Research

Josephson Current Induced by Ferromagnetic Resonance
We proposed the Josephson effect induced by spinwaves in a Ferromagnetic Josephson junction, in which two
superconductors are separated by a ferromagnetic thinlayer[1].
The
phenomenological
model
of
the
superconducting phase dynamics is extended to include
the spin wave excitation induced by ferromagnetic
resonance (FMR). The current-voltage characteristics show
step structures if the microwave frequency coincides with
the FMR. Our results show a new route to observe the spinwave excitation using the Josephson effect.
The Josephson effect is characterized by the zero voltage
current between two superconductors (SC’s) separated by athin insulating barrier as shown in Fig. 1. This effect is amacroscopic quantum phenomenon induced by phasecoherence between two superconductors. When a finite
voltage-drop (V) appears in the junction, the phase difference(θ) evolves with time. As a result, an alternating current flows
due to the time dependence of θ. When an alternating electricfield by microwave irradiation is applied to the junction, thecurrent-voltage (I-V) characteristics become like as astaircase, that is called Shapiro step.
The Josephson effect in the ferromagnetic Josephson
junction is of considerable interest in recent years. One of themost interesting phenomena is the oscillation of the criticalcurrent with the thickness of ferromagnetic layer. The originof oscillation is similar to that of Flude-Ferrel-LarkinOvchinnikov state, which is due to the exchange splitting of
the conduction band in the ferromagnetic metal (FM).
Most of studies on the ferromagnetic Josephson junction
have been focused on the effect of the exchange splitting on
Cooper pairs penetrating into the FM. On the other hand,ferromagnetic materials possess dynamical properties suchas spin wave, which is excited by using ferromagnetic
resonance (FMR). Therefore, we can think that the interaction
between Cooper pairs and spin waves in the FM plays animportant role on the transport properties. For instance, theAndreev reflection assisted by magnons has been discussed
in SC/FM junction. Up to now, however, the spin dynamics inthe ferromagnetic Josephson junction has not received muchattention.
We studied the effect of the spin dynamics on theJosephson effect in the ferromagnetic Josephson junction.
The phenomenological model is extended to include theeffect of spin wave excitations, which is a precessional motionof magnetization in the FM. As shown in Fig. 2, we found that
the I-V characteristic of this model exhibits step structures
and the dc Josephson current is induced by the spin wave
excitation due to the FMR. In particular, the step structuresappear whenever the following relation is satisfied, 2eV/ћ =2nΩ, where n is integer and Ω is the FMR frequency. We caninterpret the above equation in an analogous way to the
Shapiro step. In the present situation, a uniform mode of spinwave is excited in the FM by the FMR due to the microwave
irradiation, and Ω correspond to the energy of quantized spinwaves (magnons). If an s-wave Cooper pair penetrating into

2

Nano / Micro Structure Controlled Materials

Fig. 1 A microwave radiation hac onto the ferromagnetic Josephson
junction.

Fig. 2 Current-voltage characteristics are shown for the resonant(solid line) and non-resonant (broken line) frequencies.

the FM absorbs an odd number of magnons, the Josephson
coupling vanishes between the s-wave SCs, since theformation of spin triplet state results in no Josephson coupling.
On the other hand, if a number of magnons are even, thes-wave Cooper pair can go into the other superconductor.
Hence, the Josephson current can flow. Our results provide
a new route to observe the spin waves using the Josephsoncurrent in SC/FM/SC junctions.
References[1] S. Hikino, S. Takahashi, M. Mori, and S. Maekawa: J.
Phys. Soc. Jpn. 77, 053707 (2008).
Key Words
Josephson Effect, Ferromagnetic Resonance, Shapiro Step
Contact to
Sadamichi Maekawa (Theory of Solid State Physics Division)e-mail: maekawa@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2009

Research

Johansson-type Point-focusing Monochromator Crystal
Realized by Hot Plastic Deformation of Ge Single-crystal Wafer
(b)

0.0o

-0.5

C

L1

L2

Fo
cu
si

O

F

le

L0

P
S

Present
crystalrc

Qci

333 rocking curve
0.5o

ng

(a)

∆ω

The realization of a point-focusing monochromator toincrease the intensity of the X-ray beam from a conventional
X-ray generator without having to broaden it at the focus haslong been one of the most important aims in X-ray
instrumentation. We realize such a monochromator withmoderate angular resolution and high photon flux by using
the hot plastic deformation of a Ge single-crystal wafer.

X

S

W

F

X-ray sourceoR direction (WPQ)

L1

Focal point

L2

2R direction (SPF)

In 1933, Johansson proposed an advanced design of acurved crystal with its surface exactly on the focusing circle.
Generally, the deformation of perfect crystals into an arbitrary
shape for compact optics without deteriorating the crystal
quality is not possible. Therefore, previous work has mostly
concentrated either on small and precise deformation without
introducing lattice defects or on large deformation withrelatively broad diffraction peaks.
Recently, we developed a method of achieving the hotplastic deformation of semiconducting wafers of Si and Ge
into arbitrary shapes with high accuracy [1] (see Fig.1a)
without severely deteriorating the crystal quality [2]. Here, weapply this technique to realize a Johansson-type pointfocusing monochromator with both large crystal deformationand high crystallographic quality. Figure 1b shows thegeometric requirements for Johann- and Johansson-type
monochromators. In the Johann model, X-rays diffracted
from different points along the bent surface do not converge
on the same point on the focusing circle. On the other hand,in the Johansson model, the lattice plane should be bent to aradius of 2R and the crystal surface should coincide with thefocusing circle of radius R, which leads to the perfect focusingof the diffracted X-rays. As shown in Fig.1c, if the requirements
are satisfied in a cylinder, the diffracted X-ray beam is focused
10 mm

(a)
10 mm

Hemispherical

10 mm

Wavy

Cylindrical Hemispherical

Wafer

Prepolished wafer

Bent with 2R

Bent with 2R

Prepolish (2R)

Lattice plane:2R

(b)

R

X-Ray
sourceX-Raysource

Focus

:R

ace

Surf

2R

Focusing circle

Focusing circle

Johan model

Johansson model

Crystal surface: Cylindrical R

Lattice plane: 2R

X-ray source

(c)

Q’

Line focus

Crystal surface: Spherical R

PS
X-ray source

Lattice plane: R

W’

-10

-5

0

5

10

15

X / mm

10 mm

10 mm

Fig. 2 (a) Curvature of the (111) plane observed by rocking-curvescans of the plastically deformed Ge (111) crystal wafer, and (b)focusing of Cu Kα1 X-rays using the present crystal [3].

to a line. A point-focusing geometry also requires a curvatureradius of R for the lattice plane in the Q’PW’ direction,
perpendicular to the SPF direction. Therefore, the crystal
surface should be spherical with a radius of R, and thecurvature radii of the lattice planes should be R and 2R andorthogonal to each other.
The Ge (111) wafer was polished to the shape of acylinder with a curvature radius of 1200 mm prior to
deformation to provide an appropriate offset angle between
the surface and the lattice plane. It was then deformed
plastically into a hemispherical shape with a radius of 600
mm at a temperature close to the melting point. The curvatureof the (111) plane was evaluated by the peak shift of the 333
diffraction in an ω scan using a channel-cut incident
monochromator. Figure 2a shows that the ratio of the twoslopes of the peak shift agrees with Johansson’s condition
for lattice plane curvature. Focusing images using Cu Kα1
radiation generated as a homogeneous divergent beam areshown in Fig. 2b. The source (S) -to-crystal (C) and crystal(C) -to-focus (F) distances were 21/2R = 850 mm for thesymmetric 333 diffraction of the crystal with R = 600 mm.
Diffraction images were obtained 60 mm from the crystal (L1)and 850 mm from the crystal (L2, at the focal point). The
image at L1 shows that the diffraction condition is satisfied
simultaneously over the entire illuminated area. The patternconverges to a small circular spot of about 1.7 mm diameter at
the focal point L2, demonstrating that a point focus is realized.
We are now attempting to fabricate this type of monochromatorwith much shorter radii to realize more compact systems withmuch higher photon flux, and we have already obtained
promising results. We believe that the proposed monochromators
will be of significant use to X-ray users all over the world.
References[1] K. Nakajima, K. Fujiwara, W. Pan and H. Okudai, Nature
Materials 4, 47-49 (2005).
[2] H. Okuda, K. Nakajima, K. Fujiwara and S. Ochiai, J. Appl.
Cryst. 39, 443-445 (2006).
[3] H. Okuda, K. Nakajima, K. Fujiwara, K. Morishita and S.
Ochiai, J. Appl. Cryst. 41, 798-799 (2008).

Lattice plane: 2RF

Point focus

Fig. 1 (a) Si single-crystal wafers deformed by hot plastic
deformation, (b) geometric requirements for Johann- andJohansson-type monochromators, and (c) schematic
representation of Johansson-type line- and point-focusing
geometries.

KINKEN Research Highlights 2009 IMR-15

Key Words
X-ray-point-focusing, Single-crystal, Hot-plastic-deformation
Contact toKazuo Nakajima (Crystal Physics Division)e-mail: nakasisc@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

3Research

Electron Spectromicroscopy of Microstructures
Capable of Six-Dimensional Characterization
The integration of electron spectroscopy into microscopy
in the low energy region still remains a challenge due to thetechnical difficulty of achieving both high spatial and energy
resolution with slow electrons. We demonstrate combined
electron spectroscopy and microscopy covering the real,
reciprocal and energy spaces efficiently, enabled by a low
energy electron microscope (LEEM) equipped with a simple
energy filter.
Combined structural analysis in real (microscopy) andreciprocal (diffraction) spaces by LEEM, also capable of realtime observations, has served as a useful tool for studies of
surface processes, especially for complex systems likeorganic materials [1]. Further integration of spectroscopy into
microscopy (spectromicroscopy) will provide an even morecomprehensive level of understanding of structure andproperty in nanoscale. Using a newly-developed energy
filtering method for LEEM, efficient spectromicroscopy
experiments have been carried out on Ag microstructure,
where distributions of surface plasmon [2] and band structure[3] have been successfully resolved in two-dimensional (2D)
real, 2D reciprocal and energy spaces.
The principle of the energy filtering, which is achieved
just by adding a filter entrance slit to the LEEM, is summarized
in Fig. 1. The slit position can be precisely controlled by a
motor-driven micrometer and scanned across the full electron
angular distribution in the backfocal plane of the objective
lens. The selected electrons with different energies are
dispersed by magnetic prism array to form a spectrum imageat its exit. By obtaining such an E-k spectrum for each slit
position, the energy and angular distribution of the incoming
electrons can be measured over the full surface Brillouinzone (SBZ).
Figures 2 (a) and 2 (b) show He-I photoelectron emission
micrographs (PEEM) of a complex Ag island selecting
secondary electrons and d-band photoelectrons, respectively,
at the conjugate backfocal plane. While the missing final
states along the surface normal around the vacuum level
result in darker imaging of the Ag (001) domain in Fig. 2 (a),the d-band signal is uniformly distributed over the Ag islandin Fig. 2 (b). This indicates that they both reflect distribution

Fig. 2 Energy filtered PEEM images of a mixed Ag (111) - Ag(001) island with (a) secondary electrons and (b) d-band
photoelectrons, and reconstructed energy slice images ofphotoelectrons at the strongest resonance distributed in the
2D reciprocal space for (c) Ag (111) and (d) Ag (001) domains.
The images are symmetrised following observed symmetries
by overlapping rotated images. White lines represent SBZs.
of the density of states on the Ag island, but at different
energies [3]. Distribution of the electron energy loss signal
from the Ag plasmon is also imaged on Ag nanoislands withspatial resolution less than 35 nm [2].
By selecting any observation area of 3.5 μm in diameter, a
set of parabolic spectra covering the whole reciprocal spacecan be obtained. Photoelectron maps of the strongest d-band
resonance are reconstructed for selected Ag (111) and Ag(001) microdomains, as shown in Figs. 2 (c) and 2 (d). They
exhibit distinct 3- and 4-fold symmetries, exactly reflecting thesurface symmetries of these domains [3].
The energy-filtered LEEM technique successfully
visualizes electronic properties of micro- and nanostructured
materials in 5D spaces. With its capability of time-resolved
analysis, which adds another dimension (6D in total), it willserve as a powerful tool for structural and property control ofsuch systems.
References[1] J. T. Sadowski, G. Sazaki, S. Nishikata, A. Al-Mahboob,
Y. Fujikawa, K. Nakajima, R. M. Tromp and T. Sakurai, Phys.
Rev. Lett. 98, 046104 (2007).
[2] Y. Fujikawa, T. Sakurai and R. M. Tromp, Phys. Rev. Lett.
100, 126803 (2008).
[3] Y. Fujikawa, T. Sakurai and R. M. Tromp, Phys. Rev. B79,
121401 (R) (2009). (Editors’ Suggestion)

Fig. 1 Schematics of the LEEM optics equipped with a filter
entrance slit in the backfocal plane. Incoming electrons withdifferent energies (red, green and blue) have corresponding
Ewald spheres (shaded disks) in the plane. The position
controlled entrance slit selects a slice at a given ky, which isdispersed by the prism array to form a parabolic energy
spectrum image in the conjugate backfocal plane.

4

Nano / Micro Structure Controlled Materials

Key Words
LEEM, PEEM, Spectromicroscopy
Contact toYasunori Fujikawa (Surface and Interface Research Division)e-mail: fujika-0@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2009

Research

Optical Properties of Dislocations in Wurtzite ZnO Bulk SingleCrystals Introduced at Elevated Temperatures
PL intensity I (arb. unit)

Defects inducing long-range strain fields, as dislocations, areof considerable importance in semiconductor electronics, since
they affect the electronic properties of semiconductor devices.
Recently, ZnO is expected for UV light-emitting devices and is
supposed that dislocations are more easily introduced than GaN.
However, the optical properties of dislocations in ZnO have notbeen fully elucidated.

(a)

102

deformed
1120

10-1

400nm

R

1.5

2.5

Photon energy (eV)KINKEN Research Highlights 2009 IMR

I (arb. unit)

72meV

(b)
10-1

DAP d2
TES
d1

10-2
10-3
3.1

Photon energy (eV)

FX

(c)

D02XA

72meV

2.8

8
4
0
3.5

annealed

10-4

I (arb. unit)

Despite considerable success in optimizing the growth
conditions and structural quality, ZnO epitaxial layers still contain ahigh density of defects that influence their optoelectronic properties.
Most characteristic defects, investigated by transmission electron
microscopy (TEM), are high-density (typically 109-1011 cm-2)
dislocations passing through entire layers. In ZnO, dislocations are
introduced with lower stress and are much mobile, in comparison
with in GaN [1]. Since dislocations can influence the device
performance through non-radiative recombination, etc., theknowledge of the influence of dislocations is required for thepractical use of this material. Even though ZnO has the same
crystal structure as GaN, however, the influence of dislocations in
ZnO has not been fully elucidated. We have examined the opticalproperties of dislocations in ZnO bulk single-crystals in which anarbitrary number of dislocations were newly introduced by plastic
deformation, at temperatures comparable to the typical temperatures
for the fabrication of ZnO-based devices, by photoluminescence(PL) spectroscopy and TEM [2, 3].
Wurtzite ZnO bulk single-crystals were compressively deformed
at elevated temperatures 923 - 1123 K to introduce an arbitrary
number (109-1010 cm-2) of fresh dislocations with the Burgers vectorof a/3<1120> (the inset in Fig. 1 (a)). Deformed specimens showed
excitonic PL light emission with photon energies of 3.100 and 3.345
eV, as well as their LO phonon replicas (with the separation between
the nearest-neighbor emission lines of 72 meV) at 11 K (Figs. 1 (a)and (b)). The PL intensity for the 3.100 and 3.345 eV emission
bands increased monotonously with increasing deformation
temperature (Fig. 1 (c)). Those emission bands were, therefore,
formed as a result of the introduction of dislocations at the elevatedtemperatures. On the other hand, the introduction of the dislocations
did not influence the intensities of the other emission bands such as
the free-excitonic (FX) emission band (Figs. 1 (a) and (c)). This
characteristic of ZnO may be an advantage over GaN, since GaN
exhibits a phenomenon that all intensities decrease with introductionof dislocations [4].
The PL intensity for the 3.100 and 3.345 eV emission bands
decreased via two thermal quenching processes (Fig. 2). It was
shown that, the quenching process arose at lower temperatures(at around 30 K) was due to the transition from an excitonic
recombination at a localized level to a recombination between thelevel and a band (a bound-to-free transition), and that arose at
higher temperatures was due to a thermal escape of carriers fromthe localized level to a band. The activation energy for the latterquenching process of the 3.100 eV emission band (ΔE3.100) or thatof the 3.345 eV emission band (ΔE3.345), which corresponded tothe depth of the localized energy level associated with the 3.100 or
3.345 eV emission band, was estimated to be 0.3 or 0.05 eV,
respectively. Actually, the difference of the 0-phonon peak energies
for the FX emission band and the 3.100 eV one corresponded to
ΔE3.100; i.e., 3.378 eV - 3.100 eV = 0.278 eV at T = 11 K. Also, thedifference for the FX emission band and the 3.345 eV one wasclose to ΔE3.345; 3.378 eV - 3.345 eV = 0.033 eV at 11 K.
The Huang-Rhys factor for the 3.100 eV emission band and thatfor the 3.345 eV one were, respectively, estimated to be 1.4 and 0.2.
They differed from the factors for the other emission bands, including
the factor for the 3.2108 eV emission band (~0.5) induced by

g=2201100

R
8 10-1
4

10-2

0
3.4

・

・

3.345
+

+

3.100
・
+

・
+

109

・
+
・
+

・
+
・
+

1010

Dislocation density (cm-2)

Fig. 1 The logarithmic PL intensity versus the photon energy for a
specimen; deformed at 923 K (the red curve) or annealed at 923 K
without stress (the blue curve). The ratio of those PL intensities,
R=Ideformed/Iundeformed, is also shown (the green curve). A part of (a)is magnified in (b). The inset in (a) shows a TEM image of
dislocations in the deformed specimen. (c) The PL intensity at thepeak energy of an LO phonon line as a function of dislocation
density. The open marks, the marks with dot, those with “+”, andthose with “-” show the data for the 0-, 1-, 2-, and 3-LO phononlines, respectively.
I (arb. unit)(a) T (K)11,15, 20
30
102
40
60
101 3.100

80
100
120
150
200

FX I (arb. unit)

102 ・・・

・+++

101 ・・・

FX
3.345

+・

・
・
+++ ・ +++ +・ ・ ・+ ・
・ +
+
+ ・ +

100
2.8

(b)

3.345

3.1

Photon energy (eV)

100

+

3.4 0

100

・
+

3.100

T (K)

・
+

200

Fig. 2 (a) PL spectra for a deformed specimen (the deformed
temperature; 923 K) obtained at various temperatures T. (b)T-dependence of the peak intensity of an LO phonon line. The
meaning of the marks is the same as in Fig. 1 (c)deformation at room temperature but were not induced at elevatedtemperatures. Also, no obvious polarization was observed for theemission bands. From those results, it was suggested that point defectcomplexes involving dislocations are the candidates for the origin of
the localized energy levels associated with the emission bands.
References[1] I. Yonenaga, H. Koizumi, Y. Ohno, and T. Taishi, J. Appl. Phys.
103, 093502 (2008).
[2] Y. Ohno, H. Koizumi, T. Taishi, I. Yonenaga, K. Fujii, H. Goto,
and T. Yao, Appl. Phys. Lett. 92, 011922 (2008).
[3] Y. Ohno, H. Koizumi, T. Taishi, I. Yonenaga, K. Fujii, H. Goto,
and T. Yao, J. Appl. Phys. 104, 073515 (2008).
[4] I. Yonenaga, H. Makino, S. Itoh, T. Goto, and T. Yao, J. Electron.
Mater. 35, 717 (2006).
Key Words
Dislocations, ZnO, Electronic Structures
Contact to
Yutaka Ohno (Physics of Crystal Defects Division)e-mail: yutakaohno@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

5Research

Nano Precipitation in Fe-C-M Martensitic Steels During Nitriding
Nitriding is one of the most important surface treatments used
in industry to increase surface hardness and improve the fatigue
strength. Compared with other surface hardening treatments
accompanying quenching (carburizing, induction heating and soon), this process, which utilizes precipitation hardening by finealloy nitrides, has several advantages as listed in the following;(1) Smaller strain accompanying heat treatment(2) Higher hardness achievable(3) Excellent heat resistanceand thus, attracts more attention recently in needs of much higher
precision in size scale of structural parts after heat treatment.
Here we report our recent study on nano-sized precipitation ofalloy nitrides in Fe-C-M (M: substitutional alloying element)
martensitic steels during plasma nitriding [1]Fig. 1 shows hardness profiles of Fe-0.6C, Fe-0.6C-1Al and Fe0.6C-1Cr alloys nitrided at 823K for 57.6ks. In the plain carbon steel,surface hardness hardly changes during nitriding but largely
increases in the alloyed steels. It reaches 650HV in the Fe-0.6C-1Al
alloy and 550HV in the Fe-0.6C-1Cr alloy
Fig. 2 (a) shows SEM micrographs of the nitrided layer (depth ~200μm) in Fe-0.6C-1Cr alloy. Cementite precipitates at various
boundaries in lath martensite structure (lath, block and packet
boundaries and prior austenite grain boundaries) in the nitrided
layer as well as in the unnitrided region. The size of cementite isfiner in the alloyed steels. In the nitrided layer of Fe-C alloy, plateshape iron nitrides (Fe4N or Fe16N2) are formed during furnace
cooling in addition to cementite. In the Fe-0.6C-1Al and Fe-0.6C-1Cr
alloys, much finer alloy nitrides are formed by nitriding
Fig. 2 (b) shows the bright field image and the corresponding
selected area diffraction pattern of the surface region (~40μm) of
the Fe-0.6C-1Cr alloy nitrided at 823K for 57.6ks. Much finer plateshape precipitates lying on {100} planes of the matrix are formed
uniformly. Their length and thickness are is about 20nm andseveral atomic layers, respectively. High resolution electron
microscopy clarified that those are B1-type CrN. In the Fe-0.6C1Al alloy, on the other hand, platelet AlN of which length andthickness are about 100nm and a few nm, respectively, is lying on{100} planes of the matrix in martensite laths (Fig. 2 (c)). Both ofthem have a Baker-Nutting orientation relationship ([001]α//[001]
MN, (100)α//(110)MN) with respect to the ferrite matrix. These

HV0.1 (kgf/mm2)

700Fe-0.6C
Fe-0.6C-1Al
Fe-0.6C-1Cr

600
500
400
300
200

0 100 200 300 400 500 600 700 800
Distance from the surface (µm)

Fig. 1 Hardness profiles of Fe-0.6C, Fe-0.6C-1Al and Fe-0.6C-1Cr
alloys nitrided at 823K for 57.6ks.
Alloy
Fe-1Al (ferrite)
Fe-0.6C-1Al (martensite)
Fe-1Cr (ferrite)
Fe-0.6C-1Cr (martensite)

Increase of surface
hardness ΔHV
10
400
520
250

Table. 1 Increase of surface hardness in Fe-M ferritic alloys andFe-C-M martensitic alloys (nitrided at 823K for 57.6ks).

6Nano / Micro Structure Controlled Materials

Fig. 2 (a) SEM micrograph of Fe-0.6C-1Cr nitrided at 823K for
57.6ks. (b) TEM bright filed image and corresponding selected area
diffraction pattern of the surface region (depth~40μm) of (b) Fe0.6C-1Cr and (c) Fe-0.6C-1Al nitrided at 823K for 57.6ks. Incident
beam is parallel to [001]αprecipitates are the same as reported in nitrided Fe-M binary alloys[2]. From these results, increases in hardness in Fe-0.6C-1Al andFe-0.6C-1Cr alloys are attributed to precipitation of plate-shape
AlN and plate-shape CrN, respectively.
Increases of surface hardness in martensite and ferrite nitrided
for 57.6ks are compared in Table 2. In Al-added alloys, hardness
increase in martensite is larger than in ferrite. Since AlN nucleatesat dislocations, it is concluded that high density of dislocations inlath martensite promotes the precipitation of AlN. On the other
hand, in Cr-added alloys hardness increase in martensite is smaller
than in ferrite. This is attributed to decrease in volume fraction offine CrN precipitated in the lath. As-quenched martensite istempered during heating before CrN precipitation starts. In thisperiod, Cr is enriched into the cementite whereas Cr concentrationin the matrix decreases. Since the Cr-rich cementite is converted
to coarse massive CrN during prolonged nitriding, precipitation offine plate-shape CrN within the lath might decreases. This resultsin lower surface hardness in martensite than in ferrite of the Cradded alloys.
References[1] Y. Tomio, S. Kitsuya, G. Miyamoto and T. Furuhara: Proc. the17th International Federation for Heat Treatment and Surface
Engineering (IFHTSE) Congress, 2009, in press.
[2] Y. Tomio, G. Miyamoto, T. Furuhara and T. Maki: Proc. the 3rd
Int. Conf. on Advanced Structural Steels (2006) 460-464.
Key Words
Steel, Precipitation, Surface Heat Treatment
Contact to
Tadashi Furuhara (Microstructure Design of Structural Metallic
Materials Division)e-mail: furuhara@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2009

Research

Paving the Way to High-Quality Indium Nitride-The Effects of Reactor Geometry and PressureIndium nitride (InN), with band gap energy of about 0.7
eV, is a promising material for fabricating devices such as
laser diodes for high-performance optical communications
systems and field effect transistors. However, it is very
difficult to grow InN due to the extremely high solid-vapor
phase equilibrium pressure of nitrogen. Here, we design
some new reactors to meet this challenge and confirm thoseeffects in InN growth.
Most of the industrial semiconductor devices have beengrown by metalorganic vapor phase epitaxy (MOVPE) since it
provides large scale productions, precise composition
controls, and flexible layer designs by supplying variable MOsources. The geometry of MOVPE reactors includes verticaland horizontal types. The reactor pressure generally covers
from low pressure (LP) to atmospheric pressure (AP). The LP
MOVPE reactors are widely adopted for reducing the thermal
convection effect, and have led to the great commercial
successes in InGaAsP to InGaAlP-based devices. However,
the growth of the attractive InN film is still a tough challenge
for the LP and AP MOVPE due to the extremely high solidvapor phase equilibrium pressure of nitrogen, which limits thegrowth temperature approximately below 600°C, hindering
the growth of high-quality InN. With further increasing the
growth temperature, the surface morphology will deteriorate
and sponge-like structures are inclined to appear. In addition,
our vertical reactor provides too small InN growth rate of about
20 nm /h for device applications.
In order to increase the growth rate, we modified the vertical
reactor (Fig. 1 (a)) into a horizontal one (Fig. 1 (b)) with a narrow
flow channel [1]. As a result, the growth rate of the InN films in thehorizontal reactor significantly increases 10 to 20 times larger
than that in the vertical reactor at 550~575°C. This is due to the
reduction of the stagnant layer thickness, the increases in thesource supply to the substrate surface and the high concentrationof the source gases by narrowing the flow channel.
Here, it should be noted that the above reactors workfrom LP to AP. Although the growth rate increases for thehorizontal reactor, the problem of the sponge-like films still(a) 650 Torr(b) 1200 Torr(c) 1600 Torr
Fig. 2 Atomic force microscope images of the InN films
grown at 675°C under 650 Torr with pyramidal structures (a),1200 Torr with hexagonal platelet structures (b), and 1600
Torr with a well-coalesced surface (c). Note that all the InN
films were directly grown on sapphire substrates, which were
just pretreated in NH3 ambient at high temperature.
exists for the higher temperature growth (>600°C). Therefore,
we designed a pressurized reactor (PR) which can work
above the AP, so as to increase the nitrogen pressure andenhance the lateral growth of the InN films.
The surface morphology of the InN films is shown in Fig.
2, where pyramid-like, platelet-like, and well-coalesced
structures can be observed for the InN grown under 650,
1200, and 1600 Torr, respectively. It is indicating that the InN
surface tends to be flat under higher reactor pressures, whichfavors the growth of the low-index plane of crystals, such as(0001) plane parallel to the substrate surface [2]. This case issimilar to gallium nitride, where pyramid-like structure andplatelet-like structure correspond to the growth under 76 and
760 Torr, respectively [3].
The above results shed a new light on MOVPE growth of
InN, which has been suffering from the various difficulties
during the last decade. By stabilizing the (0001) plane growth,
PR-MOVPE is promising for the growth of high-quality dense
InN films at higher growth temperature, making a step closer
to the InN applications in the optoelectronic devices.
References[1] Y. H. Liu, T. Kimura, T. Shimada, M. Hirata, M. Wakaba,
M. Nakao, S. Y. Ji, and T. Matsuoka, phys. stat. sol. (c), 6,
S381 (2009).
[2] Y. H. Liu, Y. T. Zhang, T. Kimura, M. Hirata, Y. Ohta, S. Y.
Ji, and T. Matsuoka, The 56th Spring Meeting of Jpn. Soc.
Appl. Phys. and Rel. Soc., Tsukuba Univ., Mar. 31, 2009.
[3] T. Sasaki and T. Matsuoka, J. Appl. Phys. 77, 192 (1995).

(a) Vertical reactor(b) Horizontal reactor
Fig. 1 Schematic diagrams of the reactor geometry. Thermal
convection easily occurs for the vertical reactor with a large
diameter flow channel (a), but relieved for the horizontal
reactor with a very narrow flow channel (b). For InN growth,
trimethylindium was used as a group-III MO source, NH3 asa group-V source, and N2 as a carrier gas. The growthpressure was 650 Torr.

KINKEN Research Highlights 2009 IMR

Key Words
InN, MOVPE, PR-MOVPE Growth
Contact toTakashi Matsuoka (Physics of Electronic Materials Division)e-mail: matsuoka@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

7Research

Structure of Liquid LiNbO3 analyzed by Energy Dispersive
X-ray Diffraction

8

Nano / Micro Structure Controlled Materials

measured

Q (S(Q)-1) / nm-1

RMC

10

0

-10

-20
0

50

100

150

Q / nm-1

200

250

Fig. 1 Q (S (Q) -1)) of liquid LiNbO3. The solid line is the
EDXD result, and dotted line is the RMC simulation.
400
350

Nb-Li

Nb-Nb
Nb-Li

300

200

O-O

150

Nb-Li
O-O

250
Nb-O
Nb-O

In order to improve the single crystal quality produced by
the ordinary Czochralski method, a variety of physic-chemical
properties of liquid LiNbO3 have been measured. The
variations in such physical properties indicate an anomalous
behavior just above its melting point (1526K) and the structure
analysis of liquid LiNbO3 near the melting point is strongly
required so as to elucidate the structural origin of suchanomaly. Although the ordinary angular dispersive X-ray
diffraction technique by using Mo Kα radiation (ADXD) allows
us to obtain structural information for this purpose, theobtained structural factor less than about 150nm-1, produces
a serious limitation to make the interesting discussion on thelocal ordering structure. The energy dispersive X-ray
diffraction (EDXD) is known to be one of the most usefulexperimental methods, which serve the structural informationin the large wave vector region over 200nm-1 [1]. This strongly
prompted us to apply this advanced EDXD technique to thestructural analysis of liquid LiNbO3 [2].
The resultant Q (S (Q) -1) and corresponding RDF aregiven in Fig.1 and Fig.2, respectively. The present EDXD-RDF
data is superior to that of the conventional ADXD-RDF [3] in its
resolution because the upper limit of Fourier transformation isincreased up to 230 nm-1. In particular, a hump at larger r sideof the first Nb-O correlation peak at about 0.190nm is clearlyobserved in the EDXD-RDF, which indicates an appreciateddisparity in the nearest-neighbor Nb-O atomic pairs. Suchgeometrical features in the local ordering units of oxide
compounds are frequently realized in coordination polyhedra
linked by sharing their edges or faces. Therefore, rather
distorted NbO6 octahedron and its linkage can be suggested
as one of the common local structural image in liquid LiNbO3.
Although there is a close correspondence in the correlation
distances of the liquid and crystalline LiNbO3 up to about 0.4
nm, the RDF of liquid LiNbO3 indicates a relatively lower
probability of finding correlations at around 0.4 nm than that
expected from the first Nb-Nb (0.383 nm) and Nb-O (0.388nm) pairs in the crystalline structure. Since this range of
distances corresponds to the distribution of NbO6 octahedra in
the crystalline structure, the structure of liquid LiNbO3 can befeatured by its unique distribution of NbO6 octahedra.
In order to discuss the detailed structure of liquid LiNbO3,
we employed the reverse Monte Carlo (RMC) simulation
technique where the atomic configurations are estimatedrather simply on the grounds of their consistency with theexperimental data [4]. The present RMC simulation results
are found to reproduce the independent interference functions
Q (S (Q) -1) well, as shown in Fig.1. The obtained partial
structure factors indicate that the correlations of aNbNb (Q)agrees well with the position of the shoulder at about Q=21nm-1
enhanced near the melting point of LiNbO3 [3]. This feature, inparticular, is quite likely to correlate with the shoulder attributed
to the development of clusters originated from a clear Nb-Nb

Liquid LiNbO3 at 1550K

20

Radial distribution function

The structural analysis of the liquid LiNbO3 at 1550K has
been carried out by the energy dispersive X-ray diffraction(EDXD). The distorted NbO6 units with a disparity in the first
Nb-O atomic pairs were well confirmed as the fundamental
local structural unit in liquid LiNbO3.

100
50
0
-50

-1

EDXD(Qmax=230nm )-1
ADXD(Qmax=135nm )

0

0.1

0.2

0.3

0.4

0.5

0.6

r / nmFig. 2 The solid line is the EDXD-RDF, and dotted line is the
ADXD-RDF [2]. The arrows indicate atomic pairs in the
crystalline LiNbO3.
correlation. The development of Nb-Nb correlation may be
one of the important reasons for the origin of anomalous
behavior detected in the physical properties in the liquid statenear the melting point. It would be of great interest to apply
the present analytical approach to the structural investigationof high temperature liquids near melting points.
References[1] K.Sugiyama, V.Petkov, S.Takeda, Y.Waseda, Sci. Rep. Res.
Inst. Tohoku Univ. A38 1 (1993).
[2] K.Sugiyama, M.Saito, Y.Waseda, J. Crystal Growth.311 966(2009).
[3] K.Sugiyama, K.Nomura, Y.Waseda, P.Andonov, S.Kimura,
K.Shigematsu, Z. Naturforsch. 45a 1325 (1990).
[4] M.Saito, C.Park, K.Omote, K.Sugiyama, Y.Waseda, J. Phys.
Soc. Japan, 66 633 (1997).
Key Words
Liquid Structure, X-ray Diffraction, EDXD
Contact toKazumasa Sugiyama (Non-Crystalline Materials Division)e-mail: kazumasa@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2009

Research

Mechanically and Biologically Biocompatible Metallic Material
Fabricated by Combining Porous Titanium with Medical Polymer

A low Young’s modulus similar to that of bone is one of
the important properties of metallic biomaterials, because itis advantageous to inhibit bone atrophy and absorption due
to stress shielding. It is considered that the use of porous
materials is one of the most effective techniques for achieving
a low Young’s modulus. However, the mechanical propertiesof porous materials deteriorate with an increase in porosity.
This deterioration may be attributable to stress concentrationnear the pores of porous materials. Thus, the stressconcentration can be reduced by filling the pores with certain
materials, which in turn is expected to improve the mechanicalproperties of the porous materials. In such a case, materials
having a low Young’s modulus should be selected for filling,in order to prevent an increase in the Young’s modulus.
Furthermore, depending on the type of filled materials,
biofunctions, which are not intrinsically present in metallic
materials, can be introduced to the porous materials, inaddition to the improvement in mechanical properties. In thisstudy, pores of pTi were filled with a medical polymer, PMMA,
and then, the effects of the PMMA filling on the mechanicalproperties of pTi were investigated as a model case [1].
In order to prepare pTi, pure titanium powders weresintered at a constant temperature. Some pTi specimens
having porosities in the range of 22−50% were obtained asshown in Fig. 1. The pTi specimens are labeled according to
the maximum particle diameter of pure titanium powders andthe porosity of pTi. Some tensile specimens of pTi were
subjected to silane coupling treatment (Si treatment) before
carrying out PMMA filling. A tensile specimen of pTi was
(a)

(f)

(c)

50 µm

50 µm

(g)

(d)

(b)

50 µm

50 µm

50 µm

50 µm

(e)

50 µm

Fig. 1 SEM
of ofcrosssections(a) pTi45-22,
Fig.micrographs
1 SEM micrographs
cross sectionsof (a)of
pTi45-22,(b) pTi45-35,(c) pTi150-27,(d)(d) pTi150-38,(e) pTi150-45,(b) pTi45-35,(c) pTi150-27,
pTi150-38,(e) pTi150-45,(f) pTi250-45, and (g) pTi250-50.
(f) pTi250-45, and (g) pTi250-50.

KINKEN Research Highlights 2009 IMR

400

pTi
pTi/PMMA
Si-treated pTi/PMMA

350
Tensile strength, σB / MPa

One of the concepts for preparing novel metallic
biomaterials possessing both an ultra-low Young’s modulus
that is almost the same as that of human bone andbiofunctions usually possessed by living tissues is proposedin the present study. By introducing pores in bulk materials,
the Young’s modulus of the materials can be reduced toalmost the same value as that of bone. Further, filling poresof porous materials with a polymer possessing a certain
biofunction is expected to cause the biofunctionalization of
the porous materials. Therefore, porous titanium (pTi) filledwith polymethylmethacrylate (PMMA) was fabricated using
a new process developed in our laboratory. Then, the effectof the PMMA filling on the tensile strength and Young’s
modulus of pTi was examined. As a result, it was found that
the PMMA filling was useful for improving the tensile strength
while maintaining a low Young’s modulus of pTi.

300
250
200
150
100
50
0

pTi45-22

pTi45-35 pTi150-27 pTi150-38 pTi150-45 pTi250-45 pTi250-50

Fig. 2 Tensile strengths of pTi, pTi/PMMA, and Si-treatedFig. 2 Tensile strengths of pTi, pTi/PMMA, and Si-treated pTi/PMMA.
pTi/PMMA.
soaked in a monomer solution mixed with a polymerization
initiator. After removing air bubbles in pores under reducedpressure, the infiltrated monomer solution was polymerized
by heating it at a constant temperature under atmosphericpressure. Then, the unwanted part of PMMA was removed
by mechanical machining.
Figure 2 shows the tensile strengths of pTi, pTi filled with
PMMA (pTi/PMMA), and Si-treated pTi/PMMA; they decrease
with an increase in porosity. However, in the high porosity
range, the tensile strengths of pTi/PMMA are higher thanthose of pTi. This result indicates that the effect of the PMMA
filling on the improvement in the tensile strength of pTi islimited to the low tensile strength range. In contrast, in thecase of Si-treated pTi/PMMA, even higher tensile strength isobtained, and this improvement is not limited by porosity.
Further, Young’s moduli of pTi, pTi/PMMA, and Si-treated pTi/
PMMA were measured using the stress−strain curve obtainedby the tensile test. It was observed that the PMMA filling had
a very small influence on the Young’s modulus of pTi, because
the Young’s modulus of PMMA is lower than that of pTi.
The developed process of filling the pores of porous
materials with a polymer is expected to be applicable also topolymers other than PMMA. By selecting an appropriate
polymer, biofunctions can be added to porous materials. For
example, if pores are filled with a biodegradable polymer
mixed with an agent for promoting bone formation andgrowth, osteoconductivity is expected to improve and behigher than that in the case of pTi with unfilled pores. On thebasis of this concept, we have initiated the application of thedeveloped process to poly-l-lactic acid (PLLA), which is oneof the practical biodegradable polymers for biomedical
applications, as a next step in this study.
References[1] M. Nakai, M. Niinomi, T. Akahori, H. Yamanoi, S. Itsuno,
N. Haraguchi, Y. Itoh, T. Ogasawara, T. Onishi and T.
Shindoh, Japan Inst. Metals, 72 (2008) 839−845.
Key Words
Biomaterial, Porous Material, Metal/Polymer Composite
Contact to
Mitsuo Niinomi (Biomaterials Science Division)e-mail: niinomi@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

9Research

Highly-Glass Formable andBiocompatible Ti-based Metallic Glasses
Ti-based metallic glasses have a great potential as
structural materials due to their good mechanical properties,
chemical properties and excellent processing performanceowing to unique thermal properties [1, 2]. Meanwhile, puretitanium and some of its alloys show good performances forbiomaterials according to their mechanical and chemical
properties [3].
We designed our new system of Ti-based metallic glasses
by investigating restricted alloying elements biocompatiblefor dental implants and selected high glass formable Ti-based
metallic glasses. The Ti-based metallic glasses do not contain
Al, V and Ni elements which are well known to be neurotoxicity
and cytotoxicity for human body. Current medical reports ofimpracticability by these elements have been a hot issue inbiomaterials science. Our newly designed Ti-Zr-Pd-Cu-Sn-P
metallic glasses, namely the Ti45Zr10Pd10Cu35-x (Sn0.83P0.17)
x (x=3, 5 and 6) alloys, have a high potential to be applied forbiomaterials as dental implant. The detailed results are
summarized as follows:
Ti45Zr10Pd10Cu35-x (Sn0.83P0.17) x (x=3, 5 and 6) metallic
glasses were fabricated by mold casting technique.
Ti45Zr10Pd10Cu35-x (Sn0.83P0.17) x (x=3, 5 and 6) metallic
glasses exhibits wide supercooled liquid regions (∆Tx= TxTg) >60K, high reduced glass transition temperatures(Trg=Tg/Tl) >0.59 and Y parameters (Tx/ (Tg+Tℓ)) ≥0.40. Here,Tx, Tg and Tl are crystallization, glass transition and liquidus
temperatures, respectively. Its critical rod diameter for fully
amorphous sample is 5 mm (Figure 1).

Fig. 2 Anodic polarization curves in HBSS aqueous solutions
open to air at 310 K (the exposed area was 1 cm2)

Fig. 3
Evaluationofcytotoxicity
foras-castTi45Zr10Pd10Cu32Sn2.49P0.51 (a) and Ti45Zr10Pd10Cu29Sn4.98P1.02(b) metallic glass cylindrical rods with a diameter of 3.5mm inosteoblasts at 310 K for 7 days
Especially, the optimized Ti-based metallic glass has
high corrosion resistance with better passivity in a wide
passivation range in simulated body fluids at 310K (Figure 2).
In addition, biocompatibility of Ti-based metallic glasses was
also evaluated by cell culture in vitro. Excellent biocompatibilityof these Ti-based metallic glass shows high potentials to be
applied as biomaterials that necrosis of osteoblast (SaOS2)
was not detected in this study (Figure 3).
References[1] A. Inoue: Acta Mater., 48 (2000), p.279[2] J. -J Oak, D. V. Louzguine, A. Inoue: J. Mater. Res., 22(2007), p.1346[3] S. Steinemann. Ti ‘84 Science and Technology, G.
Lütjering, U. Zwicker and W. Bunk, Eds. (1984) p.1373.
Key Words
Metallic Glass, Biocompatibility, Dental ImplantFig. 1 The structure of the Ti45Zr10Pd10Cu30Sn4.15P0.85 bulk glassrod samples of 5.0mm (a) and 5.5mm (b) in diameter in as-caststate, TEM. (1) – Bright-field image, (2) – dark-field image and(3) – selected-area electron diffraction (SAED) patterns

10

Nano / Micro Structure Controlled Materials

Contact toHidemi Kato (Laboratory of Non-equilibrium Materials
Division)e-mail: hikato@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2009

Research

How to Push the Stubborn Protein to Appear on the Growth Scene
We can handle the nucleation of white hen-egg lysozyme
crystal in either way of enhancement or retard via applying
an external electric field. This is attributed to the manipulationof the magnitude relation of an electrical permittivity oflysozyme between crystal and solution by varying frequencyof the applied field.
Nucleation of protein crystals is never easy going. We
need to wait a week, a month or more with patience to see
their appearance in a small solution cell. This is because they
have a fairly complex structure with high molecule weights.
The enigmatic behavior of water complex surrounding thefragile protein makes its birth more difficult. Tons of refined
researches have been tried to stimulate the silent solution toinitiate the nucleation, i.e., stirring the solution, radiating laser
or ultrasonic sound on the solution, etc. Some worked andsome did not. However, such techniques can change thenucleation in only one way, i.e., enhancement or retard. And
the worse thing is that the retardation is most often
observed.
Applying an external electric field, proposed here, may
be a unique solution to this problem. The solution principle issimple. Nucleation occurs due to the difference in chemical
potential between solid and solution. This means if we canmanage the energy difference in some way, we can control

the nucleation. An external electric field adds an electrostatic
potential to the chemical potential of both solid and liquid andwould change the energetic relationship between them. The
new thermodynamic state brings about a change in driving
force for nucleation. We have demonstrated the retard of thenucleation rate of YBCO by the imposition of an electric field
as large as 104 V/cm that is sustained in the electric double
layer between the solid and liquid [1]. However, it is notpossible to enhance the nucleation of YBCO since themagnitude relation of an electrostatic potential between solidand liquid cannot be reversed with an electric field that ispractically applied.
Because the electrostatic potential added to chemical
potential is a function of an electrical permittivity, ε, and an
applied electric field, E, i.e., ∂ε/∂x E2(x denotes composition),
its magnitude relation between solid and liquid could bereversed if an electrical permittivity would change. This ispossible for organic materials such as lysozyme due to thedispersion of their electrical permittivity in the range of 0 ~ 1
MHz with several hundred applied voltage. In contrast
inorganic material requires GHz-order frequency for theconversion that is not practically possible.
Nucleation in a small droplet of lysozyme solution under anelectric field with different frequency was investigated (Fig. 1).
It is interesting to see that the nucleation of lysozyme can beeither enhanced or retarded by choosing appropriate frequencyof an applied electric field (Fig. 2). This is the first work that
demonstrates the control of nucleation of lysozyme in either
way of enhancement or retard by an applying electric field [2].
References[1] X. Huang, S. Uda and S. Koh, J. Cryst. Growth, 307, 432(2007).
[2] H. Koizumi, K. Fujiwara and S. Uda, Crystal Growth &Design, 9, 2420 (2009).
Key Words
Electric Field, Nucleation, Lysozyme

Fig. 1 Experimental set up for the nucleation in a droplet oflysozyme solution encapsulated in oil. An externally applied
electric field of 860 V/cm brings about an effective electricfield as large as ~105 V/cm in an electric double layer whichforms at the solid-liquid interface.

Contact toSatoshi Uda (Crystal Chemistry Division)e-mail: uda@imr.tohoku.ac.jp

Fig. 2 Number of lysozyme crystals appearing in a droplet depends on the frequency of an applied electric field. They reflect thefrequency-dependent nucleation rate.

KINKEN Research Highlights 2009 IMR

Nano / Micro Structure Controlled Materials

11Research

High-speed and Low Temperature Coating of
α-Al2O3 on Cutting Tools by Laser CVD
α-Al2O3 coating is essential for cutting tools to improve
performance and to elongate the lifetime. The lowtemperature deposition of α-Al2O3 films enables one to
expand wide applications of Ti (C, N) -based cermet cutting
tools. Laser chemical vapor deposition (LCVD) can coatα-Al2O3 films directly on Ti (C, N) materials at low temperatureand high deposition rate, which has never been attained byconventional coating techniques.
Tungsten (W) is an indispensable resource for cutting
tools because of the wide usage for WC-Co cemented
carbides. However, W is a localized element on earth, andthen the substitution of W for other elements is a recent key
technology for cutting tool industry. Ti (C, N) -Ni cermets canbe a promising candidate due to abundance of Ti element
and high hardness. The cutting performance of WC-Co
cemented carbide and Ti (C, N) -Ni cermets strongly depends
on the surface coating by α-Al2O3 films (Fig. 1 left).
Commercial available WC-Co cemented carbides are
generally coated with α-Al2O3 films prepared by thermal CVD(chemical vapor deposition) to enhance the hardness and toprotect from heating. The deposition temperature of α-Al2O3
by conventional thermal CVD has been commonly more than
1300 K which is too high for Ti (C, N) -Ni due to significant
diffusion of Ni into the coating layer, and its deposition rate is
too low around less than several μm/h. The thick coating
layer of α-Al2O3 more than several 10s μm is also required
due to recent severe cutting conditions. Therefore, the lowtemperature and high-speed deposition of α-Al2O3 films ishighly demanded to develop high performance Ti (C, N) -Ni
cermet cutting tools particularly for the reduction of W usage(Fig. 1 right).
We have developed new coating techniques of novel
ceramic materials with controlling nano- and macro-texture
enabling low temperature and high-speed deposition, using
high-power Nd:YAG [1, 2] and laser diode in a continuousmode significantly accelerating chemical reactions forming
plasma around a substrate zone. We have already reported
tremendously high deposition rates of YSZ (yttria stabilized
zirconia), TiO2, Y2O3 etc. around several 10 to a few mm/h.

Fig. 2 Microstructure of α-Al2O3 coatings with (006) (left) and(104) orientation (right).

We have been taking a part of national project of “Rare
Metal Substitute Materials Development Project” by the New
Energy and Industrial Technology Development Organization(NEDO) since 2007, and have revealed that laser CVD(LCVD) using Nd:YAG laser is able to lower the deposition
temperature about 200 K compared with conventional thermal
CVD [3]. Significantly oriented and columnar textured grains
have been grown with high deposition rates of 1 mm/h at
most. (006) and (104) oriented α-Al2O3 coatings are
successfully synthesized as demonstrated in Fig. 2. LCVD
using diode laser is further effective to lower the deposition
temperature move than 400 K less than conventional thermal
CVD with significant (006) orientation which is highly requiredfor practical coating for cutting tools due to much higher andsmoother surface.
References[1] T. Goto, Surf. Coat. Technol. 198, 367 (2005).
[2] T. Goto, R. Banal, T. Kimura, Surf. Coat. Technol. 201,5776 (2007).
[3] T. Goto, H. Kadokura and T. Kimura, Proceeding of 14th
International Metallurgy & Materials Congress, 921 (2008).
Key Words
Laser CVD, α-alumina Coating, Rare Metal SubstituteMaterials
Contact toTakashi Goto (Multi-Functional Materials Science Division)e-mail: goto@imr.tohoku.ac.jp

Fig. 1 Schematic of a cutting tool with α-Al2O3 hard coating(left) and laser CVD α-Al2O3 coating process by using
Nd:YAG laser and laser diode (LD) (right).

12

Nano / Micro Structure Controlled Materials

IMR KINKEN Research Highlights 2009

Research

Optimization of Hot Working Process by “Processing Maps”
We propose the optimization of hot working by processing
map, which is demonstrated by the efficiency of power
dissipation and the plastic instability parameter. In this work,
friction and adiabatic corrections are conducted on thebasis of experimental and theoretical considerations, and
the Processing map of biomedical Co-29Cr-5Mo-0.16N alloyis established.
Hot working processes are extensively used in
manufacturing products and for improving the mechanicalproperties of materials, etc. A fundamental method forinvestigating the working behavior of materials is by analyzing
the true stress–true strain curves and by observing the
microstructure of materials; this helps in understanding the
intrinsic mechanical characteristics of materials, thereby
optimizing the forging process. However, friction andtemperature rise of specimen along with other factors have to
be considered before further investigation [1, 2].
Processing maps, including both power dissipation andinstability maps, are developed on the basis of the Dynamic
Material Model. This model considers that the workpiece is a
power dissipater and that the instant power dissipated at agiven strain rate consists of two complementary parts: the G
content and the J co-content, which are related to temperaturerise and the microstructure dissipation, respectively. The
factor that partitions the power between J and G is the strainrate sensitivity (m) of the flow stress (σ) ; the J co-content isσσ ⋅έ ⋅m, where έ is the strain rate. The
given by J = ∫ έ ⋅ dσ →0

Fig. 1 3-D power dissipation map of Co-29Cr-5Mo-0.16N alloyat strain of 0.6.

Fig. 2 Contour map of power dissipation map of Co-29Cr-5Mo0.16N alloy at strain of 0.6.

m +1

condition for an ideal linear dissipater is J max =σ ⋅έ2. For anon-linear dissipater, the efficiency of power dissipation may
be expressed in terms of a dimensionless parameter η=2m/(m+1). The variation of η with the temperature and strain rate
constitutes the power dissipation map, and its domains may
be interpreted in terms of specific microstructure processes.
The occurrence of instability has been proposed by Ziegler,
and a condition for microstructure-related flow instabilities is∂ log( m / m + 1)
+ m < 0. The data obtained
obtained as ξ (έ ) =∂ log έafter the friction and adiabatic corrections are used indeveloping the processing maps-power dissipation map andinstability map. Thus, it is possible to obtain the optimum
condition of the forging process without resorting to expensive
and time-consuming trial-and-error methods.
The data used for constructing the power dissipation maps
have to be compensated in terms of friction and adiabatic
heating [1, 2]. After compensation for Co-29Cr-5Mo-0.16N alloy
as a candidate of biomedical application, the instability maps, asobtained are shown in Fig. 1-3, respectively, for a strain of 0.6 [3].
The optimum condition in the power dissipation map is closely
related to the Stacking Fault Energy of the materials, and incase of Co-29Cr-5Mo-0.16N alloy, the power efficiency in therange of 20-30% is best. By combing the results from both thepower efficiency map and the instability map (stable regime),
the best condition for hot working process can be obtained.

KINKEN Research Highlights 2009 IMR

Fig. 3 Instability map of Co-29Cr-5Mo-0.16N alloy at strain of 0.6.
References[1] Y P Li, E. Onodera, H. Matsumoto, A. Chiba, Metall.
Mater. Trans. A, 40 (2009), 982-990[2] Y P Li, H. Matsumoto, A. Chiba, Metall. Mater. Trans. A,
40 (2009), 2009 (40), 1203-1209.
[3] A. Chiba, H. Lee, H. Matsumoto, M. Nakamura, Mater.
Sci. Eng. A513–514 (2009) 286–293.
Key Words
Processing Map, Co-Cr-Mo Alloy, Dynamic Recrystallyzation
Contact to
Akihiko Chiba (Deformation Processing Division)e-mail: a.chiba@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

13Research

X-ray Analytical Approach: Change in Dislocation Density andPrecipitate Size for Copper Alloy during Aging Treatment
The correlation between growth process of precipitates
and the dislocation density in a Cu-Ni-Si alloy during aging
was investigated by small angle scattering (SAXS) and
X-ray diffraction (XRD) methods. These X-ray analyses
revealed that the rearrangement of the dislocation density ina copper matrix induced coarsening of the precipitates.

0s

copper matrix
ρ=1.7x10 15 m

precipitates

1000 sρ=1.7x10 15 m -2

1.2nm

A Cu-Ni-Si system is known as one of copper alloys forindustrial applications, because it exhibits high electricalconductivity as well as high strength by highly dispersed
nickel-silicide precipitates formed during aging. It is wellknown that the precipitate coarsening, which occurs at a
prolonged aging time, results in deteriorating the mechanicalproperties. The SAXS method can be a powerful tool for aquantitative characterization of the growth behavior of nanoprecipitates in copper alloys [1, 2]. Our study using the SAXS
method has suggested that a cold-rolling treatment prior to
the aging treatment can stimulate coarsening of the
precipitates [1]. This result indicates a high dislocation density
induced by the cold-rolling treatment could have effects on
the growth behavior of the precipitates. In order to validate
this correlation between the growth behavior of precipitates
and the dislocation of the copper matrix, the dislocation
density must be estimated accurately. The characteristics on
dislocation have been generally observed and discussed byan electron microscope. However, it becomes rather difficult
to characterize it under the existence of highly-populated
precipitates with a direct observation method. This subject
prompts us to apply the XRD line-profile analysis forcharacterizing the dislocation. In this study, we will show
several interesting behaviors of the dislocation in the copper
alloy during the aging treatment.
A Cu-2.47Ni-0.58Si (wt%) alloy was prepared by aplasma arc melting method. This alloy was homogenized at1173 K for 105 s, and was cut to sheets. The sheet samples
were solution-treated at 1173 K for 1.8 ks, and were
subsequently cold-rolled by 50 %.
Figure 1 shows the dependence of Vickers hardness of
the alloy on the aging time at 720 K. The peak hardness ofabout 220 Hv was obtained at the aging time between 5 ks
and 10 ks, and the hardness eventually dropped at the aging
time of more than 20 ks.

5000 s1.6nm

10000 sρ=2.5x10 15 m

-220000 s2.1nm
50000 sρ=2.9x10 14 m

-2

-2

Peak-aged stage

ρ=2.1x10 15 m

-2

Over-aged stage

ρ=1.9x10 14 m

-2

Fig. 2 Schematic diagram of the growth behavior of precipitates.
The dislocation density (ρ) of the copper matrix is also denoted.
The coarsening of precipitates occurred with a decrease in thedislocation density.
The SAXS measurements were carried out using a threepin-hole camera with a two-dimensional position sensitive
proportional counter. The size distribution of precipitates wasestimated from the SAXS profiles using an indirect Fouriertransform procedure. The schematic diagram of the growth
behavior of the precipitates is illustrated in Fig. 2. As expectedfrom Fig. 1, the precipitates grew homogeneously till the aging
time of 10 ks, and then inhomogeneous precipitates having
larger size were formed after the aging time of 20 ks.
The XRD profiles from the copper matrix were measured
to estimate the dislocation density of the copper matrix. The
diffraction profiles, which depend on the dislocation density
and crystallite size of the copper matrix, varied with an increasein the aging time. The dislocation density (ρ) was estimatedthrough the XRD line-profile analysis of the modified
Williamson-Hall and modified Warren-Averbach methods [3](Fig. 2). It is quite interesting that the dislocation density variedwith the aging time in accordance with the variation of thegrowth behavior of precipitates. That is, the dislocation densitychanged only a little till the peak-aged condition (10 ks) anddropped by one order at the start of the over-aged condition(20 ks). This dependence of the dislocation density on the
aging time clearly supports the mechanism we suggested [1]that the coarsening of precipitates is induced by therearrangement of the dislocation in the copper matrix.
References[1] Y. Takahashi, T. Sanada, S. Sato, T. Okajima, K. Shinoda
and S. Suzuki, Mater. Trans. 48, 101 (2007).
[2] S. Sato, J. Spectrosc. Soc. Japan 57, 294 (2008).
[3] T. Ungar and G. Tichy, Phys. Stat. Sol. (a) 171, 425 (1999).
Key Words
Small Angle X-ray Scattering, X-ray Diffraction Line-profile
AnalysisFig. 1 The dependence of the Vickers hardness of Cu-2.47Ni0.58Si (wt%) alloy on the aging time at 720 K. The maximum
hardness appears at the aging time between 5 ks and 10 ks. The
over-aged state begins at the aging time of 20 ks.

14

Nano / Micro Structure Controlled Materials

Contact toShigeo Sato (Analytical Science Division)e-mail: s.sato@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2009

Research

Atomic Structure Imaging of Nanoparticles by Aberration
Corrected HRTEM & Exit-Wave Reconstruction
The atomic structure of FePd nanoparticles has beenstudied by aberration corrected high-resolution transmission
electron microscopy. Chemical order is clearly seen as
bright contrast, free of imaging artifacts, due to the small
negative value of corrected third-order spherical aberration.
The reconstructed phase image directly shows the projected
potential distribution within the specimen and reveals theelemental differences due to chemical order.
Recent development of ultra-high density magnetic
storage technology requires novel recording media with ahigher magnetocrystalline anisotropy energy (MAE), aiming
to increase the storage density and to decrease the recording
noise. One of the candidates for future storage media is FePd
nanoparticle with the L10-type ordered structure. The hardmagnetic properties of this alloy nanoparticle originate fromthe tetragonal ordered structure with a high MAE. Therefore
the atomic ordering is a key issue of hard magnetic properties.
We used aberration corrected (AC) high-resolution
transmission electron microscopy (HRTEM) for the imagingof chemical order [1]. Besides highly improved spatial
resolution, AC-HRTEM benefits from smaller defocus values
at the optimal defocus condition, almost corresponds to theGaussian focus, due to small aberration values. This
technique has now become practically available.
The FePd alloy nanoparticles studied were formed as
epitaxial islands on a NaCl (001) substrate using sequential
deposition of Pd and Fe followed by postdeposition annealing.
HRTEM images were obtained by using an FEI Titan 80-300
TEM operating at 300 kV equipped with a field emission gun(FEG) and a CEOS image corrector [installed at Tohoku
University].
Figure 1 (a) shows HRTEM image of a 10-nm-sized FePd
nanoparticle with the c-axis of the L10 structure oriented
normal to the film plane, taken at Cs = -0.34 μm (Cs: 3rd-order
spherical aberration). The arrangement of atoms due tochemical order can be seen unambiguously as periodic brightcontrast in the HRTEM image even at the interface between
the crystal and amorphous, as well as the inside of thenanoparticle. Under the present experimental condition andwith the weak phase object approximation, the projected
potential is imaged as a bright contrast. Figure 1 (b) shows
an HRTEM image of an ordered FePd nanoparticle with its
c-axis oriented parallel to the film plane. Apparent in Fig.1 (b)Fig. 2 (a) Phase image of exit-wave reconstructed by focalseries HRTEM images, and (b) a magnified phase image. The
focal step analyzed a posteriori was 2 nm on average. The
square indicates the unit cell of the ordered structure. (c)Intensity profiles for Pd and Fe rows measured in the [110]direction on the phase image.
is bending of the lattice image along the c-axis. For example,
average (001) plane distance on lines 1-1’ and 2-2’ is distorted
by –9.2% and –5.2%, respectively, relative to those on theline 3-3’. The observed distortion may have its origin in a
possible compositional variation or strain distribution,
however, the detail still remains an open question.
Figure 2 shows the phase of the exit-wave function byreconstructing a focal-series of HRTEM images. Two kindsof bright dots, strong and weak, arise from the potentials of
Pd and Fe atoms. Figure 2 (c) shows an example of intensity
profile of the reconstructed phase both for Fe and Pd. Thus,
the reconstruction of phase image makes us possible toreveal the elemental difference as well as the clear-cut
images of atomic rows, which can be compared to the AC
HAADF-STEM images [2].
In summary, we demonstrated that the atomic structure of
the L10-FePd nanoparticles can be observed unambiguously
by AC-HRTEM. The reconstructed phase image directly shows
the projected potential distribution within the specimen. The
clear-cut long-range order is lost when the size of particles issmaller than about 5 nm, and locally ordered particles become
dominant, indicating the instability of the ordered FePd phase.
It is therefore suggested that such a size effect is likely to place
limits on industrial applications of small FePd nanoparticles.
References[1] K. Sato, T. J. Konno and Y. Hirotsu, J. Appl. Phys. 105,
034308 (2009).
[2] K. Sato, J. G. Wen and J. M. Zuo, J. Appl. Phys. 105,
093509 (2009).
Key Words
Electron Microscopy, Nanoparticle, Exit-Wave Reconstruction

Fig. 1 (a) HRTEM image of an L10-FePd nanoparticle with thec-axis oriented normal to the film plane. The defocus is estimatedto be -13 nm by analyzing the power spectra. (b) HRTEM image
with the beam incidence of [100]FePd.

KINKEN Research Highlights 2009 IMR

Contact toKazuhisa Sato (Advanced Analysis of Materials Division)e-mail: ksato@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

15Research

Malleable Hypoeutectic Zr-Ni-Cu-Al Bulk Glassy Alloys
with Tensile Plastic Elongation at Room Temperature
A new bulk glassy alloy (BGA) showing macroscopic tensile
plastic elongation at room temperature has been developed in
hypoeutectic Zr-Ni-Cu-Al alloys with high Poisson’s ratio and low
Young’s modulus. Hypoeutectic Zr70Ni16Cu6Al8 BGA exhibits thehighest Poisson’s ratio of 0.39 and the lowest Young’s modulus
73 GPa, a distinct tensile plastic elongation of about 1.7 % androom temperature malleability in compression without fracture.
The highly ductile nature seems to originate from modifications in
the medium-range atomic order in hypoeutectic region.
Figure 1 (a) shows compositional dependence of Poisson’s
ratio of as-cast hypoeutectic Zr-Ni-Cu-Al10 glassy rods of diameter
8 mm. Near the eutectic composition of Zr54Ni6Cu30Al10 [1], thePoisson’s ratio decreases monotonically with increasing Cu
content alone and increases with simultaneous rise in both Zr and
Ni contents. Zr70Ni15Cu5Al10 BGA exhibits the highest Poisson’s
ratio among hypoeutectic alloys with μ = 0.38. Figure 1 (b) shows
the compositional dependence of Young’s modulus as well as
Poisson’s ratio. Zr70Ni15Cu5Al10 BGA also exhibits the lowest
Young’s modulus of E = 78 GPa. Since the compositional
dependence of Young’s modulus in Zr-Ni-Cu-Al BGA has a positive
correlation with Al content, we modify the Zr70Ni15Cu5Al10 alloy
composition by decreasing slightly the Al concentration [1]. As aresult, we found the Zr-enriched hypoeutectic Zr70Ni16Cu6Al8
glassy alloy with the highest Poisson’s ratio of 0.393±0.003 andthe lowest Young’s modulus of 73±4 GPa. Consequently, theZr70Ni16Cu6Al8 BGA was used for subsequent mechanical testing.
Figure 2 (a) shows the corresponding stress-strain curve. One
can see a distinct linear elastic region, followed by clear yielding andthen plastic elongation. Figure 2 (b-e) shows an over view of thetensile fracture surface of the Zr70Ni16Cu6Al8 BGA with a tensile true
plastic strain of 1.7 %. Many shear band steps are observed on thesurface near the wide sliding region (Fig. 2 (b)). Near the center offracture surface, we can see the typical flower like vein patterns
even in the tensile stress mode, as shown in Fig. 2 (c). Behind thefracture surface, shear band traces are seen on the outer surface,
as shown in Fig. 2 (d), and the magnified image reveals the existenceof many shear band steps near the fracture surface (Fig. 2 (e)).
implying the formation of a single glassy phase.
The structures of bulk metallic glasses are characterized by thepacking together of efficiently packed atom clusters (chemical andFig. 2 Tensile true stress vs. true strain curve of hypoeutectic
Zr70Ni16Cu6Al8 BGA (a) and SEM image of all over tensile fractured
surface (b), center region of tensile fractured surface (c), behind thetensile fractured surface (d) and magnified image of region A in (d)(e). Arrow marks point out the surface steps of shear bands near
the tensile fractured surface.
topological short range order (SRO) on a scale of about 0.5 nm) with
the inter-cluster connectivity defining the medium range order(MRO) detectable to 1-2 nm in range by diffraction methods [2, 3].
This MRO is the same scale as the dimension of shear transformation
zones (STZs), which must be stress-activated and is connected toinitiate a shear band [4]. Furthermore, positron annihilation analysis
of coincidence Doppler broadening experiments [5] on ternary
eutectic Zr50Cu40Al10 BGA reveals that free volume is surrounded bya higher fraction of Zr atoms as compared to the nominal alloy
composition. The significant decrease in Young’s modulus in Zrenriched BGAs (Fig.1 (b)) may be caused by the formation ofbifurcated percolation path composed of weakly bonded regions
enriched in nearest Zr-Zr atom pairs.
References[1] Y. Yokoyama and A. Inoue, Mater. Trans. 48, 1282-1286 (2007).
[2] Miracle, D. B., A structural model for metallic glasses, Nature
Mater. 3, 697–702 (2004).
[3] Sheng, H. W., Luo, W. K., Alamgir, F. M., Bai, J. M. & Ma, E.,
Atomic packing and short-to-medium-range order in metallic
glasses, Nature 439, 419–425 (2006).
[4] Yavari,A.R., A new order for metallic glasses, Nature 439, 405406 (2006).
[5] T. Yano, Y. Yorikado, Y. Akeno, F. Hori, Y. Yokoyama, A. Iwase,
A. Inoue and T.J. Konno, Mater. Trans. 46, 2886-2893 (2005).
Key Words
Bulk Metallic Glass, Tensile Ductility, Hypoeutectic

Fig. 1 Compositional dependence of Poisson’s ratio (a) and
Young’s modulus (b) of hypoeutectic Zr-Ni-Cu-Al BGAs.

16

Nano / Micro Structure Controlled Materials

Contact toYoshihiko Yokoyama (Advanced Research Center of Metallic Glasses)
e-mail: yy@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2009

Research

Effect of Excess Electrons on hcp Mg andthe Model Clusters for Bulk Metallic Glasses
We report the effect of excess electrons on hexagonalclose-packed Mg and the model clusters explained by an
inflation process using density functional theory-based
calculations, in order to understand the role of conduction
electron concentration in Mg-based bulk metallic glasses.
We find the volume expansion and distortion to a higher c/a
ratio in the negative charge state.
The role of conduction electron concentration is one of
the key roles in revealing the formation mechanism of bulk
metallic glasses (BMGs). The conduction electron
concentration of Mg has recently been estimated to be larger
in Mg-based BMGs than in crystalline magnesium. The hcp
metals have a wide range in their c/a ratio from the lowest
value of 1.568 for Be to the largest value of 1.886 for Cd, andthe c/a ratio of 1.624 of hcp Mg corresponds virtually to thecase of an ideal close packing where c/a = √¯8/3 (1.633). The
structural diversity is reflected in their physical properties
such as the melting points, cohesive energies and magnetic
susceptibilities. Lowering the melting point is an issue in theformation technique of BMG. In the sense, the effect ofexcess electrons on the structure of hcp Mg is worthy ofinvestigation. We have investigated the effects of excess
electrons on hcp Mg and the model clusters explained by an
inflation process [1]. Our aim is to understand the formation
mechanism of Mg-based BMGs.
Mg-based BMGs have been formed in very wide
Mg70M1xM230-x,
composition
ranges:Mg65Ni20Nd15,
Mg75M1xM225−x, Mg77Ni18Nd5, Mg80M1xM320−x, Mg85Cu5Y10
for M1 = Ni, Cu, M2 = Nd, Y and M3 = Nd, Sn, Y, where Mgrich BMGs are excluded. The ratios of magnesium to 3n
additives (n = 0, 1, 2, . . .) are 17 : 9, 7 : 3, 3 : 1, 10 : 3, 4 : 1,
and 17 : 3, respectively. From the ratios, we consider Mg3,
Mg4, Mg7, Mg10 and Mg17 as model clusters and examine theFig. 2 Total and partial densities of states (DOS) of hcp Mg.
(a) Neutral state. (b) At the expanded lattice constants bycharging but with no excess electron.
effect of excess electrons. The inflation rule, A → AB and B→ A where A = Mg3, B = Mg4, leads to the consecutive
sequences, S0 = A (Mg3), S1 = AB (Mg7), S2 = ABA (Mg10)
and S3 = ABAAB (Mg17).
Both in the clusters (Mg7, Mg10 and Mg17) and hcp Mg,
the c/a ratio increases proportionally to the concentration ofexcess electrons (Fig. 1) and the cell volume expands. The
charged clusters up to dianion are more stable than thecorresponding neutral clusters, suggesting that the increasein the conduction electron concentration of Mg in BMGs ispossible to stabilize the cluster structures. The distortion to ahigher c/a ratio suggests that the melting point would belowered in the charge state. The pseudogap by s–p mixing at
the Fermi level is weakened in the expanded cell with adistorted c/a ratio and a new spiky structure mainly contributedby p bands appears (Fig. 2), indicating a signature of states
preferentially localized around the structural cluster which
occurs in quasicrystalline materials. It would be anticipatedthat the expanded volume and the distorted c/a ratio by the
charge transfer from additives suppress the crystallization ofsupercooled liquid. Further theoretical study on Mg–Cu–Y
systems is required.
References[1] M. Takahashi, M. Fukuhara, A. Inoue, Y. Kawazoe, J.
Phys. D: Appl. Phys 41, 155424 (2008).
Key Words
Bulk Metallic Glass, Magnesium Clusters, Density Functional
Theory Calculations

Fig. 1 Normalized c/a variation (%), Δ, vs. the excess-electron
concentration (e/atom) for Mg7 (●), Mg10 (■), Mg17 (▲), andhcp Mg (★). Δ= 1/N × ((c/a of charge state) / (c/a of neutral
state) – 1) × 100. For clusters, c = L1/N, a = L2. N = 1, 1.5, 3,
and 0.5 for Mg7, Mg10, Mg17, and hcp Mg, respectively.

KINKEN Research Highlights 2009 IMR

Contact toMasae Takahashi (Advanced Research Center of Metallic
Glasses)e-mail: masae@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

17Research

Development of Iron-group Based Glassy Alloy with HighDeformability and High Folding Endurance
Fe-, Ni- and Co-based glassy alloys have been developedwith modifying to meet the needs for application. Since folding
endurance value of the developed Ni-based glassy alloy is 10
times as large as that of stainless SUS304 steel, it become
possible to develop highly flexible precision products.
Fe-, Ni- and Co-based (Iron-group based) amorphousalloys are marketed in ribbon or wire form, and are used invarious fields. To produce the precision and small parts with
Iron-group based amorphous alloy in ribbon or wire form, theimprovement of the alloy is expected in imprintability and folding
endurance.
High plastic deformability with Newtonian viscous flow of
the amorphous alloy (glassy alloy) allows imprinting of a
precision component. Figure 1 (a) and (b) show the SEM
images of the surface morphology of the electroformed Ni dieand the imprinted (Fe0.73Ga0.04P0.11C0.05B0.04Si0.03)99Cr1 glassy
alloy ribbon [1]. The surface structure of the imprinted sample
exhibits the wave shape with 1 μm period, and that is similar tothe shape of the original Ni die. The super-cooled liquid region
defined by temperature interval between the glass transition
temperature and crystallization temperature is 45 K for (Fe0.73
Ga0.04P0.11C0.05B0.04Si0.03)99Cr1 glassy alloy. The Iron-groupbased glassy alloy with large super-cooled liquid region allows
imprinting. To raise the imprintability of glassy alloy, it isimportant to expand the super-cooled liquid region of the alloy.
Figure 2 shows the DSC curves of Co48Cr15Mo14Tm2C15B6 andNi65Cr15Si7B13 glassy alloy samples [2, 3]. It can be seen an
endothermic reaction due to glass transition, followed by asuper-cooled liquid region and then an exothermic reaction due
to crystallization. The CoCrMoCBTm and NiSiBCr glassy alloy
exhibit large super-cooled liquid region of 90 K and 63 K,
respectively. These Iron-group based glassy alloys are suitablefor imprint processing in the super-cooled liquid region tofabricate the small precision parts.
The small parts with glassy alloy are flexible and bend
easily, but the annealed parts with imprinting or coating may
lose those properties. The annealed glassy alloys need to havegood flexible durability. The flexible durability of the productswith metal and/or polymer film is evaluated by the MIT folding
endurance test. Figure 3 shows the results of the MIT folding
endurance test of the ribbons with a thickness of 0.03 mm. The
MIT folding endurance value is defined as the number of times
a ribbon can be bended without breaking on the MIT folding
endurance tester. Although the MIT folding endurance value ofthe FeSiB amorphous alloy decreases significantly withannealing because of embrittlement of the amorphous phase,the values of Ni-based alloy exhibit little change with annealing.

Fig. 2 The DSC curves of the Co48Cr15Mo14Tm2C15B6 andNi65Cr15Si7B13 glassy alloy samples

Fig. 3 The MIT folding endurance values of a Ni75Si8B17
amorphous alloy and Ni65Cr15Si10B10 glassy alloy, and thoseof a FeSiB amorphous alloy and stainless SUS304 steel are
shown for comparison
The value of the annealed Ni65Cr15Si10B10 glassy alloy withsuper-cooled liquid region is 10 times as high as that of stainlesssteel [3]. These results suggest that the NiCrSiB glassy alloy is
an imprintable and flexible material. Furthermore, the alloy
exhibits high corrosion resistance. These glassy alloys arepromising imprintable materials with high performance.
References[1] K. Amiya, Y. Saotome, N. Nishiyama, A. Urata, A. Inoue:
Proc. Asian Workshop Nano/micro Forming Tech. (2007).
[2] K. Amiya and A. Inoue, Rev. Adv. Mater. Sci., 18,
27(2008).
[3] Pat.Pend, JP-A(2008)
Key Words
Glassy Alloy, Imprint, Folding Endurance

Fig. 1 The SEM images of the surface morphology of (a) Ni
die and (b) the imprinted (Fe0.73Ga0.04P0.11C0.05B0.04Si0.03)99
Cr1 glassy alloy ribbon

18

Nano / Micro Structure Controlled Materials

Contact toKenji Amiya (Osaka Center for Industrial Materials Research)e-mail: amiya.k@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2009

Research

Discovery of Spin-Motive Force－ “Spin Battery” Using Nano-Magnets －The existence of a spin-motive force, a correction to
Faraday’s law of induction, has been proven in experimentusing a magnetic tunnel junction with nano-meter-sizedmagnetic particles [1]. This discovery enables directconversion between magnetic and electrical energies,
which opens the way for applications in ultra-high sensitivity
magnetic sensors as well as new type of “spin batteries.”
Electrical currents usually correspond to a flow ofelectrons, elementary particles that carry a tiny electricalcharge, -e. In 1930, Dirac showed that, in order to beconsistent with Einstein’s special relativity and quantum
mechanics, an electron must also have “spin,” in effect, beinga tiny magnet (Fig.1). “Spintronics” is an emerging electronics
technology in which the electron spin is used in an essential
manner [2]. The issues studied in this field include tunnelingmagnetoresistance, domain wall motion due to magnetic
fields or electric currents in ferromagnets.
One of the versatile methods to generate electric powerfor driving various electronic devices is Faraday’s law of
induction, a fundamental principle dating from 1831. The
usual Faraday’s law determines an electromotive force(e.m.f.) ε as
ε = -dΦ/dt,
where Φ is the magnetic flux. To derive this law, only forces
resulting from the charge of electrons have been accountedfor. In ferromagnetic materials, however, there exist
nonconservative “spin forces” which also contribute to thee.m.f. [3]. These might be included in Faraday’s law if Φ isreplaced by [ћ/ (-e) ]γ,
ε = -[ћ/ (-e) ]dγ/dt,
where γ = γe+γs is a “Berry phase.” γe = [ (-e) /ћ]Φ recovers
the usual charge part of Faraday’s law while γs represents
the spin contribution suitably averaged over the electron spin
direction. The latter contribution to ε refers time dependent
magnetization in general. One of the remarkable implications
is the fact that the e.m.f. can be induced even by a ‘static’magnetic field when that field drives magnetization dynamics.
The anticipated effect, however, remained elusive until
recently.
The research group consisted of Professor S. Maekawa(Institute for Materials Research, Tohoku University),
Professor M. Tanaka (Department of Electrical Engineering
and Information Systems, The University of Tokyo), andProfessor S. E. Barnes (Physics Department, University ofMiami) found that such an e.m.f can indeed be induced by astatic magnetic field in magnetic tunnel junctions containing
zinc blende-structured MnAs quantum nanomagnets (Fig. 2).
The observed e.m.f results from the conversion of themagnetic energy of the super-paramagnetic MnAs
nanomagnets into electrical energy when these magnets
undergo magnetic quantum tunneling. Due to the quantum
nature of the nanomagnets, the produced e.m.f. is over ahundred times larger and for tens of minutes, rather than for

KINKEN Research Highlights 2009 IMR

spin ћ/2-echargeFig. 1 An electron has charge and spin degrees of freedom.

A

MnAs film
AlAs barrier
GaAs matrix
MnAs
nanoparticle

Fig. 2 A schematic illustration of the device structure.

milliseconds as expected. As a consequence, a hugemagnetoresistance of up to 100,000 per cent is observed forcertain bias voltage. Their results strongly supported thecontention that, in magnetic nanostructures, Faraday’s law of
induction must be generalized to account for forces of purely
spin origin. The huge magnetoresistance and e.m.f. may findpotential applications in high sensitivity magnetic sensors, aswell as in new active devices such as “spin batteries.”
References[1] P. N. Hai, S. Ohya, M. Tanaka, S. E. Barnes & S. Maekawa,Nature 458, 489 (2009).
[2] S. Maekawa ed., Concepts in Spin Electronics (Oxford
University Press, Oxford, 2006).
[3] S. E. Barnes and S. Maekawa, Phys. Rev. Lett. 98,
246601 (2007).
Key Words
Spintronics, Spin-motive Force, Huge MagnetoresistanceContact to
Sadamichi Maekawa (Materials Property Division)e-mail: Maekawa@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

19