3FTFBSDI

Highlights of Collaborated Research
Collaborative works with outside researches are carriedout in consideration of the joint research program. Accepted
applications are performed in following each division andfacility of IMR.
1) Research Laboratories
Joint research conducted by out side researchers and
IMR staff members at each research laboratory. Three
categories; “Research in priority areas”, "General Research"and “Exploratory Research for Young Researchers” areprepared.

2) International Research Center for Nuclear Materials
ScienceThis facility is open to university scientists all over Japanto support experiments using Japan Material Test Reactor,
JMTR, the experimental breeder reactor, JOYO and the test
reactor, JRR-3 operated by JAEA (Japan Atomic Energy
Agency). The overseas reactor, BR2, located at the Belgian
Nuclear Research Center, is also used for irradiation
experiments. This facility acts as a hub for international
collaborations; specimens irradiated in overseas reactors are
accepted here for post-irradiation examination by participating
university researchers. Research subjects covered here
include fundamental studies and R&D on fusion structural
materials, high heat-flux materials, and a variety of functional
materials, as well as engineering-oriented studies for thesafety of light water reactors and basic researches supporting
them. Materials studies utilizing radio-isotopes are also being
conducted.

KINKEN Research Highlights 2008 IMR

3) Advanced Research Center of Metallic Glasses
This center was established in May 1987 as the research
center for the development of new advanced materials forthe 21st century, reorganized in May 1996 and April 2005.
The main purpose of the research in this laboratory is toestablish fundamentals and control techniques forsynthesizing artificial substances, especially the metallic
glasses, and to seek for the possibility of applying them as
multi-functional materials for high technology, such as
materials for energy saving, materials for environmental andecology, structural materials, electronic information materials,
biomaterials and materials for social welfare, etc.

4) High Field Laboratory for Superconducting Materials
This laboratory restarted in 2001, succeeding “High Field
Laboratory for Superconducting Materials” which wasestablished in 1981. The main equipment is a hybrid magnet
which generates steady high magnetic fields up to 31T. In
addition, many cryogen-free superconducting magnets whichhave been developed by our laboratory are installed.
Especially, we succeeded in developing the world’s firstcryogen-free hybrid magnet, which generates 27.5 T. The
laboratory also provides instruments for measuring various
physical properties. These facilities are open to scientists andengineers on superconductors and other materials research.

Highlights of Collaborated Research435) International Frontier Center for Advanced Materials(IFCAM)
The material sciences are fundamentally interdisciplinaryand global in nature, involving a variety of academic andengineering fields, in order to achieve better, less-expensive
and more eco-friendly functional matter. Clever development
and wise use of advanced materials are the foundation of all
improved human activity in this new century. Based on theextensive successful research record of IMR, IFCAM wasapproved and founded in 2002 with a mission to serve as aninternational think-tank in materials research. To foster
further international collaboration, liaison offices of IFCAM
have been established at Cambridge University, Harvard
University, Stanford University, the Swedish Royal Instituteof Technology, and the Institute of Physics in China.

7) Center for Computational Materials Science
This center was developed from the Laboratory ofMaterials Information Science which was established in
February 1989. Its main tasks are 1) administration andmaintenance of the supercomputing system in this Institute,2) maintenance of a supercomputing system network, 3)
general support for the usage of the supercomputing system,4) support of materials design by supercomputing simulation
with vectorization and parallelization, 5) construction of
factual database for materials in nonequilibrium phase , and6) support of activity of Nanotechnology group on SINET3
and Asian Consortium for Materials Science, ACCMS-VO.

Fig. Low energy electron microscope in IFCAM.
6) Osaka Center for Industrial Materials research
The Osaka Center was established at Osaka as a special
unit in IMR in April 2006 based on the agreement between
Tohoku University and Osaka Prefecture Government. It is alaboratory dedicated to research and development of nanostructured metallic materials from the application viewpoints
based on the basic research in the material science, chemistry
and physics associated with research organizations in the
Kansai region. Research in the Osaka Center is focused inparticular on the understanding of the fundamental propertiesof nano-structured metallic materials and the rapid realizationof their application to industry, in particular, the small andmedium enterprises in Osaka area.

44

Highlights of Collaborated Research

8) International Collaboration Center (ICC-IMR)
This Center was established in April 2008. The mission isto serve as an international center for materials research andto encourage the international collaboration on researches of
IMR. The International School on advanced materials,
international workshops are coordinated. The Center hosts
the foreign visiting scientists who collaborate on researchwith members of IMR.

IMR KINKEN Research Highlights 2008

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Dislocations of ZnO Single Crystal Examined by
X-ray Topography and Photoluminescence
Damages induced by indentation in ZnO (0001) wafers wereconsidered to lead to dislocated regions from X-ray topographic images.
Photoluminescence intensities of 2.4 eV deep emission band in ZnO
decreased drastically, with increasing in annealing temperature up to800˚C, irrespective of dislocated or nondislocated regions. The
development of a new emission band 2.8–3.0 eV was found in nondamaged ZnO through annealing at 700 and 800 ˚C, which suggests that
dislocations suppress the development of the new peak.
Wide band-gap oxide-semiconductors have attracted keen attention
for optical devices, liquid crystal displays and solar cells. ZnO possesses
a direct band-gap of 3.4 eV as well as a large exciton binding energy of 60
meV, which is about three times larger than that of ZnSe and GaN. Thus,
ZnO is one of the most promising materials for various kinds of applications
in light-emitting devices from blue- to ultraviolet wavelength range, gassensors, ltrasonic oscillators and acoustic wave transducers. For homoepitaxial growth of materials for such applications, the necessity of ZnO
substrates, of high quality with low dislocation density, is drasticallyincreasing. Such dislocations lead to inhomogeneity of their opticalproperties as well known in GaN [1]. Dislocations are considered to berather mobile in ZnO than in GaN and AlN from the difference in theirhardness. However, up to now far less has been known on dynamic andoptical properties of such dislocations in ZnO except a few papers.
X-ray topography is effective to identify dislocations and theirdistribution in materials. Indeed, we applied transmission X-ray topography
to observe dislocations and succeeded in correlating them to their opticalproperties in high quality P-doped ZnTe crystals [2]. The density of grownin dislocations was 7,000 cm–2 approximately. By the photo-luminescence(PL) measurement, it was found that the majority of the phosphorus
impurities exist in the area where many dislocations are observed in X-ray
topographic images, suggesting the probability of impurity accumulation
on dislocations during the growth process. The present study was
undertaken with the above mentioned on X-ray topographic observation
and PL characteristics of dislocations in ZnO crystals.
ZnO single crystals studied in the present work were undoped (ntype, carrier concentration of 3×1015 cm–3, resistivity of 30 Ωcm, mobility of
102 cm2/Vs), grown by the hydrothermal method (Scientific Production
Company GOODWILL, RUSSIA). The ZnO wafers with (0001) surfaces
were in square shape, approximately 10 × 10 × 0.5 mm3 in size, with sides
parallel to [1100] and [11-20]. Density of grown-in dislocations wasapproximately 500 cm–2. First, by the ZnO wafers were thinned bymechanical polishing and then by chemically polishing in a solution of HCl:
H2O = 1:5 for 60 s at room temperature to X-ray transparency. Each wafer
was damaged partly by indentation under 5 N for 5 s with a diamond
Knoop with 0.5 mm interval at room temperature in order to introduce
dislocated regions. Then, such wafers were cut into two parts for annealing
at elevated temperature 500, 600, 700 and 800 ˚C for 5 min under N2 gas
atmosphere (0.1 MPa).
Dislocations of grown-in and indentation-induced in ZnO wafers wereobserved by means of transmission X-ray topography using Borrmann
effect with 11-20 diffraction of Mo Kα1 radiation from an X-ray generator

Fig. 1 Transmission x-ray topograph of high quality undoped ZnO

(a)

(b)

Fig. 2 PL spectra of undoped ZnO single crystal (0001) at 4.2 K in noindented area (a) and indented area (b).
RIGAKU RU-500. The PL measurement was carried out using a grating
monochromator and a photomultiplier at 4.2K. The sample was excited by
a 10-mW He-Cd laser, emitting at 325 nm and density of 150 W/cm2.
Fig. 1 shows transmission X-ray topographic images of high quality
undoped ZnO single crystals. The 11-20 diffraction condition is indicatedin each figure by g vector. Upper part of each wafer was damaged byindentation. In Fig. 1(a) a few dislocations can be seen in the non-indented
area of the as-grown ZnO wafer besides some scratches generated during
polishing preparation. In the as-grown wafer, images of each indentation
are rather sharp and large around the indented center. With increasing
temperature of annealing, such images lose their sharpness and become
small slightly by annealing at 500 and 600 ˚C. In wafers annealed at 800
˚C, images are seen to be elongated along <1120> directions from theindented centers. Thus, it is supposed that some dislocations are induced
around such indented parts, though detailed observations usingtransmission electron microscopy should be conducted to clarify how
dislocations develop from such indentation during annealing.
Figs. 2(a) and (b) show PL spectra obtained from the no-indented andindented areas, respectively, of ZnO (0001) wafers measured at 4.2 K.
Deep emission band around 2.4 eV can be clearly seen in both no-indented
and indented areas in the as-grown wafer. The band is known as green
band that is related to oxygen vacancy (VO) [3]. In an increase in annealing
temperature, PL intensity of the emission band around 2.4 eV decreases inboth no-indented and indented areas. This means that the VO defects
annihilate with increasing annealing temperature. Si multaneously, a newemission band around 2.8–3.0 eV is observed to develop in PL spectra of
no-indented area by the annealing at 700 and 800 ˚C. Contrarily, there isobserved no such emission band developing around 2.8 eV in indented
area. That is, there is no deep emission observed in PL spectra of theindented area annealed at temperatures higher than 700 ˚C. The decreaseof 2.4 eV emission band may be annihilation of oxygen vacancies ormodification to some complexes promoted at elevated temperatures.
Though the origin of the 2.8–3.0 eV emission band is not identified atpresent, if they relate to radiative recombination defects formed at elevatedtemperatures, dislocations may suppress their development by absorbing
defects relevant to the emission band.
References[1] I. Yonenaga, H. Makino, S. Itoh, T. Goto, T. Yao, J. Electron.
Mater. 35, 717 (2006).
[2] K. Yoshino, T. Kakeno, M. Yoneta, I. Yonenaga, J. Mater, Sci.: Mater.
Electron. 16, 445 (2005).
[3] M. Yoneta, K. Yoshino, M. Ohishi, M. Honda, H. Saito, Phys.
Stat. Sol. (C) 3, 1185 (2006).
Contact toKenji Yoshino (Department Electrical and Electoronic Engineering
University of Miyazaki )e-mail: yoshino@pem.miyazaki-u.ac.jp
Ichiro Yonenaga (Physics of Crystal Defects Division)e-mail: yonenaga@imr.tohoku.ac.jp

single crystal (0001). The g vectors are indicated each of the figure.

KINKEN Research Highlights 2008 IMRHighlights of Collaborated Research

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Developments of Optical and Electronic Devices withHigh Performance Based on Si/Ge Heterostructures
Si-based optical and electronic devices with high performanceswere developed by Si/Ge heterostructures. Room-temperatureresonant luminescence with significantly high efficiency and spectral
purity was realized by embedding Ge quantum dots in opticalresonators. Extremely high mobilities exceeding current Si devices werealso obtained from strained Ge channels grown on SiGe virtual substrates.
Si-based light-emitting devices are highly required for realizationof silicon-based optoelectronic integrated circuit. Light emission from
Si is inevitably prevented by its indirect band gap nature. In order toovercome this inherent difficulty, Ge quantum dots (QDs) formed on Si
have been intensively studied. However, room-temperature lightemission efficiency from Ge QDs-based structures is very low and far
from practical applications. Optical resonators are well known to becapable of enhancing the quantum efficiency, where the luminescencein the optical resonator is selectively enhanced by the Purcell effect. In
this study, we combined Ge QDs with optical resonators such as
photonic crystal microcavity (Fig. 1 (a)), microdisk and microring
resonators to modify and enhance the spontaneous emission from Ge
QDs.
Fig.1 (b) shows micro-PL spectra obtained from Ge QDs
embedded microring resonators. Very sharp resonant peaks are
obviously observed in the wavelength range from 1.3 to 1.5 µm while
the reference spectrum obtained from unprocessed region reveals nosuch peak. The sharp peaks were confirmed by numerical simulations
to correspond to whispering gallery mode (WGM) and their quality
factors were found to be as high as 1800-2200. The quality factor was
further enhanced through optimization of the resonator structures up to3300 which was a surprisingly high value for Si-based light emitting
structures. Note that room-temperature photoluminescence withsignificantly high efficiency and spectral purity was obtained from Ge
QDs embedded photonic crystal microcavities and microdisc
resonators [1] as well as the microring resonator.
Channel engineering for carrier mobility enhancement has beenrecognized as essential technology for future CMOS circuits sincefurther scaling is now reaching its limitation. Although the strained-Si

Fig. 2 Strained Ge channel structures on SiGe virtual
substrates fabricated by ion implantation technique.
channel has already been involved into mass production, mobility
enhancements are still not sufficient particularly for holes. Ge channel
is one of the most attractive candidates due to its inherent much higher
mobility over strained Si. Moreover, strain introduction is known tofurther enhance the mobility. Development of strain-induced Ge
channel is, therefore, another target of our research project.
The typical structure fabricated here is illustrated in Fig.2. Ion
implantation technique was developed to produce high quality thin
SiGe virtual substrates which were indispensable for realization ofstrained-Ge channel devices. This technique was based on intentional
introduction of ion-implantation defects into the Si substrates or SiGe
virtual substrates. The epitaxial overgrowth was performed on the ionimplanted substrates, leading to effective strain relaxation of theovergrown layer. We obtained considerable relaxation enhancements
by the ion implantation with optimum conditions and realized SiGe
virtual substrates with one order of magnitude smaller surface
roughness (~0.3 nm) and thickness (~100 nm) than conventional
virtual substrates.
High quality Ge channels with extremely high compressive strains
up to 3% were formed on SiGe virtual substrates by low temperature
growth via MBE. Room temperature mobilities higher than 1500 cm2/
Vs were obtained on the developed SiGe virtual substrates withextremely small thickness, which opened the route to practical device
applications of strained Ge channels. We also showed the maximum
channel mobility exceeds 3000 cm2/Vs, the highest value everreported. Moreover, systematic study of magnetotransport properties
demonstrated that the hole effective mass can be reduced byincreasing the strain [2]. It was also found that the effective mass
sensitively changes with hole density modification through gating [3].
In conclusion, both optical and electronic devices were developed
based on Si/Ge heterostructures and excellent room-temperature
properties were presented, indicating the introduction of Si/Ge is themost promising way for realization of Si-based optoelectronic devices.
References[1] J. S. Xia, K, Nemoto, Y. Ikegami, N. Usami and Y. Shiraki: Appl.
Phys. Lett. 91, 011104 (2007).
[2] K. Sawano, Y. Kunishi, K. Toyama, T. Okamoto, N. Usami, K.
Nakagawa and Y. Shiraki: Physica E 40, 2122-2124 (2008).
[3] K. Sawano, Y. Kunishi, Y. Satoh, K. Toyama, K. Arimoto, T.
Okamoto, N. Usami, K. Nakagawa and Y. Shiraki: Appl. Phys. Express
1, 011401 (2008).

Fig. 1 (a) Schematic of Ge QDs embedded photonic crystal
microcavity. (b) PL spectra obtained from microring structure.

46

Highlights of Collaborated ResearchContact toYasuhiro Shiraki (Musashi Institute of Technology)
e-mail: yshiraki@sc.musashi-tech.ac.jp
Noritaka Usami (Crystal Physics Division)e-mail: usa@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

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Improvement of Biocompatibility of Titanium Implants by
Calcium Phosphate Coating Using RF Magnetron Sputtering
*

0.6
0.5
0.4
0.3

0.1
0

4w

0.2

2w

Titanium materials can be directly connected to living
remodeling bones at the optical microscopic level, i.e.,
osseointegration. However, it is known that more than threemonths are required to obtain a strong fixation between
titanium implants and bones and the strength of the fixation
might be influenced by the state of bones. Coating calcium
phosphate on titanium implants is one of the methods toimprove the osseointegration of titanium implants. Amorphous
calcium phosphate (ACP) fabricated by radiofrequency (RF)magnetron sputtering is a candidate coating material fortitanium implants.
An increase in the percentage of bone-implant contact
was demonstrated in titanium cylinders with an ACP coating
film by performing animal experiments in which the cylinders
were implanted into the mandibles of beagle dogs[1]. This
increase in the percentage as compared to that for uncoatedtitanium cylinders was observed 8–12 weeks after implantation.
ACP-coated titanium implants were suggested to effectively
exhibit a rapid and strong fixation to bones.
The blast treatment of titanium implants is a common
technique to improve the adhesion of the bonding between thetitanium implants and bones due to an increase in surface
roughness. Therefore, the animal experiments were performed
using ACP-coated screw-type Ti-6Al-4V implants with a blasttreated surface. The coated Ti-6Al-4V specimens were
implanted into the femur of Japanese white rabbits[2]. The
thickness of the ACP coating film was 0.5 µm and the averageroughness of the Ti-6Al-4V implants was 4.6 µm. Figure 1
shows the cross section of the ACP film fabricated on theblast-treated Ti-6Al-4V plate. The ACP coating film can be
efficiently coated over the entire blast-treated Ti-6Al-4V
surface. The surface being rough even after coating is one of
the advantages of the RF magnetron sputtering process,which can uniformly fabricate a thin coating film.

0.7

Removal torque, TR / N•m

An amorphous calcium phosphate film coated on screwtype Ti-6Al-4V implants with a blast-treated surface by using
radiofrequency magnetron sputtering was evaluated in vivo.
The coated specimens were implanted into the femur ofJapanese white rabbits. The removal torque of the implants
was statistically improved by coating the amorphous calcium
phosphate film two weeks after implantation.

Control

ACP

Fig.2 Removal torque of uncoated (control) and ACP-coatedscrew-type Ti-6Al-4V implants with blast-treated surface fromfemur of Japanese white rabbits after implantation for two andfour weeks. (*p < 0.05)

Resin
Bone

Blast-treated Ti-6Al-4V implant

2 μm

Fig.3 Scanning electron micrograph of cross-section of ACPcoated screw-type Ti-6Al-4V implant with blast-treated surface
after removal torque measurement (four week implantation).
Figure 2 shows the removal torque of the coated implantsfrom the femur of a Japanese white rabbit. It increases with theduration of implantation, and as compared to the removal
torque of the uncoated titanium implant (control), it is statistically
improved by coating the ACP films two weeks after the
implantation. The cross-section of the titanium implants coatedwith the ACP film after removal torque measurements is shownin Fig.3. A smooth and uniform interface is observed between
the new bone and the blast-treated implant.
These in vivo studies suggest that the coating of a sputtered
ACP film on titanium implants is effective for the rapid andstrong fixation of blast-treated titanium implants to bones.
References[1] K. Ueda, T. Narushima, T. Goto, M. Taira and T. Katsube,
Biomed. Mater., 2, S160 (2007).
[2] T. Narushima, K. Ueda, T. Goto, T. Katsube, H. Kawamura,
H. Nakagawa and M. Taira, J. Jpn. Soc. Power and Powder
Metall., 55,318 (2008).

Fig.1 Scanning electron micrograph of cross-section of ACP film
coated on blast-treated Ti-6Al-4V plate.

KINKEN Research Highlights 2008 IMRContact to
Takayuki Narushima (Department of Materials Processing,Tohoku University)e-mail: narut@material.tohoku.ac.jp
Takashi Goto (Multi-functional Materials Science Division)e-mail: goto@imr.tohoku.ac.jp

Highlights of Collaborated Research

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Exchange Coupling and Energy-Level Crossing of 4f-3d
Heterometallic Coordination Compounds
We have reported several 4f-3d heterometallic complexes,
such as [Dy2Cu], [Dy4Cu], and [Dy2Cu2]n, behaving as asingle-molecule magnet (SMM). Though several 4f-3d-based
SMMs are known to date, the magnitude of exchange
coupling has never been determined before our work. We
have proposed a plausible mechanism for the quantum
tunneling of the magnetization for the above compounds
based on the well-defined exchange couplings.
SMM exhibits magnetic hysteresis from the singlemolecular origin; namely, slowing down of the rate ofmagnetization reorientation due to a high energy-barrier.
Lanthanide ions seem just suitable for development of SMMs
because the barrier is related to strong magnetic anisotropy
and large spins. Actually, the above Dy-Cu compounds
showed slow magnetization reorientation, as clarified bypulsed-field and ac magnetic measurements [1-3]. Quantum
tunneling of magnetization, observable as magnetization
jumps in a hysteresis curve, is also characteristic of SMMs.
We investigated the exchange couplings and energy-levelstructures of these compounds. High-frequency (HF)-EPR
results on linear [Dy2Cu] and [Dy2Ni] are presented here forexample. For their X-ray crystal structures, magnetization
curves with jumps, and other details, see the references.
Figure 1a shows the frequency vs. magnetic field diagramfor [Dy2Cu] together with the EPR spectra superposed. The
resonance field was proportional to the frequency, as
expected from the Zeeman effect. The g value was 2.07 fromthe slope, being compatible with the Cu signal. Extrapolation
gives a critical field of 1.36 T. It is identical to the position ofa 2-µB jump due to the Cu spin-flip in the magnetization
curves. Therefore, [Dy2Cu] has an energy-level crossing fromthe ground ferrimagnetic state to a ferromagnetic one at 1.36
T. By treating the Dy moments as Ising spins with |Jz| = 15/2,
an exchange coupling model can be built involving a uniqueparameter (Fig. 2). We determined JDy-Cu/kB to be –0.126 K.
Figure 1b shows the result on [Dy2Ni]. We found a few g≈ 2 lines with zero-field splitting and “forbidden” g ≈ 4 lines. In
sharp contrast to the [Dy2Cu] case, the resonance fields are
shifted by a negative bias field, suggesting the presence offerromagnetic coupling. The data were analyzed using threeparameters (J, D, and E), to give JDy-Ni/kB = +0.031 K. The
level-crossing fields in the Zeeman diagram completely
correspond to the magnetization jump positions.
Very recently the EPR and magnetic study on [Dy2Cu2]nhas been published [3]. In spite of an apparent complex
situation (three independent exchange pathways, two Dy-Cu
and one Cu-Cu, are present), our concept totally holds forthis system. We successfully determined two JDy-Cu’s.
The HF-EPR technique is very useful for evaluation of
the exchange coupling in such compounds. We could not
directly observe Dy signals, but we clarified the environmentof Cu and Ni ions as the exchange bias field from adjacent
Dy ions. In our systematic research on 4f-3d magnets,
determination of exchange couplings of isomorphous
derivatives is of great interest. Studies on [Ln2Cu], [Ln2Ni],

48

Highlights of Collaborated Researchand [Ln2Cu2]n involving Ln = Gd, Tb, Ho, Er etc. in place of
Dy are now in progress.

（a）

（b）

Fig.1. Selected EPR spectra measured at 4.2 K as a functionof frequency for (a) [Dy2Cu] and (b) [Dy2Ni]. The spectra areoffset in a linear scale of the frequency. Dotted lines are
drawn for a guide to the eye. Structural formula of [Dy2Cu] isalso shown in (a).

Fig.2. Energy diagram at zero-field for [Dy2Cu]. Spinstructures are drawn with arrows.

References[1] F. Mori, T. Nyui, T. Ishida, T. Nogami, K.-Y. Choi, and H.
Nojiri, J. Am. Chem. Soc., 128, 1440 (2006).
[2] S. Ueki, T. Ishida, T. Nogami, K.-Y. Choi, and H. Nojiri,Chem. Phys. Lett., 440, 263 (2007).
[3] A. Okazawa, T. Nogami, H. Nojiri, and T. Ishida, Chem.
Mater., 20, 3110 (2008).
Contact to
Takayuki Ishida (The Univ. of Electro-Communications)e-mail: ishi@pc.uec.ac.jp
Hiroyuki Nojiri (Magnetism Division)e-mail: nojiri@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

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Pressure-induced Quantum Critical Phenomenain the f Electron Compound CeTX2
Effect of pressure on the electronic states has been investigatedin CePtSi2 and CeNiGe2 which crystallizes in an orthorhombic
layered structure. The external pressure enhances the Kondo
temperature TK in both compounds. A pressure induced crossover
exists in the electronic state around Pv~ 1.0 GPa in CePtSi2.
CeTX2 crystallizes in an orthorhombic CeNiSi2-type layered
structure (Cmcm), which is constructed from the deformed fragmentsof CeGa2Al2 and α-ThSi2 structures [1]. The long b-axis induces theformation the flat Brillöuin zone, which might produce quasi-two
dimensional electronic state and magnetic instability play animportant role. In these compounds, competition between the
Ruderman-Kittel-Kasuya-Yoshida (RKKY) interaction and Kondo
effect was discussed by Doniach as a function of |JN(0)|, where J
and N(0) are the exchange interaction between the conduction
electron and the localized 4f spin and the density of states at the
Fermi level, respectively. The external pressure is one of the control
parameters to tune the value J, which leads a long range magneticallyordered state, or results in a screening of the 4f moment byconduction electrons. In the present work, we describe the effect ofpressure on the electronic states of CePtSi2 and CeNiGe2.
CePtSi2 has the large electronic specific heat coefficient about
γ= 600mJ/molK2 and the antiferromagnetic order near TN ~1.5 K.
Fig. 1 shows the temperature dependence of the magnetic
contribution of CePtSi2 to the electrical resistivity ρ due to 4f
electrons, ρmag, which is estimated by using the relation, ρmag =ρ(CePtSi2) - ρ(LaPtSi2). ρ(LaPtSi2) is assumed to indicate only thephonon contribution and to be independent of pressure. ρmagincreases with decreasing temperature until it shows maximum (T2)around 28 K at ambient pressure. Below T2, a small shoulder (T1) isobserved in the ρmag at pressure below 0.4 GPa. T1 and T2 areshown by arrows in the inset of Fig. 1.

quantum critical behavior is observed and the superconducting
temperature Tsc(P) is enhanced around Pv~ 4.5 GPa [3].
Figure 3 shows the AC magnetic susceptibilities along the easy
c-axis of CeNiGe2 which orders magnetically at TN1~ 4.0 K andundergoes a spin structure rearrengement at TN2~ 3.0 K at ambientpressure. Here TN1 and TN2 are defined as the temperature wheredχc /dT shows a kink while Tmax is estimated as that where dχc /dT =0 [4]. At ambient pressure, dχc /dT shows sharp anomalies at TN1
and TN2, while TN1 is almost equal to Tmax (= 4.0 K). As the pressure
increases, some uncertainty as to the value of Tmax arises because
the magnitude of dχc /dT is suppressed and becomes zero in a wide
temperature region. On the other hand, TN1 is suppressed byapplying pressure but tends to saturate above 1.0 GPa. Thisbehavior is contrary to the case of CeAl2, CeRh2Si2, etc., in which
the antiferromagnetic ordering temperature decreases as thepressure increases and quantum critical behavior is observed at thepressure near TN ~0 K. Tmax is almost the same as TN1 up to 0.68
GPa, indicating that Tmax is related not only to the Kondo effect but
also to the antiferromagnetic ordering. However, above 1 GPa, Tmax
increases with pressure and a rate of ∂Tmax/ ∂P tends to be enhanced
by applying pressure. Assuming that Tmax indicates TK in thepressure region, Tmax is proportional to exp(1/|JN(0)|); the increasein Tmax indicates the enhancement of |JN(0)|. In other words, theelectronic state becomes unstable as pressure increases.

Fig.3 AC magnetic susceptibility of CeNiGe2 under pressure.
In summary, it is found that TK increases with pressure in bothcompounds. In CePtSi2, a pressure induced crossover exists in theelectronic state around Pv~ 1.0 GPa. In CeNiGe2, on the other hand,further measurements above 2.0 GPa are in progress in order todetermine a quantum critical behavior.
Fig.1. Temperature dependence of the ρmag of CePtSi2 at high pressure.

References[1] M. Ohashi, G. Oomi, K. Ishida, T. Komatsubara, I. Satoh, J.
Alloys and Compounds 408-412, 84 (2006).
[2] H Miyagawa, Y Fuchizaki, M Ohashi, G. Oomi, I. Satoh, T.
Komatsubara, High Pressure Research 26, 503 (2006).
[3] A. T. Holmes, D. Jaccard, K. Miyake, Physical Review B 69,
024508 (2004).
[4] M. Ohashi, G. Oomi, I. Satoh, Journal of the physical society ofJapan 76, 114712 1-3 (2007).

Fig.2. Pressure dependence of T1 and T2 of CePtSi2.

Contact to
Masashi Ohashi (Faculty of Environmental Design, Kanazawa

Fig. 2 shows T1 and T2 as a function of pressure. T2 is almost
constant below 1.0 GPa, but T1 increases with applying pressure andmerges into T2 around Pv~ 1.0 GPa, where the Kondo temperature
TK is close to TKh, a modified Kondo temperature due to the crystal
electric field [2]. Similar behavior is found in CeCu2Si2, in which a

KINKEN Research Highlights 2008 IMR

University)e-mail: ohashi@t.kanazawa-u.ac.jp
Isamu Satoh (Laboratory of α-Ray Emitters)
e-mail: satoh@imr.tohoku.ac.jp

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49

3FTFBSDI

Development of Co-Cr-Mo Alloys Containing Interstitial Elementswith High Strength and High Ductility for Biomedical
Metallic biomaterials have been widely used fororthopedics and dental applications. Among various
biomaterials, Co-Cr-Mo alloys are especially applied forload- and wear-bearing implants. The higher reliability of theimplants is strongly required from the point of mechanicalproperties as well as biocompatibility recently.
Conventional cast and wrought Co-Cr-Mo alloys containNi to improve their castability and workability. However, Ni
has been reported as one of the most common contact
allergens. The cast alloys (ASTM F-75) show low ductility
and the Co-based dental casting alloys occasionally fractureduring small plastic deformation caused by the adjustment.
Although Ni-free Co-Cr-Mo alloys have developed to avoid
the allergic problem, the alloys still show the limited ductility
owing to the shrinkage porosity and interdendritic segregation,
which mainly consists of the σ phase. In addition, Ni-free CoCr-Mo alloys contains a large amount of ε martensite (hcp) in
the γ (fcc) matrix, thereby restricting the deformation of thealloys [1,2]. Therefore, the suppression of the σ phaseformation and the stabilization of the γ phase at room
temperature are the key to enhance the mechanical propertiesof the alloy [3].
We focus boron and nitrogen as additives to the Co-CrMo alloys, which are the interstitial elements and show no
toxicity, to stabilize the γ phase (fcc). A large amount ofstacking fault and ε martensite form along (111) plane in the
Co-29Cr-6Mo alloy. Figure 1 shows TEM micrographs of Band N- added Co-29Cr-6Mo alloy. ε martensite disappears
and the γ phase is confirmed with low fault density in the Co29Cr-6Mo alloy. In addition, the σ phase precipitation was
confirmed to be suppressed by OM observation.
Figure 2 shows nominal stress-strain curves of various
Co-Cr-Mo alloys. By the addition of B, elongation is improved
although the yield strength is almost the same as Co-29Cr6Mo. When N and B added to Co-29Cr-6Mo, both yieldstrength and elongations are clearly improved compared toCo-29Cr-6Mo. Therefore, it is confirmed that the addition ofboron and nitrogen is very effective for the improvement of
mechanical properties of Co-29Cr-6Mo alloys.
References[1] S-H. Lee, E. Takahashi, N. Nomura and A. Chiba, Mater.
Trans. 46, 1790 (2005).
[2] S-H. Lee, T. Uchikanezaki, N. Nomura, M. Nakamura and
A. Chiba, Mater. Trans. 48, 1084 (2007).
[3] S-H. Lee, N. Nomura and A. Chiba, Mater. Trans. 49, 260(2008).

50

Highlights of Collaborated Research

Fig. 1 TEM micrograph of B- and N- added Co-29Cr-6Mo
alloy.

Fig. 2 Stress-strain curves of B- and /or N-added Co-Cr-Mo
alloys.

Contact to
Naoyuki Nomura (Institute of Biomaterials and Bioengineering,Tokyo Medical and Dental University)e-mail: nnomura.met@tmd.ac.jp
Akihiko Chiba (Deformation Processing Devision)e-mail: a.chiba@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

Neutron Damage Resistant SiC/SiC Composites by NITE Method
Based on the JMTR, JOYO, HFIR and other reactor
irradiation and simulation radiation damage studies, neutron
damage tolerant SiC/SiC composites have been designed
and successfully fabricated by NITE method. They provide
excellent mechanical properties, thermal and chemical
stability, as well as gas tightness. Also model components
for nuclear reactors were made and their excellent
performance was confirmed.
SiC has a long history as one of the most attractive
candidates for nuclear energy systems and many R & D has
been carried out together with strong push by the invention of
PCS-type SiC fibers by late Professor S. Yajima. Throughout
the Japan-US collaborative activities on Fusion Materials, SiC
and SiC/SiC R & D was extensively carried out and many
insights about material behavior under neutron bombardment
and reactor environments had been reported. Figure 1 is an
example showing the excellent radiation damage resistanceunder neutron bombardment, where SiC/SiC made by CVI
method kept its strength over 10dpa at elevated temperatures.
Based on these neutron radiation effect results and simulation
radiation effect study results, R & D direction of radiation
damage tolerant SiC/SiC composites design and fabrication
has been defined. The R & D direction is to fabricate (1) highcrystallinity, (2) high purity, (3) near stoichiometry and (4) nearfull density in SiC fiber and SiC matrix together with the soundC or C/SiC interphase. At the same time, an innovative
fabrication method, patented under the name of NITE (NanoInfiltration and Transient Eutectic Phase) method, has beendeveloped for realizing the R & D direction by inexpensive
process. This technology was transferred to a venture businessfrom Kyoto University for establishing strong industrial basis toprovide sufficient amount for the real utilization.
One of the important properties of NITE-SiC/SiC is theexcellent gas tightness which indicates, for the first time, thepotentiality to be applied for gas reactors and gas cooled
blanket system for fusion reactors. More than 6 orders ofmagnitude improvement in He permeability has been
confirmed, as shown in Figure 2, and the stability of thistightness even under the loading larger than proportional
limit strength of SiC/SiC has been also published.

Fig. 2 He gas permeability of SiC materials.

Fig. 3 Shape flexibility of NITE SiC/SiC.
Some of the typical SiC/SiC model components to be
applied to Generation IV nuclear reactors and fusion reactors
are shown in Figure 3, where thin tube for fuel pin applicationand its excellent machinability and excellent joint strength
are indicated.
Porous SiC/SiC for TRISO fuel compartment for lateralflow gas cooling reactor core and full dense SiC/SiC for sizemodel of intermediate heat exchanger (IHX) for many energy
conversion systems have been also introduced. The effort
starting from IMR-Oarai is one of the strong bases of thenuclear material community of Japan.
References[1] Y. Katoh, T. Nozawa, L.L. Snead, T. Hinoki, A. Kohyama,
Fusion Engineering and Design 81, 937 (2006).
[2] K. Ozawa, T, Nozawa, Y. Katoh, T. Hinoki and A.
Kohyama, J. Nucl. Mater., 367-370, 713 (2007).
[3]J.S. Park, A. KohyamaT. Hinoki and K. Shimoda, J. Nucl.
Mater., 367-370, 719 (2007).

Fig. 1 Neutron damage effect on ultimate strength of SiC
materials.

KINKEN Research Highlights 2008 IMRContact to
Akira Kohyama (IAE, Kyoto University)e-mail: kohyama@iae.kyoto-u.ac.jp
Minoru Narui (International Research Center for Nuclear
Materials Science)e-mail: narui@imr.tohoku.ac.jp

Highlights of Collaborated Research

51

3FTFBSDI

Magnetic behavior near the superconducting upper
critical field Hc2 of NpPd5Al2 has been studied at low
temperature and high magnetic field. The strong XY-typemagnetic anisotropy in the normal state suddenly
disappears below Hc2. This result demonstrates that thesuperconductivity in NpPd5Al2 is carried by electrons having
large magnetic moments with 5f character.
NpPd5Al2 is the first neptunium-based heavy fermion
superconductor [1].This compound shows superconductivity
at a fairly high superconducting transition temperature 5 Kwith typical heavy fermion behavior characterized by thelarge electronic specific heat in the normal state. It is also
surprising that the isostructural compounds with the samestructure have never been reported either in rare earth or
actinide systems. Immediately after this discovery, scientists
explored the An-Pd-Al phase again and revealed the existenceof several actinide and rare earth AnPd5Al2 compounds like
UPd5Al2 (An : rare earth or actinide elements) [2].
Compared to ordinary superconductors, the electrons in heavy
fermion superconductors, in general, have the larger effective
mass and the larger magnetic moment, both arising from thehybridization with the magnetic f-electrons. In the spin-singlet
superconducting state, the large magnetic moment disappears byforming a Cooper pair with up- and down-spin electrons. Therefore,
heavy fermion superconductivity is generally broken by magnetic
field to gain Zeeman energy of unpaired quasiparticles. In NpPd5Al2,
particularly, this occurs as a first-order phase transition [1].
In NpPd5Al2, the magnetic moments in the normal state
have the XY-type anisotropy, namely, they tend to lie in thetetragonal basal plane and cannot be oriented along the[001] direction. It is therefore interesting to detect the change
in the magnetic anisotropy through superconducting
transition. Fig 1 shows the magnetic torque with the magnetic
field applied around the [001] direction. In the normal state
above the upper critical field Hc2, large magnetic torque isobserved, which is attributed to the XY-type anisotropy. The
present torque diminishes quickly with decreasing magnetic
field through Hc2. Below Hc2 there is no sizable torquedetected, indicating quite isotropic magnetic behavior in the
superconducting state. The magnetic field where the torqueanomaly occurs exactly follows the Hc2 determined from theresistivity, as shown in Fig. 2. Combining with the previous
magnetization measurement [1], it is concluded that the large
magnetic moment existing in the normal state disappears inthe superconducting state. This result strongly suggests thatthe superconductivity in NpPd5Al2 is realized by the electrons
with extremely large anisotropic magnetic moment.
This work was performed in collaboration with Osaka
university.

52

Highlights of Collaborated Research

Magnetic Torque (arb. units)

Anomalous Behavior at the Superconducting Upper Critical Field
in Transuranium Superconductor NpPd5Al25-15.5˚

NpPd5Al2
0.1 K

-11˚
Hc2

-6.5˚
-2˚

0

2.5˚
0

50

100

150

Magnetic Field (kOe)
Fig. 1 Magnetic torque measured under magnetic field around[001] direction in NpPd5Al2. Arrows show the uppercriticalfield Hc2

Fig. 2 Left: Angular dependence of the torque anomaly iscompared to that of Hc2 determined from resistivity
measurements. Right: Crystal structure of NpPd5Al2.

References[1] D. Aoki, Y. Haga, T.D. Matsuda, N. Tateiwa, S. Ikeda, Y.
Homma, H. Sakai, Y. Shiokawa, E. Yamamoto, A. Nakamura,
R. Settai and Y. Onuki, J. Phys. Soc. Jpn. 76 063701(2007).
[2] Y. Haga, D. Aoki, Y. Homma, S. Ikeda, T.D. Matsuda, E.
Yamamoto, H. Sakai, N. Tateiwa, N.D. Dung, A. Nakamura,Y. Shiokawa and Y. Onuki, J. Alloys. Compds. (2008) inpress.
Contact toYoshinori Haga (Advanced Science Research Center, JAEA)
e-mail: haga.yoshinori@jaea.go.jp
Yoshiya Homma (Radiochemistry of Metals Division)e-mail: yhomma@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

Viscous Flow Behaviours of Zr55Cu30Al10Ni5 Bulk Metallic Glasses
over the Entire Temperature Range

Zr55Cu30Al10Ni5 bulk metallic glasses (BMGs) have especiallyattached great interest due to their excellent properties such asvery high glass forming ability that the alloys enable the formationof centimeter-sized cylindrical BMG samples and their excellent
mechanical properties [1]. It has also been reported recently thata centrifugal casting method from the equilibrium liquid state of
the BMGs is effective on fine precision molding [2]. In order toclarify the bulk glass forming process and also the fine precisionmolding process from the liquid state, the most important factor
is the temperature dependence of the viscosity over the entire
temperature range. We have reported the highly reliable viscosityof the Zr55Cu30Al10Ni5 supercooled liquids [3-5]. In the present
study, temperature dependence of the viscosity of the
Zr55Cu30Al10Ni5 BMGs has been examined over the entire
temperature range containing supercooled liquid by a penetration
viscometer and the equilibrium liquid regions by a .containerlesselectrostatic levitation method in high vacuum atmosphere [6].
Figure 1 shows the viscosity (η) of the Zr55Cu30Al10Ni5
supercooled liquids as a function of temperature at various
heating rates of 20, 200 and 400 °C/min. In each heating
condition, viscosity was measured under various applied loads
of 0.049, 0.098, 0.196 and 0.294 N on the cylindrical indentation
probe with a diameter of 1.0 mm. As shown in this figure, viscosity
was quite independent of these applied loads under the various
heating conditions. However, when the sample was slowly
heated at the rate of 20 °C/min, viscosity exhibited relatively highvalues that may be mainly due to the skin effects of oxides on thesample surface. With increasing the heating rate to 200 °C/min,viscosity (η) largely decreased and the ln η decreased linearlywith increasing the temperature up to about 470 °C and thenincreased with increasing the temperature for their crystallization.
At the heating rate of 400 °C/min, viscosity was also exhibited
relatively high values at a low temperature region between 400
and 450 °C. This may be due to the non-equilibrium liquid state
under the high-speed heating condition. In the temperature
region between 450 and 490 °C, viscosity fitted precisely to thelinear relationship measured under

the heating rate of 200 °C/min indicating the equilibrium
viscosity may be measured in this temperature region. Figure 2
shows the viscosity of the Zr55Cu30Al10Ni5 equilibrium liquid as afunction of temperature that was obtained by the conteinerlesselectrostatic levitation method. The equilibrium viscosity exhibit
very low values that are about 1x1010 times lower than that of thesupercooled liquids. Figure 3 shows a fragility plot of the viscosityof the Zr55Cu30Al10Ni5 alloys over the entire temperature range incomparison with some metallic and nonmetallic glass forming
liquids. In this case, normalized Tg = 692 K (419 °C ) was
obtained by a Vogel-Fulcher-Tammann (VFT) relationship. As
clearly shown in this figure, the Zr55Cu30Al10Ni5 alloy behavescloser to the strong glasses than the fragile glasses.
References[1] Y. Yokoyama, E. Mund, A. Inoue and L. Schlts, Mater. Trans.
48, 3190 (2007).
[2] Y. Furuya, KINKEN Research Highlights 55 (2007).
[3] T. Yamasaki, S. Maeda, Y. Yokoyama, D. Okai, T. Fukami, H.
M. Kimura and A. Inoue: Mater. Trans. 46, 2746 (2005).
[4] T. Yamasaki, S. Maeda, Y. Yokoyama, D. Okai, T. Fukami, H.
M. Kimura and A. Inoue, J. Intermetallics, 14, 1102 (2006).
[5] T. Yamasaki, S. Maeda, T. Fukami, Y. Yokoyama,. H. M.
Kimura and A. Inoue: Mater. Trans. 48, 1834 (2007).
[6] T. Ishikawa and P. F. Paradis, Proc. of the 20th Space
Utilization Symposium, 20, 263 (2004).
Contact toTohru Yamasaki (Department of Materials Science andChemistry, University of Hyogo)e-mail: yamasaki@eng.u-hyogo.ac.jp
Yoshihiko Yokoyama (Advanced Research Center for
Metallic Glasses)
e-mail: yy@imr.tohoku.ac.jp

0.121100 1200 1300 1400 1500 1600 1700 °C

0.1

Viscosity, η /Pa・s

Viscous flow behaviors of Zr55Cu30Al10Ni5 bulk metallic
glasses over the entire temperature range containing
supercooled liquid region and the equilibrium liquid region has
been examined. Viscosity has been measured by using apenetration viscometer under various heating rates in the
supercooled liquid region and a containerless electrostatic
levitation method in the equilibrium liquid region.

0.08
0.06
0.04
0.02
0
1300 1400 1500 1600 1700 1800 1900 2000 2100

Temperature,T /K

1012
400 ℃/min
20 ℃/min

1010
109

●

0.049 N■

0.098 N●

0.196 N◆

0.294 N

107
106
105

101２
1010

108

200 ℃/min
200 ℃/min

400 ℃/min

400 420 440 460 480 500 520 540
Temperature, T / ℃

Fig. 1 Viscosity (h) of the Zr55Cu30Al10Ni5 supercooled liquids
as a function of temperature at various heating rates of 20,
200 and 400°C/min.

KINKEN Research Highlights 2008 IMR

Viscosity η (η / Pa・s)s
Viscosity, η / Pa・

1011

Fig. 2 Viscosity of the Zr55Cu30Al10Ni5 equilibrium liquids asa function of temperature measured be the containerlesselectrostatic levitation method.

108
106
104

Strong

Zr55Cu30Al10Ni5
Zr46.75Ti8.25Cu7.5Ni10Be27.5
SiO2

Present study
Supercooled
liquid region

Na2O・2SiO2

102

Fragile

100
10-2
10-4

O-terphenyl

Present study
Equilibrium
liquid region

0.4

0.6

Glycerol

0.8

1.0

Tg/T

Fig. 3 Fragility plot of viscosity of the Zr-Cu-Al-Ni BMGs.

Highlights of Collaborated Research

53

3FTFBSDI

Elucidation of Magnetization Mechanism in Nd-Fe-B Thin Films
by in-field Magnetic Force Microscopy
This article describes our recent results on Nd-Fe-B “hardmagnetic” thin films. The samples were prepared onto single
crystal MgO substrates using an UHV-compatible dcsputtering system. A detailed magnetization behavior wasobserved by a magnetic force microscope (MFM) speciallyequipped with electro-magnet, which enables us to measurein-field magnetic domain structure.
Since a discovery of Nd-Fe-B permanent magnet [1], it has
been widely and practically used as magnetomotive force forapplications such as motors, actuators and sensors, becauseof its large magnetocrystalline anisotropy and high maximum
energy products. From the recent environmental and energy
issues, the consumption of Nd-Fe-B sintered magnets has
been increased by the expansion of hybrid vehicles’ (HVs)market. However, especially for HVs, in order to enhance thecoercivity Hc at high temperature environments, the additionof heavy rare earth elements such as Tb and Dy to the Nd-FeB sintered magnets is essential. Recently, another problem of
the amount of material resources has been occurred and it
urges us to reduce the amount of these elements. In addition,
the value of coercivity of Nd-Fe-B commercial magnets is still
far from the theoretical prediction of Nd2Fe14B alloy [2]. In thisstudy, in order to compensate the gap between the experimental
and theoretical value of coercivity, in other word, to elucidate
the mechanism of the coercivity of Nd-Fe-B magnets, we haveinvestigated the magnetic domain structure of the Nd-Fe-B
thin films by magnetic force microscope specially equippedwith electro-magnet [3].
All the samples were prepared using an UHV-compatible
dc-sputtering system (base pressure < 1 × 10−8 Pa). A Cr seedlayer of 1 nm and a Mo buffer layer of 20 nm were firstly
deposited onto single-crystal MgO(110) substrates at room
temperature. The nominal thickness of Nd-Fe-B layer t was 30
nm. The films were heated at 625 °C during the deposition of

Nd-Fe-B layer. The magnetic properties were measured in theapplied magnetic field perpendicular to the film plane by using
a superconducting quantum interference device (SQUID)
magnetometer in the field up to 55 kOe (for some cases, 70
kOe). The structural analysis was performed by an x-ray
diffraction with Cu-Kα radiation.
Atomic force microscope (AFM) and MFM images of NdFe-B thin film are shown in Fig. 1. Bright and dark contrast in
MFM image represent an upward and downward magnetization
direction, respectively. Magnetic fields during this measurement
were applied from 0 to 6 kOe and 6kOe to -6kOe. A remarkablemaze-like domain pattern was observed at virgin state of thefilm as shown in Fig. 1(b). With increasing an applied field, theregion of dark stripe was gradually diminished (Fig.1 (b) to (d)).
However, the stripes began to appear again after removing theapplied field (Fig.1 (e)). It is noted that this stripe region is also
observed at sintered magnet and this region is thought to be a“weak” hard magnetic region in the Nd-Fe-B thin film. The
reduction of this “weak” region is thought to plays an importantrole to elucidate the magnetization mechanism.
References[1] M. Sagawa, N. Togawa, H. Yamamoto, and Y. Matsuura, J.
Appl. Phys., 55, 2083 (1984).
[2] H. Kronmüler, Phys. Stat. Sol., 144, 385 (1987).
[3] H. Ishioka, T. Sato, H. Kato, H. Fujino, and T. Shima, Digest
book of Intermag2008, pp.230 (2008).
Contact toToshiyuki Shima (Faculty of Engng., Tohoku Gakuin
University)e-mail: shima@tjcc.tohoku-gakuin.ac.jp
Koki Takanashi (Magnetic Materials Division)e-mail: koki@imr.tohoku.ac.jp

Fig. 1. Magnetic force microscope images for Nd-Fe-B thin films. The film thickness was 30 nm. In-field magnetic fields wereapplied during the magnetic domain observation. (a) Surface morphology (AFM image), MFM image of (b) virgin state, MFM
images measured at Hex= (c) 3kOe, (d) 6kOe, (e) 0kOe, (f) -1kOe, (g) -3kOe and (h) -6kOe. Bright and dark contrast in MFM image
represent an upward and downward magnetization direction, respectively.

54

Highlights of Collaborated Research

IMR KINKEN Research Highlights 2008

3FTFBSDI

Immobilization of Poly(ethylene glycol) to Metal Surfaces with
Electrodeposition
To bifunctionalize the surface of metals, poly(ethylene
glycol), PEG, molecules are immobilized to titanium, cobaltchromium-molybdenum alloy and 316L type stainless steel
with electrodeposition. PEG molecules strongly and efficiently
immobilized to titanium surface with electrodeposition,compared than that with immersion. In addition, concentrationof active hydroxyl groups on surface oxides on metals influences
the immobilization amounts of PEG. PEG-immobilized surface
inhibited the adsorption of protein on titanium.
Recently, biofunctions are required for metals for biomedical
use. For example, stents are placed at stenotic blood vessels for
dilatation, and blood compatibility or prevention of adhesion ofplatelet is necessary. In guide wires and guiding cathetels, sliding
lubrication in blood vessel is important when those are inserted
into there. In addition, if metals were used for a sensing device,
the control of cell adhesion is necessary. For these purposes, thefundamental property is inhibition of protein adsorption.
Poly(ethylene glycol), PEG, is a biofuctional molecule on which
adsorption of proteins is inhibited. Therefore, immobilization of
PEG to metal surface is an important event to bio-functionalize
the metal surface. No successful one-stage technique for theimmobilization of PEG to base metals has ever been developed.
In this study, PEG modified both terminals or one terminal withamine bases was immobilized onto titanium surface using
electrodeposition. PEG modified at both terminals or one terminal
with amine bases was immobilized onto titanium surface.
The bonding manner of PEG to titanium surface is significant
to design PEG-immobilized materials, while characterization
techniques for the determination of immobilization manner of
PEG are a little. PEG terminated at both terminals or one terminal
with amine bases was immobilized onto titanium surface withelectrodeposition and immersion. The electrodepsition was
carried out with -5V for 300 s. Immobilization manner of PEG wascharacterized using X-ray photoelectron spectroscopy (XPS) with
angle-resolved technique and glow discharge optical emission
spectroscopy (GD-OES). As a result, not only electrodeposition
but also immersion led to the immobilization of PEG onto titanium
surface. However, more terminated amines combined withtitanium oxide as an ionic NH-O by electrodeposition, while moreamines randomly existed as NH3+ in the PEG molecule byimmersion. Moreover, the difference of amine termination led todifferent bonding manner, U-shape in PEG terminated bothterminals and brush in PEG terminated one terminal.
Characterization with XPS and GD-OES is useful to determine
immobilization mode of PEG to solid surface [1,2].
Terminated with -NH2
H H
H O C C O

PEG without
termination

H H

H
n

H
H H
H H H HH C O C C O C C C N
nH H H H
H
H H

One terminalterminated PEG－
H2Both terminalterminated PEG

＋
OH-

O－
O2

Ti

H+
CP Ti

H H H HH H
H H H HN C C C O C C O C C C N
nH H H H
H H H H
H H

NH3+

Pt

2 mass% PEG + 0.3 mol L-1 NaCl （pH11）
pH11）+
O－ NH3
O－
+
NH3
O－
－
O
+
NH3
O－
O－
PEG

Fig. 1 PEG molecules were terminated with amine bases at one
terminal or both terminals. Amine bases dissociate and are positively
charged in aqueous solution and electrically attracted to titanium
surface with cathodic charge, and eventually PEG molecules areimmobilized.

KINKEN Research Highlights 2008 IMR

Random direction

BothBoth-PEG
Immersion
NH3+
NH3+
NH3+

O-

NH3+

O-

NH3+

NH3+

NH3+

O-

NH3+

NH3+
NH3+

+
NH3+ NH3

NH3+

O-

O-

O-

O-

-NH oriented totitanium
NHNH-O
Stable bonding

ElectroElectrodeposition

NH3+…OUnstable bonding

Immersion

OneOne-PEG
Electrodeposition

BothBoth-PEG
Electrodeposition

NH

NH

NH

O

O

O O

NH

NH3+

O

NH3+

NH

NH

NH

NH

NH

NH

NH

NH

NH

O

O

O

O

O

O

O

O

O

Ti

Ti

Ti

RandomBrudhBrudh-shape

U-shapeFig. 2 Schematic model of iimobilized manners of PEG to titanium
surface with immersion and electrodeposition. In the first layer ofimmobilized PEG formed by electrodeposition, more terminatedamines exist at the interface between PEG layer and titanium oxide
and combine with titanium oxide as N-HO bond ionic bond formed
between NH3+ and OH- while more amines randomly exist as NH3+in PEG layer formed by immersion. In the both-terminated PEG, thePEG immobilized as a U-shape because both terminals combinewith the titanium surface; in the one-terminated PEG, the PEG was
immobilized as a brush because only one terminal combines withthe surface. Over second layers, nitrogen atoms mainly exist as
NH3+, while nitrogen locates a deeper site of PEG layer.

The concentrations of hydroxyl groups located inside andon the surface oxide films of a commercially pure titanium, cpTi, a type 316L austenitic stainless steel, SS, and a cobaltchromium-molybdenum alloy, Co-Cr-Mo, were evaluated using
X-ray photoelectron spectroscopy, XPS, and a zinc-complex
substitution technique. As a result, the concentrations of thehydroxyl groups detected by the zinc-complex substitution
technique, defined as active hydroxyl groups, were much larger
than those detected by other conventional techniques. The
concentration of the active hydroxyl groups on Co-Cr-Mo wassignificantly larger than those on cp-Ti and SS. Poly(ethylene
glycol), PEG, is a biofunctional molecule that inhibits theadsorption of proteins. The immobilization of PEG to metal
surfaces by electrodeposition or immersion is an important
technique to biofunctionalize the metals. The amounts of the
PEG layer immobilized on the metals were governed by theconcentrations of the active hydroxyl groups on each surface
oxide in the case of electrodeposition; it was governed by the
relative permittivity of the surface oxide in the case of immersion.
The estimation of active hydroxyl groups on the surface oxide
film with the zinc-complex substitution technique is useful for the
elucidation of reactions between metal substrates andimmobilized molecules [3].
PEG-immobilized surface inhibited the adsorption ofproteins, therefore, this electrodeposition technique is usefuland universally biofinctionalize metal surfaces.
References[1] Y. Tanaka, H. Doi, Y. Iwasaki, S. Hiromoto, T. Yoneyama, K.
Asami, H. Imai and T. Hanawa, Mater. Sci. Eng. C27, 206 (2007).
[2] Y. Tanaka, H. Doi, E. Kobayashi, T. Yoneyama and T.
Hanawa, Mater. Trans., 48, 287 (2007).
[3] Y. Tanaka, H. Saito, Y. Tsutsumi, H. Doi, H. Imai and T.
Hanawa, Mater. Trans., 49, 805 (2008).
Contact toTakao Hanawa (Institute of Biomaterials and Bioengineering,Tokyo Medical and Dental University)e-mail: hanawa.met@tmd.ac.jp
Hisamichi Kimura (Advanced Research Center of Metallic Glasses)
e-mail: hisami@imr.tohoku.ac.jp

Highlights of Collaborated Research

55

3FTFBSDI

Transport Measurements of a Heavy Fermion Compound URu2Si2
under High Magnetic Field
Much attention has been attracted in the order parameterof elusive “hidden order” state in a heavy fermion compound
URu2Si2. We have investigated the transport properties
under high magnetic fields at low temperatures using an
ultraclean single crystal of URu2Si2. The results suggest theband dependent nature of the hidden order gap.
URu2Si2 is a well known heavy fermion compound which
has elusive “hidden order” (HO) state. The HO state realized
below Th = 17.5 K with a large entropy change. However,neutron scattering experiments reported a tiny orderedmoment of 0.02 µB/U. Despite many experimental andtheoretical efforts in the last two decades, the order parameterof the HO state is still unclarified.
We measured the transport properties in the HO state athigh magnetic fields up to 27 T and low temperatures down
to 0.5 K using an ultraclean single crystal of URu2Si2. Field
dependence of the transverse magnetoresistivity ρ and Hall
coefficient RH (= ρxy/µ0H), where ρxy is the Hall resistivity, areshown in Fig. 1 (a). Superconductivity sets in below the upper
critical field µ0Hc2 = 2.5 T. Above Hc2, the magnetoresistivity
increases with increasing the field, has a plateau-like behavior
between 18 to 22 T. The Hall coefficient has a sharp dip atµ0H* = 22 T which is lower than the reported critical field (=35 T) of the HO state. It is indicative of a phase transition
inside the HO state. Above the plateau-like structure, ρ
increases without saturation up to 27 T, reaches 1000 µΩ·cm
which is 2000 times larger than that of zero field extrapolation.
This huge magnetoresistivity comes from the cleanness ofthe sample and the compensation of electrons and holes. HT phase diagram for H // [001] is shown in Fig. 1 (b). H* (T)decreases with increasing the temperature as shown in red
circles. We also observed clear oscillation in ρ and RH due tothe Shubnikov-de Haas (SdH) effect at high fields abovearound 20 T. Three fundamental branches, δ, α and β wereobserved. Figure 2 shows the angular dependence of theSdH frequency F. F (= ħcSF/2πe) is proportional to theextremal (maximal or minimal) cross-sectional area of the
Fermi surface SF. Branches α, which is corresponding to aspherical Fermi surface, and β are in good agreement withthe previous de Haas-van Alphen (dHvA) experiments.
However, δ branch, which is clearly shows an elliptical Fermi
surface, has not been reported in the dHvA experiment while
it has relatively light cyclotron effective mass mc = 2.7 m0 for
H // [001]. Branch δ begins to appear above around 20 T
which is close to the phase transition.
From these results, we conclude that the HO gap on the
newly-found elliptical Fermi surface is suppressed at highfields(HO I phase) while it is opened at low fields(HO II phase)(see Fig. 1 (b)). In analogy with the recently proposed
multiband nature of the superconducting gap in this system[1],the present results suggest the band dependent nature of
the HO gap in URu2Si2.

(a)
Hc2
T = 0.5 K
J // [100]H // [001]H*(b)

Fig. 2. Angular dependence of the SdH frequencies forbranch δ, α and β, with schematic Fermi surfaces.

HO I
HO II
SC

H*

Fig. 1. (a) Field dependence of the Hall coefficient andtransverse magnetoresistivity for H // [001] at T = 0.5 K. (b)H-T phase diagram for H // [001].

56

Highlights of Collaborated Research

Reference[1] Y. Kasahara, T. Iwasawa, H. Shishido, T. Shibauchi, K.
Behnia, Y. Haga, T. D. Matsuda, Y. Onuki, M. Sigrist and Y.
Matsuda, Phys. Rev. Lett. 99 116402 (2007).
Contact toHiroaki Shishido (Department of Physics, Kyoto University)e-mail: shishido@scphys.kyoto-u.ac.jp
Takahiko Sasaki (Low Temperature Physics Division)e-mail: takahiko@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

Characterization of Current Transport Properties in High Tc Superconducting
RE1Ba2Cu3O7-δ Coated Conductors Targeting High Field Magnet Applications
We have succeeded to describe nonlinear current
transport properties in RE1Ba2Cu3O7-δ coated conductors
as a function of temperature and magnetic field including
very high fields more than 30 T. This allows us to design ahigh field superconducting magnet taking into accountdetailed operation conditions inside the magnet.
RE1Ba2Cu3O7-δ (RE=Rare earth elements; RE-123)
coated conductor is one of the most promising superconducting
materials especially for high field magnets because of its
superior in-field critical currents and mechanical strength.
Recently, remarkable progress has been achieved in thedevelopment of long length tapes in km-length scale withhigh current carrying capability. This demonstrates a big
potential of those materials. However, fundamental current
transport properties of those new wires were not yet fully
understood for engineering applications.
Thermal stability of superconducting magnets can be
improved drastically by using high critical temperature, Tc,
superconducting (HTS) winding because 1) heat capacity
increases several orders higher than the conventional ones
at elevated operation temperature and 2) high Tc results inessential large temperature margin. This means that thedesign concept of superconducting magnet becomes quite
different from the conventional one. Namely, we need to take
into account dissipation inside the winding, which is originatingfrom nonlinear electric field-v.s.-current density (E-J)
characteristics in the superconducting strand. However, thecurrent transport characteristics vary in a complicated mannerdepending on operation conditions such as temperature,magnetic field and applied field angle to the tape surface.
Based on a physical model describing percolative
characteristics of current driven flux flow in a disordered
medium, we have established a physical model which
describes E-J characteristics in the HTS superconductors [1,2]. Schematic diagram of the percolative flux flow is shown in
Fig. 1 (a). Such percolative flux penetration has been
confirmed indeed in coated conductors by the transport
magnetic microscopy as shown in Fig. 1 (b) [2].
Within the framework of the percolation model, we
formulate the nonlinear E-J characteristics analytically, andalso predict the properties up to very high magnetic field [1,3]. As an example, transport E-J characteristics in a Gd-123
coated conductor is shown in Fig. 2, where dots are measured
results and lines are analytical expression. External field wasapplied 30 degree from tape surface as schematically shown

(a)

(b)

Fig. 1 (a) Schematic diagram of percolative flux flow and (b)visualized flux penetration into coated conductor by transportscanning SQUID microscopy.

KINKEN Research Highlights 2008 IMR

(a)

(b)

Fig. 2 Magnetic field dependent E-J characteristics in Gd-123
coated conductor (a) 4.2K and (b) 77K. Lines are analyticalexpression within the framework of the percolation model,
whereas the dots are measurements.
in the inset of Fig. 2(a). Thickness of the superconducting
layer in the tape is 1.2µm, therefore, it can be seen that morethan several-100A of transport current can be carried by
1 cm wide tape even under very high external field more than
30 T. Properties at high temperature region is also welldescribed by the same formulation as shown in Fig. 2(b).
References[1] T. Kiss, M. Inoue, S. Egashira, T. Kuga, M. Ishimaru, M.
Takeo, T. Matsushita, Y. Iijima, K. Kakimoto, T. Saitoh, S.
Awaji, K. Watanabe, and Y. Shiohara, IEEE Trans. Appl.
Supercond. 13, 2607 (2003).
[2] T. Kiss, M. Inoue, T. Shoyama, S. Koyanagi, D. Mitsui, T.
Nakamura, K. Imamura, A. Ibi, Y. Yamada, T. Kato, T. Hirayama,
Y. Shiohara, IEEE Trans. Appl. Supercond. 17, 3211 (2007).
[3] M. Inoue, T. Kiss, D. Mitsui, T. Nakamura, T. Fujiwara, S.
Awaji, K. Watanabe, A. Ibi, S. Miyata, Y. Yamada and Y.
Shiohara, IEEE Trans. Appl. Supercond. 17, 3207 (2007).
Contact toTakanobu Kiss (Research Institute of Superconductor Scienceand Systems, Kyushu University)e-mail: kiss@sc.kyushu-u.ac.jp
Kazuo Watanabe (High Field Laboratory for Superconducting
Materials)
e-mail: kwata@imr.tohoku.ac.jp

Highlights of Collaborated Research

57

3FTFBSDI

Development of a New Type of High-Temperature Structural
Material based on Dual Multi-Phase Intermetallic Alloys
Dual multi-phase intermetallic alloys, which are composedof Ni3Al(L12) and Ni solid solution (A1) phases at hightemperature and is additionally refined by a eutectoid reaction
at low temperature, according to which the Al phase isdecomposed into Ni3Al(L12)+Ni3V(D022), was developed. The
alloys with such a novel microstructure showed extremely
attractive high-temperature properties, promising for a newtype of high-temperature structural material.
A number of Ni-based superalloys currently in use arecomposed of a two-phase microstructure consisting of γ (Nisolid solution:A1 phase) and γ’ (Ni3Al:L12 phase). A1 phase ismost densely packed structure when metal or alloy exists as a
mono-component or a solid solution. On the other hand, L12
phase is categorized in geometrically close packed (GCP)
structure. It has been believed that attractive properties of Nibased superalloys are attributed to not only the presence of γ’(L12 phase) but also fine microstructure with coherency
between Ni solid solution and γ’. Like the superalloys, themulti-phase intermetallic alloy composed of the GCP Ni3X
intermetallic phases is expected to have coherent interfaces
and fine microstructure because of their closely related
crystal structures.
The so-called dual multi-phase intermetallic alloy, which iscomposed of Ni3Al(L12) and Ni solid solution (A1) phases athigh temperature annealing and is additionally refined by aeutectoid reaction at low temperature aging, according to which
the Al phase is decomposed into the Ni3Al(L12)+ Ni3V(D022)
phases, was recently developed, based on the pseudo-binary
alloy system Ni3Al-Ni3V containing Ti and Nb [1,2]. It was
shown that the dual multi-phase intermetallic alloys exhibit
highly coherent interface structure between the constituent
phases of Ni3Al(L12) and Ni3V(D022) not only at the micronscale but also at the sub-micron scale as shown in Fig.1, andalso display high microstructural and phase stability at hightemperature [1,2].

a

b

c

d

Fig. 2 Some products or parts fabricated by the dual multiphase intermetallic alloy: (a) turbine blade for jet engine, (b)ball bearing, (c) friction stirring welding tool for high melting
point materials, and (d) bolt and nut.
The dual multi-phase intermetallic alloys showed extremely
high yield and tensile strengths with good temperature retention[3,4], and also superior oxidation resistance at high temperatureand corrosion resistance in various liquids. Over a broad
temperature range, high tensile elongation and fracture
toughness were observed. The creep test conducted at hightemperature showed the presence of high threshold stress,
and also extremely low creep rate and long creep rupture time
when compared with conventional Ni-based superalloys [3,4].
The obtained results are promising for the development of thedual multi-phase intermetallic alloy as a new-type of hightemperature structural material. Figure 2 indicates someproducts and parts actually fabricated by the dual multi-phaseintermetallic alloys. In collaboration with industries not only in
Kansai area but also in Keihin area, the development based on
the dual multi-phase intermetallic alloys is currently
conducted.
References[1] Y. Nunomura, Y. Kaneno, H. Tsuda and T. Takasugi, Acta
Materialia, 54, 851 (2006).
[2] W. Soga, Y. Kaneno and T. Takasugi, Intermetallics 14, 170(2006).
[3] S. Shibuya, Y. Kaneno, M. Yoshida and T. Takasugi, Acta
Materialia, 54, 861 (2006).
[4] W. Soga, Y. Kaneno, M. Yoshida and T. Takasugi, Mater.
Sci. Eng. A, 473, 180 (2008).

Fig. 1 Microstructure of dual multi-phase intermetallic alloy
composed of geometrically close packed (GCP) structures
Ni3Al(L12) and Ni3V(D022).

58

Highlights of Collaborated ResearchContact to
Takayuki Takasugi (Graduate School of Engineering, Osaka
Prefecture University)e-mail: takasugi@mtr.osakafu-u.ac.jp
Toyohiko Konno (Osaka Center for Industrial Materials
Research)e-mail: tjkonno@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

A Review Book “Nano-and Micromaterials” Published as a New
Volume of Advances in Materials Research from Springer
As a comprehensive scientific report of collaborated
researches performed by Yokohama National University(YNU) and Instritute for Materials Research (IMR), Tohoku
University, a new review book titled “Nano- andMicromaterials” (picture) was published in February 2008
from Springer-Verlag as the ninth volume of the Springer
series, ”Advanc es in Materials Research” [1].
A fruitful IMR Collaborated Research in Priority Area
named “Nanotechnology Research by First-Principles
Calculations” has continued for three years between Kaoru
Ohno’s group at YNU and Yoshiyuki Kawazoe’s group at
IMR. Also in YNU, an interdisciplinary research project
named “Materials Science for Nano- and Micro-Scale
Control: Creation of New Structures and Functions”(Leader: Kaoru Ohno, Subgroup Leaders: Masatoshi
Tanaka and Jun Takeda) has continued at the same time
period. This book edited by K. Ohno, M. Tanaka, and J.
Takeda (YNU), and Y. Kawazoe (IMR) collects review
articles on the recent progress in variety of subjects writtenby the members in both YNU and IMR under these research
projects. The content of this book is as follows:
Chap. 1 “General Introduction” (K. Ohno, YNU) pp.1-17.
Chap. 2 “Nanometer-Scale Structure Formation on Solid
Surfaces” (M. Tanaka, K. Shudo, S. Ohno, YNU) pp.1995.
Chap. 3 “Ultrafast Laser Spectroscopy Applicable to Nanoand Micromaterials” (J. Takeda, YNU) pp.97-119.
Chap. 4 “Defects in Anatase Titanium Dioxide” (T. Sekiya
and S. Kurita, YNU) pp.121-141.
Chap. 5 “Organic Radical 1,3,5-Trithia-2,4,6-Triazapentalenyl (TTTA) as Strong Correlations Electronic Systems:
Experiment and Theory” (J. Takeda, Y. Noguchi, S.
Ishii, and K. Ohno, YNU) pp.143-169.
Chap. 6 “Ab Initio GW Calculations Using an All-Electron
Approach” (S. Ishii, K. Ohno, and Y. Kawazoe, IMR)
pp.171-188.
Chap. 7 “First-principles Calculations Involving Two- Particle
Excited States of Atoms and Molecules Using T-Matrix
Theory” (Y. Noguchi, S. Ishii, and K. Ohno, YNU)
pp.189-217.
Chap. 8 “Green’s Function Formulation of Electronic
Transport at Nanoscale” (A. A. Farajian, O. V.
Pupysheva, B. I. Yakobson, and Y. Kawazoe, IMR)
pp.219-241.
Chap. 9 “Self-Assembled Quantum Dot Structure Composedof III-V Compound Semiconductors” (K. Mukai, IMR)
pp.243-262.
Chap. 10 “Potential-Tailored Quantum Wells for HighPerformance Optical Modulators/Switches” (T. Arakawa
and K. Tada, YNU) pp.263-274.
Chap. 11 Thermodynamic Properties of Materials Using
Lattice-Gas Models with Renormalized Potentials (R.
Sahara, H. Mizuseki, K. Ohno, and Y. Kawazoe, IMR)
pp.275-290.

KINKEN Research Highlights 2008 IMR

Chap. 12 Optically Driven Micromachines for Biochip
Application (S. Maruo, YNU) pp.291-309.
Chap. 13 Study of Complex plasma (M. Shindo and O.
Ishihara, IMR) pp.311-327.

Reference[1] K. Ohno, M. Tanaka, J. Takeda, and Y. Kawazoe, edited,“Nano- and Micromaterials”, Springer Series on Advancesin Materials Research, Vol. 9, Series Editors: Y. Kawazoe,
M. Hasegawa, A. Inoue, N. Kobayashi,a nd T. Sakurai(Springer-Verlag, Berlin, Heidelberg, 2008) pp. 1-337
Contact toKaoru Ohno (Department of Physics, Graduate School ofEngineering, Yokohama National University)e-mail: ohno@ynu.ac.jp
Yoshiyuki Kawazoe (Materials Design by Computer simulation
Division)e-mail: kawazoe@imr.edu

Highlights of Collaborated Research59