfaces of the TNTZ
substrate with a mirror-like finish and that polishedwith #2400-grit abrasive paper are relatively smooth.
The surfaces of the TNTZ substrates polished with
#1500- and #1000-grit abrasive papers, on theother hand, are uniformly rough and show
nanoasperities. Microasperities are also observedon the surface of the TNTZ substrate polished with

Fig. 3 Adhesive strengths of HAp films deposited onto
TNTZ substrate with mirror-like surface finish and TNTZ
substrates subjected to mechanical polishing using#2400-, #1500-, and #1000-grit abrasive papers.

the #1000-grit abrasive paper.
Fig. 3 shows the adhesive strengths of the HAp
films deposited onto the TNTZ substrate with amirror-like surface finish and onto those whose
surfaces were polished with #1000-, #1500-, and#2400-grit abrasive papers [4]. The adhesivestrength of the HAp film deposited onto the TNTZ
substrate with a mirror-like surface finish is
approximately 10 MPa. Mechanical polishing using#1000-, #1500-, and #2400-grit abrasive papers
improved the adhesive strengths of the HAp films.
The surface of the TNTZ substrate polished with the#1000-grit abrasive paper is rougher than that of the
TNTZ substrate polished with the #1500-grit abrasive
paper. However, the adhesive strengths of the HAp
films deposited onto these two substrates are nearly
equal. These results indicate that the nanoasperitie
play a more important role in improving the adhesivestrength of the HAp films than the surface roughnessof the substrates does.
References[1] D. Kuroda, M. Niinomi, M. Morinaga, Y. Kato, and T. Yashiro,
Mater. Sci. Eng. A 243, 244 (1998).
[2] M. Niinomi, T. Hattori, K. Morikawa, T. Kasuga, A. Suzuki, H.
Fukui, and A. Niwa, Mater. Trans. 43, 2970 (2002).
[3] K. Cho, M. Niinomi, M. Nakai, J. Hieda, and X. Tao, Mater.
Trans. 54, 1295 (2013).
[4] J. Hieda, M. Niinomi, M. Nakai, K. Cho, T. Gozawa, H. Katsui,
R. Tu, and T. Goto, J. Mech. Behav. Biomed. Mater. 18, 232(2013).

Keywords: biomedical, alloy, hydroxyapatite
Mitsuo Niinomi ( Biomaterials Science Division)
E-mail: niinomi@imr.tohoku.ac.jp
URL: http://biomat.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 20138

Infrastructural Materials

Impurity Partitioning During Colloidal Crystallization

Colloidal crystal was used as a model of impurity partitioning during the growth of a bulk crystal. The detailed

partitioning behavior of the colloidal crystal was investigated for the first time. The effective partition coefficient,keff (= CS/C0), was investigated for pure polystyrene and for polystyrene dyed with fluorescent particles bychanging the ratio of the particle diameters, Dimp/Dcryst, and growth rate, V. keff was less than unity and increased
to unity with increasing V for each size of polystyrene particles, whereas keff increased to unity as Dimp/Dcryst
approached unity at a given growth rate. These results were consistent with the solute behavior analyzed usinga Burton-Prim-Slichter (BPS) model. The equilibrium partition coefficient, k0, obtained from the BPS plotincreased as Dimp/Dcryst approached unity. In contrast, although the fluorescent particles also followed the BPS
model, they showed higher k0 values than the counterpart polystyrene particles did. k0 values greater thanunity were obtained for the impurities whose sizes were similar to that of the host particles. This behavior is
attributed to the positive free energy of fusion associated with incorporating the fluorescent particles into thehost matrix. Such positive free energy of fusion implies the presence of the enthalpy associated with theinteractions among the particles.
Colloidal crystals are a promising tool forinvestigating diverse basic physical phenomena.
We applied a colloidal crystal to impurity partitioning
because impurity partitioning is one of the most
universal phenomena observed during the growthof crystals from melts, and it strongly affects theproperties of bulk crystals. Since no research has
previously been conducted on the impurity
partitioning of colloidal crystals, this work investigates
the detailed partitioning behaviors such as thegrowth rate, size of impurity particles, and types of
impurities formed during colloidal crystallization [1].
A small amount of impurities (2%) was doped

Fig. 1 A. Few tens of layers of colloidal crystals grown using
convective assembly. Cleavage surfaces were observed
using SEM to obtain concentration of impurity particles insolid (CS). B. Impurity particles substituted for constituent
particles. C. Effective partition coefficients (keff) for
various sizes of PS and w/fluor. particles. D. Equilibrium
partition coefficients (k0) for PS and w/fluor. particles.

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IMR KINKEN Research Highlights 2013

into a colloidal dispersion from which colloidal
crystals were grown. Various sizes of purepolystyrene (PS) particles (i.e., 300, 400, 600, and
700 nm) and fluorescent bearing (w/fluor.) particles(i.e., 300, 400, 520, 600, and 700 nm) were doped
as impurity particles into the colloidal crystals.
Convective assembly [2] was used to grow thecolloidal crystals. More than several tens of layers of
colloidal crystals were formed on a glass plate (Fig.
1A). The cleavage surface was observed using
scanning electron microscopy (SEM) to evaluate theconcentrations of the impurity particles in the bulk.
Most of the impurity particles had substituted forconstituent particles, as shown in Fig. 1B.
The effective partition coefficient (keff) was
measured at various growth rates (Fig. 1C) and foreach size of each impurity particle. Keff approachedunity with increasing growth rate for each sample.
Further, the closer the size of the impurity particle
was to that of the host particle, the easier it became
for keff to approach unity. In addition, the keff values
for the w/fluor. particles were larger than those forthe pure PS particles for all particle sizes. It should
also be noted that keff for the 520 nm w/fluor. particles
surpassed unity.
Fig. 1D shows k0 obtained from keff by using
BPS plot [3]. k0 decreased with increasing difference
between the size of the impurity particle and the
500-nm host particle. k0 for each w/fluor. particle
size was larger than that for the corresponding purePS particle size. k0 surpassed unity for the 520-nm
w/fluor. particle, whereas it was always less thanunity for the PS particles.
We used a Thurmond and Struthers (T&S)
model [4] to determine the difference between the
k0 values for the PS and w/fluor. particles. The T&S
model defines k0 as exp((ΔGTr–ΔH)/RT), where ΔGTr
is the change in the Gibbs free energy of fusion; thatis, the difference between the Gibbs free energiesof the solid and liquid phases of an impurity at theIMR KINKEN Research Highlights 2013transition temperature, T, ΔH is the excess enthalpy
produced by incorporating the impurity into the hostmaterial, and R is the gas constant. During typicalcrystal growth, ΔGTr is obtained at the transition
temperature of the host material. The volume fraction,
φ, which is the total particle volume divided by the
total system volume, on the other hand, is used as aparameter for the phase transition of the colloidal
dispersion into a colloidal crystal during colloidal
crystal growth. It is necessary to investigate thevolume fraction of the phase transition, φ Tr, for two
different types of impurities in order to evaluate ΔGTrand k0 for both impurities.
The value of φ Tr was determined using thenumber densities and effective diameters of theparticles. The effective diameter is the sum of thereal radius of the particle and the additional portion
owing to the thickness of the electrical double layer,
which was obtained by measuring the zeta potentialof each impurity particle. The number density wasmeasured using optical microscopy by counting thenumber of particles in the liquid phase, which
coexists with the solid phase. Consequently, the φ Tr
value for w/fluor. was less than that for PS for eachparticle size.
Fig. 2A shows a schematic of the molar Gibbs
free energy curves for the PS and w/fluor. particles.
The molar Gibbs free energy is obtained by plotting
the Gibbs free energy as functions of the mole
fractions, x, of the colloidal particles. We used thefree energy curve to represent a binary system eventhough the colloid only had one component. Counterend members are thought to be vacancy (leftdirection) and colloidal particles (right direction),
which are assumed to be at lattice points in both theliquid and solid phases. From the diagram, thechemical potentials of the solid and liquid phases
can be obtained from the lines tangent to thecurves.
We will now discuss why k0 for w/fluor. is larger
than that for PS. The molar free energy curve for the
PS impurity is the same as that for the host because
the materials are the same (Figs. 2A(a) and (b)).
Since the φ Tr for PS and that for the host are thesame, which corresponds to the same mole fraction,
the change in the free energy of the PS, ΔGTr (= μ L– μ S), is zero. This is because the common tangent
lines are the same for the host and the impurity. In
contrast, ΔGTr is positive for w/fluor. (Fig. 2A(c))

Fig. 2 A. Gibbs free energy plotted as functions of mole
fractions of host PS crystals and of PS and w/fluor.
impurity particles. B. k0 obtained from ΔGTr and ΔH for
PS and w/fluor. particles.

because the w/fluor. particle is smaller than the hostPS particle.
Therefore, ΔGTr–ΔH is larger for w/fluor. than it
is for PS, which corresponds to a larger k0 for w/
fluor. than for PS, and k0 for w/fluor. surpasses unity
in a small –ΔH range (Fig. 2B).
The different types of impurity particles had
different k0 values, which is attributed to thedifference in their φ Tr values. We have explained
impurity partitioning during colloidal crystallization
by applying the BPS and T&S models.
References[1] J. Nozawa, S. Uda, Y. Naradate, H. Koizumi, K. Fujiwara, A.
Toyotama, and J. Yamanaka, J. Phys. Chem. B 117, 5289
(2013).
[2] P. Jiang, J. F. Bertone, K. S. Hwang, and V. L. Colvin, Chem.
Mater. 11, 2132 (1999).
[3] J. A. Burton, R. C. Prim, and W. P. Slichter, J. Chem. Phys. 21,
1987 (1953).
[4] C. D. Thurmond and J. D. Struthers, J. Phys. Chem. 57, 831(1953).

Keywords: colloid, crystal growth, doping
Jun Nozawa (Crystal Chemistry Division)
E-mail: nozawa@imr.tohoku.ac.jp
URL: http://www.uda-lab.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201310

Infrastructural Materials

Cobalt- and Titanium-based Alloys for Biomedical andIndustrial Applications

Co-Cr-Mo (CCM) alloys are commonly used for artificial hip- and knee-joint prostheses. We studied thedeformation and corrosion behaviors of CCM alloys for biomedical and industrial applications. In addition,
conventional ultrafine-grained Ti-6Al-4V alloy exhibited low-temperature, high-strain-rate superplasticity.

The mechanical behaviors of Co-Cr-Mo (CCM)
alloys are dominated by strain-induced martensitic
transformation (SIMT), which not only initiates cracks
during plastic deformation but also dramaticallyenhances the wear resistance of CCM alloys inpractical use. It is essential to clarify the factors that
affect SIMT in order to develop more reliable CCM
alloys. In this study [1], we focused on the effect of
annealing twin boundaries (ATBs) on SIMT behavior.
We analyzed in detail the substructures that hadformed near a parallel pair of ATBs after deformation
under a stress preferable for a slip parallel to the
ATBs. Scanning electron microscopy/electronbackscatter diffraction (SEM-EBSD) and transmission
electron microscopy (TEM) were used to find theFig.1 High-resolution transmission electron microscope(HRTEM) images of the twin-containing CCM-alloy deformed
to 8% plastic strain. (a) ε-HCP/Twin-FCC interface. (b)Magnified image of the area indicated by a red rectangular in(a). (c), (f), (i) are HRTEM images of the upper (i.e. near M1),
middle, and lower (i.e. near Twin) regions of the ε-HCP plate.
(d), (g), (j) are the Fourier transformed patterns of (c), (f), (i),respectively. (e), (h), (k) are the magnified images of the area
indicated by dashed rectangular in (d), (g), (h), respectively.

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IMR KINKEN Research Highlights 2013

preferential formation of an ε-HCP (hexagonal closepacked) phase at ATBs in a metastable γ-FCC (facecentered cubic) phase. High-resolution TEM (HRTEM)
images (Fig. 1) indicated that the ε-HCP phase did
not regularly thicken on every second atomic plane,
which would form the perfect HCP-structured ε-phase
but irregularly left a high density of stacking faults.
Furthermore, the thickness of the ε-HCP phase wasdifferent at ATBs on both sides of the twin. The resultsof the residual stress analysis based on the EBSDWilkinson method and the results of the phase-field
simulation of solute segregation at ATBs suggest thatthe difference in the thicknesses was attributed to theinternal stress due to strain incompatibility at the
ATBs. More reliable biomedical CCM alloys can be
developed based on these results.
Co-based alloys can also be used as materials
in Al die-casting. The fundamental behaviors of theinterfacial reactions between solid Co and Fe on the
one hand and liquid Al on the other were comparedby observing the interface structures and by analyzing
the reaction kinetics [2], which laid a foundation forelucidating the mechanisms for the interfacial
reactions between Co-based and Fe-based alloys
on the one hand and liquid Al on the other. Interface
structures typical for solid Co and Fe that had
previously been immersed in liquid Al are shown in
Figs. 2(a) and (b), respectively. According to theresults of the electron probe microanalysis (EPMA)
and the Co-Al binary phase diagram, the intermediate
layers close to the Al phase and to the Co matrix wereidentified as the Co2Al9 and Co2Al5 intermetallic
compounds, respectively. It is assumed that all of theCo2Al9 formed during cooling and that only a single
Co2Al5 layer formed during the interfacial reactions.
Similarly, we believe that only FeAl2.8 inner layers
formed during the interfacial reaction between thesolid Fe and liquid Al. A general model wasconstructed based on this assumption, as shown in
Fig. 2(c), for the interfacial reactions. The solid metal(either Co or Fe) and liquid Al came into contact with
one other when the solid metal was immersed into
the liquid Al, as shown in Fig. 3(c). The two different
metals then reacted at their interface and formed acontinuous layer of an intermetallic compound (Fig.
2(d)). Al atoms that diffused through the intermetallic
layer at the interface between the solid metal and theintermetallic compound (i.e., at interface I) then

IMR KINKEN Research Highlights 2013

Fig 3 Elongation to fracture of new type of UFG-Ti-6Al-4V
alloy (a’-UFG) as compared to the result of Ti-6Al-4V alloy
with grain size of 3μm as a function of initial strain rate.

Fig. 2 Interface structure in cross section after immersion
test at 1013 K for 3600 s.

reacted with the solid metal and formed moreintermetallic compound, causing interface I to movetoward the solid metal side of the interface becausesome of the solid metal was consumed during thereaction. The intermetallic layer, on the other hand,
was unstable in the liquid Al and dissolved into it
through the interface between the intermetallic
compound and the liquid Al (i.e., through interface II),
causing interface II to move toward the solid metalside of the interface (Fig. 2(e)). The decrease in the
thickness of the solid metal is obviously due to the
movement of interface I, and the change in the
thickness of the intermetallic compound is the resultof the relative motions of interfaces I and II.
Further, the Ti-6Al-4V alloy used in this study
exhibited
low-temperature,high-strain-rate
superplasticity [3]. The mechanism for the conversionof the initial α'-martensite microstructure during hot
deformation (at 973 K-10 s-1) of the Ti-6Al-4V alloy
was studied based on detailed microstructuralobservations, kinetic analysis of the deformation of
the microstructure, and various theoretical models.
The initial α'-microstructure was compressed at 973
K-10 s-1 and at a height strain of 0.8. The initial α'microstructure whose morphology was acicular
evolved into an ultrafine-grained microstructure
whose grains were an average of 0.2 µm in diameterand that showed a large fraction of high-angle grainboundaries. The formation of martensite-variant
subgrains and new high-angle-boundary grains
located at the interfaces of the martensite variantsand {10-11} twins dominated during the initial stage of
deformation. Discontinuous dynamic recrystallization(DDRX) and heterogeneous nucleation andfragmentation of the high-angle-boundary grains
dominated when the height strain was increased to
0.8. In contrast, continuous dynamic recrystallization(CDRX) dominated throughout deformation when theinitial microstructure was (α+β). Thus, DDRX becomesdominant when the initial microstructure is changed
from the conventional (α+β) microstructure to theacicular α'- martensite one. This behavior of the α'martensite microstructure is attributed to the numerous
nucleation sites such as dislocations, to the interfaces
between the martensite variants and {10-11} twins,
and to the high-speed grain fragmentation andsubgrain formation in the initial α'- martensite
microstructure during the initial stage of deformation.
Furthermore, the Ti-6Al-4V alloy whose ultrafinegrained microstructure was produced using thismethod of processing exhibited low-temperature,high-strain-rate superplasticity.
References[1] Y. Koizumi, S. Suzuki, K. Yamanaka, B. S. Lee, K. Sato, Y. Li, S.
Kurosu, H. Matsumoto, and A. Chiba, Acta Materialia 61, 16481661 (2013).
[2] T. Tang, Y. Li, S. Kurosu, Y. Koizumi, H. Matsumoto, and A.
Chiba, Corros. Sci. 60, 32-37 (2012).
[3] H. Matsumoto, K. Yoshida, S. H. Lee, Y. Ono, and A. Chiba,
Mater. Lett., 98, 209-212 (2013).

Keywords: mechanical properties, superplasticity, biomedical
Akihiko Chiba (Deformation Processing Division)
E-mail: a.chiba@imr.tohoku.ac.jp
URL: http://www.chibalab.imr.tohoku.ac.jp/english/index.php

IMR KINKEN Research Highlights 201312

Infrastructural Materials

Atomic-Resolution Analysis of Diffusional-Displacive
Transformation of Long-Period Stacking Order phases in
Mg97Zn1Y2 Cast AlloysThe phase transformation of the long-period stacking order (LPSO) structure in aged Mg97Zn1Y2 alloys wasinvestigated using aberration-corrected high-angle angular dark field scanning transmission electron
microscopy (HAADF-STEM). The transformation of occurs through a mechanism resembling diffusionaldisplacive phase transformation. Local strain analysis on the basis of HAADF-STEM images was used toelucidate that the lattice spacing of the (0001)Mg planes in the face-centered cubic (FCC)-type enrichment
layer is shorter than that in the hexagonal close-packed (HCP)-type transition region. These structural andcompositional irregularities are elementary in the transformation of the LPSO structure in Mg97Zn1Y2 alloys.
Magnesium (Mg)-based alloys have attracted
much attention as next-generation lightweight
structural materials because of their low density,
high specific strength, damping capacity, recycling
efficiency, and other features. A series of novel
structures that show so-called “synchronized longperiod stacking order (LPSO)” in Mg-TM-RE alloy
systems (where “TM” and “RE” denote transition
metal and rare earth) has recently been reported [1].
Synchronized LPSO signifies the segregation of TM
and RE elements into stacking faults and the longrange ordering of the segregated stacking faults.
The Mg97Zn1Y2 alloy shows 0.2% proof strength
higher than 600 MPa and an elongation of 5% [1].
Many researchers have previously examined theatomic-level structure of the LPSO. The 10H, 14H,
18R, and 24R polytypes are found in the Mg-Zn-Y
alloy, depending on the thermal history of the alloy
and the concentration of solute in the alloy [2, 3].
Although the formation and growth mechanismof the LPSO have not been elucidated in a long time,
Zhu et al. recently proposed a mechanism for thetransformation from the 18R- to the 14R-type LPSO
structure in Mg96.7Zn0.8Y2.4Zr0.2 alloys [4]. They also
showed that the 24R-type irregularity plays a role in
the transformation. However, they did not consider
modulations in composition and structure at ledges,
nor did they consider the change in the sequence of
the LPSO with the 24R-type irregularity in
Mg96.7Zn0.8Y2.4Zr0.2 alloys aged for different lengthsof time.
We examined the transition state from theperspectives of structural and compositional
irregularities in the LPSO in Mg97Zn1Y2 alloys. The
irregularity of the LPSO structure in aged Mg97Zn1Y2
alloys was examined using aberration-corrected
high-angle annular dark field scanning transmission
electron microscopy (HAADF-STEM) and peak-pairanalysis (PPA) [5].

13

IMR KINKEN Research Highlights 2013

Fig. 1 HAADF-STEM image taken around transition region
between 18R- and 24R-type periodicity in LPSO band.
Stacking sequence in transition region changes from FCC
to HCP-like stacking, which is incomplete HCP stacking.

Modulations in composition and structure in
and around transition region [6]Fig. 1 shows a Bragg-filtered HAADF-STEM
image taken around a transition region in an LPSO
band. Two Burgers circuits clearly show that a pairof Shockley partial dislocations with b = 1/3[0110]
and b = 1/3[0110] exist at both sides of the hexagonalclose-packed (HCP)-like region. These partial
dislocations indicate the existence of a pair of ledge
structures. However, the structure gradually changes
from a face-centered cubic (FCC) structure to an
HCP-type one. The structure in the transition region
seems to be a strained HCP-type structure. The A'
stacking layer moves by the two close-packed(0001)Mg planes across the transition region; thestacking order locally changes from left to right
B'→C, C→A', B→B'. Here, A', B', and C' denotesolute enrichment layers with different stacking
planes in an FCC-type stacking sequence. The Z2contrast of the transition region is lower than that of
the normal enrichment layer of an FCC-type stacking
sequence, meaning that the solute elements caneasily migrate across (0001)Mg planes in thetransition region.

IMR KINKEN Research Highlights 2013

Consequently, the LPSO structure wouldchange through a type of diffusional-displacivetransformation [7].
Local strain around the transition region [6]Fig. 2 shows (a) an HAADF-STEM image and(b–d) show strain mappings around the transition
region of an enrichment layer obtained from the PPA of
the HAADF-STEM image. Here, the strain wasmeasured relative to the lattice spacing of thenonenrichment layer. The in-plane normal strain map
in Fig. 2(b) shows homogeneous contrast with anaccuracy of ±0.3%, which means that the interface
between the enrichment layer and the nonenrichment
layer is actually coherent. The out-of-plane normal
strain map in Fig. 2(c) shows that the lattice spacing of
the enrichment layer is 16% smaller than that of thenonenrichment layer. The strain is limited to within theenrichment layer including a transition region (layer (2))
and the neighboring two nonenrichment layers (layers(1) and (3)). The lattice spacing within the transition
region is about 12% smaller than that in the
nonenrichment layer. The simple shear strain map in
Fig. 2(d) also shows a feature similar to that shown in
Fig. 2(c). The signs of the strain in the enrichment layers
are identical and positive, meaning that the LPSO is
R-type and that the (0110) planes tilt to the right.
These strain maps indicate that the strain in thetransition region gradually decreased withdecreasing Z2-contrast; i.e., with decreasing
concentration of the solute elements. Figs 2(c) and(d) also show that the enrichment layers are
separated in the transition layer. On the basis ofthese results, a schematic model such as that shown

Fig. 3 Schematic model for transformation of LPSO
through transition region.

in Fig. 3 can be proposed for the region around thetransition layer. The stacking sequences are HCPand FCC-type, inside and outside the transition
region, respectively. Shockley partials exist in ledges
between those stacking sequences. The HCP-like
transition regions in front of two ledges fuse together
and form a kind of diffusion field through which thesolute elements easily migrate without changing thestacking sequence in the out-of plane direction of
the LPSO; the diffusion field in front of the ledge
serves as a buffer for the short-range migration ofsolute elements. The transition region subsequently
moves in a [0110]Mg or [0110]Mg direction on the(0001)Mg planes accompanying the movement of apair of partial dislocations. The strain accompaniedwith the transformation is localized around atransition region. Consequently, the enrichment
layer grows or shrinks in the [0110]Mg or [0110]Mg
direction, which is the elementary step in the phasetransformation of LPSO.
References[1] Y. Kawamura, K. Hayashi, A. Inoue, and T. Masumoto: Mater.
Trans. 42, 1172 (2001).
[2] E. Abe, Y. Kawamura, K. Hayashi, and A. Inoue: Acta Mater.
50, 3845 (2002).
[3] E. Abe, A. Ono, T. Itoi, M. Yamasaki, and Y. Kawamura: Philos.
Mag. Lett. 91, 690 (2011).
[4] Y. M. Zhu, A. J. Morton, and J. F. Nie: Acta Mater. 60, 6562
(2012).
[5] E. Abe, A. Ono, T. Itoi, M. Yamasaki, and Y. Kawamura: Philos.
Mag. Lett. 91, 690 (2011).

Fig. 2 (a) HAADF-STEM image and (b-d) strain mapping
around transition region of enrichment layer. (b) In-plane
normal strain, (c) out-of-plane normal strain, and (d)simple shear strain.

[6] T. Kiguchi, Y. Ninomiya, K. Shimmi, K. Sato, and T. J. Konno:
Mater. Trans. 54, 668 (2013).
[7] J. W. Christian: Metallur. Mater. Trans. A 25, 1821 (1994).

Keywords: alloy, phase transformation, scanning transmission electron microscopy (stem)
Takanori Kiguchi, Kazuhisa Sato, and Toyohiko J. Konno (Advanced Materials Analysis Division)
E-mail: tkiguchi@imr.tohoku.ac.jp
URL: http://konno-lab.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201314

Infrastructural Materials

First-Principles Study of Interface Structure and Energy ofIron/Niobium CarbideT

he Computational Materials Research Initiative (CMRI) was established as a project-oriented section at theInstitute for Materials Research (IMR) in 2011 and is currently one of the three institutions operating under theComputational Materials Science Initiative (CMSI) founded through a Grant for Field 2 “New Materials andEnergy Creation” of the High-Performance Computing Infrastructure (HPCI) Scientific Program of the Ministryof Education, Culture, Sports, Science, and Technology. As of May 2013, thirty-one researchers and supporting
staff from universities, government organizations, and companies all over Japan are participating in CMRI
along with the other two CMSI operating institutions, the Institute for Solid-State Physics (at The University ofTokyo) and the Institute for Molecular Science (at the National Institutes of Natural Sciences). CMRI conducts
various programs of computational materials science. Among the seven Priority Research Topics of CMSI for
which machine time is prioritized for using the K computer to accelerate the progress of calculations, CMRI has
been pursuing research entitled “Development of multiscale structural design and assessment techniques toimprove the performance of metallic structural materials.” The task group committed to this research is headed
by Dr. M. Kohyama, AIST, and several CMRI members are also actively involved. Introduced below is a summaryof one of the major outcomes obtained by Dr. H. Sawada, one of the core members of CMRI, and his coworkers[1]. They performed the so-called O(N) Krylov-subspace method of calculating electronic structures from firstprinciples to clarify the key factors for controlling the growth of precipitates in steel. A classical molecular
dynamics simulation was also performed to estimate the amount of strain energy around the precipitates.

The size and density of precipitates areimportant factors in controlling the strength andtoughness of steel. It is necessary to know thenumber of nucleation sites, the diffusivity of soluteatoms, and the amount of interface energy between
iron and the precipitates in order to predict the sizeand density of the precipitates. An attempt wasmade for the first time at using the O(N) Krylovsubspace method of calculating electronic structuresfrom first principles [2] instead of using theconventional diagonalization method to directly
determine the interface energies of the semicoherent
interfaces between body-centered cubic (bcc) ironand niobium carbide (NbC) arranged in a latticecontaining more than 1000 atoms. An attempt wasalso made at using classical molecular dynamics(MD) simulation to calculate the strain energy in

Fig. 1 Interface energies calculated for coherent andsemicoherent interfaces between bcc iron and NbC
and plotted as functions of length of c 0.

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IMR KINKEN Research Highlights 2013

order to estimate the size of the precipitate at thetransition from a coherent precipitate to asemicoherent one.
Two examples of a coherent interface were
calculated. In one case, a C atom existed on top of
an Fe atom at the interface and in the other, an Nb
atom existed on top of an Fe atom. These calculations
were performed for a unit cell consisting of 7 Fe and7 NbC layers, and the interface energy was definedas the difference between the total energy of thesystem including the interface and the sum of all theenergies of the individual phases. As shown in Fig.
1, the interface formed when the C atom is on top of
the Fe atom is more stable than the interface formed
when the Nb atom is on top of the Fe atom. In
addition, the interface separation in the former caseis shorter than that in the latter one.
The semicoherent interface consists of m unitsof bcc iron and n units of NbC along the [100] and[010] directions. The lattice constant of NbC,
a0(NbC), is 1⁄√2 of the conventional lattice constantbecause of the Baker-Nutting orientation relation.
The misfit, ε, between the bcc iron and NbC phases
for a semicoherent interface,
ε = |m × a0(Fe) –n × a0(NbC)| (1)was minimized to 0.03 Å by choosing 9 and 8 for m
and n, respectively. Here, a0(Fe) and a0(NbC) are
the lattice constants of bcc iron and NbC. Lengths
a0 and b0 of the unit cell for the semicoherent
interface were fixed at n × a0(NbC). The calculation
was performed for a 14-layer unit cell consisting of 7
Fe and 7 NbC layers along the [001] direction. Theinterface energy calculated for the semicoherent

IMR KINKEN Research Highlights 2013

interfaces is between those for the two coherent
interfaces, as shown in Fig. 1. The separation of thesemicoherent interface, however, is close to that of
the coherent interface when the Nb atom exists on
top of the Fe atom.
Fig. 2 shows the optimized atomic positions of
the semicoherent interface in the (110) plane. The
Fe atoms near the interface approach the C atoms,
and the lines of the Fe atoms are bent, while the Nb
atoms near the center of the interface move toward
the [001] direction to maintain their distance fromthe Fe atoms. These features are consistent with thefact that the interface where the C atom exists on topof the Fe atom is preferable to that where the Nb
atom exists on top of the Fe atom, as shown in Fig.
1. The lines through the Fe atoms, except for the twolines located near the center of the horizontal axis,
are connected to C atoms at the interface. There isa region where only one C atom exists for the twolines through the Fe atoms near the center of thehorizontal axis. In Fig. 2, this yellow region is adislocation core.
The MDs were then attempted (performed) toestimate the amount of strain energy around theprecipitate. The calculation cell was divided into two
regions—the precipitate and matrix regions—although both regions were treated using theembedded-atom method (EAM) potential of Fe. The
atomic positions were fixed so as to have the samelattice parameter as NbC in the hypotheticalprecipitate region. The strain energy was estimatedby comparing the potential energies of the unit cells
containing distorted or undistorted precipitate
regions. Fig 3 shows the sum of the interface andstrain energies plotted as functions of the length of
the side of the precipitate for both the coherent andsemicoherent precipitates and indicates that thegrowth of the precipitate leads to a transition fromthe coherent to the semicoherent state at 18- and16-Å-long sides for precipitates that are 5.36 and8.94 Å thick, respectively.
In summary, the interface and the strain energiesof the interface between bcc iron and NbC were
calculated using the O(N) method of calculating
electronic structures from first principles and usingclassical MD simulation. The lengths of the sides ofthe broad plane of the plate-like rectangularparallelepiped-shaped precipitates were also
estimated at the transition from the coherent to thesemicoherent interface.

Fig. 2 Atomic configuration of semicoherent interface in(110) plane. For easy visualization, Fe atoms areconnected by lines along [001] direction. Yellow region
corresponds to dislocation core.

Fig. 3 Sum of interface and strain energies plotted asfunctions of the side length of the precipitate for coherent
and semicoherent precipitates. Precipitates are (a) 5.36
and (b) 8.94 Å thick.

References[1] H. Sawada, S. Taniguchi, K. Kawakami, and T. Ozaki, Modelling
Simul. Mater. Sci. Eng. 21, 045012 (2013).
[2] T. Ozaki, Phys. Rev. B 74, 245101 (2006).

Keywords: first-principles calculations, high-strength steel, defects
Tetsuo Mohri (Computational Materials Research Initiative)
E-mail: tmohri@imr.tohoku.ac.jp
URL: http://www.cmri-tohoku.jp/

IMR KINKEN Research Highlights 201316

Energy-related Materials

Energy Harvesting and Transport through Spin in Insulators

Recent progress in spin injection at metal-insulator interfaces [1] enables us to include spin degrees of freedom

in insulators used in electronic [1] and energy [2] devices. Here, we demonstrate a novel concept for thermoelectric
conversion [3] based on such advances. The insulators are characterized by a simple film structure, convenient
scalability, and easy fabrication. We also present a unique unidirectional energy transport based on the chiralityof spin excitation [4], which can be used as a simple rectifier in energy-harvesting technologies.

Since heat energy is ubiquitous in a diverse
range of forms, thermoelectric (TE) technology has
tremendous opportunities for application to various
energy-harvesting and heat-sensing devices. The
most prevalent TE technology is based on theconventional Seebeck effect (Fig. 1a). Whentemperature difference ΔT is given across athermocouple (TC), it generates TE voltage, VTC, thatis fixed by material-dependent parameters (i.e.,
Seebeck coefficients). Although the output voltage
can be scaled by connecting n pairs of TCs (V =nVTC) in a TE module, fabricating such a cascaded
structure is costly.
We demonstrate another category of TE
technology based on the spin Seebeck effect (SSE)[2], where a spin current is driven by a temperature
gradient. The spin-thermoelectric (STE) coating
developed in the present study [3] consists of metallic
and magnetic insulator films directly coated onto aheat source, as shown in Fig. 1b. The STE coatingoffers the following interesting features. It isconveniently scalable because the TE output can bereadily increased simply by extending the coating
area, since the total amount of the SSE-induced spincurrent increases with increasing the film surface
area. Such straightforward scaling enables us to use
coating and printing technologies in order toimplement a simple-structure large-area TE function
into variously shaped heat sources such as electric
instruments and automobile bodies. Moreover, thethermal energy inside the heat sources can also
nonlocally contribute to the thin-film-based TE
conversion through the interplay between spin current
and phonons. These unique features may offer
ubiquitously applicable TE technology simply byappending STE coating onto various objects. The
STE coating presented here is fabricated using
magnetic insulator films fabricated using spin-coating
on nonmagnetic surfaces, which potentially enable
the production of large-area TE systems.
The STE coating was demonstrated not only on
expensivehigh-performance
single-crystalline
substrates but also on a 0.5-mm-thick synthetic fusedsilica glass substrate to check the versatility of the
STE coating in practical applications. The transmission
electron microscopy (TEM) images suggest that theSTE coating does not require specific template
17

IMR KINKEN Research Highlights 2013

Fig. 1 Concept of STE coating [3] compared with that ofconventional TE devices. Scaling of SSE voltage
enables output voltage to be accumulated without
fabricating complex structures for series connections,which are necessary for conventional TE.

surfaces for crystal growth probably because theyttrium-iron-garnet lattice structure, which isthermodynamically quite stable, can be coherently
formed over a large grain region during annealing. A
clear SSE-induced TE voltage was measured for thecoated glass device. Applying the STE coating toamorphous surfaces has a significant impact not only
from a practical perspective producing clear SSEinduced TE voltage but also in the field of materials
science and technology, where single-crystalline
films have conventionally been epitaxially grown only
on lattice-matched substrates.
In addition to the STE-coated glass, we also
coated thermally oxidized silicon substrates (SiO2/Si)
with STE, suggesting that it can be applied to various
kinds of surfaces by inserting oxide buffer layers if
necessary. This versatility will potentially create
opportunities for developing a broad range of TE
applications.
Energy transport has recently become anotherimportant technological issue because it always
generates heat, which favors randomness anddiffuses everywhere. Here, we show that nonreciprocal
spin waves; i.e., collective excitations of coupled spin
systems in magnetic lattices in the microwave
frequency range can convey energy and emit heatover long distances, becoming a “spin-wave heat

IMR KINKEN Research Highlights 2013

conveyer” that forms a negative temperature gradientdifferent from that formed in conventional microwave
heating [4].
A schematic of the experimental setup of spinwave excitation in a polycrystalline Y3Fe5O12 slab andof using an infrared (IR) camera to image the surface
temperature of the slab is shown in Fig. 2a. Different
types of spin-wave modes were excited when the
microwave frequency was swept from 1.8 to 5.2 GHz(Fig. 2b). A nonreciprocal surface spin-wave mode
(i.e., mode D in DEM in Fig. 2b) was chosen, and theincrease in temperature imaged using the IR camera
showed heat emission at the end of the sample in astatic magnetic field of +H = 630 Oe (Fig. 2c). The
direction of heat emission was reversed by changing
the polarity of the applied magnetic field –H = 630
Oe. This is evidence of the “heat conveyer effect”
due to unidirectional energy transfer through theexcitation of nonreciprocal spin-waves without time
reversal symmetry and to the conversion of spinwaves into heat at the end of the sample.
The mechanism for the heat conveyer effect may
be explained by the suppression of the reflection of
the surface spin waves (DEMs) from the bottom
surface to top one. In a sample of a certain thickness
where the amount of interaction among the surface
spin-wave modes between the bottom and top
surfaces is small, the surface spin-wave modes
coming to the end at the edge of the sample cannot
reflect back from one to the other. Furthermore, we
demonstrated that even for the thin-film of singlecrystalline Y3Fe5O12, where the surface spin-wave
interaction of the two surfaces is strong, the effective
surface spin-wave reflection can be suppressedusing a sharp 30-degree angle for the spin-wave
waveguide to introduce spin-wave multireflections.
This finding enables us to use magnetostatic spinwaves in both single- crystalline and polycrystalline
Y3Fe5O12 samples of certain thicknesses. The two
major channels of heat-conduction loss inpolycrystalline-insulator
magnetsusedinthermodevices; namely, conduction electrons andphonons, are suppressed while maintaining highly
efficient manipulation of heat because themagnetostatic spin waves also reside in an insulator
magnet in which conduction electrons are absent
and can readily pass through the grain boundaries ofpolycrystals at which phonons, in contrast, arestrongly damped. In single-crystalline insulator

Fig. 2 Spin-wave heat conveyer. (a) Schematic of setup for
spin-wave excitation, (b) obtained absorption spectrum,and (c) surface spin wave (mode D) heat conveyer effectobserved using infrared thermocamera. After ref. [4].

magnets, in contrast, the excited magnetostatic
surface spin-wave can effectively travel toward theend of the sample even at a distance of 10 mmbecause the damping coefficients for the spin-wave
are low.
The microwave-induced heat conveyer effectobserved here can be directly applied to theconstruction of a microwave heat-transport systemsuch as a heat-flow controller for manipulating heatover several-millimeter-long solids, thereby opening
the door to heat technology based on spin currents,
which is closely related to the emerging field of “spincaloritronics,” which is in turn concerned with newphenomena related to the coupling of spin, charge,
and energy transport in materials.
References[1] Y. Kajiwara, K. Harii, S. Takahashi, J. Ohe, K. Uchida, M.
Mizuguchi, H. Umezawa, H. Kawai, K. Ando, K. Takanashi, S.
Maekawa, and E. Saitoh, Nature 464, 262 (2010).
[2] K. Uchida, J. Xiao, H. Adachi, J. Ohe, S. Takahashi, J. Ieda, T.
Ota, Y. Kajiwara, H. Umezawa, H. Kawai, G. E. W. Bauer, S.
Maekawa, and E. Saitoh, Nature Mater. 9, 894 (2010).
[3] A. Kirihara, K. Uchida, Y. Kajiwara, M. Ishida, Y. Nakamura, T.
Manako, E. Saitoh, and S. Yorozu, Nature Mater. 11, 686 (2012).
[4] T. An, V. I. Vasyuchka, K. Uchida, A. V. Chumak, K. Yamaguchi,
K. Harii, J. Ohe, M. B. Jungfleisch, Y. Kajiwara, H. Adachi, B.
Hillebrands, S. Maekawa, and E. Saitoh, Nature Mater. 12, 549(2013).

Keywords: spintronics, energetic material, spin wave
Eiji Saitoh (Surface and Interface Research Division)
E-mail: saitoheiji@imr.tohoku.ac.jp
URL: http://saitoh.imr.tohoku.ac.jp/index.html

IMR KINKEN Research Highlights 201318

Energy-related Materials

Novel Material Science using Neutron ScatteringI

MR is a unique institute even in the world. The institute has a neutron facility based on the strong researchbackground in advanced material science. The IMR neutron scattering group is investigating various problems
in materials sciences, such as ferroelectric materials, high-TC superconductivity, metallic magnets, andapplication materials, as well as developing high -performance devices for neutron science. In particular, ourgroup is constructing a new type of neutron spectrometer with in Japan Proton Accelerator Research Complex(J-PARC), which will be a flagship instrument for research in material sciences in IMR.
Structural and dynamic properties of lead-freeferroelectric material
Lead-based perovskites such as Pb(Zr,Ti)O3(PZT) are widely used in everyday life as ferroelectricand piezoelectric devices. Recently, lead-freeferroelectric or piezoelectric compounds have beenstudied extensively as replacements to the leadbased ones owing to the toxicity of lead. Bismuth
sodium titanate (Bi1/2Na1/2)TiO3 (BNT) and BNTbased compounds have gained considerable
attention as lead-free alternatives owing to theirrelatively large spontaneous polarization and highCurie temperature (TC) [1]. BNT exhibits successive
structural transitions, the first from cubic to tetragonalphase at 813 K, and the second from tetragonal torhombohedral phase at ~700 K. The rhombohedralphase is ferroelectric with parallel cation shifts along
the polar <111> axes, whereas the tetragonal phaseis ferroerectric with antiparallel atomic shifts between
Bi/Na and Ti cations along the fourfold [001] axis.
Thus, the ferroelectric properties in the pure
tetragonal phase and the pure rhombohedral phasehave been clarified. However, the intermediatephase in which the rhombohedral and tetragonalphases coexist is not fully understood. For example,
the dielectric constant shows a broad maximum at
600 K, but no phase transition is observed at thattemperature. Furthermore, a dramatic degradationof the piezoelectric properties in the rhombohedralphase has been observed at 460–480 K. To
elucidate the microscopic mechanism of the broad
maximum in the dielectric constant and thedepolarization in BNT, we performed inelastic andelastic neutron scattering on single crystals of BNT.
Fig. 1 shows the thermal variations in the
dielectric constant, integrated intensity of diffuse
scattering near the Γ point, and Bragg peaks at the
M point (1/2,3/2,0) and the R point (3/2,1/2,1/2).[2]The Bragg peak at the M point becomes weak below
600 K, at which point the dielectric constant reaches
its maximum value, whereas the diffuse peak at theΓ point and the Bragg peak at the R point growbelow 600 K. Furthermore, the depolarization
temperature coincides with the disappearance ofthe M-point Bragg peak. Between the depolarization
temperature (460–480 K) and 600 K, the19

IMR KINKEN Research Highlights 2013

Fig. 1 Temperature dependence of dielectric constantε’(T) measured at 1 MHz, intensity of diffuse scattering
near Γ point (Γ diffuse), Bragg peaks at M point(1/2,3/2,0) (M-Bragg) and R point (3/2,1/2, 1/2) (R-Bragg[3]). ε’(T) shows broad maximums at 600 K (Tm).
Depolarization occurs at 460–480 K (Tdp) [2].

rhombohedral and tetragonal phases coexist. The
diffuse scattering near the Γ point indicates shortranged ferroelectric fluctuations. Because a large
hysteresis is seen in the dielectric constant in thisphase-coexistence region, we propose that that thebroad maximums in the dielectric constant anddepolarization process originate in a diffusive firstorder transition between the competing tetragonal
and rhombohedral phases.
Anomalous Spin Response in Metallic
Antiferromagnets
The discovery of iron-pnictide superconductors
reactivated studies on metallic magnets, especially,on spin-density waves (SDW) arising from instabilityof Fermi surfaces nested in multiband systems.
Mn3Si is an antiferromagnetic (AF) metal that showsincommensurate spin correlations below TN, due
probably to SDW. The characteristic magnetic
excitations, or so-called chimney structures, appears
at ω > 20 meV and around Γ points such as (2, 0, 0)
and (2, ±2, ±2,). To explore the origin of the chimney
structure, we investigated the spin dynamics of
Mn2.4Fe0.6Si, a Mn-site substituted compound; here,

IMR KINKEN Research Highlights 2013

Fe substitution is regarded to introduce ferromagnetic
interactions in antiferromagnetic Mn3Si.
Fig. 2 directly compares the magnetic scattering
intensities of the chimney structures of Fe-substituted
and pristine Mn3Si. It is clear that the cross sections
at three Γ points decrease upon Fe substitution.
Thereby, the chimney structure most likely originatesfrom the antiferromagnetic spin correlations. Basedon the different symmetry in Q space from that of thechemical unit cell in Mn3-xFexSi, the dynamic
magnetic periodicity must be changed. The can beachieved by either a switch of the nesting vector (ora switch of the electron band at the Fermi energy) or
reconstruction of the magnetic unit cell at higher ω.
We conjecture that this type of crossover indynamical spin correlation occurs in a wide variety
of itinerant-electron systems.
New Project in J-PARC
Technological advancement in neutron scienceis an important target of our group. In particular, wehave succeeded in developing a polarized neutron
diffraction system in a reactor facility,[4] and haveshown its sensitivity to magnetic moments inferromagnetic materials. Our group and the Centerof Neutron Science for Advanced Materials, IMR,
under a collaboration project with High Energy
Accelerator Research Organization (KEK), areconstructing a new polarized analysis neutron
spectrometer in J-PARC/MLF, which will, uponcompletion, be the brightest neutron facility in theworld. Given that there are very few spectrometersthat are polarization-analysis-capable, which isindispensable to investigations on pure spincorrelations and spin–lattice couplings, in the world,
this project will be a key instrument for IMR in thenear future. The details are reported in an article by
the Center of Neutron Science for Advanced
Materials in this booklet.

Fig. 2 Magnetic excitation at {2,0,0} decreases by Fe
doping. The right inset shows the scan trajectory.

Fig. 3 Main part of polarization analysis neutron
spectrometer in J-PARC. This instrument will be
important for spin correlation investigations.

References[1] T. Takenaka et al., Ferroelectrics, 106, 375 (1990).
[2] M. Matsuura et al., PRB, 87, 064109 (2013).
[3] S. B. Vakhrushev et al., JETP Lett., 35, 134 (1982).
[4] K. Ohoyama et al., Nucl. Instr. Meth. Phys. Res. A, 680, 75(2012).

Keywords: neutron scattering, high-Tc cuprate, spin fluctuation
Kenji Ohoyama, Haruhiro Hiraka, Masato Matsuura, and Toyoto Sato (Metal Physics with Quantum Beam Spectroscopy
Division)
E-mail: ohoyama@imr.tohoku.ac.jp
URL: http://www.yamada-lab.imr.tohoku.ac.jp/jp/index.html

IMR KINKEN Research Highlights 201320

Energy-related Materials

Grain Boundary Engineering in Silicon for solar cells: Use of
Functional Grain Boundaries to Suppress Multicrystallizationof Mono-like Silicon

We used functional grain boundaries artificially composed of multiseed crystals to develop a method of

suppressing the propagation of grain boundaries in Si crystals. The method is based on a fundamental studyof the configuration of grain boundaries and their reactions. We suppressed multicrystallization in a practicalscale mono-like Si ingot. The method can be used to significantly increase the yield of mono-like wafers
produced from an Si ingot and significant reduce the energy cost of producing solar cells.

Although approximately 50% of the substrates
used to produce solar cells are multicrystalline-Sibased, mono-like Si is a promising material forfurther increasing the conversion efficiency of solar
cells. Mono-like Si is a quasi-single crystalline Sigrown using single crystalline seeds in a crucible.
The same facilities and processes used for the
standard production of multicrystalline-Si-basedsolar cells can be used with minimal adjustment toproduce the mono-like Si. Therefore, research anddevelopment of mono-like Si has recently accelerated
worldwide. However, a major disadvantage of monolike Si is “multicrystallization”; that is, an increase in
the occupying portion of multicrystalline grains
nucleated at crucible side walls whose orientations
are different from that of the seed during growthfrom Si melt. This multicrystallization results in asignificant decrease in the yield of quasi-single
crystalline wafers produced from an ingot.
To solve this problem, we used the functions ofgrain boundaries artificially induced by the specificcomposition of seed crystals to develop a methodof suppressing such enlargement of the grains. Fig
1 (a) shows a schematic of the conventional methodof preparing mono-like Si (left) and our method(right). In the conventional method, some crystal
grains are nucleated at random sites on the innersurface of the crucible during growth. The area of
the generated grains increases with increasing
distance from the seed. The multicrystalline area
along the crucible walls then surrounds mono-likearea. In our method, on the other hand, grainboundaries are formed near the crucible side walls
parallel to the walls from the composed multiseed
crystal. These grain boundaries suppressed theincrease in the area of the multicrystalline grains,which resulted in a large volume of quasi-single
crystalline Si formed to the top of the ingot.
This idea originated from a fundamental studyof the reaction and crystallographic configuration ofgrain boundaries. We first investigated the grainboundaries generated at crucible side walls. In otherwords, those formed by the nucleation of themulticrystalline grains. Most of the grain boundaries
nucleated at the crucible surface were Σ3 grainboundaries. Σ represents the size of the unit cell of
the coincident site lattice (CSL) relative to that of thecrystal lattice. Similar nucleation of the Σ3 grain

Fig. 1 (a) Schematics of conventional method (left) and our method (right) of preparing mono-like Si. (b) Close-up illustrationof region near crucible side wall. Increase in area of multicrystalline grains is functionally suppressed by reaction with Σ5
grain boundaries artificially formed by the multiseed.

21

IMR KINKEN Research Highlights 2013

IMR KINKEN Research Highlights 2013

boundaries at the crucible inner surface was also
observed during the growth of multicrystalline Si [1],which should lead to the very low interfacial energyof the Σ3 grain boundaries. The Σ3 grain boundaries
in Si crystals have the lowest energy at the {111}plane, which causes almost all the Σ3 grainboundaries to have {111} planes. {111} planes incline
in the <100> direction at a 35.7° angle. Therefore,
for crystals that grow in the <100> direction, thelateral position of the Σ3 grain boundaries moves
toward the center of the crystals; i.e., the area ofmulticrystalline grains increases with increasing
distance from the seed (Fig. 1 (b)).
We chose the Σ5 grain boundary, whose
structure and electrical properties in Si crystal weinvestigated and clarified for the first time ever, as
the functional grain boundary to suppress
multicrystallization. The Σ5 grain boundary shows
low electrical property, which is an indispensableproperty for solar cell applications, and it stably
extends in the <100> growth direction. We used
multiseed crystals to artificially form Σ5 grainboundaries near the crucible side walls. As the
crystal grew, the Σ3 grain boundaries formed at the
crucible side walls intersecting with the Σ5 grainboundaries. Σ15 grain boundaries were formed
through the interaction between the Σ3 and Σ5 grainboundaries. The Σ15 grain boundaries tend toextend in the direction of crystal growth to decrease
their energy, as do the Σ5 grain boundaries, which
suppresses the increase in the area of themulticrystalline grains (Fig. 1 (b)) [2].
To demonstrate this suppression mechanism,
we grew various size ingots: columnar 3 cm indiameter, rectangular 10 cm on a side, and rectangular
40 cm on a side. The functional grain boundaries
suppressed multicrystallization in all the ingots.
Fig. 2 shows a photograph of a vertical crosssection of a 20-cm-high 70-kg ingot whose interiorside was 40 cm long. We intentionally formed two
different regions in this ingot: one with and one
without functional grain boundaries near the crucibleside wall. Fig. 2 clearly shows that while
multicrystalline grains (i.e., the Σ3 grain boundaries)
extended into the inner ingot from the side formed
without the functional grain boundaries, thefunctional grain boundaries suppressed theextension of the Σ3 grain boundaries into the inneringot from the crucible side walls formed with theFig. 2 Photograph of vertical cross-section of 20-cm-high,
70-kg ingot whose interior side is 40 cm long. Ingot wasformed without (left) and with (right) functional grainboundaries.

functional grain boundaries.
The peripheral regions of ingots near the
crucible side walls generally cannot be used to
produce solar cell wafers because of highconcentrations of impurities from the crucible.
Therefore, a 100% yield of quasi-single crystalline
wafers can be expected if the functional grainboundaries confine the multicrystalline grains in an
ingot to the area unused for solar cells. It should be
noted that functional grain boundaries can be used
to industrially produce ingots because themechanism for changing the grain boundary
characteristics is independent of the size and growth
conditions of the ingots.
In summary, we developed a method of growing
mono-like Si in which functional grain boundaries
are used to effectively suppress multicrystallization .
This method has the potential of significantly
reducing the cost of producing solar cells.
References[1] K. Kutsukake, T. Abe, N. Usami, K. Fujiwara, K. Morishita, and
K. Nakajima, Scripta Mater. 65, 556 (2011).
[2] K. Kutsukake, N. Usami, Y. Ohono, Y. Tokumoto, and I.
Yonenaga, Appl. Phys. Express 6, 025505 (2013).

Keywords: grain boundary, crystal growth, solar cell
Kentaro Kutsukake (Physics of Crystal Defects Division)
E-mail: kutukake@imr.tohoku.ac.jp
URL: http://lab-defects.imr.tohoku.ac.jp/index-e.html

IMR KINKEN Research Highlights 201322

Energy-related Materials

Calculation of Van der Waals Dispersion Coefficients with
All-Electron Mixed-Basis MethodT

he Casimir-Polder integral was used to obtain the van der Waals dispersion coefficients in order to evaluate
the parameters of weak interactions. The time-dependent local density approximation (i.e., Casida formalism)
was entirely implemented in real space within the all-electron mixed-basis approach to improve the descriptionof the oscillator strengths and true excited states. The true excitation energies of many-electron systems werecomputed from the conventional, time-independent Kohn-Sham transition energies and wave functions.
Electronic wave functions were directly expressed in real space on the basis of a regular cubic grid.
The fundamental role of van der Waals (vdW)
interactions at distances longer than interatomic
bonds between neutral atoms has recently beenrecognized in materials science. Atoms repel atshort distances where the overlap of the electronic
clouds is nonnegligible and attract through mutualpolarization of their electronic clouds at distanceswhere the overlap is negligible. There is afundamental difference between these repulsive
and attractive interactions. The rules of quantum
mechanics describe the repulsive interactions while
those of quantum electrodynamics describe theattractive or vdW ones. Within the framework of thegeneral theory of vdW forces, neutral atoms in a
vacuum and in fluctuating electromagnetic fields
absorb and emit photons. Atoms that absorb
photons are transferred into excited states, which
can lead to the appearance of virtual multipole
moments whose interactions determine themagnitude of the attractive vdW interaction between
the atoms. Therefore, long-range interactions appear
because of interactions between the electrons of
the atoms in the excited states, while short-range
interactions are due to the repulsion interaction
between the electrons of the atoms in their groundstates. The fundamental role of vdW interactions at
distances longer than interatomic bonds wasrecently recognized in nanoscience.
We implemented a linear-response theory
withinthetime-dependent
local
density
approximation (TDLDA) to study the true electron
excitations of atoms (i.e., clusters) [1]. The frequencydependent polarizabilities and vdW dispersion
coefficients for atoms (clusters) were calculatedusing the Casida method, in which TDLDA excitationenergies ΩI and oscillator strengths ƒI can be derived
from single-electron Kohn-Sham (KS) eigenvalues→
εk and eigenwave functions Ψκ( r ' ). The true
excitation energies, ΩI , which correspond to thepoles of the dynamic polarizability, can be obtainedby solving the Casida equation:
where matrices Q are given as:

23

IMR KINKEN Research Highlights 2013

(1)


(2)
where λlk = nτ − nk is the difference between theoccupation numbers, ωlk = εk − εl is the difference
between the eigenvalues of the lth and kth singleparticle states, and Kij, kl is the adiabatic coupling
matrix. The analytic expression for the adiabatic
coupling matrix is written in the form(3)

The oscillator strengths, ƒI , are given as
(4)
where FI are the eigenvectors of the Casida Eq. (1)shown as
(5)and where {r1, r2, r3} = {x, y, z}.
The single-electron KS eigenvalues, εk, and→eigenwave functions Ψκ( r ' ) obtained within the allelectron mixed-basis approach (TOMBO) were used
to find the elements of the coupling matrix [2]. The
characteristic feature of the all-electron mixed-basis
approach is that the Kohn-Sham wave functions are
expanded in a mixed basis consisting of the planewaves (PWs) and well-localized functions centered→at the atomic nuclei positions, R=0, the so-calledatomic orbitals (AO) (Fig. 1).
The Kohn-Sham wave functions for atoms are
expressed as follows:(6)
It is necessary to evaluate the eigenvectors, FI ,of the Casida Eq. (1) and the tensor of dipole moment(Bβ)ij in order to use Eq. (4) to find the oscillator
strengths, ƒI .
The exact expression (i.e., the Casimir-Polder
integral) for the leading isotropic dispersion
coefficient (C6 (A, B)) term describing the vdW
interaction between two atoms or molecules A andIMR KINKEN Research Highlights 2013

B is written as
(7)where αA/B (iω) are the frequency dependent
polarizabilities of A and B evaluated at imaginary
frequencies. The average dynamic polarizability ofatoms or molecules can be expressed as:(8)where ΩI is the true excitation energy, and ƒI isthe oscillator strength determined from the Casida
equations; i.e., Eqs. (1) and (4).
The dispersion coefficients, C6, for hydrogenatoms and for hydrogen, oxygen, and watermolecules are listed in Table 1. The C6 coefficientfor hydrogen molecules was 29.25 atomic units at a
distance where the orientation of the molecules canbe neglected. This value is consistent with thosepreviously calculated based on long-range density
functional theory.
TABLE 1. Dispersion coefficient, C6 (A-B), (measured in
Rydberg atomic units). SPCE, TIP4, ST2, and BNS are
semiempirical potentials commonly used for describing
interaction between water molecules.

C6 (H-H) C6 (H2-H2) C6 (O2-O2) C6 (H2O-H2O)
14.6 (a)29.25 (a)94.7 (a)83.25 (a)12.0 (b)29.0 (b)71.6 (c)89.8 (SPCE)
13.0 (d)34.3 (e)97.48 (f)88.5 (TIP4P)
23.0 (f)80.5 (g)38.96 (ST2)
22.28 (BNS)(a) Our study [1]; (b) Y. Anderson et al., Phys. Rev. Lett.
76 102 (1996); (c) J. C. Platteeuw and J. H. van der
Waals, Mol. Phys. 1 91 (1958); (d) D. M. Bishop and J.
Pipin, Int. J. Quant. Chem. 45 349 (1993); (e) K. Murata
et al., J. Phys. Chem. B 106 11132 (2002); (f) U. Hohm,
Chem. Phys. 179 533 (1994); (g) B. Bussery and P. E. S.
Wormer, J. Chem. Phys. 99 1230 (1993).

Fig.1 Schematic of one-electron wave function expressed
in TOMBO program.

obtained using our method was also higher thanthat obtained using the Hartree-Fock molecularorbital calculations (80.5 atomic units), but wasclose to the value obtained using quantum
mechanical simulation (97.48 atomic units). We also
directly calculated the C6 coefficients for O-O andH-O dimers as 47.3 and 30.4 atomic units,
respectively. On the basis of the calculated O-H and
H-H dispersion coefficients, the dispersion
coefficient of the vdW interaction between watermolecules at long distances can be calculated using
the standard summation rule. Table 1 compares theC6 for water with those for various atom-atomic
interaction potentials. The results are consistent withthe dispersion coefficients of the vdW interaction
within the SPCE and TIP4P potentials, whichsuggests that the proposed formalism is valid forusing the Casida equation with TOMBO program.
References[1] R. V. Belosludov, H. Mizuseki, R. Sahara, Y. Kawazoe, O. S.
Subbotin. R. K. Zhdanov, and V. R. Belosludov, Computational
Materials Science and Computer Aided Materials Design &Processing. In: K.-M. Lee, J. Kauffman (Eds.) Handbook of
Sustainable Engineering. Springer, New York, pp. 1215-1247(2013).
[2] M. S. Bahramy, M. H. F. Sluiter, and Y. Kawazoe, Phys. Rev. B
73 045111 (2006).

The C6 interaction coefficient for oxygen
molecules was 94.7 atomic units when measured atlong distance. This value overestimates thesemiempirical C6 parameter of the Lennard-Jones
potential (71.6 atomic units) usually used for theinteraction between oxygen molecules. The value

Keywords: dispersion coefficient, all-electron mixed basis approach, Casida equation
Yoshiyuki Kawazoe (Materials Design by Computer Simulation Division)
E-mail: mizuseki@imr.edu
URL: http://www.mdcs.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201324

Energy-related Materials

In situ Observation of Radioluminescence from IrradiatedLithium Tantalate

Ceramic materials are expected to play an important role in developing nuclear systems. Examples of suchmaterials are the diagnostic materials used in nuclear energy systems and the solid tritium breeders used innuclear fusion systems. Measuring the radiation-induced luminescence (i.e., radioluminescence) from irradiatedceramic materials is an effective method of monitoring the changes in the properties of the materials in aradioactive environment. A method of measuring the luminescence from ceramic materials was developed
using the tandem accelerator at the Institute for Materials Research, and the radioluminescence behaviors ofvarious ceramic materials subjected to various ion radiations were systematically studied. The results of recent
research on the candidate solid tritium breeder LiTaO3 is reported here.

Many irradiated ceramic materials emit photons
whose wavelengths are in the range 300–1000 nm
through electronic excitation. Typical examples ofsuch emission are the 692-nm emission from ruby(i.e., Cr-doped alumina) and the 420-nm emission(i.e., F-center emission) from sapphire. The featuresof the luminescence spectra (i.e., the shape, position,and intensity of the peaks, etc.) give information
about the microstructures of materials, such as thechemical composition of a material, the concentrationof defects and impurities in the material, etc. [1-4].
We developed an in situ method of observing
the radioluminescence emitted from irradiatedceramic materials in the 1.7-MeV tandem accelerator
at the Institute for Materials Research, as shown in

Fig. 1 Apparatus used to measure radioluminescence
under ion irradiation in 1.7-MeV tandem accelerator at
IMR.

25

IMR KINKEN Research Highlights 2013

Fig. 1. The method can be used to measure radiationinduced luminescence (i.e., radioluminescence) in
the range 300–1850 nm and can be used with othermethods of analyzing materials, such as Rutherford
Back Scattering (RBS, which gives information
about the change in crystal structures, impurity
distributions, chemical compositions, etc.), andelastic recoil detection analysis (ERDA, which gives
information about the distribution of hydrogen atoms
in materials), to elucidate how nuclear radiation
affects the strength and integrity of materials.
The correlations between the radioluminescence
spectra for and observed changes in the structuresof various ceramic materials for application to
nuclear systems (estimated using RBS, ERDA, andother methods) have previously been studied. One
example is the strong interaction between the Cr+
ions in sapphire (Al2O3) and the radiation-induced
oxygen vacancy (i.e., the F-center.) The decrease in
the 692-nm radioluminescence is strongly correlatedwith the increase in the 420-nm luminescence (i.e.,
the F-center created from oxygen vacancies) andrate of decrease of the luminescence strongly
depended on the rate at which vacancies were
produced (according to the calculation). Anotherexample is the interaction between the local structure
responsible for the radioluminescence from andhydrogen concentration in ceramic-based materials.
Some Li-containing ceramics, which are candidates
for tritium-breeding material, showed decreases in
the intensity of radioluminescence with increasinghydrogen concentration, implying that the localized
structure preferentially accommodates hydrogenatoms and that in situ measurement of the intensity

IMR KINKEN Research Highlights 2013of the associated radioluminescence is an effective
and convenient method of estimating theconcentration of hydrogen in the ceramics.
Fig 2 [4] shows the rates of annihilation andrecovery of the local structure responsible for the
540-nm radioluminescence of LiTaO3. The rates areplotted as functions of the nuclear stopping power
for the particular ions. The rates of annihilation andrecovery were estimated from a simple model ofproducing defects and from the change in theintensity of the 540-nm radioluminescence . Fig 2
clearly shows that the behavior of defect production
is strongly correlated with the nuclear stopping
power; namely, the rate at which point defects areproduced through atomic displacement. The socalled radiolysis; namely, the dynamic reduction of
the ceramics through electronic excitation and theresulting formation of structural point defects, isinconsequential for LiTaO3.
Materials absorb energy from the actual
radioactive environments of nuclear systems mainly
through electronic excitation. For example, energy
is absorbed through electronic excitation more than
100 times faster than through atomic displacement
in fission- and fusion-based energy systems. These
results imply that LiTaO3 will be relatively stable inactual radioactive environments.

Fig. 2. Rates of annihilation and recovery calculated fromchange in intensity of 540-nm radioluminescence in
LiTaO3.

References[1] K. Hoshi, S. Nagata, M. Zhao, and T. Shikama, Nucl. Instr.
Methods, accepted for publication in 2013.
[2] H. Katsui, S. Nagata, B. Tsuchiya, and T. Shikama, Nucl. Instr.
Methods, B272 275-279 (2012).
[3] M. Malo, S. Nagata, B. Tsuchiya, A. Morono, T. Shikama, andE. R. Hodgson, Fusion Eng. Des., 86 2470-2474 (2011).
[4] H. Katsui, S. Nagata, B. Tsuchiya, M. Zhao, and T. Shikama, J.
Nucl. Mater., 417 753-755(2011).

Keywords: ceramic, ion beam analysis, luminescence
Tatsuo Shikama (Nuclear Materials Science Division)
E-mail: shikama@imr.tohoku.ac.jp
URL: http://www-lab.imr.tohoku.ac.jp/~wshikama/

IMR KINKEN Research Highlights 201326

Energy-related Materials

Effects of Displacive Radiations in Fe and its Alloys

To develop materials for advanced fission and fusion reactors, we investigated irradiation effects in iron andiron-based alloys with special emphasis on modeling the microstructural evolutions and resulting changes in
the mechanical properties of the alloys. A533B alloy irradiated with ions at does up to 1 displacement per atom(dpa) revealed irradiation-induced hardening, and microstructural observation and electrical conductivity
measurements were used to determine the mechanism by which the hardening had occurred. Molecular
dynamics calculation was applied to investigate the behavior of defect clusters. The solute-defect interaction
affected the migration of the defect clusters, giving the characteristic length of the one-dimensional migration.

The mechanical degradations of nuclear
materials strongly depend on the environments ofthe materials. The alloys used in light-water reactors
are affected by irradiation-induced defects andpossibly by simultaneous interaction with coolant
water (i.e., corrosion). Further, the properties of thealloys used in fusion reactors deteriorate throughthe accumulation of both defects and helium.
Therefore, it is indispensable to understand themicrostructural evolutions and resulting mechanicaldegradations of the materials under the environmental
conditions the materials are located. We employ
irradiation techniques such as ion accelerators andhigh voltage electron microscopes, conducted withmicroscopic, spectroscopic, and mechanicalproperty measurements to reveal irradiation effects
in Fe- and Zr- based alloys in order to determine
how the radiation affected the alloys. The results ofthis research are further extended to materials
development.
This report describes irradiation-induced
phenomena in iron- based alloys.
(1) Irradiation Embrittlement of Reactor Pressure
Vessel (RPV) Steels [1]As one of the projects for the Enhancement,
Management and Maintenance of Aging Nuclear
Power Plants mandated by the Nuclear and Industrial
Safety Agency (NISA), the irradiation- induced
embrittlement of RPV steels was investigated using
ion beams, nano- hardness tester and electricalconductivity measurements to elucidate fundamental
insights into the radiation damage, as well as theembrittlement in the heat-affected zones (HAZs) in
weld joints.
Ion beam irradiation is one of the methods
capable for the precise correlation of the degree ofradiation damage and resulting hardening. The
correlation of irradiation hardening and embrittlement
is also desired in terms of microstructure. Bainitic
phase is a typical RPV steel (A533B) consisting ofcolonial distribution of carbides embedded in aferrite matrix. The mechanical properties of the steel
and their susceptibility to irradiation weresystematically investigated. The nano- hardness of27

IMR KINKEN Research Highlights 2013

Fig. 1. Nano-hardness of A533B steel irradiated with 2.8
MeV Fe2+ ions. Dose rate was set at 10 -5 and 10 -4 dpa/s.
Irradiation temperature was set at 563 K.

the A533B steel irradiated with 2.8 MeV Fe2+ ions at
563 K is shown in Fig. 1 Irradiation-induced
hardening is evident, and the hardening is
consistently greater in the carbide colony regions
than in the ferrite matrix. This finding suggests thatthe volume which contains carbide colonies is more
susceptible to irradiation hardening than the ferrite
matrix. The degree of radiation hardening wasroughly proportional to the square root of the dosein both regions up to 1 displacement per atom (dpa).
Significant hardening in the carbide colony regions
can be explained by the effect of the carbide/ferriteinterface, which works as sink for point defects. The
effect of the radiation dose rate on hardening was
also observed. The degree of hardening increased
with decreasing dose rate.
The electrical resistivity was measured todetermine how the solute atoms diffused throughout
the alloy matrix. The solute atoms trapped vacancies,
reducing the rate of vacancy annihilation andenhancing solute atom diffusion.
In order to obtain the technical basis to judge
the necessity of the in-reactor surveillance
specimens from the HAZ of weld joints, the neutron

IMR KINKEN Research Highlights 2013

irradiation was performed. The fracture toughnessof the irradiated specimens and the probabilistic
fracture mechanics analysis revealed that each HAZ
was susceptible to irradiation embrittlement. The
results suggest that fragmented micro-cracks wereformed at the fine grained HAZ (FGHAZ).
(2) Modeling glide motion of interstitial clusters
in random alloys [2, 3]
Clusters of up to a few tens of interstitial atoms
directly formed through large collision cascades
under high-energy neutron irradiation. In a recent
study of developing materials for future nuclear
power applications, the nature and behavior of smallinterstitial clusters were found to be practicallyimportant because small irradiation-induced
interstitial clusters and precipitates serve as
obstacles to dislocation motion and are responsiblefor irradiation-hardening, and because smallinterstitial clusters are glissile at a low activationenergy along the direction of the Burgers vector.
This type of gliding is called “one-dimensional (1D)
migration,” and it affects the evolution of defectstructures.
Fluctuation in the microscopic distribution ofsolute atoms serves as a barrier to the 1D migrationof interstitial clusters in random alloys. To better
understand and model 1D migration in dilute andconcentrated alloys, we studied how the energy offormation of the gliding interstitial clusters fluctuatedfor various alloys. We developed an analytical modelin which the total energy for the interaction between
an interstitial cluster and individual solute atoms wasa superposition of the potential for the interaction
between the cluster and individual solute atoms. We
then examined how the total interaction energy of agliding cluster fluctuated. The average amplitude of
the fluctuation was directly proportional to the squareroot of both the solute concentration and the cluster
radius. Although the distance separating the local
peaks in the fluctuation was virtually independent of
the solute concentration and the cluster radius, it
was only dependent on the range of the interaction
potential. We developed a model for anotherfluctuation in the interaction energy because ofsolute-solute interaction, which is effective at soluteconcentrations higher than a few percent.
The models were used to interpret the results ofthe molecular statics simulations in which the energy

Fig. 2. Schematic of present model describing interaction
between an interstitial cluster and solute atoms in
random binary alloys.

for the interaction between interstitial clusters in
dilute and concentrated Fe-Cu alloys on the one
hand and randomly distributed solute atoms on theother hand fluctuated under systematically varied
conditions. The typical height of the local energy
peaks in the spectra for concentrated alloys
exceeded a few electron volts even though theenergy for the individual solute-defect interaction
was around 0.1eV. The distance between the local
peaks was up to 4 nm. We proposed that the
distance was responsible for the short-range 1D
migration observed for the various alloys used in theelectron irradiation experiments [3]. The amplitudeof the fluctuation is attributable to the solute-defectinteraction, and gives the characteristic length of 1D
migration in alloys.
References[1] H. Abe , K. Onizawa, J. Katsuyama, K. Murakami, Y. Nagai, T.
Toyama, T. Iwai, T. Iwata, Y. Katano, S. Tamura, and N. Sekimura,
IAEA-CN-194, 072 (2012).
[2] Y. Satoh, H. Abe, and T. Matsunaga, Philos. Mag. 93, 1652
(2013).
[3] T. Hamaoka, Y. Satoh, and H. Matsui, J. Nucl. Mater. 433, 180(2013).

Keywords: radiation effects, electron microscopy, mechanical properties
Hiroaki Abe and Yuhki Satoh (Nuclear Materials Engineering Division)
E-mail: abe.hiroaki@imr.tohoku.ac.jp
URL: http://www-lab.imr.tohoku.ac.jp/~abeken/

IMR KINKEN Research Highlights 201328

Energy-related Materials

Three-dimensional Open-cell Nanoporous Iron-based Alloys

Three-dimensional open-cell nanoporous Fe-Cr alloys were prepared by dealloying Fe-Cr-Ni solid-solution
precursors in a metallic Mg melt. Ni was selectively dissolved into Mg, and the remaining Fe and Cr formed abicontinuous nanostructure when the precursors were immersed into the melt. The size of pores and ligaments
were well-tuned by varying the dealloying time and temperature.
Three-dimensional
open-cell
nanoporous
metals have attracted considerable attention for
their excellent functional properties including highcatalytic activity, sensing capabilities, and surfaceenhanced Raman scattering. The properties ofnanoporous metals differ significantly from those of
their dense, bulky counterparts. Such nanoporous
metals have conventionally been prepared bydealloying nanoalloys with aqueous solutions.
Dealloying is when the less-noble components in analloy are etched away and selectively dissolved andthe remaining noble components such as Au, Pt, Pd
etc. self-organize into a three-dimensional
nanoporous structure. The formation of a nanoporous
structure is highly desirable not only in noble metals
but also in less-noble metals used for various
functional applications including energy-harvestingFig. 1 (a) SEM micrograph and elemental mappings for (b)Fe and (c) Ni for cross-section of as-rolled Fe30Ni70
precursor. (d) SEM micrograph and elemental mappingsfor (e) Fe and (f) Mg for cross-section of Fe30Ni70
immersed into Mg melt at 1023 K for 1800 s.

29

IMR KINKEN Research Highlights 2013devices because less-noble metals are abundantand inexpensive. However, metal oxides usuallyform instead of nanoporous metals during
conventional dealloying because of the low stabilityof the less-noble metals. To overcome this limitation,
we have recently used a metallic melt instead of anaqueous solution to develop method of dealloyingthat increases the yield of nanoporous metals [1-3].
Selective dissolution originates from the interactions
among the atoms in metallic melts; therefore,
dealloying in metallic melts can even be used toprepare nanoporous less-noble metals. We used
dealloying in a metallic melt to prepare a nanoporous
Fe-based alloy.
Fig. 1a–c show the scanning electron
microscopy (SEM) image and energy dispersive
X-ray (EDX) mappings for the cross-section of the
as-rolled Fe30Ni70 precursor. The cross-section
shows a featureless contrast, and the Fe and Ni
elemental mappings show no fluctuations inconcentration. These findings are consistent withthe formation of a uniform solid solution. The
microstructure drastically changed into a finebicontinuous composite structure, as shown in Fig.
1d, after the precursor was immersed into a Mg melt
at 1023 K for 1800 s. The elemental mappings for Fe
and Mg shown in Figs. 1e–f reveal that the composite
is composed of Fe (bright region) and Mg (darkregion). From the equilibrium phase diagrams, theNi-Mg system is miscible while the Fe-Mg system is
immiscible. From the enthalpy of mixing, theinteraction between Fe and Mg is repulsive whilethat between Ni and Mg is attractive. These atomic
interactions induced the selective dissolution of Ni
from the Fe30Ni70 precursor into the Mg melt, andenabled the remaining Fe atoms to spontaneously
form a unique interconnected structure.
Porous Fe-based alloys were formed byremoving the Mg phase from the composites. Fig
2a–d show the SEM micrographs of the porous Febased alloys prepared by immersing Fe30Ni70,(Fe0.8Cr0.2)30Ni70, and (Fe0.2Cr0.8)30Ni70 into the Mg

IMR KINKEN Research Highlights 2013

melt at 1023 K for 10 min. As shown in Fig. 2a, thehomogeneous bicontinuous structure spans theentire sample in both the transverse and longitudinal
directions. The chemical compositions for these
porous samples are Fe99.8Ni0.2, Fe79.0Cr20.5Ni0.5,
and Fe19.4Cr80.2Ni0.4 (at.%). Almost no Ni remainedin any of the dealloyed porous samples, and the Fe/
Cr ratios for the dealloyed porous samples wereconsistent with those for the precursors. The size ofthe pores changed depending on the dealloying
temperature and immersion time; i.e., lower
temperatures and shorter immersion times resulted
in structures containing finer pores. The timedependence of pore and ligament size followed apower law whose component was nearly 4,
suggesting that surface diffusion plays an importantrole in forming nanoporous structures. Therefore,
the underlying mechanism for forming the porous
structure through dealloying in a metallic melt is thesame as that for forming the porous structure throughdealloying in an aqueous solution.
The preparation of nanoporous less-noble Febased alloys whose porous structures are tunable isa significant breakthrough for the cost-efficientdevelopment of highly functional materials.

Fig. 2 SEM micrographs for surface of porous samples
preparedbyimmersing(a,b)

Fe30Ni70,(c)(Fe0.8Cr0.2)30Ni70, and (d) (Fe0.2Cr0.8)30Ni70 precursors
immersed into Mg melt at 1023 K for 600 s and by
subsequently removing Mg. Chemical compositions
analyzed using SEM-EDX are labeled on eachmicrograph.

References[1] T. Wada, K. Yubuta, A. Inoue, and H. Kato, Mater. Lett. 65,
1076 (2011).
[2] T. Wada, A. D. Setyawan, K. Yubuta, and H. Kato, Scripta
Mater. 65, 532 (2011).
[3] T. Wada, H. Kato, Scripta Mater. 68, 723 (2013).

Keywords: foam, porosity, microstructure
Hidemi Kato (Non-Equilibrium Materials Laboratory Division)
E-mail: hikato@imr.tohoku.ac.jp
URL: http://www.nem.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201330

Energy-related Materials

Changes in Lattice Dynamics of Aluminum-based Complex
Hydrides during Hydrogen release

Aluminum-based complex hydrides (e.g., LiAlH , NaAlH ) are candidate materials for hydrogen-storage
44technologies because of their high gravimetric and volumetric hydrogen densities and their reversible hydrogen
release (dehydrogenation) and uptake (rehydrogenation) reactions. However, the atomic-level mechanism fordehydrogenation has not yet been well-understood. We used inelastic neutron scattering to study the changes
in the lattice dynamics of LiAlH4 with increasing temperature. The librational modes of the [AlH4]- anion showed
not only obvious broadening but also obvious softening in the range 250–300 K, which is interestingly much
lower than the dehydrogenation temperature for LiAlH4 (>426 K). The results suggest that dehydrogenation
originates from low-temperature bulk instability/anharmonicity because of the weakened bonds between thecomplex anions and their surroundings.
Hydrogen is expected to be an alternative clean
energy carrier for mobile applications. The
developmentof
efficienthydrogen-storage
technologies is thus a very important issue. There
are several candidates for storing hydrogen, suchas compressed gaseous and liquid hydrogen andhydrides. Hydrides store hydrogen as hydrogenatoms in a crystal structure and have higher
volumetric densities than compressed gaseous andliquid hydrogen. For this reason, many studies havepreviously been reported on using hydrides ashydrogen storage materials.
There are several types of hydrides in which
complex hydrides are ionically bonded to metal

cations and hydrido-complex anions. For instance,lithium aluminum hydride, LiAlH4, is composed of a
Li+ cation and four [AlH4]- complex anions withcovalent Al-H bonds (i.e., the so-called “alanates”).
They show much higher gravimetric hydrogendensities (e.g., LiAlH4: 10.6 mass%) than conventional
metal hydrides (e.g., LaNi5H6: 1.4 mass%) [1]. In
addition to the studies on using complex hydrides
for hydrogen storage materials, some studies show
lithium- and sodium-ion conduction in complex
hydrides containing [BH4]-, [NH2]-, or [AlH6]3complex anions [2, 3]. Therefore, complex hydrides
are interesting energy-related materials as not onlyhydrogen storage materials but also lithium andsodium-ionic conductors.
Aluminum-based complex hydrides
Since the discovery of reversible hydrogen
release (i.e., dehydrogenation) and uptake (i.e.,
rehydrogenation) in Ti-doped NaAlH4 undermoderate conditions (i.e., temperature: <473 K andhydrogen gas pressure: <10 MPa) [4], aluminumbased complex hydrides containing aluminum
hydrido-complex anions (i.e., alanates) have been
intensively studied [1]. Their decomposition reactions
are described by
3NaAlH4 ↔ Na3AlH6 + 2Al + 3H2 (g)↑↔ 3NaH + 3Al + 9/2H2 (g)↑.

Fig. 1 Crystal structure of LiAlH4 (top) and local structureof [AlH4]- surrounded by Li+ (bottom). Li, Al, and H
atoms and [AlH4]- are displayed as red, yellow, andblue circles and cyan tetrahedra, respectively.
LiAlH4 takes monoclinic structure with a = 4.8254 Å, b
= 7.8040 Å, c = 7.8968 Å, and β = 112.268º in P21/c
space group (No. 14) [6].

31

IMR KINKEN Research Highlights 2013

Such reversible reactions were recently
achieved for isomorphic LiAlH4 [5], which has ahigher gravimetric hydrogen density and a lower
dehydrogenation temperature than NaAlH4. The
crystal structure of LiAlH4 is illustrated in Fig 1.
Many researchers have attempted toexperimentally and theoretically elucidate themicroscopic mechanism for the dehydrogenation ofalanates on the basis of precursory phenomena.
However, the mechanism has not yet been wellunderstood.

IMR KINKEN Research Highlights 2013

Lattice dynamics of aluminum-based complex
hydrides [7]Inelastic neutron scattering (INS), on the other
hand, has often been used to investigate atomic
motions (i.e., lattice dynamics) in which the intensity
is proportional to inelastic neutron-scattering crosssections. Interestingly, hydrogen has a large inelastic
neutron-scattering cross-section (82.02 barns; 1barn = 1.0 × 10 -24 cm2), which is approximately 10
times larger than those for most other elements. In
addition, although INS has no selection rules,
studies on Raman and infrared spectroscopies use
them. For those reasons, numerous lattice dynamics
of hydrides have been elucidated using INS [8].
We therefore used INS spectroscopy toinvestigate the temperature dependence of thelattice dynamics of LiAlH4 below the dehydrogenation
temperature (426 K) in order to understand themicroscopic origins of the dehydrogenation reaction[7]. We particularly focused on the translational andlibrational modes (i.e., the external [AlH4]- motions)rather than the Al-H bending and stretching modes(i.e., the internal [AlH4]- motions) because theexternal [AlH4]- motions are expected to show moretemperature dependence than the internal ones
from previously reported Raman spectroscopy
results for NaAlH4 [9]. The INS spectra for LiAlH4
measured from 20 to 300 K and the characteristic
schematic figures of the translational and librationalmodes of [AlH4]- are shown in Fig. 2.
The overall spectra start to broaden withincreasing temperature in the range 100–150 K. In
addition, the librational modes of the [AlH4]- anion,
exhibit not only broadening but also softening in therange 250–300 K. This phenomenon occurs attemperatures surprisingly much lower than thedehydrogenation temperature. This result suggeststhat the dehydrogenation stems from lowtemperature bulk lattice instability/anharmonicity, inparticular, because of the weakening of the binding
between the [AlH4]- anions and their surroundings.
Further research is needed to understand thisphenomenon from the combined perspectives ofthe local structure and the present dynamic study.
This article was produced according to ref. 7.

Fig. 2 INS spectra for LiAlH4 measured in 20–300-K range
and characteristic schematic figures of translational(<400 cm-1) and librational (400–500 cm-1) modes [7].
Pink dotted line in spectra indicates softening above
250 K. Black arrows in schematic figures indicate
directions of atomic motions

References[1] S. Orimo, Y. Nakamori, J. R. Eliseo, A. Züttel, and C. M. Jensen,
Chem. Rev. 107, 4111 (2007).
[2] H. Oguchi, M. Matsuo, S. Kuromoto, H. Kuwano, and S. Orimo,
J. Appl. Phys. 111, 036102 (2012).
[3] M. Matsuo, S. Kuromoto, T. Sato, H. Oguchi, H. Maekawa, H.
Takamura, and S. Orimo, Appl. Phys. Lett. 100, 203904(2012).
[4] B. Bogdanović and M. Schwickardi, J. Alloys Compd. 253254, 1 (1997).
[5] X. Liu, H. W. Langmi, S. D. Beattie, F. F. Azenwi, G. S. McGrady,
and C. M. Jensen, J. Am. Chem. Soc. 133, 15593 (2011).
[6] B. C. Hauback, Z. Kristallogr. 223, 636 (2008).
[7] K. Tomiyasu, T. Sato, K. Horigane, S. Orimo, and K. Yamada,Appl. Phys. Lett. 100, 193901 (2012).
[8] T. Sato, A. J. (Timmy) Ramirez-Cuesta, K. Ikeda, S. Orimo, andK. Yamada, Inorg. Chem. 50, 8007 (2011).
[9] E. H. Majzoub, K. F. McCarty, and V. Ozolinš, Phys. Rev. B 71,
024118 (2005).

Keywords: energy storage, neutron scattering
Toyoto Sato and Shin-ichi Orimo (Hydrogen Functional Materials Division)
E-mail: toyoto@imr.tohoku.ac.jp, orimo@imr.tohoku.ac.jp
URL: http://www.hydrogen.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201332

Energy-related Materials

Growth of Shaped Crystals and their Application in Device
Fabrication

Device size–shaped crystal growth technology has been established using the micro-pulling-down method.

Devices using these shaped crystals have started to become commercially available. Shaped sapphire
crystals, langasite-type piezoelectric crystals for combustion sensors, scintillator crystals for pocket real-time
dosimeters, and Ir-based alloy crystals for spark plugs have been grown in specially shaped crucibles.
Introductions
Various large bulk crystals have been grownfrom melts with the Czochralski and Bridgeman
methods; these grown crystals are used in manyindustrial and research fields after forming processes
such as cutting from the bulk crystals and polishingof the cut crystals. However, the forming processes
after crystal growth greatly affect the manufacturing
cost and increase the prices of the final products
using the crystal elements.
The micro-pulling-down (μ-PD) method is a
crystal growth method that uses a crucible with a
die at the bottom. Compared to conventional
methods, the μ-PD method has a great advantage
in that a shaped single crystal can be grown with a
specially shaped crucible [1]. Because of its potential
for the near–net shape growth of a single crystal,
the μ-PD method is expected to be applied in mass
production.
Development of Shaped Crystals
Shaped sapphire crystals with various
configurations and multiple shaped sapphire
crystals were grown by the μ-PD method using
molybdenum crucibles with various shaped dies[2]. All shaped crystals had high transparency andFig. 1 Shaped sapphire crystals grown by μ-PD method.

33

IMR KINKEN Research Highlights 2013visible inclusions, and no cracks were observed in
the crystals. The crystallinity of each shaped crystal
was evaluated by X-ray rocking curves, and all
crystals indicated a high crystallinity of less than
100 arcsec.
We developed a novel Pt alloy crucible with asuitable wetting angle for a langasite-type melt.
Columnar, plate, and tube-shaped La3Ta0.5Ga5.5O14(LTG), Ca3NbGa3Si2O14 (CNGS), Ca3TaGa3Si2O14(SNGS),
and(CTGS),
Sr3NbGa3Si2O14
Sr3TaGa3Si2O14 (STGS) piezoelectric crystals weregrown using Pt alloy crucibles with special-shaped
dies [3]-[8]. The shaped langasite-type crystals
indicated comparable piezoelectric properties topreviously reported results. In addition, a combustion
sensor test device using the shaped langasite-type
crystal was developed.
After the nuclear accident in Fukushima, thedemand for dosimeters has greatly increased.
However, previous dosimeters using a semiconductor
had low sensitivity to radiation, and those using
scintillators with high sensitivity were expensive.
Therefore, we developed a growth method forshaped Ce-doped Y3Al5O12 (Ce:YAG) scintillator
crystals to decrease the price of dosimeters. Facetsof as-grown crystals can be controlled by the crystal
orientation of the seed crystal; the facet surfaces
had high transparency without polishing. A dosimeterusing the shaped Ce:YAG crystals without polishing
was developed, and it indicated good sensitivity toradiation.
Recently, a novel Ir alloy with high oxidation
resistance was developed and used for a spark
plug in an engine. However, the Ir alloy is unworkable,
and the thin wire had to be made from the Ir alloy
cylinder by scraping. In order to develop a shaping
technology for Ir alloy wire with one operation, wedeveloped a ceramic crucible that can be used at
2400 °C without evaporation or breaking, and ashaped Ir alloy wire was grown in it.

IMR KINKEN Research Highlights 2013

Fig. 2 Shaped piezoelectric, scintillator and Ir alloy crystals grown by μ-PD method and developed combustion sensor test
device and dosimeter using shaped crystals.

References

and A. Yoshikawa, IEEE Trans Ultrason Ferroelectr Freq Contr.

[1] A. Yoshikawa, M. Nikl, G. Boulon, and T. Fukuda, Opt. Mat. 30,

59, 1864 (2012).

6 (2007).
[2] A. Yoshikawa and V. Chani, MRS Bull. 34, 266 (2009).

[8] Y. Yokota, M. Sato, Y. Futami, K. Tota, T. Yanagida, K. Onodera,
A. Yoshikawa, J. Cryst. Growth 352, 147 (2012).

[3] Y. Yokota, V. Chani, M. Sato, K Tota, K Onodera, T. Yanagida,
and A. Yoshikawa, J. Cryst. Growth 318, 983 (2011).
[4] Y. Yokota, M. Sato, K. Tota, Y. Futami, T. Yanagida, K. Onodera,and A. Yoshikawa, Jpn. J. Appl. Phys. 50, 09ND03 (2011).
[5] Y. Yokota, M. Sato, V. Chani, Y. Futami, K. Tota, S. Kurosawa,
K. Onodera, and A. Yoshikawa, Sensors and Actuators A (inpress) doi.org/10.1016/j.sna.2012.10.017.
[6] Y. Yokota, A. Yoshikawa, Y. Futami, M. Sato, K. Tota, K.
Onodera, and T. Yanagida, IEEE Trans Ultrason Ferroelectr
Freq Contr. 59, 1868 (2012) .
[7] Y. Yokota, Y. Futami, M. Sato, K. Tota, K. Onodera, T. Yanagida,

Keywords: crystal growth, piezoelectric, radiation effects
Akira Yoshikawa (Advanced Crystal Engineering Laboratory)
E-mail: yoshikawa@imr.tohoku.ac.jp
URL: http://yoshikawa-lab.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201334

Energy-related Materials

Dissolution Behavior of REEs from Coal Fly Ash by
Extraction Using Dilute H2SO4 Solvent

In order to clarify the partitioning and dissolution behavior of rare earth elements (REEs) in coal fly ash particles,

tests of acid digestion and acid dissolution using ICP-Mass Spectrometry (ICP-MS) were conducted. The REE
concentrations followed the Oddo–Harkins rule. The unreacted core model fitting revealed that REEs exist bothon the surface of coal fly ash particles and inside them.
1. Introduction
In recent years, rare earth elements (REEs)
have become recognized as important materials
since they are widely utilized in the production ofmagnets, fluorescent materials, laser crystals, etc.
[1]. Since Japan is the largest consumer of REE
products, exploration of alternative supply routes isimportant for resource security.
Coal fly ash contains relatively highconcentrations of REEs at several hundred parts per
million weight. REEs are known to be soluble in anacidic solvent. In this study, we examined thedissolution behavior of REEs in dilute acidic solventsin order to determine the feasibility of REE recovery
using an acid-washing process.

digestion. The equipment for the dissolution test
comprised a Pyrex beaker, magnetic stirrer, heaterwith a stirring magnet, electrodes (to monitor thepH, redox potential, and temperature of thesolutions), and laptop computer. ICP-MS
measurements were then taken for the solutions
against the rhenium internal standard.

2. Materials and Methods
Three types of coal fly ash specimens wereused: Ash-A, Ash-B, and IRANT EOP. Ash-A and
Ash-B were provided by an electrical powercompany in Soma, Fukushima, in Japan. In order todetermine the REE concentrations in Ash-A and
Ash-B and verify the accuracy of the measurements
obtained by inductively coupled plasma mass
spectrometry (ICP-MS), the REE concentrations inthe specimens were determined by means of acid

4. Discussion
There seem to be two types of REE speciationin coal fly ash. Some of the REEs in the original coaltend to volatilize during coal combustion and arethen deposited onto the surface of coal fly ash
particles [2]. The other type of REE transition is theaccompaniment of blown-up ash components fromthe original coal. In this case, REEs are melted withash components and distributed into the coal fly ash
particles. If their dissolution is assumed to becontrolled by interlayer diffusion, the unreacted coremodel [3] can fit the dissolution rate very well with avariance of greater than 0.87. The intercepts for all of
the fitting lines were non-zero. Similar trends were
found for the rest of the REEs considered in thisstudy. Therefore, these fitting lines suggest two
types of REE occurrence. In the early stages ofdissolution, REEs at the surface of the coal fly ash
dissolve immediately to result in a drastic increaseof (3(1 − f)2/3 − 2f). REEs inside the coal fly ash
particles then dissolve gradually, which corresponds
to the linearity of the fitting line [4].

Fig. 1 Change in lanthanum dissolution rate over time.

35

IMR KINKEN Research Highlights 2013

3. Results
Fig. 1 displays the dissolution behavior oflanthanum in a dilute H2SO4 solvent as an example.
The dissolution rates increased with the temperatureof the H2SO4 solution, and Ash-B had a higher
dissolution rate than Ash-A. The other REEs show
dissolution behavior quite similar to lanthanum.

IMR KINKEN Research Highlights 2013

5. Conclusion
Laboratory-scale dissolution experiments wereconducted to determine the REE content in coal fly
ash particles and investigate their dissolution
behavior in H2SO4 solvents. In the dissolution tests,
dissolution rates of REEs increased with thetemperature of the H2SO4 solvent. The transition of
the REEs was attributed to surface deposition andaccompaniment, and the probabilities of bothprocesses could be explained by applying theunreacted core model to the dissolution behavior of
REEs from coal fly ash in H2SO4 solvent.

Fig. 2 Dissolution rates of REEs in a dilute H2SO4 solvent
fitted by unreacted core model assuming that dissolution

References

is controlled by interlayer diffusion.

[1] JOGMEC, “Material Flow of Mineral Resources,” 2010
[2] R. Meij. 39 199 (1994).
[3] S. Yagi and D. Kunii, Chemical Engineering Science, 16, 364(1961).
[4] S. Kashiwakura, Y. Kumagai, H. Kubo, and K. Wagatsuma.
Open Journal of Physical Chemistry, 3 69 (2013).

Keywords: mass spectroscopy, waste management
Kazuaki Wagatsuma, Shigeo Sato, and Shunsuke Kashiwakura (Analytical Science Division)
E-mail: shunsuke-k@imr.tohoku.ac.jp
URL: http://wagatsuma.imr.tohoku.ac.jp/index.html

IMR KINKEN Research Highlights 201336

Electronic Materials

Entangled Electromagnetic and Thermal Responses ofTopological Insulators and Superconductors

We studied the nontrivial responses of topological insulators and superconductors to determine theirelectromagnetic and thermal potentials. Magnetization was generated by application of an electric field, which
is known as the topological magnetoelectric (ME) effect. We investigated the stability of the ME effect andclarified its robustness with regard to randomness and a strong magnetic field. We generalized this crosscoupling to the case of thermal response and predicted that the temperature gradient generates orbital angularmomentum.
In 2005, the theoretical prediction of a newclass of materials known as topological insulators[1] promised novel physics and the potential forvarious applications. The key distinction fromordinary insulators lies in the fact that topologicalinsulators support stable gapless boundary modes.
Topological insulators owe their unique properties totime-reversal symmetry: if the flow of time is reversed,
the material would behave in the same way. One ofmost striking foundations is that the surface states
are protected from any perturbation, which preserves
the time-reversal symmetry when the bulk energy
gap remains intact. Anderson localization is
topologically forbidden on the surface of topologicalinsulators even in the presence of disorders.
On the other hand, magnetic impurities cause
an effective mass gap in the Dirac spectrum on thesurface. When the Fermi level resides in the surface
gap, a quantized anomalous Hall effect is realized.
As depicted in Fig. 1 (a), applying an electric field
produces a Hall current on the surface of the
topological insulator, which generates a magnetic

field, or magnetization in the bulk. This electricallyinduced magnetization without dissipation is referred
to as the topological magnetoelectric (ME) effect [1,2]. This type of cross-coupling is described by theenergy density functional(1)
where θ=π for a topological insulator and θ=0 for an
ordinary insulator.[2] The topological ME effect is
not only a useful signature of a topological insulator
but has a variety of potential applications. Therefore,
the stability of this effect is of central interest to
researchers in this community.
We studied the effect of magnetic disorder on
the surface of a topological insulator and clarified
that randomness unexpectedly enhances thesurface quantized anomalous Hall effect that helps
realize the topological ME effect [3]. Using scaling
analysis, we clarified that all surface states are
localized because of magnetic impurities, while thequantized Hall plateau width becomes enlarged

Fig. 1 Illustration of a magnetically doped topological insulator with (a)a cylindrical geometry and (b)a slab geometry.

37

IMR KINKEN Research Highlights 2013

IMR KINKEN Research Highlights 2013

from the original surface gap to all energy regions.
We recently studied the effect of bulk impurities [4]and electron–electron [5] interaction. The topologicalphase withstands such perturbations to someextent.
By using the quantized Hall surface states,
surface magnetization can be controlled throughthe application of both electric and magnetic fields:
to get the energy gain in Eq. (1), all surface spinstend to point out perpendicular to the surface. In thepresence of only the magnetic field, on the other
hand, all surface spins point in the same direction to
the field because of Zeeman coupling.
We recently examined this topological ME effectin a thin-film topological insulator [6]. When the widthof the film was not very thin, we found that the ME
effect worked even beyond the linear response
regime.
Topologicalsuperconductors
aresuperconductor analogs of topological insulators[1]. The bulk has a finite superconducting gap, while
there are gapless Andreev bound states on thesurfaces. Topological superconductors intrigue
physicists because they are expected to host exotic
particles on the surface known as Majorana fermionsthat differ fundamentally from any other known
particles in the sense that the particles are their own
antiparticles [1]. Since Majorana fermions are
expected to have applications in quantum
computation, detecting and manipulating them are
the great challenges in the field.
We predicted the existence of a unique andunexpectedsignatureof
topologicalsuperconductors: a temperature gradient applied toa topological superconductor generates mechanicalrotation, i.e., orbital angular momentum[7].
The magnetoelectric coupling in topologicalinsulators is because electrons moving along thesurface of a sample under the influence of an electricfield generates a magnetic field or magnetization in
the bulk. Similarly, Majorana fermions on the surfaceof topological superconductors moving under the
influence of a temperature gradient contribute to theorbital angular momentum. Simultaneously, thethermal Hall conductivity, which characterizes theheat transport of Majorana fermions on the surface,is quantized in terms of universal constants and themean temperature. This feature indicates that theeffect should be stable over extrinsic perturbations.
We obtained the free energy describing this nontrivial
thermomechanical coupling as [7](2)where v is the Fermi velocity and Ω is the angularvelocity. The Einstein–de Haas effect reveals therelationship between magnetism and angularmomentum of the system and has been used for the
study of ferromagnetic materials. We assumed acylindrical three dimensional (3D) topologicalsuperconductor suspended by a thin string andapplied a thermal gradient. This induced a surface
energy current with angular momentum Lz. By theconservation law of total angular momentum of thematerial, which can be directly measured in principle,
the inverse effect is the generation of thermal
polarization induced by rotating the material [7].
References[1] M. Z. Hasan and C. L. Kane, Rev. Mod. Phys. 82, 3045
(2010).
[2] X.-L. Qi, T. L. Hughes, and S.-C. Zhang, Phys. Rev. B 78,
195424 (2008).
[3] K. Nomura and N. Nagaosa, Phys. Rev. Lett. 106, 166802(2011).
[4] S. Ryu and K. Nomura, Phys. Rev. B 85, 155138 (2012).
[5] A. Sekine, T. Z. Nakano, Y. Araki, and K. Nomura, Phys. Rev.
B 87, 165142 (2013).
[6] D. Baasanjav, O. A. Tretiakov, and K. Nomura, in preparation.
[7] K. Nomura, S. Ryu, A. Furusaki, and N. Nagaosa, Phys. Rev.
Lett. 108, 026802 (2012).

Keywords: insulator, electrical properties, magnetic properties
Kentaro Nomura (Theory of Solid State Physics Division)
E-mail: nomura@imr.tohoku.ac.jp
URL: http://www-lab.imr.tohoku.ac.jp/~nomura/index.html

IMR KINKEN Research Highlights 201338

Electronic Materials

Bench-Top Pulsed Magnet for Tabletop SpectrometerAbench-top pulsed magnet generator was developed for spectroscopies such as photo luminescence andTHz spectroscopy. Magnetic fields of up to 30 T can be generated on a standard optical bench. The
compactness of the system is useful for a wide range of magneto-spectroscopy experiments. The system
extends the range of high magnetic fields in small laboratories by offering an extremely strong magnetic fieldof 30 T.
spectrometer. Even so, extremely strong magnetic
A magnetic field is a powerful and precise
fields of 30 T are easily obtained. The instrument
tuning knob for matter such as magnetic andand operational costs are quite small. For example,
semiconducting materials. It induces various new
the full set including a magnet shown in Fig. 1 costs
states with unexpected magnetic and electronic
properties. Photospectroscopy is widelyused as one of the most powerful means toinvestigate the states of materials. However,the combination of these two tools has
mostly been limited to magnetic fields ofless than 20 T at university laboratories.
This is because a superconducting magnet
is used in most cases; such magnets areeasy to handle, and a continuous magnetic
field is supplied, but the maximum field islimited. Resistive water-cooling magnets
and pulsed magnets that generate magnetic
fields of above 20 T are operational in
dedicated high magnetic field laboratories.
Pulsed magnets are running in smaller
laboratories having specialists for highmagnetic fields.
Besides the limited maximum magnetic
field, there are some technical difficulties incombining a superconducting magnet with
optical spectroscopy experiments. A large
stray field has a strong influence on
detectors and optical systems driven by
motors. The optical pass can be as long asa few meters in 20-T magnets; in manycases, an optical fiber is used in such
experiments. This is not compatible withtechniques such as wide-band THz time
domain spectroscopies, which need adirect optical pass.
Compact bench-top pulsed highmagnetic field generators, which weredeveloped by the magnetism division of
IMR, are a new breakthrough in this field.
Fig. 1(a) View of bench-top 30 T pulsed magnet installed in Rice
The idea is very simple. The energy needed
University. The magnet is installed inside the tall cryostat (leftfor a pulse magnet is known to scale withpanel). The capacitor bank is installed close to the optical table
both the magnet volume and the square of(right upper panel). The wave form is recorded by an oscilloscope
the magnetic field intensity. When a magnet(right middle panel). A sample is cooled by the micro-spectroscopy
is reduced in scale, the energy can becryostat in the center. In this photo, the cryostat has been temporarily
reduced, and the power supply can beremoved from the bam axis shown by arrows. An ultrafast laser,compacted [1]. Because of the compactnesslens system, and detector are arranged on the optical bench forof the magnet, the system can easily benormal zero field experiments (bottom panel).
placed on the optical bench of the39

IMR KINKEN Research Highlights 2013

IMR KINKEN Research Highlights 2013

only a few million JPY, and cheap liquid nitrogen isused for magnet cooling instead of expensive liquid
helium.
Fig. 1 shows the bench-top spectroscopy
setup. The magnet is installed in the compact liquid
nitrogen cryostat, and the magnet inner bore is 13
mm in diameter. The distance from the optical
window to the magnet center is only 35 mm. The
cryostat can be put on the optical bench, and no
special probe is needed for 30 T experiments. The
power supply stores an energy of 11.2 kJ andcharged with 2 kV. The energy required for 30 T isabout 9 kJ. The rise time and total pulse width are 2
ms and 10 ms, respectively. By combining thissystem with a compact cryostat used for microspectroscopy, a sample can be cooled down to
5–10 K. The experimental setup in Fig. 1 is at Rice
University, Houston. The first results of fluorescencein a semiconductor superstructure have beenreported elsewhere [2].
The system is quite versatile in terms of bothgenerating magnetic fields and application tospectroscopies. The 16 kJ system designed for the
Photon Science Institute of Manchester University,
as shown in Fig. 2, can be upgraded to charging at
4 kV. The system will be used for THz spectroscopies[3].
In summary, the bench-top 30 T pulsed magnet
system developed at IMR will extend the frontiers ofmagneto-spectroscopies
andincreasetheopportunities for new findings in research on
magnetic and semiconducting materials.
References[1]Y. Narumi, T. Nakamura, T. Kinoshita, Y. H. Matsuda, and H.
Nojiri, SRN 25, 12 (2012).
[2]H. Nojiri and Z. W. Ouyang, Terahertz Sci. Technol. 5, 1
(2012).
[3]G.Noe II, H. Nojiri, J. Leotin. G. Woods, A. Belyanin and J.
Kono, Proc. of CLEO2013, San Jose, June 9-14, (2013).

Fig. 2 16 kJ capacitor bank exported for the Photon
Science Institute at Manchester University using theMaterial Transfer Agreement protocol of Tohoku
University. The dimensions are 760 mm wide × 1000
mm long × 1400 mm high (upper panel). The capacitor,
charge-discharge, and high-voltage power supply
components are assembled as independent units; thus,
a flexible upgrade is possible. The inside view of thecharge-discharge unit. The high-voltage terminal is on
the right front side. The matching filter is on the left front
side (lower panel).

Keywords: high magnetic field, spectroscopy, instrumentation
Hiroyuki Nojiri (Magnetism Division)
E-mail:nojiri@imr.tohoku.ac.jp
URL: http://www.hfpm.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201340

Electronic Materials

Relaxor Ferroelectricity Induced by Electron Correlation in
Molecular Dimer Mott InsulatorT

he charge degrees of freedom in the molecular dimer Mott insulator β′-(BEDT-TTF)2ICl2 were investigated.
An anomalous peak structure with strong frequency dependence on the temperature dependence of thedielectric constant and corresponding glassy polarizations along the dimer direction were observed. These
results suggest that the charge degrees of freedom in dimers fluctuate and survive even in a dimer Mott
insulator.
The ferroelectricity induced by electrons [1], notconventional ionic displacements, is currently being
explored in the field of strongly correlated electron
systems, which is related to the charge-orderedstate of Mott insulators. Molecular dimer Mott
insulators are a good candidate for ferroelectricity
and are constructed from dimers of organic
moleculessuch
asbis(ethylenedithio)
tetrathiafulvalene (BEDT-TTF). When each dimer ina crystal possesses a charge carrier, particularly forthe half-filled case of a dimer molecular orbital, theelectron system often shows a Mott transition for
relatively large (onsite) Coulomb interactions U at
the dimer site compared to the electron hopping
energy t. An important physical insight is that a
dimer can be viewed as the primitive constituent ina crystal rather than as a molecule, as shown in
Figs. 1(a) and 1(c). In this manner, the understandingof strongly correlated electron systems based on
the dimer picture within the Hubbard model has
progressed [2].
However, recent reports on the dielectric
properties of several dimer Mott insulators [3-5]have implied that the charge degrees of freedom ina dimer survive even in the dimer Mott insulatingstate [6,7]. If a charge disproportionation occurs in
the dimers, electric dipole moments appear as
shown in Figs. 1(b) and 1(d). The fluctuation of
charges in dimers can be observed in the dielectric
properties. This discovery will enable a deeper
understanding of organic Mott insulators beyond
the successful dimer lattice model.
β′-(BEDT-TTF)2ICl2 is an anisotropic quasi twodimensional molecular dimer Mott insulator with a
one-hole carrier in the dimer of a BEDT-TTF
molecule. The molecular dimer forms the squarelattice shown in Figs. 1(c) and 1(d). The dimers arealigned along the b axis, arrayed in a column on thebc plane; the layers are stacked along the a* axis (⊥bc plane) and are well separated by ICl2 anion
layers. Dimerization was supported by NMR [8] andcalculation of the intradimer transfer integral. The
long-range antiferromagnetic order [9] below TN =22 K, as shown in Fig. 2(c) indicated that this system
is located inside the well-defined dimer Mott
insulating region with a large U/t value.
41

IMR KINKEN Research Highlights 2013

Fig. 2(a) shows the temperature dependence ofthe dielectric constant ε for β′-(BEDT-TTF)2ICl2 in
E||b parallel to the dimer array at each frequency[6]. An anomaly with a strong frequency dependence
was discerned at around 80–150 K. With decreasing
frequency, the anomaly in the dielectric constantbecomes large and forms a peak structure. A similarfrequency-dependent dielectric constant wasobserved for κ-(BEDT-TTF)2Cu2(CN)3 [3], which istypically seen in relaxor ferroelectric materials withrandomness and frustration.
The dielectric constant obeys the Curie–Weisslaw with a constant term: ε/ε 0 = C/(T – TC) + εconst./ε 0.
Here, ε 0, C, and TC are the vacuum permittivity,
Curie constant, and Curie temperature, respectively.
The Curie constant C was obtained according to thefit to the Curie–Weiss law. The dipole moment
parallel to the b axis was estimated to be 0.13ed,
where e and d are the electron charge and distance
between the BEDT-TTF molecules in a dimer,
respectively.

Fig. 1 Schematic crystal structure of κ-(BEDTTTF)2Cu2(CN)3 in (a) and (b) and β′-(BEDT-TTF)2ICl2 in(c) and (d). The upper two panels represent the dimer
lattice model, and the lower two represent the dipole
dimer model. The circles in the dimer lattice model are
the sites of the dimers where the molecules in the
dimers are ignored. The ovals in the dipole dimer model
represent the dimers. The circles in the oval aremolecules where a dipole moment can emerge if acharge disproportionation occurs.

IMR KINKEN Research Highlights 2013

The inset of Fig. 2(b) shows the fit to the data for
E||b of β′-(BEDT-TTF)2ICl2 to the Vogel–Fulcher (VF)relation for ε: f = f0 exp{−EVF/(Tmax − TVF)}, where f0 ,EVF , and Tmax are the representative flipping
frequency of electric dipoles, average random
energy barrier, and temperature at which ε takes amaximum, respectively. The empirical VF relation
accurately represents the frequency dependence ofthe peak in ε with TVF = 23 K, EVF = 804 K, and f0 =2.17 × 107 Hz. TVF is very close to the Neel temperatureof 22 K. This implies a possible relation between theglassy dielectric behavior and the magnetic
property.
Fig. 2(b) shows the electric polarization P with
E||b obtained by pyrocurrent measurements, and (c)shows the temperature dependence of the magnetic
susceptibility in a magnetic field with H||a*, b, and cfor β′-(BEDT-TTF)2ICl2. A small but measurable
polarization was observed below the temperature of
TFE = 62 K, which is nearly identical to the estimatedTC = 67 K from the Curie–Weiss behavior. However,
the experimentally obtained polarization of 6 nC/
cm2 at Epol = 1.2 kV/cm was ~100 times smaller thanthe expected full polarization using the Curie
constant. A fully polarized state has not yet beenobtained under poling electric fields of 1.2 k V/cm.
Thus, such a polarization indicates the formation ofglassy polar domains, which is also consistent withthe strong frequency dependence of ε.
The temperature dependence of the magnetic
susceptibility shown in Fig. 2 (c) is explained by a
two-dimensional Heisenberg model with an
antiferromagnetic interaction JAF of 59 K [9]. Aroundthis temperature, the magnetic susceptibility
significantly decreased, which indicates a shortrange spin correlation in the low-dimensional
system. This energy scale is comparable to TFE,
which also implies a relation between the glassy
nature of the dielectric and antiferromagnetic
interaction. Several scenarios based on the spincharge correlation in the dielectric anomalies havebeen proposed [3-7].
However, deepening understanding of thecharge-spin correlation and the dielectric is a topic
for future research.

Fig. 2 Temperature dependence of (a) dielectric constantfor frequencies from 20 Hz to 200 kHz in E||b, (b)polarization by pyrocurrent measurements along the baxis for poling electric field of +/-1.2 kV/cm, and (c)magnetic susceptibility in H||a*, b, and c for β′-(BEDTTTF)2ICl2. The inset in (b) shows the Vogel-Fulcher plot
for the frequency dependent dielectric constant. The
vertical dashed lines show the temperatures of TFE = 62
K and TN = 22 K.
H. Seo, C. Hotta, and H. Fukuyama, Chem. Rev. 104, 5005
(2004).
[3] M. Abdel-Jawad, I. Terasaki, T. Sasaki, N. Yoneyama, N.
Kobayashi, Y. Uesu, and C. Hotta, Phys. Rev. B 82, 125119(2010).
[4] P. Lunkenheimer, J. Mueller, S. Krohns, F. Schrettle, A. Loidl,
B. Hartmann, R. Rommel, M. de Souza, C. Hotta, J. A. Schlueter,
and M. Lang, Nat. Mater. 11, 755 (2012).
[5] S. Iguchi, S. Sasaki, N. Yoneyama, H. Taniguchi, T. Nishizaki,
and T. Sasaki, Phys. Rev. B 87, 075107 (2013).
[6] M. Naka and S. Ishihara, J. Phys. Soc. Jpn. 79, 063707(2010).
[7] C. Hotta, Phys. Rev. B 82, 241104(R) (2010).
[8] K. Miyagawa, K. Kanoda, and A. Kawamoto, Chem. Rev. 104,

References[1] S. Ishihara, J. Phys. Soc. Jpn. 79, 011010 (2010).
[2] H. Kino and H. Fukuyama, J. Phys. Soc. Jpn. 64, 2726 (1995),

5635 (2004).
[9] N. Yoneyama, A. Miyazaki, T. Enoki, and G. Saito, Bull. Chem.
Soc. Jpn. 72, 639 (1999).

Keywords: dielectric properties, electronic material, organic
Satoshi Iguchi and Takahiko Sasaki (Low Temperature Condensed State Physics Division)
E-mail: satoshi.iguchi@imr.tohoku.ac.jp
URL: http://cond-phys.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201342

Electronic Materials

Three-Dimensional Elemental Analysis in SemiconductorBased Actual Devices with Atom Probe Tomography

Large-scale integration and high performance of metal-oxide-semiconductor field-effect transistors havebeen achieved with continuous shrinkage and the emergence of complicated architectures (from planar-type
to multi-gate-type). The electrical properties rely on the element distribution. Hence, there is a growing needfor characterizing the elemental distributions in three dimensions (3D). Laser-assisted atom probe tomography(APT) has been attracting attention for visualizing elements in semiconductor materials with atomic-scale
resolution. The process of sample preparation using a focused ion beam (FIB) with a high-resolution scanningelectron microscope has been improved for site-specific analysis, which has contributed to nanoelectronicsrelated application studies. Herein, we show dopant distribution in planar-type actual devices for validating theorigin of electrical variability. Non-planar transistors can also be characterized using APT.

Metal-oxide-semiconductor field-effect transistors(MOSFETs) have supported advances in thecontemporary information society. The large-scale
integration and performance of these transistors
have been developed rapidly. Controlled dopant
position and density would gain prominence, as
device scales decrease to less than 30 nm. Since
the electrical properties of these devices are
sensitive to the spatial distribution of dopants inthem, an analytical technique for determining theexact location of each dopant is desired.
Atom probe tomography (APT) can be used for
determining the three-dimensional (3D) distributionof individual dopants in materials with nearly atomicscale resolution. In laser-assisted APT, a pulsed
laser combined with high standing voltages isapplied to a needle-shaped specimen, resulting in
the ionization and extraction of the atoms on theapex, also known as field evaporation. The elemental
distribution and mass-to-charge-state ratio of each
atom are determined using spatially resolved timeof-flight analysis.
The development of sample preparation
techniques using a focused ion beam (FIB) equippedwith a high-resolution scanning electron microscope(SEM) has greatly improved sampling productivity
for site-specific analysis. A combination ofsophisticated FIB techniques and APT instruments
has enabled us to obtain 3D dopant distribution in
device structures (Figs. 1(a) and 1(b)) and contributeto broad areas of research in device physics [1-4].
The characteristic variability in MOSFETs has
increased substantially owing to their continuously
decreasing feature sizes. In particular, the variability
in n-MOSFETs is larger than that in p-MOSFETs.
Through statistical analysis using APT, we foundthat an anomalous redistribution of dopant (boron)
occurs in n-MOSFETs owing to ion irradiation during
source/drain extension (SDE) formation (Fig. 1(c))[2]. To clarify the electrical variability, we directly
obtained the channel dopant concentration of actual
n-MOSFETs whose threshold voltage (V T) widely
43

IMR KINKEN Research Highlights 2013

Fig. 1 (a) Cross-sectional transmission electron microscope(TEM) image of our device structure. (b) 3D elemental
map of dopants in 65-nm-node n-MOS device
structures. (c) Normal probability plots for n- andp-MOSFETs with/without SDE formation.

ranging from ±4.0σ deviation from the median V T in
one million (1024 × 1024) transistors. We attempted
to fabricate a needle-shaped specimen (Fig. 2(a)).
Fig. 2(b) shows the elemental map of the main
dopants (boron, phosphorus, and arsenic), nickel,
and oxygen atoms. APT assessments revealed thatV T is positively correlated with the channel boron
concentration (Fig. 2(c)) [3].
With respect to multi-gate (non-planar) device
structures, dopant distribution analysis of finFET, apotential future device structure, was carried out.
Fig. 3(b) exhibits the elemental map of a prototype
boron-doped fin array, tilted at ~45° [4]. The point isthat near-surface dopant distribution in non-planar

IMR KINKEN Research Highlights 2013

Fig. 2 (a) Cross-sectional TEM image of 65-nm-node
MOSFET. (b) Elemental map of n-MOSFET (V T = 0.421
V). Nickel, phosphorus, arsenic, oxygen, and boron
atoms are shown. Silicon atoms are not shown. (c)Correlation between V T and channel boron
concentrations in 25 × 25 × 30 nm3 box just below gate
oxide.

Fig. 3 (a) Cross-sectional scanning TEM image of silicon
fin arrays. Observed region is indicated by broken
curve. (b) Elemental map of specimen tilted at ~45°. (c)One-dimensional concentration profiles of oxygen andboron across (i) top, (ii) sidewall, and (iii) bottom
interfaces.

References[1] K. Inoue, H. Takamizawa, Y. Shimizu, T. Toyama, F. Yano, A.
Nishida, T. Mogami, K. Kitamoto, T. Miyagi, J. Kato, S. Akahori,
N. Okada, M. Kato, H. Uchida, and Y. Nagai, Appl. Phys.
Express 6, 046502 (2013).

devices can be obtained using APT after being filledwith amorphous silicon by low-energy silicon beam
deposition without a significant temperature rise.
Oblique sample preparation is required for obtaining
a uniform spatial resolution across the edge regions
because the spatial resolution of APT in the lateral
direction is lower than that in the depth direction [5].
We found that the boron concentrations in the top
and bottom of the sample are higher than that insidewall. Nevertheless, the amount of boron in the
sidewall is higher than the amount of activated boron
resulting from thermal annealing.
Thus, the combination of a sample preparation
technique and APT enables us to characterize
dopant distribution in devices having complicatedstructures.

[2] H. Takamizawa, Y. Shimizu, K. Inoue, T. Toyama, N. Okada, M.
Kato, H. Uchida, F. Yano, A. Nishida, T. Mogami, and Y. Nagai,
Appl. Phys. Lett. 99, 133502 (2011).
[3] H. Takamizawa, Y. Shimizu, K. Inoue, T. Toyama, F. Yano, A.
Nishida, T. Mogami, N. Okada, M. Kato, H. Uchida, K. Kitamoto,T. Miyagi, J. Kato, and Y. Nagai, Appl. Phys. Lett. 100, 253504(2012).
[4] H. Takamizawa, Y. Shimizu, Y. Nozawa, T. Toyama, H. Morita,
Y. Yabuuchi, M. Ogura, and Y. Nagai, Appl. Phys. Lett. 100,
093502 (2012).
[5] Y. Shimizu, Y. Kawamura, M. Uematsu, M. Tomita, T. Kinno, N.
Okada, M. Kato, H. Uchida, M. Takahashi, H. Ito, H. Ishikawa,
Y. Ohji, H. Takamizawa, Y. Nagai, and K. M. Itoh, J. Appl. Phys.
109, 036102 (2011).

Keywords: atom probe tomography, 3D structure, electronic material
Yasuo Shimizu, Koji Inoue, and Yasuyoshi Nagai (Irradiation Effects in Nuclear and Their Related Materials Division)
E-mail: yshimizu@imr.tohoku.ac.jp
URL: http://wani.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201344

Electronic Materials

Anisotropic Features of m-Plane InGaN Growth

Nitride semiconductors such as InN, GaN, AlN, and their alloys are very promising materials for fabricating

low power consumption, high power, high-frequency devices. InGaN-based blue light-emitting diodes (LEDs)
and laser diodes (LDs) are already commercially available. New trends in nitride semiconductors in the regionof deep ultraviolet, green, and infrared wavelengths will promote energy-saving device applications andcreating new energies. Our studies are related to these new technology trends, including applications such as
phosphor-free white LEDs, brilliant green LEDs, high-efficiency solar cells, and the use of nitride semiconductorbased LDs in optical fiber communications systems.

Further progress in InGaN-based light-emitting
devices is required for their widespread use.
However, overcoming the decrease in the radiative
recombination rate due to the quantum-confined
Stark effect is a serious challenge in the longer
wavelength region, where In concentration andmisfit strain increase. To this end, non-polar andsemi-polar InGaN growth is an attractive alternative.
Furthermore, these surfaces have in-planeanisotropy in atomic structures depending on the
direction of the c-axis. The efficiency of In
incorporation into an InGaN film grown on a miscut[1010] m-GaN substrate depends on the sign of theoff-angle toward the [0001] (Ga-polar) or [0001](N-polar), as shown in Figs. 1 (a) and (c) [1].
For clarifying the origin of incorporation
asymmetry, the c-polarity was precisely determinedusing anomalousdispersion X-ray diffraction(ADXRD). The crystal structure factor of wurtzite
GaN hkh + kl reflection is given as follows: ,and its kinematic intensity for the same h, k,
and |l |is proportional to,

where fGa and fN are the atomic scatteringfactors of Ga and N, respectively, and u = 0.376 is
the N position’s internal parameter.
When a Cu-K X-ray is used, fGa and fN arepractically real. Then, the diffraction intensities of
the wurtzite hkh + kl and hkh + k l reflections are
identical. However, closely above the K (1s)absorption edge of Ga, the imaginary part of fGa issufficiently large (anomalous dispersion), and thedifference in the intensities of the two reflections is
detectable. It was estimated from the crystal
structure factor that the intensity of 2021 is larger
than that of 202 1, while the intensity of 2022 issmaller than that of 202 2. In addition, it was
confirmed that the kinetically and dynamicallycalculated intensities were nearly identical bymeasuring the reflections from a 330-μm-thick
m-GaN for a skew-symmetric geometry.
Three InGaN films, each approximately 150-nm
45

IMR KINKEN Research Highlights 2013

Fig. 1 2θ-ω scan spectra of (1010)-symmetric X-ray
diffraction using Cu Kα line from InGaN films grown on
m-GaN substrate. The numerical values in the figure
indicate the miscut angle of the substrate in the +cdirection. Thickness of the InGaN films is about 150 nm
for (a), (b), and (c), and 300 nm for (d).

thick, were grown at 750 °C by metalorganic vapor
phase epitaxy on +5°, 0°, and -5° miscut m-GaN
substrates. The conventional 2θ-ω X-ray-diffraction
spectra of these films are shown in Figs. 1(a), (b),and (c). Anomalous-dispersion x-ray diffraction
measurements were performed at beamline 13XU in
SPring-8. The c-polarity of the samples wasdetermined such that agreement between thecalculated and measured values was obtained forintensity variation along the 202 2, 202 1, 2021, and2022 series of reflections. As a result, it was
confirmed that the In incorporation efficiency at thepositive off-angles, where +c [0001]- direction sticks
out of the miscut m-surface (Fig. 2(a)), is higher thanthat at the negative off-angles (Fig. 2(b)).
Additionally, the anisotropy of the m-plane
inducesthecharacteristic
elongated-island
morphology and phase separation into different Incontent regions. Fig 3(a) shows a differentialinterference-contrast microscopy image of theislands elongated about 10 μm along the a-axis,
observed after the partial strain relaxation of the
InGaN film grown on a 0°-off m-plane GaN substrate.

IMR KINKEN Research Highlights 2013

The 2θ-ω X-ray diffraction spectrum of this film isshown in Fig. 1(d). In this case, an approximately
300-nm-thick InGaN film was grown under atriethylgallium flow rate that was 3× higher than thatfor the other three films shown in Fig. 1 for inducing
the desired island morphology. The phase separationof two different In concentrations can be inferredfrom the two InGaN peaks seen in Fig. 1(d).
To clarify the phase-separation mechanism, themicroscopic two-dimensional distribution of In
concentration in the InGaN film was observed withmicrobeam X-ray diffraction. The X-ray beam was
focused using a zone plate and slits to be 0.2 μm
and 0.5 μm in the vertical and horizontal full-width
half-maximum, respectively [2]. The In mole fraction,x, was estimated from the peak position in thereciprocal-space map by considering strain
relaxation due to misfit dislocations [1,3]. Fig 3(b)shows the real space mapping of x, where theincident position of the microbeam X-ray is scanned
in 0.2-μm steps along the c-axis and in 1.0-μm steps
along the a-axis by moving the sample. The sizeand shape of the high-x region near the center ofFig. 3(b) is comparable to that of the [0001]-side
slope (dark contrast area) of the elongated islands
in Fig. 3(a). Therefore, it is considered that the In
incorporation efficiency is high on the [0001] side of
the island surface. Lymperakis et al. calculated the
Ga adatom diffusion barrier of m-GaN as 0.93 eV in
the c-direction and 0.21 eV in the a-direction [4].
These anisotropic diffusion barriers are consistent
with the elongated-island morphology along thea-direction. Furthermore, the small diffusion length
in the c-direction induces the nucleation of islandson the m-plane terraces, thus increasing the stepdensity along the c-axis. The In incorporation
efficiency increases on the resulting slope toward[0001] owing to the nitrogen-terminated step
structure on this side (Fig. 2(a)).
In summary, the in-plane c-polarity of m-GaN
was determined using ADXRD. The In incorporation
efficiency increases when the +c-direction sticks
out of the miscut m-surface. In addition, it was foundthat the In mole fraction is high on the slope toward[0001] of the islands extended along the a-axis. The
formation of these islands is induced by theanisotropic features of m-plane growth and is
expected to be applicable for fabricating microsized structures.

Fig. 2 Atomic structure of miscut m-plane of InGaN, where
miscut angle is 5° (a) and -5° (b).

Fig. 3 Differential-interference-contrast- microscopy
image (a) and in-plane distribution of In mole fraction(b) of a partly strain relaxed InGaN layer grown on 0°-off
m-GaN substrate.

References[1] T. Hanada, T. Shimada, S. Y. Ji, K. Hobo, Y. H. Liu, and T.
Matsuoka, Phys. Stat. Sol. (C), 8, 444 (2011); 9, 1856 (2012).
[2] Y. Imai, S. Kimura, O. Sakata and A. Sakai, X-ray Optics andMicroanalysis: Proc. 20th Intern. Congress AIP CP1221, 30(2010).
[3] K. Shojiki, T. Hanada, T. Shimada, Y. H. Liu, R. Katayama, andT. Matsuoka, Jpn. J. Appl. Phys., 51, 04DH01 (2012).
[4] L. Lymperakis and J. Neugebauer, Phys. Rev. B, 79, 241308(R)(2009).

Keywords: nitride, vapor phase epitaxy (vpe), x-ray diffraction (xrd)
Takashi Matsuoka (Physics of Electronic Materials Division)
E-mail: matsuoka@imr.tohoku.ac.jp
URL: http://www.matsuoka-lab.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201346

Electronic Materials

Low Field Magnetization Switching of Hard Magnetic L10FePt Triggered by Spin Wave Excitation

Recent rapid progress in spintronic and magnetic storage devices has required nanomagnets to balancecompeting goals for high coercive field and low switching field. In this study, we experimentally demonstrated
extremely low field magnetization switching in a highly coercive L10 -FePt using a spin wave excited in a soft
magnetic Permalloy (Ni81Fe19; Py), where Py was exchange-coupled to L10 -FePt through the interface. Becausea significant decrease of one order of magnitude was achieved in the switching field when the spin wave was
excited, this spin wave–assisted magnetization switching is a promising technique for for manipulating themagnetization direction with ultralow energy.

47

IMR KINKEN Research Highlights 2013

magnetic moments in L10 -FePt, which isschematically shown in Fig. 1a. [3] The Py layer
thickness (tPy) was varied from 40 nm to 120 nm,
while the FePt layer thickness was fixed at 10 nm.
Fig. 1b shows the FMR spectrum of L10 -FePt/Py fortPy = 100 nm. We observed three resonance peaks
at the frequencies (f) of 8.0 GHz, 9.5 GHz, and 13.5
GHz, as indicated by the red, blue, and green
triangles, respectively. This experimental FMR
spectrum is well reproduced by numerical simulation.
From the numerical simulation, we found thatperpendicular standing spin wave (PSSW) modes
with n nodes were excited in the L10 -FePt/Pya

Saturated State

Twisted State

H

H

Py

Py

Fe

Fe

Pt

Pt

tPy = 100 nmbn=0Experiment
H = -700 Oe

Re[∆S11] (a.u.)

Spintronic and magnetic storage devices store
information using the magnetization directions ofmagnetic bits. For maintaining its magnetization
along a certain direction, a nanometer-sized magnet
should possess high magnetic anisotropy energythat can overcome thermal energy. However, highmagnetic anisotropy leads to a large external energy
for writing information onto a bit, i.e., for switching
the magnetization direction. Thus, in a magnet, thehigh thermal stability of magnetization, namely, highcoercive field (Hc), and the low switching field (Hsw)
are competing properties, both of which are crucial
for the advancement of spintronic and magnetic
storage devices. Therefore, we need to find asolution for this dilemma.
An energy-assisted switching method using
external microwave power, which is calledmicrowave-assisted magnetization reversal (MAMR)[1], is considered as a promising method for reducing
the switching field of a highly coercive magnet. It
applies a radio-frequency magnetic field (Hrf) to amagnet for inducing uniform magnetization
precession, thus resulting in decreased Hsw.
However, MAMR utilizes the ferromagnetic
resonance (FMR) mode, and its frequency (fFMR) isof the order of 100 GHz for a highly coercive magnet.
This high fFMR is a serious obstacle for practicalapplications. Thus, another method for reducing
Hsw while maintaining high Hc is eagerly desired.
The concept of magnetization switching we propose
here is based on the use of spin wave modes in asoft magnet for switching the magnetization in a
hard magnet [2].
We prepared exchange-coupled bilayer films
consisting of a hard magnetic L10 -FePt layer and asoft magnetic Permalloy (Fe19Ni81; Py) layer by
employing sputtering systems. In this L10 -FePt/Pybilayer, the L10 -FePt (110) layer was epitaxially grown
on a MgO (110) substrate, resulting in a uniaxial
easy magnetization axis along the in-plane [001]direction. Py was exchange-coupled with L10 -FePt
through the interface. The magnetic moments in Py
are spatially twisted when the magnetic field (H) isapplied in the direction opposite to that of the

n=1

n=2

n=0

Simulation
H = -700 Oe

n=1

0

2

4

6

8

f (GHz)

10

n=2

12

14

Fig. 1 a Magnetic structures in L10 -FePt/Py bilayer. The
left (right) panel depicts the direction of magnetic
moments in the saturated (twisted) state. A magnetic
field (H) was applied along the in-plane [001] directionof the MgO (110) substrate, which corresponds to theeasy magnetization direction of FePt. b FMR spectra oftPy = 100 nm for experiment (upper panel) and simulation(lower panel) at H = -700 Oe. The triangles indicate theresonant frequencies of the PSSW modes with n = 0(red), n = 1 (blue), and n = 2 (green).

IMR KINKEN Research Highlights 2013atPy = 100 nm∆R (Ω)

0-0.05 Hsw
-0.10
-0.15

Py

Py

FePt

FePt

Hrf = 0 Oe

-2000 -1000b∆R (Ω)

bilayer.
We measured the anisotropic magnetoresistance(AMR) curves of the microfabricated L10 -FePt/Py
element to examine magnetization switching under
PSSW-mode excitation. Fig. 2a shows the AMR
curve for tPy = 100 nm without the application of Hrf.
When H was swept from positive to negative(denoted by red circles), the device resistancechange (∆R) showed an abrupt decrease at H =-175 Oe, reflecting the formation of a twisted structureof magnetic moments in Py, as shown schematically.
R returned to the baseline of ∆R = 0 Ω at H = Hsw =-2050 Oe, where all magnetic moments of the bilayer
were switched. Fig. 2b shows the AMR curve under
the application of Hrf = 145 Oe with f = 8.0 GHz. Thisf corresponds to the resonance frequency of the
PSSW mode with n = 0. A decrease in Hsw down to- 325 Oe was obtained, indicating that magnetization
switching was induced by the PSSW-mode
excitation.
We achieved the desired large reduction in Hsw
for tPy = 120 nm. Hsw was reduced from 2000 Oe to
250 Oe at f = 7 GHz and Hrf = 134 Oe. This Hsw
reduction is significant compared to that achievedby conventional MAMR techniques, and theexcitation frequency remained less than 10 GHz,even for switching in the highly coercive L10 -FePt.
Because we simultaneously achieved the competing
goals of high Hc and low Hsw, spin wave-assisted
magnetization switching can be a new technology
for information writing across a wide range of
spintronic and magnetic storage applications.

0

H (Oe)

1000

2000

0
-0.05
-0.10

Hsw
Hrf = 145 Oe
f = 8 GHz

-0.15
-2000 -1000

0

H (Oe)

1000

2000

Fig. 2 Device resistance change (∆R) as a function of H
for microfabricated L10 -FePt/Py element a without
applying rf magnetic field (Hrf) and b under theapplication of Hrf = 145 Oe. The frequency (f) of Hrf was
8.0 GHz, and tPy was 100 nm. ∆R originates from theanisotropic magnetoresistance effect, where the device
resistance depends on the relative angle between the
directions of electric current flow and magnetic field.
The green arrows indicate the bilayer’s switching field(Hsw).

References[1] J.-G. Zhu, X. Zhu, and Y. Tang, IEEE Trans. Magn., 44, 125(2008).
[2] T. Seki, K. Utsumiya, Y. Nozaki, H. Imamura, and K. Takanashi,
Nature Commun., doi: 10.1038/ncomms2737 (2013).
[3] K. Utsumiya, T. Seki, and K. Takanashi, J. Appl. Phys., 110,103911 (2011).

Keywords: spin wave, magnetic properties, spintronic
Koki Takanashi (Magnetic Materials Division)
E-mail: koki@imr.tohoku.ac.jp
URL: http://magmatelab.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201348

Electronic Materials

Lithium-containing Oxide Film for All-solid-state Lithiumion Battery by Chemical Vapor Deposition

For the development of all-solid-state thin-film lithium-ion battery, the oriented textures and chemicalcompositions of the electrolytes and active materials should be controlled precisely for fast transportation oflithium ions. Furthermore, the deposition rate is required to be high for practical processes. Our group has
been developing a process for the chemical vapor deposition (CVD) of the oxide-based cathode and electrolyte
film. LiCoO2 cathode films with off-c-axis orientation and garnet-type, cubic Li7La3Zr2O12 electrolyte films weresuccessfully fabricated at high deposition rates by metalorganic and laser CVD.

The development of all-solid-state thin-filmlithium-ion batteries is required for reducing battery
size, mitigating safety issues associated with them,and fulfilling the power requirements of recent
electronic devices such as cellphones, electricvehicles, and energy storages modules. Oxidebased electrolytes and active materials have great
advantages because of their chemical stability inair, and the thin-film processing of these materials isnecessary for assembling all-solid-state lithium-ion
batteries.
Lithium cobaltate (LiCoO2) is widely used as acathode material in commercially available lithiumion secondary batteries. In all solid-state batteries,
optimization of the solid/solid interface between an
electrode and an electrolyte with high intercalation/
deintercalation of lithium ions in the cell is an
important issue. The lithium ion migrates in two
dimensions along a layer of lithium atoms within thelayered LiCoO2 structure (Fig. 1(a)). Therefore, theorientation and morphology of the LiCoO2 films

should be controlled for ensuring good transportationof lithium ions at a cathode–electrolyte interface.
The development of an electrolyte with high Li-ion
conductivity is another important issue. Garnet-type
cubic Li7La3Zr2O12 is a promising electrolyte material
because of its high lithium conductivity, wide
potential window, and chemical stability in air. Solidstate reaction mainly involved in the synthesis of
Li7La3Zr2O12, in which precise control over chemicalcomposition and temperatures above 1500 K are
required. The film fabrication of these active andelectrolyte materials has been studied by employing
physical vapor deposition (PVD) techniques suchas pulsed laser deposition (PLD) and sputtering. To
date, only a few research groups have reported
control over oriented texture under low deposition
rates, and the fabrication of well-formed Li7La3Zr2O12
films with good crystallinity has not been reported.
Chemical vapor deposition (CVD) allows the
preparation of films under high deposition rates,
while affording control over their morphology andFig. 1 (a) crystal structure of LiCoO2, (b)–(d) surface morphologies of (003)- and (104)-oriented LiCoO2 films by CVD, andof (101)-oriented LiCoO2 film by laser CVD, (e), (f) crystal structures of tetragonal and cubic Li7La3Zr2O12, (g)–(h) surface
morphologies of tetragonal and cubic Li7La3Zr2O12 by CVD.

49

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IMR KINKEN Research Highlights 2013

orientation. We have developed a laser CVD [1,2]process that could be an effective approach topreparing highly oriented LiCoO2 cathode films and
Li7La3Zr2O12 electrolyte films under high deposition
rates.
Preparation of LiCoO2 and Li7La3Zr2O12 films
A highly (003)-oriented LiCoO2 film was grownon polycrystalline Al2O3 substrate at a deposition
temperature of 923 K by CVD (Fig. 1(b)) [2].
Hexagonally faceted flat surface reflected c-planesin the rhombohedral LiCoO2 lattice. Rounded
platelet grains were grown in the (104)-oriented
LiCoO2 at 1123 K (Fig. 1(c)); the platelet facets inthese were (003) LiCoO2 planes tilted 55° relative tothe (104) planes. Fig.1(c) depicts the LiCoO2 films
by laser CVD under laser power of 150 W. This film
showed (101)-orientation, and the deposition rate
was 30 μm h−1, which is 100 times greater than thatreported in the literature.
Fig 1(g) and (h) shows Li7La3Zr2O12 films
prepared by CVD. A film comprising tetragonal
Li7La3Zr2O12 having polygonal facets was depositedat 1073 K (Fig.1 (g)), while a cubic Li7La3Zr2O12 filmwith fine granular surface morphology was formed
at a relatively high temperature of 1223 K. The
deposition rates of the tetragonal and cubic
Li7La3Zr2O12 films were higher than 20 μm h−1. In
addition to the Li7La3Zr2O12 crystalline film,
amorphous Li-La-Zr- O films with smooth morphology
were deposited using laser CVD. The deposition
rates of these Li7La3Zr2O12 films were more than 100μm h−1.
Oriented growth of LiCoO2 films
Hetero-epitaxial growth is a powerful method
for fabricating films oriented along a specific planeand direction on single crystals. LiCoO2 films
oriented along (104) and (018) were hetero-epitaxiallygrown on (100) and (110) MgO substrates by CVD[3, 4]. Fig. 1 shows the surface and cross-sectional
morphologies of the (104)- and (018)-LiCoO2
epitaxial films. The surface morphology of the (104)
LiCoO2 film contained diagonally arrayed rectangularfacets (Fig.1 (a)), whereas the (018) LiCoO2 film
exhibitedaroof-likestructureextendedunidirectionally (Fig.1 (b)) toward [110] MgO. Twinstructures were formed with in-plane relationships of[010] LiCoO2//[010] LiCoO2 and [421] LiCoO2//[010]Fig. 2 Oriented textures of LiCoO2 epitaxial films. The left
SEM images are the surface and the cross section ofthe (104)-oriented film, and the right images are thoseof (018)-oriented film. Schematic illustrations of the twin
structures are also included (i)-(iiii).

LiCoO2 in the (104) LiCoO2 film and with in-plane
relationships of [100] LiCoO2//[100] LiCoO2 in the(018) LiCoO2 film, thus reflecting the striations on
cross-sectional textures. The deposition rates of the(104)- and (018)-oriented LiCoO2 films were about
10 μm h−1, almost 100 times greater than thosereported for LiCoO2 epitaxial films fabricated bypulsed laser deposition.
References[1] T. Goto and T. Kimura, Thin Solid Films, 515, 46 (2006).
[2] H. Katsui, Y. Yamashita, R. Tu and T. Goto, J. Ceram. Soc.
Jpn., 121, 406 (2013).
[3] H. Katsui and T. Goto, Surf. Coat. Technol., 218, 57 (2013).
[4] H. Katsui, Y. Yamashita, and T. Goto, Key Eng. Mater., 508,
300 (2012).

Keywords: chemical vapor deposition (cvd), microstructure, electronic material
Takashi Goto, Hirokazu Katsui and Akihiko Ito (Multi-Functional Materials Science Division)
E-mail: goto@imr.tohoku.ac.jp
URL: http://www.goto.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201350

Electronic Materials

Unveiled Fermi surfaces in hidden order state of URu2Si2A

ctinide compounds are attracting interest in condensed matter physics because of their unique properties.
The 5f electrons are intermediate between itinerant and localized electrons, and show strong spin–orbit
interaction. URu2Si2 is a well-known uranium-based heavy fermion compound with a “hidden-order” transition,
the origin of which remains unknown. Herein, we present our recent results on quantum oscillation measurements
under extreme conditions with ultra-clean single crystals of URu2Si2. The properties of Fermi surfaces and theirinstabilities under strong magnetic fields are clarified, thus providing a clue to the origin of the hidden orderstate.
URu2Si2 is one of the most attractive actinide
compounds in the heavy fermion system. It shows asecond-order transition at 17.5 K; this is accompanied
by a large jump in specific heat. However, more than
25 years since the discovery of this transition, theorder parameter is remains unknown. Thus, thisordered state is called the “hidden order.”Unconventional superconductivity appears at low
temperatures (<1.5 K) as well, coexisting with thehidden-order state.
Recently, we clarified the Fermi surface
properties of URu2Si2 in the hidden-order statethrough quantum oscillation measurements under
extreme conditions such as at temperatures as low
as 30 mK and under magnetic fields as high as 35
T [1].
The Fermi surface is one of the most fundamental
physical quantities that characterize the metal. Thus,
it is important task for researchers to elucidate Fermi
surfaces experimentally. We have developed a highsensitivity quantum oscillation technique for the deHaas-van Alphen (dHvA) and Shubnikov-de Haas
(SdH) effect, according to which the magnetic
susceptibility or resistivity oscillates periodically as
a function of the inverse of magnetic field.
A high-quality single crystal is essential forthese measurements. Using the Czochralski method
in a tetra-arc furnace, as shown in Fig.1, we grewultra-pure single crystals of URu2Si2 with a residual
resistivity ratio (RRR) > 500, which is the best
possible quality of URu2Si2.
Fig 2 shows the angular dependence of the
dHvA/SdH frequencies in the hidden order state of
URu2Si2.
With regards to the low carrier numbers of
URu2Si2, all detected Fermi surfaces have low
volume, and the topology of the main Fermi surfaces
51

IMR KINKEN Research Highlights 2013is of the closed-type with a nearly spherical shape.
However, the cyclotron effective masses are
extremely large, indicating a heavy fermion system.
An anomalous split of the main branch, α, is observedin the basal plane in the tetragonal structure, andthe split immediately disappears when the field istilted to the c-axis, implying a fingerprint of thehidden-order state. High-magnetic-field experiments
reveal the Fermi surface instabilities with the fielddependent effective mass. Near 30 T, just below thecritical field of the hidden-order state, a new Fermi
surface with heavy quasi-particles was found. Theinterplay between the Fermi surface instabilities dueto the Zeeman spin-split leading to the so-called
Lifshitz transition and the field-dependent
characteristic gap energy plays an important role inthe collapse of the hidden-order state. Similar Fermi
surface reconstructions under high fields have been
detected in our thermopower measurements as well[2]. The polarized neutron experiments reveal theredistribution of magnetization below and above theFig. 1 High-quality single-crystal growth of URu2Si2 using
Czochralski method in tetra-arc furnace.

IMR KINKEN Research Highlights 2013

hidden-order transition, implying a fingerprint ofrank 5 multipoles [3]. In addition, a key experimenton thermal expansion under uniaxial stress was
performed [4].
This work was done in collaboration withmembers of CEA-Grenoble (Commissariat àl'énergie atomique et aux énergies alternatives),
LNCMI (Laboratoire National des Champs
Magnétiques Intenses), ILL (Institut Laue-Langevin),
and JAEA (Japan Atomic Energy Agency).
References[1] D. Aoki, G. Knebel, I. Sheikin, E. Hassinger, L. Malone, T.D.
Matsuda, and J. Flouquet: J. Phys. Soc. Jpn., 81, 074715(2012).
[2] A. Pourret, A. Palacio-Morales, S. Krämer, L. Malone, M.
Nardone, D. Aoki, G. Knebel, and J. Flouquet: J. Phys. Soc.
Jpn., 82, 034706 (2013).

Fig. 2 Angular dependence of dHvA/SdH frequencies in[3] E. Ressouche, R. Ballou, F. Bourdarot, D. Aoki, V. Simonet,

hidden-order state of URu2Si2. The frequency is

M.T. Fernandez-Diaz, A. Stunault, and J. Flouquet: Phys. Rev.

proportional to the extremal cross-sectional areas of theLett., 109, 067202 (2012).

Fermi surfaces [1].

[4] S. Kambe, D. Aoki, B. Salce, F. Bourdarot, D. Braithwaite, J.
Flouquet, and J.P. Brison: Phys. Rev. B, 87, 115123 (2013).

Keywords: actinide, Fermi surface, heavy fermion
Dai Aoki (Radiochemistry of Metals Division)
E-mail: aoki@imr.tohoku.ac.jp
URL: http://actinide.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201352

Electronic Materials

High-energy Neutron Scattering Study of Spin Excitations
in New Frustrated Spin Ladder System

Quantum spin systems have been studied extensively owing to their exotic ground states and non-trivial
phenomena. A spin-ladder system is one of ideal systems for testing a theory that bridges a one-dimensional
spin chain and a two-dimensional square lattice. To characterize the ground state and magnetic excitations of
the newly found two-leg spin ladder BiCu2PO6, we performed a time-of-flight inelastic neutron scattering
measurement on its single crystal. Our result indicates triplon excitation with the minimum energy of 2 meV at
the incommensurate momentum position. This indicates the formation of a rang-singlet at the ground state andthe existence of strong competing interactions along the ladder.
The physics of low-dimension quantum
antiferromagnets (AF) is fascinating because nontrivial behaviors emerge at low temperatures. The
spin-ladder system is a characteristic system thatbridges a one-dimensional (1D) spin chain and a
two-dimensional (2D) CuO2 layer. A two-leg spinladder system is indeed realized with the copperoxide families SrCuO3 [1] and Sr14−x (La,Ca)x Cu24O41[2]. Because superconductivity emerges in the lattersystem [3], the nature of antiferromagnetic (AF) spincorrelation in a spin-ladder system has been
investigated in connection with the mechanism ofhigh-Tc superconductivity [4]. However, the quite
large nearest-neighbor AF exchange interaction (J ≈2000 K) of these compounds prevents observationof the entire spin excitation spectrum [5]. No
experimental evidence of two-triplon (coupled
triplets on the lungs of ladder) continuum excitation,
which is similar to two-spinon continuum excitationin one-dimensional quantum magnets [6], has been
detected thus far. To make progress on this issue,
we have started an experimental as well as atheoretical study on BiCu2PO6, a new model system,
in which the S = 1/2 quantum spins of Cu2+ ionsconstitute two-leg spin ladders with relatively smaller
J (~100 K) [7]. Recently, we grew a high-quality
large single BiCu2PO6 crystal (Fig. 1). This progresshas enabled us to elucidate the full excitation
spectrum of the spin-ladder system through neutron
scattering measurements.
Sample preparation and characterization
A single crystal of BiCu2PO with a diameter of 8mm and length of 100 mm was grown using thefloating-zone method. The phase purity and sample
mosaic of the grown crystal were confirmed throughpowder X-ray measurements and the X-ray back
Laue method (Fig. 1), respectively. Our sample
shows a spin-gap behavior in magnetic susceptibility.
Two 35-mm-long samples were co-aligned for the
inelastic neutron scattering measurement. To study
the overall spin excitation spectrum of BiCu2PO6, in
which S = 1/2 quantum spins from Cu2+ ions form a
two-leg spin-ladder, we performed the first high53

IMR KINKEN Research Highlights 2013

Fig. 1 Grown single crystal of BiCu2PO6 (upper figure) andthe Laue pattern taken for a part of crystal (lower
figure).

energy neutron scattering measurement at the
chopper spectrometer 4SEASONS [7] installed at
J-PARC
Magnetic excitation spectrum in BiCu2PO6
As shown in Fig. 2, the magnetic excitation of
BiCu2PO6 was successfully observed in wide
momentum and energy spaces. The excitation along
the crystallographic b-axis shows clear dispersion
with the energy band between 2 meV and 28 meV,while that along c-axis and a-axis (not shown) is
non-dispersive. Therefore, the spin correlation is
one-dimensional, which is consistent with the resultof the antiferromagnetic two-leg spin-ladder model,
and a rung singlet is formed at the ground state.
Importantly, the spin gap opens at theincommensurate (0, 1 ± 0.45, 0.75) momentum
position along the leg direction. This result suggests
the existence of a large next-nearest-neighborinteraction, namely, spin frustration along the leg.
Thus, the BiCu2PO6 [8] system was confirmed to be

IMR KINKEN Research Highlights 2013

Fig. 2 Excitation spectrum of BiCu2PO6 along leg direction
at l = 0 (upper left figure), l = 0.75 (upper right figure),
and along rung direction at k = 0 (lower left figure) andk = 1 (lower right figure). Clear dispersion was observed
along the b-direction with the maximum energy of ~28
meV. The minimum energy of ~1 meV was observed atthe (0, 1 ± 0.45, 0.75) incommensurate position,suggesting the existence of frustration between theantiferromagnetic nearest and next-nearest neighborinteractions.

a frustrated spin-ladder system, which affords aunique opportunity to study the novel nature ofquantum spin systems. Furthermore, we investigated
the temperature dependence of the entire spin
excitation for characterizing the thermal evolution ofspin correlation in this frustrated spin-ladder system.
It was found that the intensity at 28 meV decreases
with increasing temperature, and the peak
disappears around 70 K. This temperature isconsistent with the temperature at which themagnetic susceptibility shows a broad maximum.
Therefore, the thermal evolution of the magnetic
branch extending at 28 meV closely correlates withthe change in magnetic ground state and thedegradation of spin-singlet formation at hightemperatures.

Fig. 3 Excitation spectrum of BiCu2PO6 along leg-direction
measured at 45 K (upper figure) and at 125 K (lower
figure). The intensity of excitation at ~28 meV decreases
at high temperatures, suggesting rung-singlet formation
degradation.

References[1] Z. Hiroi et al., J. Solid State Chem., 95, 230 (1991).
[2] E. M. McCarron et al., Mat. Res. Bull., 23, 1355 (1988).
[3] M Uehara et al., J. Phys. Soc. Jpn., 65, 92764 (1996).
[4] E. Dagotto, Phys. Rev. B, 45, 5744 (1992).
[5] S. Notbohm et al., Phys. Rev. Lett., 98, 027403 (2007); R.S.
Eccleston et al., Phys. Lev. Lett., 81, 1702 (1998).
[6] M. Arai et al., Phys. Rev. Lett., 77, 3649 (1996).
[7] M. Nakamura et al., J. Phys. Soc. Jpn., 78, 093002 (2009).
[8] B. Koteswararao et al., Phys. Rev. B, 76, 052402 (2007).

Keywords: neutron scattering, spin dynamics, quantum spin system
Motoki Fujita (Materials Processing and Characterization Division)
E-mail:fujita@imr.tohoku.ac.jp
URL: h ttp://www.imr.tohoku.ac.jp/en/org/research/27.html
http://www.yamada-lab.imr.tohoku.ac.jp/jp/index.html

IMR KINKEN Research Highlights 201354

Electronic Materials

Electrical Transport Properties of Ni-Nb-Zr-H Amorphous
Alloy Ribbons

E

lectrical resistivity measurements were carried out for investigating the effects of hydrogen absorption on theelectrical transport properties of Ni36Nb24Zr40 amorphous alloy ribbon. The amount of hydrogen absorbed canbe controlled by changing the current density of electrochemical charging. The electrical resistivity showednegative temperature dependence, and absolute value of the temperature coefficient of resistivity, TCR,
increased as the amount of hydrogen absorbed increased. The negative enhancement of TCR could haveoriginated from an increase in the residual resistivity owing to the induced hydrogen and vacancy.
Ni-Nb-Zr ternary alloys have been known toform an amorphous phase in wide concentration
regions [1]. The hydrogen permeability of Ni-Nb-Zr
amorphous alloy ribbons was investigated by
Yamaura et al., and an excellent permeability valueof about 1.59 × 10 -8 mol/m-s-(Pa)1/2 was reported
for melt-spun amorphous (Ni0.6Nb0.4)50Zr50 alloy[2,3]. In addition, the mechanical properties ofhydrogen-absorbed amorphous alloy ribbons were
investigated by Kawashima et al [4]. Recently,
peculiar behaviors of the electrical transport
properties such as superconductivity, electron
avalanche behaviors, and Coulomb-blockade
oscillations of hydrogen-absorbed [(Ni0.6Nb0.4)1x/100Zrx/100] (30 < x < 50) amorphous alloy ribbons
were elucidated by Fukuhara et al. [5,6]. They
pointed out that the localization effect of hydrogen in
outside and inside spaces of the distorted
icosahedral Zr5Ni5Nb3 clusters plays important rolesin their various electrical transport properties. In
other words, the state of hydrogen absorption isthought to be sensitive to the difference in the local
structure of the glassy state [3]. Therefore, asystematic study of the effects of hydrogen
absorption on the electrical transport properties of
the Ni-Nb-Zr amorphous alloy is important forclarifying the mechanisms underlying such peculiar
behaviors.
We made amorphous alloy ribbons of
Ni36Nb24Zr40, and controlled the amount of theabsorbed hydrogen by changing the current densityof electrochemical charging. The amorphous alloy
ribbons were fabricated by the single roller meltspinning technique, where the surface velocity of
the copper wheel was about 50 m/s. The thickness
and width of the fabricated ribbon were about 25
μm and 1 mm, respectively. Hydrogen-absorbed
specimens were prepared by electrochemicalcharging in 0.5 M H2SO4 + 1.4 g/L thiourea(H2NCSNH2) at room temperature with a current
density of about 10–120 A/m2 for 2 h. The amount ofabsorbed hydrogen was measured using the inert
gas carrier melting-thermal conductivity method
several times under each condition. Thermal
analyses were performed with a differential scanning55

IMR KINKEN Research Highlights 2013

calorimetry under a heating rate of 0.33 K/s. Phasestate was identified using x-ray diffraction with CuΚα radiation. The electrical resistivity was measured
using a conventional four-probe method in the
6–300 K temperature range with cooling and heating
rates of about 0.017 K/s.
Hydrogen absorption effects on electricalresistivity
Fig. 1 shows mass concentration of thehydrogen absorbed by the electrochemical method
as a function of current density. Here, the charging
time was fixed to 2 h. The amount of absorbedhydrogen in the ribbon specimens was found toincrease with increasing current density and to havea tendency to saturate. The maximum value ofhydrogen absorption was thought to be limited to a
maximum of 0.25 mass% (15 at.%), because thespecimens turn very brittle if they are charged withthe higher current density of 120 A/m2.
Fig. 2 indicates the x-ray diffraction patterns ofa Ni36Nb24Zr40 melt-spun ribbon and specimens
charged electrochemically for 2 h under current
densities of about 40 and 70 A/m2. The amounts ofabsorbed hydrogen in these specimens were foundFig. 1 Mass concentration of absorbed hydrogen as afunction of the current density in the electrochemicalcharging of Ni36Nb24Zr40 amorphous alloy ribbons. The
dashed curve is a visual guide.

IMR KINKEN Research Highlights 2013

Fig. 2 X-ray diffraction patterns of Ni36Nb24Zr40 melt-spun
amorphous alloy ribbon and hydrogen-absorbed
specimens.

Fig. 3 Temperature dependence of normalized electricalresistivity, ρ/ρ 300K, for (Ni0.36Nb0.24Zr0.40)100-y Hy with y =0.0, 2.0, 3.3, 9.2 and 14.7.

to be about 9.2 and 14.7 at.%, respectively. The
x-ray diffraction patterns show typical halo-type
patterns, indicating an amorphous structure. In Fig.
2, it can be seen that the main peak around 40°
shifts to a slightly smaller angle with increasinghydrogen absorption, thus indicating volume
expansion.
Fig. 3 shows the temperature dependence of
normalized electrical resistivity, ρ/ρ 300K, for y = 0.0,
2.0, 3.3, 9.2, and 14.7 in (Ni0.36Nb0.24Zr0.40)100-y Hy.
All resistivity exhibits negative temperaturedependence, according to the Baym-Meisel-Cotetheory based on the Boltzmann transport equation
for amorphous alloys with non-periodic crystal
structure, as follows [7]:absolute value of temperature coefficient ofresistivity, TCR, increases with increasing absorbedhydrogen, y, clearly indicating the effect of hydrogen
absorption on electrical resistivity. With regards theorigins of this effect, there are many possibilities,e.g., the change in the mean free path associated towith volume expansion, change in the Debye
temperature and electronic state, and the existenceof scattering related to absorbed hydrogen.
Clarifying this mechanism is the next objective of
our study.
References[1] H. Kimura, A. Inoue, S. Yamaura, K. Sasamori, M. Nishida, Y.
Shinpo, and H. Okouchi, Mater. Trans., 44, 1167 (2003).
[2] S. Yamaura, Y. Shimpo, H. Okouchi, M. Nishida, O. Kajita, H.

(1)

Kimura, and A. Inoue, Mater. Trans., 44, 1885 (2003).
[3] S. Yamaura, M. Sakurai, M. Hasegawa, K. Wakoh, Y. Shimpo,

where ρ 0 is the residual resistivity, exp[-2W(T)] is the
Debye-Waller factor, ρph is the phonon term, Λ is themean free path, and q is the wave vector. For theamorphous and quasi-crystals with non-periodic
structures, Λ is small, and, consequently, the second
term in Eq. (1) becomes negligible. Therefore, theelectrical resistivity is governed by the temperaturedependence of the Debye-Waller factor and not by
the electron-phonon scattering. In such a case, theelectrical resistivity shows T2-dependence in the
low-temperature range and T-linear dependence
close to room temperature [8,9]. As shown in Fig. 3,

M. Nishida, H. Kimura, E. Matsubara, and A. Inoue, Acta Mater.,
53, 3703 (2005).
[4] A. Kawahima, S. Yamaura, N. Ohtsu, H. Kimura, and A. Inoue,
Mater. Trans., 6, 1523 (2006).
[5] M. Fukuhara, A. Kawashima, S. Yamaura, and A. Inoue: Appl.
Phys. Lett., 90, 203111 (2007).
[6] M. Fukuhara, H. Yoshida, K. Koyama, A. Inoue, and Y. Miura,
J. Appl. Phys., 107, 033703 (2010).
[7] L.V. Meisel and P.J. Cote, Phys. Rev. B, 16, 2978 (1977), ibid.
17, 4652 (1978).
[8] U. Mizutani, Prog. Mat. Sci., 28, 97 (1983).
[9] U. Mizutani, Phys. Stat. Sol. (b), 176, 9 (1993).

Keywords: amorphous, electrical resistivity
Rie Y. Umetsu (Advanced Materials Development and Integration Novel Structured Metallic and Inorganic Materials)
E-mail: rieume@imr.tohoku.ac.jp
URL: http://amdi-pro.imr.tohoku.ac.jp/en/index.html

IMR KINKEN Research Highlights 201356

Research Centers

Recent Facility Progress at Oarai Center for Material
Irradiation Testing and Actinide Materials Science
International Research Center for Nuclear Materials Science (Oarai Center)

Facilities of the Oarai Center are open to experimental studies that involve the use of nuclear reactors and

α-emitters materials such as actinides. Recent facility progress is described as follows. A new irradiation rig
system that provides several different neutron doses under the same irradiation conditions such as neutron-flux
and irradiation-temperature is being developed. The old tetra-arc furnace with a glass bell jar has been renewed
for growing high-quality single crystals of uranium and thorium compounds. Low-temperature measurementsof nuclear magnetic resonance and Mössbauer spectroscopy, as well as those of electrical resistivity havebeen achieved by installing a helium-free refrigerator.

Neutron irradiation using material-testing
reactors is common in research on the effects ofirradiation on materials. For a more precise
understanding of irradiation dose dependency, themaintenance of identical irradiation conditions, i.e.,
irradiation rate (neutron flux), neutron spectrum, andirradiation temperature, are necessary for different
doses. We have started to develop new irradiation
rig for providing such series irradiation tocollaborators.
Materials irradiation at the Oarai center has
been performed using the BR2 reactor of SCK∙CEN,
Belgium Nuclear Research Center, after thesuspension of JMTR in 2006. A new irradiation rig
system, LIfting Basket in the Experimental Rig forBR2 Thimble tube sYstem or “LIBERTY,” is being
developed under the framework for collaboration
between IMR and SCK∙CEN. LIBERTY’s irradiation
specifications are approximately the same as thoseof JMTR’s rig system in terms of maintaining samples
with references under identical conditions of neutron
flux and temperature [1].
The new rig system has been improved for
fulfilling the specifications and incorporating thefollowing three features. 1) Individual needles canbe lifted independently above the reactor core level
when the specified fluence is reached, while theother needles remain in the neutron flux. 2) Individual
needle can be instrumented separately. Therefore,
electrical heating wires and thermo-couples can beinstalled into the needles for controlling specimen
temperature during irradiation. 3) Relatively large
specimens such as CT specimens (10 × 10 mm²)
can be tested.
The
concept
underlying
LIBERTY
isschematically shown in Fig. 1 (a); LIBERTY consistsof one shroud tube containing up to five needles
with samples inside. Initially all needles are inserted
at the bottom position of the shroud, which is in thereactor’s core region. When the desired irradiationdose is achieved, the needle is lifted from the bottom
to the top of the shroud. Each needle can be lifted
57

IMR KINKEN Research Highlights 2013individually, thus allowing for five different doses. In
each needle, we can install holders with samples,thermocouples, and electric heating elements(heating wires). A cross-sectional view of LIBERTY’s
core region is shown in Fig. 1 (b); it is worth notingthat it is made of the stainless steel shroud tube, of
the five larger needles and of the center tube
containing the gamma thermometer. The coolingwater first flows downward around the needles
through the shroud tube, turns upward, and flows
between the outer and inner walls of the shroud
tube. The rig itself is reusable, and it does not need
to be sent to the hot cell for needle loading andunloading; these operations can be performed in
the water pool. LIBERTY is now being developed,
as shown in Fig. 1 (c), and will be available in FY2013
for collaborative studies.
It is now widely recognized that actinides

Fig. 1 (a) Concept of “lifting-needle irradiation,” (b) crosssection of core region, and (c) photograph of LIBERTY
prototype.

IMR KINKEN Research Highlights 2013

represent a unique part of the periodic table ofelements. Many interesting physical and chemical
properties of actinide materials in the solid and liquidstates are derived from the special nature andcharacteristics of their open shell of 5f electrons.
Both sample preparation for high quality, advanced
microscopic probe measurement such as dHvA,
and NMR and Mössbauer spectroscopies are very
important for elucidating the intrinsic role of 5f
electrons.
The aging tetra-arc furnace with defects in both
ultimate base vacuum and stability of electrodecurrents was being used for limited sample
preparation of polycrystalline intermetallics andlanthanide-based alloys. A new tetra-arc furnacewith a Czochralski puller on top of the stainlesschamber was installed in the Actinide Laboratory in2012 to grow high quality single crystals (see Fig.2).
High base vacuum after 1 h of evacuation (5 × 10 -6Torr) and agitation of the titanium getter by the fifthelectrode prevent oxidation of the active melting
button. Each arc current can be controlled preciselyfrom 3 to 20 A for several hours for growing acylindrical sample of the single crystal that has atypical diameter of 3–4 mm and is few-millimeterslong. This machine allows for the melting and pullingof the uranium- and thorium-based systems with low
vapor pressures and rather high melting
temperatures.
Low-temperature physical properties werebeing measured by conventional helium-type
cryostats at the Oarai center. A cryogen-freesuperconducting magnet with a bore size of 100
mm, homogeneity less than 0.01 %, and maximum
field of 6 T (JASTEC JMTD-6T100HH) has been
installed for NMR spectroscopy. Another closedcycle refrigerator for sample cooling from 4.5 K toroom temperature (JANIS SHI-950) has beenarranged for the superconducting magnet. We canperform not only NMR spectroscopy but alsomeasure electrical resistivity with this cryostat. The
high-temperature system with the electric furnaceand the very-low-temperature system using ADR in
this magnet station are under construction.
Mössbauer spectroscopic measurements at
low temperatures are possible using the closedcycle refrigerator with beryllium windows (Taiyo
Nissan corp., Model-331). Temperatures of the
Mössbauer source and absorber (sample) arecontrolled individually for enhancing recoil-lessFig. 2 Photograph of tetra-arc furnace.

nuclear resonance, even those of the high-energy
γ-rays associated with 237Np and 197Au Mössbauer
effects. Conventional Mössbauer measurements of
57Fe and 151Eu nuclides have been also performed
using this cryostat. Mössbauer spectra reflect thesource conditions, especially those of the 241Am
source for the 237Np Mössbauer effect, where alphadecay and strong recoil-less by high-energy γ-rays
damages the atomic arrangement of the source
matrix. A fine splitting spectrum resulting from thehyperfine field could not been obtained using theold 241Am metal source [2]. We prepared the Pd10%(Pu,Am) and Pd-5%Am alloys following acoupled reduction process starting with a mixture of
Pd and dioxides in He-4%H2 atmosphere at 950 °C.
It was demonstrated that the Pd-Am alloy is an ideal
source for 237Np Mössbauer spectroscopy owing to
its narrow half width and oxidation resistance.
References[1] A. Kuramoto, T. Toyama, T. Takeuchi, Y. Nagai, M. Hasegawa,T. Yoshiie, Y. Nishiyama, J. Nucl. Mater., 425, 65 – 70 (2012).
[2] Y. Homma, M. Nakada, A. Nakamura, S. Nasu, D. Aoki, H.
Sakai, S. Ikeda, E. Yamamoto, Y. Haga, Y. Onuki, Y. Shiokawa,
Hyperfine Interact., 168, 1 175-1179 (2006).

Tatsuo Shikama (International Research Center for Nuclear Materials Science (Oarai Center))E-mail: shikama@imr.tohoku.ac.jp
URL: http://www.imr-oarai.jp/

IMR KINKEN Research Highlights 201358

Research Centers

Recent Progress in Development of Advanced Materials
Cooperative Research and Development Center for Advanced Materials(Formerly Advanced Research center of Metallic Glasses)T

he center has a critical role to play in the investigation and development of new materials with importantfunctional properties; these materials range from metallic compounds to oxides and from amorphous structures
to bulk single crystals. Researchers/students can collaborate with the various research divisions of the centerand use its experimental facilities under the inter-university research program.
In April 2013, the center was reorganized andrenamedas"Cooperative
ResearchandDevelopment Center for Advanced Materials" formaking a fresh start in promoting research,principally, on advanced materials. The underlying
philosophy of research in the center is to find andestablish the fundamental principles and technology
for synthesizing and industrializing highly and multifunctional next-generation materials throughcollaborative research.
Recently, the density functional theory (DFT)
simulation has been found to be an effective
approach for explaining the mechanism of adsorption
on various metal and metal oxide catalyst surfaces.
An iron catalyst is commonly used in the FischerTropsch synthesis for fuel production from a gaseous
mixture of CO and H2. The feed gas, however, also
contains H2S, which can shorten the active life of
the Fe catalyst. We have investigated the influenceof Na and K on the properties of the CO and H2S
molecules adsorbed on a Fe(100) surface [1]. It has
been found that the presence of Na and K atoms on
the Fe(100) surface promotes the adsorption of CO
molecules and hinders the adsorption of H2S
molecules, and, in particular, the potassium additiveFig. 1 (a) Adsorption energies of CO molecule on Fe(001)
and alkali/Fe(100) surfaces; (b) Charge density of
CO/K/Fe(001) at preferred site; (c) Selected surface
model and most stable configuration of CO2 on
CaO(001) surface; (d) Density of states and side view ofexcess (red) and depletion (blue) charges of CO2 on
Ca1-xBexO (001) surface (at x=0.02).

59

IMR KINKEN Research Highlights 2013

has been found to induce the largest variations inadsorption energies (Fig. 1(a)). Through examinationsof the local density of states and charge analysis,
we have shown that electrostatic interactions
between the alkali atoms and the molecules aremainly responsible for this enhancement effect (Fig.
1(b)). With the aim of developing a novel absorbent
for effective CO2 sequestration, DFT calculations
have been employed for studying the adsorption
process of CO2 on a CaO (001) surface (Fig. 1(c)).
The effect of adding other alkaline earth metaloxides to CaO on the adsorption energy of CO2 hasbeen analyzed (Fig. 1(d)). The structural deformationof the CaO surface due to the formation of acompound or solid solution would result in variationsin the adsorption energy, and in the case of CaO/
BeO, carbon dioxide capture during adsorption is
improved significantly [2].
The handling of glass-forming molten alloys is avery important issue in both basic and appliedresearch on noncrystalline materials. In our center,
we proposed the following press release: (i)“Development of Compact Automatic Arc-Furnace”[3] and (ii) “Development of New Atomization
Process using High-Velocity Air Fuel Flame” [4].
First, good reproducibility of cast amorphous alloy
would be achieved using the automatic fabrication
system with the optimized alloying process. The
developed compact automatic arc furnace is mainlycomposed of three components: weighing, alloying,
and casting. Because an amorphous material has afeatureless structure without a second phase, grainboundaries, and so on, the prealloying process is
an important for achieving sufficient master alloy
homogeneity by avoiding the formation of
unexpected
high-melting-pointintermetallic
compounds. Full use of the automatic arc-furnace
could lead to meaningful progress in the advanced
quality control of cast amorphous alloys. Second, toproduce fine amorphous alloy powder, weinvestigated high-temperature atomization using ahigh-velocity air fuel (kerosene) flame (HVAF) (seeFig. 2(a)) because the diameter of atomized powder
particles is related to the surface tension andviscosity of the molten alloy. At high temperatures,

IMR KINKEN Research Highlights 2013

Fig. 2 (a) View of HVAF atomization and (b) SEM image ofatomized Fe-based amorphous powder.

the molten alloy possesses both the attributes. In
contrast, conventionally, water or gas is used as theworking fluid for powdering in atomization. Because
the temperature of the HVAF is approximately 1873
K, a suitable cooling system should be developed
for obtaining dry, quenched atomized powder. As
shown in Fig. 2(b), we successfully obtained a fineamorphous powder of Fe73.2Cr2.2Si11.1B10.8C2.7
alloy.
Rare-earth aluminoborides REAlB4 with layered
structures have attracted considerable attention
from researchers. Fig. 3 shows a lattice image andan electron diffraction pattern of an α-TmAlB4 single
crystal synthesized under very rapid cooling (> 200K/h) [5]. All diffraction spots can be indexed as
belonging to the α-type phase without additionalspots and diffuse streaks, which suggest no
structural modulations. The magnetic susceptibility
χ of α-TmAlB4 along the c-axis under different fields
is shown in Fig. 4(a). Peaks of χ can be observed
underlower-magnitude
fields,indicating
antiferromagnetic (AF) transitions. A sharp λ-likepeak of specific heat, C, is observed at around 5.5
K in Fig. 4(b). The behavior of this peak is different
compared with the C of slow-cooled (5–50 K/h)
α-TmAlB4 containing building defects. The present
result shows how the large peak structure in thespecific heat curve of the high quality sample canbe destroyed by the presence of a small amount ofbuilding defects.
The center will continue to promote thedevelopment of advanced materials based on
fundamental and industrial research using cuttingedge technology.

Fig. 3 Lattice image and electron diffraction pattern of
α-TmAlB4 crystal taken along c-axis.

Fig. 4 Temperature dependence of (a) magnetic
susceptibility χ and (b) specific heat, C, of α-TmAlB4.

References[1] M. H. Mahyuddin, R. V. Belosludov, M. Khazaei, H. Mizuseki,
Y. Kawazoe, J. Chem. Phys. C, 115, 23893 (2011). (with Cover
page)[2] K. Sasaki, K. Wakuta, N. Tokuda, R. V. Belosludov, S. Ueda, Y.
Kawazoe, T. Ariyama, ISIJ International, 52, 1233 (2012).
[3] http://www.tohoku.ac.jp/japanese/2012/11/press20121105-01.
html[4] http://www.tohoku.ac.jp/english/2013/01/press20130125-01.
html[5] T. Mori, I. Kuzmych-Ianchuk, K. Yubuta, T. Shishido, S. Okada,K. Kudou, Y. Grin, J. Appl. Phys., 111, 07E127 (2012).

Akihiro Makino (Cooperative Research and Development Center for Advanced Materials)
E-mail: amakino@imr.tohoku.ac.jp
URL: http://www.crdam.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201360

Research Centers

High-strength Nb3Sn Rutherford Cable for
47 T Hybrid Magnet
High Field Laboratory for Superconducting MaterialsOur group, the High Field Laboratory for Superconducting Materials, is planning to develop a 25 T cryogenfree superconducting magnet, a 47 T hybrid magnet, and a 30 T superconducting magnet. We have been
investigating the superconducting wire and cables needed for development of these magnets.
The High Field Laboratory for Superconducting
Materials (HFLSM) has concentrated on the key
research fields of “High field magnet technology,”“Materials science in high fields,” and “Basic
research under multiple extreme conditions” ascooperative research in the Institute for Materials
Research, Tohoku University. The HFLSM has
demonstrated some high field superconducting
magnets such as a 30 T hybrid magnet, a 20 T
cryogen-free superconducting magnet [1], and a 28
T cryogen-free hybrid magnet [2].
As the next project in the construction of highfield facilities in Japan, our group is planning todevelop a wide-bore 20 T superconducting magnet
with a 400 mm room-temperature bore for a 47 T
hybrid magnet [3, 4], a 25 T cryogen-freesuperconducting
magnet,anda30
Tsuperconducting magnet. Three types ofsuperconducting wires [NbTi, Nb3Sn, andREBa2Cu3O7 (RE : Y or rare-earth elemtnt)] will be
used for the three large magnets, and we have beenplanning to use Nb3Sn Rutherford cables.
We carried out a performance test of a singlelayer test coil with five turns using a CuNb-reinforced
Nb3Sn (CuNb/Nb3Sn) Rutherford flat cable (Fig. 1).
Fig. 2 shows the transport current properties of
CuNb/Nb3Sn Rutherford coil. The wide-bore coilwith a 270 mm outer diameter was examined intransport currents up to 1500 A at 8 T and 4.2 K.
Strain gauges attached to stainless steel tape thatwas co-wound around the Rutherford flat cable forreinforcement exhibited a strain of 0.25%. These
results suggest that we can use the CuNb/Nb3Sn
Rutherford cable under such a large electromagnetic
force.
We also developed a new fabrication process
for the CuNb-reinforcing stabilizer for Nb3Sn wires.
One of the serious problems for Nb3Sn wires is that
their superconducting properties deteriorate
because of the applied strain [5] when a large
electromagnetic force is applied in a high field
superconducting magnet. In-situ CuNb-reinforced
Nb3Sn superconducting wires have been developed
to enable the wire to withstand such a large
electromagnetic force [6].
Because a large-scale superconducting
61

IMR KINKEN Research Highlights 2013

Fig. 1 Photograph of the CuNb-reinforced Nb3Sn
Rutherford coil. The single-layer test coil composed ofa five-turn Rutherford flat cable was made with sixteen
0.8 mm diameter CuNb/Nb3Sn strands. Stainless steel
SUS316L tape was co-wound in order to provide
reinforcement against a large electromagnetic force.

Fig. 2 Transport current properties of CuNb/Nb3Sn
Rutherford coil at 8 T and 4.2 K. The Rutherford coil was
examined in transport currents up to 1500 A.

magnet has a significant amount of stored magnetic
energy, we developed a CuNb-reinforced stabilizer
with high electrical conductivity to ensure thecryogenic stability of the large-bore superconducting
magnet. Unfortunately, in-situ CuNb has a relatively
low electrical conductivity because of the very low

IMR KINKEN Research Highlights 2013

solubility of Nb in pure Cu.
In order to improve the solubility of Nb in pureCu and thus increase the electrical conductivity, wedeveloped a new Nb-rod process in order tofabricate the CuNb-reinforcing stabilizer [7]. Fig. 3
shows the cross section of the Nb-rod-processed
CuNb/Nb3Sn wire. The wire is 0.8 mm in diameter.
The superconducting property of the CuNb/
Nb3Sn wire under tensile stress at 4.2 K and 14.5 T
is shown in Fig. 4. The Nb-rod-processed CuNb/
Nb3Sn wire had a maximum critical current Ic at atensile stress state of 230 MPa, whereas the peak of
the Ic stress curve for the in-situ-processed wire
appeared at 210 MPa. Therefore, the tensile stressdependence of Ic was improved using the Nb-rodprocessed CuNb instead of the in-situ-processed
CuNb as a reinforcement material.
References

Fig. 3 Cross-sectional view of the 0.8-mm-diameter Nbrod-processed CuNb/Nb3Sn superconducting wire.
The Cu-20vol%Nb-reinforcing stabilizer was placed
between copper outside layer and the Nb3Sn
superconducting layer.

[1] S. Hanai, T. Tsuchihashi, Y. Minemoto, S. Ioka, K. Watanabe,
S. Awaji, H. Oguro, Abstract of CSSJ Conference, 87, 148(2013). (in Japanese)[2] K. Watanabe, G. Nishijima, S. Awaji, K. Takahashi, K. Koyama,
N. Kobayashi, M. Ishizuka, T. Itou, T. Tsurudome, J. Sakuraba,
IEEE Trans. Appl. Supercond., 16, 934 (2006).
[3] K. Watanabe, S. Awaji, H. Oguro, K. Takahashi, K. Minegishi,
T. Suwa, Y. Sasaki, T. Kiyoshi, T. Asano, S. Hanai, H. Tsubouchi,
I. Inoue, IEEE Trans. Appl. Supercond., 22, 4300804 (2012).
[4] K. Watanabe, S. Awaji, Y. Hou, H. Oguro, T. Kiyoshi, H.
Kumakura, S. Hanai, H. Tsubouchi, M. Sugimoto, I. Inoue, IEEE
Trans. Appl. Supercond., 23, 4300304, (2013).
[5] J. W. Ekin, Cryogenics, 20, 611 (1980).
[6] H. Sakamoto, S. Endoh, Y. Nagasu, K. Wada, A. Kimura, S.
Meguro, S. Awaji, K. Watanabe, M. Motokawa, IEEE Trans.
Appl. Supercond., 12, 1067 (2002).
[7] H. Oguro, S. Awaji, K. Watanabe, M. Sugimoto, H. Tsubouchi,
J. Supercond. Nov. Magn., 26, 2099 (2012).

Fig. 4 Applied tensile strain dependence of the criticalcurrent density for the Nb-rod-processed CuNb/Nb3Sn
wire and the in-situ-processed CuNb/Nb3Sn wire at 4.2
K and 14.5 T. The Nb-rod-processed CuNb/Nb3Sn has
the maximum Ic at a tensile stress state of 230 MPa,
whereas the peak of the Ic stress curve for the in-situprocessed wire appears at 210 MPa.

Hidetoshi Oguro (High Field Laboratory for Superconducting Materials)
E-mail: h-oguro@imr.tohoku.ac.jp
URL: http://www.hflsm.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201362

Research Centers

Scientific Breakthrough toward Creation of New Industrial
Materials
Kansai Center for Industrial Materials ResearchIndustrial Materials Research (IMR) has launched the Kansai Center, a comprehensive Institute-wide effort thatpairs IMR’s research groups across the innovation spectrum to solve today’s industrial challenge and transform
tomorrow’s global energy problems. The center was established in April 2011 based on an agreement between
IMR, Osaka Prefecture Government, and enterprises in the Kansai area sponsored by government, taking over
Osaka Center project.
Kansai Center was established in Osaka as a
special unit within the Institute for Materials Research(IMR), Tohoku University in April 2011 based on anagreement with Osaka Prefecture Government. The
Center is sponsored by the government (Ministry of
Education, Culture, Sports, Science and Technology)
and took over the Osaka Center, who carried out its
mission from 2006 to 2011. The mission of the Osaka
Center was to support small and medium enterprisesin the Osaka area. The area covered by the Kansai
Center expanded from the Osaka area to the whole
Kansai area. The Center has three missions. First, tosolve technical problems that industries havestruggled to resolve. Solving such problems is
occasionally accomplished through collaboration
across universities, research institutions, fields, andinstitutional boundaries. Some of these activities
have developed in collaborative partnerships withindustry. Second, the Center introduces academic
output to industries, with the aim of applying suchoutput to society. Currently, about 50 enterprises
have contracted to engage in joint research on
projects of mutual interest, and the Center has
vibrant patenting/licensing activities with 23 newinvention disclosures. Third, the Center helps
educate next-generation materials scientists andresearchers in universities and enterprises. To thisend, we have organized a bimonthly forum named“Monodukuri Kisokoza,” which focuses on special
topics regarding materials and processing.
Projects are conducted through wide
collaborations between the government, universities,
research institutions, and other organizations suchas Osaka Bay-area Consortium, Kansai Bureau
Economy Trade and Industry, Research Society for
Global Materials Crisis, and General Production
Company in order to achieve results efficiently.
Among them, the Development Center for New
Metallic Material in Osaka Prefecture University andthe Research Center for Nano-Micro Structure
Science and Engineering at the University of Hyogo
are close partners in collaborative projects, andCenter staff participates in educating students at
both universities. We have seven groups at theCenter: Environmental Protection and Energy
63

IMR KINKEN Research Highlights 2013

Conversion Materials, Functional Materials for Future
Industries, Analytical Science and Technology,
Structural Materials for a Low-Carbon Society,
Nanostructure Control for Engineering Materials,
Advanced Biomaterials, and Structure Design ofGreen Metallic Materials. The Center has four venues
to cover the Kansai area: the Osaka office at Osaka
Prefecture University, the Hyogo office at Universityof Hyogo, the Sendai office in IMR, and Monodzukuri
Business Information Center Osaka (MOBIO) in theCreation Core Higashi-Osaka, a governmental bodywhere sixteen local universities and one college arelocated to facilitate collaboration with various
industries.
Technical inquiries from industries are handledsolely by seven professors in Kansai Center, whichhelps save investigation and solving time. Fig. 1
shows the number of inquiries from industries toCreation Core Higashi-Osaka. It is apparent that
Kansai Center occupies the largest portion among
the universities and the number of inquiries increases
each year. According to the statistics, about 30% of
all inquiries are from enterprises located in areas
other than Kansai.
Two special “Monodukuri Kisokoza” series havebeen held so far (Fig. 2). The first was aimed towardlearning more about metals in terms of plasticity,
corrosion, and characteristic functions, in addition

Fig. 1 Fraction of inquiries (2009–2012) from industries atCreation Core Higashi-Osaka, where sixteen local
universities and one college are located. The area in
red corresponds to Kansai Center.

IMR KINKEN Research Highlights 2013

Fig. 2 Open seminar “Monodukuri Kisokoza” provides upto-date information on materials and process
technologies to engineers and researchers in industry.

Fig. 4 Trial products resulting from collaboration withindustries (displayed in Business Encourage Fair 2012
in Osaka).

Fig. 3 Thread-rolled metallic glass bolt (left) and the SEM
micrograph of the screw thread of the bolt (right).

Fig. 5 Newly selected Kansai Center logo advocating thethree missions of the Center.

to providing background such as the history andproduction of metals. Each seminar is composed ofbasic information and its application, lectures
presented by academia and engineers, respectively.
So far, six seminars focusing on Ti, Cu, Ni, Al,
stainless, and Mg have been held. The second
series is organized by the young academic staff at
the Center in order to foster collaborative researchand development with industries. Each topic is
determined by considering the industrial needs andtheir own specialty. Further, Center professors offer
lectures on materials design, technical skills, andrelated fundamentals of materials science to thestudents in Osaka Prefecture University and Hyogo
Prefecture University for the purpose of demonstrating
the importance of material science contributing toindustrial applications.
The Center has developed various materials andprocesses under collaboration with industry. One ofthem is metallic glass bolts fabricated by squeeze
casting and cold-thread rolling (Fig. 3). These bolts
have 1.3 times the strength of high-strength steel andare 2 times more difficult to loosen than steel. It is
expected that these bolds can be used in biometallics
or earthquake-resistant parts in structures [1].
In another successful collaboration, bonding
technology using spark plasma sintering (SPS) wasapplied to fabricate a bimetallic composite of steel
and a Cu alloy for mechanical parts in an axial piston
pump. Interdiffusion promotion by SPS strengthens
the bonding cohesion through transformation in thesteel [2]. Further, high-performance analysis
equipment in IMR contributes to new materials
fabrication. Nano-sized metallic powders expected
to be used for electric devices such as electrodes
and metal lines tend to aggregate, a characteristicthat needs to be suppressed. A systematic study byusing TEM contributed to a new method of fabricating
nano-sized powders. The trial products resulting
from collaborative projects are displayed in various
exhibitions or MOBIO in order to appeal the visitors.
(Fig. 4).
The Center recently selected a logo (Fig. 5).
The Center’s three missions are embodied by threeelements in the logo, and the green leaf in the top
right corner symbolizes green innovation. This logo
was selected from a design yearbook competitionof Japanese typography, yearbook 2013.
References[1] S. Yamanaka, K. Amiya, Y. Saotome, A. Inoue, Mater. Trans.
52, 243 (2011).
[2] N. Masahashi, S. Semboshi, K. Watanabe, Y. Higuchi, H.
Yamagata and Y. Ishizaki, J. Mat. Sci., 48, 5801 (2013).

Naoya Masahashi (Kansai Center for Industrial Materials Research)
E-mail:masahasi@imr.tohoku.ac.jp
URL: http://www.kansaicenter.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201364

Research Centers

First-principles Study on Mechanical Properties of Fe–Si Alloy
Center for Computational Materials Science

Most of the work performed in Computational Materials Science (CMS) has been centered around the subjectof functional properties that can be analyzed directly by the electronic structure of the material. While various
mechanical properties of alloys, such as strength and ductility, are basically multiscale phenomena, and few
attempts have been made to simulate and analyze the deformation behavior of an alloy from first-principles.
Introduced below is a challenging attempt in this direction performed by Prof. Chen’s group, in which they
predict the ductile–brittle transition of a Fe–Si alloy. Computer simulation of the mechanical properties isbelieved to become a key area of CMS performed by supercomputers.

Silicon steel is one of the most important
industry materials used in transformers, motors, andmagnetic coils because of its excellent mechanicaland magnetic properties. A higher Si content in steel
improves effectively the magnetic properties suchas smaller hysteresis, higher permeability, and alower magnetostriction constant, whereas beyond acertain Si concentration, the mechanical properties
degrade rapidly. In order to investigate the origin of
the effect of Si concentration on the mechanicalproperties of high-silicon steel, we employed firstprinciples calculations for dilute Si in a bcc-Fe solid
solution. Carbon and other elements present in thesilicon steel were omitted from the calculations inorder to reduce the complexity of the calculations.
Electronic structures were calculated using density
functional theory with Projector augmented
wave(PAW)-Generalized gradient approximation(GGA) potential [1-4] on supercell models thatsimulate various Si concentrations.
First, we investigated the ground-state structuresof bcc-based Fe–Si in order to get overall informationon the stability of the system. The cluster expansion
method [5, 6] has been carried out to predict the 0
K ground state structures in the bcc Fe–Si system.
Fig 1 shows the enthalpy of formation versus Si
concentration for the various predicted structuresby using 41 structures and 34 types of clusters,
which led to a prediction of five ground-state
structures throughout the entire range of Siconcentrations. An Fe-rich D03 -ordered bcc
structure, which is one of the two bcc-based
structures present on the Fe-rich side of the Fe–Si
system at finite temperatures, was correctly predicted
to be stable at 0 K. We also compared the enthalpy
of formation of stable binary compounds in the Fe–
Si system (Fe2Si, Fe5Si3, α-FeSi2, β-FeSi2) to confirm
the validity of the calculations.
Further work focused on the Si-concentration
dependence of various mechanical properties inthe dilute Si region of the Fe–Si system [7]. The
elastic constants of the dilute Si in bcc Fe structures
were calculated using the stress-strain method [8]based on electronic-states calculations on the fully65

IMR KINKEN Research Highlights 2013

relaxed Si-doped structures. Fig 2(a) shows thecalculated elastic properties, including c11, c12,
c44, bulk modulus (B), shear modulus (G), andYoung’s modulus (E) at different Si concentrations.
A steady change in these properties with increasing
Si up to 8.3 at% is shown. These properties
decreased clearly insert at 10.9 at% and thenFig. 1 Ground-states diagram of Fe–Si bcc- based
structures: Formation enthalpy vs. Si concentration(at%).

(a)(b)

Fig. 2 Calculated elastic properties of dilute Si in bcc-Fe.
(a) Elastic constants and modulus. (b) Ratio of bulk
modulus and shear modulus as empirical criteria.
between ductility and brittleness of alloy.

IMR KINKEN Research Highlights 2013

rebounded to a certain extent as the Si concentration
increased further. This trend is well consistent withexperimental results [9]. Fig 1(b) further plots theratio of the bulk modulus B and the shear modulus
G with respect to various Si concentrations.
According to Puge’s empirical rule [10], a B/G ratioof 1.75 can be taken as the criterion between ductile
and brittle for metal. Thus, the decrease in B/G value
from 2.0 to 1.5 at 8.3–10.9 at% Si clearly shows aductile-to-brittle transition. The origin of the Siconcentration dependence of the elastic properties
is discussed as following in terms of the evolution ofelectronic structures.
Fig 3 shows the total electronic density-ofstates and the partial electronic density-of-states of
Si atoms for pure Fe and 5 Si-doped systems. For all
Si concentrations from 0 to 8.3 at% Si, the density ofstates (DOS) distributions were found to sharesimilar characteristics, whereas noticeably different
distributions between 8.3 and 10.9 at% Si were
found. It can be seen that for minority-spin at 8.3
at% Si, the Fermi level is at the dip of the two peaks,
whereas at 10.9 at% Si, the Fermi level falls just at
the peak. This increase in minority-spin led to an
obvious increase in the total DOS and resulted in asignificant drop of magnetic moment as comparedto lower Si concentrations. This indicates that thecharacteristics of bonding between Fe atoms
surrounding Si atoms vary, which is thought to be
the direct reason for the change in the mechanicalproperties with increasing Si concentration.
Further analysis on the redistribution of theelectronic structures due to the presence of Si atoms
revealed a relationship between magnetism and theelastic properties of the system. Combined with thefinite temperature behavior, we reached theconclusion that the interplay between magnetismand the D03 ordering effect resulted in the ductile–
brittle transition for Si concentrations of about 8.3
at%. This example shows a modern computational
materials science approach based on large-scale
electronic structure simulations for investigating,
explaining, and predicting the fundamental
properties of real materials.

Fig. 3 (a) Total electronic density of states of Fe-bcc, Sidoped bcc Fe at 2.7, 5.6, 8.3, 10.9, and 12.5 at% Si. (b)Si partial electronic density of states in Si-diamond, Sidoped bcc-Fe at 2.7, 5.6, 8.3, 10.9, and 12.5 at% Si.

[4] P. Perdew, K. Burke, M. Ernzerhof, Phys. Rev. Lett. 77, 3865
(1996).
[5] J. M. Sanchez, F. Ducastelle, and D. Gratias, Physica, 128A,
334 (1984).
[6] W. Connolly and A. R. Williams, Phys. Rev. B, 27, 5169
(1983).
[7] A. Saengdeejing, Y. Chen, K. Suzuki, H. Miura, and T. Mohri,Comp. Mater. Sci., 70, 100-106 (2013).
[8] Y. Le Page and P. Saxe, Phys. Rev. B, 65, 104104 (2002).
[9] A. Machova and S. Kadeckova, Czech. J. Phys. 27, (5), 555-

References[1] G. Kresse, VASP, Vienna Ab initio Simulation Package, 2004.

567 (1977).
[10] S. Pugh, Philos. Mag. Ser. 7, 45, (367), 823–843 (1954).

<http://www.vasp.at>.
[2] G. Kresse, D. Joubert, Phys. Rev. B, 59, 1758 (1999).
[3] P.E. Blochl, Phys. Rev. B, 50, 17953 (1994).

Tetsuo Mohri (Center for Computational Materials Science)
E-mail: tmohri@imr.tohoku.ac.jp
URL: http://www-lab.imr.edu/~ccms/Eng/index.php

IMR KINKEN Research Highlights 201366

Research Centers

Crystal Growth and Spectroscopic Properties of Laser
Single Crystals/Ceramics
International Collaboration Center (ICC-IMR)

I

have continued fruitful cooperation with Prof. Yoshikawa’s and Prof. Goto’s labs between November 13 andDecember 16, 2012, mainly on (Nd3+/Yb3+ )-doped transparent sesquioxide and spinel ceramics or crystals forlaser application. I was also invited to participate in the Materials Science Week 2012, November 28–30,
2012.
It has been a great privilege to cooperate with
IMR Laboratories since a long time. I have proposed
applying our fruitful relationships with Laboratoirede Physico-Chimie des Matériaux Luminescents(LPCML, nowadays ILM at the Université Claude
Bernard Lyon 1) and Institute for Materials Research(IMR, Tohoku University) with the natural continuationof the laser materials research program on rare earth(Nd3+/Yb3+)-doped transparent sesquioxide andspinel ceramics/crystals as laser materials with highmelting points. I want to express my thanks to Prof.
Niinomi, head of IMR, Prof. Nojiri, head of ICC-IMR,
Prof. Konno, TEM Lab, Prof. Goto, SPS Ceramics
Lab, and Prof. Yoshikawa, Advanced Crystal
Engineering Lab.
The project ELyT lab: M12–LASMAT was started
in March, 2012, and included Yoshikawa’s andGoto’s teams. Lu2O3 was selected from amongsesquioxides for the following reasons: it is aninactive lattice in optics such as Sc2O3 and Y2O3,
Lu3+ can easily be substituted by other rare earth
cations as dopants, it has the highest thermal
conductivity (12.5 W/m/K) and a lower phonon
energy (391 cm-1) as compared with YAG (700 cm1). We first focused our attention on Lu2O3 ceramics,
and then studied Lu2O3 crystals grown by µ-PD
using a special crucible. Indeed, Lu2O3 has a highmelting point up to 2490 °C, which makes it extremely
difficult to grow by conventional crystal growthmethods. Such difficulties encourage researchers
seek other ways to obtain Lu2O3 as a ceramic lasermaterial.
Research activity on advanced laser materials
is indeed increasing with the availability of transparent
sintered polycrystalline ceramics, which present
some advantages as compared to single crystals interms of size, mechanical strength, overall production
cost, and cost-effectiveness [1]. After the pioneering
work of Dr. A. Ikesue [2] at World Lab. Co. Ltd. and
Dr. H. Yagi at Konoshima Company [3] on ceramics
fabricated by conventional methods, original andpromising research is in progress with the synthesisof Nd3+:Lu2O3 ceramics by a nonconventional spark
plasma sintering (SPS) technique developed by
Prof. T. Goto’s team [4]. The spectroscopic properties
67

IMR KINKEN Research Highlights 2013

were analyzed at LPCML (UCBLyon1) and will becompared with the rare earth-doped crystals grownby Prof. A. Yoshilawa’s team. Regarding the spatial
distribution of rare earth dopants, we have alreadypublished results with A. Yoshikawa’s team on therare earth segregation phenomenon in YAG laser
ceramics, especially from TEM analysis associatedwith an EDX probe in cooperation with Dr. T. Epicier
at the Institut National des Sciences Appliquées deLyon lab [5-8]. We program the same evaluation
with Lu2O3 ceramics.
Another area of interest of Nd3+:Lu2O3 single
crystals stems from a 2011 report on the laser output
of such crystals [9], which points out that specificdual Nd3+ wavelengths lasing at 1076.4 and 1080.5
nm (see Fig.1) are very promising for applications incoherent terahertz (THz) generation by difference
frequency and ultrahigh repetition rate pulse by
optical beating.
The fabrication, measurement of thermal
conductivity, and structural characterization by TEM
will be done at IMR, whereas spectroscopic andlaser properties will be analyzed and optimized at
ILM. Our first analysis of the absorption spectra,
emission spectra, and lifetimes of Nd3+ 4F3/2 level
under selective wavelengths made at LPCML (Lyon)
have been submitted and discussed with themembers of this cooperative project.
I also had the opportunity to attend ICC-IMRELyT workshop on materials science on November
27, and I was invited to participate in the Materials
Science Week 2012 (MSW2012), Summit of Materials
Science, November 28–30, 2012. The intenseinterest in Materials Science Week is to have accessto the large program of IMR on materials. My Invited
talk was given for the introductory session (I2) andwas titled ”Laser materials: From single crystals topolycrystalline ceramics and distribution of rareearth dopants,” which is directly connected to my
IMR cooperation [5-8]. The goal of this lecture wasto present the evolution of laser materials fromdoped single crystals to doped ceramics. Anotherimportant cooperation between IMR (Prof. T. Konno,
Prof. K. Sato), INSA-Lyon (Dr. T. Epicier), and us in
UCBLyon1 involved detection of Yb dopants at theIMR KINKEN Research Highlights 2013

Fig. 1 IR Emission spectra of a 1%Nd-doped Lu2O3 ceramic showing the two 1076.4 and 1080.5 nm laser lines of the main
C2 crystallographic site.

atomic level in a YAG optical ceramic by scanningtransmission
electronmicroscope-high-angle
annular dark-field imaging (STEM-HAADF), which
was presented by T. Epicier (P 94).

Q. Feng, Y. B. Pan, V. Chani, A. Yoshikawa, Opt. Mater., 33,
684-687 (2011).
[7] G. Boulon, T. Epicier, W. Zhao, V. I. Chani, T. Yanagida, A.
Yoshikawa, Jpn. J. Appl. Phys., 50, 9, Article ID: 090207,
published on 2011/09/20 (online).

References[8] V. I. Chani, G. Boulon, W. Zhao, T. Yanagida, A. Yoshikawa,
Jpn. J. Appl. Phys., 49, 075601 (2010).

[1] G. Boulon, Opt. Mater., 34(3), 499-512 (2012).
[2] A. Ikesue, I. Furusato, K. J. Kamata, Am. Ceram. Soc., 78,[9] L. Hao and al. Opt. Express, 19, No. 18, 17774-17779 (2011).

225–228 (1995).
[3] J. Lu, M. Prabhu, J. Song, C. Li, J. Xu, K. Ueda, H. Yagi, T.
Yanagitani, A. A. Kaminskii, Jpn. J. Appl. Phys. 40, L552
(2001).
[4] L. An, A. Ito, T. Goto, J. Am. Ceram. Soc., 94, 695-698 and
31(9), 1597-1602 (2011).
[5] W. Zhao, C. Mancini, D. Amans, G. Boulon, T. Epicier, Y. Min,H. Yagi, T. Yanagitani, T. Yanagida, A. Yoshikawa, Jpn. J. App.
Phys., 49, 022602 (2010).
[6] W. Zhao, S. Anghel, D. Amans, G. Boulon, T. Epicier, Y. Shi, X.

Georges Boulon (Institute Light Matter (ILM), UCBLyon1-CNRS)
E-mail: georges.boulon@univ-lyon1.fr
URL: h ttp://georges-boulon.univ-lyon1.fr
http://ilm.univ-lyon1.fr
Hiroyuki Nojiri (International Collaboration Center)
Email: icc-imr@imr.tohoku.ac.jp
URL: http://www.icc-imr.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201368

Research Centers

Promoting Materials Research for Low-Carbon Society
Integrated Materials Research Center for a Low-Carbon Society (LC-IMR)

The Integrated Materials Research Center for a Low-Carbon Society (LC-IMR) was founded to promote
integrated research in materials science to contribute to achieving a low-carbon society. The main aim of LCIMR is to promote outstanding projects in the field of materials research, which will make low-carbon emissions
a reality through various perspectives, from energy saving to energy harvesting and storage. We are also
making active collaborations with research groups outside our institution. Cooperation between industry, thegeneral public, and the university will be sought through workshops and by actively making information
available.
It is important for the field of materials science
to contribute to the achievement of low carbon
emissions and a sustainable society. The Integrated
Materials Research Center for a Low-Carbon Society(LC-IMR) was founded in April of 2010 to promote
integrated research in materials science towards theachievement of a low-carbon society. Individual
members of LC-IMR belong to various materials
research fields such as structural materials,
semiconductors,
energystoragematerials,
electronics, and spintronics. We believe that
integration of research in different fields is very
important for making a breakthrough and thus weseek collaborations both among our members andwith outside centers/institutes/universities.
In the past three years, we have been working
on the following three subjects:
1. Conducting strategic research by individuals.
2. Supporting new research projects throughmanagement of internal research funding.
3. Having discussions with industry andgovernment for a better understanding of theneeds of our society.
Japanese funding agencies are finding thatvarious strategic research programs in materials
science are necessary for our future society. We arebeing well funded and are making good progress inthose research programs. In 2012FY, we conducted

Fig. 1 Outlook of LC-IMR

69

IMR KINKEN Research Highlights 2013

IMR KINKEN Research Highlights 2013

Fig. 2 One scene from the 3rd LC-IMR workshop

four projects in energy-saving technology, seven in
renewable/clean energy, and three in element
strategy.
LC-IMR is also managing a program on internal
research funding in IMR to support both promisingand challenging research projects, with particularemphasis on the center's mission. In 2012FY,
projects included clathrate hydrates, silicon solar
cells,
nitridesemiconductors,
germanium
superlattices, and spin Seebeck effects. In theselection of projects, contributions to low-carbon
emissions and a sustainable society, as well as
research integration among different research fields,
are being emphasized.
To publicize the work of our center, we organize
an annual workshop and introduce our research
activities. We also invite key persons from outside
the university such as from industry or government
to discuss how to develop suitable systems for asustainable society and the expected roles ofmaterials science. Such collaborations are anotherimportant part of our activities.
After three years, it is time for LC-IMR to focus

on activities based on the center’s mission. We willcontinue to make efforts to progress in our missions
and ask collaboration with individual researchersand organizations all around the world.

Tadashi Furuhara (Integrated Materials Research Center for a Low-Carbon Society)E-mail: lc-imr@imr.tohoku.ac.jp
URL: http://www.low-carbon.imr.tohoku.ac.jp/eng/index.html

IMR KINKEN Research Highlights 201370

Research Centers

Neutrons for Advance Material Science
Center of Neutron Science for Advanced Materials

The Center of Neutron Science for Advanced Materials (NC-IMR) is a unique neutron facility that has the

background of advanced material science in IMR. NC-IMR operates two neutron spectrometers for materialscience in a reactor facility. It is also the core of a construction project for a new neutron spectrometer in the Japan
Proton Accelerator Research Complex (J-PARC/MLF), which is the world brightest neutron source. Moreover,based on crystal growth technology developed by IMR, NC-IMR is developing novel neutron devices.

New Project at J-PARC
The Center of Neutron Science for Advanced
Materials (NC-IMR) is constructing a new neutron
spectrometer at the world brightest neutron facilityof Japan Proton Accelerator Research Complex(J-PARC) by collaboration between Tohoku
University, High Energy Accelerator Research
Organization (KEK), and other universities/institutes.
Funded by the government, the construction will
start in 2013FY, and the first beam is expected by
the early 2015FY. NC-IMR is the core of this nationwide project.
This neutron spectrometer, named POLANO(POLarization Analysis Neutron spectrOmeter), ischaracterized by polarized neutrons, which are very
sensitive for observing pure magnetic correlations,and lattice-spin couplings in magnets. From anapplication viewpoint, POLANO will also be
important in the development of rare earth freepermanent magnet materials.
POLANO aims at observations of spin-latticecorrelations in wide-Q and ω space up to 30 meV in
the first phase and up to ~150 meV in the final phase.
Below 30 meV, low-lying magnetic scattering will beobserved, which is quite important for clarifying theroles of spin-lattice couplings in novel magnetic
materials such as multiferroic materials, in whichspin orderings are strongly coupled with dielectric
polarization and multipolar ordering behavior in rareearth compounds. In the final phase, metallic

magnets such as Fe, Cr, and Mn3Si will be important
targets, because at the final stage POLANO canobserve spin correlations beyond the 100 meV
region where characteristic magnetic excitations
coupled with band excitations exist. Understandingof the higher-energy magnetic excitations is a key toclarifying the origin of magnetism in metals, which
has been a secret since the discovery of iron inancient Greece. Because there exist just a few
polarized neutron spectrometers in the world,
POLANO will be a key instrument to generate
breakthroughs in novel materials science.
IMR spectrometers in a reactor facility
NC-IMR is operating a neutron spectrometerand neutron diffractometer in the JRR-3 reactor in
the Japan Atomic Energy Agency (Ibaraki, Japan).
The neutron powder diffractometer, HERMES, candetermine the structures of spins and nuclei inmaterials. In particular, HERMES is actively used forvisualization of ion distributions in ion conductors
and Li-ion batteries by the maximum entropy
method. In 2012, Yashima (Tokyo Institute ofTechnology) and his colleagues succeeded inclarifying the roles of doped Ga and Cu in enhancing
ion conductivity in a Pr-oxide based on data fromHERMES [2]. The triple-axis spectrometer AKANE
is an instrument to observe time and spatial
fluctuations of spins and nuclei. AKANE isindispensable for investigations of high-TC

Fig. 1 A rendering view of the polarization analysis neutron spectrometer POLANO in J-PARC.

71

IMR KINKEN Research Highlights 2013

IMR KINKEN Research Highlights 2013superconductors, geometrical spin frustration
systems, and so on. Note that observation of spincorrelations under high-pulsed magnetic fields,
which is one of the characteristic technologies in
IMR, has already been realized.
As a new neutron device, a neutron polarizer
system has been developed using the polarized
3 He gas filter method [1]. The polarizer will enhance
the sensitivity of the observations of spins inferromagnetic applications of HERMES andAKANE.
In 2012, NC-IMR completed construction of a
new compact diffractometer, HERMES-E, as a subinstrument of HERMES by competitive funds (Fig.
2), which is a quite rare neutron instrument focusedon educational experiments. Because opportunities
for neutron scattering experiments are quite limited,
HERMES-E will be an important instrument for
students and young researchers.
KINKEN-WAKATE2012
As an educational activity, NC-IMR organized
the 9th Materials Science School for Young Scientists
and Students (WAKATE2012) in November 2012.
Four of the world’s leading scientists gave lectures
introducing neutron and X-ray scattering and its

Fig. 3 KINKEN-WAKATE2012 organized by NC-IMR in
November-2012. Prof. Sunghun (U. Virginia USA) gave
a lecture on importance of dynamics investigations byneutrons.

applications in novel materials science to young
people in Tohoku University.
Developmentof
single-crystal
neutronmonochromators in IMR
A neutron monochromator is an importantcomponent for neutron spectroscopy. NC-IMR andCooperative Research and Development Center for
Advanced Materials are collaborating on the
manufactureofhigh-performance
crystalmonochromators, and they succeeded in producing
mosaic-controlled crystals of hot-pressed Ge andcold-pressed Cu for high-luminance and highenergy monochromators. In 2012, sets of the hotpressed Ge monochromator were installed on
HERMES (HERME-E) and AKANE. In particular, themonochromator for HERMES can provide two
separate neutron beams with different wave lengths(see Fig. 2). Moreover, we succeeded in developing
plastically deformed Ge single-crystal wafers. This
technique is quite ambitious for the generation ofhigh-luminance neutron beams a few millimeters insize [2]. These new crystal devices will be provided
to other facilities in Japan and abroad. Collaborations
with Korea Atomic Energy Research Institute are
now in progress.
References[1] K. Ohoyama et al. Nucl. Instr. Meth. Phys. Res. A, 680, 75(2012).
[2] Yashima et al. Chem. Mater. 24, 4100 (2012);[3] H. Hiraka et al., Nucl. Inst. Meth. Phys. Res. A 693, 166

Fig. 2 New diffractometer HERMES-E constructed in 2012
for educational purposes.

(2012).

Shin-ichi Orimo (Center of Neutron Science for Advanced Materials)
E-mail: imr_neutron@imr.tohoku.ac.jp
URL: http://neutron-center.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201372

Research Centers

High-Sensitivity Magnetometer Operating below 1 K
Laboratory of Low Temperature Materials ScienceAhigh-sensitivity magnetometer employing the Faraday balance was fabricated in the Laboratory of Low
Temperature Materials Science. In combination with a 3He cryostat, this handmade device enables us tomeasure the DC magnetization at low temperatures below 1 K. The apparatus is now becoming popular incollaborative works at Tohoku University because of its ease of operation, in addition to the commerciallyavailable magnetometers operating above 1 K which are also installed in this laboratory.
The Laboratory of Low Temperature Materials
Science (LLTMS) has a variety of cryogenic
equipment to study the basic properties of materials
as a function of temperature T and magnetic field H
and is open for collaborative works among
researchers at Tohoku University. During
collaborations, the most frequent experiments are
DC magnetization measurements of superconducting
and magnetic materials. If the desired experimental
conditions are not so extreme, for example attemperatures above 1 K (achievable with liquid 4He)
and magnetic fields below 7 T, one can easily
perform the experiments by using commerciallyavailable
magnetometers
such
as
the
superconducting quantum interference device(SQUID) magnetometers, which have been installedin LLTMS. However, if the required temperature forthe experiments is lower than 1 K, the situation isdifferent. Because the measurement principles of
the above mentioned magnetometers are based onelectromagnetic induction, the signal related to themagnetization is detected by moving the sample in
the pick-up coil. The heat generation because of thefriction accompanied with the sample motion
becomes a severe problem at low temperatures. In
addition, it is technologically difficult to move thesample without leakage in the 3He bath producing
the environment below 1 K or in the adiabatic space
thermally connected to the bath.
As a possible technique to avoid this problem,
we adopted a method using a Faraday balance, inwhich the magnetization M is measured without
moving the sample through the force F = MdH/dz
generated in the magnetic field, with z being the
distance from the center of the magnet coil. To
detect F, we made a compact capacitance load cell[1] which can be attached to most cryostats. Figure
1 shows the basic concept of the load cell where acapacitor is composed of an upper electrode,located at the bottom surface of the sample holder
suspended by the wire springs, and a fixed lower
electrode. When a sample with magnetization M
experiences a force F in a constant field gradientdH/dz, it causes a displacement of the sample ∆dwhich can be detected by a capacitance bridge.
Knowing the spring constant k of the wires, M can73

IMR KINKEN Research Highlights 2013

be directly derived by k∆d/(dH/dz).
Fig. 2 Shows the capacitance load cell attachedto the 3He cryostat, both of which are handmade bystaff at LLTMS. In combination with thesuperconducting magnet and the field gradient coil,which can generate a field gradient dH/dz of 12 T/m
in homogeneous background fields up to 8 T, we
can perform the magnetization measurements at T
= 0.3 - 1.4 K with a resolution of 10 -5 emu comparable to that of the SQUID magnetometer.

Fig. 1 Schematic for the principle of magnetization
measurement using a capacitance load cell.

Fig. 2 The Faraday-type magnetometer employing acapacitance load cell attached to the handmade 3He
cryostat.

IMR KINKEN Research Highlights 2013

Because the advantage of this magnetometer is thatthe load cell can be easily attached to other cryostats
such as a dilution refrigerator, we are planning toextend the measurement temperature range towards
the lower end.
Recently, this easy-to-operate magnetometer
employing the Faraday balance has been frequently
used in collaborative works to examine the phasetransition or crossover phenomena occurring below
1 K. Last year, a magnetic-field-induced transition
was successfully observed in the YbAl3C3
compound, which is known be a rare 4f-electron
system showing a spin-gap state [2]. The
investigation of magnetic phases around thequantum critical point in other Ce- and Ybcompounds is an ongoing subject.
As well as the Faraday balance magnetometer,
the SQUID magnetometers covering a higher
temperature range from 2 K to 400 K in magnetic
fields below 7 T are also popular apparatuses in
LLTMS because of their good accessibility. These
have contributed to the publication of more than 20
papers in one year alone. Recent topics include (i)Annealing effects on the magnetic properties of
Fe2CrAl Heusler alloy [3] and (ii) Logarithmic
behavior in the coercive force of ferromagnetic CoAl-O nano-granular films [4], as mentioned below.
Fe-based Heusler alloys are thought to be goodcandidates for ferromagnetic electrodes in spintronic
devices. In order to investigate the effects of thedegree of order on the magnetic properties of
Fe2CrAl, magnetic measurements were carried out
for the specimens obtained by annealing at several
temperatures. As shown in Fig. 3, it is clearly
confirmed that the magnetic properties such as the
spontaneous magnetization Ms and the Curie
temperature TC,indicate annealing effects. Together
with X-ray diffraction measurements, it is concluded
that the specimens with a lower degree of orderexhibit higher values of Ms and TC.
In Co-Al-O granular films, ferromagnetic nanoCo particles are dispersed in the insulating Al-O
matrix. These films show ferromagnetism andmetallic transport properties when the Coconcentration is higher than 60 atomic percent.
Recently, we have found that the coercive force of
these soft magnetic films shows a characteristic
behavior. As shown in Fig. 4, the coercive force of
these films increases logarithmically when lowering
the temperature below ~ 50 K, which is quite a new

Fig. 3 Magnetization curves measured at 5 K for Fe2CrAl
alloy annealed at 1573, 873 and 673 K. Inset are
thermomagnetization curves measured in a magnetic
field of 1 kOe.

Fig. 4 Temperature dependence of the coercive force ofthe ferromagnetic Co-Al-O and Fe-Al-O granular films
using a logT plot.

phenomenon.
References[1] T. T. Sakakibara, H. Mitamura, T. Tayama and H. Amitsuka,
Jpn. J. Appl. Phys. 33, 5067 (1994).
[2] K. Hara, S. Matsuda, E. Matsuoka, K. Tanigaki, A. Ochiai, S.
Nakamura, T. Nojima and K. Katoh, Phys. Rev. B 85, 144416(2012).
[3] R.Y. Umetsu, N. Morimoto, M. Nagasako, R. Kainuma and T.
Kanomata, J. Alloys Compds. 528, 34 (2012).
[4] S. Nakamura, T. Nojima, A. Yoshihara and S. Ohnuma, J.
Korean Phy. Soc., 63, 773 (2013).

Tsutomu Nojima (Laboratory of Low Temperature Materials Science)
E-mail: nojima@imr.tohoku.ac.jp
URL: http://www.ltsd.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201374

Research Centers

Research Facility for Studying Physical and Chemical
properties of Radioactive and Nuclear Materials
Laboratory of Alpha-Ray EmittersOur laboratory facilitates the use of more than 170 different radioisotopes and nuclear materials. This facility
accommodates physical, chemical, and radiochemical researches on actinides compounds, and materialscientific researches on the irradiation effect by using measurements system of positron lifetime and electron
microscopes. Successive introductions of several advanced systems, such as a physical properties
measurement system (PPMS; the lowest limit for measurements is 100 mK) and a tetra-arc-furnace for single
crystal growth, integrated the leading-hub laboratory in actinide physics and chemistry, under a close
cooperation with International Research Center for Nuclear Materials Science.
The present building, built in 1978 as a
successor to the former facility at IMR (located in
the basement of the old Build. 1 since 1960), isapproved to use more than 170 radionuclides andelements, a specific example being α-ray emitterssuch as actinide. The capability of this facility isdeserving of attention because of its scarcity inJapan; others are the JAEA (Japan Atomic Energy
Agency) and the International Research Center for
Nuclear Materials Science (Oarai), also at IMR.
Therefore, this laboratory occupies an important
position for the study of actinide elements not only inJapan, but also in international research fields.
This facility is composed of two areas: theradiation-controlled area and the non-controlled
area. The controlled area consists of three chemistrylab rooms, three physical lab rooms, a room forradiation measurements, a contamination test roomand storage rooms for radioisotopes and nuclear
materials. Each chemistry lab is equipped with local
exhaust ventilation systems which enable treatments
using various types of chemicals. In the radiation
measurement room, two γ-ray spectrometers (highpurity germanium semiconductor detector), a liquid
scintillation counter and an α-ray spectrometer(silicon detector) installed for radiation detection are
in operation.
For the preparation of alloys and compounds
such as UCoGe, a mono-arc furnace, a tetra-arc
furnace (Fig. 1), an electrical discharge cutting
machine and a Laue camera with an IP reader areused. In terms of chemistry, two inert-gas gloveboxes,
various types of electrochemical measurement
systems, a CHNS element analyzer, an ICP emission
spectrophotometer, a FT-IR spectrometer, a UVvisible spectrophotometer, a scanning electron
microscope (SEM), a powder X-ray diffractometer(XRD) and an X-ray single-crystal structure analysis
system are used. A physical property measurement
system (PPMS) installed is useful for various physicalexperiments. A transmission electron microscope(TEM) has been used to study the effect of irradiationon the nuclear materials.
75

IMR KINKEN Research Highlights 2013

The Laboratory of Alpha-Ray Emitters supports
research activities using radioactive elements. One
such example is the investigation of the lowtemperature properties of uranium (U) compounds.
As is already well known, the U compounds occupy
an important place in condensed-matter physics,
since novel phenomena have been discovered, e.g.:“ferromagnetic superconductivity,” in which mutually
exclusive ferromagnetism and superconductivity
coexist in the same compound; “hidden order,” inwhich an order parameter of the ordered state has
not been identified; “metamagnetic” transitions fromthe paramagnetic state to the highly polarized
ferromagnetic state by sweeping the magnetic field,
and so on.
To clarify the intrinsic properties of the Ucompounds, the preparation of high-quality single
crystals is necessary. For this purpose, the “tetraarc furnace” in our lab (Fig. 1) has been used bymany groups. Prof. N. K. Sato (Nagoya Univ.), one of
the users of the facility, succeeded in preparing
high-quality single crystals of UCoGe using thisfurnace and provided them to collaborators. Prof. K.
Ishida (Kyoto Univ.) performed 59Co-nuclear
magnetic resonance (NMR) and showed the intimate
relationship between ferromagnetic fluctuations andFig. 1 Photograph of the tetra-arc furnace

IMR KINKEN Research Highlights 2013

Fig. 2 (a) Magnetization along the c-axis (Mc) and Magnetic fluctuation along the c-axis (Sc) are shown on a contour plot in
the Hc-T phase diagram. The star symbol denotes the CEP determined as (μ 0Hc,T)CEP ~ (1 T, 12 K), and the Sc at thestar symbol is enhanced, indicative of the development of critical fluctuations[2]superconductivity, which is the first evidence thatferromagnetic
fluctuationsarerelatedtosuperconductivity and provides a strong constraint
to the mechanism of superconductivity in UCoGe[1].
They also collaborated with N. Kimura (Tohoku
Univ.) to study the critical behavior of magnetism
near the metamagnetic critical endpoint (CEP) in
itinerant-magnet UCoAl. They performed 27Al-NMR
and developed a contour plot of the c-axis
magnetization (Mc) and fluctuations (Sc) against thec-axis field (Hc) and temperature (T) using
measurements of the Knight Shift and 1/T1, as shownin Fig. 2. Remarkable critical behavior was observedin Sc around the CEP. They pointed out that thecritical phenomena at the itinerant-electron MM CEPin UCoAl exhibited a common feature as a “gasliquid” transition, from the critical exponent of Mc
and Sc near the CEP[2].
α-particle emitting radionuclides are gaining
attention in the field of nuclear medicine. Becauseof their high linear energy transfer and the shortrange of α-particles in tissues, α-emitters (211At,
223 Ra, and 225Ac) show promise in treating tumors
such as leukemia, melanoma, and bone metastases.
Recently, we labeled a peptide DOTA-Re(Arg11)
CCMSH (Fig. 3), which bound specifically to a
receptor expressed in melanoma cells, with an
α-emitter 225Ac. Preliminary results using the γ-ray
emitter 111In showed a high uptake of DOTARe(Arg11)CCMSH in melanoma tumors (Fig. 4).
Therapeutic studies using 225Ac are currently underway.

Fig. 3 Structure of DOTA-Re(Arg11)CCMSH labeled with
an α-particle emitter, 225Ac

Fig. 4 Melanoma imaging of
111In-labeled
DOTARe(Arg11)CCMSH
ina
B16F1 murine melanomabearing C57 mouse

References[1] T. Hattori, Y. Ihara, Y. Nakai, K. Ishida, Y. Tada, S. Fujimoto, N.
Kawakami, E. Osaki, K. Deguchi, N. K. Sato, I. Satoh, Phys.
Rev. Lett. 108, 066403 (2012) Selected for a Viewpoint inPhysics[2] K. Karube, T. Hattori, S. Kitagawa, K.Ishida, N. Kimura, T.
Komatsubara, Phys. Rev. B 86, 024428 (2012) Selected for
Editors’ suggestion

Tomoo Yamamura (Laboratory of alpha-Ray Emitters)
E-mail: yamamura@imr.tohoku.ac.jp
URL: http://alpha.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201376

Research Centers

Chemical Separation and Pre-concentration Techniques for
Determination of Trace Impurities in Metallic Materials
Analytical Research Core for Advanced MaterialsO

ur research group conducts elemental analysis of various materials to determine major, minor and tracelevels of metallic as well as nonmetallic elements. This report introduces the analysis of trace elements inmetallic materials as one of our recent activities. In order to develop analytical procedures for trace impurities
in high-purity metallic materials, we have investigated several separation and pre-concentration techniques.
Less than ppm levels of many impurity elements in metallic materials could be detected.
High-purity metallic materials are of broadinterest because of their unique mechanical andchemical properties. Recent advances in refining
technology have enabled the content of impurity
elements in various metals to be drastically reduced.
To evaluate the purity and elucidate the inherent
properties of these materials, it is essential todetermine their trace impurities.
It is necessary to separate and pre-concentratethese impurity elements from the matrix metals
because the metals may cause adverse effects
during the determination of trace-level elements [1].
Furthermore, it is impossible to apply conventional
analytical methods to high-purity materials becauseof their poor detection sensitivity as the amounts ofimpurities have been further reduced. Moreover,
several chemical procedures using harmful organic
solvents have been employed to separate and preconcentrate the trace elements. Because suchorganic solvents might be hazardous to theenvironment and to humans, analytical procedures
without harmful reagents are needed. Therefore, thedevelopment of new high-sensitivity and highprecision analytical methods are required.
For trace analysis of metal samples, several
chemical procedures for separation and preconcentration have been employed. They include
co-precipitation methods (Fig. 1 (a)) [2-5], liquidliquid extraction methods (Fig. 1 (b)) [6], gasgeneration methods (Fig. 1 (c)) [7, 8] and ionexchange methods (Fig. 1 (d)) [9].
For the identification of impurity elements,
inductively coupled plasma-optical emission
spectrometry (ICP-OES, Fig. 2), electro-thermal
atomic absorption spectrometry (ET-AAS, Fig. 3),
inductively coupled plasma-mass spectrometry(ICP-MS, Fig. 4) and spectrophotometry were
employed.
In order to develop analytical procedures fortrace impurities in high-purity metallic materials, wehave investigated several separation and preconcentration techniques. The detection limits
achieved using these methods are under μg/g (ppm)levels.

77

IMR KINKEN Research Highlights 2013

Fig. 1 Chemical separation methods for the determinationof trace impurities in metallic materials: (a) Coprecipitation, (b) Liquid-liquid extraction, (c) Gas
generation and (d) Ion exchange

Fig. 2 Echelle-type inductively coupled plasma opticalemission spectrometer

Reductive co-precipitation method withpalladium [2-4]: a sample was weighed anddecomposed in HNO3 -HCl on a heating plate.
H2SO4 was then added and the solution was heatedto evaporate HNO3. Palladium solution and ascorbic
acid powder were then added and the resulting
solution was left with a ground stopper at room

IMR KINKEN Research Highlights 2013

temperature for more than 3 hours. After being
allowed to rest, the precipitate was collected using amembrane filter. Immediately thereafter, theprecipitate was dissolved in a HNO3 -HCl mixture atroom temperature. Finally, the solution was diluted
with water.
Liquid-liquid extraction method using
4-methyl-2-pentanone [6]: a sample was weighed
and decomposed in HNO3 -HCl on a heating plate.
After this, the solution was heated until dry. Aftercooling, the dried sample was dissolved andtransferred into an extraction funnel using 5.4 mol/Lof HCl. Then, 4-methyl-2-pentanone was added andthe extraction funnel was agitated for 1-2 minutes.
After separation of the organic and inorganic
phases, the inorganic phase was transferred into aquartz beaker and evaporated on a heating plate. In
order to decompose the 4-methyl-2-pentanone
remaining in the inorganic phase completely, HNO3
was added and the solution was heated until dry.
Finally, the solution was diluted with water.
Anion exchange method [9]: a sample wasweighed and decomposed in HF-H2O2 on a heatingplate. Next, the solution was diluted with 1 mol/L of
HF and introduced into a column packed with anionexchange resin. After 2 mol/L of HF was introduced
into the column to remove the matrix iron, analytes
were eluted using an eluant (HNO3 -H2O2). H2SO4
was added to the eluant, subsequently heating tofumes. Finally, the solution was diluted with water.
By using the separation and pre-concentration
method, the influence of matrix metals on theanalysis by the analytical instruments describedabove is removed and the analytes in the sample
solution can subsequently be concentrated.
Consequently, it is possible to determine traceelements in metallic materials more sensitively andprecisely.

Fig. 3 Continuous-light-source
spectrometer

Fig. 4 Double-focused
spectrometeratomicinductively

absorption

coupledmass& TECHNIQUES, Numéro Hors Série-Décembre 31 (2003).
[7] M. Ishikuro and J. kimura, Bunseki Kagaku 37, 498 (1988).
[8] T. Ashino, K. Tozawa and K. Wagatsuma, Bunseki Kagaku 60,

References[1] K. Takada, T. Ashino, T. Syoji, T. Itagaki (Sato) and K.
Wagatsuma, Tetsu-To-Hagane, 93, 190 (2007).

175 (2011).
[9] F. Sakamoto, K. Takada and K. Wagatsuma, Bunseki Kagaku
54, 1037 (2005).

[2] T. Ashino , K. Takada and K. Hirokawa, Anal. Chim. Acta 297,443 (1994).
[3] T. Ashino and K. Takada, Anal. Chim. Acta 312, 157 (1995).
[4] T. Ashino and K. Takada, J. Anal. Atom. Spectrom. 11, 577(1996).
[5] T. Itagaki, T. Ashino, K. Takada and K. Wagatsuma, Bunseki
Kagaku 59, 43 (2010).
[6] T. Ashino, K. Takada, K. Wagatsuma and K. Abiko, MATÉRIAUX

Tetsuya Ashino (Analytical research Core for Advanced materials)
E-mail: ayustet@imr.tohoku.ac.jp
URL: http://bunseki-core.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201378

Research Centers

Significant Reduction in Transmission and Power Losses
Through Development of Nanocrystalline Soft Magnetic
Materials
Research and Development Center for Ultra High Efficiency Nanocrystalline Soft Magnetic Materials

T

he Research and Development Center for Ultra High Efficiency Nanocrystalline Soft Magnetic Materials wasestablished in June 2012. This is one of the three programs that were started by the national project of MEXT
and the Reconstruction Agency for Tohoku Innovative Materials Technology Initiatives for Reconstruction. We
aim to develop novel nanocrystalline magnetic alloys, "NANOMET™," made from nanoheterostructures obtained
after the crystallization of the amorphous phase. These alloys will be mass produced at low cost and are
expected to exhibit superior magnetic properties. A platform is to be built to enable collaborative research.
Professors, researchers and engineers with different skills from academia and industry will work together to findsolutions to a number of basic scientific issues and technical problems involved in the commercialization of
new magnetic materials with extremely low core loss. This center will also contribute to essential design
guidelines in the field of nanocrystalline soft magnetic materials. Moreover, the multifaceted collaboration
between government, industry and academia motivated by this center can lead to the blooming of industries
in the Tohoku region.

79

Necessity of a breakthrough in soft magnetic
materials
There is an urgent demand for a soft magnetic
material with a high saturation magnetic flux density(Bs) and low coercivity (Hc). However, it is hard toincrease Bs without increasing Hc at the same time
because Hc changes inversely to Bs. The soft
magnetic Fe-Co and Fe-Ni based amorphous alloys
are well known. These amorphous alloys exhibit
permeabilities of an order of magnitude higher thanpure iron but their saturated magnetic flux density isrelatively low.
The present project aims to develop new Febased soft magnetic materials with the same Bs as
for Fe-Si steels but with a lower magnetic core loss
and a higher permeability.

The "NANOMET™": a new nanocrystalline soft
magnetic material"NANOMET™," a new soft magnetic alloy
developed by our group, is made up of heterogeneous
microstructures consisting of α-Fe grains (with size
~ 10 nm) which are surrounded by thin iron-metalloid
phases. The production of "NANOMET™ " is very
cost-effective because of the low materials cost (norare earth elements) and the material's processing
in air.
It is highly anticipated that "NANOMET™ " canbe very effective at preventing energy loss inapplications such as high-current electricaltransformers and motors. At a saturated magnetic
flux density of 1.5 T, "NANOMET™ " is able to reducethe magnetic core loss by over 90% in comparison
to the commercial Fe-Si magnetic steels.

Fig. 1 Roadmap of innovation for "NANOMET™."

Fig. 2 Dependence of W on Bm of "NANOMET™."IMR KINKEN Research Highlights 2013

IMR KINKEN Research Highlights 2013

Fig. 3 Self-Assembly of heteroamorphous microstructuresof FeSiBPCu nanocrystalline alloys

Fig. 4 Dedication of "NANOMET™ " to reduce power loss

Creationofthenanoheteroamorphousmicrostructure
Single-phase Fe-based amorphous alloys havealready been commercialized and are used as the
precursors for the development of Fe-based
nanocrystalline soft magnetic alloys. However, theFe content of amorphous precursors is restricted to
below 78 at%. A higher concentration of Fe in thealloy reduces its glass-forming ability. This low
glass-forming ability leads to the precipitation of
α-Fe grains with sizes ranging from a few tenths ofnanometers to hundreds of nanometers, even in theas-quenched state. A nonuniform precipitation of
crystalline grains leads to a significant deterioration
in soft magnetic performance. By the carefulselection and adjustment of constituent elements(Fe, Si, B, P, Cu; based on interfacial energy andheats of mixing) in the alloy, our group has created anewnanoheteroamorphous
structure.
The
nanocrystallization of the alloy has been successfully
controlled to form uniform nanocrystalline (grain size
~ 10-20 nm) soft magnetic Fe-Si-B-P-Cu alloys,"NANOMET™," with Fe content > 80 at%. This new
alloy revealed superior soft magnetic performancecompared to the Fe-Si magnetic steels.

Expected achievements for "NANOMET™"
The present project aims to explore the minimum
core loss, commercialization of the wide
"NANOMET™ " thin ribbons, development of thecommercial process for the thick, bulk alloys using
gas atomization and stacking-layered methods.
Research into the basic magnetism, relationshipbetween nanocrystalline microstructure andmagnetic properties, mechanism of reduction in thecore loss and optimization of metalloid contents is tobe conducted to elucidate the theoreticalunderstanding and the material designs of the newalloy systems. The mass production of the"NANOMET™ " materials is to be achieved tocontribute to future energy savings.

Reduction of core loss
The magnetic core loss in transformers andmotors is approximately 3.4% of the total electrical
power consumption (33.5 billion Kwh per year inJapan). This energy loss can be reduced by ~ 72%with the replacement of commercial Fe-Si magnetic
steels by "NANOMET™ " as a core material.

Akihiro Makino (Research and Development Center for Ultra High Efficiency Nano-crystalline Soft Magnetic Materials)
E-mail: nanoc@imr.tohoku.ac.jp
URL: http://nanoc.imr.tohoku.ac.jp/english/index.html

IMR KINKEN Research Highlights 2013

80

2013

Index

Author Index
Name

page

Name

page

Nozawa, Jun AAbe, Hiroaki 
Aoki, Dai 
Arima, Hiroshi 
Ashino, Tetsuya

28
52
6
78

B
Boulon, Georges

68C
Chiba, Akihiko

12

10

O
Oguro, Hidetoshi
Ohoyama, Kenji 
Orimo, Shin-ichi 622032, 721842
14
36
20, 3228
26, 5844
6G
Goto, Takashi 

50

S
Saitoh, Eiji 
Sasaki, TakahikoSato, Kazuhisa
Sato, Shigeo
Sato, ToyotoSatoh, Yuhki Shikama, Tatsuo Shimizu, Yasuo 
Sugiyama, Kazumasa H
Hiraka, Haruhiro

20

T
Takanashi, Koki 

48

IIguchi, Satoshi 
Inoue, Koji Ito, Akihiko 

42
44
50

U
Umetsu, Y. Rie

56

W
Wagatsuma, Kazuaki 

36

K
Kamikawa, Naoya 
Kashiwakura, ShunsukeKato, Hidemi Katsui, Hirokazu Kawazoe, Yoshiyuki 
Kiguchi, Takanori
Konno, J. Toyohiko 
Kutsukake, Kentaro 

4
36
30
50
24
14
14
22

Y
Yamamura, Tomoo 
Yoshikawa, Akira 

76
34

F
Fujita, Motoki 
Furuhara, Tadashi 

54
4, 70

M
Makino, Akihiro Masahashi, Naoya
Matsuoka, TakashiMatsuura, Masato Miyamoto, Goro 
Mohri, Tetsuo 60, 8064
46
20
4
16, 66

N
Nagai, Yasuyoshi 
Niinomi, Mitsuo Nojima, Tsutomu 
Nojiri, Hiroyuki 
Nomura, Kentaro 

448
74
40, 6838

IMR KINKEN Research Highlights 201382

Keyword Index
Keyword

page

A
actinide
all-electron mixed basis approachalloy
amorphous
atom probe tomography

52
24
8, 14
56
44

B
biomedical

8, 12

C
Casida equation
24
ceramics
26
chemical vapor deposition (cvd) 50colloid10
cryptocrystalline
6
crystal growth
10, 22, 34
D
3D structure
defectsdeformationdielectric properties
dispersion coefficients
dopingEelectrical properties
electrical resistivity
electron microscopy
electronic material
energetic material
energy storage

83

44
16
4
42
24
10
38
56
28
42, 44, 50
18
32

F
Fermi surface
first-principles calculations
foam

52
16
30

G
grain boundaries

22

H
heavy fermion
high magnetic field
high strength steel
high-Tc cupratehydroxyapatite

52
40
16
20
8

I
instrumentationinsulator
ion beam analysis

40
38
26

IMR KINKEN Research Highlights 2013

KeywordL
luminescence

page
26

M
magnetic propertiesmass spectroscopy
mechanical properties
microstructure38, 4836
4, 12, 28
30, 50N
neutron scattering
nitride20, 32, 54
46

O
organic

42

P
pair distribution function
phase transformation
piezoelectric
porosity

6
14
34
30

Q
quantum spin system

54

Rradiation effects

28, 34S
scanning transmission electron microscopy(stem)
14
solar cells
22
spectroscopy
40
spin dynamics
54
spin fluctuation
20
spin wave
18, 48spintronic
48
spintronics
18
steel
4superplasticity
12Vvapor phase epitaxy (vpe)

46

W
waste management

36

X
XAFS
x-ray diffraction (xrd)

646

Editors:
Hiroaki Abe (IMR), Takashi Matsuoka (IMR),
Gerrit E.W. Bauer (IMR), Atsushi Tsukazaki (IMR),
Kozo Fujiwara (IMR), Yuichiro Koizumi (IMR), andSatoshi Iguchi (IMR)

Editorial Staff:
Yoshiko Takahashi (IMR), Miori Kikuchi (IMR)

Printing:
HOKUTO Corporation

Organization
Institute for Materials Research
Tohoku University
2-1-1 Katahira, Aoba-ku, Sendai 980-8577, Japan
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URL: http://www.imr.tohoku.ac.jp/
E-mail: pro-adm@imr.tohoku.ac.jp© IMR October 2013