Research

Effect of Structural Parameter on Superconductivity of IronBased Superconductor FeSe1-xTex
By systematic crystal structure refinements ofcharacteristic superconductor, FeSe1-xTex (x= 0.5, 0.625,
0.75, 0.875) and FeSe by the powder neutron diffraction
technique, we studied the relationship between crystal
structre and superconductivity. In this work, we found therole of Fe-X network (X; calcogens, pnictogens, ) for themechanism of the iron-based superconductor. This work is
a result of collaboration investigations between IMR and
WPI-AIMR groups.
The discovery of superconductivity in the iron oxypnictide
family of compounds has generated considerable interest in
the mechanism of superconductivity.[1] The understandingof electronic, magnetic and structural properties of thesecompounds is the key to elucidation of this mechanism. Lee
et al. have reported that Tc becomes maximum when FeAs4lattices form a regular tetrahedron (As-Fe-As angle
α=109.47°).[2] This result indicates a relationship between
crystal structure and superconductivity. However, it has notbeen clarified whether this relationship is a universal featureof the iron superconductor because there are only a few
reports on the crystal structure analysis of materials withlower α angle (α>105°). Recently, superconductivity was
discovered in α-FeSex system. The crystal structure of FeSe
is the most simplest, and the Se-Fe-Se angle of FeSe4tetrahedra at ~105° is considerably smaller than that of theiron oxypictide family. Therefore, we attempt to determine
the crystal structural parameters of FeSe1-xTex by using theneutron diffraction technique.
We performed neutron powder diffraction experiments on
the Kinken powder diffractometer for high efficiency and highresolution measurements, HERMES, of Institute for Materials
Research, Tohoku University, installed at the JRR-3 reactor
in Japan Atomic Energy Agency, Tokai.
The relationship between Tc and X-Fe-X (X=As, P, Se

Fig. 2 Superconducting transition in FeSe0.5Te0.5 under highpressures. Inset shows the temperature dependence ofelectrical resistivity under 14GPa.

and Te) angle is shown Fig.1. The Tc of the FeSe1-xTex
system increases with increasing α, indicating that thisrelationship is universal feature and crystal structure plays animportant role in the iron superconductor.
From the structural analysis, Fe(Se,Te)4-tetrahedrons
distort greatly from the regular shape implies the improvementof Tc upon applying pressure. Thus, we have performed on
the high-pressure resistivity experiments in FeSe0.5Te0.5,
which shows the maximum Tc(=14K) under ambient pressure.
Figure 2 shows the temperature dependence of the resistivityof FeSe0.5Te0.5 at various pressure up to 7.5GPa. The onset
temperature of Tc increases rapidly from 14K to 26K applying
up to 2GPa and the metallic phase was observed at P=14GPa.
This experiment presents the first observation of a phasetransition from a superconducting to metallic phase on the
P-Tc phase diagram.[3]This work has been selected as Editor’s choice Paper of
J. Phys. Soc. Jpn. in 2009.
References[1] Y. Kamihara et al., J. Am. Chem. Soc. 130 (2008) 3296.
[2] C. H. Lee et al., J. Phys. Soc. Jpn. 77 (2008) 083704.
[3] K. Horigane et al., to be published in JPSJ
Key Words
Superconductivity, Structural Effect, Neutron Diffraction
Contact toKenji Ohoyama (Metal Physics with Quantum Beam
Spectroscopy Division)e-mail: ohoyama@imr.tohoku.ac.jp

Fig. 1 X-Fe-X (X:As, P, Se and Te) angle α dependence ofsuperconducting transition temperature Tc.

20

Future Energy / Environmental Materials

IMR KINKEN Research Highlights 2009

Research

Theoretical Research for Improved Hydrogen Storage Materials
Hydrogen is attractive as a fuel because its use createsneither air pollution nor greenhouse-gas emissions. The useof hydrogen requires an effective, safe, and stable storage
medium. However, how to store hydrogen easily and cheaply
is still a challenging problem. The current methods of storinghydrogen as compressed gas or in the liquid form do notmeet the industrial requirements because the energy densities
are much lower than that in gasoline. Moreover, there areissues of safety and cost involved in compressing hydrogen
under high pressure or liquefying it at cryogenic temperatures.
Although storage of hydrogen in solid-state materials offers
an alternative, there are no current solid-state storage
materials that meet the industry requirements. Doping oftransition and alkali elements increases the hydrogen storage
capacity of many materials. Inspired by these finds, we
explore the hydrogen storage properties of transition metal
doped boron nitride sheets, lithium doped Metal Organic
Frameworks (MOFs) and Lithium doped calix[4]arenes. All
the numerical calculations have been conducted by using the
HITACHI SR11000 K2/51 supercomputing system at theCenter for Computational Materials Science, IMR.
Theoretical calculation shows that BN sheets storage
capacity can be improved from > 1 to over 2.5 wt.% by dopingwith transition metal atoms (Ni and Rh). The site of metal
doping was found to play a major role in the capacity ofhydrogen storage. Both Ni and Rh atoms are capable toabsorb up to three hydrogen molecules chemically and themetal atom to BN sheet distance increases with the increasein the number of hydrogen molecules. Our calculations offer
explanation for the nature of bonding between the metal atomand the hydrogen molecules, which is due to the hybridizationof metal d orbital with the hydrogen s orbital [1].
Li adsorption on isoreticular MOFs with metal Fe, Cu, Co,
Ni and Zn was studied using density function theory. Li
functionalization shows a considerable structural change
associated with a volume change in isoreticular MOF-5
except for the Zn metal center. Hydrogen binding energies
on Li functionalized MOFs are seen to be in the range of 0.2
eV, which is the desired value for an ideal reversible storage
system. Each Li atom doped was found to hold three H2molecules firmly due to the charge induced dipole interaction.
The most stable position for Li atom was found was on thebenzene ring, This study has clearly shown that Li doping ispossible only in Zn-based MOF-5, which would be better
candidate to reversibly store hydrogen [2].
Density functional theory (DFT) was performed toinvestigate hydrogen storage in Li-functionalized p-tert-butyl

Fig. 2 Optimized geometries of adsorbed hydrogen molecules
on Li functionalized Zn-MOF-5 with four hydrogen molecules.

Fig. 3 Optimized geometries for Li-functionalized calixarene
with four hydrogen molecules.
calixarene (LTBC). The pure organic compound, p-tertbutylcalix[4]arene (TBC), has the capability to store hydrogenup to 0.6 wt.%. TBC was found to hold one H2 molecules
inside its cavity and the Li-functionalized calixerene was able
to hold 4 H2 molecules inside its cavity. Calculations basedon DFT show that Li-functionalized calixarene significantly
improves the average binding energy of hydrogen molecules.
The mean distance between the Li and hydrogen molecules
increases with the increase in the number of hydrogenmolecules. Ab initio molecular dynamic studies show that Li- functionalized TBC is stable up to 200K while hydrogen
incorporated material was stable up to 100K [3].
References[1] N. S. Venkataramanan, M. Khazaei, R. Sahara, H. Mizuseki,
and Y. Kawazoe, Chem. Phys. 359 (2009) 173-178.
[2] N. S. Venkataramanan, R. Sahara, H. Mizuseki, and Y.
Kawazoe, Int. J. Mol. Sci. 10 (2009) 1601-1608.
[3] N. S. Venkataramanan, R. Sahara, H. Mizuseki, and Y.
Kawazoe, J. Phys. Chem. C. 112 (2008) 19676-19679.
Key Words
Hydrogen Storage Material, First Principles Calculations,
Metal Organic Frameworks

Fig. 1 Optimized geometric structures for the hydrogen
adsorbed on the metal doped BN sheet.

KINKEN Research Highlights 2009 IMR

Contact toYoshiyuki Kawazoe (Materials
Simulation Division)e-mail: kawazoe@imr.edu

DesignbyComputer

Future Energy / Environmental Materials

21Research

Aluminum Hydride for Hydrogen Storage

Complex hydrides with high hydrogen density are being
extensively studied for developing advanced hydrogenstorage materials [1]. Aluminum trihydride (AlH3, alane) also
has high gravimetric and volumetric hydrogen densities (10mass% and 149 kgH2/m3, respectively). Equilibrium hydrogenpressure of α-AlH3 (the most stable phase among the variantsof the crystalline structure) is around 1 GPa at 298 K, although
AlH3 undergoes a simple hydrogen desorption reaction (AlH3→ Al + 3/2H2) at 370–470 K [2].
Fig. 1 shows the TG profile and SEM images during thehydrogen desorption reaction of the AlH3 particles [3]. The
TG profile shows that the hydrogen desorption reaction starts
at approximately 370 K and the amount of desorbed hydrogenis 9.6 ± 0.2 mass% (corresponding to 96% of the ideal
amount in AlH3, 10.1 mass%). From the XRD profiles, we
confirmed the existence of single-phase α-AlH3 with amolecular volume of 33.4 Å3/α-AlH3 and metallic Al with amolecular volume of 16.6 Å3/Al before and after the hydrogen
desorption reaction, respectively. The two-fold shrinkage inmolecular volume from α-AlH3 to metallic Al was highly
expected to cause a change in the morphologies of the
primary particles. The SEM images, however, indicate thatthe morphologies of the particles of size 100 nm–1 μm do notchange at all during the hydrogen desorption reaction, as
shown in Figs. 1 (a) – (c).
A TEM image of the primary AlH3 particles before thehydrogen desorption reaction is shown in Fig. 2 (a) [3].
Electron diffraction patterns suggested that the particles
before the hydrogen desorption reaction are single crystalsof α-AlH3. In addition, electron energy-loss spectroscopy and
X-ray photoelectron spectroscopy revealed that the primary
AlH3 particles are covered by a surface layer of amorphousaluminum oxide with a thickness of 3–5 nm [3, 4].
Furthermore, the TEM images taken during the hydrogen
desorption reaction (heating achieved by electron irradiation)
are shown in Figs. 2 (b) and (c). Clearly, nanoscale (~1 nm)
precipitations can be observed inside the primary AlH3
particles of size 100–200 nm; then, these precipitations
exhibit a continuous grain-growth (20–50 nm). Simultaneously,the ‘boundary space’ in the primary AlH3 particles increases
because of the shrinkage in the molecular volume from
α-AlH3 to metallic Al, as described above. The electron
diffraction pattern indicates the existence of polycrystalline
metallic Al after the hydrogen desorption reaction; it may be
noted that the polycrystals were formed from a single crystalof AlH3 during the hydrogen desorption reaction. On thesurface layer, neither the formation of cracks nor a change inthickness was observed during the hydrogen desorption
reaction. It was, therefore, found that the surface morphologyof the primary AlH3 particles was not affected by the hydrogen
desorption reaction; however, the nanostructure inside theparticles was found to change drastically because of the
precipitation and continuous grain-growth of metallic Al.
Thus, the surface layer on the primary AlH3 particles should

22

Future Energy / Environmental Materials0

TG (mass%)

Hydrogen desorption reaction of aluminum hydride,
AlH3, was investigated by in situ microscopic observations
combined with thermal and surface analyses. Before thehydrogen desorption reaction, primary AlH3 particles of size
100 nm–1 μm were thought to be covered by an oxide layer
with a thickness of less than 5 nm. The morphologies of theparticles covered by the layer did not change during thehydrogen desorption reaction.

-2

(a)

-4

(b)

(c)

-6
2 µm

-8

-10

300

350
400
Temperature (K)450Fig. 1 Thermogravimetry (TG) profile and in situ SEM imagesof the AlH3 particles [3]. (a) before the hydrogen desorption
reaction, and during the reaction at (b) 370 K and (c) 420 K.

(a) Before hydrogen desorption(b)(c) After hydrogen desorption

100 nmFig. 2 TEM images of the primary AlH3 particles (a) before,(b) during and (c) after the hydrogen desorption reaction [3].
play a dominant role against the occurrence of the hydrogen
desorption reaction (thereby retaining the ‘AlH3’ structure
inside the particles) at room temperature.
References[1] S. Orimo, Y. Nakamori, J.R. Eliseo, A. Züttel and C.M.
Jensen, Chem. Rev. 107, 4111 (2007).
[2] S. Orimo, Y. Nakamori, T. Kato, C. Brown and C.M. Jensen,
Appl. Phys. A 83, 5 (2006).
[3] K. Ikeda, S. Muto, K. Tatsumi, M. Menjo, S. Kato, M.
Bielmann, A. Züttel, C.M. Jensen and S. Orimo, Nanotechnology
20, 204004 (2009).
[4] S. Muto, K. Tatsumi, K. Ikeda and S. Orimo, J. Appl. Phys.
105, 123514 (2009).
Key Words
Aluminum Hydride, Hydrogen Storage Material, Transmission
Electron Microscope Observation
Contact to
Kazutaka Ikeda (Hydrogen Functional Materials Division)e-mail: kikeda@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2009

ResearchMultiple Superior Properties of New Au-based Bulk Glassy
Alloys with Ultralow Glass Transition Temperature
New Au-based Au-Cu-Ag-Si bulk glassy alloys (BGAs)
with high Au contents exhibit ultralow glass transition
temperature (Tg), high glass-forming ability (GFA), wide
supercooled liquid region (ΔTx), good mechanical properties,
excellent corrosion resistance, high thermodynamic andkinetic stability in supercooled liquid state, strong oxidation
resistance, and excellent thermoplastic formability[1].
The Au-based BGAs contain high Au content of ~70 at.%
(89 wt.%), which is much higher than that of 18-kirat Au. The
alloys show wide ΔTx of~36-50 °C and ultralow Tg of ~ 66
-102 °C (see Fig.1), which is the lowest in the Tg among
known BGAs except Ca-based alloys. The alloys also
possess high GFA, and the full glassy sample diameters arein the range of ~2-5 mm. The glassy alloys can be preparedby melt spinning and copper mold casting in air. Their luster
and thermal properties are the same as those of the samples
prepared in an argon atmosphere, indicating the strongoxidation resistance even at high temperature, which issimilar to gold.
Although the Au-based BGAs show ultralow Tg, they
have high stability at ambient temperatures. These BGAs
stored at room temperature (RT) for more than 4 months arestill in a fully glassy state. Their stability is also testified bytemperature-time-transformation (TTT) diagram shown in
Fig. 2. Arrhenius extrapolation with a predicted lifetime at RT(20°C) is over of 4.9x107s. It is estimated that the fully glassy
state can be kept at RT for at least 18 months.
Mechanical tests on these BGAs exhibit quite highcompressive fracture strength of ~ 643-935 MPa, Young’s
modulus of ~ 46.6-65.4 GPa and high Vicker’s hardness of ~264-353. Their corrosion resistance in 1N HCl and 1N H2SO4
solutions is similar to that of Ni-based BGAs, and much better
than those of other BGAs and SUS316L.
The Au-based BGAs have low viscosity of ~107 Pa.s in
the supercooled liquid state, making them easy to deformthermoplastically at a temperature below 100°C. Figure 3 (a)shows the glassy rods formed into U shapes by using two
pairs of tweezers in near-boiling water. A small dish shape(Fig. 3 (b)) is pressed from a glassy plate using a pair of pliers.
The pattern on the surface is copied from the engaging teethof the pliers. It has been demonstrated that the glassy pillars
with nanoscale diameters can be formed directly by imprinting.
Fig. 3 (c) shows SEM image of glassy pillars with diameters of
about 200 nm fabricated by isothermal processing of
Au60Cu15.5Ag7.5Si17 alloy on porous alumina at 92 °C for 180

Fig. 2
TTT diagrams of Au60Cu15.5Ag7.5Si17 and
Au65Cu10.5Ag7.5Si17 glassy alloys. The solid lines extrapolated
to RT (~20°C) show their onset time for crystallization at RT
are 8.1x108 and 4.9 x107 s, respectively.

Fig. 3 (a) Au60Cu15.5Si17Ag7.5 glassy rods formed into U
shapes by using two pairs of tweezersp in near-boiling water(97°C). (b) Pattern formed by a pair of pliers on the surface ofthe glassy plate (Φ5x1.5 mm2) in near-boiling water. (c) SEM
image of the nano-glassy pillars formed by imprinting on
porous alumina. The pillars with diameters of about 200 nm
are formed when the alloy is heated through its supercooled
liquid region (92°C ) for 180 s under an applied pressure of
100 MPa.
s under an applied pressure of 100 MPa. The results indicatethat the shaping, imprinting and other thermoplastics
processes can be conveniently applied to the BGAs.
The combination of the many superior properties makes
the new Au-based BGAs excellent prospects for scientificand many applications, including micro/nano-machinery,bioengineering, high density data recording media, andjewelry.
References[1] W. Zhang, H. Guo, M. W. Chen, Y. Saotome, C. L. Qin
and A. Inoue, Scripta. Mater, 61 (2009) on line.
Key Words
Metallic Glass, Glass Transition Temperature, Thermoplastic
Formability

Fig. 1 DSC curves of Au50-70Cu5.5-25.5Ag7.5Si17 glassy alloys.

KINKEN Research Highlights 2009 IMR

Contact to
Wei Zhang (Research and Development Project on Advanced
Metallic Glasses, Inorganic Materials and Joining Technology)
e-mail: wzhang@imr.tohoku.ac.jp

Future Energy / Environmental Materials

23Research

Development of the Ni-based Metallic Glassy Bipolar Plates forProton Exchange Membrane Fuel Cell (PEMFC)
The Ni65Cr15P16B4 glassy alloy was found to be an
optimal alloy for bipolar plate of PEMFC. A bipolar plate was
successfully produced by hot-pressing the glassy alloy
sheet in a supercooled liquid state. The I-V characteristics ofa single cell with the glassy bipolar plates were measured.
Proton exchange membrane (PEM) fuel cells have been
receiving considerable attention as clean energy sourcesthat may be used in the near future. Bipolar plates areconventionally made of carbon graphite. However, it cannot
be used to produce thin bipolar plates because of its
brittleness. Thinner bipolar plates will result in lighter andmore compact fuel cell stacks, saving costs and enhancing
the usefulness of fuel cells.
Recently, many research groups have tried to develop
metallic materials for bipolar plates. The corrosion-resistant
stainless steel SUS316L is recognized as a standard material
for comparison. Many different types of metallic bipolar plates
have subsequently been developed, including those basedon stainless steels or other alloys with coatings.
Metallic glassy alloys are also potential materials forbipolar plates because they have many advantageous
characteristics over crystalline alloys [1, 2]. For example,
glassy alloys exhibit higher mechanical strengths and higher
corrosion resistance. Furthermore, metallic glasses can show
viscous flow (plastic) deformation in a supercooled liquidstate at temperatures between the glass transition
temperature, Tg, and the crystallization temperature, Tx.
Recently, we found that a Ni-Cr-P-B glassy alloy could be
melt-spun even in air without any degradation in its thermalor other properties and that it had a wide supercooled liquid
region ΔTx.
In this work, we adopted the above mentioned Ni-Cr-P-B
quaternary system and optimized the alloy compositions
suitable for bipolar plate production. After the screening, theNi65Cr15P16B4 metallic glass was selected in this study as an
optimal alloy having as wider ΔTx, and lower Tg and Tx as

Fig. 2 Comparison of I-V characteristics between cells with
Ni65Cr15P16B4 glassy alloy bipolar plates and SUS316 plates.
possible. The Tg, Tx and ΔTx of the Ni65Cr15P16B4 glassy alloy
are 659 K, 702 K and 43 K, respectively. A bipolar plate wasthen produced by hot-pressing in a supercooled liquid state.
Finally the I-V characteristics were tentatively obtained forthe glassy alloy.
Figure 1 shows a bipolar plate produced by hot-pressing
the Ni65Cr15P16B4 glassy alloy sheet at 679 K (=Tg+20 K) for
300 s in a dilute Ar atmosphere in the supercooled liquidstate. The groove depth is 0.7 mm. Excellent viscous flow
deformation was obtained as clearly demonstrated in thisfigure.
Figure 2 indicates the I-V curves of single cells with
Ni65Cr15P16B4 glassy and SUS316 bipolar plates. The asprocessed glassy plates were used for these measurements
without any further surface-treatment. The SUS316 bipolar
plates were produced by machining. The same flow field
design was adopted for both bipolar plates. The metallic
glassy bipolar plates were mounted on a carbon frame andused for measurement. In the potentiodynamic corrosion
tests, the current density of the non-polished Ni-Cr-P-B
sample is smaller than that of the SUS316L alloy. This is also
reflected in the better performance of the single cell with
Ni65Cr15P16B4 glassy plates. However, the carbon bipolar
plates seem to have a better I-V performance. We are
currently conducting further study to improve the performanceof metallic glassy bipolar plates.
References[1] S. Yamaura, M. Yokoyama, H.M. Kimura and A. Inoue, J.
Phys.:Conf. Ser. 144, 012001 (2009).
[2] M. Yokoyama, S. Yamaura, H.M. Kimura and A. Inoue,
Int. J. Hydrogen Energy 33, 5678 (2008).
Key Words
Metallic Glass, Bipolar Plate, Fuel Cell

Fig. 1 Optical micrograph of the Ni65Cr15P16B4 glassy alloy
bipolar plate produced by hot-pressing in the super-cooled
liquid state.

24

Future Energy / Environmental Materials

Contact toShin-ichi Yamaura (NEDO Laboratory for Metallic Glass
Research Collaboration and Promotion)e-mail: yamaura@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2009

Research

3DAP and Positron Annihilation Study of Precipitation Behavior
in Cu-Cr-Zr Alloy
Precipitation behavior in a Cu–0.78%Cr–0.13%Zr alloy
during aging and reaging has been studied by laser-assisted
local electrode three-dimensional atom probe (Laser-LEAP)and positron annihilation spectroscopy (PAS). PAS experiments
demonstrate the formation of vacancy-like defects and theirassociation with Cr during the heat treatments.
Cu–Cr–Zr is considered as one of the candidate materials
for the heat sink of the International Thermonuclear
Experimental Reactor (ITER) divertor. Thermal aging is used
to achieve high strength by precipitation hardening andsimultaneously high thermal conductivity of Cu–Cr–Zr alloys.
However, there has not been a unanimouse conclusion on
the microstructure of Cu-Cr-Zr. To clarify the precipitation
behavior we have employed Laser-LEAP and positronannihilation spectroscopy (PAS).
The test specimens of Cu-0.78wt% Cr, 0.13wt% Zr,
0.003wt% Si, 0.008wt% Fe were solution annealed (SA) at
960°C for 3 h and then quenched, followed by prime aging at
460°C (PA) and reaging (overaging) at 600°C for 1 (RA1)and 4h (RA2), respectively [1].
Figure 1 shows atom maps around each Cr-rich precipitatein the PA and reaged states. In the precipitate of the PA state
shown in Fig. 1 (a), Zr, Fe and Si impurities are found to beslightly segregated on the surrounding region of the Cr-richcore. In the RA1, 2 state, the Cr-rich precipitates coarsened
as state above and the Zr, Si and Fe impurities furthermore
segregated to form enriched layer regions as shown in Fig. 1(b, c) ; the Cr-rich precipitates are covered by enriched
regions.
Positron lifetime spectra of the alloy after the heat
treatment were well described single lifetime components [2];these lifetimes are longer than the positron lifetime of bulk Cu(110 ps) but shorter than that of monovacancies in Cu(173 ps). Thus, they are given by averaged positron lifetimes
between the bulk and vacancy-like defects. Fig. 2 (a) shows
the ratio curves of the coincidence Doppler broadening (CDB)
spectra, normalized to that of pure Cu, of the heat treated
Cu–Cr–Zr alloy together with those of pure Cr and Zr as
references. Furthermore, that of neutron-irradiated Cu (C-80T)
is also shown; its average positron lifetime is about 176 ps
which exhibits that almost of all the positrons are trapped at
irradiation-induced monovacancies. Small increases at thelower momentum regions accompanied sluggish decreases
at higher momentum regions are seen for the curves of the
SA and PA states. However, these changes in the ratio curves
are markedly enhanced in the curves for the reaged states of
Ra1 and RA2, which means the introduction of ample
vacancy-like defects by reaging and is consistent with positron
lifetime observation stated above. To see the origin of thevalley, the CDB spectrum of the RA1 is normalized to neutron-

Fig. 2 Ratio curves of the CDB spectra for the Cu–Cr–Zr
alloy. CDB spectra are normalized to defect free pure Cu (a)and neutron-irradiated pure Cu (C-80T) (b).
irradiated pure Cu, together with that of pure Cr and Zr asshown in Fig. 2 (b). A broad valley centered around
20 × 10−3m0c is observed. The ratio curve multiplied by afactor of 6 is found to be very close to that of pure Cr. Thisshows that about 1/6 (15%) of positrons annihilate with Cr
electrons and the remaining 5/6 (85%) of positrons annihilate
with Cu electrons. Furthermore, the fraction of positrons whichannihilate with Cr electron, ICr, is estimated for the RA1 andRA2 states, using the fitting of high-momentum part of the
CDB ratio curves: 15% for the RA1 and 17% for RA2,
respectively. The above positron lifetime and CDB experiments
demonstrate that more than 80% of positrons are trapped atvacancy-like defects induced by reaging and about 15% ofthem annihilate with surrounding Cr electrons in the reaged
states. This suggests that the vacancy-like defects are
induced in the interface regions between the Cr-richprecipitates and the Cu matrix in the reaged states.
References[1] M. Hatakeyama, T. Toyama, Y. Nagai, M. Hasegawa, M.
Eldrup and B.N. Singh, Mater. Trans. 49, 518 (2008).
[2] M. Hatakeyama, T. Toyama, J. Yang, Y. Nagai, M.
Hasegawa, T. Ohkubo, M. Eldrup and B.N. Singh, J. Nucl.
Mater. 386-388, 852 (2002).
Key Words
Future Energy/Environmental Materials, Cu-Cr-Zr, Precipitate

Fig. 1 Atom maps of the solute and impurity distributions in
the Cu–Cr–Zr alloy with different heat treatments: (a) PA, (b)RA1 and (c) RA2.

KINKEN Research Highlights 2009 IMR

Contact toMasahiko Hatakeyama (International Research Center for
Nuclear Material Sciense)e-mail: hatake@imr.tohoku.ac.jp

Future Energy / Environmental Materials

25Research

Variety of the Flux Pinning State Diagrams for Various c-axis
Correlated Pinning Centers
In order to understand the flux pinning mechanism withthe different variety of the c-axis-correlated disorder, theangular dependence of critical current density Jc wasmeasured in detail for an Y123 film with BaZrO3 nano-rods,
a Sm123 film fabricated by the low temperature growth(LTG) technique and an Y123 film introduced the columnar
defects by the heavy ion irradiation parallel to the c-axis.
We found that the flux pinning state is similar in low fields
below the matching field, but it shows different behaviors inhigh fields above the matching field for the variety of thec-axis correlated pinning. It is considered that the interstitialvortices interacted with the c-axis-correlated disorders
through the vortex interaction may play an important role forthe appearance of the c-axis-correlated pinning behavior ina high-field region above the matching field.
Practical high temperature superconducting (HTS) tapes
and wires are developing under the national projects in theworld. However one of the serious issues is the large
anisotropy of the HTS materials. To overcome the large
anisotropy of the superconducting parameters such as a
critical current density Jc and an irreversibility field Bi, c-axis
correlated pinning centers, i.e. columnar shaped precipitates
or disorders along c-axis of the superconducting matrix, areintroduced. Typical c-axis correlated pinning centers are thefission tracks by the heavy ion irradiation, dislocation andcolumnar-shaped precipitates along c-axis. It is consideredthat the flux pinning behaviors, which determine the Jc and Bi
properties, are related to those density, size and growth
directions. In order to clarify the detailed flux pinning behavior
under the various c-axis correlated pinning centers, wemeasured the transport Jc in wide range of a temperature, amagnetic field and a magnetic field angle carefully anddiscuss the flux pinning behavior.
We selected three different RE1+xBa2-xCu3Oy (RE: rareearth) films with different c-axis correlated pinning centerssuch as the nano-scale BaZrO3 columnar-shaped precipitates,
i.e. nanorods (BZO-Y123), columnar defects by the irradiationof the Au-ions with 240 MeV (CD-Y123), and dislocations
along c-axis (LTG-Sm123). Typically, the peak appears on
the angular dependence of Jc for θ = 0˚ (B//c), if c-axis
correlated pinning works effectively. Hence, the c-axis
correlated pinning state can be determined by the existenceof the Jc peak for B//c. Fig. 1 is the flux pinning state diagramsfor three different films, which were determined by the angulardependence of Jc. In the orange-hatched region, the Jc (θ)
peak for B//c can be observed, and then the c-axis correlated
pinning behavior is dominant. On the other hand, in the yellow
region, the random pinning is dominant because the Jc peak
attributed to the c-axis correlated pinning vanishes. We cansee the c-axis correlated pinning dominant region in low
fields for all samples. However the pinning behaviors in ahigh field region are different among a sort of the c-axis
correlated pinning [1]. It is suggested that the size anddirection of the c-axis correlated pinning centers affect theflux pinning state.
According to the theory, the strong Bose glass state
exists below the matching field [2]. Here, the matching field isa corresponding magnetic field estimated by the number ofc-axis correlated pinning centers and a flux quantum.
However it changes to the weak Bose glass state due to theexistence of the interstitial vortices in between vortices pinned
directly by the c-axis correlated disorders [3]. In this case,those interstitial vortices should interact with either the c-axis
correlated disorders through the vortex interaction or the
random pinning directly. Therefore, we have to consider theelastic behavior of the vortices under the competitive state
between the random and the correlated pinning, in order tounderstand the complex vortex pinning state diagram.
References[1] M. Namba, S. Awaji, K. Watanabe, T. Nojima, S. Okayasu,
M. Miura, Y. Ichino, Y. Yoshida, Y. Takai, T. Horide, P. Mele,K. Matsumoto, Appl. Phys. Exp. 1 (2008) 031703.
[2] D. R. Nelson and V. M. Vinokur: Phys. Rev. B 48 (1993)
13060.
[3] V. M. Pan and A. V. Pan: Low Temp. Phys. 27 (2001) 732.
Key Words
Superconductor, Critical Current Density, High Magnetic Field
Contact toSatoshi Awaji (High Field Laboratory for Superconducting
Materials)
e-mail: awaji@imr.tohoku.ac.jp

Fig. 1 Schematic flux pinning state diagrams at B//c for (a) BZO-Y123, (b) LTG-Sm123, and (c) CD-Y123. B*, B’ are thecharacteristic magnetic fields at which Jc peak for B//c vanishes and appears with increasing a magnetic field, respectively. Bi is
the irreversibility field, Tk is an onset temperature of the pinned flux liquid state.

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