KINKEN

Research Highlights
2012

1986. February
Courtesy of Tohoku University Historical Museum.

Institute for Materials Research,Tohoku University

KINKEN Research Highlights 2012
Contents
Preface���������������������������������������������������������������������������������������������������������������������������������������������������������� 1

Research Highlights

1. Infrastructural Materials
Microstructure Design of Structural Metallic Materials Division�������������������������������������������������������������������� 3
Chemical Physics of Non-Crystalline Materials Division������������������������������������������������������������������������������ 5
Bio-Materials Science Division��������������������������������������������������������������������������������������������������������������������� 7
Multi-Functional Materials Science Division������������������������������������������������������������������������������������������������� 9
Deformation Processing Division����������������������������������������������������������������������������������������������������������������11
Advanced Materials Development and Integration ofNovel Structured Metallic and Inorganic Materials������������������������������������������������������������������������������������� 13

2. Energy-related Materials
Crystal Physics Division����������������������������������������������������������������������������������������������������������������������������� 15
Surface and Interface Research Division��������������������������������������������������������������������������������������������������� 17
Metal Physics with Quantum Beam Spectroscopy Division����������������������������������������������������������������������� 19
Materials Design by Computer Simulation Division������������������������������������������������������������������������������������ 21
Irradiation Research on Nuclear Materials Division������������������������������������������������������������������������������������ 23
Nuclear Materials Science Division������������������������������������������������������������������������������������������������������������ 25
Nuclear Materials Engineering Division������������������������������������������������������������������������������������������������������ 27
Non-Equilibrium Materials Division������������������������������������������������������������������������������������������������������������� 29
Crystal Chemistry Division������������������������������������������������������������������������������������������������������������������������� 31
Hydrogen Functional Materials Division����������������������������������������������������������������������������������������������������� 33
Advanced Analysis of Materials Division���������������������������������������������������������������������������������������������������� 35
NEDO Laboratory for Metallic Glass Research Collaboration and Promotion������������������������������������������� 37

3. Electronic Materials
Theory of Solid State Physics Division������������������������������������������������������������������������������������������������������� 39
Magnetism Division������������������������������������������������������������������������������������������������������������������������������������ 41
Low Temperature Condensed State Physics Division�������������������������������������������������������������������������������� 43
Physics of Crystal Defects Division������������������������������������������������������������������������������������������������������������ 45
Physics of Electronic Materials Division����������������������������������������������������������������������������������������������������� 47
Magnetic Materials Division������������������������������������������������������������������������������������������������������������������������ 49
Radiochemistry of Metals Division������������������������������������������������������������������������������������������������������������� 51
Analytical Science Division������������������������������������������������������������������������������������������������������������������������ 53
Materials Processing and Characterization Division���������������������������������������������������������������������������������� 55
Advanced Crystal Engineering Laboratory������������������������������������������������������������������������������������������������ 57

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

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

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

Preface

Dear colleagues,

Around one and half years have passed since the Great East Japan Earthquake happened on 11
March 2011. The Institute for Materials Research, Tohoku University (KINKEN) has already recovered
and the people at KINKEN are working under ordinary conditions. We would like to express our sincere
thanks to you for helping us to come back to ordinary working conditions after the horrible disaster.

We are pleased to bring out KINKEN Research Highlights 2012, which is the sixth annual report witha collection of research output during the past year from the Institute for Materials Research (IMR),
Tohoku University. KINKEN (short Japanese name of IMR, which is familiar to the materials community,
particularly in Japan) determined in 2007 to launch this report annually so that our colleagues aroundthe world can recognize our research activities. In this report, the research activities of individual researchlaboratories and research centers, as well as joint projects with other academic institutions, are provided,
especially for the following three key research fields: (1) Infrastructural Materials, (2) Energy-relatedMaterials, and (3) Electronic Materials.
The Institute was established at Tohoku Imperial University in 1916 by Professor Kotaro Honda as asteel research organization known officially as the 2nd Division of the Provisional Institute of Physical andChemical Research. In 1987, the Institute was reorganized into its present form as a national collaborative
research institute attached to Tohoku University. The name of the institute was consequently changed to
the Institute for Materials Research (IMR). Since its establishment, IMR has become known for its
excellence in both basic and applied research on metals and a wide range of new materials. A variety
of functional materials have been studied and developed at the Institute, including the KS magnet steel,new KS magnet steel, SiC fiber, compound alloys, metallic glasses, and others.

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IMR KINKEN Research Highlights 2012

Director

IMR has greatly contributed to the advancement of materials science. As one example, our study ofmagnetism resulted in the invention of the strongest permanent magnet, the KS magnet, which
dramatically improves the performance of electrical machines. IMR has been paying great attention tobasic research that opens up the most advanced areas of applications. More recently, IMR has createda wide array of new materials, including high-performance, high-quality, and multifunctional materials
such as amorphous alloys; bulk metallic glasses; intermetallic compounds; quasicrystals; oxides;
ceramics; nanostructured metals; Si; Ge; III-V, II-VI, and oxide semiconductors; optical and electronicmaterials; solar cell crystals; biomaterials; organic materials; hydrogen storage alloys; and shaped
crystals.
Now, we are facing the serious issues of the deteriorating global environment and the depletion ofworldwide natural resources and energy sources. IMR will further promote materials research to address
the problems associated with the global environment and energy resources, aiming at sustaining humandevelopment and ensuring a high standard of living for all people. In this way, we will offer a brighter
future for the next generation.
We earnestly hope that KINKEN Research Highlights will help you recognize and support our
research activities and will serve as a medium to promote worldwide collaborations in materials researchwith our institute. We sincerely ask for your continuing support and welcome any suggestions.

Sincerely yours

IMR KINKEN Research Highlights 20122

Infrastructural Materials

Strengthening of Low-Alloy Hot-Forging Steel with Nanosized Precipitation

In a recent national research project on high-performance steels, the strengthening of medium-carbon lowalloy steel for automobile parts was extensively studied. The key technology for strengthening is the use of finedispersions of alloy carbides several nanometers in size. Here, we describe briefly the microstructure–property
relationships from our research conducted in this project.

To reduce the fuel consumption by automobiles,
medium carbon steels for forging parts, such as
shafts and connecting rods, must be further
strengthened. When carbide-forming elements suchas Ti, Nb, and V, are alloyed with carbon steels, afine dispersion of alloy carbides and large
precipitation hardening can be achieved. Continuous
cooling after austenitization has been frequently
reported to produce repeated nucleation of alloy
carbides at ferrite/austenite interphase boundaries
during the growth of ferrite or pearlite, leading toformation of precipitates in regularly spaced rows.
This peculiar precipitation phenomenon is known asinterphase precipitation [1].
In medium- and high-carbon steels containing
vanadium, interphase precipitation of VC is used forstrengthening. However, there are few studies oninterphase precipitation in medium- and highcarbon steels, and the precipitates formed throughinterphase precipitation in those steels have notbeen characterized quantitatively. In this study [2],Fig. 1 Effects of V content on the maximum hardness ofthe specimens isothermally transformed at various
temperatures.

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IMR KINKEN Research Highlights 2012

Fig. 2 Dark-field TEM of VC dispersion formed via
interphase precipitation in low alloy 0.4% C steels
containing 0.5% V (transformed at 923 K).

the variations in hardness and microstructure,including the size and number density of VC with V
addition, have been investigated.
Fig. 1 shows the variations in peak hardness atvarious temperatures as a function of V content.
Peak hardness increases linearly with V content at
all three temperatures, while the slope at 973 K issmaller than those at 873 K and 923 K. As thetemperature at which the ferrite/pearlite transformation
occurs decreases, greater hardening is observed.
Fig. 2 is a dark-field TEM micrograph of theinterphase precipitation of VC in the 0.4% C alloy
containing 0.5% V. It is clearly seen that fine VC
particles (imaged as bright particles) is aligned inrows parallel to the interphase boundary. Such VC
precipitation was also observed in low-alloy eutectoid
steel containing vanadium.
Precipitation/dispersion strengthening can beclassified into the following two categories: a cutting
mechanism and a by-pass (Orowan) mechanism. A
larger particle size leads to higher strength in theformer case and vice-versa in the latter case. The

IMR KINKEN Research Highlights 2012

Fig. 3 Comparison of the increase in strength by precipitates between calculation based on the Ashby–Orowan model andexperiments: (a) original Ashby–Orowan model where λ = x and (b) modified Ashby–Orowan model assuming λ to be
3x.

increase in yield strength in low-carbon ferritic steel
containing fine alloy carbide particles is well
explained by the Orowan-type dispersion hardening
model modified by Ashby [3].

References[1] R. W. K. Honeycombe: Phase transformation in ferrous alloy,
TMS-AIME, Pennsylvania, 259 (1984).
[2] G. Miyamoto, R. Hori, B. Poorganji, T. Furuhara: ISIJ Int., 51,
1733 (2011).
[3] T. Gladman: The Physical Metallurgy of Microalloyed Steels,
Maney Publishing, London p. 53 (2002).

Here, L is the average spacing between particles inthe slip plane, G is the shear modulus, b is themagnitude of Burgers vector, and λ is the outer cutoff distance.
Fig. 3 shows the relationship between theestimated and observed values. The original Ashby–
Orowan model underestimates the strength
significantly, as shown in Fig. 3(a). After somemodification of the cut-off distance, a morereasonable agreement between observation andestimation was obtained (Fig. 3(b)). These findings
were used in the first systematic, physically basedmodel proposed to predict the microstructure andyield strength of vanadium-microalloyed mediumcarbon steels [4].

[4] H. -H. Kuo, M. Umemoto, K. Sugita, G. Miyamoto, T. Furuhara:
ISIJ Int., 53, 669 (2012).

Keywords: steel, phase transformation
Goro Miyamoto (Microstructure Design of Structural Metallic Materials Division)
E-mail: st-mat@imr.tohoku.ac.jp
URL: http://www.st-mat.imr.tohoku.ac.jp/F-top-E.html

IMR KINKEN Research Highlights 20124

Infrastructural Materials

Structural Analysis of Amorphous Zr50Ni50

The structure of Zr Ni amorphous alloy was analyzed by anomalous X-ray scattering (AXS) coupled with50

50

reverse Monte-Carlo (RMC) simulation. Topological features in the nearest-neighbor region suggest that thestrong chemical interaction between Ni and Zr pairs breaks the icosahedral-like local ordering units commonly
realized in the dense random packing of hard sphere (DRPHS) model [1].

Zr-based amorphous alloys have attracted
many researchers’ interests because of theircharacteristic physico-chemical properties [2-3].
Among them, Zr–Cu and Zr–Ni amorphous alloys
have been extensively studied to elucidate therelationship between thermal stability and atomic
scale structure. In the case of Zr–Ni amorphousalloy, the short distance of the Zr–Ni pair in thenearest-neighbor region was demonstrated to be aconsequence of a strong chemical interaction, andthis suggests crystal-like chemical short-rangeordering (CSRO) around Ni [4]. Nevertheless, theavailable structural information is insufficient toexamine and clarify the possible CSRO around Ni.
This prompted us to perform AXS-RMC analysis of
the Zr–Ni amorphous alloy sample. AXS-RMC
analysis is suggested to be one of the most powerful
methods for elucidating stereoscopic atomic
configurations.
The detailed procedure of AXS-RMC analysis
has been discussed previously [5,6]. For theconvenience of discussion, only the important points
are given below. AXS measurements at the Ni K
absorption edge and the Zr K absorption edge werecarried out at beam line stations 7C and NW10A of
the Institute of Materials Structure Science, HighEnergy Accelerator Research Organization,
Tsukuba, Japan. Pairs of incident energies
corresponding to 25 eV and 300 eV below each of
the Ni and Zr K absorption edges were employed in
the present AXS measurements. The RMC modelingstarted with an initial model of 2,000 atoms (Zr:
1,000 and Ni: 1,000) with a body-centered cubic(BCC) arrangement in a cubic hyper-cell (density =7.36 Mg/m3).
Fig. 1 shows the ordinary RDF measured at E =17.699 keV (Zr K edge - 300 eV) and the environmental
RDFs around Ni and Zr obtained by AXS
measurements. The vertical dashed lines indicate
the distances for the Ni–Ni, Ni–Zr, and Zr–Zr pairs
estimated from Goldschmidt radii. The ordinary RDF
suggests three kinds of correlations in the 0.2–0.4
nm range. The hump at around 0.27 nm is a harmonyof the Ni–Ni and Zr–Ni pairs, and the remaining
peak at 0.32 nm corresponds to the Zr–Zr pair. On
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IMR KINKEN Research Highlights 2012

the other hand, the environmental RDF around Ni
showed only the harmony of the Ni–Ni and Zr–Nipairs; the Zr–Zr correlation in the ordinary RDF was
not observed here. A similar discussion is possible
for the case of AXS analysis for Zr, where thestructural information on the Zr–Ni and Zr–Zr pairs is
included. It should be noted that the first-neighbor
Zr–Ni correlation in the ordinary RDF is obviously
shorter than that for the Zr–Zr correlation. Thisfeature could be considered as an indication of thestrong chemical interaction between Zr and Ni.

Fig. 1 Ordinary and environmental RDFs for Zr50Ni50
amorphous alloy. Vertical dashed lines indicate atomic
distances estimated from the Goldschmidt radii.

The AXS-RMC analysis reproduces three
independent interference functions QΔiNi (Q),QΔiZr (Q), and Qi(Q), and we can discuss the threegij (r) functions for Zr50Ni50 amorphous alloy, as
shown in Fig. 2. Nevertheless, it should be notedthat the partial structure factors and correspondinggij (r) functions estimated from AXS-RMC analysis
are necessary conditions at best, not sufficientconditions. The maximum peak positions for Ni–Ni

IMR KINKEN Research Highlights 2012

and Zr–Zr like-atom pairs indicate almost equal
distance, respectively, in comparison with thoseestimated by Goldschmidt radii. On the other hand,the distance of the unlike-atom Ni–Zr pair is shorter
than the estimated value. This feature again suggests
the strong bonding between the unlike atomic pair.
The calculated coordination numbers in the nearestneighbor region indicate no strong preference forthe correlation of unlike-atom pairs, as suggested in
the previous report.

Fig. 3 Preferred Voronoi indices around Ni and Cu foundin amorphous alloys. Filled and hatched bars correspondto those for Zr50 Cu50 [5] and Zr50Ni50, respectively.

values for Cu and Ni show no significant difference[5]. This suggests that the strong chemical interaction
between Zr and Ni breaks the icosahedral-like local
ordering structure common in the DRPHS model.
References[1] T.Kawamata, K.Sugiyama and Y.Yokoyama: Key Eng. Mater.

Fig. 2 Three partial gij (r) functions of Zr50Ni50 amorphousalloy.

Fig. 3 shows the preferred Voronoi index aroundNi found in Zr50Ni50 amorphous alloy. The ideal
icosahedron with a Voronoi index ( 0 0 12 0 0 ) was
found as one of the typical environments around Ni,
with a 4.1 % rate of occurrence. This fraction is muchlarger than that around Zr (0.3 %). This result isreadily explained by considering the atomic size
ratio R*: the value of R* for Ni was much closer tothat for an ideal icosahedron (R* = 0.90) than that of
Zr. It may be worthwhile to note that, in Zr–Ni alloys,
the fraction of icosahedral structure around Ni ismuch smaller than that around Cu, although R*

508, 146–151 (2012).
[2] W. H. Wang, C. Dong and C. H. Shek: Mater. Sci. Eng. R 44,
45–89 (2004).
[3] A. Inoue and A. Takeuchi: Mater. Sci. Eng. A 375–377, 16–30(2004).
[4] X. J. Liu, X. D. Hui, G. L. Chen and T. Liu: Phys. Lett. A 373,
2488–2493 (2009).
[5] T. Kawamata, Y. Yokoyama, M. Saito, K. Sugiyama and Y.
Waseda: Mater. Trans. 51, 1796–1801 (2010).
[6] K. Sugiyama, T. Muto, T. Kawamata, Y. Yokoyama and Y.
Waseda: Phil. Mag. 91, 2962–2970 (2011).

Keywords: structure, amorphous, diffraction
Kazumasa Sugiyama, Toru Kawamata (Chemical Physics of Non-Crystalline Materials Division)
E-mail: kazumasa@imr.tohoku.ac.jp
URL: http://www.xraylab.imr.tohoku.ac.jp/index.htm

IMR KINKEN Research Highlights 2012

6

Infrastructural Materials

Characteristics of Biomedical β-type Titanium Alloy
Subjected to High-Pressure Torsion

β-type titanium alloy, T–29Nb–13Ta–4.6Zr (TNTZ), developed for use in biomedical applications, exhibits nontoxicity and a low Young’s modulus similar to that of bone. Severe plastic deformation, such as high-pressure
torsion (HPT), is effective for grain refinement and dislocation strengthening. In this study, the effect of HPT on
the microstructure and hardness of TNTZ was systematically investigated to improve its mechanical strength
while maintaining the low Young’s modulus. In the cross sections of TNTZ subjected to HPT, a heterogeneous
microstructure consisting of a matrix and an unetched band not corroded by an etching solution were observed.
Both the matrix and the unetched band comprise a single β phase. Furthermore, submicron-sized grains or
subgrains were formed by HPT, and grain misorientation increased with the number of HPT rotations. The
hardness distribution in the cross section of TNTZ subjected to HPT was also heterogeneous; the hardness
was higher in the matrix than in the unetched band when the number of HPT rotations was small. However, withincreasing number of HPT rotations, the hardness of the unetched band became comparable to that of thematrix.
Stress transfer between a metallic implant andbone is non-homogeneous when the Young’s
modulus of the metallic implant and that of the bone
are different. The Young’s moduli of the metals andalloys used for fabricating metallic implants aremuch higher than that of bone; therefore, stressstimulation to the bone is reduced (stress shielding
effect), with the result of inducing bone atrophy andthen loosening of the metallic implant and refracturing of the bone. Thus, to mitigate the stressshielding effect, metals and alloys with Young’s
moduli equal to that of bone (10–30 GPa) are
believed to be ideal for fabricating metallic implants[1]. Consequently, extensive efforts have been madeto develop β-type titanium alloys with Young’s moduli
nearly equal to that of bone. One such alloy, Ti–
29Nb–13Ta–4.6Zr (TNTZ), has been developed by
the authors [2]. This alloy has a low Young’s modulus
of around 60 GPa under a solutionized condition.
However, the mechanical strength of this alloy under
the solutionized condition is less than that of theconventional titanium alloy, Ti–6Al–4V ELI, used for
fabricating metallic implants. Thus, various
thermomechanical treatments have been examined
to improve the mechanical strength of TNTZ [3].
These treatments improve the mechanical strength,
but they also increase the Young’s modulus becausea large amount of precipitates is formed. Therefore,
a method that affords TNTZ with both low Young’s
modulus and high mechanical strength is required.
Grain refinement is generally an effective
method for strengthening metals and alloys. As amethod for obtaining submicron-sized grains, severeplastic deformation has attracted considerable
attention. In comparison with other strengthening
methods, grain refinement is expected to achieve
high mechanical strength while maintaining a low
Young’s modulus for TNTZ because it retains the βphase, which ensures the desired low Young’s
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IMR KINKEN Research Highlights 2012

modulus. Therefore, by employing high-pressure
torsion (HPT) as a representative severe plastic
deformation technique, we recently attempted to
improve the mechanical strength of TNTZ while
keeping its Young’s modulus low [4, 5].
A hot-forged bar of TNTZ was used in this study.
The bar was subjected to solution treatment, coldrolling, and high-pressure torsion under an appliedpressure of around 1.25 GPa with a rotation speedof 0.2 rpm for 1, 5, 10, or 20 rotations (N) at room
temperature. The microstructures were evaluated by
opticalmicroscopy,transmission
electronmicroscopy (TEM), and X-ray diffraction (XRD). The
hardness was measured by subjecting thespecimens to the Vickers hardness test.
Fig. 1 shows optical micrographs of the TNTZ
cross section subjected to HPT at N = 10, as arepresentative result. The white band (unetchedband), which was not corroded by the etchingsolution, is observed horizontally in the middle ofthis TNTZ cross section. The unetched band is
expected to have higher corrosion resistance thanFig. 1 Optical micrographs of cross section of TNTZ
subjected to HPT at N = 10.

IMR KINKEN Research Highlights 2012

(b)

(a)

100 nm

100 nmFig. 2 Results of TEM observation at a plan view from thecenter of the thickness direction of TNTZ subjected to(a) cold rolling and (b) subsequent HPT at N = 5.

Fig. 3 Hardness distributions on cross sections of TNTZ
subjected to HPT at N = (a) 1 and (b) 20.

the other regions (matrix). Such unetched bands
were observed in the other TNTZ cross sections
subjected to HPT under all specified rotations, butthey did not form after cold rolling. Furthermore, thethickness of the unetched band increased with thenumber of rotations and the center-to-edge distance.
The equivalent strain increased with the number ofrotations and the radius of the specimen. Therefore,
this result indicates that the unetched band wasformed in the region where a large strain was
imposed.
The XRD profiles of the TNTZ surfaces subjected
to cold rolling and subsequent HPT at each rotation
showed only the β phase for every TNTZ specimen;
the matrix in TNTZ subjected to HPT for any numberof rotations is probably composed of a single βphase.
Fig. 2 shows the TEM results obtained from aplan view from the center of the thickness direction
TNTZ subjected to cold rolling and subsequent HPT
at N = 5. These regions correspond to the unetchedbands observed in the TNTZ cross sections
subjected to HPT. In the selected area electron
diffraction patterns, spots derived from the β phaseare observed in TNTZ subjected to cold rolling,whereas Debye rings appear for TNTZ subjected to
HPT at N = 5. This result indicates that the unetchedband comprises a single β phase as well as thematrix. Furthermore, small grains or subgrains areformed by HPT, and grain misorientation increases
with the number of HPT rotations. In other words,
the number of high-angle grain or subgrainboundaries increases with strain.
Fig. 3 shows the hardness distributions in crosssections of TNTZ subjected to HPT at N = 1 and 20.
Overall, the hardness values tend to increase with
the number of rotations. Furthermore, the hardness

values increase gradually with the center-to-edge
distance. These results indicate that the hardness
increases with the equivalent strain. However, thehardness at the center of the TNTZ cross section
subjected to HPT at N = 1 is lower than that near its
surface, whereas the hardness of the TNTZ
specimen subjected to HPT at N = 20 is almost
uniform throughout the thickness direction. Unetchedbands were formed by HPT at the middle of thecross sections of every TNTZ. Therefore, thehardness of the unetched band is found to be lessthan that of the matrix in TNTZ subjected to HPT at
N = 1 but comparable to that of the matrix in TNTZ
subjected to HPT at N = 20. The misorientation ofgrains or subgrains in the unetched band increased
with the number of rotations, as shown in Fig. 2.
Thus, it is considered that the hardness increases
with the number of high-angle grain or subgrainboundaries.
References[1] M. Niinomi, J. Mech. Behav. Biomed. Mater. 1, 30 (2008).
[2] D. Kuroda, M. Niinomi, M. Morinaga, Y. Kato and T. Yashiro,
Mater. Sci. Eng. A 243, 244 (1998).
[3] T. Akahori, M. Niinomi, H. Fukui, M. Ogawa and H. Toda,
Mater. Sci. Eng. C 25, 248 (2005).
[4] H. Yilmazer, M. Niinomi, M. Nakai, J. Hieda, Y. Todaka, T.
Akahori and T. Miyazaki, J. Mech. Behav. Biomed. Mater. 10,
235 (2012).
[5] H. Yilmazer, M. Niinomi, M. Nakai, J. Hieda, T. Akahori and Y.
Todaka, Key Eng. Mater. 508, 152 (2012).

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

IMR KINKEN Research Highlights 20128

Infrastructural Materials

Transparent Lutetium-based Oxides Prepared by Spark
Plasma SinteringTransparent polycrystalline ceramics are promising optical materials exhibiting excellent properties that
surpass single crystals. Spark plasma sintering is a versatile technique to obtain fully dense ceramics with finegrains and high transparency because of its fast densification. We demonstrated the preparations of transparent,
practical, and novel Lu-based oxides by spark plasma sintering.
Rapid progress in the development of opticaldevices, such as lasers and scintillators, requires
novel materials and processing techniques. Opticalmaterials are commonly used in a single crystal
form grown from melt. Therefore, a high melting
point of a potential compound often prevents a new
optical material design.
Lutetium (Lu)-based oxides, Lu2O3 andLu3Al5O12 (LuAG), have recently attracted muchattention for their potential applications in lasers,
scintillators, and phosphors because of lutetium’s
small ionic radius and high atomic number (Z = 71)
among rare earth elements. Conversely, Lu-based
complex oxides, such as Lu2Ti2O7 and Lu3NbO7,
have rarely been studied in either single crystal or
polycrystalline form, due to their high melting points(over 2300 K).
Transparent polycrystalline ceramics are analternative to single crystals because of their shortproduction time, feasibility to produce large andcomplicated shapes, and high-concentration,
uniform doping of active ions. The main issue inconventional sintering techniques is the requirementof homemade nanopowders. Moveover, the hightemperature and prolonged time causes graincoarsening, which results in low transmittance andb
Uniaxialpressure
Punch
Mold

Pulsedcurrent

Lutetium sesquioxide: Lu2O3
Lu2O3 exhibits high thermal conductivity among
rare-earth sesquioxides, which makes it a promising
optical medium for high power lasers. We proposeda two-step pressure sintering combined with a low
heating rate to prepare transparent Lu2O3 ceramics(Fig. 1b) [1–3]. The preload pressure affected theintermediate relative density, and the transmittancein the visible range was more sensitive to preloadpressure than that in the infrared range. A highheating rate resulted in a significant increase indensity but a large grain size and porosity, which
degraded transparency. The combination of a low
heating rate and the two-step pressure profile
resulted in a high transparency by inhibiting defectformation and eliminating pores. High transmittances

Pressure Densificationalow mechanical properties.
Spark plasma sintering (SPS) is a versatiletechnique to prepare transparent polycrystalline
ceramics because of its fast densification by the
combination of pulsed current and uniaxial pressure(Fig. 1a). A short holding time (several minutes) anda low sintering temperature prevent grain coarsening
and favor the fabrication of fine-grained ceramics
with high transparencies.

I

II

i

III

ii

Sintering stages:
I: Packing
II: Grain sliding, diffusion
III: Pore elimination
Merit of two-step pressure:i: Particle rearrangementii: Rapid densification

Temperaturec
PowderPunch

Fig. 1 a: SPS apparatus. b: Densification curve and sintering stages plotted against temperature: (stage I) particle packing,(II) grain sliding and grain boundary diffusion (III) pore elimination and grain creep. The advantages of a two-steppressure profile are enhancement of particle rearrangement by a low pressure at stage I and rapid densification under ahigh pressure at stages II and III. c: Transparent Lu2O3 ceramic 10 mm in diameter prepared by SPS.

9IMR KINKEN Research Highlights 2012

IMR KINKEN Research Highlights 2012of 60% at 550 nm and 79% at 2000 nm wereobtained under a preload pressure of 30 MPa at aheating rate of 0.17 K s−1 (Fig. 1c). Transparent Y2O3
ceramics can also be prepared by SPS using amoderate temperature and pressure profile [4].
Lutetium aluminum garnet: Lu3Al5O12
Lutetium aluminum garnet (Lu3Al5O12, LuAG) isa promising host medium for scintillation materials
owing to its high density (6.7 Mg m−3) and higheffective atomic number (Zeff = 63). LuAG can be an
efficient and fast-response scintillator when dopedwith rare-earth ions such as Ce3+ and Pr3+. Singlephase LuAG bodies with high densities (>99.5%)
were obtained at sintering temperatures of 1573–
1923 K. The average grain size of the LuAG bodies
was 0.18–0.52 μm at 1573–1773 K and the optimal
sintering temperature was 1773 K, which showed atransmittance of 77.8% at 2000 nm after annealing
at 1423 K in air for 43.2 ks (Fig. 2a). Grain growth
became significant at 1823 K and pores wereobserved at the triple junctions of the grainboundaries at high sintering temperatures, which
decreased transparency [5].
Lutetium fluorite-related oxides: Lu2Ti2O7,
Lu2Hf2O7, and Lu3NbO7
Lutetium titanate (Lu2Ti2O7) has a pyrochlore
structure, consisting of two interpenetrating networks
of corner-sharing Ti2O6 octahedra and Lu2O chains,which provides unique properties. Therefore,transparent Lu2Ti2O7 is an interesting multifunctional
material that combines optical and other properties,
such as ion conduction, dielectric behavior, catalytic
activity, and radiation resistance. Single-phasepyrochlore Lu2Ti2O7 ceramics were prepared atsintering temperatures above 1273 K. A
homogeneous microstructure was observed at
1573–1723 K, and grain coarsening occurred at1823 K. The transmittance showed the highest
values, 57% and 74% at wavelengths of 550 and2000 nm, at a sintering temperature of 1723 K, whichmight be related to the homogeneous microstructure(Fig. 2b) [6,7]. The effect of calcination of the starting
powder was also investigated [8].
Lutetium niobate (Lu3NbO7) has a cubic defectfluorite structure, where Lu and Nb cations randomly
occupy the 4a site, and 1/8 of the 8c sites for the O
anion are statistically defective. Such disordering of

a

b

c

Fig. 2 a: Transparent LuAG ceramics sintered by SPS atvarious temperatures. b: Transparent Lu2Ti2O7 ceramics
prepared by SPS before (right) and after (left) annealing.
c: Transparent Lu3NbO7 ceramics prepared by SPS (10mm in diameter).

cations and O defects make Lu3NbO7 an interesting
multifunctional material. We demonstrated the
preparation of transparent Lu3NbO7 ceramics forthe first time. The transparent Lu3NbO7 body had auniform microstructure with an average grain size of
0.77 μm. The transmittance at 550 nm reached 68%(Fig. 2c) [9].
Lu2Hf2O7 is a promising scintillation material, as
it has the highest density (9.95 Mg m−3) and higheffective atomic number (Zeff = 68.9) among rareearth hafnates; however only the powder synthesis
has been reported because of its high melting point(2700 K). A transparent Lu2Hf2O7 body was preparedby reactive SPS using Lu2O3 and HfO2 powders.
The transmittance reached 78% at 2000 nm after
annealing at 1123 K for 21.6 ks [10].
References[1] L.Q. An, A. Ito, T. Goto, J. Am. Ceram. Soc. 94, 695 (2011).
[2] L.Q. An, A. Ito, T. Goto, J. Eur. Ceram. Soc. 31, 1597 (2011).
[3] L.Q. An, A. Ito, T. Goto, Ceram. Int. 37, 2263 (2011).
[4] L.Q. An, A. Ito, T. Goto, J. Eur. Ceram. Soc. 32, 1035 (2012).
[5] L.Q. An, A. Ito, T. Goto, J. Eur. Ceram. Soc. 32, 3097 (2012).
[6] L.Q. An, A. Ito, T. Goto, J. Eur. Ceram. Soc. 31, 237 (2011).
[7] L.Q. An, A. Ito, T. Goto, J. Am. Ceram. Soc. 94, 3851 (2011).
[8] L.Q. An, A. Ito, T. Goto, Ceram. Int. 38, 4973 (2012).
[9] L.Q. An, A. Ito, T. Goto, Mater. Lett. 65, 3167 (2011).
[10] L.Q. An, A. Ito, T. Goto, Opt. Mater. (2012) accepted.

Keywords: ceramics, sintering, microstructure
Takashi Goto, Liqiong An, 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 201210

Infrastructural Materials

Origin of Metastability and Strain-induced Martensitic
Transformation in Biomedical-grade Co–Cr–Mo Alloys with
Negative Stacking Fault Energy

B

iomedical-grade Co–Cr–Mo alloys are used as artificial hip or knee joints, which need significantly high wear
resistance as well as good biocompatibility and mechanical properties. The phase stability of the alloys is affectedby a small amount of nitrogen doping. The metastable γ-fcc phase exhibits negative stacking fault energy at room
temperature. We have revealed the mechanism of meta-stability for the γ phase, which is thermodynamicallystable at temperatures higher than approximately 1223 K. The strain-induced martensitic transformation has been
investigated to elucidate the plastic deformation of Co–Cr–Mo alloys at room temperature.
The Co–29Cr–6Mo (CCM) alloy has two phases:a γ phase (face-centered cubic (FCC) structure) anda ε phase (hexagonal close packing (HCP) structure),
which are stable at high and low temperatures,
respectively. Recent research indicates that a small
amount of N addition (0.1–0.16 mass%) to CCM alloyleads to the Co–Cr–Mo–N (CCMN) alloy, and full γ
phase microstructure is obtained after even a
moderate rate cooling from high temperature.
Nevertheless, the γ phase in CCMN alloy is still
metastable even at room temperature.
In our study [1], the distribution of the N atoms
with respect to those of the other three elements inthe alloy was observed using a 3D atom probe, as
shown in Fig. 1 a, for Cr and N atoms. The N atoms(red circles) were randomly distributed about the Co
and Mo, indicating that there was no obvious affinity
between the N atoms and the Co or Mo atoms.
However, each N atom was surrounded by Cr atoms,
implying that the N atoms selectively existed in thevicinity of the Cr atoms. The density distributions of
the Cr and Mo atoms are shown as functions of the
distance from a N atom in Fig. 1 b. The normalized Cr
atom density was markedly higher than the Mo atom
density in the vicinity of the N atoms, implying that the
N atoms were closer to the Cr atoms. Thus, it isapparent that the Cr atoms interact more strongly with
the N atoms than do the Mo and Co atoms, which
form an inert matrix in this alloy. The atomic radius of
N (0.80 Å) is slightly larger than one-half of the atomic
radii of Co, Cr, and Mo (1.26, 1.27, and 1.40 Å,
respectively). The N dissolved in the CCMN alloy is
speculated to occupy the octahedral interstitial sitesof the matrix. Since an additional driving force isconsidered necessary to destroy or alter the existing
Cr–N cluster during the γ→ε phase transformation,the γ phase is indirectly stabilized. In other words, theformation of Cr–N clusters increases the energy
barrier for the γ→ε phase transition, although thistransition remains a G-decreasing process.
On the basis of the abovementioned results, we
propose a schematic model of the positioning of theelements in the CCMN alloy (Fig. 1 c). The Cr and
Mo atoms are thought to be randomly distributed in
11

IMR KINKEN Research Highlights 2012

Fig. 1 (a) 3-D atom probe tomography for Cr–N. (b) Atomic
density of Cr and Mo atoms as a function of distancefrom the N atom; the inset graph is a schematic modelof the positioning of N atoms in matrix, in which N atoms
selectively occupy the octahedral interstitial sites whereat least one Cr atom is present in the vicinity. (c)Schematic representation of γ-phase stabilization by N
addition in CCMN alloy.

the Co matrix, while the N atoms selectively occupy
the octahedral interstitial sites where at least one Cr
atom is present in the vicinity. In CCMN alloy, N
addition accelerates the γ→ε phase transition bymaking G of the ε phase higher than that of the γphase, but it also increases the energy barrier forIMR KINKEN Research Highlights 2012

the γ→ε phase transition by the formation of Cr–N
clusters. Therefore, the stability of the γ phase iscontrolled by these two opposite and competitive
processes, and the energy barrier resulting from the
Cr–N cluster formation appears to have a greater
effect on the phase transition since the γ phase iseffectively stabilized by N addition in CCMN alloy.
Biomedical Co–Cr–Mo (CCM) alloys have beencommonly used for artificial hip and knee joint
prostheses, but the need to improve their biomedical
inertness and wear resistance has become widely
recognized. The mechanical behaviors of CCM
alloys are dominated by strain-induced martensitic
transformation (SIMT), which causes crack initiation
during plastic deformation but enhances the wear
resistance dramatically in practical use. To develop
more reliable CCM alloys, it is essential to clarify thefactors affecting the occurrence of SIMT. In thepresent study, we investigated the effect of annealing
twin boundaries (ATBs) on SIMT behavior.
We analyzed in detail the substructures near aparallel pair of ATBs after deformation under a stressthat was preferential towards a slip parallel to the
ATBs. Preferential formation of the ε phase at the
ATBs was found by both SEM-EBSD analysis andTEM observation. A phase-field simulation was
conducted by utilizing our original code developed
for studying Suzuki segregation [2]. HRTEM images
indicated that the thickening of the ε phase does notproceed irregularly, leaving a high-density ofstacking faults. In addition, the thickness of the εphase was found to be different even at ATBs on
both sides of the twin. This difference was attributed
to the internal stress due to strain incompatibility at
the ATBs on the basis of a residual stress analysis
by the EBSD–Wilkinson method. Fig. 2 shows theresult of the EBSD–Wilkinson analysis. Fig. 2 a, b, c,and e are an image quality (IQ) map, kernel average
misorientation (KAM) map, phase map, and inverse
pole figure (IPF) map. SIMTed ε phase is clearlyobserved on the upper ATB. IQ and KAM aremeasures of plastic strain, and their maps indicatethat plastic strain is relatively high in the region
immediately above the SIMTed ε-phase. Fig. 2 eshows the distribution of internal shear stress for
Shockley partial dislocation slips. The thickest
ε-phase plate (dt ~ 1.5 µm) is formed in the region
where both the external applied stress and internal
stress promote the SIMT, and thinner ε plates (dt ~Fig. 2. SEM-EBSD micrographs of the twin-containing
Co–Cr–Mo alloy. (a) Image quality (IQ) map, (b) kernel
average misorientation (KAM) map, (c) phase map, (d)inverse pole figure (IPF) map, and (e) shear residual
stress for (1-1-1)[-11-2]M leading partial slip.

0.5 µm) were formed in the region where the external
applied stress promotes the SIMT but the internal
stress suppresses it. Interestingly, very thin ε plates(dt < 0.1 µm) were formed even in the region within
the twin, near the boundary, despite the suppressive
external stress for the SIMT and the absence of aslip trace parallel to the ATB. This is attributed to theinternal stress. Thus, the effects of ATBs on SIMT
have been clarified. Based on this result, morereliable biomedical CCM alloys will be developed.
References[1] Yunping Li, Jinshan Yu, Shingo Kurosu, Yuichiro Koizumi,
Hiroaki Matsumoto, Akihiko Chiba, "Role of Nitrogen Addition in
Stabilizing the γ Phase of Co-29Cr-6Mo Alloy", Materials
Chemistry and Physics 133, 29 (2012)[2] Y. Koizumi, T. Nukaya, S. Suzuki, S. Kurosu, Y.P. Li, H.
Matsumoto, K. Sato, Y. Tanaka and A. Chiba, Acta Materialia
60, 2901 (2012).

Keywords: Co-Cr-Mo alloy, mechanical properties, phase transformation
Akihiko Chiba, Yuichirou Koizumi, Yumping Li, Hiroaki Matsumoto (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 201212

Infrastructural Materials

Development of Ni-free Ti-based Metallic Glasses andComposites with Potential for Biomedical ApplicationsT

he effects of minor additions on the glass-forming ability (GFA), thermal stability, and mechanical andcorrosion properties of Ni-free Ti-based (Ti-Zr-Cu-Pd) metallic glasses were investigated. The substitution of
Cu by 2% Sn significantly improved the GFA. Minor additions of Nb, Ta, and noble metal elements improved
the plasticity and corrosion resistance. Using a spark plasma sintering process, large-size Ti-based bulk
metallic glasses, composites, and porous glassy alloys having the approximate Young’s modulus of bone weredeveloped.

13

Ti-based bulk metallic glasses (BMGs) arepromising materials for functional applications suchas biomedical materials due to their high corrosion
resistance, excellent mechanical properties, andgood biocompatibility. Many Ti-based BMGs havebeen developed in the framework of the Ti–Ni–Cu
and Ti–Zr–Cu–Ni alloy systems. However, these Tibased BMGs contain Ni, Be, etc., which are notsuitable for contact with the human body becauseof cellular toxicity, limiting their applications inmedical fields. Recently, we developed Ni-free Ti–
Zr–Cu–Pd BMGs with high strength and good
corrosion resistance that make it possible to create
novel Ti-based BMG implants [1]. However, arelatively low glass-forming ability (GFA) (with critical
diameter of 7 mm) restricted the biomedical
applications. A large critical diameter and excellent
mechanical properties of BMGs are the fundamental
requirements for these applications. Minor additions
are fundamental to controlling the formation,
manufacture, and properties of metallic materials bycontrolling nucleation during solidification. In thisstudy, we investigated the effect of minor additions
on GFA, thermal stability, and the mechanical andcorrosion properties of the Ti–Zr–Cu–Pd glassy alloy
system [2,3].
The substitution of Cu by 2% Sn significantly
improved the GFA. A Ti40Zr10Cu34Pd14Sn2 bulk

glassy alloy rod with a diameter of 12 mm has been
fabricated by a copper mold casting technique (Fig.
1). The addition of 2–4% Sn enlarged the supercooled
liquid region, indicating good thermal stability.
TiZrCuPdSn bulk glassy alloys exhibited highcompressive strengths of about 2000–2050 MPa
[2].
We
also
implanted
theTi-based(Ti40Zr10Cu34Pd14Sn2) BMG bars under the skin andin the bone of rats. The Ti-based BMG bars exhibited
excellent biocompatibility in both soft and hardtissue. They also showed nice osteoconductivity
when implanted in bone tissue, and no metal ion
diffusion was found up to one month after
implantation.
High strength and distinct plastic strain wereobserved in the stress–strain curves for Nb-added
Ti–Zr–Cu–Pd alloys. A yield strength exceeding
2050 MPa, low Young’s modulus of about 80 GPa,and distinct plastic strains of 6.5% and 8.5%
corresponding to serrated flow sections wereobtained for the 1% (Fig. 2) and 3% Nb-added
alloys, respectively. These minor additions of Nb or
Ta also exhibit higher corrosion resistance, as shownin Fig. 3. The minor addition of noble metal elements(Au, Pt) can significantly improve the plastic stain of
the Ti–Zr–Cu–Pd glassy alloys (Fig. 2) due tonanoparticles dispersed in the glassy matrix
blocking the propagation of shear bands [3].

Fig. 1 XRD patterns of TiZrCuPdSn alloys.

Fig. 2 Compressive stress–strain(Ti40Zr10 Cu36Pd14)99M1 alloys.

IMR KINKEN Research Highlights 2012

curvesoftheIMR KINKEN Research Highlights 2012

spark plasma sintering a mixture of gas-atomized
Ti-based glassy alloy powders and solid salt (NaCl)
powders, followed by a leaching treatment into waterto eliminate the salt phase. Pores were
homogeneously distributed in all of the sinteredsamples. The porosity can be controlled bycontrolling the volume fraction of the added salt
phase. In this way, high-strength porous Ti-based
BMGs with low Young’s modulus, equal to that ofbone, can be obtained.

Fig. 3 Potentiodynamic polarization curves of the(Ti40Zr10 Cu36Pd14)99M1 alloys at 310 K. The curves ofpure Ti and Ti–6Al–4V alloy are also shown forcomparison.

It is known that BMGs are commonly produced
using solidification techniques such as copper moldcasting or water quenching. However, when using asolidification technique, a rather high cooling rate isrequired to suppress the formation of more
thermodynamically stable crystalline phases. The
sample size and shape are seriously limited, thereby
limiting the range of applications. On the other hand,by using a consolidation technique, it is possible toproduce larger metallic glassy alloy parts in a greater
variety of shapes than by using a solidification
technique. Spark plasma sintering (SPS), developed
as a rapid consolidation technique, has great
potential for producing large-size glassy samples ina short sintering time [4,5].
We prepared Ti-based metallic glassy powders
using a high-pressure argon gas atomization
method. Using the glassy alloy powders, or theirmixed powders blended with hydroxyapatite (HA)
powder, we produced Ti-based BMGs andcomposites with diameters over 15 mm by the SPS
process [6,7]. The HA particles were uniformly
distributed in the glassy matrix. No crystallization of
the glassy matrix was observed. The thermal stability
was almost independent of the addition of the HA
particles. The produced composites exhibited lower
Young’s moduli than those of the as-cast Ti-based
bulk glassy alloys.
Furthermore, we also produced porous Tibased BMGs (Fig. 4) with diameters over 15 mm by

Fig. 4 SEM images of sintered porous Ti-based(TiZrCuPdSn) samples, after leaching NaCl, with various
porosities produced by the spark plasma sintering
process at 643 K, 600 MPa, 10 min.

References[1] S.L. Zhu, X.M. Wang, F.X. Qin, M. Yoshimura and A. Inoue,
Mater. Trans., 48, 2445 (2007).
[2] S.L. Zhu, G.Q. Xie, F.X. Qin, X.M. Wang and A. Inoue, Mater.
Trans., 53, 500 (2012).
[3] F.X. Qin, G.Q. Xie, Z.H. Dan and A. Inoue, Applied Mechanics
and Materials, 148-149, 241 (2012).
[4] G.Q. Xie, D.V. Louzguine, H. Kimura and A. Inoue, Appl. Phys.
Lett., 90, 241902 (2007).
[5] G.Q. Xie, D.V. Louzguine, H. Kimura, A. Inoue and F. Wakai,Appl. Phys. Lett., 92, 121907 (2008).
[6] G.Q. Xie, F.X. Qin, S.L. Zhu and A. Inoue, Intermetallics, 29, 99(2012).
[7] S.L. Zhu, X.J. Yang and Z.D. Cui, Intermetallics, 19, 572(2011).

Keywords: metallic glass, biomedical, powder metallurgy
Guoqiang Xie (Advanced Materials Development and Integration of Novel Structured Metallic and Inorganic Materials)
E-mail: xiegq@imr.tohoku.ac.jp
URL: http://bmg-pro.imr.tohoku.ac.jp/index.html

IMR KINKEN Research Highlights 201214

Energy-related Materials

Control of Domain Size and Doping of Barium Disilicide
Epitaxial Thin Films for Earth-abundant High-efficiency
Solar Cells

Barium disilicide (BaSi ) is regarded as a promising candidate for a novel earth-abundant material for high2efficiency thin-film solar cells owing to its large absorption coefficient and ideal bandgap. We attempted torealize high-quality BaSi2 epitaxial thin film crystals with large crystal domains and controlled doping. Intentional
introduction of a miscut to the Si(111) substrate selectively enlarged one of the three equivalent epitaxial variants
by modifying the nucleation sites; this led to extraordinarily large crystal domains. Ion-implantation of BF2 andsubsequent annealing were found to be useful to obtain p-type BaSi2 with a carrier concentration of around1018 cm-3.
Orthorhombic barium disilicide (BaSi2) has
been proposed as a promising, novel candidate forhigh-efficiency thin-film solar cells owing to its large
absorption coefficient of 105 cm-1 at 1.5 eV [1] andbandgap that is tunable to 1.4 eV by substituting halfof the Ba sites with isoelectronic Sr atoms [2]. Abase layer only 1 μm thick is considered sufficient toabsorb all the photons with energies higher than the
bandgap. Importantly, all the constituent atoms areearth-abundant and suitable for the global
deployment of photovoltaic technology.
To fabricate a solar cell based on BaSi2, it is of
crucial importance to realize high-quality thin film
crystals with controlled doping. So far, epitaxial
BaSi2 thin films with (100) orientations have beengrown on Si(111) substrates. However, the epitaxial
relationship allows formation of three epitaxial
variants that are equivalent in terms of 60° in-planerotations. As a consequence, many crystal domains
and lattice defects exist in the films. A possible route
toward realization of high-quality BaSi2 with large
crystal domains is to reduce the symmetry of thesurface to selectively grow one of the epitaxial
variants.
For this purpose, we investigated the impact ofa miscut given to the Si(111) substrate on the
microstructures of epitaxially grown BaSi2 thin films[3]. BaSi2 epitaxial films were prepared by reactive

deposition epitaxy followed by molecular beam
epitaxy. First, Ba atoms were deposited on theSi(111) substrate at a deposition rate of 0.9–1.0 nm/
min for 5 min, thereby producing a thin template
BaSi2 layer. An a-axis-oriented BaSi2 epitaxial filmwith a thickness of 80–90 nm was formed
subsequently by depositing Ba and Si atoms on thetemplate layer. On-axis and vicinal (with a 2° miscut
towards the [112] direction) Si(111) substrates wereused.
Fig. 1 compares the electron backscatter
diffraction (EBSD) orientation maps of the BaSi2
films on (a) on-axis and (b) vicinal Si(111) substrates.
In Fig. 1(a), it is seen that three epitaxial variants
randomly appear and that the domain shape iscomplicated. The size of most domains is less thanseveral micrometers. On the other hand, one of thethree epitaxial variants was dominant when the
BaSi2 film was grown on the 2°-inclined substrate,
as shown in Fig. 1(b). In the dominant variant, the baxis was found to be parallel to the miscut
direction.
Atomic force microscopy clarified that, on the
2°-inclined substrate, barium silicide platelets arefound next to the step-bunched regions and the
proportion of one epitaxial variant is increased. This
result suggests that lattice matching with the stepedge is an important factor for orientation selection.

Fig. 1 EBSD crystal orientation maps of BaSi2 epitaxial films grown on (a) on-axis and (b) 2°-inclined substrates.

15

IMR KINKEN Research Highlights 2012

IMR KINKEN Research Highlights 2012(c)

Intensity (arb. unit)

Assuming that the steps are strictly straightforward,
the step edge is parallel to the Si<110> direction.
The BaSi2[001] is accordingly parallel to Si<110>.
The repeating-unit length of BaSi2[001] is 0.12%
smaller than three Si<110> repeating-unit lengths.
The other two orientations differ by 60° around the
BaSi2 a axis from the large-domain orientation. In
such orientations, the BaSi2[031] direction is almost
parallel to Si<110>; the angle between these
directions is 0.25°. The repeating-unit length of
BaSi2[031] is 0.88% smaller than six Si<110>repeating-unit lengths. Thus, the lattice mismatch tothe step edge for the small-domain orientations islarger than that for the large domain. The orientationof the large domain is, therefore, energetically morefavorable than the orientations of the small domains.
This would explain why the large domain is produced
when the 2°-inclined Si(111) substrate is used.
In regards to doping, both n-type and p-type
BaSi2 thin films have been obtained by in-situ doping
during molecular beam epitaxy. The electron
concentration of n-type BaSi2 has been widely
controlled in the range of 1016 –1020 cm-3. On theother hand, the carrier concentration in p-type BaSi2
is still limited to the order of 1016 –1017 cm-3. It istherefore desired to discover a p-type dopant thatbrings about a high carrier concentration of around1018 cm-3 or more.
Boron is expected to be a p-type dopant, as is
the case for the other group-III atoms Al and In. We
adopted BF2 ion implantation and subsequentannealing [4], since it is time-proven to produce
p-type Si. In general, to recover the damaged layer,
annealing at (2/3)Tm is necessary, where Tm is themelting point. Since the melting point of BaSi2 is1180 °C, annealing at 696 °C is necessary.
Fig. 2 shows the Raman spectra of 120 nm
BaSi2 films implanted with BF2 ions at doses of (a)1.5×1012 cm-2, (b) 1.5×1013 cm-2, and (c) 1.5×1014
cm-2. The spectra of the as-implanted samples andthose after rapid thermal annealing (RTA) at 800 °C
for 30 s are shown in solid and dotted lines,
respectively. The spectrum of the as-implanted film
agrees with that reported for BaSi2 [5]. According to
the previous report, the vibrations with wave numbers
above 250 cm-1 can be interpreted based on the Td
symmetry of the discrete [Si4]4- anion. The full widthat half maximum (FWHM) of the Raman peak at
485–490 cm-1 increased with increasing dose,

(b)

(a)

300

400

500
−1

Raman shift (cm )Fig. 2 Raman spectra of BaSi2 films with BF2 implantation
doses of (a) 1.5×1012, (b) 1.5×1013, and (c) 1.5×1014,cm-2.
Solid and dotted lines show the spectra of the asimplanted and annealed samples, respectively.

indicating that the as-implanted samples containdamage. After RTA, a Si peak at 520 cm-1 is observedthat originates from the oxidation of BaSi2, whichmust be somehow suppressed. Importantly, the
FWHM value decreases to the as-grown level,
indicating that the implantation damage is removed
almost completely by RTA. Furthermore, Hallmeasurement revealed that the majority carriers in
the B-doped BaSi2 films are holes and that theirconcentration reaches 2×1018 cm-3 at the highest
implantation dose of 1.5×1014 cm-2. This is, so far,the highest observed carrier concentration in p-type
BaSi2. We have thus revealed that B is a promising
acceptor dopant for BaSi2.
References[1] K. Morita, Y. Inomata and T. Suemasu, Thin Solid Films 508,
363 (2006).
[2] K. Morita, M. Kobayashi and T. Suemasu, Jpn. J. Appl. Phys.
45, L390 (2006)[3] K.O. Hara, N. Usami, K. Toh, K. Toko and T. Suemasu, Jpn. J.
Appl. Phys. in press.
[4] K.O. Hara, N. Usami, Y. Hoshi, Y. Shiraki, M. Suzuno, K. Tokoand T. Suemasu, Jpn. J. Appl. Phys. 50, 121202 (2011).

Keywords: semiconducting, solar cells, ion implantation
Noritaka Usami (Crystal Physics Division)
E-mail: usa@imr.tohoku.ac.jp
URL: http://www.xtalphys.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201216

Energy-related Materials

Dynamical Generation of Spin CurrentsG

enerating spin currents in nonmagnetic materials is crucial for exploring spin physics in condensed matter.
We demonstrate that dynamical spin current generation driven by magnetization precession offers an easy andversatile pathway for spin injection into semiconductors. This dynamical spin injection method enables efficient
spin injection into high resistivity materials, opening a route for accessing spin–charge conversions in a wide
range of materials.
Spintronics is a promising technology in which
the spin states of electrons are utilized as anadditional degree of freedom for device operation[1]. There has been significant progress in recent
years in spin injection into semiconductors, whichuses tunnel barriers to circumvent the impedancemismatch problem; the impedance mismatch
between ferromagnetic metals and high-resistivity
materials drastically limits the spin-injection
efficiency [2,3]. However, because of this problem,

microwave
V

Ni81Fe19

M(t)

＋
GaAs

－
σ

V( θ = 0) (nV)dI(H)/dH (arb. unit)

EISHE

H

js

1
0

V

＋

-1

－

H

200
Ni81Fe19
100 /n-GaAs

Ni81Fe19/p-GaAs

0
-20

-10

0
10
H − HFMR (mT)20

Fig. 1 Spin pumping in a NiFe/GaAs junction. I and V are
the microwave absorption intensity and electromotive
force, respectively. H is the external field.

17

IMR KINKEN Research Highlights 2012

there is no route for spin injection into semiconductors
through low resistivity interfaces, i.e., ohmic contacts,
even though it promises an easy and versatilepathway for spin injection into high-resistivity
materials.
The impedance-mismatch problem arises fromthe fact that spins are injected by carrier transport
across a ferromagnetic metal/semiconductor
interface. Thus, it seems natural to consider that thisproblem disappears when spins are injected directly
into the semiconductor layer without using chargetransport across the interface; the driving force forthe spin flow is expected to offer a way for versatile
spin injection free from the impedance-mismatch
problem. In this work, we experimentally demonstrateefficient spin injection into semiconductors throughohmic contacts using dynamical spin injection: spin
pumping. Spin pumping, the generation of pure spincurrents from magnetization precession [4], provides
a powerful way for direct spin injection. The spin
angular momentum of the precessing magnetization
in a ferromagnetic metal layer is transferred to thecarriers in the semiconductor layer via dynamical
exchange interactions at the ferromagnetic metal/
semiconductor interface, inducing a pure spin
voltage (the potential acts on spins, not on charges)
in the semiconductor layer. This enables spin
injection into both p- and n-type GaAs from NiFe
through both ohmic and Schottky contacts in a NiFe/
GaAs interface even at room temperature (see Fig.
1). Furthermore, as shown in Fig. 2, we demonstratethat the spin exchange interaction can be controlled
electrically by applying a bias voltage across the
NiFe/GaAs interface, enabling electric tuning of thespin pumping [5].
Dynamical spin injection is a powerful method
for exploring the physics of spin currents in highresistivity materials. A spin current is coupled with acharge current through the spin–orbit interaction inIMR KINKEN Research Highlights 2012

EF

eVin

EF

φS
Vin < 04

Vin > 0Ni81Fe19/p-GaAs
PMW = 200 mW

100
90
80
70
60
50
40
30
20
10
0

200 mW
190
180
170
160
150
140
130
120
110

3

F

V (µV)

gr↑↓ (1017 m-2)5

Vin = 0eVin E

2

1

3
2

0
-0.6 -0.3

00.3
Vin (V)

0.6

Fig. 2 Electric tuning of spin exchange and spin pumping

90

100 110H (mT)

120

130

Fig. 3 Observation of the inverse spin Hall effect in Si.

at a Schottky NiFe/GaAs interface.

solids. This coupling gives rise to the conversion ofspin currents into charge currents (the inverse spin
Hall effect) providing a route for electric detection ofspin currents [6]. However, so far, the inverse spin
Hall effect has only been accessible in a limited
number of materials, excluding a wide range ofindirect-bandgap semiconductors and precluding
further development. Notable in this group is Si,
currently the most fundamental material in electronic
chips. Accessing the inverse spin Hall effect in Si
has been believed to be difficult because of the very
weak spin–orbit interaction. Here, we show thatdynamical spin injection allows access to the inverse
spin Hall effect in Si at room temperature thanks tothe ability to generate giant spin currents in amacroscopic area. In spite of the small spin–orbit
interaction in Si, we found a clear electric voltage
due to the inverse spin Hall effect in a p-type Si film,
as shown in Fig. 3 [7], demonstrating that Si can beused as a spin current detector.
References[1] I. Zutic, J. Fabian, and S. Das Sarma, Rev. Mod. Phys. 76, 323(2004).
[2] G. Schmidt, D. Ferrand, L. W. Molenkamp, A. T. Filip, and B. J.
van Wees, Phys. Rev. B 62, R4790 (2000).
[3] E. I. Rashba, Phys. Rev. B 62, R16267 (2000).
[4] Y. Tserkovnyak, A. Brataas, and G. E. W. Bauer, Phys. Rev.
Lett. 88, 117601 (2002).
[5] K. Ando, S. Takahash, J. Ieda, H. Kurebayashi, T. Trypiniotis,
C. H. W. Barnes, S. Maekawa, and E. Saitoh, Nature Mater. 10,655 (2011).
[6] E. Saitoh, M. Ueda, H. Miyajima, and G. Tatara, Appl. Phys.
Lett. 88, 182509 (2006).
[7] K. Ando and E. Saitoh, Nature Commun. 3, 629 (2012).

Keywords: spintronics, spin current
Kazuya Ando, Eiji Saitoh (Surface and Interface Research Division)
E-mail: ando@imr.tohoku.ac.jp
URL: http://saitoh.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201218

Energy-related Materials

Novel Material Science Using Neutron ScatteringI

MR is a unique institute, even worldwide, that has a neutron facility based on a strong background in advanced
material science. The IMR neutron scattering group is investigating issues in materials science, such as
ferroelectric materials, high-TC superconductivity, hydrides, and application materials, as well as developing
high-performance devices in neutron science. In particular, our group is aiming to construct an originative
neutron spectrometer with next-generation technologies in J-PARC, which will be a flagship instrument ofmaterial science in IMR.
1) Anomalous Spin Response in Metallic
Antiferromagnets
Recent inelastic neutron-scattering experiments
performed on Fe-based superconductors [1] give aclose-up view of the spin dynamics of spin-density
waves (SDW) in metallic antiferromagnets, which
would be relevant to high-Tc pairing formation. Mn3Si
and Mn2.8Fe0.2Si are antiferromagnetic (AF) metals
with TN of about 20 K. The AF magnetic fluctuations
show a steeply dispersive relation (or a continuum)[2]. These behaviors appear similar to those in the
SDW state of the AF metal Cr [3].
We succeeded in observing magnetic
excitations at a wide Q–ω phase using a large single
crystal of Mn2.8Fe0.2Si on a time-of-flight chopper
spectrometer (SEQUOIA) at the SNS of Oak Ridge
National Laboratory in the U.S.A. Salient data areshown in Fig. 1. While AF low-energy spin-wave
excitations are located around L positions [Fig. 1(b)],the magnetic symmetry in Q space is different above
50 meV and the cross section concentrates at some(not all) zone-center Γ positions [Fig. 1(a)], i.e., at(2,2,2) and (2,0,0) but not at (0,2,2) or (4,0,0).
Moreover, as shown in Fig. 1(c), such a dynamical
response at Γ points persists vertically to more than
100 meV in the magnetic dispersion.
The qualitative change upon ω in the magnetic

excitations is distinct, but no reasonable explanation
has been established yet. Further studies on thesecomplicated spin dynamics are now in progress toreveal the spin correlations in magnetic metals,
where localized spins and itinerant-electron spins
coexist.
(H. Hiraka)2) Impurity Effect in High Tc Superconductors
The role of magnetism in high transition
temperature (Tc) superconductors has been argued,
since the parent compounds of high-Tc cuprates are
antiferromagnetic (AF) Mott insulators. Studying theeffect of impurities in Cu sites is an effective way toassess the correlation between magnetism andsuperconductivity, because small amounts ofimpurities can induce remarkable changes in bothsuperconductivity and magnetic properties withoutmajor changes to the bulk lattice framework. It is
known that a non-magnetic Zn impurity suppresses
superconductivity more severely than a magnetic Ni
impurity, which is in contrast to the impurity effects
on BCS superconductors. Studies on μSR haveelucidated the microscopic picture for Zn doping;
Zn locally kills superconductivity and Tc issuppressed by the reduction in superconducting
carriers. However, there exists no microscopic
50

E comme (meV)

40
30
20
10
0
0

10

20

30

40

T c (K)
Fig. 1 Neutron intensity contours using a single crystal of
Mn2.8Fe0.2Si measured by a chopper-type time-of-flight
spectrometer. (a, b) Q-slices and (c) ω–Q slice.

19

IMR KINKEN Research Highlights 2012

Fig. 2 A plot of Ecomme versus Tc upon Ni-doping. The
horizontal error bars represent broad superconducting
transition, which is characteristic to Ni-doped high-Tc
cuprates [ref 2-1]

IMR KINKEN Research Highlights 2012

explanation for the weak suppression by Ni doping.
To solve this issue, we have performed inelastic
neutron scattering (INS) measurements on single
crystals of La1.85Sr0.15Cu1-yNiyO4 over a wide doping
range from y = 0.009 to 0.04.
The INS studies elucidate the anomalousenergy and Ni dependences of the broadening of
low-energy incommensurate peaks, which isevidenced by the appearance of a commensurate
peak at the AF zone center (π,π) and finite energy
Ecomme. Fig 2 shows a plot of Ecomme vs Tc upon Ni
doping. With increasing Ni content, Tc decreases
from 36.8 K (y = 0) to 22.9 K (y = 0.015), 11.6 K (y =0.029), and 0 K (y = 0.04). The horizontal bars
represent the transition width of ΔTc determined fromthe shielding signal of the magnetic susceptibility
measurement. We have clarified that Ecommedecreases in proportion to Tc. The linear relationshipbetween Ecomme and Tc explains the weak
suppression of Tc for Ni doping by reduction in thecharacteristic energy of spin fluctuations.
(M. Matsuura)
3) Vibrational Properties of Hydrogen Storage
Materials
Hydrogen has been identified as an alternative
clean energy carrier for mobile applications. For
instance, hydrogen storage technologies are onenecessity for the widespread use of hydrogen. Since
the discovery of the reversible hydrogen storage
properties of Ti-doped NaAlH4 [4], Al-based
hydrides have been extensively studied bothexperimentally and theoretically [5]. In order to gain
further insight into their hydrogen storage properties,
it is important to investigate the vibrational propertiesof a series of Al-based hydrides. Inelastic neutron
scattering (INS) can be a powerful technique
because hydrogen has a large inelastic neutron
scattering cross-section (82.02×10 -24 cm2) that is
approximately 10 times as large as those of mostother elements.
Using first-principles calculations, we elucidated
the vibrational properties of Al-based hydrides withdifferent AlH6 networks: Na3AlH6 with isolated AlH6
units, CaAlH5 with a one-dimensional AlH6 network,and α-AlH3 with a three-dimensional AlH6 network(Fig. 3) [6]. The INS spectrum of Na3AlH6 was mainly
divided into four regions: (i) the translational modesat < 218 cm-1, (ii) the librational modes of the AlH6

Fig. 3 Contour plot of INS spectra on Na3AlH6, CaAlH5
and α-AlH3. The reft side indicates AlH6 network foreach structure. Yellow and blue circles are Al and H
atoms, respectively.

units at 339–515 cm-1, (iii) the H–Al–H bond-bending
modes at 715–994 cm-1, and (iv) the Al–H bondstretching modes at 1166–1534 cm-1. However, thespectra of both CaAlH5 and α-AlH3 were mainly
divided into three regions because there were nogaps between regions (ii) (the external motions) and(iii) (the intramolecular motions) due to the networked
nature of the AlH6 units. Although both CaAlH5 andα-AlH3 showed qualitatively similar INS spectra, thelibrational modes of AlH6 on α-AlH3 exhibited higher
frequencies than CaAlH5 because of tighter bonding
between the AlH6 units.
(T. Sato)
References[1] For example: M.D. Lumsden et al., Nature Phys. 6, 183(2010).
[2] S. Tomiyoshi et al., Phys. Rev. B 36, 2181 (1987).
[3] Y. Endoh and P. Boni, J. Phys. Soc. Jpn. 75, 111002 (2006).
[4] B. Bogdanović and M. Schwickardi, J. Alloys Compd. 253254 (1997) 1.
[5] S. Orimo, et al., Chem. Rev. 107 (2007) 847.
[6] T. Sato, et al., Inorg. Chem. 50 (2011) 8007.

Keywords: neutron scattering, spin lattice dynamics, novel magnetic materials
Kazuyoshi Yamada, Kenji Ohoyama, Haruhiko Hiraka, Masato Matsuura, and Toyoto Sato
(Metal Physics with Quantum Beam Spectroscopy Division)
URL: http://www.yamada-lab.imr.tohoku.ac.jp/jp/index.html

IMR KINKEN Research Highlights 201220

Energy-related Materials

Atomistic Level Description of Phase Diagram of Clathrate
HydratesIn order to accurately estimate the thermodynamic properties of hydrogen clathrate hydrates, we developed amethod that allows one to take into account multiple cage occupancy, host lattice relaxation, and the descriptionof the quantum nature of guest behavior. Starting from first-principles estimation of the guest-host interaction
with a TOMBO approach, the thermodynamic properties of clathrate hydrates and the hydrogen concentration
at various pressures can be evaluated without recourse to any empirical parameter fitting. Thus, the present
methodology can support the experimental exploration of novel storage materials.
The reduction of the use of fossil fuels and theirreplacement with alternative and sustainable energy
sources is one of the present grand energy
challenges. Hydrogen is an ideal candidate for thissubstitution and, therefore, the technology of usinghydrogen as an environmentally clean and efficientfuel is an active research area worldwide. However,
it is important to overcome the many scientific andtechnological hurdles that exist between the present
state of hydrogen production, utilization, and storage
capabilities and those required for a competitiveand sustainable hydrogen economy.
Clathrate hydrate is one potential candidate forhydrogen storage. This is a guest–host, or inclusion,
compound in which water molecules form a threedimensional host structure where the molecules(guests) can be encaged in the empty voids. Recently,
it has been reported that the clathrate hydrate structure
II (see Fig. 1) can store around 4.96 wt% hydrogen at
220 MPa and 234 K. Following this work, manystudies of hydrogen hydrates using different methods
have been conducted. Much of this work focused on
the formation of hydrogen hydrates at lower pressuressince the high-pressure conditions for pure hydrogenhydrate formation place a limiting constraint on its
practical application. A reduction by nearly two ordersof magnitude in the hydrate formation pressure wasfound in the tetrahydrofuran (THF) – hydrogen – water
system. Thus, the possible formation of hydrogenhydrates with different structures by adding various
guest molecules suggests the practical feasibility of abinary hydrate as a hydrogen storage candidate.
Therefore, it is important to study the phase diagramsof multicomponent clathrate hydrates in detail, whichis still a challenging experimental task due to theircomplexity.
There are several types of gas hydrate structureswith different cage shapes, and some of thesehydrate structures can hypothetically store morehydrogen than the hydrate of structure CS-II. Taking
into account the same cage occupancy as in thecase of CS-II hydrate, larger values of hydrogenmass (6.33 wt%) can be achieved with the CS-I
structure than with the proposed hydrogen mass
(4.96 wt%) for the CS-II structure.
21

IMR KINKEN Research Highlights 2012

Fig. 1 Hydrogen CS-II clathrate hydrate. (Inset) Fragmentof CS-II hydrate structure in which the two fundamental
cages (512 and 51264) are connected directly.

From these points of view, theoretical
approaches can be useful for understanding thephysical and chemical properties of the binary
hydrates, and they can support the experimental
exploration of novel hydrogen storage materials
based on clathrate hydrates. Therefore, a novel
approach based on the solid solution theory of vander Waals and Platteeuw has been developed. Thismodel allows one to take into account the influenceof guest molecules on the host lattice and guest–
guest interactions, especially when more than one
guest molecule occupies a cage. We used thequasiharmonic lattice dynamics (QLD) method toestimate the free energies, equations of state, andchemical potentials [1]. The QLD approach includes
the quantum effects in the energy and entropy
calculations. These effects are significant in water
ice at low temperatures and no procedure has been
established to correct these effects within molecular
dynamic calculations. The short-range repulsion
part of the H2–H2 interaction was estimated using
the all-electron mixed-basis method (TOMBO code),which allowed us to better describe the interactions
between hydrogen molecules in water cavities.
Using this approach, the phase diagrams ofpure hydrogen and mixed H2 + C3H8 [2], H2 + CH4[3] and H2 + C2H6 [4] with multiple cage filling byhydrogen were constructed. The obtained results
were in agreement with known experimental data.

IMR KINKEN Research Highlights 2012

Moreover, a structural transformation in propane
hydrate was propose to explain the formation of CSIV hydrates with greater storage capacities thanthose observed in CS-II hydrate [5].
At low temperature, the mass percentage ofhydrogen in pure CS-II hydrogen hydrate can reachup to 3.8 wt% for pressures of 16–18 MPa. At higher
pressure, the hydrogen storage capacity canincrease to 4 wt%. Increasing the temperature to260 K reduces the amount of stored hydrogen. The
maximum amount of hydrogen stored at thistemperature and high pressure is about 3.5 wt%.
The introduction of propane as a second guest
component in the binary hydrogen hydrate not only
reduces the formation pressure but also reproduces
the hydrogen storage ability of the pure clathrate
hydrate [2].

to form both the CS-I and CS-II hydrates, and theCS-I hydrate with this guest is more stable than the
CS-II structure. Based on these conditions, themixed H2 + CH4 and H2 + C2H6 hydrates with CS-I
and CS-II structures have been studied with different
gas concentrations.
The stabilization of the CS-I hydrate can berealized for both the H2–CH4 –H2O and H2–C2H6 –H2O systems with a small concentration of methaneand ethane, respectively, as shown in Fig. 3.

Fig. 3 Pressure p(T) of monovariant equilibrium: gas phaseof CH4+H2 – ice Ih – CH4+H2 CS-I (squares) and CS-II(circles) hydrates at 6% methane concentrations in thegas phase.

Fig. 2 Pressure dependence of hydrogen storage capacityof pure CS-I and CS-II hydrates at T = 250 K.

Fig. 2 shows the hydrogen storage capacity ofhydrogen hydrates with CS-I and CS-II structures
as a function of pressure at T = 250 K. The
calculation results confirmed that, for pure hydrogenhydrate, the CS-I structure can store more hydrogenthan the CS-II structure. It can be seen that in low
pressure regions the CS-I structure can store up to
3.1 wt% at T = 250 K. However, it has been foundthat the CS-I structure is thermodynamicallyunstable compared to CS-II [4]. This result can
explain the experimental realization of the hydrogen
CS-II hydrate only.
In order to stabilize the CS-I structure, a second
guest component was selected according to thefollowing conditions. In pure form, this guest is able

The amount of hydrogen storage depends onthe concentration of the second component in thegas phase. Thus, at a low concentration of ethane,
2.5 wt% hydrogen storage can be achieved in the
CS-I hydrate at T = 250 K. However, for hydrogenstorage applications, realization of ethane–hydrogen hydrate with the CS-II structure ispreferable because, based on the recent
experimental analysis, 3.5 wt% of hydrogen storage
can be achieved by this structure [4].
References[1] R. V. Belosludov et al., in Handbook of Nanophysics, edited by
K. D. Satler, (CRC Press, Taylor&Francis Group, Boca RatonLondon-New York 2010). pp. 2-1 – 2-24.
[2] R. V. Belosludov et al., J. Chem. Phys. 131, 244510 (2009).
[3] V. R. Belosludov et al., Comput. Mater. Sci. 49, S187 (2011).
[4] R. V. Belosludov et al., Mol. Simul. 38, 773 (2012).
[5] M. Souissi et al., J. Incl. Phenom. Macrocycl. Chem. 69, 281(2011).

Keywords: clathrate hydrate, hydrogen storage, phase diagram
Hiroshi Mizuseki (Materials Design by Computer Simulation Division)
E-mail: mizuseki@imr.edu
URL: http://www.mdcs.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201222

Energy-related Materials

Study of Microstructural Changes in Thermally Aged
Stainless Steel Weld Overlay Cladding of Nuclear Reactor
Pressure Vessels by Atom Probe TomographyThe effect of thermal aging on microstructural changes in stainless steel weld overlay cladding of nuclear
reactor pressure vessels composed of austenite and δ-ferrite phases was investigated using atom probe
tomography. We found that thermal aging at 400 °C for 10,000 h caused a large fluctuation in the Cr concentration
due to spinodal decomposition as well as the precipitation of G phases in the ferrite phase. It was suggested
that the hardening by thermal aging was mainly caused by the spinodal decomposition of Cr.
Stainless steel weld overlay cladding is used asa protective barrier against corrosion on the innersurface of commercial water-cooled reactor pressure
vessels (RPVs). The cladding materials are subjected
to neutron irradiation and prolonged thermal aging
at a service temperature of about 300 °C. Thesethermal and irradiation environments may cause
microstructural changes, resulting in the degradationof mechanical properties and corrosion behavior.
Thus, a detailed analysis of the microstructural
evolution of the cladding is necessary to assure theintegrity of RPVs.
The cladding material is composed of about
90% austenite phase and 10% δ-ferrite phase with
net-like structures [1]. The mixed ferrite and austenitic
phases bring about enhanced corrosion resistanceof the claddings [2]. On the other hand, it is known
that the ferrite phase is significantly hardened during
thermal aging, giving rise to a toughness loss [3]. As
reported for duplex stainless steels (DSSs) [4], thespinodal decomposition of the ferrite phase into Ferich α and Cr-rich α’ phases, together with theformation of Ni-rich G phase and Cr-rich M23C6 -type
carbide, are considered to play a key role in thehardening. However, the detailed mechanisms ofthese changes in the weld overlay cladding and DSS
have yet to be understood because advanced
experimental tools have not yet been applied toreveal the fluctuation of chemical compositions withnanometer- scale spatial resolution.
In this work, we applied atom probe tomography(APT), a unique method of mapping the alloying
elements in three-dimensional real space with nearlyatomic scale resolution [5], to investigate thespinodal decomposition of Cr and the formation of
the G phase in the δ-ferrite in weld overlay claddings
subjected to thermal aging [6,7].
The material studied was stainless-steel overlay
cladding made by electroslag welding on A553B
steels. The as-received material was subjected topost-welding heat treatments (PWHT) at 615 °C for
7 h. The microstructure was composed of net-like δ
ferrite phases formed along some grain boundariesof the austenite matrix, as shown in Fig. 1. The
23

IMR KINKEN Research Highlights 2012

material was thermally aged at 400 °C for 10,000 h
as an accelerated aging for practical conditions.

30µm
Fig. 1 Microstructure of the overlay cladding with net-likeδ-ferrite phases observed by optical microscope.

The microstructural analysis was carried out
with a laser-pulse APT (LEAP 3000X-HR by
CAMECA) equipped with an energy- compensatingreflectron lens. The samples for the APT analysis
were first electrochemically polished, and finallysharpened by a dual-beam FIB/SEM ion millingsystem in order to select the ferrite phase at the tipof the needle specimens.
Fig. 2 shows a comparison of several atom
maps of the as-received and thermally aged ferrite
phases. The atom maps were constructed fromslices 8 nm in thickness to clearly illustrate detailed
chemicalcomposition
fluctuations.
Largerfluctuations in the Cr and Fe concentrations overnano-scale areas were observed in the aged ferrite
phase than in the as-received one. In the aged ferrite
phase, Si, Mn, P, Ni, and Cu-concentrated
precipitates (G precipitates) were also observed.
The Cr concentration fluctuated between about
20% and 35%, and the wavelength of the fluctuation
was estimated to be about 8–10 nm for the asreceived phase. The fluctuation in the Crconcentration was magnified by the aging to
between about 10% and 45%, while the wavelengthof the fluctuation, estimated to be around 10 nm,

IMR KINKEN Research Highlights 2012

remained almost unchanged. These findings indicatethat the fluctuation in the Cr concentration is causedby spinodal decomposition in both the ferrite phases
subjected to the cooling process after the PWHT andprolonged thermal ageing.

Mn

C

P

100

Fe

Si

Concentration [%]

Fe

80

Si

P

Ni

Cr

Mo

Cu

Fe
Si
Ni
Mn
Cr
P

60
40

Mn20

Ni

C

0
0

50
100
Distance [nm]150

Fig. 3 1D concentration profiles penetrating the G
precipitates at the ferrite phase in the aged material.

30nm

Cr

Mo

Cu

Finally, we found a linear relation between theferrite hardening and a scalar parameter representing
the spinodal decomposition for the as-received andaged ferrite phases. It was suggested that thehardening by thermal aging was mainly caused by
the spinodal decomposition of Cr.

(b)

Fe

Si

C

Si

References[1] J. A. Brooks, A.W. Thompson, Inter. Mater. Rev. 36, 16 (1991).
[2] I. Alvarez-Armas, S. Degallaix-Moreuil, Duplex Stainless Steels,
Wiley-ISTE, New York, 2009.
[3] F.M. Haggag, R.K. Nanstad, NUREG/CR-6363 ORNL/TM-

Ni

u

(b)

(a)

Fig. 3 shows the 1D concentration profiles atand near the G precipitates in the aged ferrite phase.
The chemical composition at the central part of thelarge precipitate was 53% Ni, 23% Si, 20% Mn, 1%
Fe, Cr, P, and slight Cu and Mo. The ratios of Ni, Si,
and Mn were Ni:Si:Mn = 16:7:6. This type ofprecipitate is called a G phase [4]. Although thesmall precipitates contained more Fe and Cr thanthe large one, the ratios of atoms were almost
Ni:Si:Mn = 16:7:6 as well.

Mn

P

Ni

13047 (1997).
[4] N.M. Chung, O.K. Chopra, International symposium on
environmental degradation of materials in nuclear powersystems, Traverse City, MI, USA, 30 Aug 1987.
[5] M. K. Miller, Atom Probe Tomography, Kulwar Academic/
Plenum Publishing, New York, NY, 2000.

30nm

Cr

Mo

Cu

Fig. 2 Solute atom maps of the ferrite phase in the asreceived material (a) and in the aged material (b).

[6] T. Takeuchi, J. Kameda, Y. Nagai, T. Toyama, Y. Nishiyama, K.
Onizawa: J. Nucl. Mater. 415, 198 (2011).
[7] T. Takeuchi, J. Kameda, Y. Nagai, T. Toyama, Y. Matsukawa, Y.
Nishiyama, K. Onizawa: J. Nucl. Mater. 425, 54 (2012),

Keywords: nuclear materials, steel, atom probe tomography
Yasuyoshi Nagai (Irradiation Research on Nuclear Materials Division)
E-mail: nagai@imr.tohoku.ac.jp
URL: http://wani.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201224

Energy-related Materials

Behaviors of Hydrogen Isotopes in Silicon Carbide for
Evaluation of Hydrogen Permeability and Retention in
Nuclear Fusion Systems

The behaviors of hydrogen isotopes in silicon carbide were studied after hydrogen isotopes were injected into
the silicon carbide (SiC) by exposure to a hydrogen plasma from the fusion plasma machine of GAMMA-10,
irradiation in the high flux fission reactor of HFIR with lithium containing oxides (tritium injection through (n, α)
nuclear reaction of lithium), and ion accelerator irradiation with hydrogen and deuterium.
The behaviors of hydrogen isotopes in siliconcarbide (SiC) were studied after hydrogen isotopes
were injected into the SiC by exposure it to ahydrogen plasma from the fusion plasma machineof GAMMA-10, irradiation in the high flux fission
reactor of HFIR with lithium containing oxides (tritium
injection through (n, α) nuclear reaction of lithium),and ion accelerator irradiation with hydrogen anddeuterium. The GAMMA-10 exposure showed ahigh retention of tritium in the SiC of about 5×1020 H/
m2 even after a short exposure period of about 30
seconds. The HFIR irradiation revealed that theretention of tritium in the SiC is substantial and themobility of tritium is limited even at 800 °C [12,3],
while the accelerator irradiation with 10 keV
deuterium showed that deuterium has a good
mobility at about 800 °C [4]. The comparison among
the data obtained shows the importance of thetrapping of hydrogen isotopes by radiation-induced
defects. Silicon carbide has some attractive material
properties, such as a low Z, low induced radioactivity,
high strength, good radiation resistance, andrefractoriness. It is a strong candidate for nuclear
fusion materials in several applications, such as
thermal and tritium barriers, and as a plasma facing
material. There, the mobility and the retention

54

2.8 MeV MeV ERD

PSL intensity/ a. u.

200

HRecoil yield

SiC single crystal
100

exposedas prepared

Surface in contact
with breeder
SiC with LAO

0

100

200

300

400

500

2
Back face0

Channel number

Bulk region

0.5
1.0Distance/ mm

Fig. 1 Hydrogen concentration in the SiC exposed to theFig. 2 Tritium profile in the SiC irradiated in the HFIR at 800

plasma from GAMMA-10. (Exposure to 28 shots of 0.39

°C. The cross section of a 1 mm SiC disk was exposed

second duration with average ion energy of 150 eV.)to the imaging plate (IP) after the HFIR irradiation. LTO:

Recoiled hydrogen by 2.8 MeV helium was detected as

Lithium titanate (Li2TiO3); LAO: Lithium aluminate

a function of its energy (ERD). The horizontal axis

(LiAlO2)

corresponds to the depth of the hydrogen in the SiC.

25

SiC without breeder

3

1

0

SiC with LTO

IMR KINKEN Research Highlights 2012

behavior of the hydrogen isotopes injected into SiC
are very important engineering factors. The SiC will
not work as a tritium barrier if the mobility is high,and the tritium inventory will be seriously high if theretention is high.
Fig. 1 shows the hydrogen retention in single
crystal SiC exposed to the GAMMA-10 edge plasma
for 29 shots of 0.29 second pulse. [5] The averageion energy was about 150 eV, with a total fluence of
1×1022 H/m2. The elastic recoil detection (ERD)
analysis by 2.8 MeV helium revealed that thehydrogen retention is about 5×1020 H/m2. Exposure
up to (3–6)×1023 H/m2 did not increase the amountof retention, indicating that the retention was efficientin the early stage but it saturated soon. Fig. 2 shows
the tritium concentration profile in single crystal SiC
that was irradiated in the HFIR [1]. Tritium was
injected into the SiC from the oxide through thenuclear reaction of Li6(n, α)T during the reactor
irradiation. Thus, the whole of the SiC was irradiatedby high-flux neutrons during the tritium injection, but
the radiation damage in the tritium-trapped area
was dominated by the damage caused by theenergetic tritium. The irradiation temperature wasabout 800 °C. The retention of the tritium wasmeasured by the image plating (IP) technique andwas found to be substantial. It was estimated thatthe whole of the tritium injected into the SiC wastrapped and that it did not move along the SiC matrix
during heavy irradiation in the HFIR at 800 °C
extending for more than three months.
Ion acceleration studies up to now have shownthat the injected hydrogen isotopes will start to moveat just below 800 °C, as shown in Fig. 3 [4], which iscontradictory to the results obtained in the HFIR
mentioned above. These results imply that theradiation-induced microstructures in the SiC are
acting as strong trapping sites but saturate easily.
Thus, the saturation amount of the retention will bestrongly dependent upon the defect concentration
introduced by radiation damage. SiC is a covalentmaterial and its radiation damage structures andRetained Deuterons (%)

IMR KINKEN Research Highlights 2012

100

50

10 keV D+

16

2x10 ions/cm
0

0

200

400

2

600

SiC
800

1000

1200

Temperature (℃)
Fig. 3 Release of deuterium from SiC after accelerator
injection. Deuterium concentration was measured by
the elastic recoil detection (ERD) method.

tritium trapping defects will be different from those inmetallic and ionic materials. The tritium trapped by
the defects in the SiC, probably Si–T and C–T, arestrong candidates for the trapped hydrogen isotopes
observed in the present study, which are strongly
bound and easily saturate the concentration.
References[1] H. Katsui, et. al., presented at the 15th Int. Conf. on Fusion
Reactor Materials, Charleston, Nov., 2011.
[2] H. Katsui, et. al., J. Nucl. Mater., 417, 753-755 (2011).
[3] H.Katsui, et.al., Fus. Sci. Technol., 60, 288-291 (2011).
[4] S. Nagata, et al., J. Nucl. Mater., 128&129, 760-763 (1984).
[5] S. Nagata, et. al., To be presented at The 24th IAEA Fusion
Energy Conference, 8-13 Oct., 2012, San Diego.

Keywords: nuclear materials, neutron irradiation, ceramics
Tatsuo Shikama (Nuclear Materials Science Division)
E-mail: shikama@imr.tohoku.ac.jp
IMR KINKEN Research Highlights 201226

Energy-related Materials

Effects of Displacive Radiation, Hydrogenation, and Creepin Zr-Based Alloys and FCC Metals

To develop materials for advanced fission and fusion reactors, we investigated Zr- and Fe- based alloys bymeans of a TEM interfaced with an ion accelerator, an HVEM, and mechanical testing techniques. Irradiationinduced hardening in Zr–Nb alloys was evident at a dose of less than 1 dpa (displacement per atom).
Microstructural evolutions under hydrogenation were clarified by in-situ observations. Creep in Zircaloy-4 wasinvestigated to modify the deformation mechanism map, and a new low-temperature creep behavior without
the assistance of diffusion processes was found. In fcc metals, the microstructural configuration of subcascades
in extremely high-energy cascades were characterized.

The mechanical degradations of nuclear
materials strongly depend on their environments.
For instance, Zr alloys for cladding in LWRs areaffected by the distribution of hydrides as well asradiation defects, and Fe alloys for fusion blankets
are affected by both radiation and helium
accumulation. Therefore, it is indispensable tounderstand the microstructural evolutions andresulting mechanical degradations and to clarify
their correlations. We investigate hardening inducedby ion irradiation and/or He/H implantation in Zr- andFe-based alloys by nano-hardness testing andelectron microscopy, as well as fundamental study
on the stability of defect clusters under electron
irradiation in an HVEM.
(1) Effects of hydrogenation in Zr alloys [1,2]Three kinds of tubular-shaped alloys, namely
Zr–1.8Nb, Zr–1.6Nb–0.1Cr, and Zr–2.5Nb (wt. %),were employed in this work. The samples were
subjected to corrosion and hydrogenation
experiments. The corrosion experiments were
performed in PWR primary water conditions, and thehydrogenation was done in a hydrogen-gas-flowchamber. The disk-shaped samples were irradiatedat 570 K with 4 MeV Ni or Zr ions up to 10 dpa at a

depth range from the surface to 500 nm. Conventional
measurements of hydrogen absorption and corrosion
as well as SEM, TEM, EBSD, and XRD were
performed.
The characteristics of hydrogenation in Zr–Nb
alloys are summarized as follows: a linear increasein hydrogen absorption to the duration of exposure
to hydrogen gas, the formation of hydride mostly
parallel to the clad surface, and the formation of highdensity hydride at the rim of the specimen athydrogen absorption above 500 ppm. In-situ
observations under 20 keV H2+ ion implantations atroom temperature in the TEM-accelerator facility at
JAEA revealed that the microstructural evolution in
Zr–Nb alloys was dominated by hydride nucleation,
especially at the beginning of the implantation, while
the growth of hydride dominated in zircaloy over thewhole period of irradiation. Crystallographic
information on the hydrides was clarified by TEM,
EBSD, and XRD. In the corrosive environment, thedifference between the Zr–Nb alloys and zircaloy wasevident. In the Zr–Nb alloys, the weight gain curve
did not show a transition within the experiments, andindeed, a trivial absorption of hydrogen wasobserved. This suggests a higher sustainability of thealloys under in-reactor environments.

Fig. 1 Microstructural evolutions in Zr–1.8Nb under irradiation with 20 keV H2+ ions at room temperature.

27

IMR KINKEN Research Highlights 2012

IMR KINKEN Research Highlights 2012

(2) Creep behavior in Zircaloy-4
Creep tests have been performed in Zircaloy-4
to revise the deformation mechanism map at low
temperature, because the alloy has shown a new
creep mechanism with an apparent activationenergy (Q) of about 20 kJ/mol at less than 373 K. As
shown in Fig. 2, after the creep test, the dislocations
were straightly aligned and were merely tangled.
This indicates that the dislocations are mobile
without evident trapping and tangling sites, that
dislocation-core diffusion is inactivated in graininteriors, and that dislocations pile up at grainboundaries. Due to the extra-low Q values, rather
than the conventional diffusion processes, slipinduced grain boundary sliding (GBS) works as a
rate-controlling process by absorbing the piled-up
dislocations. The GBS can be activated throughdislocation shuffling, which requires a shorter lengthof atomic jump and does not assist the diffusion
process. This is a new creep mechanism in Zircaloy-4
at low temperature.
(3) Defect clusters in fcc metals irradiated withspallation neutrons
Correlation of defect structural development
under high-energy neutron irradiation among
different energy spectra is a key issue for researchand development of future nuclear materials. To
examine the defect reactions involved in extremely
large collision cascades, we inspected defectclusters in fcc pure metals irradiated with spallation
neutrons (energies up to 500 MeV) at room
temperature at KENS, High Energy Accelerator
Research Organization (KEK).
In gold, large defect groups included more than
10 clusters, and the damage region extended over
50 nm, which was larger than that induced by 14
MeV D–T fusion neutron irradiation (<20 nm) [3].
Although small stacking fault tetrahedra (SFTs) wereformed at subcascades by fission and fusion neutron
irradiation [3], dislocation loops were also observedin the present experiments, as shown in Fig. 3. Large
dislocation loops (>10 nm) were identified as
vacancy type by the conventional inside–outsidecontrast method. Large vacancy loops are formed
through cooperative reactions among subcascades
in a single collision cascade with a large recoil
energy.

Fig. 2 Dislocations were straightly aligned after creep testat 294 K.

Fig. 3 TEM images of defect cluster groups in Au irradiatedwith (a)–(f) spallation neutrons 1 × 1018 n m-2, (g) fission
neutrons 4 × 1021 n m-2, and (h) fusion neutrons 2 ×1020 n m-2.

References[1] H. Abe et al, Proc 1st Asian Nuclear Fuel Conference (ANFC)(2012).
[2] T. Maruyama et al, ibid.
[3] Y. Satoh, M. Tsukada, H. Matsui, T. Yoshiie, J. Nucl. Mater.
329-333, 1185 (2004).

Keywords: hydrogen implantation, neutron irradiation, electron microscopy
Hiroaki Abe, Yuhki Satoh and Tetsuya Matsunaga (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 201228

Energy-related Materials

Preparation of Nanoporous Metals by Dealloying With
Metallic Melts

Dealloying, which commonly involves corrosion processes in aqueous solutions, is a promising technique forpreparing functional nanoporous metals. While this technique is ideal for preparing nanoporous noble metals
such as of Au, it is not readily applicable to less-noble metals. Here, we propose a novel dealloying method
employing a metallic melt, instead of an aqueous solution, as the dealloying liquid for the preparation of
nanoporous base metals.
Nanoporous metals have attracted considerable
attention for their excellent functional properties,
including high catalytic activity, sensing capabilities,
and surface-enhanced Raman scattering, which
differ significantly from those of their dense, bulky
counterparts. Such nanoporous metals have mainly
been prepared by dealloying in aqueous solution.
Dealloying is a selective corrosion phenomenon
wherein the base metal in an alloy is selectively
etched by corrosion to form a three-dimensional
porous structure within. However, this technique isonly applicable to the noble metals such as Au. The
less-noble metals cannot form nanoporous
structures because they are easily oxidized in the
aqueous solution. To overcome this, we haverecently developed an alternative dealloying method
using a metallic melt instead of an aqueous
solution[1]. The origin of the dealloying phenomenon
in the metallic melt is not corrosion but the attractive/
repulsive forces among the atoms. The heat of
mixing is the quantitative indicator of the atomic

interactive force. A positive heat of mixing indicatesa repulsive force between atoms and a negative one
indicates an attractive force. To confirm the feasibilityof this new dealloying method, the preparation of
nanoporous Ti was performed based on the above
strategy. Ti is a typical base metal with a very low
standard electrode potential, and it has never formed
nanoporous structures by the conventional
dealloying technique with aqueous solution.
A Cu–Ti alloy and pure Mg were employed as
the precursor and metallic melt, respectively. From
the aspect of the atomic interaction, the positive
heat of mixing between Mg–Ti and negative one for
Mg–Cu should result in the selective dissolution of
Cu from the precursor to the metallic melt, namely,
the dealloying of Cu by the Mg melt.
Fig. 1a shows the microstructure of the Cu–Ti
precursor immersed in the Mg melt at 973 K for 5 s,followed by an etching treatment in 3 mol/l nitric acidaqueous solution for 30 min at room temperature.
The figure shows that an open-cell nanoporous

Fig. 1 SEM images of the (a) surface and (b) cross section of the porous sample prepared by immersing Ti–Cu precursor
into Mg melt at 973 K for 5 s followed by etching treatment[1]29

IMR KINKEN Research Highlights 2012

IMR KINKEN Research Highlights 2012

Fig. 2 TEM image and SAED pattern for the porous sample prepared by immersing Ti–Cu precursor into Mg melt at 973 K
for 5 s followed by etching treatment[1]structurecomposedof
three-dimensionallyinterconnected ~200 nm granules is present. The
cross section of the porous sample in Fig. 1b
indicates that the porous structure is uniformly
spread throughout the entire section. The TEM
image (Fig. 2) and corresponding selected-area
electron diffraction (SAED) pattern confirm that thenanoporous sample has hcp structure, as no oxide
or other compound phase was observed.
Based on the heat of mixing among the atoms
in the precursor and the liquid metal, the detailed
formation process of the nanoporous Ti can be
explained as follows. The Mg–Cu and Ti–Cu bonds
are attractive, whereas the Mg–Ti bond is repulsive.
In the pure Mg melt, only the Cu atoms in the Ti–Cu
alloy precursor are attracted by Mg atoms and,
therefore, are eluted into the Mg melt. The Ti atoms
left at the interface between the alloy precursor andthe melt agglomerate and enrich to form fine solid
granules, and finally, these granules spontaneously
organize to yield a nanoporous structure. This
process is essentially similar to dealloying inaqueous solution, in which the base metal in an alloyof noble and base metals is selectively etched by anelectrolyte, and the noble metal atoms left at thesurface form continuous islands due to surface

diffusion. By changing the temperature of the melt
and immersion time in the melt, the diffusivity canbe controlled and the characteristic lengths of theporous structure, such as pore size or ligament size,
can be tailored over a wide range[2].
As demonstrated, dealloying with a metallic
melt will be a breakthrough for the preparation of
nanoporous structures of base metals for futurefunctional materials.
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).

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

IMR KINKEN Research Highlights 201230

Energy-related Materials

Control of Helmholtz Free Energy Relationship between
Protein Polymorphs under an External AC Electric FieldT

he distribution of the bulk tetragonal phase and spherulitic crystals in hen-egg white lysozyme was controlled
by applying an external alternating current electric field. The phase distribution differed depending on
differences in the magnitude of the electrostatic energy contribution to the respective chemical potentials of thetwo solid phases. Consequently, the Helmholtz free energy relationship between the two solid phases could becontrolled by varying the frequency of the applied external electric field. This method for controlling the Helmholtz
free energy relationship among polymorphs is expected to be applicable to different types of proteins.
Structural crystallography is generally used todetermine the 3D structures of protein molecules.
However, since crystallography requires suitable
crystals, much research has focused on preparing
high-quality single crystals of proteins of interest.
Despite significant efforts in this regard,
crystallization of proteins remains a considerable

Fig. 1 (a) Optical micrograph of bulk and spherulitic
crystals of HEWL coexisting in a droplet of solution. (b)Schematic illustration of the change in the stability
relationship induced by applying an external electricfield. With increasing magnitude of the electric field E,
the line along which bulk and spherulitic crystals coexist
shifts downward when the temperature and soluteconcentration are kept constant. Consequently, thestable phase changes from spherulite to the bulk crystal
with increasing applied electric field E.

31

IMR KINKEN Research Highlights 2012

challenge, and it is still difficult or impossible tocrystallize many proteins. Moreover, spherulitic
crystals rather than bulk crystals are obtained under
growth conditions that provide a high driving force
for nucleation. Since spherulitic crystals arecomposed of thin needles that grow radially outward
from a center (Fig. 1(a)), they are generally unsuitablefor 3D structure analysis.
We have recently been able to increase or
reduce the nucleation rate of hen-egg whitelysozyme (HEWL) crystals by varying the frequencyof an applied electrostatic field [1–4]. The ability tocontrol the nucleation rate via an external electricfield is attributed to the electrostatic energy
contribution to the chemical potentials of both theliquid and solid phases. Whether the nucleation rateincreases or decreases depends on the magnitudeof the difference between the electrical permittivitiesof the liquid and solid phases [5]. A large dispersion
in the electrical permittivity is observed in protein
crystals when an electrical field with a frequency
below 1 MHz is applied [6]. The magnitude of thedifference in the electrical permittivities of the liquidand solid phases can be regulated by exploiting thislarge dispersion in the electrical permittivity, andthus the nucleation rate in HEWL crystals can becontrolled by applying an external AC electric field[1]. This novel approach represents a breakthroughin controlling the nucleation of protein crystals.
Protein crystals contain large amounts of water,
the amount of which varies among different crystal
polymorphs. For protein crystals, it has beensuggested that the frequency at which dispersion of
the electrical permittivity occurs increases withincreasing water content [7]. Furthermore, because
different crystal polymorphs have different dielectric
properties, it may be possible to control (or evenreverse) the difference in the electrical permittivities
of different crystal polymorphs. That is, as shown in
Fig. 1(b), the phase stability relationship between thebulk and spherulitic crystals can be controlled bytailoring the relative contribution of the electrostatic
energy to the respective chemical potentials of thetwo solid phases such that the phase stability isreversed. In the following, we demonstrate activeIMR KINKEN Research Highlights 2012

Fig. 2 shows the phase distributions in dropletsof the solution with and without an external electricfield. They can be classified into four types: (i) no
nucleation, (ii) bulk crystal, (iii) bulk crystal +spherulite, and (iv) spherulite. The number ofdroplets observed by optical microscopy is indicatedin each figure. For a 1 MHz applied field, theproportion of bulk crystals formed in the solution
droplet was larger than that without an applied field,
whereas the proportion of spherulitic crystals
decreased when a field was applied (Fig. 2(b)). Thisimplies that the driving force for nucleation of thebulk crystal was larger than that of the spherulitic
crystal when a 1 MHz external electric field wasapplied. In contrast, when a 3 MHz electric field wasapplied, the proportion of spherulitic crystals formed
increased, whereas the proportion of bulk crystals
decreased (Fig. 2(c)). This indicates that the driving
force for the nucleation of bulk crystals is muchsmaller than that for spherulitic crystals when a 3
MHz external electric field is applied.
Such a large electrical dispersion is considered
to be a common property of almost all protein
crystals, and thus this method for controlling theHelmholtz free energy relationship among
polymorphs should be applicable to many proteins.
References[1] H. Koizumi, K. Fujiwara and S. Uda, Cryst. Growth & Des. 9,2420 (2009).
[2] H. Koizumi, K. Fujiwara and S. Uda, Cryst. Growth & Des. 10,2591 (2010).
[3] H. Koizumi, S. Uda, K. Fujiwara and J. Nozawa, J. Crystal
Growth 312, 3503 (2010).
[4] H. Koizumi, S. Uda, K. Fujiwara and J. Nozawa, Langmuir 27,
8333 (2011).
[5] S. Uda, X. Huang and S. Koh, J. Cryst. Growth, 281, 481(2005).
[6] S. Takashima, J. Polym. Sci. 62, 233 (1962).

Fig. 2 Phase distributions in solution droplets with andwithout the application of an external electric field. (a)No electric field and applied fields at (b) 1 and (c) 3
MHz.

[7] S. Bone, J. Eden and R. Pethig, Int. J. Quantum Chem. 20, 307(1981).
[8] Y. Tomita, H. Koizumi, S. Uda, K. Fujiwara and J. Nozawa, J.
Appl. Crystallogr. 45, 207 (2012).

control of the phase distribution between the bulk
tetragonal phase and spherulitic crystals of HEWL
under the application of an external AC electric field[8].

Keywords: phase equilibria, nucleation and growth, crystal growthHaruhiko Koizumi and Satoshi Uda (Crystal Chemistry Division)
E-mail: h_koizumi@imr.tohoku.ac.jp
URL: http://www.uda-lab.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201232

Energy-related Materials

Complex Hydrides as Advanced Lithium Fast-Ionic
Conductors

In this report, we briefly summarize the fast-ionic conduction in LiBH , a typical complex hydride, and then we
4

explain the conceptual development of complex hydrides as a new type of solid-state lithium fast-ionic conductor
using LiBH4-based and LiNH2-based complex hydrides [1].

LiBH4 –LiX (X = Cl, Br and I) system [4,5]Fig. 1 shows the electrical (here, simply ionic)
conductivities of (1 - x)LiBH4 + xLiI/xLiCl. Thechange in the activation energy due to the structuraltransition is observed at around 340 K for LiI with x= 0.13, whereas the conductivities for x = 0.25
exhibit Arrhenius behavior throughout the measured
temperature range because of the HT-LiBH4
stabilization. As expected, the conductivity at RT
increases by three orders of magnitude without adecrease in the high conductivity of the HT-LiBH4.
Furthermore, the activation energy decreases withincreasing LiI content due to the faster mobility of Li+ions caused by the high polarizability of the I – ionssubstituted for [BH4]− ions. The minimum value isobtained at x = 0.13 (0.39 eV). The replacement of[BH4]− ions with Cl− or Br− ions enhances the
conductivity.
LiBH4 –LiNH2 system [6]As shown in Fig. 2, Li2(BH4)(NH2) exhibits afast-ionic conductivity of 1 × 10 –4 S cm−1 even at
33

IMR KINKEN Research Highlights 2012

Temperature (K)
0

10
-1

Lithium ionic conductivity (S cm )Complex hydrides exhibit various energyrelated functions, such as hydrogen storage,
microwave absorption, and neutron shielding.
Furthermore, we have reported another novel
energy-related function: lithium fast-ionic conduction[2,3].
The conductivities of the low-temperature phaseof (LT-)LiBH4 are very low, between 10−8 and 10−6
S cm−1, and they increase monotonically uponheating. Surprisingly, at approximately 390 K, i.e.,the structural transition temperature, the conductivity
jumps up drastically by three orders of magnitude.
As a result, the high-temperature phase of (HT-)
LiBH4 exhibits high conductivity, on the order of 10−3
S cm−1. The activation energies for conduction havebeen evaluated as 0.69 eV and 0.53 eV for the LT
and HT phases, respectively.
The lithium fast-ionic conduction in LiBH4 could
potentially contribute to the development of solid
electrolytes in all-solid-state batteries. For related
applications, however, it is highly desirable toenhance the conductivity at room temperature (RT).
To achieve this, we now present our recent
conceptual material developments of fast-ion
conductors from LiBH4-based complex hydrides.

450

400

350

300

1) 0.75LiBH4+0.25LiI-1102) 0.87LiBH4+0.13LiI-2

103) 0.75LiBH4+0.25LiCl

-3

10

-4

1)

10

-5

10

2)

-6

10

3)

-7

LiBH4

10

-8

10

2.2

2.4

2.6

2.8

3.0

-1

3.2

3.4

1000/T (K )Fig. 1 shows temperature dependence of the ionic
conductivities of the LiBH4 –LiI/LiCl system.

RT, which is four and five orders of magnitude
higher than those of the host hydrides LT-LiBH4 and
LiNH2, respectively. Moreover, the conductivity
increases monotonically upon heating. The
activation energy for conduction decreases
significantly at around 368 K from 0.66 eV (303–348
K) to 0.24 eV (above 368 K) as a result of the meltingof Li2(BH4)(NH2). The total ionic conductivity
reaches 6 × 10 –2 S cm−1 after melting at the highest
temperature measured, 378 K. This result suggeststhat Li2(BH4)(NH2) could be used as a new type ofionic liquid as well as a solid-state fast-ionic
conductor. Li4(BH4)(NH2)3 also shows a highconductivity of 2 × 10 –4 S cm−1 at RT, and the value
reaches 2 × 10 –1 S cm−1 at 513 K after melting.
Furthermore, the activation energy for conduction
before melting is determined to be 0.26 eV. It isnoteworthy that this value is less than half the valueof Li2(BH4)(NH2) before melting or LiBH4 (LTphase: 0.69 eV; HT-phase: 0.53 eV), indicating that
Li4(BH4)(NH2)3 has higher Li+ ion mobility.
LiNH2-LiI system [7]Fig. 3 shows the ionic conductivities for x =0.06, 0.13, and 0.25 of (1 - x)LiNH2 + xLiI. For all
values of x, the conductivities exhibit Arrhenius
behavior throughout the measured temperature
range. The maximum conductivity of 2 × 10 –5 S cm –1

IMR KINKEN Research Highlights 2012

Temperature (K)
450

400

350

Tm

Tm

-1

10

-2

10

Li2(BH4)(NH2)

-3

10

-4

Li4(BH4)(NH2) 3

10

-5

10

-6

10

LiBH4

-7

10

LiNH2

-8

10

1.8

2.0 2.2 2.4

2.6

2.8 3.0-1

Temperature (K)

300

450

0

-1

10

550 500

Lithium ionic conductivity (S cm )-1

Lithium ionic conductivity (S cm )0

3.2 3.4

10

400

350

300

-1

10

-2

10

(1–x)LiNH2+xLiI

-3

10

-4

10

x = 0.25

-5

10

0.13

-6

10

0.06

LiI

-7

10

LiNH2

-8

10

2.2

2.4

1000/T (K )

2.6

2.8

3.0

-1

1000/T (K )

3.2

3.4

Fig. 2 shows temperature dependence of the ionic
conductivities of the LiBH4 –LiNH2 system.

Fig. 3 shows temperature dependence of the ionic
conductivities of LiNH2–LiI system.

at 300 K, which is four and three orders of magnitude
higher than those of the host materials LiNH2 and
LiI, respectively, is obtained at a composition of x =0.25. At this composition, the activation energy for
conduction is determined to be 0.58 eV, comparable
to that of HT-LiBH4 (0.53 eV). The notable lithium
fast-ionic conduction around x = 0.25 is probably
caused by a specific crystal structure of the newly
found single-phase Li3(NH2)2I [7], as shown in Fig.
4; that is, many intrinsic interstitial sites appear at thecenter of clusters of six tetrahedral (T1, T2) andbetween the clusters. Li+ ions may migrate throughthese spaces.

Li3(NH2)2I

T1
T2

N
I

Li
H

[600]

Further research is expected to yield practicalapplications of the complex hydrides as solid-statelithium fast-ionic conductors (namely, solid
electrolytes for batteries) and to demonstrate newphenomena that are characteristic of complex
hydrides. This article was produced according to
the review paper [1].

[5] M. Matsuo, H. Takamura, H. Maekawa, H.-W. Li, S. Orimo,

References[6] M. Matsuo, A. Remhof, P. Martelli, R. Caputo, M. Ernst, Y.

Fig. 4 Crystal structure of Li3(NH2)2I viewed along [600].
Please, see the details in Ref. [7][4] H. Maekawa, M. Matsuo, H. Takamura, M. Ando, Y. Noda, T.
Karahashi, S. Orimo, J. Am. Chem. Soc. 131, 894 (2009).
Appl. Phys. Lett. 94, 084103 (2009).

[1] M. Matsuo and S. Orimo, Adv. Energy Mater. 1, 161 (2011).

Miura, T. Sato, H. Oguchi, H. Maekawa, H. Takamura, A.

[2] M. Matsuo, Y. Nakamori, S. Orimo, H. Maekawa, H. Takamura,

Borgschulte, A. Züttel, S. Orimo, J. Am. Chem. Soc. 131, 16389

Appl. Phys. Lett. 91, 224103 (2007).
[3] T. Ikeshoji, E. Tsuchida, T. Morishita, K. Ikeda, M. Matsuo, Y.
Kawazoe, S. Orimo, Phys. Rev. B 83, 144301 (2011).

(2009).
[7] M. Matsuo, T. Sato, Y. Miura, H. Oguchi, Y. Zhou, H. Maekawa,
H. Takamura, S. Orimo, Chem. Mater. 22, 2702 (2010).

Keywords: hydride, ionic conductor, energy storage
Motoaki Matsuo, Shin-ichi Orimo (Hydrogen Functional Materials Division)
E-mail: mmatsuo@imr.tohoku.ac.jp, orimo@imr.tohoku.ac.jp
URL: http://www.hydrogen.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201234

Energy-related Materials

Electron Microscope Study of the Formation of Graphitic
Nanostructures in Nickel-loaded Wood Char

Graphitization of carbonized larch wood chars impregnated with an aqueous solution of nickel acetate has

been investigated by using scanning electron microscopy (SEM), both in secondary and transmission modes,
and high-resolution transmission electron microscopy (HRTEM). Graphitization sporadically started at thistemperature, and some of the Ni particles were aggregated. SEM observations on chars heat-treated at 900 °C
suggested that graphitic nanoshells about 50–200 nm in diameter formed by the catalytic effects of the Ni
particles.

Renewable carbon-neutral wood resources, orwoody biomass, are promising sources for energy
conversion in view of global warming and petroleum
depletion; therefore, they are the subject of a numberof research and development activities. For example,
Suzuki et al. have reported an integrated refinery
process for woody biomass, which accounts forabout 90 % of the total biomass, by using nickeland/or iron-catalyzed carbonization [1, 2]. Theirresearch started by recognizing the need for highly
efficient but environmentally undemanding energy
conversion and production techniques by employing
low-cost wood resources and common chemicals
as starting materials. One promising process iscarbonization of nickel-acetated wood flakes, where
impregnation of raw wood materials with a nickel
acetate solution, followed by their carbonization at ahigh temperature, is the key treatment. The mostdistinctive feature of this method is the efficient
production of high-quality wood charcoal or
crystallized mesoporous carbon, which has the dual
functionality of yielding high conductivity and a highmolecular liquid-phase adsorption capacity. These
functions make them potentially importantcandidates for electrode materials for lithium-ion

batteries and electrical double layer capacitors.
These industrially important functions areclosely related to the microstructure of the wood
charcoals, and several studies have already beenmade. For example, Suzuki et al. used transmission
electron microscopy (TEM) to show that nickelacetated wood charcoal carbonized at 900 oC
consists of a crystal part, which looked like a ribbonshaped spiral, and of amorphous carbon. Theirobservation led them to develop a new wood
charcoal, in which nickel catalysts and amorphous
carbon are removed from Ni-catalyzed wood
charcoals by acid treatment and oxidization. The
new wood charcoal, which had improved
conductivity, was not ribbon-shaped but consistedof shell-shaped particles covered with graphitic
stacks, which they termed "graphitic shell chains"(GSCs). The purpose of this study was hence tofollow the evolution of Ni atoms inside nickelacetated wood charcoal and to elucidate theunderlying mechanisms of the formation of a threedimensionally extended GSC structure. We
employed scanning electron microscopy (SEM),
equipped with a scanning transmission electron
microscopy (STEM) detector, to investigate

Fig. 1 Stereo pair of SEM image of wood carbon prepared with a metal catalyst, showing three-dimensionally entangled
graphite shell units.

35

IMR KINKEN Research Highlights 2012

IMR KINKEN Research Highlights 2012

Fig. 3 HRTEM image (80 kV with a Cs value of -12 µm) ofgraphitic layer structure and FFT of the region indicatedby the arrow.

Fig. 2 (a) SE-SEM (30 kV); (b) BF-STEM; (c) DF-STEM
images (150–380 mrad) of Ni-loaded wood char, heattreated at 900 °C

morphological changes, and we used highresolution TEM for observations at the atomic scale.
[3]
Fig. 1 is a stereo pair of SE-SEM images taken
at 3 kV, showing that GSCs have grown outwards ina winding manner, giving rise to a three-dimensionallyintertwinedcomplexmicrostructure.
These
observations thus confirm that acid treatments andpost-oxidization are effective means of removing
metallic catalysts and residual amorphous carbon,
leading to a drastic increase in surface areas, asreported previously.
Fig. 2 shows (a) SE-SEM (30 kV), (b) BF-STEM,
and (c) DF-STEM images (150–380 mrad) takenfrom a thin area of 900 oC NiWC flake (all taken fromthe same position). (a) The arrows show that the
weakly contrasted areas are coincident with thestrong contrasts in b and c. In the transmission
mode (b and c), gray winding contrasts are also
seen clearly. These contrasts originate from thegraphitic shell chain structure. Most of the Ni
particles are surrounded by GSCs, whereas thearrowed particle in (b) is "intact." The arrow in theupper left in (c) shows an aggregate of small Ni
particles.
Finally, in Fig. 3, we show an HRTEM image ofthe GSC structure. Here, not only are the (0002)
planes seen, but graphitic layers comprising an
ABAB... stacking sequence can also be confirmed,
even though the graphite atoms at the 2b and 2c
positions, whose distance is 0.067 nm, are notresolved.
To summarize, we have demonstrated that a
combination of low voltage SEM, both in conventional
SE and transmission (STEM) modes, and HRTEM isa versatile technique to elucidate the mechanism ofmetal-catalyzed graphitization.
References[1] T. Suzuki, K. Suzuki, Y. Takahashi, M. Okimoto, T. Yamada, N.
Okazaki, Y. Shimizu, M. Fujiwara, J. Wood Sci. 53, 54-60(2007).
[2] T. Suzuki, K. Suzuki, Y. Takahashi, M. Okimoto, T. Yamada, N.
Okazaki, Y. Shimizu, M. Fujiwara, Chem. Lett. 34, 870-871(2005).
[3] Y. Kodama, K. Sato, K. Suzuki, Y. Saito, T. Suzuki, T.J. Konno,
Carbon 50, 3486-3496 (2012).

Keywords: graphitization, SEM, TEM
Toyohiko J. Konno, Yumiko Kodama, and Kazuhisa Sato (Advanced Analysis of Materials Division)
E-mail: tjkonno@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201236

Energy-related Materials

Production of Ni-Based Amorphous Alloy-Coated Bipolar
Plates for Proton Exchange Membrane Fuel Cells by
Electrochemical PlatingIn this study, Ni–P amorphous alloy-coated bipolar plates were produced by electroless plating, and Ni–Cr–P
amorphous alloy-coated bipolar plates were produced by electroplating on metal plates with a flow field. Powergeneration tests were conducted at 353 K on single fuel cells with these bipolar plates. It was found that thesingle fuel cell with the Ni–Cr–P-coated bipolar plates showed better I–V performance than that with the Ni–Pcoated bipolar plates. This is because the addition of Cr to the Ni–P amorphous alloy improves its corrosion
resistance. The single fuel cell with the Ni–Cr–P bipolar plates showed excellent I–V performance, performing
as well as a cell with carbon bipolar plates.

Production of Ni-based amorphous alloy-coatedbipolar plates [2, 3]
Fig. 1 shows the outer view of the Ni–P

amorphous alloy-coated bipolar plate produced byelectroless plating. Amorphous alloy thin films were
deposited on both sides of Al substrates with a flow
field.
Fig. 2 shows the XRD pattern of the Ni–P
amorphous alloy thin film deposited on an Al plate.
As clearly seen in the figure, the film showed a broad
halo peak without any distinct peaks from crystalline
phases. The P content of the film was 11.0 mass%
P.
Fig. 3 shows the outer view of the Ni–Cr–P

Fig. 1 Outer view of the Ni–P amorphous alloy-coatedbipolar plate produced by electroless plating.

Fig. 3 Outer view of the Ni–Cr–P amorphous alloy-coatedbipolar plate produced by electroplating.

Fig. 2 XRD pattern of the electroless plated Ni–P alloy thinfilm deposited on an Al plate.

Fig. 4 XRD pattern of the electroplated Ni–Cr–P alloy thinfilm deposited on a Cu plate.

The bipolar plates are one of the most importantcomponents of proton exchange membrane fuelcells (PEMFC) [1]. In this study, we produced Ni–P
[2] and Ni–Cr–P [3] amorphous alloy-coated bipolar
plates by electrochemical plating and examined theI–V performance and long-time durability of a single
fuel cell with these plates.

37

IMR KINKEN Research Highlights 2012

IMR KINKEN Research Highlights 2012

amorphous alloy-coated bipolar plateproduced by electrochemical plating. A Cu
plate was covered with polymer film, exposing only
the flow field area to the chemicals. So, the Ni–Cr–P
alloy was deposited only around a flow field on bothsides of the plate.
Fig. 4 shows the XRD pattern of the Ni–Cr–P
amorphous alloy thin film deposited on a Cu plate.
The broad halo peak comes from a single amorphous
phase and the distinct peaks may come from the Cu
plate outside the area around the flow field.
According to the analysis, the alloy produced in thisstudy was Ni52Cr25P23 in at%.
Power generation tests [2, 3]
Fig. 5 shows the I–V curves of single fuel cellswith the Ni–P and Ni–Cr–P bipolar plates. The I–V
curve measured with carbon graphite issuperimposed on the figure for comparison. As
seen in the figure, the single fuel cell with the Ni–
Cr–P bipolar plates shows excellent I–V performance,
performing as well as that with the carbon bipolar
plates. Furthermore, it was found that the single cellwith the Ni–Cr–P bipolar plates showed better I–V
performance than that with the Ni–P bipolar plates.
This is because the former includes a high content
of the corrosion resistant element Cr.
Fig. 6 indicates the results of the long-time
power generation test conducted at a current densityof 200 mA/cm2 at 353 K for 440 h. The final productinside a fuel cell is water. Water should be removed
from the fuel cell quickly so as not to degrade thecell performance, especially over a long operation
time. Therefore, the gas flow rates of H2 and O2 wereincreased to 0.3 L/min for this test in order to drain
off the water produced inside the fuel cell completelyfrom the flow field of the fuel cell. As seen in the
figure, the voltage decreased gradually with timefrom the beginning of the measurement until 300 h,
and then the voltage was almost constant after 300
h. This may be because a protective passive film
was formed on the surface of the Ni–Cr–P bipolar
plates during the test, leading to fully passivatedplate surfaces. Although some small microcracks
were observed in the microstructure, they did not
significantly affect the I–V performance shown in
Fig. 5. The results of the long-time operation test in
this work were not affected significantly by the
microcracks. However, it is not clear whether the
microcracks would degrade the cell performanceFig. 5 I–V performance of single fuel cells with the Ni–P
and Ni–Cr–P amorphous alloy-plated bipolar plates.

Fig. 6 The result of long-time power generation at aconstant current density of 200 mA/cm2 for 440 h.

after longer operation. We need to investigate
further. It is also important to optimize the conditions
to prepare electroplated amorphous alloy films withsmooth surfaces and without microcracks.
At any rate, the Ni-based amorphous alloycoated bipolar plates were produced byelectrochemical plating. The potential of amorphousalloy-coatedbipolarplatesproducedbyelectroplating was successfully shown in this work.
References[1] S.C. Kim, S. Yamaura, Y. Shimizu, K. Nakashima, T. Igarashi,
A. Makino and A. Inoue, Mater. Trans. 51, 1609 (2010).
[2] S.C. Kim, S. Yamaura, A. Makino and A. Inoue, Mater. Trans.
52, 709 (2011)[3] S. Yamaura, S.C. Kim and A. Inoue, J. Phys.: Conf. Ser. (2012)
in press.

Keywords: fuel cell, amorphous, plating
Shin-ichi Yamaura (NEDO Laboratory for Metalic Glass Research Collaboration and Promotion)
E-mail: yamaura@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201238

Electronic Materials

Electrically Activated Magnetic Insulators

We study the excitation of spin waves by the current-induced spin-transfer torque in a magnetic insulator. We
predict the preferential excitation of surface spin waves induced by a perpendicular surface anisotropy, whichstrongly reduce the critical currents and enhance the excitation power of the current-induced magnetization
dynamics.
By the discovery of electrical [1] and thermal [2]excitation of yttrium iron garnet (YIG) films, theSaitoh group established magnetic insulators as animportant material class for spintronics and spin
caloritronics [3]. The demonstration of the electricalgeneration of spin waves by the spin Hall effect(SHE) and their detection by the inverse SHE adds
credibility to the spin-wave bus concept for lowdissipation interconnects between spintronics
circuits [4]. Here, we discuss a theoretical treatmentof spin wave generation by the spin Hall effect in
Pt|YIG bilayers [5] (Fig. 1) and address the physics
behind the observed low critical currents andcomplex power spectrum.
The magnetization dynamics of YIG are
described by the Landau–Lifshitz–Gilbert (LLG)
equation with boundary conditions that take into
account injection of a spin current Js at the x = 0
interface and free boundary conditions at x = -d.
The dynamics under excitation are governed by themagnetic anisotropy energy of the magnetization M
as a function of its direction relative to the atomic
lattice. The magnetic anisotropy can be very different
at the surface, often favoring a ground statemagnetization normal to the interface, which
competes with the magnetostatic anisotropy that
prefers in-plane magnetization. Both the spin current
Js and the surface anisotropy Ks can beparameterized in terms of the effective wave
andks=2γKs/AexMs,
numberskj=2γJs/AexMs
respectively, where Aex is the exchange constant or
spin wave stiffness, Ms is the saturation magnetization,and γ is the gyromagnetic ratio.
We solve the above problem numerically after
linearization for in-plane wave vectors q, leading tocomplex dispersions ωn (q) of spin waves with mode
index n as a function of kj and ks. The imaginary partof ωn (q) is the effective damping, which is positive in
the absence of a current and is governed by the
Gilbert parameter α. The torque due to the absorbed
39

IMR KINKEN Research Highlights 2012

spin current can either increase or decrease thedamping, depending on its sign and polarization. At
currents that render Imωn (q) = 0, the respectivemode becomes instable, which defines the thresholdfor (mode-dependent) current-induced spin wave
excitation [6]. In the following, we discuss the results
obtained using established parameters for YIG anda magnetic film thickness of 0.61 μm. The magnitudeof the surface anisotropy is not known and is treatedas a parameter. The main results are shown in Fig.
2, in terms of the spin wave power spectrum
approximated as

P (ω ) = ∑n∫ Im ω ( q ) δ [ω − Re ω ( q )] d q.
nn

Im ω n < 0

x
Js

Pt
YIG

M
z

Fig. 1 The samples are bilayers of electrically insulating
magnetic films (YIG) of thickness d with magnetization
M and normal metal (Pt). A spin current Js generated in
the normal metal, e.g., by the spin Hall effect in Ref. [1],excites spin waves in the ferromagnet by spin-transfer
torque [6].

In Fig. 2, the power spectrum is plotted with andwithout magnetic surface anisotropy at various
current levels. From these results, we can draw
several conclusions:
1) Perpendicular surface anisotropy induces a new
type of surface spin wave that is excited moreeasily than bulk or magnetostatic surface spinwaves.
2) Multiple modes at various frequencies can be
excited almost simultaneously. We explain thefeatures of the experiments by Kajiwara et al. [1],

IMR KINKEN Research Highlights 2012

P (a.u.)j c

0.8

0.9 1 1.1Re(ω/ω )

1.2

8

0
−0.2

P(ω) (a.u.)P(ω) (a.u.)0−0.2 k /k 000.7

ks = 25.0/µm

0.5

0.04
P (a.u.)0.025

ks = 0

0
0.7

k /k

j c

0.8

0

0.9 1 1.1Re(ω/ω )

1.2

M

M

Fig. 2 Power spectrum (resolution δω/ωM = 0.01, where ωM = γμ 0Ms) at various current levels (kj = 0.2kc at the top, decreasingby Δkj = 0.01kc), without (left: ks = 0) and with (right: ks = 25.0/μm) surface anisotropy [5]. Inset: the integrated power
versus kj.

such as very small critical currents and abroadband spin wave power spectrum only in thepresence of surface anisotropy, possibly inducedby the Pt overlayer.
3) Agreement with observed critical currents can beachieved by assuming highly efficient spincurrent-to-magnon
conversion
at
YIG|Ptinterfaces, as expressed by the spin-mixing
conductance. This is consistent with predictions
from first-principle calculations [7].
4) The easily excitable surface modes havesignificant group velocities.
We conclude from these studies that a magnetic
surface anisotropy should be invoked to explain theexperiments by Kajiwara et al. [1]. Our results
indicate that YIG thin films are suitable materials toallow fast and efficient information transmission overlarge distances that can be applied as interconnects
for power-saving future electronic devices.

References[1] Y. Kajiwara et al., Nature 464, 262–266 (2010).
[2] K. Uchida et al., Nature Mater. 9, 894–897 (2010).
[3] G.E.W. Bauer, E. Saitoh and B.J. van Wees, Nature Mater. 11,
391–399 (2012).
[4] A. Khitun and K. L. Wang, J. Appl. Phys. 110, 034306 (2011).
[5] J. Xiao and G.E.W. Bauer, Phys. Rev. Lett. 108, 217201(2012).
[6] J. C. Slonczewski, J. Magn. Magn. Mater. 159, L1-L7 (1996).
[7] X. Jia, K. Liu, K. Xia, and G. E.W. Bauer, Europhys. Lett. 96,
17005 (2011).

Keywords: magnetic, insulator, electrical spin injection
Jiang Xiao (Fudan U., Shanghai) and Gerrit E.-W. Bauer (IMR, Theory of Solid State Physics Division)
E-mail: g.e.w.bauer@imr.tohoku.ac.jp
URL: http://www.bauer-lab.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201240

Electronic Materials

High Resolution of THz ESR Uncovers the Mystery of Spin
PolyhedronM

athematicians have investigated large and complex polyhedrons intensively for centuries because of theirunique geometrical properties. Recently, coordination chemists found that the idea of an icosidodecahedron
is realized in magnetic clusters. Today, the quantum truncated Keplerate {Mo75V20} is examined by the highresolution ESR of the magnetism division. The origin of the residual magnetic moments and the exchange
network are determined by the microscopic separation between the intrinsic and the extrinsic magnetic
responses.
The Keplerate molecules {Mo72Fe30}, {Mo72Cr30}
and {Mo72V30} are state-of–the-art objects
synthesized by coordination chemists[1]. Each has
30 magnetic sites sitting on an icosidodecahedron,
which is made up of pentagons and triangles. Whenmagnetic couplings are antiferromagnetic, thesystem is strongly frustrated and the energy levels
are highly degenerate. An S = 0 non-magnetic
ground state is expected for the system because
the number of magnetic sites is even. However, thenumber of states is huge for longer site spin such as
Fe3+, and a proximity to the classical ground statemight be realized with the recovery of the local
moments[2].
On the other hand, the existence of an S = 0
singlet ground state is plausible when the site spin isshort and is associated with a strong quantum
fluctuation. The cluster of {Mo72V30} with S = 1/2, theV4+ spins is expected to be a “quantum Keplerate.”
When magnetic susceptibility is examined, however,a non-zero value is found at low temperatures. It is
proposed that paramagnetic impurities remaining in
the clusters are the origin of the residual magnetic
components. Attempts have been made to separate
the intrinsic and the extrinsic components of themagnetic susceptibilities, but it is difficult becauseof the very low density of magnetic moments andlarge diamagnetic backgrounds.
To overcome this difficulty, we use a microscopic
probe of electron spin resonance (ESR) to separate
the different components. The intensity of the ESR
signal is proportional to the product of the populationdifference and the transition matrix element between
multiplet levels. Thus, the temperature dependenceof the intensity is useful to determine the energy
levels of the clusters. For example, ESR of aparamagnetic impurity increases in intensity steadily
toward lower temperatures. On the other hand, a
Schottky-peak-like behavior is expected for theclusters with singlet ground states. The question isthen how these two components can be separated.
The answer is that the two components can beseparated by their difference in line widths with theaid of the ultimate resolution of THz ESR.
Fig. 1 shows the schematic structure of a41

IMR KINKEN Research Highlights 2012

Keplerate cluster. At the bottom, the equivalentmagnetic model for a truncated Keplerate is shown[3]. This is the very rare realization of a saw-tooth
chain with a periodic boundary condition. The
ground state of the system is known to be singlet.
There are excited triplets and other multiplets above
the ground state, and those excited states contributeto the ESR transitions.

Fig. 1 The schematic structure of a Keplerate cluster{Mo72Fe30} is shown. In a quantum version, Fe3+ ion is
replaced with V4+ ion. Further modification by truncating
two pentagon rings at the top and bottom leads to a
{Mo75V20} ring cluster. The system can be regarded asa saw-tooth chain with a periodic boundary condition.

Fig. 2 shows an example of the ESR spectrumat 190 GHz. The signal consists of two components
with different line widths. When the temperaturedependence is assessed, the two peaks showcompletely different behaviors. The sharp signal
grows in intensity towards lower temperatures, while
the intensity of the broad peak decreases rapidly
when the temperature is lowered. With thesetemperature dependences, the origins of the two
peaks can be assigned as follows: peak B is caused
by a paramagnetic impurity and peak A is caused
by a transition between thermally excited states ofIMR KINKEN Research Highlights 2012

the {Mo75V20} cluster. One of the origins of the broadline width of the intrinsic signal is the strong exchange
couplings among the magnetic ions. The frustrationof the triangular network also causes broadening.
The present analysis verifies the idea of twomagnetic components that was speculated uponfrom magnetic susceptibility measurements. The
components can be separated easily by ESR fromthe difference in line width. It should be noted thatthe separation is possible because of the highresolution and high sensitivity of the THz ESR
system.

legs. The absolute value was also determined to be
150 K. As shown here, ESR gives not only aqualitative view but also a quantitative understandingof the magnetic parameters of the system.

Fig. 3 Temperature dependence of the intensity of the
intrinsic ESR signal A. The onset in the low temperatureside is a measure of the singlet–triplet gap, and theabsolute intensity and temperature of the broad peak
are sensitive to the low-energy magnetic levels. These
characteristics enable us to evaluate the exchange
couplings.
Fig. 2 An example ESR spectrum of a {Mo75V20} cluster.
The sharp and broad peaks are marked as B and A. B
is assigned to a paramagnetic impurity signal. Signal A
originates from the body of the cluster. The broad widthof A is caused by the large exchange coupling and by
the strong frustration.

The experimental and calculated temperaturedependence of the intensity is plotted in Fig. 3. In
this curve, there are several experimental fingerprintsof the energy levels, such as the onset temperature,
the peak position, and the peak intensity. These
parameters are related to the singlet–triplet gap aswell as the density of low energy levels and theirabsolute spacing. The structure of the energy levels
depends on the exchange couplings of the triangularunits. We found that the exchange coupling in thebase of the triangular ring is half of those of the two

In summary, we succeeded in separating the
intrinsic ESR signal of the molecules from the signalof the magnetic impurities in the spin polyhedroncluster {Mo75V20}. The parameters of the underlying
Heisenberg Hamiltonian were then deduced.
We would like to stress that ESR is a powerfultool to evaluate the low-energy spectrum of frustrated
magnetic clusters. It is much more effective thanother methods, such as specific heat.
References[1] C. Schroder, H. Nojiri, J. Schnack, P. Hage, M. Luban and P.
Kogerler, Phys. Rev. Lett. 94, 017205(2005).
[2] O. Cepas, T. Ziman; Prog. Theor. Phys. Supplement 159, 280(2005).
[3] Y. Oshima, H. Nojiri, J. Schnack, P. Kogerler and M. Luban,
Phys. Rev. B 85, 024413(2012).

Keywords: nanocluster, electron spin resonance, frustration
Hiroyuki Nojiri (Magnetism Division)
E-mail: nojiri@imr.tohoku.ac.jp
URL: http://www.hfpm.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201242

Electronic Materials

Electron Localization in Cooperation with Mott Insulator
and Anderson LocalizationT

he Mott–Anderson transition has been known as a metal–insulator transition due to both strong electron–
electron interactions and the randomness of the electrons. Recently, we demonstrated the randomness effectof the strongly correlated electrons of BEDT-TTF molecule-based organic conductors. X-ray irradiation of thecrystals introduces molecular defects, which cause random potential modulation for the correlated electrons.
In combination with hydrostatic pressure, we are able to control the parameters of randomness and correlations
for electrons approaching the Mott–Anderson transition.
BEDT-TTF-based molecular conductors is a goodcandidate for studying the randomness effects inthe correlated electrons. The Mott transition andsuperconductivity in the half filled dimer–Mott
system κ-(BEDT-TTF)2X have been investigatedactively for a typical bandwidth-controlled metal–
Mott insulator transition.
Randomness can be introduced by X-ray
irradiation of the molecular crystals, which does notmodulate the electronic states but only the scattering
for conduction electrons [4] because the molecular
defects are introduced mainly in the anion molecule
layers. Random potential modulation in space forthe conduction electrons in the BEDT-TTF molecule
layers is induced by the anion molecules defects.
Fig. 1 shows the temperature dependence ofthe resistivity in κ-(BEDT-TTF)2X with X = (a)Cu(NCS)2, (b) Cu[N(CN)2]Br, (c) Cu[N(CN)2]Cl, and

The metal–insulator transition is an importantresearch subject on electronic properties in solidstate physics. Among the various types of metal–
insulator transitions, the Mott transition due toelectron–electron interactions is one of the most
attractive phenomena [1,2]. Another type of metal–
insulator transition is caused by electron localizationthat originates from the interference of electron wave
functions due to randomness. This results in anAnderson localization insulator, derived byintroducing disorder into the crystal [1,3]. For morethan half a century, much active research has been
conducted theoretically and experimentally since P.
W. Anderson propounded the concept in 1958 [3].
Since randomness in the correlated electron systemis essentially important in real materials, systematicstudies of disorder and randomness effects aredesired in systems near a Mott transition. A series of
105

(a)

104

(b)

(c)

103

500h

102

ρ (Ωcm)

50h
100h
150h
225h
300h

400h
300h
250h
200h
160h
120h
80h
40h

101
100
10-1
10-2
10-3
10-4
10-5
1

(e)

324h
167h

0h

5 10
50 100T (K)

X=Cu(NCS)2

1

5 10

(d)

0h

0h

880h

170h
730h

400h

0.064
0.062

0h

0.06

50 100
T (K)

1

3

10 30
T (K)

100 3001

X=Cu[N(CN)2]Br X=Cu[N(CN)2]Cl

metal/superconductor

critical point

880h

200

250

300

5 10
50 100T (K)

X=Cu2(CN)3

Mott insulator U /W
dim

Fig. 1 Temperature dependence of the dc resistivity of κ-(BEDT-TTF)2 X irradiated by X-rays, (a) X = Cu(NCS)2, in which theresidual resistivity increases with irradiation dose; (b) Cu[N(CN)2]Br, in which an Anderson localization insulating state
appears from a strongly correlated metal state after a critical irradiation dose; (c) Cu[N(CN)2]Cl, in which the Mott insulatingstate changes to an Anderson localization insulating state; and (d) Cu2(CN)3, in which the Mott insulating state is robust
to introduction of molecular defects. The time indicated in each panel is the total X-ray exposure time at room temperature.
(e) Schematic phase diagram of κ-(BEDT-TTF)2 X for Udim/W.

43

IMR KINKEN Research Highlights 2012

IMR KINKEN Research Highlights 2012(d) Cu2(CN)3, which were irradiated by X-rays atroom temperature. The series is arranged in order ofthe correlation strength Udim/W, as shown in Fig.
1(e), where Udim is the onsite Coulomb interactionenergy of a BEDT-TTF dimer and W is the bandwidth.
The metal–Mott insulator transition occurs between
X = Cu[N(CN)2]Br and Cu[N(CN)2]Cl. X-ray
irradiation of the samples changes the resistivity
drastically. In the case of the superconductor X =Cu(NCS)2, the characteristic hump structure around
100 K is suppressed. In contrast to X = Cu(NCS)2, a
rather low irradiation dose for the X = Cu[N(CN)2]Br
sample induces a drastic change in the temperaturedependence of the resistivity [5]. The resistivity
hump at 100 K is completely suppressed, and thenthe resistivity curves show an insulating behavior at
low temperature. The change in the resistivity clearly
indicates that the observed metal–insulator transition
is induced by disorder, i.e., an Anderson localization
insulator by X-ray irradiation. On the other hand, in
the Mott insulators X = Cu[N(CN)2]Cl and Cu2(CN)3,
the resistivity decreases but shows an insulating
behavior with X-ray irradiation. This change
demonstrates a transition from Mott insulator toAnderson localization insulator. In this manner, therandomness from the molecular defects strongly
affects the electronic states of κ-(BEDT-TTF)2X. The
effect becomes especially significant in X =Cu[N(CN)2]Br and Cu[N(CN)2]Cl, which are in thevicinity of the critical point for the Mott transition.
A schematic electronic phase diagram with theparameters of temperature, band width, and disorderis depicted in Fig. 2 [5,6]. The present X = Cu[N(CN)2]
Y (Y = Br and Cl) samples are located near the Mott
transition on the metallic and Mott insulator sides,
respectively. The introduction of disorder changes a
Fermi liquid (FL) metal to an Anderson-type
localization insulator at a critical disorder value
characterized by the resistivity value of ρc. In thecase of the superconductor X = Cu(NCS)2, theelectron correlation is not strong enough to cause
localization of carriers at the same amount ofdisorder as in X = Cu[N(CN)2]Br. In the Mott insulator
X = Cu[N(CN)2]Cl, the randomness turns the Mott
insulating state to an Anderson localization insulatingstate at a similar ρc value. The most importantsuggestion in the present phase diagram is thecritical behavior at the Mott transition point for thedisorder. A finite small value of disorder in the system

Fig. 2 Schematic electronic phase diagram of κ-(BEDTTTF)2X for temperature T, irradiation time tirr
corresponding to disorder, and pressure P to bandwidth.
The thick dashed red lines indicate the critical resistivity
ρc between the metal (FL) and the localization insulator.
The temperature dependence of the resistivity shows ametallic ρ ~ T2 dependence in the FL metal. Note that a
finite value of disorder suppresses the first order Mott
transition.

modifies the pure Mott transition to Anderson
localization. Several phase diagrams calculated on
the basis of the disordered Hubbard model havebeen reported. For improved understanding of thecorrelated electronic states with disorderapproaching the Mott critical point, theoreticalsuggestions in addition to experimental investigations
are urgently needed in order to obtain a proper
physical picture for the future.
References[1] N. F. Mott, Metal–Insulator Transitions; Taylor and Francis:
London, UK, (1990).
[2] M. Imada, A. Fujimori and Y. Tokura, Rev. Mod. Phys. 70 1039
(1998).
[3] P. W. Anderson, Phys. Rev. 109, 1492 (1958).
[4] T. Sasaki, H. Oizumi, Y. Honda, N. Yoneyama and N. Kobayashi,
J. Phys. Soc. Jpn. 80, 104703 (2011)[5] K. Sano, T. Sasaki, N. Yoneyama and N. Kobayashi, Phys.
Rev. Lett. 104, 217003 (2010).
[6] T. Sasaki, K. Sano, H. Sugawara, N. Yoneyama and N.
Kobayashi. Phys. Status Solidi B 249, 947 (2012)

Keywords: organic, electronic properties, defects
Takahiko Sasaki (Low Temperature Condensed State Physics Division)
E-mail: takahiko@imr.tohoku.ac.jp
URL: http://cond-phys.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201244

Electronic Materials

Detection of Misoriented Domains with a Preferential
Orientation in Nonpolar GaN Layers

Controlling microstructure, including misorented domains, in nonpolar GaN layers is an important issue inrealizing highly efficient optoelectronic devices. We investigated the microstructure of nonpolar GaN layers bytransmission electron microscopy (TEM). Investigating the crystallographic orientation relationship between
misoriented domains and the surrounding matrix, we efficiently distinguished misoriented domains distributed
in nonpolar GaN layers.
GaN-based wide-bandgap semiconductors
are favorable for short-wavelength light-emitting
devices. In GaN layers grown along the [0001] polar

Fig. 1 (a) Bright-field TEM image of the GaN layers. (b) A
diffraction pattern obtained from the area surrounded
by a circle in (a), showing the orientation relationshipbetween the misoriented domains and the surrounding
matrix. (c) Dark-field TEM image of the same area as(a), exhibiting the distribution of misoriented domains.
Arrows in (a) and (c) indicate the position of the interface
between the GaN layers and sapphire substrate.

45

IMR KINKEN Research Highlights 2012

axis of the wurtzite (WZ) structure, electric fields are
generated due to spontaneous and piezoelectric
polarization, resulting in detrimental effects on
device properties. In order to prevent or reducesuch effects, the growth of nonpolar or semipolar
layers has been developed. However, unlike (0001)
c-plane GaN layers, nonpolar or semipolar GaN
layers can easily contain domains with different
orientations [1,2]. Although controlling the formationand evolution of misoriented domains is an important
issue in the growth of nonpolar or semipolar GaN
layers, such misoriented domains are often
overlooked.
Here,byinvestigating
thecrystallographic orientation relationship between
misoriented domains and the surrounding matrix,
we efficiently revealed misoriented domains
distributed in nonpolar GaN layers.
GaN layers were grown by hydride vapor phase
epitaxy. After the nitridation of r-plane sapphire
substrates, 410 nm thick low-temperature (LT) GaN
layers were grown at 550 °C. Then, the LT GaN
layers were heated to 1040 °C. The annealed LT
GaN layers were confirmed to be dominantly a-planeoriented by X-ray diffraction [1,2]. Cross-sectional
specimens were investigated by TEM with various
incident beam directions.
Fig. 1(a) shows bright-field TEM images of the
GaN layers with a properly selected incident beam
direction, the [2203] direction of the a-plane oriented
GaN layers. Fig. 1(b) shows a diffraction patternobtained from the area surrounded by the circle inFig. 1(a). In Fig. 1(b), the diffraction spots linked bywhite lines are from the a-plane oriented GaN layers.
Besides those spots, an extra diffraction pattern
was detected. The diffraction pattern indicated bytriangles corresponds to that obtained with theincident beam parallel to the [1 120] direction (a-axis)of WZ GaN. Fig. 1(c) shows dark-field TEM images

IMR KINKEN Research Highlights 2012

Fig. 2 Schematic illustration of the experimentally observed preferential orientation relationship between the a-planeoriented layers and the misoriented domain. A green arrow parallel to the edge of a tetrahedron indicates the observation
direction to efficiently distinguish the misoriented domains.

of the same area as Fig. 1(a) with the incident beam
direction parallel to the [2203] direction of thea-plane oriented GaN layers. The dark-field images
were taken using the 1100 diffraction spot shown in
Fig. 1(b). As a result of the dark-field TEM
observations with this incident beam direction, it
was revealed that the misoriented grains were
distributed in the vicinity of the interface between
the GaN layers and the sapphire substrate.
Fig. 2 shows a schematic of the misoriented
domain and the a-plane oriented layers. In WZ GaN,
every atom is bonded to four neighboring atoms inthe shape of a tetrahedron. In the tetrahedron, thec-plane is distinguished from the other three planes
because the bond parallel to the c-axis (bc) and theother three bonds (ba) are not equivalent in the WZ
structure. Considering the orientation relationshipbetween the misoriented domains and the a-planeoriented layers by using the structure model, apossible formation mechanism for the misoriented
domains can be suggested as follows: the c-planesof the misoriented domains are formed on a “pseudoc-plane” of the a-plane-oriented layers. Here, the“pseudo c-plane” is the atomic plane perpendicular
to a ba bond. It has been indicated that theinequivalence between bonds ba and bc can be
negligible under a thermally unstable condition,
which is the origin of the misoriented domains [3].
Based on the formation mechanism of the misoriented
domains, we suggest a method to efficiently
distinguish misoriented domains: observing along a
direction parallel to an edge of the tetrahedron (oneof the a-axes of the misoriented domains, green
arrow in Fig. 2), we can find the misoriented domains
with strong electron diffraction intensity from low
index planes.
References[1] H. -J. Lee, J. -S. Ha, T. Goto, T. Yao, C. Kim, S. -K. Hong and
J. Chang, Appl. Phys. Lett. 94, 102103 (2009).
[2] H. -J. Lee, T. Goto, K. Fujii, T. Yao, C. Kim and J. Chang, J.
Cryst. Growth 312, 198 (2010).
[3] Y. Tokumoto, H. -J. Lee, Y. Ohno, T. Yao and I.Yonenaga, J.
Cryst. Growth 334, 80 (2011).

Keywords: film, nitride, TEM
Yuki Tokumoto (Physics of Crystal Defects Division)
E-mail: y.tokumoto@imr.tohoku.ac.jp
URL: http://lab-defects.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201246

Electronic Materials

Temperature Stability of Bandgap Energy of InN

Nitride semiconductors such as InN, GaN, AlN, and their alloys are very promising for fabricating devices
with low power consumption, high power, and high frequency. InGaN-based blue light-emitting diodes (LEDs)
and laser diodes (LDs) are already commercially available. New trends of nitride semiconductors that extend
into the deep ultraviolet, green, and infrared wavelength regions will promote device applications that saveenergy and create new energy. Our research is related to these new technology trends, challenging the nitride
semiconductor-based LDs for use in optical fiber communications systems, phosphor-free white LEDs, brilliant
green LEDs, and high-efficiency solar cells.

The temperature dependence of bandgap
energy is determined by the electron–latticeinteraction, according to the Bose–Einstein model:

47

IMR KINKEN Research Highlights 20128

α2 ( x 10 9 cm-2 )InN is a difficult material to grow owing to thehigh equilibrium vapor pressure between the solidand vapor phases of nitrogen. Recently, wedeveloped a pressurized-reactor Metalorganic
vapor phase epitaxy (MOVPE) (PR-MOVPE). We
found that the pressurized growth is beneficial for
facet control of the InN growth and implementationof dense films at high growth temperature [1], which
generally favors the formation of a stable crystalline
structure. In this report, we investigate thetemperature dependence of the bandgap energy of
InN films grown by PR-MOVPE.
The InN film was grown on a polished (bothsides) sapphire substrate at 625 °C under 800 Torr.
The flow rates of NH3 as a nitrogen source andtrimethylindium as an indium source were 9 slm and16 μmol/min, respectively. The InN film thickness
was about 300 nm. The carrier density was estimatedto be about 1×1019 cm-3 from the plasmon frequency
in the IR reflection spectra.
The absorption spectrum of the InN film isshown in Fig. 1. The squared absorption coefficient
can be well fitted by a linear function of the incident
photon energy in the region from 0.9 to 1.2 eV, basedon a direct interband transition. This leads to an
experimental value of the InN bandgap energy, 0.78
eV. In the inset of Fig. 1, it can be observed that a
red shift of the bandgap energy occurs withincreasing sample temperature. This shift is shownin detail in Fig. 2, with a relatively small temperaturedependence compared with other semiconductors,
such as InP, GaAs, and InAs [2].

3.1K
100K
200K
300K

6
4
2
0

0.75 Eg 0.85

0.6

0.8

1.0

Photon Energy (eV)

1.2

Fig. 1 Temperature dependence of the optical absorption
spectra of the InN film.

E g (T ) = E g (0) − 2α / [exp(θ / T ) − 1] ,where α is the electron–phonon interaction
strength and θ is the average phonon energy. By
fitting the experimental data, these parameters wereobtained and listed in Table 1. Compared with GaAs
and InP, since α is smaller and θ is larger for InN, thevalue of [Eg (T) - Eg (0)] / Eg (0) is smaller. The bandgap
energy and its temperature stability in InN are very
promising for fabricating laser diodes for opticalfiber communications systems.
Table 1 Electron–phonon interaction strength and phonon
energy of InN bandgap energy compared with othermaterials for infrared lasers.

α (meV)
θ (K)

InN
12.8
266

GaAs
57.0
240

InP
35.0
209

IMR KINKEN Research Highlights 2012

InN (Eg~2 eV)
AlN
GaN
InN
(ｔhis work)
InP
GaAs

Eg(T)-Eg(0)
Eg(0)

-0.05

-0.10

1.0
CB+BGR(WZ)

0.8

with ZB
inclusion
CB+BGR
(ZB)

0.6

InAs
-0.15

reflection
transmission

1.2

E g (eV)

0.00

0

100

200

300

400

Temperature (K)Fig. 2 Temperature dependence of the InN bandgap
energy compared with other materials.

10

18 2

4 6

19 2

4 6

10
-3
n e (cm )

10

20 2

4

Fig. 3 Comparison between the residual electron density
ne dependence of the InN bandgap energy E from IRtransmittance (open circles) and reflectance (closed
circles) spectra, where ne is a function of growthtemperature.

The absorption edge increases with the residual
electron density ne of an InN film by the Burstein–
Moss effect [3]. The relationship between ne and theabsorption edge energy was measured by opticaltransmission (open circles) and reflection (closed
circles), as shown in Fig. 3 [4]. Measurements weremade for InN films grown at different temperatures
Tg. In the case of InN grown above 600 °C, ne values
estimated by both transmission and reflection
methods are in good agreement with each other.
This result indicates that both the optical absorption
edge measured by transmission and the plasmon
frequency measured by reflection are directly
connected to ne if the crystallographic phase is awurtzite single phase. Moreover, ne is found to
decrease as Tg increases, indicating the improvementof the structural quality. However, in case of the InN
films grown below 575 °C, a large discrepancy
appeared between the ne values obtained from thetwo methods; this is probably explained by theincorporation of a zincblende phase and related
defects at the boundary between the two phases[5,6].
References[1] Y. Zhang, Y. Liu, T. Kimura, M. Hirata1, K. Prasertsuk, S. Ji, R.
Katayama, and T. Matsuoka, Phys. Stat. Sol. (c) 8, 482 (2011).
[2] Q. Guo and A. Yoshida, Jpn. J. Appl. Phys. 33, 2453 (1994).
[3] J. Wu, W. Walukiewicz, W. Shan, K. M. Yu, J. W. Ager III, E. E.
Haller, H. Lu, and W. J. Schaff, Phys. Rev. B 66, 201403(2002).
[4] K. Prasertsuk, M. Hirata, Y. Liu, T. Kimura, Y. Zhang, T.
Iwabuchi, R. Katayama, and T. Matsuoka, Phys. Stat. Sol. (c) 9,
681 (2012).
[5] T. Kimura, K. Prasertsuk, Y. Zhang, Y. Liu, T. Hanada, R.
Katayama, and T. Matsuoka, Phys. Stat. Sol. (c) 9, 654 (2012).
[6] T. Iwabuchi, Y. Liu, T. Kimura, Y. Zhang, K. Prasertsuk, H.
Watanabe, N. Usami, R. Katayama, and T. Matsuoka, Jpn. J.
Appl. Phys. 51, 04DH02 (2012).

Keywords: nitride, energy gap, optical properties,
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 201248

Electronic Materials

Magnetic Anisotropy and Chemical Order for L10-Ordered
FeNi Films Grown by Alternate Monatomic Layer Deposition

L

10 -FeNi films with large uniaxial magnetic anisotropy energy (Ku) were prepared by alternate monatomic
layer deposition, and Ku reached 7.0 × 10 6 erg/cm3. Long-range order parameters (S) were evaluated by X-ray
diffraction using synchrotron radiation, and the relation between Ku and S was investigated. It has been foundthat Ku is roughly proportional to S.
A perpendicularly magnetized film is used forthe recording layer in a hard disk drive, and is also
useful for spintronic devices such as magnetic
random access memory. Large uniaxial magnetic
anisotropy energy (Ku) is necessary for such aperpendicularly magnetized film to realize highdensity and stable recording. Fig 1 shows therelation between the uniaxial magnetic anisotropy
energy (Ku) and the saturation magnetization (Ms) of
typical materials for perpendicularly magnetizedfilms.
L10 -ordered CoPt and FePt alloys are
attracting attention due to their very large Ku andadequate saturation magnetizations (Ms). However,Pt is very expensive and deposits are restricted.
L10 -FeNi consisting of only ubiquitous elements also
has a large Ku of 1.3 × 10 6 erg/cm3, which wasreported in the bulk form by Paulevé et al. [1]; thus,
L10 -FeNi is a candidate for a noble-metal-freeperpendicularly magnetized film. However, thepreparation of L10 -FeNi is very difficult because
atomic diffusion is quite slow below the order–
disorder transition temperature (320 °C). Therefore,
L10 -FeNi with a large Ku had been obtained only byannealing with neutron irradiation under magnetic
fields in a nuclear reactor [1]. Recently, we havesucceeded in fabricating L10 -FeNi films by alternatemonatomic layer deposition, because the L10
structure consists of alternate stacking of two
different monatomic layers along the c-axis, whichcorresponds to the easy magnetization axis [2,3].
However, there were two major problems. One
problem was a ferromagnetic buffer layer thathindered precise evaluation of the Ku of L10 -FeNi.
The other was the difficulty in evaluating of the longrange order parameter (S) of L10 -FeNi by X-ray
diffraction (XRD) because of very weak superlattice
peaks due to the close atomic scattering factors of
Fe and Ni.
A buffer layer should not only be nonferromagnetic but also have a flat surface and goodlattice matching to L10 -FeNi (a≈c=0.358 nm). We
have found that the slight addition of Au for Cu
growth leads to the formation of an atomically flat
surface [4]. The lattice parameters of Au, Cu, and Ni
are 0.408, 0.362, and 0.352 nm, respectively.
Therefore, a non-ferromagnetic Au–Cu–Ni ternary
alloy with a small Au composition is suitable for the49

IMR KINKEN Research Highlights 2012buffer layer. The composition was preliminarily
optimized to Au0.06Cu0.51Ni0.43 by comparison of Ku
for FeNi films on Au–Cu–Ni buffers with various
compositions.
In order to obtain high intensities for thesuperlattice peaks of L10 -FeNi in XRD, the use ofsynchrotron radiation is favorable, and grazing
incidence geometry can exclude unnecessarydiffracted X-rays from the substrate and buffers.
Thus, grazing incidence XRD (GI-XRD) usingsynchrotron radiation was used to observe
superlattice peaks and evaluate S. In this study[5,6], we prepared L10 -FeNi films on Au0.06Cu0.51Ni0.43
buffer layers at several substrate temperatures (Ts),and evaluated S by GI-XRD. The relation between
Ku and S was examined by changing Ts.
Films were deposited by molecular beam
epitaxy. An Fe seed layer (1 nm) and Au (20 nm)
were deposited on a MgO(001) substrate at 80 °C,
and Cu (50 nm) was grown at 500 °C. The Cu andAu layers alloyed into Cu3Au [7]. Then the
Au0.06Cu0.51Ni0.43 buffer (typically 50 nm) was grown
at 100 °C by ternary co-deposition. Finally, Fe andNi monolayers were alternately grown at Ts = 40, 67,
100, 127, 157, 187, and 217 °C. The Fe and Ni
layering was repeated 50 times. Magnetic properties
were investigated by magnetization measurements
using a super-conducting quantum interference
device (SQUID) magnetometer at room temperature.
The GI-XRD measurements to estimate S werecarried out at the BL46XU beamline in SPring-8.
Epitaxial growth for all the films was confirmedby reflection high-energy electron diffraction(RHEED) and GI-XRD. The (110) L10 -FeNi superlattice
peak was clearly observed by the in-plane GI-XRD.
From the XRD patterns, S was evaluated by(1)where Isup and Ifund are integrated intensities forthe superlattice peak and the fundamental peak,
respectively. The denominator of eq. (1) is thecalculated value when S = 1. The largest S of all the
FeNi films prepared was found to be 0.48.
Magnetization curves were measured when themagnetic field was applied in parallel andperpendicular to the film plane. Although the easymagnetization axis for all the films was the in-plane
direction because of a large demagnetizing field
due to the shape anisotropy of the film (2πMs2 =˜ 9×
10 6 erg/cm3), the Ts = 187 °C film showed large
magnetic anisotropy in the perpendicular direction.
From the magnetization curves, Ku was estimated
by


(2)

where H and M are the magnetic field and themagnetization, respectively. For the Ts = 187 °C film,
the largest Ku obtained was 7× 10 6 erg/cm3, whichis larger than those of artificially synthesized FeNi
films reported previously.
Fig. 2 shows the relation between Ku and S. Ku
is roughly proportional to S. By extrapolation, it isconsidered that Ku would be beyond the shapeanisotropy around S > 0.65, showing perpendicularmagnetization.
In summary, we prepared L10 -FeNi films on
Au0.06Cu0.51Ni0.43 buffer layers at substratetemperatures in the range from Ts = 40 to 217 °C.
The largest Ku and S obtained were 7.0 × 10 6 erg/

Ku [erg/cm3]

108

FePt
MnAl

107

CoPt

Ku [10 6 erg/cm3]

IMR KINKEN Research Highlights 2012

Shape anisotropy
2π Ms2S
Fig. 2 The relation between Ku and the long-range orderparameter S for L10 -FeNi films grown by alternatemonatomic layer deposition. The shape anisotropy
˜ 9 × 10 6 erg/cm3. Extrapolationenergy is 2πMs2 =indicates that Ku would be beyond the shape anisotropy
aroundS>
0.65,
showing
perpendicularmagnetization.

cm3 and 0.48, respectively. Ku was roughly
proportional to S. Perpendicularly magnetized L10 FeNi films could be obtained with a further increaseof S.

L10-FeNi

Co-Cr-Pt

References[1] J. Paulevé, A. Chamberod, K. Krebs, and A. Bourret, J. Appl.

106

Phys. 39, 989 (1968).

Ordering

105[2] T. Shima, M. Okamura, S. Mitani, and K. Takanashi, J. Magn.
Magn. Mater. 310, 2213 (2007).

Ni[3] T. Kojima, M. Mizuguchi, and K. Takanashi, J. Phys.: Conf. Ser.

Fe

104
0

266, 012119 (2011).

fcc-FeNi
5

10

15

20

Ms2 [105 (emu/cm3)2]
Fig. 1 The relation between the uniaxial magnetic
anisotropy energy (Ku) and saturation magnetization(Ms) of typical materials for perpendicularly magnetizedfilms. The inset shows the structure of L10 -FeNi. Ku of
FeNi is strongly enhanced by ordering to L10 structurefrom fcc. L10 -FeNi has large Ku and Ms even without
noble metals.

[4] T. Kojima, M. Mizuguchi, and K. Takanashi, to be submitted.
[5] M. Mizuguchi, T. Kojima, M. Kotsugi, T. Koganezawa, K.
Osaka, and K. Takanashi, J. Magn. Soc. Jpn. 35, 370 (2011).
[6] T. Kojima, M. Mizuguchi, T. Koganezawa, K. Osaka, M.
Kotsugi, and K. Takanashi, Jpn. J. Appl. Phys. 51, 010204(2012).
[7] M. Mizuguchi, S. Sekiya, and K. Takanashi, J. Appl. Phys. 107,
09A716 (2010).

Keywords: magnetic properties, spintronics, thin film
Koki Takanashi (Magnetic Materials Division)
E-mail: koki@imr.tohoku.ac.jp
URL: http://www.magmatex.imr.tohoku.ac.jp/J0_index.html

IMR KINKEN Research Highlights 201250

Electronic Materials

Field-reinforced Superconductivity in Uranium
Ferromagnets

F

erromagnetism and superconductivity have been thought to be mutually competitive phenomena, since thelarge internal field easily destroys the Cooper pair needed for conventional superconductivity. We discovered
a new class of superconductivity in uranium compounds, such as URhGe, UCoGe, and UGe2, which can
microscopically coexist with ferromagnetism. The triplet-state of pairing allows superconductivity with a large
upper critical field HC2. Surprisingly, superconductivity is reinforced by a magnetic field, in which both
ferromagnetic fluctuations and Fermi surface instabilities are induced.
is very large, and all curves show unusual S-shapes
or even field-reentrant behaviors. In general, Hc2 isgoverned by two different effects. One is the socalled Pauli paramagnetic effect, which is related to
the Zeeman spin splitting. The other is an orbital
effect that is related to the cyclotron motion of theelectrons. In conventional superconductors basedon the spin-singlet state, Hc2 is usually limited by thePauli paramagnetic effect, which suppresses Hc2 toa rather low value. However, in our ferromagnetic
superconductors, Hc2 markedly exceeds the Pauliparamagnetic limit since the spin-triplet state isrealized. Thus, only the orbital effect, which can be
expressed by the Fermi velocity or the effective
mass of the conduction electrons, governs Hc2 at
0 K. In this sense, a heavy fermion system, in whichstrong Coulomb interactions yield a large effective
mass, is a goldmine for realizing a huge Hc2. In fact,
all ferromagnetic superconductors in Fig. 1 reveal

The discoveries of superconductivity in the
ferromagnetic phases UGe2 [1], URhGe [2], andUCoGe [3] were spectacular events in the field ofunconventional superconductivity, although thepossibilityofsuperconductivity
nearthe
ferromagnetic quantum critical point was theoretically
predicted a long time ago. A spin-triplet pairing ispresumably generated by the ferromagnetic
fluctuation. It is established that in UGe2, URhGe,and UCoGe an equal spin pairing is realized. Moresurprisingly, field re-entrant (or reinforced)
superconductivity was found in UCoGe [4] andURhGe [5]. When a magnetic field is applied to thesample, the superconductivity is easily destroyed,
and this critical field is known as the upper criticalfield Hc2. Fig. 1 shows the temperature dependencesof the upper critical fields Hc2 in UGe2, URhGe, andUCoGe. Although the superconducting transition
temperature Tc is below 1 K, the value of Hc2 at 0 K

20
UGe 2
H // a-axis
1.35 GPa

Hc2 (T)

15

UCoGe
H // b-axis

URhGe
H // b-axis
SC

10

SC
5
0

SC
0

0.2

0.4

T (K)

0.6 0

SC

0.2

0.4

T (K)

0.6 0

0.2

0.4

0.6

T (K)

Fig. 1 Temperature dependence of the superconducting upper critical field Hc2 in the ferromagnetic superconductors,
UGe2, URhGe and UCoGe. All Hc2 curves reveal very unusual behaviors, owing to the spin-triplet state, with ferromagnetic
fluctuations at high fields [7]. "sc" represents the superconducting phase.

51

IMR KINKEN Research Highlights 2012

IMR KINKEN Research Highlights 2012

URhGe
UCoGe

50

(b)
Ferromagnetic
Superconductors

30

Hc2/Tc (T/K)(a)

CeRhSi3
CeIrSi 3
CeCoGe3

Noncentrosymmetric

20

UBe13

Heavy fermion

10
UGe2

ChevrelMxMo6S8

URhGe

UCoGe

Type II
Pauli limit

0

Hg Pb

1920

CePt 3 Si

CeCu2Si 2

Nb3Sn

Nb

1940

1960

1980

2000

Year

Fig. 2 History of the superconducting transition temperature Tc (a) and the upper critical field Hc2 scaled by Tc (b) [8].
Ferromagnetic superconductors based on the spin-triplet state show huge Hc2. The large values of Hc2 /Tc are realized by
the strong mass enhancement of conduction electrons and the longitudinal magnetic fluctuations.

enhancements of their effective masses caused byferromagnetic fluctuations under a magnetic field,
which has been clarified by our precise experiments
under extreme conditions using very high-quality
single crystals. Interestingly, the ferromagnetic
fluctuation is quite anisotropic with strong Ising
behavior; thus, Hc2 reveals a strong anisotropy aswell [6,7].
Fig. 2 shows the history of Hc2 scaled by Tc.
Thanks to the discovery of the first heavy fermion
superconductor almost 30 years ago, Hc2 isdramatically increased because the superconducting
Cooper pairs are mediated by magnetic fluctuations.
Presently, a new stage of superconductivity withhuge Hc2 has just started with actinide ferromagnetic
superconductors.

References[1] S.S. Saxena, P. Agarwal, K. Ahilan, F.M. Grosche, R.K.W.
Haselwimmer, M.J. Steiner, E. Pugh, I.R. Walker, S.R. Julian, P.
Monthoux, G.G. Lonzarich, A. Huxley, I. Sheikin, D. Braithwaite
and J. Flouquet: Nature 406, 587 (2000).
[2] D. Aoki, A. Huxley, E. Ressouche, D. Braithwaite, J. Flouquet,
J.-P. Brison, E. Lhotel and C. Paulsen: Nature 413, 613 (2001).
[3] N.T. Huy, A. Gasparini, D.E. de Nijs, Y. Huang, J.C.P. Klaasse,
T. Gortenmulder, A. de Visser, A. Hamann, T. Görlach and H.v.
Löhneysen: Phys. Rev. Lett. 99, 067006 (2007).
[4] D. Aoki, T.D. Matsuda, V. Taufour, E. Hassinger, G. Knebel and
J. Flouquet: J. Phys. Soc. Jpn. 78, 113709 (2009).
[5] F. Lévy, I. Sheikin, B. Grenier and A.D. Huxley: Science 309,
1343 (2005).
[6] D. Aoki, F. Hardy, A. Miyake, V. Taufour, T.D. Matsuda and J.
Flouquet: C. R. Physique 12, 573 (2011).
[7] D. Aoki and J. Flouquet: J. Phys. Soc. Jpn. 81, 011003 (2012).
[8] D. Aoki, W. Knafo, and I. Sheikin: arXiv:1204.5128

Keywords: actinide, ferromagnetic, superconducting
Dai Aoki (Radiochemistry of Metals Division)
E-mail: aoki@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201252

Electronic Materials

Detection of Two-photon Transitions of Argon by
Optogalvanic and Laser-induced Fluorescent Spectroscopy

By employing a simple experimental setup, two-photon excited optogalvanic and laser-induced fluorescence(LIF) spectra of argon were successfully observed in a Grimm-style glow discharge tube, which has been
widely applied to depth profiling of the elemental composition of various film-like samples by emission
spectrometry.

Optical excitation of nonmetallic atoms, such as
O, N, and H, in the ground state is very difficult
because the lowest resonance absorption of suchelements is in the deep vacuum ultraviolet (VUV)
region and the generation of tunable laser radiationin the VUV region is very difficult. However, nonlinear two-photon absorption occurs when an atom
is exposed to high peak-power radiation generated
from a conventional laser.
If two-photon excitation of atoms could be
easily detected, conventional tunable lasers would
be extensively applicable to spectrochemical
analysis.
In this study, optogalvanic (OG) and two-photon
laser-induced fluorescence (LIF) signals were used
to detect the two-photon transition of argon atoms,
as an example.
The optogalvanic (OG) effect originates fromthe change in impedance in a steady-state glow
discharge when the discharge medium resonantly
absorbs radiation. The radiation rearranges theexcited-state population, which in turn alters the net
ionization rate of the plasma. The change in the
discharge impedance may be detected as avariation in the discharge current (i.e., a voltagechange at the discharge electrode). This OG signal
can be used to detect a two-photon transition. The
fluorescence signal from the two-photon excitedenergy level can also be utilized in the detection ofa two-photon transition.
After evacuating a discharge chamber to below
3 Pa (2 × 10 -2 Torr), high purity argon was introduced
to 4 × 102 Pa (3 Torr). The pressure of the chamber
was measured with a Pirani gauge and a Baratronpressure gauge installed between an evacuation
port and a rotary pump. The discharge current was
adjusted to 3 mA using the constant current mode
53

IMR KINKEN Research Highlights 2012of a DC power supply (the discharge potential
ranged over 400–450 V). A pulsed Ti:sapphire laser
was pumped by the second harmonic (532 nm) of a
Nd:YAG laser. The line-width of the Ti:sapphire laser
was about 0.1 nm and the oscillation wavelengths
were between 735 and 850 nm. The output energyof the pulsed laser was between approximately 7
and 10 mJ, and the glow discharge plasma was
directly illuminated by the laser beam. Resonant
absorption of the laser beam in the argon plasma
was confirmed by observing an optogalvanic (OG)signal. The OG signal was detected as a voltage
drop through a register that was inserted in the
discharge circuit. No focusing of the laser beam
was needed to observe the two-photon excited
optogalvanic (OG) and laser-induced fluorescence(LIF) spectra of argon. To detect two-photon LIF,
firing of the laser pulse was monitored by a biased
silicon photo detector, and the gate of an ICCD
detector equipped with a spectrograph was opened
to overlap the exciting laser pulse. A high-energy
Ti:sapphire laser mirror reflecting between 740 and
860 nm was used to remove the strong backscattered
Ti:sapphire laser radiation from the Cu cathodesurface of the Grimm-style glow discharge tube.
Two photon excited optogalvanic spectra ofargon [1]About 29 two-photon optogalvanic peak lines
were observed in the wavelength region of 735–850
nm. The two-photon OG spectra of argon were
easily observed without focusing the pulsed laser
irradiation to more than 1 mJ. The excitation
mechanisms of the observed two-photon OG
spectra could be classified into 2 types: two-photon
direct transitions and two-photon stepwise
transitions.
A partial energy-level diagram of several

IMR KINKEN Research Highlights 2012

16

16

Ionization limit

Ionization limit

2
2
1

5

3p 4p

12

3p 4s

13

641.63
1

12
5

0

3p 4s

1

2

5

3p 4p
753.39

1

0

2

14

839.64802.96834.54760.21

13

J

5

3p 6s

795.66

14

Energy / eV

Energy / eV

3p 4d

5

15

2
3
1
0

5

641.63J

15

2Metastable

Metastable

Fig.1 Partial energy-level diagram of the two-photon

Fig.2 Partial energy-level diagram of the two-photontransition optogalvanic spectra. J is total angulartransition and optically allowed transition of the laser-

momentum.

inducedfluorescence(LIF).

J

is

total

angularmomentum.

observed strong two-photon OG peaks (two-photon
direct transitions) is shown in Fig. 1.
Two photon laser-induced fluorescence spectraof argon [2]The
two-photon
excitedlaser-induced
fluorescence (LIF) of argon at 641.63 nm wasobserved with laser excitation at 753.39 or 795.66
nm. After the two-photon excitation, the opticallyallowed transition at 641.63 nm took place. A partial
energy-level diagram related to the LIF peak at
641.63 nm is shown in Fig. 2.
Although many two-photon transitions of argon
were observed in the optogalvanic study of theargon glow plasma [1], only the 641.63 nm LIF peak
could be observed in this study. This is probably
due to insufficient sensitivity of the present
experimental setup to observe the other LIF spectraof argon.
In conclusion, two-photon transitions of argon

atoms could be detected by optogalvanic (OG) andlaser-induced fluorescence (LIF) spectroscopy. In
this case, an ultraviolet (UV) transition could be
detected when using near-infrared (near-IR) laser
radiation.
By utilizing a two-photon excitation of an atom,
conventional tunable lasers will be more applicableto spectrochemical analysis.
References[1] H. Matsuta, K. Wagatsuma and K. Kitagawa, Anal. Sci. 26, 25(2010).
[2] H. Matsuta and K. Kitagawa, Spectrosc. Lett. 45, 13 (2012).

Keywords: optical absorption, electronic structure
Hideyuki Matsuta (Analytical Science Division)
E-mail: matsuta@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201254

Electronic Materials

Spin Degrees of Freedom in Carrier-Doped
Antiferromagnetic Mott InsulatorW

hether high-Tc superconductivity arises from a Mott insulator or a Fermi-liquid metal is a vital issue instrongly correlated electron systems. To clarify the nature of spin correlations in cuprate oxide superconductors,
neutron scattering measurements give indispensable information. Our recent study on two single-layer systems,
Bi2+xSr2-xCuO6+y and La2-xSrxCuO4, revealed the variation of two spin degrees of freedom with doping: localized
spin and itinerant spin natures. Therefore, there exists a transition or crossover between the strong and weakcouplings against doping, and the optimal superconductivity appears in the intermediate region.

Universal doping-evolution of spin correlations
in single-layer cuprate oxides
The doping evolutions of spin correlations insingle-layer LSCO and double-layer YBa2Cu3O6+d(YBCO) have been extensively studied, due to theavailability of high-quality and sizable single crystals.
However, there are only two systems for which themagnetic properties can be measured as a functionof doping energy and temperature. Therefore,
appropriate reference systems are highly required.
Quite recently, we have succeeded in growing alarge singe crystal of Bi2+xSr2-xCuO6+y (Bi2201) andin observing the spin fluctuations in this system [2,4].
Figure 1 shows the constant energy spectra at 4–6
meV for four Bi2201 samples with different hole55

IMR KINKEN Research Highlights 2012

/5.0minBi2+xSr2-xCuO6+yω=4-6 meV
200
100

150
100

(a) p=0.01

T=15K(b) p=0.06

T=40 K

x=0.5

x=0.4

50
0

/22.0minCounts/43.6min.

0

T=40 K

500
400
300
0

/21.6min.

Mechanisms of exotic superconductivity are afascinating issue in modern physics research. In
high-transition-temperature (high-TC) cuprate oxide,
the role of spin fluctuations in the emergence ofsuperconductivity has been intensively discussed.
Therefore, clarifying the nature of spin fluctuations in
doped Mott insulators is important. [1] The conceptof fluctuating spin and charge stripes, which isbased on the localized spin picture, provides anappealing explanation of the magnetic fluctuationsin the doped system. However, there is an alternative
school of thought that the spin fluctuations are
explained through Fermi-surface nesting. Althoughthese pictures provide different ways to understandthe physical properties in high-Tc superconductors,
there has not yet been a successful experimental
discrimination between them.
To gain a deep insight into the nature of spin
fluctuations, we have performed neutron scattering
measurements on a couple of new target systems.
Here, we introduce the recent progress in our studyof spin correlations in single-layer Bi2+xSrxCuO6+y(Bi2201) [2] and Fe-doped La2-xSrxCuO4 (LSCO)[3].

(c) p=0.09
x=0.3

T=70 K

500
400
0
-0.4

(d) p=0.12
x=0.2
-0.2

0

0.2

0.4

Mod-q (r.l.u.)
Fig. 1 Inelastic neutron scattering spectra at 4–6 meV in
Bi2.2Sr1.8CuO6+y with p = (a) 0.01, (b) 0.06, (c) 0.09, and(d) 0.12. Each spectrum was measured along the [1,-1,0] and [1,0,0] directions for (a,b) and (c,d),respectively.

concentrations (p). With increasing p, a single peak
splits into two incommensurate (IC) peaks, and the
distance between the peaks (2δ) increases while
maintaining a linear relation between δ and p.
Furthermore, the incommensurate modulation wave
vector (Q δ) was found to change from tetragonal[1,1,0] to the [1,0,0]/[0,1,0] direction, and we
confirmed that the rotation of Q δ occurs around thespin-glass and superconducting phases. These
results are quantitatively consistent with theobservations in LSCO. Therefore, we concluded
that the doping evolution of IC spin correlation is

IMR KINKEN Research Highlights 2012

(a)

2δSO

k
1/2

0.15

δSO (r.l.u.)

Cross over from localized spin picture toitinerant electron picture with hole doping
The impurity effect on spin correlations has
been extensively studied in connection with themechanisms of suppression of superconductivity. It
is known that by impurity doping the existing
magnetic order in the underdoped (UD) LSCO (x <0.16) can be enhanced with suppression of thesuperconductivity [7]. The simple explanation forthis stability of magnetic order is that the stripe
fluctuations are pinned statically by local defects.
However, the existence of stripe fluctuations in theentire SC phase is still under debate.
Taking advantage of the remarkable effect on
the stability of the magnetic order by Fe doping, weinvestigated the static spin correlation in Fe-doped
LSCO (Fe-LSCO). In the heavily overdoped (OD)
region, we clarified that only 1% Fe doping induces
a static IC magnetic order with a relatively large δ
below ~20 K. Figs. 2(a) and (b) show the holeconcentration dependence of δ and the magnetic
ordering temperature in the LSCO family [8, 9]. Bluecircles represent the results for Fe-LSCO. The overall
δ–p relation exhibits a kink at p ~ 0.12, indicating a
different nature of spin correlations separated by p
at ~0.12. Our angle-resolved photoemission
spectroscopy (ARPES) experiments on Fe-LSCO
revealed that the doping dependence of δ in the OD
region is well reproduced in term of Fermi-surface
nesting [3]. This agreement between neutron
scattering and ARPES results supports the itinerantspin picture, and spins of itinerant doped holes
define the magnetic ordering ground state in the OD
region.
Since the spin correlation in the UD region iswell characterized by the localized spin picture, theevidence of itinerant-spin character in the OD region
suggests a transition or crossover between the two

0.20

1/2

h
LSCO
LBCO
Nd-LSCOFe-LSCOδCO LBCO

0.10
k

0.05

2δSO

1/2
1/2

0

Tm (K)

common over a wide hole–concentration range ofsingle layer cuprate. These characteristic features ofspin correlation for x ≤ 0.12 are well reproduced by
the stripe and spin spiral models based on thelocalized spin picture [5,6].

h

60

(b)

40
20
0

0

0.1

0.2

p (x-y)0.3

Fig. 2 (a) Hole concentration (p) dependence of (a) the
incommensurability of magnetic order evaluated at low
temperature and (b) the magnetic ordering temperature.
The gray area represents the p region where CDW order
was observed in the LSCO family. Results for La2x(Sr,Ba)xCuO4 and La1.6-xNd 0.4 SrxCuO4 aretaken from

Refs. 8 and 9, respectively. In both figures, solid lines
are guide to the eye.

spin degrees of freedom with increased hole
doping.
References[1] M. Fujita et al., J. Phys. Soc. Jpn. 81 01107 (2012).
[2] M. Enoki et al., arXiv1205.3301 (submitted to Phys. Rev. Lett.)[3] H. Ruihua and M. Fujita et al., Phys. Rev. Lett. 107 127002(2011).
[4] M. Enoki et al., J. Phys. Soc. Jpn. 80 SB026 (2011).
[5] G. Seibold et al., Phys. Rev. B 83 205108 (2011).
[6] Q. P. Sushkov et al., Phys. Rev. B 79 174519 (2009).
[7] M. Fujita, Physica C 481 23 (2012).
[8] K. Yamada et al., Phys. Rev. B 57 6165 (1998); M. Fujita et al.,Phys. Rev. B 70 104517 (2004); S. R. Dunsiger et al., Phys. Rev.
B 78, 092507 (2008).

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

IMR KINKEN Research Highlights 201256

Electronic Materials

Development of Ce-doped Gd3(Al,Ga)5O12 (GAGG)
Scintillator and its ApplicationsGAGG crystals have been developed in our laboratory through collaboration with industry. This scintillator

has a high light output of more than 60,000 photons/MeV, a decay time of approximately 80 ns, and no
hygroscopic nature. This scintillator has been successfully applied to a dosimeter and it will be used in Compton
cameras or other detectors and monitors.

Introduction
A radiation detector consists of a scintillator andphoto-detector. The scintillator is the key material todetermine the performance of the radiation detector;
it works as an energy converter of keV or MeV
radiation particles (alpha rays, gamma rays) into
photons with energies of a few eV (UV, vis. light).
Radiation detectors are applied in X-ray
computed tomography (CT), positron emission
tomography (PET) and other medical imaging
techniques, high energy and nuclear physics
detectors, natural resource surveys, and so on. In
order to obtain high detection sensitivity, high-lightoutput scintillators are required. Ce:LaBr3 hasrecently been proposed and showed a high light
output of ~60,000 ph./MeV; however, it is strongly
hygroscopic [1]. A Ce:Lu2SiO5 (Ce:LSO) scintillator
also has a good light output of 30,000 ph./MeV,
while Ce:LSO has an intrinsic background from Lu176 (natural abundance of 2.59%) [2].
Crystal growth and its scintillation properties
We have investigated a novel scintillator with ahigh light output and no hygroscopic nature using a
micro-pulling-down (μ-PD) method [3, 4] developedin our laboratory. This method enables the growth of
one crystal within approximately 10 h, whileconventional methods such as Czochralski (CZ) or
Bridgeman techniques require much longer growthtimes of about one week or more. Optical properties(transmittance, emission, and excitation spectra),and scintillation properties (radio-luminescence,pulse height, decay time, etc.) were measured with
our equipment [5, 6]. This prompt and complete
cycle allows for quick material screening.
On the way to improve the light yield of Pr:LuAG,
we found Ga substitution is effective. We haveexpected similar positive effect on Ce doped
sample. Consequently, Ce:GAGG crystals achieving
a high light output and no hygroscopic nature weredeveloped. Additionally, the maximum emission
wavelength of GAGG was 520 nm, which is in thehigh-detection efficiency region of Si detectors,
such as avalanche photodiode (APD). This emission
is ascribed to the 5d–4f transition of Ce3+, and the
decay time was 80 ns [7].
After the study of GAAG prepared by the μ-PD
57

IMR KINKEN Research Highlights 2012

method, a floating zone method was applied toinvestigate the effect of the growth atmosphere [8].
A bulk crystal growth process was also developed.
Until now, the CZ method was used to grow 2 inch
size GAGG crystals (Fig 1, b), which showed a highlight output of 60,000 photons/MeV and FWHM
energy resolution of 5.4% at 662 keV, as measured
by a Si-APD (S8664, Hamamatsu).
Applications
Using a GAGG scintillator, a novel dosimeter
was developed as a product named “gamma
spotter” by FURUKAWA Co., Ltd (Fig. 1, c). Since
the GAGG scintillator has high sensitivity, thisdetector can measure a dose of less than 1 s. The
gamma spotters are used by many consumers after
the FUKUSHIMA accident.
The GAGG scintillator also could be applied in
medical imaging or astronomy. The GAGG isespecially advantageous for application in aCompton camera consisting of a Compton scatterpart and a scattered gamma-ray absorber, because
the angular resolution of a Compton camera
depends on the energy resolution of the absorber.
The scintillator could be used as the absorber due
to its superior detection efficiency over semiconductors. Additionally, due to the lack of an
intrinsic background, a Compton camera mountedwith a GAGG scintillator would obtain high quality
images. Recently, a national project supported by
JST has been started with Univ. of Tokyo, JAEA andFurukawa Co.Ltd., in order to develop a GAGG
Compton camera. Pollution mapping in the
Fukushima area could be conducted by a GAGG
Compton camera mounted on an unmanned
helicopter.
References[1] E. V. D. van Loef, P. Dorenbos, C. W. E. van Eijk, K. W. Kramer
and H. U. Gudel. Scintillation Nuclear Instruments & Methods
in Physics Research Section A 486, 254-258 (2002).
[2] C. L. Melcher and J. S. Schweitzer. Cerium-Doped Lutetium
Oxyorthosilicate - a Fast, Efficient New Scintillator. IEEE
Transactions on Nuclear Science, 39, 502-505 (1992).
[3] A. Yoshikawa, T. Satonaga, K. Kamada, H. Sato, M. Nikl, T.
Fukuda, J. Cryst. Growth 270, 427-432 (2004).
[4] A. Yoshikawa, M. Nikl, G. Boulon, T. Fukuda., Opt. Mater. 30,

IMR KINKEN Research Highlights 2012

Fig. 1 Schematic view of the development cycle of the scintillator (a), picture of a GAGG scintillator with a 2 inch diameter(b), and gamma spotter (c).

6-10 (2007).
[5] K. Kamada, T. Yanagida, J. Pejchal, M. Nikl, T. Endo, K.
Tsutumi, A. Yoshikawa, Cryst. Growth & Design 11, 4484–4490(2011).
[6] K. Kamada, T. Yanagida, J. Pejchal, M. Nikl, T. Endo, K.
Tsutumi, Y. Fujimoto, A. Fukabor and A. Yoshikawa, J. Phys. D:

Appl. Phys. 44, 505104 (2011).
[7] K. Kamada, T. Yanagida, T. Endo, K. Tsutumi, Y. Usuki, M.
Nikl, Y. Fujimoto, A. Yoshikawa, J. Cryst. Growth, in press.
[8] A. Yoshikawa, Y. Fujimoto, A. Yamaji, S. Kurosawa, J. Pejchala,
S. Wakahara, Y. Yokota, K. Kamada, T. Yanagida, K. Yubuta, T.
Shishido, M. Nikl, Opt. Mat. accepted.

Keywords: crystal growth, luminescence, micro-pulling-down method
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 201258

Research Centers

Advanced Evaluation of Radiation Effects on Fusion Materials
International Research Center for Nuclear Materials ScienceAthermal desorption spectrometer with an ion-implantation gun and 15.5 T superconducting magnet systemwere installed in a radiation-controlled laboratory in IRCNMS. Samples irradiated by nuclear reactors are testedby these devices to study radiation effects on fusion materials. These devices are open for collaborative works
not only with Japanese researchers but also with foreign researchers.
A new framework of bidirectional collaborative
research program has been started in 2010 for the
study of nuclear fusion reactor materials. By utilizing
the environment of the research center at each
university, joint research is conducted interactively
between the centers and the National Institute of
Fusion Science (NIFS). Two topics that wereconducted in this framework are described in thispaper.
The first topic is a study of a trapping mechanismof hydrogen Isotopes in neutron-irradiated plasmafacing materials [1]. One feature of fusion reactor
environments is that helium and hydrogen are
produced by nuclear transmutation, and thetransport and retention of hydrogen isotopes andhelium from the core plasma and tritium generated
at the blanket occur under neutron irradiation. It is
hence indispensable to clarify the effects of neutron
irradiation on the behavior of hydrogen isotopes andhelium in the candidate materials to assess thefeasibility of their use in fusion reactors.
For this, in 2010 a thermal desorption
spectrometer (TDS) apparatus (Fig. 1) was installed
in a radiation-controlled laboratory at the International
Research Center for Nuclear Materials Science(IRCNMS). This apparatus allows us to obtainthermal desorption spectra, thereby enabling theidentification and quantification of hydrogen isotopes
and helium contained in radioactive materials
following reactor or accelerator irradiation. In 2011,
the TDS apparatus was equipped with an ion gun to
inject hydrogen isotopes or helium into neutronirradiated specimens.
By using this apparatus, the emission andretention behavior of D was measured in recrystallized
tungsten (W) specimens that were first irradiatedwith 2.8 MeV Fe ions to 3 dpa (displacement peratom) and then exposed to D atoms with very low
energy and low flux (fluence: 6.2×1022 Dm-2); theresult is shown in Fig. 2. For comparison, the resultsfor W specimens irradiated with neutrons or Fe ions59

IMR KINKEN Research Highlights 2012

Fig. 1 Thermal desorption spectrometer (TDS) apparatus

Fig. 2 TDS results on D desorption rate in recrystallized
pure W with different irradiation (Fe ions or neutrons)
and D exposure (atoms or plasma (TPE))

and then exposed to TPE (tritium plasma experiment:high energy, high fluence 1025 D m-2) are included.
In the case of TPE exposure, neutron irradiatedspecimens exhibit D desorption from relatively low
temperatures to as high as above 1000 K, whereas
Fe-ion-irradiated specimens exhibit the completionof D desorption at around 700 K. In the case of
specimens with Fe-ion irradiation and D atom
exposure, D desorption starts at around 500 K andcontinues even above 1000 K. The D desorption
behavior suggests the existence of trapping sites forIMR KINKEN Research Highlights 2012hydrogen isotopes that are as strong as those inneutron irradiated specimens (de-trapping energy:~2 eV). These results indicate that 2.8 MeV Fe-ion
irradiation can induce strong trapping sites (detrapping energy: ~2 eV) for hydrogen isotopes andthat such strong trapping sites survive under lowflux D exposure but disappear under high-flux D
exposure. The de-trapping energy of ~2 eV
corresponds to chemisorption entrapment ofhydrogen isotope atoms at the inner surface of voids
in W.
The second topic is the neutron-irradiation
effect on the superconducting magnet systems of
fusion devices [2]. A fusion reactor yields manyhigh-energy neutrons, and some of them stream outof the plasma vacuum vessel, penetrate the blanket
system, and reach the superconducting magnets
that provide high magnetic fields to confine highenergy ionized particles. Under neutron irradiation,
the magnet materials are activated and theirproperties may change.
Changes in the superconducting properties ofthe superconducting wires and degradation of theinterlaminar shear strength (ILSS) of electric
insulation materials are the most concerning possible
irradiation effects. The superconducting magnet
system (Fig. 3), which was installed in the radiation
control area of IRCNMS, is a so-called dry magnet
working without liquid helium. The magnet generates
15.5 T as the maximum field in a 52 mm RT bore,
and it is equipped with Gifford–McMahon (GM)refrigeration with an air-cooled compressor . The
magnet system requires only electric power.
A variable temperature insert (VTI) included in
the system is a kind of sample holder that has acapacity of 500 A for a sample current. The sample
holder is cooled down by thermal conduction withGM refrigeration, and the sample temperature canbe controlled from 4.5 K to around 20 K at 500 A.
The control system consists of amplifiers and acomputer. The control computer is connected toanother computer located in a non-radiation-control
area by a dedicated internet line. Users are able tochange the test conditions from the non-radiation
room once an irradiated sample is set up. The 15.5

Fig. 3 15.5 T superconducting magnet (lower) and variabletemperature insert (upper) in a magnetic field shieldstructure.

T superconducting magnet is placed and the VTI is
set up on the magnet. Samples irradiated in nuclear
reactors will be tested with the new facility, and very
important data on the critical current, criticalmagnetic field, and critical temperature will beobtained.
The two experimental devices described above
are open for collaborative works not only withJapanese researchers but also with foreign
researchers.
References[1] Masashi Shimada, Y. Hatano, P. Calderoni, T. Oda, Y. Oya, M.
Sokolov, K. Zhang, G. Cao, R. Kolasinski, J.P. Sharpe. J. Nucl.
Mater., 415, S667–S671 (2011).
[2]A. Nishimura, T. Takeuchi, S. Nishijima, K. Ochiai, G. Nishijima,
K. Watanabe, M. Narui, H. Kurishita, T. Shikama, J. Nucl.
Mater., 417, 1-3, 842-845 (2011).

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

IMR KINKEN Research Highlights 201260

Research Centers

Materials for Highly Energy Efficient Electrical Devices/machines
Advanced Research Center of Metallic Glasses

We at the Advanced Research Center of Metallic glasses (ARCMG) are focusing on the development of newmaterials that have important functional properties, such as soft/hard magnetism, excellent mechanicalproperties, and easy fabrication abilities. The basic route to obtain these new materials is from the molten state
to solidification in different forms, such as meta-stable glassy, nano-crystalline, and even single crystal form.
The aim is to develop new materials that can benefit our society and environment. To do so, ARCMG is wellequipped with advanced manufacturing and testing facilities for different materials. Researchers/students also
have opportunities to understand their experimental results with simulations using a supercomputer. Additionally,
ARCMG is open to various researchers working in the same or other universities/institutes, so the students
have opportunities to widen their horizons.

The most important problem/issue in today's
world is the creation and management of energy
without affecting our environment. Electrical energy
is the most important form of energy, and without it
our society will collapse. Electrical energy is mostly
obtained at the cost of our Earth's natural resources(coal, oil, etc.), and it sometimes involves dangerous
methods of production (nuclear power plants).
Therefore environmentally friendly methods ofproduction and proper management of electrical
power are very important. For example, in power
applications such as transformers and motors, nonoriented silicon steel and oriented silicon steel aremainly used because of their high magnetic flux
density (B), but the electrical power loss in the Si–
steel cores is still high. Core loss (W) in silicon steels
has been improved, but further improvement seems
impossible. With the newly developed soft magnetic
amorphous/glassy materials (in our and otherlaboratories), it is possible to further decrease theelectrical power loss, but these materials have lower
B. High B is important for reducing the size andweight of the electrical devices (which is necessary
to make them energy efficient). Therefore,
development and mass production (at low cost) of
nano-crystalline soft magnetic alloys that cansimultaneously exhibit high B similar to Si steel (~2.0
T) and low magnetic core loss (W) as amorphous
metals is ideal.
We at ARCMG are close to the development of
this dream material. The new Fe-rich nano-crystalline
soft magnetic material Fe85-86Si1-2B8P4Cu1 iscomposed of ~95 mass% Fe without rare-earth
metals [1]. This low-cost soft magnetic material
exhibits high magnetic flux density (1.8–1.9 T) alongwith extremely low core losses (1/2–1/3 of ordinarymaterials). Fig. 1 shows the changes in magnetic
core loss (W) at 50 Hz for the Fe85-86Si1-2B8P4Cu1
nanocrystalline alloys as a function of maximum
magnetic induction (Bm). The data for amorphous
Fe78Si9B13, highest-grade oriented Fe–3 mass% Si,
and non-oriented Fe–3 mass% Si and Fe–6.5
61

IMR KINKEN Research Highlights 2012

Fig. 1 Magnetic core losses (W) of nanocrystalline
FeSiBPCu alloys as a function of maximum magnetic
induction (Bm). The data for Fe78Si9B13 amorphous and
Fe–Si crystalline alloys are also shown for comparison.

mass% Si alloys are also shown for comparison.
These special nano-crystalline soft magnetic
materials are produced by the careful crystallizationof amorphous/hetero-amorphous metallic alloys.
Amorphous/glassy alloys are attractive because of
their unique random arrangements of atoms, which
result in excellent properties such as soft magnetism,high mechanical strength, and corrosion resistance.
Additionally, these metallic glasses have precision
micro-/nano-machining capabilities, which makethese materials very important for future
nanotechnology applications [2]. We have beendeveloping new Fe-based, Co-based, Cu-based,Ni- based, Ti-based, Zr-based, and Mg-based
glassy/amorphous alloys for different required
functionalities. In the development of these special
materials for energy saving or precision machining,alloy design and processing condition optimization
are very important. At ARCMG, we have state-ofthe-art facilities (automatic, semi-automatic, or
manual) to develop such materials (Fig. 2)
Besides the reduction of power loss in electricalIMR KINKEN Research Highlights 2012

Fig. 3 (a) SEM and (b) TEM images of the Li4Ti5O12 single
crystals grown from NaCl flux.

Fig. 2 (a) Automatic arc-melting furnace at ARCMG and(b) typical shape of the Fe-rich Fe85-86Si1-2B8P4Cu1 alloy
ribbon.

devices, storage of electrical power is also very
important for the next generation of hybrid/electricalcars. Lithium-ion rechargeable batteries are very
important. We at ARCMG are also focusing on thegrowth of single crystals for Li-ion batteries. Spineltype Li4Ti5O12 (m.p. 1288 K) single crystals fornegative electrodes in lithium-ion rechargeable
batteries were grown by a NaCl (m.p. 1074 K) flux
growth method in air [3]. As shown in Fig. 3, thegrown crystals had an octahedral shape with welldeveloped facets and a flat surface. The TEM image
in Fig. 3 indicated that the Li4Ti5O12 crystals were ofvery high crystallinity without point defects or
dislocations. Environmental damage and productcosts are thought to be greatly reduced in the flux
growth technique using NaCl flux.
In addition to experimental studies, it is also
possible to gain a theoretical understanding ofvarious amorphous and nanocrystalline alloys/
materials using computer simulations.
For example, a detailed study of the coolingprocess of Pd42.5Cu30Ni7.5P20 glassy alloy has been
done using first-principles calculations, and theresults were compared with the data obtained froma synchrotron X-ray diffraction experiment. The
calculated densities of states show noticeablechanges in the electronic structures ofPd42.5Cu30Ni7.5P20 compositions during a coolingprocess. The detailed analysis of the atomic andelectronic structure of the alloy in the liquid andglassy states reveals the formation of short-range
chemical order in the temperature range
corresponding to a non-Arrhenius temperaturedependence of viscosity [4]. The theoretical groupat ARCMG has a great deal of expertise in theapplication of modern theoretical methods foranalyzing important properties of experimentallydeveloped alloys or suggesting new alloys toexperimentalists, based on their theoreticalunderstanding.
ARCMG not only is engaged in new cuttingedge scientific research but also focuses on
industry-related issues in the development of these
new materials. The new Fe-based soft magnetic
nano-crystalline alloy developed in ARCMG is an
excellent example. Besides the excellent soft
magnetic properties (Bs ~1.8T; Hc ~3–5 A/m), thisalloy has low materials (only needs industry-grade
materials, without expensive materials such as rareearths) and production costs (can be produced inair), which are extremely attractive for thecommercialization of these alloys.
References[1] A. Makino IEEE Trans. Magn. 48, 1331 (2012).
[2] Y. Wang, P. Sharma and A. makino, J. Phys.: Cond. Mat. 24,
076004 (2012).
[3] K. Teshima et al., Cryst. Growth Des. 11 (2011) 4401.
[4] D. V. Louzguine-Luzgin, et al., J. Appl Phys. 110 043519(2011).

Akihiro Makino (Advanced Research Center of Metallic Glasses)
E-mail: amakino@imr.tohoku.ac.jp
URL: http://www.arcmg.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201262

Research Centers

In-field Heat-treatment Processing for New Materials
Development Using the Cryogen-free Hybrid Magnet
High Field Laboratory for Superconducting Materials

In order to realize a hybrid magnet with no need of liquid helium for operation, we have developed a cryogenfree hybrid magnet, consisting of an outer cryogen-free superconducting magnet and an inner water-cooled
resistive magnet. This easy-to-operate cryogen-free hybrid magnet is now operating for application-oriented
research in high fields up to 27.5 T and at high temperatures up to 1400 K in combination with an electric
furnace.
The High Field Laboratory for Superconducting
Materials (HFLSM) has concentrated on the three
key research fields of "High field magnet technology,""Materials science in high fields," and "Basic researchunder multiple extreme conditions" as cooperative
research in the Institute for Materials Research,
Tohoku University. The HFLSM has demonstrated
the first instances of various kinds of cryogen-freesuperconducting magnets in the world [1]. A newresearch field that is called magneto-science has
made great progress using a static high field
generated by one such unique cryogen-freesuperconducting
magnet.
Magneto-science
encompasses crystal growth, chemical reactions,biological reactions, and materials processing inhigh fields. It is especially important to focus onfunctional materials development from the melt inhigh fields.
In order to make an easy-to-operate hybrid
magnet system, we have constructed a cryogenfree 27.5 T hybrid magnet (28T-CHM), consisting ofan outer cryogen-free 8.5 T superconducting
magnet and an inner traditional water-cooled 19 T
resistive magnet. Fig. 1 shows 28T-CHM, which
generated 27.5 T in a 32 mm room-temperatureexperimental bore [2], and this cryogen-free hybrid
magnet surely demonstrates quite flexible machinetime for use in long-term experiments.
We also installed a differential thermal analysis(DTA) combined with 28T-CHM. The DTA
measurement furnace system consists of Pt crucibles
and an alumina support equipped with two pairs ofPt–PtRh thermocouples allowing the temperature
difference to be measured [3]. Sample powders anda reference sample of Al2O3 powders were
investigated. The maximum temperature wasassumed to be 1200°C when the temperature
increased in the argon atmosphere.
The motor systems of evolving hybrid electricvehicles require a working potential at hightemperatures above 200 °C. The coercivity decreaseof a NdFeB magnet has been eased by the additionof Dy [4]. However, this rare-earth metal is now
restricted as a strategic material, and in the future it
will become quite difficult to obtain important rare63

IMR KINKEN Research Highlights 2012

Fig. 1 World’s first cryogen-free 27.5 T hybrid magnet,composed of a 360 mm wide bore cryogen-free 8.5 Tsuperconducting outsert magnet and a 32 mm roomtemperature experimental bore 8 MW-19 T water-cooled
resistive insert magnet.

earth metals such as Dy. Therefore, the developmentof rare-earth-free bulk magnets at the temperature
range of 500–600 K is needed as innovative
materials science.
On the other hand, it is generally known that
MnBi has a high coercivity even at 500 K [5].
However, since Bi and Mn usually remain as
individual foreign phases, it is difficult to synthesize
single-phase MnBi, and as a result, the developmentof a high-quality MnBi synthesis process is required.
The first-order phase transition through the liquid
phase in high fields and the magnetic anisotropy on
the crystallographic axis for the MnBi system are
expected to produce various functional compositions
with in-field heat treatment. This means that themagnetic characteristics can be controlled by thein-field heat treatment. Therefore, it is necessary to
make an equilibrium phase diagram for the efficientinvestigation of materials development in highfields.

IMR KINKEN Research Highlights 2012

Bi–Mn samples with different compositions
were made by the method of arc-melting followed
by heat treatment. The sample composition wasdetermined by inductively coupled plasma opticalemission spectroscopy (ICP-OES). Diffraction peaks
from the low temperature phase of ferromagnetic
MnBi and Bi were observed in the Bi-richcomposition. On the other hand, diffraction peaks
from MnBi, Mn, and unreacted Bi were observed in
the Mn-rich composition.
The DTA curves for each composition were
measured under a zero field. The endothermic peak(TBi) from the appearance of the Bi–Mn liquid phasefrom the ferromagnetic MnBi phase and Bi wasobserved in all the compositions. This result
corresponds to the diffraction peak from Bi observedin the X-ray diffraction measurement. Since
unreacted Bi remained in the sample, theendothermic peak from Bi was also observed in the
Mn-rich composition. Moreover, an endothermic
peak (Tt) was observed due to a decomposition
reaction accompanied by a magnetic transformationof ferromagnetic MnBi + Bi–Mn liquid phase to
paramagnetic Mn1.08Bi + Bi–Mn liquid phase. In
addition, an endothermic peak (Tm) due to theperitectic reaction of Mn1.08Bi + Bi–Mn liquid phase→ Mn + Bi–Mn liquid phase was observed.
The DTA curve measured when a magnetic
field was applied to Bi–53at.%Mn is shown in Fig. 2.
One notes that the endothermic peaks of TBi and Tm
are unchanged in fields up to 18 T, while Tt increases
with applied magnetic field. However, the Tt increase
clearly deviates from linear behavior above 20 T, as
shown in Fig. 3 [6]. We found that the magnetic
energy of Mn1.08Bi at Tt increases nonlinearly in highfields above 20 T and becomes large, comparable
to that of MnBi. These results show that we cangenerally control the composition, decomposition
temperature, and the equilibrium state offerromagnetic magnetic materials by a high magnetic
field if the materials have a large magnetic
moment.
References[1] K. Watanabe, Y. Yamada, J. Sakuraba, F. Hata, C.K. Chong, T.
Hasebe and M. Ishihara, Jpn. J. Appl. Phys. 32, L488 (1993).
[2] K. Watanabe, G. Nishijima, S. Awaji, K. Takahashi, K. Koyama,
N. Kobayashi, M. Ishizuka, T. Itou, T. Tsurudome and J.
Sakuraba, IEEE Trans. Appl. Supercond. 16, 934 (2006).
[3] S. Awaji, K. Watanabe and M. Motokawa, J. Crys. Growth 226,

Fig. 2 DTA curves in magnetic fields up to 18 T for Bi–
53at.%Mn.

Fig. 3 The straight line is a fitting of Tt by the minimum
mean square method in magnetic fields up to 26 T. The
dotted line is a fitting of Tt in a magnetic field range
below 18 T.

83 (2001).
[4] M. Sagawa, S. Fujimura, H. Yamamoto, Y. Matsuura and K.
Hiraga, IEEE Trans. Magn. MAG-20, 1584 (1984).
[5] X. Guo, X. Chen, Z. Altounian and Strom-Olsen, J. Appl. Phys.
73, 6275 (1993).
[6] K. Koyama, Y. Mitsui, E.S. Choi, Y. Ikehara, E.C. Palm and K.
Watanabe, J. Alloys and Compounds 509, L78 (2011).

Kazuo Watanabe (High Field Laboratory for Superconducting Materials)
E-mail: kwata@imr.tohoku.ac.jp
URL: http://www.hflsm.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201264

Research Centers

Challenge for Metallic Material Innovation by Collaboration
between IMR, Industry, and Government
Kansai Center for Industrial Materials ResearchAcollaboration project (Kansai Center) between IMR, Osaka Prefecture Government, and enterprises in the
Kansai area, sponsored by the Ministry of Education, Culture, Sports, Science, and Technology, was establishedon April 2011. The Kansai Center strives to create a new value in manufacturing in the metallic material industry
and to contribute to human resource development through collaborations with industry and government.

The Kansai Center was established in Osaka
as a special unit of the Institute for Materials
Research, Tohoku University in April 2011, based on
an agreement with Osaka Prefecture Government.
The center is sponsored by the government (Ministryof Education, Culture, Sports, Science, andTechnology), and it took over the project of the
Osaka Center, which was carried out from 2006 to2011 aiming to support small and medium enterprisesin the Osaka area. There are three primary scopesof the Kansai Center. First, we respond to technicalproblems in enterprises related to metallic materials
and their process technologies, and we work with
the industry to solve them. Second, the Center
introduces academic research output to industry
and strives to apply collaborative research withcompanies. Third, the Center educates nextgeneration materials scientists and researchers inuniversities and enterprise about material science
applications. The project is conducted under various
collaborations with other organizations in order toachieve efficient activities, especially with theDevelopment Center for New Metallic Material at
Osaka Prefecture University and the ResearchCenter for Nano-Micro Structure Science andEngineering at the University of Hyogo. There areseven groups in the Center: Environmental Protection
and Energy Conversion Materials, Functional
Materials for Future Industries, Analytical Scienceand Technology, Structural Materials for a LowCarbon Society, Nanostructure Control forEngineering Materials, Advanced Biomaterials, andStructure Design of Green Metallic Materials.
The Center has some unique features andactivities. First, the Kansai Center has six full-time
university academic staff members (four in Osakaand two in Hyogo) to achieve the work, which isdifferent from other similar organizations at otheruniversities. This enables us to concentrate on
collaborative activities with industry and government.
Second, the Center has various collaborations withother organizations and supporting industries, suchas the Osaka Bay-area Consortium, ResearchSociety for Global Materials Crisis, and General65

IMR KINKEN Research Highlights 2012

Production Companies, which enables efficient
support of industries. Third, seven professors in theCenter respond to technical inquiries from industries
themselves, and this saves investigation andproblem-solving time. Fig. 1 shows the number ofinquiries from industries at Creation Core Higashi
Osaka (a governmental body where sixteen
universities and one college are located to respondto inquiries from industries). It is clear that the Kansai
Center occupies the largest portion among theuniversities, and the number of inquiries to the
Kansai Center increase each year.
The Kansai Center has opened technicalseminars named “Monodukuri Kisokoza” on special
topics to provide scientific knowledge to technicians
and researchers in industry. The seminars areorganized by the professors in the Center, and the
topics were selected based on industrial needs.
Since fall of 2011, the new seminar has started toprovide fundamentals of individual metals, such as
plasticity, corrosion, and characteristic functions, inaddition to history and production. Furthermore, theCenter professors provide lectures on materials
design, technical skills, and related fundamentals of

Fig. 1 Fraction of inquiries from industries at Creation Core
Higashi Osaka (where sixteen universities and one
college are located); the red region corresponds to the
Osaka Center.

IMR KINKEN Research Highlights 2012

materials science to the students in Osaka Prefecture
University and Hyogo Prefecture University todemonstrate the importance of material science
contributions to industrial applications.
The Kansai Center developed structural andfunctional materials under collaboration withindustries in the Kansai area. Last July, the first
application
using
academic
outputs
wascommercialized by an enterprise in Osaka: Ni-based
two-phase intermetallic alloys composed of Ni3Al
and Ni3V [1] were applied to the tool materials offriction stir welding (FSW) equipment (Fig. 2). This
alloy exhibits extraordinarily high hardness above
700 °C compared to commercial WC hard materials,
which is beneficial for heat-resistant materials.

Fig. 3 A sink roll with a Ni–Mo–B amorphous spray coating
film (lower) and its hardness (upper).

Fig. 2 FSW (friction stir welding) equipment (left) and a
non-consumable tool composed of Ni-based twophase intermetallic alloys (right).

Fig. 3 shows a sink roll with a Ni–Mo–B
amorphous spray coating film that exhibits anticorrosion behavior in hydrochloric acid [2]. The
corrosion resistance of the developed Ni–Mo–B
amorphous coating film is three times higher thanthat of Hastelloy C276 in concentrated hydrochloric
acid, and its hardness is about the same as Cermet(WC-12Co). The Ni-based amorphous coating is
expected to be applied to machine parts underheavy corrosive atmospheres.
Photocatalytic TiO2 prepared by anodic
oxidation [3] for water-purification is underdevelopment with an enterprise in Osaka. Fig. 4
shows a prototype device (left) and an incorporated
Ti lath coated with anodic oxide (right), which could
purify the contaminated water.
The Kansai Center strives to innovate in themetallic material industry through collaborations
between industry, academia, and government andto promote material science.

Fig.4 Prototype device (left) and incorporated Ti lathcoated with anodic oxide (right).

References[1] T. Moronaga, S. Ishii, Y. Kaneno, H. Tsuda, T. Takasugi, Mater.
Sci. Eng.A 539, 30(2012)[2] M. Komaki, T. Mimura, R. Kurahasi, H. Odahara, K. Amiya, Y.
Saotome, T. Yamasaki, Mater. Trans. 53, 681(2012).
[3] Y. Mizukoshi and N. Masahashi, Chem. Lett. 41 545 (2012).

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

IMR KINKEN Research Highlights 201266

Research Centers

Invention of New Materials by Supercomputer
Center for Computational Materials Science

Simulation is becoming more and more important as a third tool, along with experiment and theory. The

research fields of materials design covered by computer simulations are expanding rapidly. The newly
implemented supercomputer system HITACHI SUPER TECHNICAL SERVER SR16000 Model M1 has a 300
TFLOPS processing speed, 42.2 TB of memory, and 1.4 PB of disk storage capacity. We are planning toestablish the basis to serve this large-scale supercomputer resource as one of the centers within a nationwide
project headed by the K computer that form a network of researchers in materials design in Japan as well as
worldwide.

The aim of computational materials science isto design and expedite the synthesis of new materials
with novel or extreme properties. One of theadvantages of computational materials science isthat it can guide this exploration by providing
efficient strategies. Mutual collaboration between
experiments, theories, and computer simulations isimportant for materials science. Providing stimulating
and useful computational environments is the keyrole of the Center for Computational Materials
Science (CCMS). It includes high-performancesupercomputers, the latest versions of software, anda high-speed network.
CCMS was developed from the Laboratory ofMaterials Information Science that was establishedin February 1989. Its main tasks are as follows:
1. Administration and maintenance of thesupercomputing system in this Institute andmaintenance of the supercomputing system
network.
2. Support the usage of the supercomputing
system.
3. Support materials design by supercomputing
simulation with parallelization.
4. Support HPCI (High-Performance Computing
Infrastructure) and CMRI (Computational Materials
Research Initiative). A brief introduction to CMRI
is given in the last paragraph.

Supercomputing system for computational materials science
Supercomputer

10Gbps / 1Gbps

Storage area network( Fibre channel )File server

8Gbps FC×16

10Gbps / 1Gbps

Application server

Storage system

8Gbps FC×8

8Gbps FC×12

Network system
10Gbps / 1Gbps

8Gbps FC×4
10Gbps / 1Gbps
Server area network( Ethernet )1Gbps

10Gbps / 1Gbps

StarTAINS

Fig. 1 System configuration.

67

IMR KINKEN Research Highlights 2012

IMR KINKEN Research Highlights 2012

The supercomputing system at CCMS aims toserve as a solution for researchers in materials
design to realize a sustainable world by computer
simulation. Fig. 1 shows the system configuration of
the supercomputing system at CCMS, which was
newly installed in April 2012. The details of eachsystem are as follows:
1. The supercomputer is the HITACHI SUPERTECHNICAL SERVER SR16000 Model M1, which
is a shared-memory type parallel machine (watercooled system). It has a total processing power of
300 TFLOPS composed of 320 nodes eachhaving 980.48 GFLOPS. This is a 40 times speed
increase from the previous machine. It has 42.25
TB of main memory, which serves enoughcomputational space for the large-scalecomputing required by the institute.
2. The application server is an IBM Blade Center
HS22, which is a distributed memory parallelmachine. It consists of 50 computing nodes and
4 GPFS servers. It is mainly used for software
applications.
3. The file server is a HITACHI ENTERPRISE SERVER
EP8000, which manages each user's home
directory.
4. The disk storage system is a Hitachi Adaptable
Modular Storage2500, which consists of a 1.4 PB
storage area, sufficient for intermediate storage oflarge scratch files.
Fig. 2 shows an example of the results of largescale simulations for the theoretical design of new
conductive oxides. Transparent conducting oxides(TCOs) are a unique class of materials that are
optically transparent and electrically conductingwith a wide variety of optoelectronic applications,such as transparent electrodes in flat panel
displays, photovoltaic cells, electrochromic
devices, and transparent semiconducting devices.
One common TCO is tin-doped indium oxide (ITO),
owing to its good conductivity and high opticaltransparency. At present, joint experimental andtheoretical ab initio investigations are being
performed to predict a new TCO with the prescribedcharacteristics of interest. [1]Finally, we introduce CMRI. CMRI wasestablished in 2011 in IMR as a project-oriented
section and is one of the three operating institutionsof the Computational Materials Science Initiative

Fig. 2 Unit cell of indium oxide.

(CMSI), which was founded through a Grant for Field
2, “New Materials and Energy Creation,” of the HPCI
Scientific Program of the Ministry of Education,
Culture, Sports, Science, and Technology.
The goal of CMRI is to use supercomputers,
among which the K computer boasts the world’s top
performance, and the supercomputer system toinnovate state-of-the-art computational materials
research. Current research topics include firstprinciples calculations of mechanical properties,
interface mechanical/functional relationships, andevolution of microstructure, etc. of various metallic
alloy systems. CMRI researchers have beendeveloping their own program codes for theseprojects, based on electronic structure and statisticalmechanics theories.
In recognizing the importance of forming thebasic infrastructure that will lead to the next
generation of computational materials science,
CMRI offers seminars and symposia on parallel
computing and collaborative meetings with company
researchers. CMRI also promotes training andeducation in computational materials science forstudents and researchers in universities andcompanies. [2]References[1] M. N. Tripathi, K. Shida, R. Sahara, H. Mizuseki, and Y.
Kawazoe, J. Appl. Phys. 111, 103110 (2012).
[2] http://cmri-tohoku.jp/

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

IMR KINKEN Research Highlights 201268

Research Centers

Growth and Characterization of Alloys of GaN, InN, and AIN
for Optoelectronics by ICC-IMR Visiting Professorship
International Collaboration Center (ICC-IMR)I

CC-IMR is a gateway to IMR from the worldwide community of material science. Those who want to perform
collaboration with the groups in the institute can apply for a variety of programs. Here is an example of theachievement on the growth of nitride semiconductors conducted as a guest professorship.
ICC-IMR was founded in April 2008 to succeed
the International Frontier Center for Advanced
Materials (IFCAM). The mission of the center has
extended from an international think-tank of materials
science to the center for international researchcollaborations of the Institute for Materials Research.
To achieve this purpose, the program of the centerhas expanded to include the research collaboration
projects by international teams.
As one of the centers of excellence in materialscience, IMR holds 27 research groups and six
research centers. ICC-IMR is working as the gateway
for collaborations between international researchers
and IMR. At present, the following seven different
programs are carried out in ICC-IMR. 1: international
project research, 2: visiting professorship, 3: short
single research visits, 4: international workshops, 5:fellowship for young researchers, and 6: coordinaton
of international research collaboration. 7: Material
Transfer Program. We welcome applications for
those international programs from all over the world.
Prof. Baskar, Crystal Growth Centre, Anna
University worked at ICC-IMR as a guest professor
in FY2011 at the Prof. Matsuoka laboratory. Here is areport of his activities, impressions about IMR, andmessage for future applicants.
Upon arrival to IMR, Prof. Baskar arranged tohave detailed discussions about the facilities
available for the growth of InGaN materials on
sapphire substrates. The works here have been
made very effective by using the two MOVPE
systems of the laboratory. He has grown InN and
InGaN materials. The epitaxial layers were analyzed
by using high-resolution X-ray diffraction (HXRD),
photoluminescence (PL), infrared reflectancespectroscopy, scanning electron microscopy (SEM),
atomic force microscopy (AFM), and a Hallmeasurement system. As such, the instruments at
IMR are well prepared. The special growth
conditions, such as temperature, high V/III ratio, andhigh reactor pressure, required for the growth ofgood quality indium-based nitrides have been easily
achieved.
Furthermore, in-situ monitoring systems, suchas the Laytech tool, required for understanding of
the growth mechanisms of indium nitride in MOVPE,69

IMR KINKEN Research Highlights 2012

particularly the shape of the initial stage nucleationand the coalescences at higher temperature toobtain two-dimensional growth, have been
discussed in detail. The possible use of InGaN for
high-efficiency solar cells was also discussed,
particularly the issues related to p-type doping, theadvantages for concentrated solar cell applications
due to strong bonding between nitrogen and indium,and small changes in the band gap with temperature.
Those discussions with experts are very useful for
progress studying the crystal growth of nitrides.
AlGaN samples with aluminum contents from
20 to 70% were prepared by MOVPE on a two-inch
sapphire substrate and cut into 1 cm2 and 1 × 2 cm2size samples using a diamond scriber. The samples
were cleaned in acetone and isopropyl alcohol.
Indium contacts were used to measure the carrier
concentration and mobility by a probe method.
Though the samples are un-doped, they exhibited
n-type conduction with a carrier concentration on
the order of 1017 cm-3 up to an aluminum content of
35%. The higher aluminum content samples, above
40%, were highly resistive and the conduction typewas difficult to measure. The AlGaN/GaN samples
with aluminum content of 22% and thickness ofaround 80 nm have shown a two-dimensional
electron gas (2-DEG) with a mobility of 980 cm2 V/
sec. The growth of high-aluminum-content samples
on GaN surfaces without cracks and with good
crystalline quality is challenging due to a 3% latticemismatch between AIN and GaN. The samples
were analyzed by scanning probe microscopy, andsmooth atomic surface steps, pinholes, and microcracks were observed depending on the aluminium
content. A typical microstructure of an AlyGa1-yN/
AlxGa1-xN/AlyGa1-yN double heterostructure (DH) on
GaN/Al 2O3 observed by AFM is shown in Fig. 1.
Though the thicknesses of the GaN buffer layerand AlyGa1-yN barrier are the same, a small thickness
variation of 5 nm resulted in micro-cracks, which
terminate at pinholes. The micro-cracks are formed
due to the relaxation of strain once the critical
thickness of the active layer of AlxGa1-xN exceeds
the limit. Attempts have to be made to grow highaluminum-content (Al > 40%) AlGaN with highcrystalline quality and smooth morphology using an

IMR KINKEN Research Highlights 2012

AlN buffer with a better lattice match. Once theconditions are optimized for the AlGaN layers,
resolution of the doping issues, particularly p-type
doping, will be undertaken to develop high-efficiency
ultraviolet light-emitting diodes. Prof. Bakar has hadmuch success in these crystal growth studies in thevery short term of his visit due to the excellent
research conditions of IMR.
During his stay, he visited Nagoya Institute ofTechnology, Hiroshima University, Kyushu University,and Hokkaido University over 14–20 October 2011.
He found that all universities are eager to establish
and strengthen the bilateral academic and researchcooperation with Anna University, Chennai, India.
As there are a large number of Japanese companies
in Chennai and discussions are occurring between
the government of Japan and Government of Tamil
Nadu in Chennai, India, to establish a Japaneseindustrial park in Chennai, many institutions arelooking for a trilateral cooperation between industry
in Japan or India, academia in Japan, and Anna
University.
Extensive discussions were made to strengthen
the semiconductor research at the Crystal GrowthCentre of Anna University. It has been proposed tohave an India–Japan joint workshop on nitride
semiconductors for electronics, photonics, andphotovoltaic applications in the year 2012 so that
researchers in Japan and India can join the workshop
at Anna University to bring forth joint deliberations
on future strategies to establish strong collaborations
betweenIndianuniversities/institutions
andJapanese universities. The deliberations will also
be useful to strengthen the research on nitride
semiconductors for the development of novel
devices.
Prof. Baskar said that the ceremony for thedeclaration of international material science weekon 11th October 2011 at IMR, Sendai was a wonderful
opportunity to share some of his thoughts about the
recent disaster at Fukushima, the collaborative
research between IMR and Anna University, the
new initiatives in the development of materials at
Anna University, and other topics for a better andsustained future for mankind.
He has expressed thanks to IMR, particularly
Prof. T. Matsuoka, and Anna University, India, fortheir support in establishing active collaboration.

Fig. 1 The surface morphology of AlGaN-DH with various
active-layer thicknesses. From top to bottom: 20, 15,
and 10 nm QW.

Hiroyuki Nojiri (International Collaboration Center)
E-mail: icc-imr@imr.tohoku.ac.jp
URL: www.icc-imr.imr.tohoku.ac.jp
K. Baskar, Crystal Growth Centre, Anna University
E-mail: baskar@annauniv.edu

IMR KINKEN Research Highlights 201270

Research Centers

Activity of the Integrated Materials Research Center for a
Low-Carbon Society
Integrated Materials Research Center for a Low-Carbon Society (LC-IMR)

T

he 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 contribute to making lowcarbon 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, the general 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. 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 seek
collaborations both among our members and withoutside centers/institutes/universities.
In the past two years, we have has been working
on the following three subjects:

Fig. 1 Outlook of LC-IMR.

71

IMR KINKEN Research Highlights 2012

IMR KINKEN Research Highlights 2012

Fig. 2 One scene from the 1st LC-IMR workshop.

1. Conducting strategic research by individuals.
2. Supporting new research projects throughmanagement of internal research funding.
3. Having discussions with industry and government
for a better understanding of the needs of oursociety.
Japanese funding agencies are determiningthat various strategic research programs in materials
science are necessary for our future society. We arebeing funded and making good progress in thoseresearch programs.
LC-IMR is also managing an internal research
program funded by IMR to support both promisingand challenging research projects, particularly
emphasizing on the center's mission. Examples ofthe projects selected so far include the strengtheningof structural steels for automobiles and thedevelopment of new soft magnetic materials andsensor materials.
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 industry or government, todiscuss how to develop suitable systems for asustainable society and the expected roles ofmaterials science. Such collaborations are anotherimportant part of our activities. We will continue tomake efforts to progress in our missions and ask
support from everyone in 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 201272

Research Centers

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

T

he Center of Neutron Science for Advanced Materials (CNSAM) is a unique neutron facility that has abackground of advanced material science at IMR. CNSAM operates two neutron spectrometers for materialscience in a reactor facility and is also the core of a project to construct a new neutron spectrometer in J-PARC/
MLF, which is the world’s brightest neutron source. Moreover, based on the crystal growth technology of IMR,CNSAM is developing novel neutron devices.

Fig. 1 IMR neutron diffractometer, HERMES (upper), andspectrometer AKANE (bottom) installed in the JRR-3
reactor.

73

IMR KINKEN Research Highlights 2012

11

Fe52Pt30 B18 (475°C, 900 s)
CNT / 14.4 M cnt~ 36 min

4000

(a)

B = 0.8T
+
I
I

3000
2000
1000
0

+ I - I /48 min

1000
800
600

800

(b)

Fe2B
L10-FePt

B =0.8T
B ~ 5mT

600
400

400
200

200

0

0
40

+ I - I / 36 min

Material science using IMR spectrometer anddiffractometer in a reactor facility
CNSAM operates a neutron spectrometer andneutron diffractometer for materials science in the
JRR-3 reactor in the Japan Atomic Energy Agency(Ibaraki, Japan). The powder diffractometer,
HERMES, (Fig. 1 upper panel) is an instrument todetermine the structures of spins and nuclei inmaterials with high efficiency. HERMES is one of themost active diffractometers for material science inJapan because HERMES is a powerful tool forinvestigationsofmagnetic
materials
andenvironmentally applicable materials.
Using HEMRES, we have succeeded inobserving the microscopic magnetic characteristics
of a composite magnet, Fe52Pt3011B18, based on ametallic glass, by polarized neutron diffraction (PND)
experiments. PND is sensitive to small magnetic
moments in magnetic materials. Fig. 2(a) shows the
PND patterns of annealed Fe52Pt3011B18; I+ (red) andI- (blue) indicate the PND patterns for parallel andantiparallel neutron spins and sample magnetization,

45

50

55

60

65

Scattering Angle / °Fig. 2 PND patterns of I+ (parallel) and I− (antiparallel) for
Fe52Pt3011B18 composite magnet for B=0.8T (a), and thedifference patterns (b) for B=0.8T and B~5mT [1].

respectively. The small difference in peaks meansthat the peaks include weak ferromagnetic
components. Fig. 2(b) shows the difference, I+ - I-,
which is proportional to the ferromagnetic
components, for different magnetic fields applied tothe sample. Note that the ferromagnetic components
of FePt remain even at 5 mT, while those of Fe2B
disappear, meaning that the difference in soft andhard magnet characters can be directly distinguishedby PND experiments.
The IMR neutron triple-axis spectrometer,
AKANE, (Fig. 1, bottom panel) is an instrument toobserve time and spatial fluctuations of spins andnuclei in materials with high accuracy. AKANE isindispensable for investigating the main issues ofnovel material science: high-TC superconductors,
metallic
antiferromagnets,
geometricalspin
frustration systems, and so on. Note that observationof spin correlations under high pulsed magnetic
fields, which is one of the characteristic technologies
at IMR, has already been realized on AKANE.
Spin dynamics in quantum spin systems
Quantum spin systems have attracted muchinterest due to their non-trivial phenomena. Neutron

IMR KINKEN Research Highlights 2012

study between experiment and theory on quantum
magnets proceeds under an operation of J-PARC.

Fig. 3 Inelastic neutron scattering spectrum of the quasione-dimensional spin system CuGeO3 measured in awide energy-momentum space.

scattering experiments provide direct information,
which connects experimental and theoreticalstudies. In the quasi-one-dimensional system
CuGeO3, frustration between antiferromagnetic
exchange couplings within the spin-chain is reported
to be responsible for the exotic physical properties.
To elucidate the nature of the frustration in the spin
system, we measured the full excitation spectrum of
CuGeO3 in collaboration with a group at J-PARC.
Fig. 3 shows the inelastic neutron scattering intensity,
which is directly related to the spin–spin correlation
function. Clear continuum excitation exhibiting aquantum spin nature was observed. An intriguingfeature was found in the form of a relatively strong
intensity on the upper boundary of the continuum.
This feature is not expected in a system without
frustrated couplings, and the finding triggered arevisiting of the theoretical study of one-dimensional
magnets. The observed spectrum was well
reproduced by a recent model calculation that
included frustrated couplings [2]. Further cooperative

New projects in J-PARC
To make a jump in material science, CNSAM istrying to realize a new project: a second-generation
spectrometer named POLANO at the world’s
brightest neutron beam facility at the Japan Proton
Accelerator Research Complex (J-PARC), based oncollaborations between Tohoku Univ. and KEK(project leader: K. Ohoyama) (Fig. 4). CNSAM is at
the core of the POLANO project. The neutron
spectrometer POLANO is characterized by polarized
neutrons, which are very sensitive to pure magnetic
correlations in magnets. Although we have alreadyrealized PND experiments on HERMES, the neutron
flux for polarized neutron experiments is not sufficientfor novel material science. To overcome this point
and to make breakthroughs in material science
using polarized neutrons, POLANO will be
indispensable. From an applications viewpoint,
POLANO will be important to develop rare-earthfree permanent magnet materials and forinvestigations of the fundamental physics in novel
magnets. Since only a few polarized neutron
spectrometers exist in the world, this spectrometer
will be a key instrument to generate breakthroughs
in novel material science in Japan.
In 2011, this project succeeded in gaining
authorization by the J-PARC Center. With a goal of
beginning construction in 2012FY, detailed
engineering designs have already begun.
References[1] K. Ohoyama et al. Nucl. Instr. Meth. in Physics Res. A, 680, 75(2012).
[2] T. Sugimoto et al., (unpublished, arXiv:1109.2410)

Detector Bank

Spin Analyser

Stair to Upper Floor

Entrance

Beam Stopper

Spin Polariser
Radiation Shield

Beam Transport Tube
Fermi Chopper
Neutron Beam

Sample Position
Radiation ShieldFig.4 A rendering view of the Tohoku Univ. spectrometer POLANO in J-PARC.

Kenji Ohoyama (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 201274

Research Centers

Open Laboratory at Cryogenic Center for Wide Research Fields
from Frontier Material Science to Low-Temperature Physics
Laboratory of Low Temperature Materials Science

The Laboratory of Low Temperature Materials Science supplies the platforms to perform basic research on

materials though the measurement of transport and thermodynamic properties over a wide range of temperatures,from 350 K to 0.01 K. Using a variety of cryogenic equipment, a number of collaborative research projects with
scientists from various research fields are ongoing under the technical and educational support of the staff inthis laboratory.
Recently, cryogens, such as liquid helium, havebeen widely used by scientists not only in the fields
of low-temperature condensed matter physics andapplied physics but also in other fields that are notdirectly related to low-temperature matter. Most
generic equipment for characterizing materials with
automatic PC control (for example, SQUID
magnetometers, NMR systems, etc.) is now available
as commercial products, and one can easily operate
it from room temperature down to 1 K without special
low-temperature techniques. However,theysometimes require the support of experts on lowtemperature science for maintenance, containing
initial cooling, keeping the equipment cooled, andtroubleshooting.
On the other hand, the special apparatuses for
producing very low temperature environments below
1 K have also attracted many researchers working
on low-temperature physics and frontier materialscience, since in many cases, novel quantum
phenomena in new devices, materials, andcompounds appear in the extreme absence ofthermal effects. Experiments below 1 K usuallyrequire relatively large-scale facilities, particulartechniques, expensive cryogens such as 3He, and alarge amount of liquid helium (4He) for long-termcooling. Therefore, it is effective and convenient tokeep such apparatuses together with the genericequipment at one division, which works as an open
cryogenic center and consists of low-temperatureexperts.
On the basis of the concepts mentioned above,
the Laboratory of Low Temperature Materials
Science (LLTMS) possesses many kinds ofcryogenic equipment (as listed below) that meet thedemands of most researchers and enable them tomeasure the basic properties of materials as afunction of temperature T and magnetic field H.
(1) S QUID magnetometers (× 2 sets):
for magnetization measurements at T = 1.8–350
K in H up to 7 (5.5) T.
75

IMR KINKEN Research Highlights 2012(2) V SM magnetometer (under reconstruction):
for magnetization measurements at T = 4–300 K
in H up to 14 T.
(3) Superconducting magnet (SM) and variabletemperature insert (VTI) system (× 2 sets):
mainly for transport measurements at T = 1.2–
320 K in H up to 11 (10) T.
(4) S QUID microscope system:
for observation of magnetic field distributions on
sample surfaces at 2–100 K(5) Top Loading 3He cryostat and SM:for transport measurements at T = 0.26–300 K in
H up to 8 T.
(6) 3He cryostat and SM with field gradient coils:
for magnetization measurements by Faraday
methods at T = 0.3–2 K in H up to 8 T.
(7) D ilution refrigerator and SM system:
for transport and ultrasonic measurements at T =0.02–1.4 K in H up to 10 T.
(8) S uperconducting vector magnet system withdilution refrigerator or VTI:
for transport measurements as a function ofmagnetic field direction in H up to 7 T
Taking advantage of the plentiful liquid helium
made in the Center for Low-Temperature Science(cryogenic center), much of the above equipment isalways kept cooled and open to all scientists of
Tohoku University. In addition to in-house research,a number of cooperative projects are now in progressunder the careful technical and educational supportof the staff of LLTMS. LLTMS contributes to the
publication of more than 20 papers per year. The
recent topics that the staff of LLTMS have deeply
contributed to are:Discovery of superconductivity in KTaO3 byelectrostatic carrier doping [1]By employing an electric double layer transistor
configuration for electrostatic carrier doping,superconductivity is successfully induced at thesurface of KTaO3, in which superconductivity has

IMR KINKEN Research Highlights 2012

Preparation of n-type metallic YBa2Cu3Oy using
an electrochemical technique [2]Realizing an ambipola cuprate superconductor
is an important issue. Using an electrochemical
technique, we succeeded in converting the sign ofconduction carriers in YBa2Cu3Oy, which is known
to be a typical p-type cuprate superconductor with
Tc = 90 K, from holes to electrons (Fig. 2), to theextent that the new n-type YBa2Cu3Oy shows ametallic transport property.
Observation of field-induced anomalous magnetic
states in a spin-gap material YbAl3C3 [3]YbAl3C3 is known to be a rare 4f-electroncompound that shows a spin-gap state, originatingfrom dimerization of Yb3+ pairs. In the detailed
measurement of magnetization curves below 1 K,
the authors found that the 4f-electron dimer state in
YbAl3C3 transfers to an unusual field-induced state(Fig. 3), which reveals different properties from thatreported in d-electron dimer systems, through acrossover transition.

istance ( Ω )Sheet resis

30

104

n-type103

YBa2Cu3Oy

2

10

101
100

-0.02

Gate

Tc =
45 mK

0

0

0.02 0.04 0.06

Carrier number / Cu
Fig. 2 Evolution of electrical resistivity with hole reductionand electron doping by an electrochemical reaction in
YBa2Cu3Oy films. The horizontal axis is the carrier
number per CuO2 plane.

YbAl3C3
H//a
dimer
phase

KTaO3 EDLTVG = 5V
I = 50 nA

10p-typefieldinduced
disorderphase
polarized
phase

20

0

Resistivity (300 K)

never before been observed (Fig.1).

Fig. 3 Field dependence of magnetization of YbAl3C3 inunits of μ B per Yb atom. Two kinds of field-induced
transition from the spin dimer state at low fields areobserved around H = 4.5 T and 6.6 T at T= 0.47 K.

Ionic liquid

References

Electric
double layer

Semiconductor[1] K. Ueno, S. Nakamura, H. Shimotani, H. T. Yuan, N. Kimura, T.

0.1
T(K)

0.2

[2] T. Nojima, H. Tada, S. Nakamura, N. Kobayashi, H. Shimotani

Nojima, H. Aoki, Y. Iwasa and M. Kawasaki, Nature Nanotech.
6, 408 (2011).

Fig. 1 Superconducting transition observed in thetemperature dependence of resistance in KTaO3 with
an electric double layer transistor (EDLT) configuration
at a gate voltage VG = 5 V. The inset shows the conceptof EDLT.

and Y. Iwasa, Phys. Rev. B 84, 020502(R) (2011).
[3] K. Hara, S. Matsuda, E. Matsuoka, K. Tanigaki, A. Ochiai, S.
Nakamura, T. Nojima and K. Katoh, Phys. Rev. B 85, 144416(2012).

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

IMR KINKEN Research Highlights 201276

Research Centers

Research Facility for Studying Physical and Chemical
Properties of Radioactive and Nuclear Materials
Laboratory of alpha-Ray Emitters

More than 170 species of radioisotopes and nuclear materials can be used in the Laboratory of Alpha-Ray

Emitters. This facility is one of the important centers in Japan for studying the physical and chemical propertiesof radioactive materials, especially for alpha-ray emitters such as actinide elements, as well as fission andfusion reactor materials. The facility has been used by many researchers from all over the country for preparationof samples and for chemical and physical measurements of a variety of radioactive materials.

In March 1960, a radioisotope (RI) handlingfacility was officially approved by the Science andTechnology Agency (currently, Ministry of Education,
Culture, Sports, Science, and Technology) as‘SHIDAI 337’ and installed in the basement of Build.
#1, IMR. The facility was relocated to the present
building in 1978. The Laboratory of Alpha-Ray
Emitters is currently approved for the use of morethan 170 species of radioisotopes, including alpharay emitters such as actinide elements and nuclear
materials (Fig. 1 shows the uranium metal stored in
the facility). The number of such facilities is quite
limited in Japan; the others are JAEA (Japan Atomic
Energy Agency) and the International ResearchCenter for Nuclear Materials Science (Oarai) of IMR.
The Laboratory of Alpha-Ray Emitters has, therefore,
an important role in the study of actinide elements inJapan, and this activity and capacity extends
international research collaborations.
The facility is composed of two areas: theradiation-controlled area and the non-controlled
area. The controlled area consists of three rooms forchemical experiments, three rooms for physicalexperiments, one room for radiation measurements,
one contamination test room, and rooms for storage
of RI and nuclear materials. Each chemical room
has local exhaust ventilation systems for thetreatment of various chemicals.
The radiation-detection equipment that can beused in this laboratory are two gamma-ray
spectrometers (high-purity Ge semiconductor
detector), a liquid scintillation counter, and an alpharay spectrometer (silicon detector). A mono-arc
furnace, tetra-arc furnace, electrical dischargemachine, X-ray diffractometer, and IP reader areused for the preparation and identification of nuclear
material alloys, such as UCoGe. Two inert-gas glove
boxes, a CHNS element analyzer, an ICP emission
spectrophotometer, an FT-IR, a UV–visible
spectrophotometer, a powder X-ray diffractometer,
and an X-ray single crystal structure analysis system
are used for the treatment of nuclear materials andother chemicals. A scanning electron microscope,
automatic titration system, and electrochemical
77

IMR KINKEN Research Highlights 2012

measurement system are installed in the laboratory.
A transmission electron microscope has been used
to study the irradiation effects of nuclear materials.
In recent years, a physical property measurement
system (PPMS) was installed and used for thephysical study of actinide compounds (Fig. 2).
In resent years, this laboratory has produced
several outstanding achievements. The noted work
introduced here includes (i) a study on the physicalproperties of uranium metallic compounds basedon preparation of very fine single-crystals, and (ii)Fig. 1 Uranium metal rods.

Fig. 2 Physical property measurement system (PPMS)
installed in the radiation-controlled area. PPMS can beused to measure heat capacity as well as AC and DC
susceptibilities.

IMR KINKEN Research Highlights 2012

the synthetic study of molecular-based magnets
including f inner-transition elements, i.e., actinides
and lanthanides.
The interplay of ferromagnetism andsuperconductivity has been a central topic ofcondensed matter physics. Recently, we presented
the first example in which superconductivity is
induced by ferromagnetic spin fluctuations in the
ferromagnet UCoGe [1-2], and we further suggested
the possibility that a spontaneous vortex state
emerges as a compromise of the two competitive
orders; a vortex induced by spontaneous
magnetization can exist even at zero external
magnetic field (right panel of Fig. 3), in contrast to aregular type-II superconductor (left panel of Fig. 3).
Magnetic field

Magnetic field

Hc2

Hc2

Mixed state
Hc1

Spontaneous
vortex stateFig. 4 Molecular structure of the Er(III)Zn(II)3 wheelshaped tetranuclear complex generated on the basis of
X-ray data. Green and yellow spheres represent Er(III)and Zn(II) ions, respectively. Er(III) ion is located in anine-fold coordination, of which six equatorial oxygen
donor atoms are represented as red spheres. This typeof anisotropic coordination geometry is essential for
Ising-type magnetic anisotropy for Er(III) ion.

Meissner state
Temperature

Lanthanide and actinide complexes were widely
investigated as candidates for quantum nanomagnets, such as a single-molecule magnet (SMM)or a single-chain magnet (SCM), in which eachmolecule shows superparamagnetic behavior below
Tc. We have synthesized the first example of astructurally designed Er(III) SMM (Fig. 4 and Fig. 5)[3].
The fission and fusion reactor materials
intensively investigated are (1) Zr–Nb binary alloy
development for advanced nuclear fuel rods, (2)irradiation behavior of ferritic–martensitic steels forfusion first walls, (3) vanadium-based alloy
development for fusion (collaborating with NIFS),
and (4) microstructural evolution and mechanicalproperty change in steels irradiated with highenergy ions, performed at the Cyclotron Radioisotope
Center (collaborating with School of Engineering).

0.8

10000 Hz
5600
3200
1800
1000
560
320
180
100

-1

Fig. 3 Schematic illustration of superconducting phase
diagrams.

M'' / emu mol

Temperature

1.0

0.6
0.4
0.2
0.0

2.0

3.0

4.0

5.0

T/K
Fig. 5 Temperature dependence of the out-of-phasecomponent of the AC magnetic susceptibility (χM’’)
measured for various frequencies. Below 4 K, slow
relaxation of the magnetization occurs and frequencyand temperature-dependent χM’’ signals were observed,
which is typical for single-molecule magnets.
Rev. Lett. 108, 066403 (2012).
[2] K. Deguchi, E. Osaki, S. Ban, N. Tamura, Y. Simura, T.
Sakakibara, I. Satoh, and N. K. Sato, J. Phys. Soc. Jpn. 79,
083708 (2010).

References[1] T. Hattori, Y. Ihara, Y. Nakai, K. Ishida, Y. Tada, S. Fujimoto, N.
Kawakami, E. Osaki, K. Deguchi, N. K. Sato, and I. Satoh, Phys.

[3] A. Yamashita, A. Watanabe, S. Akine, T. Nabeshima, M.
Nakano, T. Yamamura, T. Kajiwara, Angew. Chem. Int. Ed., 50,
4016 (2011).

Kenji Shirasaki (Laboratory of alpha-Ray Emitters)
E-mail: kshira@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201278

Research Centers

Determination of Released Trace Elements in Static
Immersion Test for Characterization of Metallic Biomaterials
Analytical Research Core for Advanced MaterialsO

ur research group conducts elemental analysis of various materials to determine major, minor, and tracelevels of metallic and non-metallic elements. This report introduces the analysis of metallic biomaterials, one of
our recent activities. The amounts of metallic ions released from the materials using static immersion into
simulated body fluids (SBF) were determined. Sensitive and accurate determination of trace metallic elements
in SBFs could be performed by inductively coupled plasma optical emission spectrometry (ICP-OES). By usinga H2SO4-fume pre-treatment method, it was possible to determine concentrations of elements in SBFs on a
ppb order.
Metallic biomaterials have been widely
employed by orthopedists or dentists in clinicalpractice. Each of the metallic biomaterials has been
applied to appropriate parts of the human body.
However, when the metallic biomaterials are
implanted for a long time, metallic elements may bedissolved into the surrounding tissues. The resulting
toxic influence of metals released from metallic
implants has been demonstrated [1, 2]. It has been
recently reported that not only allergies but also
several medical disorders have been caused bythese metallic elements. Allergy responses areespecially sensitive to trace amounts of the metallic
elements. Accordingly, metallic biomaterials thathave high corrosion resistance (high safety for thehuman body and low metal release) are required;
therefore, the evaluation of corrosion resistance isimportant for their development.
The corrosion resistance of metallic biomaterials
is evaluated by using static immersion into asimulated body fluid (SBF) and then assessing themetallic ions in the SBFs. Therefore, trace metallic
elements in the SBFs should be noted wheninvestigating the properties of the metallic
biomaterials. Moreover, in order to evaluate the traceelements released by the static immersion test, it isnecessary to determine μg L-1 (ppb) levels of trace
metallic elements in the SBFs. The concentrationsof various elements in SBF solution samples havebeen determined by inductively coupled plasma
optical emission spectrometry (ICP-OES, Fig. 1),
inductively coupled plasma mass spectrometry(ICP-MS, Fig. 2) and atomic absorption spectrometry(AAS, Fig. 3) [3, 4].
ICP-OES has analytical advantages, such as
the simultaneous determination of elements with awide dynamic range. However, the operating
conditions of the ICP are influenced by the matrix
elements included in a sample solution [5, 6] or theviscosity of the solution. SBFs include a large
amount of alkali metal salts (NaCl, KCl, etc.) and/ororganic compounds (glucose, amino acids, etc.)
79

IMR KINKEN Research Highlights 2012

Fig. 1 Echelle-type inductively coupled plasma opticalemission spectrometer.

Fig.2 Double focused inductively coupled plasma mass
spectrometer.

and have a high viscosity. This could alter the signal
intensities of the analytes, and thus it is difficult toobtain precise and accurate results in the analysis
of elements in the SBFs. Accordingly, it is vital toselect an optimal sample pre-treatment method
before analysis of the elements using ICP-OES.
In order to develop a simple and easy procedurefor high-sensitivity, precise, and accurate

IMR KINKEN Research Highlights 2012

determinations of trace metallic elements in the
SBFs by ICP-OES, we investigated a H2SO4-fume
pre-treatment method [7 - 9]. By using this method,the viscosity of a sample solution can be decreased
and the solution can be stabilized in acidic condition.
Moreover, the organic compounds contained in the
SBFs can be removed. Thus, it is possible todetermine the concentrations of elements in the
SBFs to a ppb level
SBFs include high concentrations of salts, suchas NaCl, KCl, and several alkali or alkaline earth
metals. The authors investigated the influences ofthese matrix elements on measurements by ICPOES. Solutions containing analytes (2 mg L-1) and
NaCl (8 g L-1) were prepared, and the emission
intensities of the analyte elements were measured
by ICP-OES. Their intensities and intensity ratios(elements/Y) are shown in Fig. 4. Both the intensities
and the intensity ratios are normalized to unity whencontaining no NaCl. When containing NaCl, theintensities for Al and Pd were more than 1, but theintensities for the other elements investigated wereless than 1. The intensity ratios for Au, Ca, Cr, Ti, V,
and Zn were nearly 1. In this case, an internal
standard correction method by using Y can beapplied for these elements; however, it is not effective
for the other elements such Al and Pd. Na has aserious influence on the measurement of suchelements by ICP-OES, and therefore, it is necessary
to prepare matrix-matched standard solutions forcalibration to ensure accurate determination.
We determined the elements released fromSUS 316 or Ti–6Al–4V alloys by static immersion
tests. These materials were respectively immersed
into lactic acid in a PFA bottle at 310 K for 20 daysin ambient atmosphere. After removing the materials,
the elements released into the lactic acid were
determined by ICP-OES with the H2SO4-fumetreatment. It was possible to determine concentrationsof the released elements at a ppb level by using theH2SO4-fume treatment method.

Fig.3 Continuous-light-source atomic absorption spectrometer

Fig. 4 Intensity and intensity ratio of analyte elements by
ICP-OES in 8 g dm-3 NaCl solution. Both the intensities
and intensity ratios are normalized to unity whencontaining no NaCl.

[6] M. Ishikuro and K. Wagatsuma, Bunseki Kagaku 56 (2007)933[7] N. Ohtsu, T. Ashino, K. Takada, H. Kimura, T. Hanawa, J.

References[1] F. F. Hennig, H. J. Raithel, K. H. Scaller, J. R. Dohler, i,Electrolytes. Health. Dis. 6, 239 (2011).
[2] K. L. Waper, Clin. Orthop. Relat. Res. 271 12 (1991)[3] Y. Okazaki and E. Gotoh, Biomaterials 26 11 (2005)

Mater. Sci. Mater. Med. 18 429 (2007)[8] T. Ashino, N. Ohtsu, K. Wagatsuma, Mater. Trans. 49 (2008)
1342[9] T. Ashino, N. Ohtsu, K. Wagatsuma, J. Japan Inst. Materials 73(2009) 728[4] Japan Industrial Standard (JIS) T 0304 (2002)[5] T. Ashino and K. Takada, Anal. Sci. 9 737 (1993)

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 2012

80

2012

Index

Author Index
Name

page

A
Abe, HiroakiAndo, KazuyaAn, LiqiongAoki, Dai
Ashino, Tetsuya

28
18
10
52
80

B
Bauer, E.-W. Gerrit

40

C
Chiba, Akihiko

12F
Fujita, Masaki
Furuhara, Tadashi

56
72

G
Goto, Takashi

10

H
Hiraka, Haruhiro

20

I
Ito, Akihiko

10K
Kato, HidemiKawamata, Toru
Kodama, YumikoKoizumi, Haruhiko
Koizumi, Yuichiro
Konno, J. Toyohiko

30
6
36
32
12
36

L
Li, Yunping

12M
Makino, Akihiro
Masahashi, Naoya
Matsumoto, HiroakiMatsunaga, TetsuyaMatsuo, MotoakiMatsuoka, TakashiMatsuta, Hideyuki
Matsuura, MasatoMiyamoto, Goro
Mizuseki Hiroshi

62
66
12
28
34
48
54
20
4
22

N
Nagai, Yasuyoshi

24

Name

page

Niinomi, Mitsuo
Nojima, Tsutomu
Nojiri, Hiroyuki

876
42, 70O
Ohoyama, Kenji
Orimo, Shin-ichi

20, 7434

S
Sahara, Ryoji
Saitoh, Eiji
Sasaki, TakahikoSato, Kazuhisa
Sato, ToyotoSatoh, Yuhki
Shikama, Tatsuo
Shirasaki, Kenji
Sugiyama, Kazumasa

68
18
44
36
20
28
26, 60
78
6

T
Takanashi, Koki
Tokumoto, Yuki

50
46

U
Uda, Satoshi
Usami, Noritaka

32
16

W
Watanabe, Kazuo

64X
Xiao, Jiang
Xie, Guoqiang

40
14

Y
Yamada, Kazuyoshi
Yamaura, Shin-ichi
Yoshikawa, Akira

20
38
58

IMR KINKEN Research Highlights 201282

Keyword Index
KeywordA
actinide
alloy
amorphous
atom probe tomography

52
8
6, 38
24

B
biomedical

8, 14

C
ceramics
clathrate hydrate
Co-Cr-Mo alloy
crystal growth

83

page

10, 2622
1232, 58D
defects
diffraction

44
6E
electrical spin injection
electron microscopy
electron spin resonanceelectronic properties
electronic structureenergy gapenergy storage

40
28
42
44
54
48
34

F
ferromagnetic
film
foam
frustration
fuel cell

52
46
30
42
38

G
graphitization

36

H
hydridehydrogen implantationhydrogen storage

34
28
22

I
insulator
ionic conductor
ion implantation

40
34
16

L
luminescence

58

IMR KINKEN Research Highlights 2012

KeywordM
magnetic
magnetic properties
mechanical properties
metallic glass
micro-pulling-down method
microstructureNnanoclusterneutron irradiation
neutron scattering
neutron scattering measurementnitridenovel magnetic materials
nuclear materials
nucleation and growth

page
40
50
12
14
58
8, 10, 3042
26, 2820, 5656
46, 4820
24, 2632

O
optical absorption
optical properties
organic

54
48
44

P
phase diagram
phase equilibria
phase transformation
platingporosity
powder metallurgy

22
32
4, 12
38
30
14

S
SEM
semiconducting
sintering
solar cells
spin correlation
spin current
spin lattice dynamics
spintronics
steel
structure
superconducting

36
16
10
16
56
18
20
18, 50
4, 24
6
52

T
TEM
thin film

36, 4650

Editors:
Hiroaki Abe (IMR), Takashi Matsuoka (IMR),
Gerrit Ernst-Wilhelm Bauer (IMR),
Yuichiro Koizumi (IMR), Shigeo Sato (IMR), andSatoshi Semboshi (IMR),

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

Printing:
HOKUTO Corporation

Organization
Institute for Materials Research
Tohoku University
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URL: http://www.imr.tohoku.ac.jp/
E-mail: pro-adm@imr.tohoku.ac.jp© IMR September 2012