3FTFBSDI

Spin-Motive Force― Giving Faraday's Law a New Spin ―Faraday's law of induction is a basic principle of physics
which explains an electromotive force that drives electricalcurrents in generators and transformers. Recently, it has
been revealed that, in magnetic materials, there is acorrection to this law, which can be elegantly restated interms of the time rate of change of the so-called “Berryphase” [1].
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 [3], and the spinHall effect in paramagnetic metals with strong spin-orbit
interaction such as Pt and Au, aiming at controlling the spin
degrees of freedom.
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 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 ε.
These might be included in Faraday’s law if Φ is replaced by[ћ/(-e)]γ,
ε = -[ћ/(-e)]dγ/dt,
where γ = γe+γs is a “Berry phase.” γe = [(-e)/ћ]Φ recovers theusual charge part of Faraday’s law while γs reflects the spin
contribution suitably averaged over the electron spin direction.
These contributions to ε represent the requirements of energy
conservation in itinerant ferromagnets with time dependent
magnetization. This is an extremely common situation, forexample, the situation in an every-day power transformer. It
has more important implications in connection with
spintronics.
A Berry phase is a geometric phase, which appears
ubiquitously in the context of condensed matter physics,
optics, quantum computation, etc. The spin Berry phase γs

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

Nano / Micro Structure Controlled Materials

(a)
∆Ω
n(t)
θ
n(t+∆t)O(b)

Border paramtermetern(z,t)
z
θ

x
⋅
φ
v

y
voltmeter

Fig. 2. (a) A geometrical aspect of a spin Berry phase. (b) A
ferromagnetic wire containing a domain wall. Black arrows
indicate the direction of the order parameter (magnetization).
for an electron is –Ω/2, where Ω is the solid angle subtended
by the path of the order parameter n in spin space, as
illustrated in Fig. 2(a). The simple example, see Fig. 2(b), isa ferromagnetic wire which lies along the z direction andwhich contains a single domain wall, a transition region
connecting two magnetic domains. The ferromagnetic orderparameter is indicated by the black arrows and the unit vectorn(z,t). This defines the direction of the internal exchange field
which is specified by two angles, θ and φ. For a DC magnetic
field B along the z axis, the wall starts to precess so that it
develops a time dependent solid angle, dΩ/dt = ±4μB/ћ, for
conduction electrons traversing the wire. The + (-) sign
corresponds to spin-up (-down) of the electrons. This
generates a “spin-motive force” εs± = ±2μB/e, and hence, anelectromotive force ε = 2pμB/e, where p is a spin polarizationof the ferromagnet. The magnitude is evaluated for typicalitinerant ferromagnets, e.g., Ni81Fe19, ε ~ 100 [µV/T].
While the corrections are small for traditional electricmachines, they embody the requirements of energy
conservation for the spintronic devices being developed byIBM, Freescale, and Hitachi. In more general terms, thecorrection to Faraday's law can be demonstrated conclusively
for the first time in terms of the Berry phase.
References[1] S. E. Barnes and S. Maekawa, Phys. Rev. Lett. 98,
246601 (2007).
[2] S. E. Barnes, J. Ieda, and S. Maekawa, Appl. Phys. Lett.
89, 122507 (2006).
[3] M. Yamanouchi, J. Ieda, F. Matsukura, S. E. Barnes, S.
Maekawa, and H. Ohno, Science 317, 1726 (2007).
Contact to
Sadamichi Maekawa (Theory of Solid State Physics Division)e-mail: maekawa@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

Direct Observation of Field-induced Variant Transformation in
Fe3Pt Using Pulsed Magnetic Field X-ray Diffraction
A field induced variant transformation and its dynamical
response have been measured by combining a fast
sweeping high magnetic field and a synchrotron X-ray. It is
confirmed that a variant can rotate in a short time of msec
range.
A time resolved synchrotron X-ray diffraction is one of
powerful tools to investigate a material response triggered byextreme parameter such as high magnetic fields. A dynamical
behavior of martensitic transformation has attracted muchattention because of the possible application as a fast
actuator. An example is found in the giant magnetic-fieldinduced strain of the ordered Fe3Pt alloy. However, the direct
observation of the variant transformation has been difficult
and thus the response of the variant has not been known sofar.

magnetic-field-induced strain in the martensitic phase ofordered Fe3Pt alloy. Figure 2 shows the time dependence ofthe ω-scan profile of the sample. Two different variants areobserved as two distinct peaks. The intensity of the rightpeak decreases with increasing magnetic fields. It is the first
direct evidence of the field induced rearrangement of thevariants (crystallographic domains) in pulsed magnetic fields.
The decrease of the right peak intensity is caused by therotation of the variants from the c axis perpendicular to themagnetic field to the c axis parallel to the magnetic field.
Moreover, we found that the reorientation behaviors aredifferent for different variants with the same crystallographic
direction. The results demonstrate the crucial importance of
microscopic investigation, such as X-ray diffraction analysis,of the martensitic transition.

Fig. 2 The field dependence of ω-scan profile of Fe3Pt alloyin fast sweeping magnetic fields. The right peak looses theintensity for variant transformation.

Fig. 1 A small pulsed magnet used to observe the variant
transformation in fast sweeping high magnetic field. A
magnetic field of msec pulse duration can be generated.

In the present work, we have developed a novel technique
to investigate the field induced martensitic phase transition.
Figure 1 shows the mini magnet used for time resolved X-ray
diffraction experiment in a fast sweeping magnetic field.
Thanks to the extremely small size, the magnet and thecapacitor bank are portable and thus the combination with
X-ray diffractometer becomes possible. In fact, the capacitor
bank is as small as a small refrigerator and the storing energy
is only 2 kJ. An advantage of such small system is that amagnetic field sweep rate can be varied up to 1.9 x 103 T/s.
The first experiment wad made to study the giant

KINKEN Research Highlights 2008 IMR

We examined the sweeping speed dependence of thevariant rotation and found that the variant follows the pulsed
field of up to 1.9 x 103 T/s. It indicates that the Fe3Pt has the
sub-msec response to the magnetic fields.
Reference[1] Z. W. Ouyang, Y. H. Matsuda, H. Nojiri, T. Inami, K.
Ohwada, M. Tsubota, T. Sakon, T. Fukuda, and T. Kakeshita,
J. Appl. Phys. 102, 113917 (2007).
Contact toHiroyuki Nojiri (Magnetism Division)e-mail: nojiri@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

3FTFBSDI

GaN Integration on Si via Symmetry-Converted Silicon-on-Insulator
Integration of metals and semiconductors having 3- or 6fold symmetry on device-oriented (i.e., (001)) silicon wafers,
which have 4-fold symmetry, has been a longstanding
challenge. We demonstrate that, by using symmetry-converted(111) silicon-on-insulator, we can integrate wurtzite-structure
gallium nitride, which has 3-fold symmetry, with Si(001). This
technique would enable the commercial integration of wideranging important materials onto Si (001) base wafers.
Silicon serves as the fundamental material for thesemiconductor industry because of its superior processability
in the fabrication of various device structures, and especiallybecause of its ability to form a high-quality oxide. Si(001),
with a square surface lattice, has been used for actual device
fabrication primarily because of the lower interface state
density. A major and important class of materials having 3- or
6-fold symmetries has intrinsic difficulty growing on Si(001),
because the symmetry mismatch at the interface inducesroughness and polycrystallization, factors that seriously
degrade electronic and optoelectronic performance. A
general solution for this long-standing problem is given byutilizing a silicon-on-insulator (SOI) structure where a thinSi(111) layer is bonded to Si(001) via the oxide layer. A
wurtzite GaN film is directly grown on this SOI structure todemonstrate the validity of our approach. [1]The wurtzite structure of GaN has a freedom in its polarity,
which in turn has direct influence on surface structure, growth
process, and its quality. In case of non-polar substrate like Si,
it is especially important to control the nucleation stage toaccomplish mono-polar film with atomically smooth surface.
The GaN films grown on Si(111) substrates with the N-richnucleation condition exhibit 3x3, 6x6, and c(6x12)reconstructions in their scanning tunneling micrographs (STM),
all typical of the surface structures for N-polar GaN films, as
shown in Fig. 1(a). [2] This indicates that the surface is covered
uniformly with N-polar film without forming inversion domain
wall. On the other hand, the nucleation under Ga-rich conditions
results in formation of a rougher surface after the completion of
the film growth, and very flat Ga-fluid areas over Ga-polar
crystal coexisting with the N-polar c(6x12) reconstructedareas can be seen in the STM images (Fig. 1(b)). [2] These
results indicate that N-rich condition in the nucleation layer
growth stage is critical for achieving the growth of mono-polar
GaN films on Si(111) without inversion domain boundaries.
An improved polarity control has been achieved by insertinga ZrB2 buffer layer at the interface. [3]Making a good use of above-mentioned polarity control
technique in GaN growth on Si, we integrated GaN on symmetryconverted SOI structure, as is shown in Fig. 2. [1] Ordered
STM images can be observed for the Si(111)-SOI substrate
and the GaN surface, which assures their microscopic uniformity
as well as their uniform connectivity on the insulating oxide
layer to the end of the specimen. The stability of the orientation
converted SOI on Si(001) observed here would enablegeneral integration of the intriguing class of materials with 3or 6-fold surface symmetries on actual Si devices.

Fig. 2. Schematics of symmetry converted growth of GaN on
Si(001) with superimposed STM images (25 x 25 nm2)
observed in this work. GaN is grown on a 7x7 clean surface
formed on a Si(111) template layer bonded with Si(001)
handle wafer via SOI structure. An ordered 6x6 reconstructedsurface appears as a consequence of excess Ga terminationof grown film, showing uniform formation of a N-polar GaN
film. Removal of the SOI structure by heating at 1300 °C
results in the formation of the Si(001) dimer-row structure,having square-base surface lattice.

References[1] Y. Fujikawa, Y. Yamada-Takamura, G. Yoshikawa, T.
Ono, M. Esashi, P. P. Zhang, M. G. Lagally, and T. Sakurai,
Appl. Phys. Lett. 90, 243107 (2007).
[2] Z. T. Wang, Y. Yamada-Takamura, Y. Fujikawa, T. Sakurai,
and Q. K. Xue, Appl. Phys. Lett. 87, 032110 (2005).
[3] Y. Yamada-Takamura, Z. T. Wang, Y. Fujikawa, T.
Sakurai, Q. K. Xue, J. Tolle, P.-L. Liu, A. V. G. Chizmeshya,J. Kouvetakis, and I. S. T. Tsong, Phys. Rev. Lett. 95, 266105(2005).
Fig. 1. Scematic drawings together with corresponding STM
images (25 x 25 nm2) of wultzite GaN films grown directly on
Si(111) substrates with (a) N-rich and (b) Ga-rich nucleation
conditions. The N-rich condition is critical to achieve uniform
film growth without forming inversion domain walls.



Nano / Micro Structure Controlled Materials

Contact toYasunori Fujikawa (Surface and Interface Research Division)
E-mail: fujika-0@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

Development of Scanning Neutron Diffraction and Local
Structures Investigations in Stainless SteelWe applied a spatial scanning measurement technique toneutron diffraction on a powder diffractometer, and performed
scanning neutron diffraction experiments to investigate local
structures of materials. By this technique, we succeeded inobserving distribution the strain induced martensite phase in a
type-304 stainless steel plate. The comparison with the
distribution of magnetization gives us information of local
structures in internal regions. This method is particularly
important for investigation of industrial materials because it canobserve internal regions of materials,
From industrial view points, assessment of local structures inapplication materials are important; such as grain distribution,
microcrystals, partially deposed second phases, and so on. To
investigate local structures of materials, we have put the scanningneutron diffraction (SND) technique to practical use on the
KINKEN neutron powder diffractometer HERMES [1]. The SND
technique can obtain information on local structures of materials
and its position dependence. Because neutrons have a larger
penetration depth (a few cm) than X-rays, the SND techniqueprovides structural information even in internal regions ofindustrial applications, while X-ray can observe only near surface
regions. Note that neutrons are the unique prove to observe
structures at internal regions of materials with a cm size by nondestructive observation.
Figure 1 shows a schematic view of the alignment of theapparatus of the SND experiments on HERMES, which isbasically same as the scanning X-ray diffraction technique. The
x-y scanning sample stage was controlled by the data acquisition
system of HERMES so that a diffraction pattern at each point on
the sample can be observed automatically. Since 150 3He
detectors are installed on HERMES, one can obtain diffraction
patterns for a short time. The beam size was reduced with thesecond beam narrower with a beam cross section of 2 x 5mm2
for the present experiments of a type-304 stainless steel. The
position resolution on the sample was determined by the beam
size. Note that the positional resolution can be easily changed bychanging the beam size.

alloy austenite phase, and local distribution of the stress induced
martensite phase. Figure 2 shows comparison of distribution ofmagnetization observed by a scanning SQUID magnetometer(SSM) [2], and integrated intensity of the 211 Bragg peak fromthe martensite (α’) phase observed by SND. Note that themartensite phase enhances mechanical hardness, and may
increase the risk of cracks. Thus, assessment of the induced α’phase is important for applications. As shown in Fig.2, the larger
magnetization (circles) was basically disposed around thedeposition of the martensite phase (arrows), indicating that themagnetization is caused by the martensite phase. On the other
hand, some regions with deposition of the martensite phase but
without magnetization were observed exists; for instance
x=10mm, y=25mm (indicated with the question mark). Thismeans that the martensite phase exists in the internal region of
the sample, where SSM cannot observe magnetization at theregions from the surface.
The SND technique is a particularly important probe for
estimating the stress-induced phase, because it non-destructively
observes even the inside of bulk materials. Moreover, by the
SND technique, a single measurement can detect even
unpredictable precipitation of an unknown secondary phase, inwhich the positions of Bragg reflections cannot be predicted.
Moreover, SND on HERMES will be applicable for texture
measurements of materials. The details of the SND techniquehave already been published elsewhere [3]. We also haveapplied for a patent of the SND technique on HERMES.

Fig.2 Distributions of magnetization observed by SSM(left),and integrated intensity of a Bragg peak from the martensite
phase observed by SND of a type-304 stainless steel sample.
In the magnetization (left), red and blue indicate larger
positive and negative magnetizations, respectively. In theneutron intensity (right), red indicate stronger intensity.

Fig.1 Schematic view of the layout of the SND experiments
on HERMES [1]. The distance between the second narrower
and the sample was 15mm, and between the sample andeach detector was 1350mm. The sample could be automatically
moved along x- and y- directions against the fixed beam.
We succeeded in observing positional dependence ofdiffraction patterns from small areas (approximately 2 x 5 mm2)of a type-304 stainless steel plate with a dimension of 19.5 x 50
x 4.7mm3 subjected to a mechanical tensile strain of 90% of the
0.2% offset yield strength. In the SND experiments, we confirmedthat the crystallographic circumstance in the strain-loaded type304 stainless steel is not homogeneous. For instance, we
observed positional deviation in the grain direction of the master-

KINKEN Research Highlights 2008 IMR

References[1] K. Ohoyama, T. Kanouchi, K. Nemoto, M. Ohashi, T. Kajitani
and Y. Yamaguchi, Jpn. J. Appl. Phys. 37, 3319 (1998).
[2] K. Isawa, Superconductivity, Magnetism and Magnets (Nova
Science Publishers, New York, 2006) Chap. 10.
[3] K. Ohoyama, K. Isawa, K. Yamada, Jpn.J. Appl. Phys. 46,
7925 (2007).
Contact toKenji Ohoyama (Neutron and γ-Ray Spectroscopy on Condensed
Matters Division)e-mail: ohoyama@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

3FTFBSDI

Defect Dynamics in Silicon under High Magnetic Field
Dislocations lead to spatial variations in electrical andoptical properties and also degradation of semiconductor
devices. Thus, a great deal of effort is being made inclarifying dislocation properties and dislocation-impurity
interactions under various applied fields as temperature,stress, current, light, and so forth. However, far less is knownof the influences or modifications of dislocation activities bya magnetic field.
Utilization of a high magnetic field has received keen
attention for development of various devices incorporating
magnetic, superconductive, spintronic materials, and also
some semiconductors. Knowledge of the influences against
or modifications of defect properties and structures
accompanying atomic displacements by an external magnetic
field, however, is even now quite limited except for someworks on a concept of magneto-plastic phenomena in impure
ionic crystals, metals, and alloys. In such crystals, dislocations
become mobile, resulting in reduction of their mechanicalstrength. Herein, such phenomena and relevant mechanisms
are discussed in terms of dislocation motion and solid-state
reaction, i.e., the spin-dependent release of dislocations fromsegregated paramagnetic impurity centres/complexes andthe subsequent relaxed motion of dislocations in matrix
crystals, leading to enhancement of plasticity.
Impurity oxygen (O) in semiconductor silicon (Si) is wellknown to have a strong effect of pinning/locking of dislocations
due to preferential segregation and formation of complexes
at elevated temperatures. Thus, it is quite interesting to clarify
how high magnetic field affects on the dislocation-impurity
interaction in the most typical semiconductor Si. Indeed, weinvestigated magnetic effects on dislocation-oxygen impurity
interaction in Czochralski (CZ) grown Si containing oxygen in
a concentration of ~1018 cm-3 under a high magnetic field up
to 10 T at elevated temperatures.
Specimens prepared from an undoped CZ-Si were preannealed at 650˚C for 1 hour and subsequently were treatedat the same temperature under an application of magnetic
field up to 8 T for duration from 15 min to 2 hours in a furnace
installed into a cyclo-cooled superconducting magnet (11TCSM) in the High Field Laboratory for Superconducting
Materials, Institute for Materials Research, Tohoku University.
After removed from the furnace at RT, the specimens were
stressed at 650˚C in a vacuum by means of three-point
bending.
Fig. 1 shows the variation in the critical stress for
dislocation generation against the duration of the magnetic
field application of 1 and 8 T at 650˚C together with that ofcritical stress for dislocation generation in specimens treated
without magnetic field 0T. The dashed green line shows thecritical stress level for dislocation generation in the as-grownspecimen. Under the both intensities of the magnetic field,
the magnitude of the critical stress first increases to a
maximum and then decreases probably to the stress level ofas-grown specimen, while the specimens treated withoutmagnetic field show a little increase in the critical stress,Nano / Micro Structure Controlled Materials

probably due to additional annealing effect. The maximum of
the critical stress ≈ 15 MPa is comparable for the bothintensities of magnetic field. It is seen that the treatment
under application of the higher intensity of magnetic fieldleads to the more prompted variation of the critical stress.
These results imply that the magnetic treatment under ≈ 1 T
for 1 hour and under ≈ 8 T for 15 min at 650˚C most effectively
suppresses dislocation generation from the scratch. However,the prolonged magnetic- treatment leads to disappearance ofsuch effect.
The critical stresses for dislocation generation of floatzone (FZ) grown Si specimens are almost zero, irrespectiveof the magnetic-treatments. This implies that there is noeffect of magnetic treatment in FZ-Si. Thus, it is supposedthat the observed variation in the critical stress is originatingfrom the magnetic field application on dislocation generationin only CZ-Si, i.e., containing O impurity.
It is found that generation of dislocations was effectively
suppressed under certain conditions of the magnetic
treatments. The underlying microscopic nature induced byapplication of high-magnetic field is complicated at present.
The details of dislocation-impurity interaction in Si should beclarified experimentally as well as theoretically, forunderstanding basic natures and kinematics of the spindependent solid-state reaction in a future. Here, it should beemphasized that spin-dependent reaction should be usefulfor improving semiconductor device efficiency throughcontrolling atomic configuration, displacement of defects,
accumulation of harmful metallic impurities.

Fig. 1. Variation in the critical stress for generation of 60˚
dislocations against the duration of the magnetic treatment at
650˚C in Si crystals. The critical stress for dislocation
generation in the as-grown Si crystal is superimposed.
References[1] I. Yonenaga and K. Takahashi, J. Appl. Phys. 101, 053528(2007).
[2] I. Yonenaga, K. Takahashi, T. Taishi, and Y. Ohno, Physica
B 401-402, 148 (2007).
Contact toIchiro Yonenaga (Physics of Crystal Defects Division)e-mail: yonenaga@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

Manifestation of Precipitation Induced Recrystallization During
Plasma-Nitriding of Fe-Cr Alloys
This article reviews precipitation induced recrystallization(PIR) phenomenon in plasma-nitriding of Fe-Cr alloys
discovered recently. When Fe-Cr alloys are nitrided in an
appropriate condition, new ferrite grains composed of (ferrite+ CrN) lamellae holding specific orientation relationship withoriginal ferrite grains are formed in the nitrided zone without
any mechanical-deformation treatment.
Fig. 1 shows a cross-sectional SEM image and thecorresponding ferrite orientation map of the Fe-18Cr alloy
nitrided at 843K for 18ks [1]. The area with bright contrast in
a left hand side corresponds to the nitrided zone in (a). In theorientation map, columnar-shaped ferrite grains which areelongated along the nitriding direction and exhibit orientations
different from that of the original ferrite grains are seen.
Crystallographic analysis reveals that most of newly formed
ferrite grains hold Σ9 coincide site lattice relation, represented
by 38.9º rotation around <110>α, with respect to the original
ferrite grains within a deviation of 5º. Those new ferrite grainsconsist of very fine (ferrite + CrN) lamellar structure. Basedon detailed TEM observation, we proposed that ferrite
formation is caused by specific recrystallization process
induced by volumetric mismatch between CrN continuously
precipitated and the ferrite matrix. We named thisFig. 2 Condition for PIR manifestation in nitriding of Fe-Cr
alloys.

phenomenon as precipitation induced recrystallization (PIR).
Subsequent discontinuous precipitation at the grain boundary
between the recrystallized and original ferrite grain forms(ferrite + CrN) lamellar structure. Fig. 2 summarizes the effectof a Cr content and a nitriding temperature on the appearanceof PIR in plasma-nitrided Fe-Cr alloys [2]. Clearly, PIR tends
to appear at a higher Cr content and a lower nitriding
temperature.
As shown above, PIR has been first clarified. Thisphenomenon has potential for being new surface modification
method to obtain fine-grained structure composed of special
grain boundary without any mechanical-deformation
treatment.
References[1] G. Miyamoto, A. Yonemoto, Y. Tanaka, T. Furuhara, T.
Maki, Acta Mater., 54, 4771 (2006)[2] G. Miyamoto, A. Yonemoto , Y. Tanaka , T. Maki, T.
Furuhara, ISIJ Inter., 47, 1491 (2007)
Contact to
Tadashi Furuhara (High Purity Metallic Materials Division)
E-mail: Furuhara@imr.tohoku.ac.jp

Fig. 1 SEM image and corresponding ferrite orientation mapof a cross section of the Fe–18Cr alloy nitrided at 843 K for
18 ks.

KINKEN Research Highlights 2008 IMR

Nano / Micro Structure Controlled Materials

3FTFBSDI

Reduction of Dislocations in GaN films on AlN/sapphire
Templates Using CrN Nanoislands
Significant reduction of threading dislocations wasachieved in GaN films grown by hydride vapor phase epitaxy(HVPE) on AlN/sapphire templates by employing CrN
nanoislands on AlN layer. We demonstrate that the densityof threading dislocations can be efficiently reduced on AlN/
sapphire templates. The CrN nanoislands play a key role in
reducing threading dislocations by masking and bending
the propagation of dislocations.
GaN based III-nitride materials and devices have been
one of the hottest issues in compound semiconductor
researches and related industries. Recently, studies on highbrightness deep UV LEDs or LDs emitting at wavelengths
shorter than 360 nm have been receiving increasing interests[1]. In order to realize high-brightness deep-UV LEDs, thegrowth of high quality AlGaN/AlN templates is important, so
several methods of high quality buffers such as AlN/AlGaN
superlattices, AlGaN buffer, and GaN/AlN superlattices havebeen reported [2]. For the performance of GaN-based lightemitters, the reduction of dislocation density is crucial for highinternal quantum efficiency, since electron-hole pairs canrecombine through non-radiative processes at dislocations.
However, AlN films grown on sapphire substrates generally
have high density of edge type threading dislocations, whilesmall density of screw type threading dislocations. So there
have been some efforts of reducing the dislocation density by
special growth method such as epitaxial laterally overgrowthof GaN[3] (ELOG), which provides a dislocation density of~106/cm2. As reported by Mukai et al.[4], a UV GaN LED with
7 × 106/cm2 dislocation density has about two fold higher
output power than that with 1 × 1010/cm2 dislocation density.
In this study we report significant reduction of threading
dislocations in GaN films grown by HVPE on AlN/sapphire
templates by employing CrN nano-masks on the AlN. Highquality GaN films with very small twist mosaics as well as
small tilt mosaics have been grown on AlN/sapphire templates,
though the AlN/sapphire templates have small tilt but verylarge twist mosaic. Fig. 1(a) shows a XRD θ-2θ scan of a Cr/
AlN/sapphire sample after nitridation. We can see the (111)
CrN peak in addition to the (0002) AlN, which indicatescomplete nitridation of the Cr layer into CrN. Fig. 1(b) showsa SEM micrograph a nitrided Cr layer, which consists of nanosized triangular CrN hillocks with sizes ranging from 15 to
90nm.

Fig. 1. XRD θ-2θ scan (a) and SEM microsgraph (b) of the
Cr/AlN/Sapphire sample after the nitridation

Nano / Micro Structure Controlled Materials

In order to estimate dislocation densities, plan view TEM
observations were performed, which are shown in Fig. 2(a)and (b) with the <0001> zone axis. Since there is a large
lattice misfit of 17 % between GaN and sapphire, a GaN filmgrown on AlN templates without CrN contains majority ofedge type dislocations which form low angle grain boundaries
with a misangle of 0.7◦ as shown in Fig. 2(a). The estimatedtotal dislocation density is 6.4 x109cm-2. On the other hand, a
GaN film grown with CrN buffer shows a significantly reductionof dislocation densities as shown in Fig. 2(b). The total
dislocation is estimated to be 2.7 x108cm-2. It should be notedthat most of the dislocations observed in our samples wereedge-type dislocations as shown in Fig. 2.

Fig. 2. Plan view bright-field TEM micrographs for the GaN
films on the AlN/sapphire templates without (a) and with (b)the CrN.

In summary, we have demonstrated that the density of
threading dislocations can be effciently reduced by employing
CrN nano-masks on AlN/sapphire templates. The total
dislocation density of a GaN film as evaluated by plan view
TEM observation is reduced down to 2.7 x108cm-2. The CrN
nanoislands play a key role in reducing threading dislocations
by masking the propagation of dislocations as well as bybending the dislocations. The studies on the effects of CrN
nanoislands coverage i.e. the density of nanoislands will be
the next topic for investigation.
References[1] H.Hirayama, T.Yatabe, N. Noguchi, T.Ohashi, and N.
Kamata, Appl. Phys. Lett. 91, 071901 (2007)[2] J. P. Zhang, A. Chitnis, V. Adivarahan, S. Wu, V.Madavilli,
R. Pachipulusu, M. Shatalov, G. Simin, J. W. Yang, and M.
Asif Khan, Appl. Phys. Lett. 81, 4910 (2002).
[3] A. Usui, H. Sunakawa, A. Sakai, and A. A. Yamaguchi,
Jpn. J. Appl. Phys., 36, L899 (1997)[4] T. Mukai, S. Nagahama, N. Iwasa, M. Senoh and T.
Yamada, J. Phys.: Condens. Matter 13 7089 (2001)
Contact toTakafumi Yao (Electronics Materials Division)e-mail: tyao@cir.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

MOVPE Growth Mechanism of High-Quality N-polar GaN
on C-Plane Sapphire
The MOVPE growth mechanism of the high-quality Npolar GaN with a smooth surface grown on c-plane sapphire
has been clarified. The key issues cover sapphire nitridation,
GaN buffer layer annealing, GaN growth temperature, andthe supply of hydrogen and ammonia. These technologies
will pave the way for industrial production of N-polar GaNbased devices.
During recent years, nitride semiconductors have been
explored extensively, especially GaN and its related materials.
Most of the nitride devices are fabricated with Ga-polar GaN
grown on a c-plane sapphire substrate by MOVPE, which isa very good choice for large scale manufacturing. On theother hand, N-polar GaN is also a promising material giving alarge flexibility to device design. However, the growth of Npolar GaN is difficult due to the low lateral growth speed,which leads to pyramidal surfaces. We have solved thisproblem and reported the first success in MOVPE growth ofN-polar GaN with smooth surface [1]. The full width at half
maximum of the X-ray diffraction spectra in ω-2θ and ω scan
is 19 and 35 arcseconds, respectively, which is the smallest
in the reported ones to our knowledge. Here we report the
MOVPE growth mechanism of N-polar GaN [2].
GaN was grown using the two-step MOVPE growthtechnique with a low-temperature buffer layer on a (0001)
sapphire substrate. Just before the growth, the sapphire
substrate was treated with nitridation, followed by the growthof a 20 nm thick GaN buffer layer at 550°C using triethylgallium(TEG) and ammonia. The buffer layer was annealed at
1030°C for the crystallization. Finally, the 1~2μm thick GaN
film was successively grown at 1020°C using trimethylgallium(TMG). To understand the growth mechanism, we investigatedthe effects of nitridation, annealing of a buffer-layer, hightemperature growth, and flow rates of ammonia and hydrogen.
Fig. 1 shows the surface morphologies of the GaN grownwithout and with nitridation of sapphire, observed by adifferential interference contrast microscope. Pyramidal
hilllocks dominate the GaN surface grown without nitridation(Fig. 1(a)), while a smooth surface was obtained withnitridation for 6 min. (Fig. 1(b)). This proper nitridation time isindispensable for acquiring N-polar GaN with a smoothsurface, which is owing to a uniform N-polar AlN layer
formation during the nitridation process.
From the comparison between as-grown and annealed
GaN buffer-layers in observation of the atomic-force
microscope (Fig. 2), the nuclei density is found to decrease
from 11.4 ×108/cm2 to 4.8 ×108/cm2. Using this buffer layer,
the grain in high-temperature grown GaN becomes large and(a) without nitridation(b) with nitridation

Fig. 1 Surface morphology of GaN epitaxial layer.

KINKEN Research Highlights 2008 IMR

as a result the dislocation density small. In addition, theanneal process changes spherical nuclei (Fig. 2(a)) to plain
one (Fig. 2(b)) due to the etching of the GaN facets with highplane indexes in an ambient with hydrogen from the crackingof ammonia during annealing process. The plain nuclei are
believed to enhance the step-flow growth of GaN grown athigh temperature.

(a) as-grown(b) after annealing

Fig. 2 AFM image of GaN buffer.
In the high-temperature growth process, the growthtemperature, the hydrogen and ammonia flow rates play
important roles in GaN quality. The morphologies of the GaN
grown at 1010°C, 1030°C, and 1050°C show a hillock, a flat,
and a rough surface, respectively. At 1010°C, the migration
length of Ga atoms is not enough for the step-flow during
growth. Hence, 3-dimensional growth appears. At 1030°C,
the migration length is enough and the step-flow occurs. At
1050°C, the nitrogen desorbs and GaN re-evaporation
becomes strong. For the ammonia flow rate, an excess flowof ammonia gas forms the surface with hexagonal hillocks
due to the poor migration of Ga-atoms, while its less flow can
not supply enough nitrogen reactant and results in nitrogenvacancies generation. With increasing the hydrogen gas flowrate from 1.8slm to 7.5slm, the surface morphology was
greatly improved as the 3-dimensional nuclei on the terrace
were etched and the step-flow was enhanced.
These results are promising for establishing a standard
for industrial fabrication of N-polar GaN with high quality
grown by MOVPE. It can also be extended to grow N-polar
InN [3]. Using N-polar growth technique, it is possible toimprove the crystal quality of InN, thus making one step
closer to implement an InN-based distributed-feedback laser
for low-cost and high-capacity optical communications
systems.
References[1] Matsuoka, T. Mitate, H. Takahata, S. Mizuno, Y. Uchiyama,
A. Sasaki, M. Yoshimoto, T. Ohnishi and M. Sumiya, phys.
stat. sol. (b) 243, 1446 (2006).
[2] N. Motegi, T. Kimura, M. Nakao, T. Toda, M. Hirata, Y. Liu
and T. Matsuoka, ICNS7, X-7, Las Vegas, USA (2007).
[3] T. Matsuoka, Proceeding of SPIE, "Gallium Nitride
Materials and Devices II", 6473, 1 (2007)
Contact toTakashi Matsuoka (Advanced Electronic Materials Research
Division)e-mail: matsuoka@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

3FTFBSDI

Neutron Holography with Atomic Resolution
Atomic-resolution neutron holography was carried out
utilizing the powder diffractometer in JRR-3M reactor.
Spherically emitted neutrons around a PdH0.78 single crystal
produced by incoherent neutron-proton scattering, which
formed a hologram, were efficiently collected using multiarray 3He detectors. In addition to the hologram data, strong
undulation due to the elastic thermal scattering wasobserved. To exclude this effect, the sample was cooled to
100 K. The thermal scattering was properly removed.
Atomic resolution X-ray and electron holography has
been developed in the last two decades.1-3) Since thesetechniques can visualize 3D local atomic arrangements
around selected elements in solids, they have been attracted
as local structural analysis of dopants and adsorbates.
However, the photoelectron or X-ray holography has a
disadvantage that they cannot apply to hydrogen systembecause hydrogen dose not emit any fluorescent X-rays andphotoelectrons. In order to overcome this difficulty, neutron
holography was proposed in 2001. First demonstration of theneutron holography was realized using a single crystal ofsimpsonite. They measured angular anisotropy of incoherently
scattered neutrons from hydrogen nuclei, and reconstructedsurrounded oxygen nucleus images. In 2004, Cser et al.
applied this technique to a palladium hydride single crystal,and determined the location of hydrogen in the PdH. We alsomeasured efficiently the neutron hologram of a palladium
hydride single crystal using the powder diffrcatometer of
HERMES in JRR-3M reactor, and found the thermal diffuse
scattering effect in the measured hologram.4) However,strong undulation due to the elastic thermal scattering wasobserved. To exclude this effect, the thermal scattering pattern
was calculated and subtracted from the measured data. From
the resulting hologram data, neighboring Pd and H nuclei
images were successfully reconstructed, as shown Fig. 1.
In this study, we used again a palladium hydride single
crystal as the measured sample and evaluated thetemperature dependence of the sample on the thermal
diffuse scattering in the hologram. The sample dimension
was 10 mm in a diameter and 5 mm in a thickness, and it wascoated by Cu not to desorb the hydrogen gas. Wavelength of
the neutron was 0.182 nm.
6
4

y(Å)

2
0

1
00
2

100

2

4

-2

(a)

150
120

θ(º) 90
60
30

(b)

0

90

0

90

180

270

360

180

270

360

φ(º)

150
120

θ(º) 90
60
30

φ(º)

Fig. 2 Holograms of PhH0.78 measured at room temperature (a)and 100 K (b), respectively.
Sample was rotated in the range of 0° ~ 359° with thestep of 1°, where rotation axis was parallel to the incoming
beam direction. Neutrons incoherently scattered by hydrogen
nuclei in the sample were detected by HERMES 150 He3
counters at each sample rotation angle. Polar angle of themeasured hologram defined by the counter ranged from 7° to
157°. Dwelling time at each angle was 480 sec. The average
neutron intensity at each pixel was about 500 counts. To
reduce the effect of thermal diffuse scattering the cryostream
cooler (Oxford: cryostream 70 series) was used for keeping
the sample temperature at 100 K. Figures 2 (a) and (b) show
the 2D angular distributions of the neutron intensities aroundthe sample at room temperature and 100 K, respectively.
The observed data were four-fold symmetrized using sample
crystal symmetry in order to suppress statistical error, andthen they were low-pass filtered. The displayed patterns
exhibit strong spots reflecting the sample crystal symmetry,
due to thermal diffuse scattering. However, the contrast of
the pattern in Fig.2 (b) is weaker than that in Fig. 2 (a).
Especially at polar angle of 80˚, the amplitude of the intensitychange at room temperature was twice as much as that at
100 K. These results revealed that the effect of the thermal
diffuse scattering effect was reduced by the sample cooling.
References[1] K. Hayashi, Advances in Imaging and Electron Physics
140, 119 (2006).
[2] K. Hayashi, T. Matsushita and E. Matsubara, Journal of
the Physical Society of Japan 75, 053601 (2006).
[3] K. Hayashi, T. Hayashi, T. Shishido, E. Matsubara, H. Makino,T. Yao and T. Matsushita, Phys. Rev. B 76, 014119 (2007).
[4] K. Hayashi, K. Ohoyama, S. Orimo, Y. Nakamori, H.
Takahashi and K. Shibata, Japanese Journal of Applied
Physics 47, 2291 (2008).

-4
-6
-6

-4

-2

0

x(Å)

6

Fig. 1 Reconstructed real space-image of (001) plane around
H in PdH0.78.

10Nano / Micro Structure Controlled Materials

Contact toKouichi Hayashi (Chemical Physics of Non-Crystalline
Materials Division)e-mail: khayashi@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

Biofunctional Porous Titanium Filled with Medical Polymer
through Impregnated Monomer Polymerization Technique

Low Young’s modulus similar to that of human bone is
one of the important properties for metallic biomaterials,
because it is advantageous to inhibition of bone atrophy andabsorption due to stress shielding. In order to obtain lower
Young’s modulus, one of the effective ways is to use porous
materials. However, the mechanical properties of porous
materials deteriorate with an increase in porosity. This
deterioration may be attributable to the stress concentrationnear the pores of porous materials. Thus, the inhibition of thestress concentration due to filling up of the pores with certain
materials is likely to improve their mechanical properties. In
such a case, materials with low Young’s modulus should beselected for filling in order to prevent an increase in Young’s
modulus. Furthermore, depending on the type of the filledmaterials, biofunctionalities, which are not intrinsically present
in the metallic biomaterials, can be imparted to the porous
materials in addition to the improvement of mechanicalproperties. In this study, the pores of porous pure titanium(pTi) were filled with a medical polymer, polymethylmethacrylate(PMMA), and then the effects of PMMA filling on the mechanicalproperties of pTi were investigated as a model case [1].
Fig. 1 shows the new process of polymer filling in poresof porous metallic materials developed in this study. A tensile
specimen of pTi with the porosity of 22−50%, which is aporous sintered compact made of gas-atomized pure titanium
powders, was soaked in monomer solution mixed with apolymerization initiator. After removing air babbles in pores
under a reduced pressure, the impregnated monomer waspolymerized by heating at a constant temperature underatmospheric pressure. Then, the extra part of PMMA wasremoved by mechanical machining.
The PMMA filling rate was evaluated by an image analysis
technique, and a fairly high filling rate, greater than 98%, wasobtained. Further, for one of the safety evaluations, theconcentration of residual monomer eluted from PMMA into
water at room temperature was analyzed by liquid
chromatography. It was found that the residual monomer issuppressed by using relatively high concentration ofpolymerization initiator in monomer solution.
Fig. 2 shows the tensile strengths of both pTi and pTi
filled with PMMA. They decrease with an increase in porosity.
However, the tensile strengths of pTi filled with PMMA are
higher than those of pTi in the high porosity range. This result
indicates that the effect of the PMMA filling on the improvementof the tensile strength of pTi is likely to appear in the low
tensile strength range of pTi. Moreover, Young’s moduli of
pTi and pTi filled with PMMA were measured using thestress−strain curve obtained by tensile test. It was observed

KINKEN Research Highlights 2008 IMR

that the PMMA filling hardly affected Young’s modulus of pTi
because Young’s modulus of PMMA is lower than that of pTi.
The process of polymer filling in pores of porous metallic
materials developed in this study is likely to be applied todifferent polymers other than PMMA. By selecting a proper
polymer, biofunctionalities are able to be provided to porous
metallic materials.
MMA solution
mixed with AIBN
powderTensile specimen of pTi

Glass tube
(2) A tensile specimen of pTi
is placed in the glass tube.

(1) MMA solution is poured
into a glass tube.

Evacuation(3) Air babbles in the pores of
the tensile specimen of pTi areremoved in a chamber under areduced pressure.

Water
bath

Heater(4) MMA solution is polymerized in awater bath at a constant temperature.

(6) Extra part of PMMA is removed
from the tensile specimen of pTi
filled with PMMA by mechanicalmachining.

(5) PMMA solidified into the shapeof the glass tube, which includes
the tensile specimen of pTi, istaken out from the glass tube.

Fig. 1 Schematic drawing of process of PMMA filling inpores of porous pure titanium.

Fig. 1 Schematic drawing of process of PMMA filling in poresof porous pure titanium.
250

Tensile strength, σB/MPa

A new process of polymer filling in pores of porous
metallic materials was developed in this study. Such a
polymer filling is expected to improve mechanical properties
and biofunctionalities of porous metallic materials. Therefore,
it was applied to polymethylmethacrylate (PMMA) filling inpores of porous pure titanium (pTi), and tensile strength andYoung’s modulus of pTi filled with PMMA were examined.

pTi
pTi/PMMA

200

150

100

50

0

15

20

25

30

35

40

45

50

55

Porosity, (%)
Fig. 2 Tensile strengths of pTi and pTi filled with
PMMA as a function of porosity.

Fig. 2 Tensile strengths of pTi and pTi filled with PMMA as afunction of porosity.
Reference[1] M. Nakai, M. Niinomi, T. Akahori, Y. Shinozaki, H. Toda,
S. Itsuno, N. Haraguchi, Y. Itoh, T. Ogasawara and T. Onishi,
Proc. Int. Conf. on Advanced Technology in Experimental
Mechanics 2007 (ATEM'07), CD-ROM, (2007).
Contact to
Mitsuo Niinomi (Biomaterials Science Division)e-mail: niinomi@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

11

3FTFBSDI

Fundamental Aspects of Bulk Metallic Glasses
This article reports our recent research results of fundamental
aspects of bulk metallic glasses (BMGs) in terms of (1) a newly
developed molecular dynamics simulation method to analyze
the formation of liquid phase and (2) mechanisms of the stressovershoot phenomenon and non-Newtonian viscous flow
around Tg
A computational method using molecular dynamics
simulations to create a noncrystalline structure from a crystalline
compound related to metallic glass has been proposed based on
the idea that groups of atoms in a compound can be treated as
hypothetical clusters [1]. This method was hinted by an illustrationin an old literature in 1960 [2] as shown in Fig. 1(a), showing thata C6C23 crystalline structure is drawn as an ensemble of clusters
and atoms. When we saw Fig. 1(a) at the first time, we hit on an
idea that "What would be the structure if the clusters are allowed
random rotations around the center of gravity?", which state isillustrated in Fig. 1(b). The computation in terms of pair-distribution
revealed that atomic arrangements expressed by Fig. 1(b) exhibita sharp peaks corresponding to a crystalline phase at aninteratomic distance shorter than 1 nm. However, the atomic
arrangements shown in Fig. 1(b) lose long-range periodicity at along interatomic distance of about 1 nm, which is a characteristicof a crystalline phase. Hence, we understood that the atomic
arrangements shown in Fig. 1 (b) can be an intermediate state totransform from a crystalline state (Fig. 1 (a)) to a liquid phase.
After that we confirmed that atomic arrangements similar to aliquid phase are obtainable by structural relaxation to a stateillustrated in Fig. 1 (b). The following conclusions have beenachieved by the analysis [1]. A computational method to create
noncrystalline structure directly from a crystalline structure, whichis related to metallic glasses, has been proposed. The basis of
this method is that it handles groups of atoms as hypotheticalcluster. To verify its applicability, this scheme was applied to a
C6C23 structure type, which revealed the similarity in the local
atomic arrangements between the computationally relaxed
C6C23 structure and the liquid phase when the clusters are
permitted to rotate before the C6C23 structure relaxes. The
C6Cr23 structure can be regarded as an ensemble of atoms andclusters, which leads to the C6Cr23 structure to be equivalent to
Cu2AlMn structure type in bcc derivative structure family. The
noncrystalline forming-ability of the C6Cr23 structure is due to thelow packing factor of the C6C23 structure in the bcc derivative
structure family, which provides extra freedom by permitting
random cluster-rotations. The significance of the present scheme
is that clusters are used to analyze the local atomic arrangements
in a metallic system and that a noncrystalline structure is
produced from an actual alloy compound with the aid of computer
calculations.
One of the most attractive properties of bulk metallic glass(BMG), whose fracture strength at the ambient temperature isusually in GPa order, is its excellent workability driven by low

Fig. 2 Dependence of true stress and temperature rise of sample
due to the viscous frictional loss at a strain rate of 2.2 ×10 -2 s-1 onthermal conductivity of the die-plates. The initial testing temperature
was set at 690 K.

applied stress in kPa~MPa order at the supercooled liquid state(SCL). We can deform BMGs easily for more than 10,000 % ofelongation and imprint nano-sized patterns on the surface of
BMGs very precisely [3]. These can be performed with thehomogeneous viscous flow. The stress-overshoot phenomenon
is known to be a nonlinear viscoelastic behavior, which brings
the steady-state homogeneous non-Newtonian viscous flow
around the glass transition temperature (Tg). Therefore, it isimportant to understand the mechanism of the stress-overshootphenomenon. In-situ observations of the stress-overshootphenomenon and enthalpy evolution related to the structural
relaxation driven by the stress as well as temperature have been
carried out in a Zr-based BMG via direct measurements of stressas well as sample temperature under compressive constantram-velocity deformations [4]. The results indicate that thestructural relaxation related to the softening phenomenon istriggered by the applied stress, and the increase in the sample
temperature caused by the viscous frictional loss of the nonNewtonian flow also contributes to the softening phenomenon of
the sample. This indicates that sample temperature, thus,
viscosity during the working is affected by the heat conductivityof the surrounding materials, i.e., die, punch for the shaping or
mold for the imprinting. Figure 2 fairly demonstrates thedependence of the steady -state flow stress observed after the
stress-overshoot phenomenon in the Zr-based BMG on the diematerials with different heat conductivity, i.e., WC/Co (71 W/mK), alumina (13 W/m K) and ZrO2 (3 W/m K). This is telling usthat a die (or mold) with lower heat conductivity is more
appropriate for shaping (or imprinting) BMGs.
References[1] A. Takeuchi and A. Inoue, Intermetallics, 16, 283 (2008)[2] W. Hume-Rothery and G.V. Raynor, "The structure of metals
and alloys", The Institute of Metals, Richard Clay and Company,
U.K., 1962.
[3] Y. Saotome, K. Imai, S. Shioda, S. Shimizu, T. Zhang and A.
Inoue, Intermetallics, 10, 1241 (2002).
[4] H. Kato, A. Inoue and H. S. Chen, J. Non-cryst. Solid, 353,
3764 (2007).

Fig. 1 Schematic illustrations showing (a) C6C23 crystalilne structuredrawn as an ensemble of clusters and atoms, and (b) a state where
hypothetical cluster rotations around the center of gravity were allowed
to the crystalline structure.

12

Nano / Micro Structure Controlled Materials

Contact toHidemi Kato (Non-Equilibrium Materials Division)e-mail: nano2000@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

Giant Spin Hall Effect in Gold with a Perpendicular Spin Injector

(a)

js

(b)

je

je
(c)

js

FePt perpendicularspin injector
M

V
Current source

Au Hall crossVoltmeter

RISHE (mΩ)

200

100

-32.0

0
-32.5
-100
-33.0
-10

The SHE is generally considered to be the anomalous
Hall effect (AHE) in a nonmagnetic metal or semiconductor.
The most distinguished point of the SHE is that it does not
require ferromagnetic order. When an unpolarized charge
current flows in a nonmagnetic material having a large spinorbit coupling parameter, the electrons are scattered in thetransverse direction. The scattered directions are opposite
between the up- and down-spin electrons (Fig. 1 (a)). Thismeans that the flow of spins, i.e., a spin current appears,
which is called the direct SHE (DSHE). On the other hand, aspin current induces the transverse charge current (Fig. 1(b)), which is called the inverse SHE (ISHE). The SHE
enables the conversion between charge and spin currents innonmagnetic materials. Although it fascinates us toincorporate this effect into the applications, the complicateddevice structures or the sophisticated measurement
techniques were required to detect the SHE in previous
studies. In addition, the small magnitudes of spin Hall signals
has limited the possibility of device applications of SHE.
We developed a multi-terminal device showing a “giant
SHE” at room temperature [1]. A present multi-terminal
device has a nano-sized Au Hall cross and a FePt
perpendicular spin injector, which is schematically shown in
Fig. 1 (c). Perpendicularly magnetized FePt generates ordetects perpendicularly polarized spin current without
external magnetic field, which simplifies the device structure.
A great advantage of the present multi-terminal device is thatthree kinds of Hall effects, i.e. DSHE, ISHE, and local Hall
effect (LHE) are measured in the same device. We employed

ISHE
LHE

-31.5

RLHE (mΩ)

In the field of "spin-electronics", which takes advantageof the freedom of spins and provides electronic devices withmulti-functionality, the generation and detection of the flowof spins (so-called “spin current”) are keys. Recently, spin
Hall effects (SHEs) in nonmagnetic materials have attracted
much attention because this enables us to generate or
detect the spin current without a ferromagnetic material.
This new type of Hall effect opens the way for the further
advances of spin-electronic devices.

-200
-5

0

H (kOe)510Fig. 2 Resistance of the inverse spin Hall effect (RISHE) as a function of
external magnetic field (H) measured at room temperature. For
comparison, the resistance of the local Hall effect (RLHE) is also shown(red lines)

a non-local spin injection technique to detect DSHE and
ISHE. In the case of the ISHE, a spin-polarized electric
current is injected from FePt to Au, and the spin current flows
in Au due to the spin accumulation. Then, the non-local Hall
voltage induced by the ISHE is detected by the Au Hall cross.
For the DSHE, on the other hand, the FePt spin injector act
as a spin detector measuring the spin-dependent
electrochemical potential at the FePt/Au interface. When an
electric current flows in the Au Hall cross, the spin current isgenerated by the DSHE, and the generated spin current is
detected by the FePt spin injector. An important point is that
the SHE is detected even at zero external magnetic fieldowing to the perpendicular magnetization of FePt.
Figure 2 shows the non-local Hall resistance in the ISHE
geometry (RISHE) as a function of external magnetic field (H).
A clear hysteretic transition of RISHE is observed, and theshape of its hysteresis loop was the same as that of the LHE(red line), which means that the hysteresis loop of RISHE
reflects the magnetization reversal of FePt. A similar
hysteresis of the non-local resistance was also observed in
the DSHE geometry. From the resistance change of the SHE,
the spin Hall angle, which is the ratio of the spin Hall
conductivity to the electric conductivity (σSH/σ), was calculated
to be 0.113. This large spin Hall angle in Au is interpreted by
the skew scattering mechanism. The maximum resistancechange (∆RISHE) at room temperature is obtained to be 2.9mΩ for the device with d = 70 nm, which is significantly larger
than those in previous studies. We believe that the large spin
Hall signal opens the way for the further development of spinelectronic devices.
Reference[1] T. Seki, Y. Hasegawa, S. Mitani, S. Takahashi, H. Imamura,
S. Maekawa, J. Nitta and K. Takanashi, Nature Mater. 7, 125(2008).

Fig. 1 Schematic illustrations of the concepts of (a) direct spin
Hall effect and (b) inverse spin Hall effect. Je and Js are thecharge current and the spin current, respectively. (c) A
schematic illustration of the multi-terminal device consistingContact toKoki Takanashi (Magnetic Materials Division)e-mail: koki@imr.tohoku.ac.jp

of the Au Hall cross and the FePt perpendicular spin injector.

KINKEN Research Highlights 2008 IMR

Nano / Micro Structure Controlled Materials

13

3FTFBSDI

Low-Temperature and High-Speed Deposition of α-alumina
Coatings by Laser Chemical Vapor Deposition
α-Al2O3 (α-alumina) coating is of great interest as wareand heat- resistant coating on cutting tools due to their highhardness and chemical stability at high temperatures.
Conventional thermal chemical vapor deposition (CVD) has
been commercially applied on WC-Co-base cutting tools. In
order to enable to coat α-alumina on TiNi-base cermets, thedecrease of deposition temperature is essential. In thisstudy,α-Al2O3 coating succeeded at a low deposition
temperature by developing laser chemical CVD.
Cemented carbides such as WC-Co have been widelyused as cutting tool material. Nowadays, heat-resistant coatedcemented carbide cutting tools are developing for high-speed
cutting and dry cutting without cooling lubricants. α-Al2O3
coating is crucial to commercialize cermet-based cutting
tools such as TiN-Ni. However, no study has succeeded tocoat α-Al2O3 on cermet-based cutting tools due to too highdeposition temperature causing outward diffusion of Ni.
Chemical vapor deposition (CVD) can produce highquality crystalline Al2O3 coatings. However, a conventional
thermal CVD technique requires high deposition temperature
over 1300 K to obtain α-Al2O3. The deposition rate of α-Al2O3
by conventional thermal CVD is too low around several µm/h.
A thickness of coating should be 20 to 50 µm to elongate thelife-time, and then a high deposition rate has been
demanded.
Last year, the Ministry of Economy, Trade and Industry(METI) has started “Development of alternative materials torare metal project”, where W (tungsten) has nominated as
one of four rare metals. A new hard material for cutting tools
with a small amount or without W should be developed. We
are going to conduct the project to produce high-performance
thick Al2O3 coating on cutting tools.
We have invented a new laser CVD process[1-3] tofabricate α-Al2O3 coating at a low deposition temperature
and a high deposition rate. We have been systematicallyinvestigating the optimum deposition conditions to obtainhigh quality α-Al2O3 coating.
An experimental procedure is as follows: A substrate isplaced on a heating stage in the CVD chamber and pre-heated. Tri-acetyiacetonate aluminum (Al(acac)3) precursor
vapor and O2 gas are supplied to the substrate through adouble nozzle, and Nd:YAG laser (wavelength:1063 nm,
laser power(PL): 0-250 W) is irradiated to the whole substratesurface by expanding a laser beam. The substratetemperature during deposition (Tdep) is measured by athermocouple. The chemical reaction between precursor
vapor and O2 gas is significantly enhanced by plasma formed
around the substrate.

Fig. 2 Surface morphologies of α-Al2O3 prepared at (a) PL =172 W, Tvap = 443 K and (b) PL = 192 W, Tvap = 413 K
Fig.1 shows the effects of Tdep and PL on crystalline
phases of Al2O3 coatings. Crystalline phases have changed
from γ to γ+α to α with increasing Tdep and PL. The lowest
Tdep for α-Al2O3 is 1122 K at PL = 160W, which is 150 K lower
than that of commercial thermal CVD. The deposition rates of
Al2O3 have ranged from 250 to 320µm/h at a precursor
vaporization temperature of 443 K, and are several 10s to
several 100s times higher than those of thermal CVD.
The crystalline orientation of α-Al2O3 has been controlled
by deposition conditions. The (300), (104), (006) and nonoriented α-Al2O3 coatings have been fabricated. The
microstructures of α-Al2O3 coating have well correlated withthe crystalline orientations. Fig.2 demonstrates typicalsurface morphologies of α-Al2O3 coating with different
crystalline orientations. Fig.2 (a) shows a surface SEM imageof (006) oriented α-Al2O3, in which hexagonal plate-likecrystal grains can be observed, while Fig.2 (b) shows (300)oriented rectangular morphology.
References[1] T. Goto, R. Banal, T. Kimura, Surf. Coat. Technol.,201,5776 (2007).
[2] T. Goto, T. Kimura, Thin Solid Films, 515, 46 (2006),.
[3] T. Goto, T. Kimura, Key Engineering Materials, 317-318,
495 (2006).
Contact toTakashi Goto (Multi-functional Materials Science Division)e-mail: goto@imr.tohoku.ac.jp

Fig. 1 Effects of Tdep and PL on crystalline phases

14

Nano / Micro Structure Controlled Materials

IMR KINKEN Research Highlights 2008

3FTFBSDI

Development of Co Based and Ti Based Alloys for Highlyfunctional Structural Materials

Co-Cr-Mo alloys are one of the most useful alloys forbiomedical applications such as dental and orthopedic
implants. However, cast Co-Cr-Mo alloys exhibit low ductility
due to the formation of shrinkage porosity, interdendritic
segregation, and intermetallic inclusion of σ at the grainboundaries. Stabilizing the γ phase (F.C.C.) is thought to beeffective for improving the low ductility in cast Co-Cr-Mo
alloys. Ni, C and N are recognized to stabilize the γ phase in
Co alloys. Among these γ stabilizing elements, the effect of N
addition into Co-Cr-Mo alloy on microstructure and mechanicalproperties is not well understood. Therefore, this work aims
at examining these properties of N-added Co-Cr-Mo alloys
with various Cr contents.
Co-(29~34) mass%Cr-(5~5.8) mass%Mo-(0.01~0.6)
mass%N alloys are prepared using a high-frequency vacuum
induction furnace in this study. Since compositional changes
before and after the casting are negligible, therefore thecomposition will be denoted hereafter by nominal composition
abbreviated as (29~34) Cr. After casting, no apparent change
in dendritic structure among alloys of (29~34) Cr can be
seen. Concerning the evolution in phase constituent, N
addition into Co-Cr-Mo alloys stabilizes γ phase without
precipitation of brittle σ phase even in the enrichment in Cr
content from 29 to 34 mass%. Accompanying with theincrease in Cr content, the solubility of N into γ phaseincreases, thereby, resulting in the significant improvement
in mechanical properties especially for ductility with increasing
the (Cr+N) content (Figure 1)[1].
Thus this study reveals that the Ni-free Co-Cr-Mo alloys
with enriched Cr content up to 34 mass%, modified by N
addition, are suitable as materials for artificial hip and knee
joints and dental implants that are produced on the basis of
the casting process such as investment casting.
And recently, we have presented a new type of Ti alloys
with low Young’s modulus, high strength and high ductility forstructural applications such as springs, screws and flexible
fibers. The developed alloys are composed of fully α’martensite (H.C.P.) and hence their microstructure is stable
during heating up to 600K. This thermal stability of α’ Ti alloyis advantageous for the structural applications as comparedwith conventional β type Ti alloys. However, the characteristic
causing the high ductility in α’ alloy as compared with (α+β)
alloy is uncertain, therefore, this work focuses on theexamination of the deformation behavior in new type α’ Tialloy.
Acicular α’ martensitic structure of quenched Ti alloy
evolves by cold rolling into extremely refined dislocation-cellstructure with an average grain size of 60nm, accompaniedby the development of rolling texture in which the basal plane
is tilted ±40~50 degrees for nominal direction(ND) to

KINKEN Research Highlights 2008 IMRtransverse direction (TD), and rolling direction (RD) is close
to <1 100> orientation [2][3]. Fracture mode after tensile test isdependent on the testing orientation from RD to TD (Figure2). Supposing that the quasi-cleavage plane corresponds tothe primary active slip plane, fracture manner of cold rolled α’Ti alloy can be reasonably explained in terms of the primaryactivation of <a+c> dislocation[2]. This result implies theexcellent ductility in α’ alloy as compared with (α+β) alloy iscaused by the activation of <a+c> slip. Also, this deformation
behavior in α’ Ti alloy suggests the possibility of the material
production (for sheet, bar and wire of Ti alloy) under cold
working (so called α’ processing).
1200

Nominalσ/MPaNominalstrain,stress, σ/MPa

New type of biomedical Co-Cr-Mo-N alloy and structural
α’ type Ti alloy with excellent ductility are proposed in theseresearches. In these developed alloys, high-function isachieved by controlling the alloying element, content,
microstructure and thermo-mechanical condition.

1000
800
600

30Cr

34Cr

29Cr

400
200

0.1

0
Nominalstrain,strain, ε εNominalFig.1 Stress strain curves of cast Co-Cr-Mo-N alloys.

Fig.2 Fractographs of cold rolled (Ti-8%V)-4%Sn alloy after
tensile test. Testing direction: (a)(d, e) RD, (b)(f) 45°from RD,(c)(g) TD.

References[1] S. H. Lee, N. Nomura, A. Chiba, Mater. Trans. 49, 260(2008).
[2] H. Matsumoto, A. Chiba, S. Hanada, Mate. Sci. Eng. A
486, 503 (2008).
[3] H. Matsumoto, S. Watanabe, S.Hanada, Mate. Sci. Eng.
A 448, 39 (2007).
Contact to
Akihiko Chiba (Deformation Processing Devision)e-mail: a.chiba@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

15

3FTFBSDI

Quenched Structure in Tetragonal-Monoclinic Phase Transitionof Non-doped ZrO2 Nano-Layer
Although zirconia (ZrO2) has attractive properties forhigh-k gate dielectrics and buffer layers in semiconductor
devices, its application is not possible due to the tetragonalmonoclinic (t-m) phase transition. In the present study, it
was elucidated that the phase transition was quenched in
the ultra-thin ZrO2 layer deposited on a Si wafer. The layer
had an atomically sharp coherent phase boundary, flat
surface and interface. This means that a way to circumventthe harmful t-m phase transition in ZrO2 has been found.
Yttria stabilized zirconia (YSZ) is an attractive material asa buffer layer for metal-insulator-semiconductor (MIS)-type
Si devices such as ferroelectric gate transistors andsuperconductive devices. One of the critical issues concerning
YSZ is the instability of the threshold voltage of an MIS-field
effect transistor (MIS-FET) due to the large hysteresis of theircapacitance-voltage (C-V) characteristics. The hysteresis is
ascribed to mobile ions or polarization caused by oxygen
vacancies and related defects [1]. Furthermore, the diffusionof doping elements into the Si substrate increases interfacetrapped charges. Use of pure ZrO2 is effective to avoid theC-V hysteresis and interface-trapped charges attributable tothe doped elements such as Y. However, ZrO2 has amonoclinic phase with large spontaneous strain, thus
resulting in a complicated domain (twin) structure with large
roughness at surfaces and interfaces [2]. Therefore,
monoclinic ZrO2 itself cannot be used for the gate dielectrics
and buffer layers. We succeeded in elucidating (1) the film
morphologies of a normal monoclinic phase ZrO2 thin film,and (2) the size effect in the film thickness that controls thetetragonal-monoclinic phase transition of ZrO2 [2].
Un-doped ZrO2 films were deposited on a Si(001) wafer
with thin SiO2 layers by the Pulsed-Laser Deposition (PLD)
technique. The nanostructure and interface of the films were
investigated using high-resolution transmission electron
microscopy (TEM). Figure 1 (a) shows a plan-view image ofthe 17-nm ZrO2 film, where 90° and 180° domain structuresare separated, respectively, with (110) and (100) planes.
Figure 1(b) shows a plan-view image with diffractograms ofeach area of the image of 3-nm ZrO2 film, where theprecipitate (B,C) exists in the matrix (A). Diffractograms,
which are two-dimensional Fourier transformed patterns ofregions A, B, and C, indicate that the matrix is the tetragonalor cubic phase and that the precipitate is the monoclinic
phase. Figure 2 (a) shows a cross-sectional image of the film
observed in the [010] direction of 17-nm ZrO2 film. The
monoclinic phase has a coherent interface with tetragonal
matrix. The (002) planes are tilted about 8.1° from the Si
surface in the inclined area. The interface roughness and thesurface roughness are large due to the lattice tilt of the ZrO2
buffer layer, which causes the local misorientation of theoxide layers grown on the buffer layer. It is also unsuitable forapplication in high-k gate dielectrics. Figure 2 (b) shows across-sectional image of the 3-nm ZrO2 ultra-thin film
projected in the [110] direction. The projected lattice image ofthe ZrO2 layer has orthogonal symmetry, and the lattice16

Nano / Micro Structure Controlled Materials

parameter in the out-of plane direction is larger than that of
the in-plane direction. Therefore, 3-nm ZrO2 ultra-thin film ismainly composed of the tetragonal phase with nano-scaleprecipitates of the monoclinic phase. These results indicatethat the tetragonal-monoclinic phase transition was quenched
due to a type of size effect in the direction of film thickness.
Two possible reasons are considered: the increased
contribution of surface stress induced in the ultra-thin film
compared with thicker films, and compressive stress inducedby the dilative phase transition from the tetragonal to themonoclinic phase.
In summary, we found ultra-thin ZrO2 film with a thicknessof only 3-nm is adequate for buffer layers and gate dielectrics.
The present results show that an ultra-thin film can overcome
the structural issue of the monoclinic phase ZrO2 thin films,
which can contribute to solving the electrical issue of aconventional YSZ buffer layer and gate dielectrics resulting
from dopant cations and oxygen vacancies.

Figure 1 Plan-view image of ZrO2 film of (a) 17nmn and (b)3nm in thickness.

Figure 2 Cross-sectional image of ZrO2 film of (a) 17nmn and(b) 3nm in thickness.
Referrences[1] T. Kiguchi, N. Wakiya, K. Shinozaki, N. Mizutani, Mat.
Res. Soc. Symp. Proc 748, 171 (2003).
[2] T. Kiguchi, N. Wakiya, J. Tanaka, and K. Shinozaki, Mater.
Sci. Eng. B 148, 30 (2008).
Contact toTakanori Kiguchi (Advanced Analysis of Materials Division)e-mail: tkiguchi@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

Coulomb Oscillation of a Proton in a Ni-Nb-Zr-H Glassy Alloywith Multiple Junctions
Electric current-induced voltage oscillation at 500-560
kHz was observed in the current-voltage curves ofnanoscopic size (~0.9-nm) tunnel junctions arranged in a
low-capacitance (~1aF), multiple-junction configuration of(Ni42Nb28Zr30)100-xHx (5.2≤x≤15.2) glassy alloys in thetemperature range of 205 K to 6 K. This behaviour appeared
to be derived from Coulomb oscillation resulting from thetunnelling of individual protons charging and discharging
the vacancy capacitance of Zr-H-□-H-Zr atomic bond arrays
among Zr-tetrahedral clusters, where □ is the vacancy
barrier, termed the free volume, in the glassy alloys [1].

Resistivity (μΩcm)

25000
20000

Temperature (K)Following the discovery of the single-electron transistor[2],many research groups have studied the ‘Coulomb blockade’
effect, in which charge transport through the device occurs inan electron-by-electron manner at low temperatures. Thismeans that a single electron on a metallic island can block
the flow of another electron if the charging energy of theisland greatly exceeds the thermal energy kT. The Coulomb
blockade effects in mesoscopic systems with nanostructures
show nonlinear responses to the externally connected
macroscopic system.
In the present report, in the light of new developments,
we consider the possibility of a nonlinear proton mode, using
a glassy alloy that contains hydrogen. Glassy alloys are
peculiar metallic alloys in that they lack the short-, medium-or
long-range cyclic order of crystalline alloys, on the nanoscale.
Therefore, the metallic glass is considered to be a macroscopic
material with a mesoscopic system that consists ofnanostructures.
For metallic glass that consists of antinomic affinity Zr
and anti-affinity Ni elements for hydrogen[3], a melt-spun
flexible amorphous Ni42Nb28Zr30 membrane with excellent
hydrogen permeability was selected. The Ni42Nb28Zr30
material is a typical metal-metal type alloy that consists offamiliar transition elements. Our interest lies in investigating
the temperature-dependent hydrogen effect of electric
conduction of the (Ni42Nb28Zr30)100-xHx glassy alloys in termsof Coulomb oscillation. However, these are no previous
reports on this subject for glassy alloys with hydrogen.
The electrical resistances of (Ni42Nb28Zr30)1-xHx

250
200
150
100
50
0

heatingcooling

15000

5 10 15 20 25
Content of H (at%)

Zr

(a)
Zr

vacancy

H

RT

(b)

Zr

Zr

Zr

H
Zr

Zr

RS

C

10G

Zr

Zr

electrodeHvacancyH

ZrelectrodeFig.2 Hydrogen tunnelling model (a) for two tetrahedralclusters and blocking oscillation circuit (b). The solid circles
represent is Zr atoms (purple), hydrogen atoms (red),vacancies (yellow).
References[1] M.Fukuhara, A.Kawashima, S.Yamaura and A.Inoue,
Appl.Phys.Lett., 90, 203111 (2007).
[2] M.A.Kastner,P.F.Kwasnick and J.C.Licini, Phys.Rev.
B,36,8015(1987).
[3] M.Fukuhara, A.Kawashima, S.Yamaura and A.Inoue,
Phys.Stat. Sol. (RRL), 1,R50 (2007).

cooling0(x=0~0.22) alloys under 1 mA DC current with positive sign
were measured during cooling and heating runs. We observed
the abnormal resistivity variations (Fig.1). The abnormal
variations in the two runs occurred at restricted hydrogenconcentrations of 5.2 to 15.2 at% H and within the temperature
range of 200 K to 6 K (insert, Fig.1). The wave pattern andpower spectrum showed an AC saw wave with voltage
oscillation of 560 kHz. The desorbed hydrogen from
specimens during DC operation could not be measured by aquadrupole mass spectrometry. Thus it is clear that thisabnormal behaviour is a DC current-induced AC oscillation,
which is associated with the quantum mechanism of solutehydrogen. This oscillation is suggested to arise fromsequential quantised charging of the sample.
Based on the present results, it is possible that theabnormal behaviour of interest is a Coulomb oscillation thatarises from the tunnelling of an individual proton charging
and discharging the capacitance, thereby producing discrete
voltage jumps with 2nd -3rd order of e/C in Ni-Nb-Zr-H glassy
alloys. From the 4-coordination sites that are surrounded in atetrahedral pattern by four Zr atoms, we propose a schematic
representation of a microscopic junction between the Zrtetrahedral clusters in the Ni-Nb-Zr-H glassy alloy ( Fig.2a).
In metallic glasses, an atomic Zr-H-□-H-Zr array corresponds
to a single Coulomb blockade tunnelling junction, where □ isa vacancy barrier with electrostatic capacitance C.

heating

10000
5000
0
0

50

100
150
Temperature (K)200

Fig.1 Abnormal resistivity of the (Ni42Nb28Zr30) 94.8H 5.2 alloy.
Insert: Voltage variation with temperature and hydrogen content.

KINKEN Research Highlights 2008 IMRContact toMikio Fukuhara (Research and Development Project on
Advanced Metallic Glass, Inorganic Materials and Joining
Technology)
e-mail: fukuhara@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

17

3FTFBSDI

Nanowire and Nanotube Made of Metallic Glass
Metallic glass nanowires were spontaneously createdon the fracture surfaces that were produced by a conventional
mechanical test. The presence of the nanowires is directly
related to the one-dimensional meniscus configuration witha small viscosity at high temperatures and to the wide
supercooled liquid region of the metallic glass. The electron
microscopic observations demonstrate the diameters, thelengths, and the amorphous structural states, and the energy
dispersive X-ray reveals the chemical components. In
addition, we found that round ridges are constructed fromnanotubes. The finding of amorphous nanostructures
provides not only fundamental understanding of fracture
processes but also give a new insight into nano-engineering
constructions [1].
Bulk metallic glass (BMG) received attention because
they offer unique mechanical properties such as ultrahighstrength, high hardness, and large elastic strain due to theabsence of crystal slip. There have been extensive studies
regarding the mechanical behavior of the BMGs that is
associated with the nanoscale structural evolution including
the dynamic crack propagation, the nano-crystallization along
shear bands, and the microstructure formation. However, farless attention has been paid for the formation of the metallic
glass nanostructures though they are expected to inherit thesuperior mechanical properties. Nevertheless, the metallic
amorphous or glassy nanowires have not been achieved
experimentally, in particular, formed by mechanical processes
during fracture.
Figure 1(a) represents the SEM image of the fracturesurface of the Zr50Cu40Al10 BMG. The morphology like cellular“vein pattern” consists of the network of round ridges. The
red arrows indicate nanostructures that can be recognizedeven in the low magnification SEM image. The image in Fig.
1(b) shows a nanowire having 100 nm in diameter that has
droplets at each end. To identify the structural phases ofnanowires, we employed the FIB technique and a vein texture
was sliced for TEM observations. To avoid excessive millingof the nanowire, a W(CO)6 gas was exposed to the surface
and the side walls of the sample was milled until the thickness

(a)

(b)

(c)

(d)

Fig. 1 Nanowires on the fracture surfaces of Zr based BMG.
(a) Vein Pattern. (b) 100 nm diameter nanowire. (c) FIB
fabrication for TEM observations. (d) TEM image of nanowire
cross section.

18

Nano / Micro Structure Controlled Materialsof ~0.1 µm, as depicted in Fig. 1(c). Figure 1(d) is the brightfield TEM image showing the cross section of the nanowire thathas 240 nm in diameter. The inset shows the electron diffraction
pattern that is obtained from the center of the nanowire. The
diffused pattern can be recognized, thus we conclude that
there is no trace of the crystallization in the nanowire.
On the fracture surface, there are not only the nanowires
but also the round ridges that configure vein patterns, as seenin Fig. 1(a). To understand structural consequences of theridges, we used the Ar ion milling to prepare an obliquesectional surface, as illustrated in Fig. 2(a). Figure 2(b) is theSEM image showing both the milled and fractured surfaces.
For the left side of the image, the milled surface has featurelessand the EDX obtained from this area represents the bulk
chemical components of the BMG. For the right side, the priorvein features can be recognized. Remarkably, at the boundary
between the milled and fractured surfaces, the ridges aslabeled by the white arrows have a tubular structure. Figure2(c) is the expanded image of the white box in (b). It showsthat the inner diameter is 350 nm. Although the cylindricalfeature is not a perfect round shape, the image provides theexistence of the nanotube on the fracture surface.

Fig. 2 (a) Schematic diagram of the tilt Ar ion milling method.
The Ti shield plate (green) located in front of the fractured
sample (blue). The oblique sectional surface was preparedby removing the uncovered area of the shield with Ar ion
irradiation (pink). (b) The SEM image represents the boundary
between the milled surface (left) and the fracture surface(right). The white arrows indicate the nanotube locations. (c)Higher magnification SEM image of the nanotube wasobtained from the white box in (b). (d) 2D meniscus models.
Reference[1] K. S. Nakayama, Y. Yokoyama, G. Xie, Q. S. Zhang, M. W.
Chen, T. Sakurai and A. Inoue, Nano Lett. 8, 516-519 (2008).
Contact toKoji S. Nakayama (Research and Development Project on
Advanced Metallic Glasses, Inorganic Materials and Joining
Technology; Present: WPI-AIMR)
e-mail: kojisn@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

Real-space Structural Studies of Cu-Zr-Ti Glassy Alloy
Structural features of a glassy Cu60Zr30Ti10 alloy are studied
by using the real-space pair distribution and radial distributionfunctions. A certain degree of medium-range order in this alloy
maintains up to about 2 nm distance. The interatomic distances
closely correspond to those of oC68 Cu10Zr7 compound whichallows obtaining high relative density of the glassy alloy.
Cu60Zr30Ti10 bulk glassy alloy exhibits a high mechanicalstrength exceeding 2 GPa. In the present work we study the
structure of the Cu60Zr30Ti10 alloy by using real-space distributionfunctions. Synchrotron radiation X-ray diffraction in transmission
was carried out using a high energy monochromatic beam of theEuropean Synchrotron Radiation Facility equipped with a nitrogencooled double-silicon monochromator. High-resolution transmission
electron microscopy (HRTEM) investigation was also carried out.
The interference function was calculated from the XRD
intensity profile after the necessary corrections while pair
distribution PDF (r) and radial distribution RDF(r) functions wereobtained by Fourier Transformation (Fig. 1).

According to Fig. 2 detectable medium-range order maintains
until about 2 nm of the interatomic distance independently on thetruncation distance of the Fourier Transformation in the reciprocal
space. It allows us to anticipate that at the maximum Q value of 160
nm-1 used there is no significant influence of the truncation value ofthe Fourier Transformation of the interference function on the shapeof the PDF(r) function and the observed medium-range order until
about 2 nm is rather true. Such a high degree of the medium rangeorder is well consistent with recent models which predict that metallic
glasses are not random packing of atoms but dense packing ofclusters. It is also deduced from the small volume difference between
glassy and the corresponding crystalline phases.

Fig. 2. Two parts of the HRTEM image in Fig. 1: (a) about 16 nm2and (b) about 4 nm2 in size. Area (b) is a part of (a). (c) and (d) are
the corresponding FFT images.
r,Fig. 1 PDF (r) and RDF(r). The insert – HRTEM image
This is a Cu-rich alloy, the lowest possible Goldschmidt bondlength observed in Cu-Cu pair of 0.256 nm is also reflected in theposition of the main maximum. However, taking into account large
negative mixing enthalpy of the Cu-Zr interatomic pair one canassume that this atomic pair corresponds to the most probableinteratomic distance (left shoulder of the PDF function) in thepresent alloy. The maximum of the left shoulder (0.273 nm)
corresponds to somewhat lower value than the Cu-Zr interatomic
distance (0.289 nm) predicted by the Goldschmidt atomic radii
owing to negative mixing enthalpy in this atomic pair.
Such effect is quite common in the intermetallic compounds. oC68
Cu10Zr7 phase is a stable phase forming in the Cu60Zr30Ti10 alloy after
crystallization of the glassy phase. It is isomorphous to oC68 Ni10Zr7 one,has close lattice parameters and the Goldschmidt atomic radius of Ni is
close to that of Cu. In some clusters this phase shows Ni-Zr distances as
small as 0.238 nm and typical Ni-Zr interatomic distances are about 0.260.28 nm which is less than 0.289 nm predicted by the Goldschmidt atomic
radii but corresponds well to the PDF(r) peak position.
Different Cu-Cu and Cu-Zr interatomic distances observed in
the oC68 Cu10Zr7 compound varying in wide ranges from about
0.24 to 0.31 nm for Cu-Zr pair and from about 0.22 to 0.31 for CuCu pair explain the shape of the main PDF(r) maximum (Fig. 1)
responsible for the Cu-Cu, Cu-Ti and Cu-Zr atomic pairs. It makes
clear why partial distances could not be resolved as the shoulders
and disproves the applicability of the Goldschmidt atomic radii
approach for complex glassy alloys. The composition of the oC68
Cu10Zr7 phase is different from that of the Cu60Zr30Ti10 alloy but theCu10(Zr,Ti)7 phase is the main structure constituent in this alloy
after complete crystallization and can be used as a good
approximant to explain structure features of this glassy alloy.

KINKEN Research Highlights 2008 IMR

A question may arise if the medium-range order maintains
until about 2 nm why no clearly ordered zones are seen in the
HRTEM image in Fig. 1. If one takes a Fast Fourier Transformation(FFT) of two parts of the HRTEM image like shown in Fig. 2 one
can see that although the FFT image shows a diffuse halo only
from the area of about 16 nm2 it has some features if taken fromthe area of about 4 nm2. At the same time an icosahedral-type
medium range order which is expected to be a dominant local
configuration in liquids and metallic glasses may be a good
explanation why no crystal-like atomic arrays are observed in the
HRTEM images (Figs. 1 and 2) of these glasses [1]. This idea is ina good consistency with the formation of the nanoscale icosahedralphase observed in a Cu-Zr-Ti-Pd alloy on devitrification [ 2,3 ].
Thus, the structure of the Cu60Zr30Ti10 alloy was studied byreal-space distribution functions. The influence of the truncation
distance of Fourier Transformation procedure is also studied. It issuggested that the medium-range order in this alloy maintains up
to about 2 nm distance. It is also shown that the interatomic
distances correspond to those of oC68 Cu10Zr7 compound andillustrated inapplicability of Goldschmidt atomic radii approach to
the length of Cu-Zr atomic pair in the glassy phase.
References[1] S. Ranganathan and A. Inoue: Acta Materialia 54 3647 (2006).
[2] D. V. Louzguine and A. Inoue: Scripta Mater. 48, 1325 (2003).
[3] D. V. Louzguine, A. R. Yavari and A. Inoue: Phil. Mag. 83, 2989
(2003).
Contact toDmitri V. Louzguine (R&D Project on Advanced Metallic Glasses,
Inorganic Materials and Joining Technology; WPI AIMR. )e-mail: dml@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

19

3FTFBSDI

Synthesis of New Cu-Zr-Based Bulk Glassy Alloys with Unusual
Glass-Forming Ability
The eutectic Cu36Zr48Ag8Al8 alloy exhibited unusualglass-forming ability (GFA), the glassy samples withdiameters of over 20 mm were fabricated by Cu mold casting
and water-quenching [1,2]. The GFA was further enhancedby addition of a small amount of Ni or Pd [3].
The thermal stability of supercooled liquid and GFA of CuZr-Ag alloys remarkably increased by addition of Al. Figure 1
shows the composition dependence of the critical sample
diameters (dc) for the quaternary Cu-Zr-Ag-Al glassy alloys.
The dc above 15.0 mm are obtained by Cu mold casting in awide composition range of 32 to 42 at% Cu, 42 to 50 at% Zr,
5 to 10 at% Ag, and 5 to 8 at% Al. The quaternary eutectic
Cu36Zr48Ag8Al8 alloy exhibits the largest supercooled liquid
region（ΔTx ）of 108 K and the best GFA. Its dc reaches 25
mm.
Figure 2 (a) shows the outer shape and surface appearanceof the Cu36Zr48Ag8Al8 alloy samples with diameters of 16 and 20
mm produced by water-quenching in no flux melting condition.
The surfaces did not exhibit a metallic luster because the surface
reacted slightly with water or quartz tube. However, all thesurfaces were smooth and neither protuberances nor cavities
were observed. The XRD patterns, which were taken from thecentral region of the cross-section of the rods, consisted of only
broad peaks, indicating that a single glassy phase was formed(see Fig. 2(c)). A full eutectic structure consisting of CuZr2,
AlCu2Zr, AgZr2 and other unknown phases are formed forwater-quenched Cu36Zr48Ag8Al8 samples with a diameter of
25 mm (see Fig. 2(b, c)). The eutectic structure is very
uniform and fine. The average spacing of the eutectic
structure in the central region is less than 500 nm. These
results suggest that solidification of the eutectic structure
occurs via homogeneous nucleation and nuclei growth at alimited rate. This solidified mode of the eutectic alloy seems
to promote glass-formation.
The addition of a small amount of Ni or Pd into the
Cu36Zr48Ag8Al8 alloy is effective for further increase in thethermal stability and GFA. Figure 3 shows the outer shapeand XRD patterns of the as-cast Cu32Zr48Ag8Al8Ni4 and
Cu34Zr48Ag8Al8Pd2 alloy samples with a diameter of 30 mm.
The surface of the samples exhibits a mirror-like luster. The
XRD patterns consist of only broad peaks, indicating that the
samples are in a glassy state.

Fig. 2 (a) Outer shape of water-quenched Cu36Zr48Ag8Al8 alloy
samples with diameters of 16 and 20 mm. (b) Cross-sectional
structure in the central region of water-quenched alloy sample
with a diameter of 25 mm. (c) XRD patterns of the water-quenchedalloy samples with diameters of 16, 20 and 25 mm. The data ofthe melt-spun ribbon are also shown for comparison.

Fig. 3 ((a) and (b)) Outer shape and (c) XRD patterns of ascast Cu32Zr48Ag8Al8Ni4 and Cu34Zr48Ag8Al8Pd2 alloy samples
with a diameter of 30 mm.
Compared to other bulk glassy alloys with critical
diameters above 20 mm, the present alloys can be easily
prepared, and exhibit simultaneously large ΔTx, high corrosion
resistance and excellent mechanical properties in conjunction
with a lower alloy cost, and thus may enable usefulengineering applications in the near future.
References[1] Q. Zhang, W. Zhang, and A. Inoue, Mater. Trans., 48, 692(2007)[2] W. Zhang, Q. Zhang and A. Inoue, J. Mater. Res., 23, 1452
(2008).
[3] Q. Zhang, W. Zhang and A. Inoue, Mater. Trans., 48, 3031
(2007).

Fig. 1 Composition dependence of the critical sample
diameter (dc) for Cu-Zr-Ag-Al glassy alloys.

20

Nano / Micro Structure Controlled Materials

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

IMR KINKEN Research Highlights 2008

3FTFBSDI

Structural Evolution of Cr-rich Precipitates in a Cu-Cr-Zr Alloy
During Heat Treatment Studied by 3 Dimensional Atom Probe
Cr

Si

Zr

Fe

（b）RA1

Yaxis

-1

5

Xaxis

1nm

X axis（d）RA2 (view from another direction)

region2

1nm

Yaxis

（c）RA2

region1

Zaxis

Xaxis

5nm

Xaxis

Cu–Cr–Zr alloy is expected as potential heat sink material
for the International Thermonuclear Experimental Reactor(ITER) divertor [1]. The main problem to be solved is its
thermal stability, since the heat treatment during joining
procedures of divertor fabrication may cause solution
annealing or reaging of the alloy. Furthermore, the effect of
Zr on the precipitation behavior is not well understood for the
Cu-Cr-Zr system. The question of chemical compositions
and the structures around the precipitates in Cu-Cr-Zr also
remain unresolved. To clarify the structural evolution of
precipitates, we have employed the Laser-LEAP [2].
The test specimens of Cu-0.78 wt% Cr, 0.13 wt% Zr, 0.003
wt% Si, 0.008 wt% Fe were solution annealed (SA) at 960 °C for
3 h and then quenched, followed by prime aging at 460 °C (PA)and reaging at 600 °C for 1(RA1) and 4h (RA2), respectively.
Atom maps of distributions of the solute (Cr and Zr) and
impurity (Si and Fe) atoms in the heat-treated (PA, RA1 andRA2) states of the alloy are shown in Figure 1. Cr-richprecipitates are formed in the PA state and coarsen during
reaging at 600°C. High number density of fine Cr-richprecipitates were observed by Laser-LEAP.
In the RA1 state the Cr-rich precipitates were coarsened
and almost spherical in shape. Zr, Si and Fe atoms segregatedaround the Cr-rich precipitates [Fig. 1(b)].
During the RA2 reaging the spherical precipitates werecoarsened furthermore and Zr, Si and Fe atoms segregatedin the interface regions between the precipitates and thematrix [Fig. 1 (c)].

（a）PA

Yaxis

Structural evolution of nano-precipitates in a Cu-Cr-Zr
alloy was studied after aging by laser assisted local electrode3D atom probe (Laser-LEAP). It clear that after prime aging
at 460 ºC, Cr precipitates enriched with Zr. Further overaging(reaging) at 600 ºC caused segregation Zr and impurities ofSi and Fe around the Cr precipitates.

5nm

Figure 2. Atom maps along the thinnest direction of the
precipitates in the Cu-Cr-Zr alloy. In (d) the projection is done
along the other direction shown by the arrow in (c).

(b)RA1Figure 2 shows detailed atom maps around each Cr-richprecipitate in the reaged states marked by the arrows in Fig.
1. In Fig.2 (a), showing a precipitate in the PA state, Zr, Fe
and Si atoms were found to be slightly segregate on theinterface regions of the Cr-rich core and form an atmosphereof these atoms. In the RA1 state, not only Cr-rich core but
also the Zr and Si rich interface layer grew as shown in Fig. 2(b) The spherical Cr-rich precipitates were sandwiched
between two enriched shell regions with Zr, Si and Fe.
Concentrations of Zr and Si are about 8 at% at the interfaces,
which corresponds to a chemical composition of about
Cu7Cr3ZrSi.
Figure 2 (c) and (d) shows the atom map of the precipitatein the RA2 state. In some cases, the enriched shell formed
only on one side of the Cr core, but almost not on the otherside. The concentrations of Zr, Si and Fe in the surrounding“region 1” outside of the enriched shell are lower than that of
the “region 2” in Fig. 2(c) due to the segregation to theenriched shell. This segregation of Zr and Si on one side
seems to be affected by the strain field around the precipitates.
In Cu-Cr-Zr systems, one would expect the incoherency
between the Cr-rich precipitates and the matrix to increase
with growth of the Cr-rich precipitates. If so, Zr and impurities ofSi and Fe might segregate on the interfaces between the Cr-richprecipitates and the matrix and then, form an enriched shell on
the other side which relaxes the strain induced by the misfit
between the BCC Cr-rich precipitates and the FCC matrix.

(c)RA2

References[1] M. Hatakeyama, H. Watanabe, M. Akiba and N. Yoshida,
J. Nucl. Mater. 307-311, 444 (2002).
[2] M. Hatakeyama, T. Toyama, Y. Nagai and M. Hasegawa,M. Eldrup and B. N. Singh, Mater. Trans. 518-521, 49 (2008).

X

Cr

Zr

Si

Fe

Z
(a)PA

10nm

Figure 1. Atom maps of the solute and impurities in theheat-treated states of the Cu-Cr-Zr alloy. The atom mapsof the precipitates marked by the arrows are enlarged inFig. 2.

KINKEN Research Highlights 2008 IMRContact
Masahiko Hatakeyama (International Research Center for
Nuclear Material Sciense)e-mail: hatake@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

21

3FTFBSDI

Ductility Improvement of a Zr-Based Bulk Glassy Alloyby Cold Rolling
New ductile bulk glassy alloy (BGA) with macroscopic
tensile elongation has been developed by control of preintroduced shear band structure with cold rolling. Coldrolled structure is characterized by the lamellar structurewith combination of deformed shear band regions and nondeformed block regions. Since the deformed shear band
region can be deformed by lower shear stress than that ofnon-deformed region, cold rolled bulk glassy alloy candeform homogeneously in macroscopic. By the control ofshear band structure, we can obtain ductile BGAs withmacroscopic tensile elongation at room temperature.
Cold rolling of BGA is usually required sufficient ductility
to introduce high-dense shear bands even at a few percent of
reduction ratio. Insufficient ductility brings about accidentallythe serious fracture during the cold rolling. Therefore, wehave to determine the estimation method of the maximum
reduction ratio of cold roll [1]. As a result, the density of castZr-based glassy alloy has a good linear relation with maximum
reduction ratio of cold rolled samples, as shown in Fig. 1.
The density is useful information to estimate the ductility of
the cast Zr-based glassy alloy.

Fig.1 The relationship between the maximum reduction ratio ofcold roll and density of cast Zr50Cu30Ni10Al10 glassy alloys.
Furthermore, the shear band structure of cold rolled
glassy alloy is effective to enhance the ductility and toughness[2], sometimes, which factors are considered as drawbacks
in glassy alloys. Especially, little tensile elongation is usuallyseen in glassy alloys at room temperature, whereas thesuperior elongation can be seen in compression test ofductile glassy alloys. However, the tensile plastic elongation
at room temperature can be obtained by the cold rolling, asshown in Fig. 2. The maximum plastic strain was seen in
10 % cold rolled specimen as about 0.3 % [3]. Figures 3 (a-f)
show the over view, magnified fractured surface and side
images of the same specimens in Fig. 2. A number of finesteps were seen on the fractured surface (Fig. 3(e)) of thecold rolled glassy alloy, which suggests the reoperation ofpre-introduced shear band during the tensile test [4]. This
fact means the co-operative shear bands movement
originates to the tensile plastic elongation without apparent
work softening. In conclude, co-operate movement of anumber of pre-introduced shear bands, whose morphology

22

Nano / Micro Structure Controlled Materials

was controlled to be crossed to each other, reveals thedistinct tensile plastic elongation.
As the result, the composite-like structural control by coldrolling is a useful process to enhance the toughness andductility of glassy alloys. The intrinsic ductility of the Zr-based
bulk glassy alloy is extremely high [5], that phenomenon canbe enhanced much more by the control of shear band
structure with the cold rolling process.

Fig.2 Tensile stress and strain curves of cold rolled
Zr50Cu30Ni10Al10 glassy alloys with reduction ratios of 0 and
10 %.

Fig.3 Tensile fractured surface SEM images of non rolled(a~c) and 10% cold rolled (d~f) Zr50Cu30Ni10Al10 glassy
alloys.
References[1] Y. Yokoyama, Y. Akeno, T. Yamasaki, P. K. Liaw, R. A.
Buchanan, and A. Inoue Mater. Trans., 46, No12, 2755 (2005).
[2] Y. Yokoyama, K. Yamano, K. Fukaura, H. Sunada and A.
Inoue, Mater. Trans., 42, 623 (2001).
[3] Y. Yokoyama, K Inoue and, K. Fukaura, Mater. Trans., 43,
3199 (2002).
[4] Y. Yokoyama, The World of Bulk Metallic Glasses andtheir Composites, ISBN:978-81-308-0196-4 Edited by Can
Fang, pp.141 (2007).
[5] N. Yoshida, K. Fujita, Y. Yokoyama, H. Kimura and A.
Inoue, J. Jpn. Inst. Metals, 71, 730 (2007).
Contact toYoshihiko Yokoyama (Advanced Research Center of Metallic
Glasses)e-mail: yy@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

Superhydrophilicity and Photocatalytic Activity of TiO2 Preparedby Anodic Oxidation in Sulfuric Acid Aqueous Solution(a)Intensity / arbitary units

TiO2 photocatalyst has been practically applied and has
potentially broad applications in quantitative decomposition
or sterilization of chemical substances like endocrine
disruptors, injurious volatile compounds, and organic
pollutants in water and air. Furthermore, the self-cleaning
property of TiO2 due to superhydrophilicity is useful forphotodegrading various noxious or malodorous chemicals.

UHV under UV
exposure to Airhydroxyl groups

Intensity / arbitary units

0

Anodized TiO2 oxide on titanium and its alloys has several
advantages for controlling the photocatalytic activities by
doping through controlling the electrolyte composition,
forming stable oxide with adhesion strength at the interface,
fabricating porous microstructure with tubular pores and cost
performance. The electrochemical parameters such as
conversion potential and current density could control the
microstructure and constituent phase of the oxide, andinfluences the photocatalytic activities and superhydrophilicity.
High crystallinity of anatase and/or rutile is expected to
improve photocatalytic activities because it reduces a numberof re-combination sites of electron and hole. An increase of
potential and current density generates a sparking due toelectric arcs through dielectric breakdown and concurrently
porous morphology is produced.
The effects of sulphuric acid concentration on the anatase
phase formation and superhydrophilicity revealed that anataseformation is promoted either by an increase in conversion
potential or by annealing and that superhydrophilicity is also
improved by the above [1]. The mechanism of superhydrophilicity
was due to a newly formed oxygen defect; this oxygen defect isformed after the surface oxygen traps two holes that are
generated by the photocatalysis reaction as shown in Fig. 1 [2].

before illumination
exposure to O2

540

535

530

525

Binding energy / eV(b) hydroxy groups

520

before illumination
exposure to O2
UHV under UV
exposure to Air

540

535

530

525

Binding energy / eV520

Fig.2 O 1s XPS spectrum of the as-anodized (a) andannealed (b) oxides: before UV illumination, exposure to O2
for 5 min under illumination, illumination for 30 min andexposure to air for 1 h followed by illumination.
Further increase of the sulfuric acid concentration in theelectrolyte generates nano-sized pores (Fig.3), and the
predominant oxide phase is replaced from anatase to rutile.
The bleaching of methylene blue over the anodized oxide
increases with the concentration of sulfuric acid and byannealing. Studies have been under progress [3].

Fig.3 Microstructure of the oxide anodized in 0.1 M thesulfuric acid concentration in the electrolyte.
Fig.1 Proposed mechanism of the photoinduced conversionof TiO2 [2].
All the oxides show an asymmetric O 1s XPS spectrum,
and a shoulder band on the higher binding energy side of themain peak originated from the adsorption of hydroxyl groups isobserved (Fig. 2). The intensity of the shoulder in the annealedoxide (b) is higher than that in the as-anodized oxide (a).
Moreover, the intensity of the shoulder peak increases with UV
illumination when it is exposed to oxygen under UV illumination,and it is to air and subsequently subjected to UV illumination,
implying that the annealed oxide interacts strongly with thehydroxyl groups. The results suggest that the annealed oxideinteracts strongly with the hydroxyl groups.

KINKEN Research Highlights 2008 IMR

References[1] N. Masahashi, S. Semboshi, N. Ohtsu, and M. Oku, to be
published in Thin Solids Film (2008).
[2] N. Sakai, A. Fujishima, T. Watanabe, K. Hashimoto, J.
Phys. Chem. B 105, 3023 (2001).
[3] N. Masahashi, Y. Mizukoshi, S. Semboshi and N. Ohtsu,
unpublished work (2008).
Contact toNaoya Masahashi (Osaka Center for Industrial Materials
Research)e-mail: masahasi@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

23

3FTFBSDI

Nanoimprint of Glassy Alloy for Interference
Optical Components
Glassy Alloys (Metallic glasses) are useful for fabricating
nanodevices due to their viscosity in the supercooled liquidstate and homogeneity on nanometer scales. In this study,
we developed interference optical components of thematerial such as RGB diffraction grating, a hologram andnon-reflecting surface by nanoimprint (superplastic
nanoforming) with dies fabricated by LIGA process and
FIB-nanomachining.
Interference optical components have been fabricated bymolding polymer materials because of their precision
formability and productivity. However, the materials are notuseable at elevated temperatures, and are poorer inmechanical properties and corrosion resistance than metallic
materials. Glassy alloys (metallic glasses) are fundamentallyamorphous metals and exhibit nanoimprintability due to the
Newtonian viscous flow in the supercooled liquid temperature
range and homogeneity on nanometer scales. Furthermore,
the materials reveal superior mechanical and functional
properties, since the materials are promising nanomaterials
for mass producing various nanodevices [1]. This study used
Pt-based metallic glass with a glass transition temperature of
502 K for fabricating reflective RGB diffraction grating by
nanoimprint [2]. A diffraction grating has a fine convexconcave pattern on the surface and is fabricated bynanoimprint with a die. The die was made by Ni-electroformingwith master models fabricated by photolithography with oneshot exposure of the interference fringe. An interference
optical exposure system has great advantages over the
fabrication of nanometer-scale dies by large-area lithography.
The light source for the optical system was a solid-state laser(375 nm wavelength), and the fringe pattern interval was
determined by the system arrangement. Nanoimprint was
carried out with the die. The working temperature of Ptbased metallic glass was 550 K under a compressive stressof 10 MPa. Figure 1 shows the photographs of (a)the die,(b)nanoimprinted RGB diffraction grating and (c),(d),(e)SEM
microphotographs of the surface of the grating. The results
demonstrated the advantages of the nanoimprint techniqueon glassy alloys and for mass production of opticalcomponents.
Similar nanoimprint was performed for
fabricating reflection holograms. A hologram sticker wasused as the master pattern for fabricating a die. The master
pattern was replicated onto the three-dimensional surface
microstructure of the die by Ni-electroforming.
Another application of nanoimprint of glassy alloy is a
non-refrecting surface [3]. The die was fabricated bynanomachining with focused-ion beam (FIB). FIB-machining
characteristics of glassy carbon and Zr-based metallic glass
have been studied and are useful for fabricating nano-scaledies because of the isotropic homogeneity of their amorphous
structures. Figure 2 (a) illustrates the nanoimprinted surfaceof Pt-based metallic glass, observed with an optical
microscope. We can observe four small areas where there
are differences in brightness. Figure 2(b),(c),(d), and (e)present SEM microphotographs of the nano-structured

24

Nano / Micro Structure Controlled Materials

surfaces associated with each small area. The reflectanceratio of these nano-structured surfaces has been measured
at λ=533nm and is represented in the figure. The ratio for thesurface of Fig.2(c) is only 1 %, since the surface is almost
non-reflecting surface.
These nano-imprinted, opticalcomponents of metallic glasses are useful for numerous
applications and under various conditions.

Figure 1 Photographs of (a)electroformed Ni-die, (b)fabricated
RGB diffraction grating of Pt-based glassy alloy (metallic glass)by nanoimprint with the die(a), and (c)(d)(e)SEM microphotographsof nanoimprinted surface of the grating(b)Figure 2 (a)Imprinted surface of Pt-based glassy alloy(metallic
glass) observed with an optical microscope and (b)(c)(d)(e)SEM
microphotographs of nanostructured surface imprinted on thesurface(a). Reflectance ratio was measured at λ=533 nm.

References[1] P.Sharma, N.Kaushik, H.Kimura, Y.Saotome and A.Inoue,
Nanotechnology, 18, 035302 (2007).
[2] Y. Saotome, Y. Fukuda, I. Yamaguchi and A. Inoue, J.
Alloys and Compounds, 434- 435, 97 (2007).
[3] Y.Saotome, S.Okaniwa, H.Kimura, and A. Inoue, J. Mat.
Sci. Forum, 539- 543, 2088 (2007).
Contact toYasunori Saotome(Osaka Center for Industrial Materials
Research)e-mail: saotome@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008

3FTFBSDI

Ultra-large Room-temperature Plasticity ofa Nanocrystalline Metal
Ultra-large room-temperature plasticity up to 200% true
plastic strain is achieved in nanocrystalline (nc) Ni. The lowtemperature, high strain rate and high flow stress, in conjunction
with significant ‘cold’ nanograin growth, demonstrate that theobserved ultra-large plasticity in nc Ni is mainly performed bya grain-boundary-mediated process that is driven by highstresses rather than by thermal diffusion, which is intrinsicallydissimilar to that in traditional superplastic materials
deformed at high temperatures and low strain rates.
Ultra-large plasticity often occurs at elevated temperatures
above 50% of the melting point (TM) of materials with a slow
strain rate (10-5 – 10-4 s-1). Generally, ultra-large plastic flowof a material is associated with its grain boundary (GB)activities and smaller grains promote the ultra-large plastic
flow at lower temperatures and higher strain rates. Thus, asa reasonable extrapolation, room-temperature ultra-large
plasticity of nanocrystalline metals, which possess acharacteristic grain size less than 100nm, has been expectedmany years ago. However, nc metals often fail at a tensile
strain typically less than 2%, a definitive evidence of whether
the ultra-large plastic flow of nc metals can be intrinsicallysustained at room temperature still remains missing.
Compression testing of eletrodeposited nc Ni was
performed by a newly-developed micro-compression
technique. The micro-pillars with nominal diameters ranging
from 3 µm to 8 µm and an aspect ratio of 2:1 were prepared bya focused ion beam (FIB) system. A nanoindentation apparatus,
equipped with a flat-end diamond indenter, was employed tocarry out the room temperature microcompression tests withloading rates ranging from 0.132 mN/s to 13.2 mN/s (thecorresponding strain rates range from 4x10-3 s-1 to 5.4x10-1 s-1during steady plastic flow).
Up to 200% true strain, or 85% loss of its initial height,
was surprisingly obtained in nc Ni without neither material
nor structural failure (Fig. 1). A typical true stress (σ) vs. truestrain (ε) curve of nc Ni strained up to ~100% is shown in Fig.
1b. Steady flow stresses of 2.0 to 2.4 GPa were obtainedwith loading rates ranging from 0.132 mN/s to 13.2 mN/s.
The altering mechanical performances with deformation
amounts, i.e. strain softening up to the deformation of about
60% during the steady plastic flow and followed by strain
hardening, are found to be associated with the deformationinduced nanostructural evolution.
Visible nanograin growth occurs in the specimens(a)Fig. 2. When the grain size increases due to continuous and/or discontinuous grain growth (a) caused by thermal and/or
mechanical activation, yield strength of the material decreases
when grain size is larger than 10nm (b).
deformed in the steady plastic flow stage and nanograins
remain nearly equiaxial with ~40% plastic deformation. The
deformation-induced nanograin growth can lead to enhanced
ductility and obvious strain softening in nanocrystalline metals(Fig. 2). In contrast, the microscopic images taken from thesample deformed in the strain-hardening regime (~140% truestrain) show coarsened and elongated nanograins that
contain a high density of dislocations, stacking faults anddeformation twins, indicating intragranular straining of thecoarsened nanograins and, thus, strain hardening takes place.
GB sliding takes place mainly via a large number of small
local sliding events of atomic clusters that comprises a few orat most a few tens of GB atoms. This stress-enhanced masstransport along GBs, in nature, can lead to ultra-large plastic
flow in the absence of enough thermal assistance by hightemperatures, accompanied by other atomic-scale GB
processes such as nucleation of dislocations. The observed
ultra-large plastic flow of nc Ni is thus nearly rate-independent
since the dominating deformation mechanism, i.e. stressenhanced mobility of GB atoms, mainly depends on stresslevels rather than time in the case of thermally activated GB
diffusion, facilitating the ultra-large plastic flow at high strain
rates. In addition, this collective atoms motion at GBs
predicted by MD simulations fairly coincides with themeasured activation volumes for ultra-large plastic flow anddeformation-induced nanograin growth. Accordingly, thenanograin growth and plastic deformation of nc Ni undercompression most likely take place simultaneously throughstress-enhanced GB motion because both sliding for plastic
deformation and migration for nanograin growth are motivatedby the mobility or diffusivity of GB atoms.
References[1] D. Pan, S. Kuwano, T. Fujita and M.W. Chen, Nano Lett.
7, 2108 (2007).
[2] M. W. Chen and X. Q. Yan, Science 308, 356c (2005).
[3] D. Pan, T.G. Nieh and M.W. Chen, Appl. Phys. Lett. 88,
161922 (2006).

(a)

5 µm
Fig. 1. (a) FIB micrographs of microcompression sample under
200% true strain. The inset is an as-fabricated micropillar. (b)Representative true stress vs. true strain curve of nc nickel
subjected to microcompression.

KINKEN Research Highlights 2008 IMRContact to
Mingwei Chen (International Frontier Center for Advanced
Materials)
e-mail: mwchen@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

25

3FTFBSDI

Spontaneous Aggregation of Pentacene Molecules on SAMtreated SiO2 and Its Influence on FET Characteristics
It has been found that pentacene molecules aggregate
spontaneously on the substrates having lower surface
energy than that of pentacene (001) surface under ultrahighvacuum and ambient temperature conditions. In situ AFMFET measurements have proved that the FET characteristics
are significantly affected by this aggregation.

we have also observed significant degradation of FET
performance even in the case of thicker pentacene films, this
spontaneous aggregation should be one of the major causesof the instability and irreproducibility of pentacene-based
devices.

Pentacene (C14H22, Fig. 1) is one of the most promisingorganic semiconductors for organic field effect transistors(OFETs). In order to improve the FET performance, thesurfaces of gate insulator, such as SiO2, are commonly
modified by the self-assembled monolayers (SAMs), such as
hexamethyldisilazane (HMDS) or octadecyltrichlorosilane(OTS) (Fig. 1(c)). Owing to utilization of these SAMs, theperformance of the pentacene FET has already exceededthat of amorphous Si FET. There remain, however, crucial
problems for practical application, namely, stability andreproducibility of their characteristics. In the present study,
we have found spontaneous aggregation of pentacenemolecules deposited on SAM-treated SiO2 substrates underultrahigh vacuum and ambient temperature conditions (Fig.
2) [1]. In situ and real-time observation of growth and theaggregation of pentacene molecules with low-energy electron
microscopy (LEEM) revealed that the balance of surface
energy between pentacene film and substrate is an origin of
this aggregation. Pentacene films grown on the HMDS- or
OTS-treated surfaces, which have lower surface energy thanthat of pentacene (001) surface, tend to aggregate, while
those grown on the clean-SiO2 surfaces, which have higher
surface energy, do not aggregate (Fig. 2(a-c)). We constructed
an in situ atomic force microscopy (AFM)-FET measurement
system to find that the FET mobility is significantly affectedby the aggregation of pentacene molecules (Fig. 2(d)). Since
Fig. 2. Morphological change of 1.3 ML-thick pentacene films
grown on (a) clean SiO2, (b) HMDS, and (c) OTS observedwith LEEM under UHV and ambient temperature conditions.
Elapsed time after deposition is indicated at the top of eachfigure. Thickness of pentacene film in each region is indicatedin the figures. (d) Change of morphology and FET characteristics
of 1.3 ML-thick pentacene film measured with in situ AFM-FET
system under high vacuum and ambient temperature
conditions. The size of AFM images is 4.8x4.8 µm2.

Reference[1] G. Yoshikawa, J. T. Sadowski, A. Al-Mahboob, Y.
Fujikawa, T. Sakurai, Y. Tsuruma, S. Ikeda and K. Saiki,
Appl. Phys. Lett. 90, 251906 (2007).

Fig. 1. (a) Molecular structure of pentacene. (b) Crystal
structure of the pentacene thin film. (c) Schematics of thesample used for in situ AFM-FET measurements.

26

Nano / Micro Structure Controlled Materials

Contact toGenki Yoshikawa (SAKURAI Project, International FrontierCenter for Advanced Materials)
e-mail: yoshik-0@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2008