3FTFBSDI

Strained Surface Structure in the Growth of Quantum Dots
The formation of Ge self-organized quantum dots (huts) on
the Si(001)-2x1 surface has attracted much attention due to its
status as the premier model system for investigating straininduced semiconductor nanostructure growth. Our findings on
this strained surface structure imply that both the structure andthe strain of semiconductor nanostructures may be controlled
through the understanding of their surface properties.
Although the hut structures are known to have a highlyordered structure defined by four Ge{105} facets, little was
known about the structure until our investigation. We foundusing field ion scanning tunneling microscopy (FI-STM) that theGe(105)-1x2 surface structure exhibits significant bias
dependent on its filled or empty states. (Fig. 1(B), Fujikawa etal., Phys. Rev. Lett. 88, 176101 (2002).) With the help of firstprinciple calculations, we proposed a new model consisting ofrebonded steps (RS model). Additionally, direct observations of
the RS structure were achieved by a high-performance noncontact atomic force microscope (AFM), resolving all individual
dangling bonds (Fig. 1(A)) [1]. This surface structure ischaracterized by strong surface strain coming from the rebonded
structure unit, which was found to relax by lattice mismatchstrain applied from the Si substrate. Distorted bonding
configuration of the surface atoms, which resulted from strong
charge transfer, is found to also be responsible for surface
strain relaxation. This charge transfer was documented at an
atomic level by Kelvin-force microscopy (KFM) [1] and scanningtunneling spectroscopy (STS) [2]. We successfully obtained an
atomically-resolved quantitative surface potential map by KFM
for the first time, which agrees well with the theoretical
calculations of the surface electronic states. The STS results
indicate that the surface charge transfer is explained by theuneven distribution of surface bands, rather than the separated
electron filling of localized dangling bonds.
The RS model reveals the importance of tensile surface

Fig. 1. (A) Non-contact AFM and (B) STM images of theGe(105)-1x2 surface. The surface unit cell of Ge{105}-1x2
reconstruction with a size of 10.86 Å x 13.84 Å is described asa broken rectangle, and the surface atomic structure of the RS
model is superimposed on the AFM image. The dangling bonds
on atoms a, b and dimer c are clearly resolved as separate
protrusions in our AFM observation. The filled- and empty-state
STM images exhibit distinctive differences, which are wellunderstood by the calculated density of states based on our RS
structure model accompanied with charge transfer from dimer c(marked by circles) to atoms a and b (marked by crosses).



IMR KINKEN Research Highlights 2007

strain on the rebonded steps. For example, hydrogen adsorption
dramatically destabilizes the Ge/Si(105) surface, which can beunderstood by the fact that hydrogen adsorption increases thesurface strain by causing rebonded surface dimers to be placed
at asymmetric positions (Fig. 2) [3]. Our STM observations
reveal that hydrogen adsorption causes defect formations on
the SiGe(105) surface formed on the Si(105) substrates, which
proves the strain enhancement effect described above. This
result explains the previously observed phenomenon ofhydrogen suppressing Ge hut formation on Si(001) (surfactant
effect).

Fig. 2. Simulated atomic structures of (A) clean and (B)hydrogen-terminated Ge/Si(105) surfaces. The solid rectangles
show 1x2 unit cells of the surface structure. Hydrogen atoms(white) change the bonding configuration of the surface atoms
by causing the rebonded dimers (red) to be placed on thesurrounding Ge atoms (blue) in a stressed manner.
Our recent observations of Si island growth on Ge(001)reveal that all of the {105}-related features observed in Ge/Si
island formation completely disappear in the oppositely strained
Si/Ge islands, while some similarity remains between themreflecting the symmetric response of the bulk strain. This factindicates that the tensily-strained {105} facet structure plays acritical role in structure determination of the strained island,
reacting asymmetrically to the lattice mismatch strain from thesubstrate.
The strained facet structure on quantum dots, establishedby our detailed atomic-scale characterizations using scanningprobe techniques, was found to play an important role in theirstructure determination through its unique response to latticemismatch strain as well as surface conditions such as adsorption.
This unique response would make the structure as well as thestrain of the quantum dots controllable.
References[1] T. Eguchi, Y. Fujikawa, K. Akiyama, T. An, M. Ono, T.
Hashimoto, Y. Morikawa, K. Terakura, T. Sakurai, M. G. Lagally,
and Y. Hasegawa, Phys. Rev. Lett. 93, 266102 (2004).
[2] Y. Fujikawa, T. Sakurai and M. G. Lagally, Appl. Surf. Sci.
252, 5244-5248 (2006).
[3] Y. Fujikawa, T. Nagao, Y. Yamada-Takamura, T. Sakurai, T.
Hashimoto, Y. Morikawa, K. Terakura and M. G. Lagally, Phys.
Rev. Lett. 94, 086105 (2005).

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

Nano / Micro Structure Controlled Materials

3FTFBSDI

Superelasticity in Ti–10V–2Fe–3Al Alloys with Nitrogen Addition
Recently, much attention has been paid to Ni-free βtitanium shape memory/superelastic alloys for biomedical
use. For wider application, however, reduction of production
cost is important. Interstitial elements, which were not usedpreviously due to the deterioration of ductility/toughness,
exhibit large solution strengthening and were recently applied
in industrial development of low-cost titanium alloys. We
reported new superelastic alloys by addition of nitrogen to
Ti-10V-2Fe-3Al which is originally used for airplanes [1,2].

Contact to
Tadashi Furuhara (High Purity Metallic Materials Division)e-mail: furuhara@imr.tohoku.ac.jp

Fig.1 schematically describes the concept of alloy
development. Nitrogen suppresses plastic deformation of β
by increasing the yield strength of β (σβy). Further, it
suppresses β→ω and β→α” martensitic transformations (seethe change of transformation starting temperatures, ωs andMs). As a result, the stress-induced β→α” martensitic
transformation occur in a wider temperature range, and thepotential of superelasticity is increased (highlighted in purple).
In-situ observation of surface relieves during tensile
deformation (Figs. 2(a)-(c)) shows clearly that stress-induced
α” martensite is formed by deformation and reverselytransformed to β matrix by unloading in the alloy containing
0.2mass%N. This is a direct evidence of the occurrence ofsuperelasticity. Fig. 2(d) shows the recovered strain obtainedin cyclic tensile tests. Total elastic strain reaches to 2% withthe addition of nitrogen by increasing pure elastic strain
through strengthening and enhancement of superelasticity.
References[1] T. Furuhara, S. Annaka and T. Maki, J. Mater. Eng.
Perform. 14, 761 (2005).
[2] T. Furuhara, S. Annaka, Y. Tomio and T. Maki, Mater. Sci.
Engng. A, 438-440, 825 (2006).

Fig. 1: Improvement of superelasticity with nitrogen addition.
(a) nitrogen free alloy, (b) nitrogen-added alloy; σSIM : stressrequired to induce the martensitic transformation. As and Af :temperatures where the reverse transformation starts andends, respectively.

Fig. 2: (a)-(c) Optical micrographs showing the formation and reversion of surface relieves during 1.5% tensile deformation andunloading at room temperature in the Ti-10V-2Fe-3Al-0.2N alloy, respectively.; (d) variation of total elastic and superelastic strains
recovered with total applied strain in cyclic tensile tests.

Nano / Micro Structure Controlled Materials

KINKEN Research Highlights 2007 IMR

3FTFBSDI

Effects of Morphology and Doping on the Electronic andStructural Properties of Hydrogenated Silicon Nanowires
Using first principles based calculations, we explore thedifferent avenues of applications of Si-nanowires. We show
the novel reconstructions on the surfaces of the nanowires,interpretation of STM images, sensors capabilities of bioand chemical species and designed p-, n-doped and opticalnanowires. This study has been one of the major contributions
to the area of Si-nanowires and their applications.
Si-nanowires (SiNWs) are considered as the strongest
candidates for nanoelectronics due to their electrical andoptical properties, controlled production [1], and promising
new applications such as molecular scale sensors. These
applications rely heavily on the electronic structure of SiNWs
and require precise control of the surface atomic andelectronic structures. For the rapid development of this field,
advances in the understanding of the electronic structure of
SiNWs and the various possible ways to modify it, are very
important [2, 3].
C

D

2.4

NW1
NW2s
NW3s
NW4s
NW5

Band gap (eV)

B

2.0

1.6

1.2

1.2
1.4
1.6
Mean diameter (nm)1.8

FEnergy (eV)Energy (eV)

E

1.0

X Γ

Γ

Γ

X

X

n-typeΓ

p-typeX

Fig 1. a) Novel reconstructions on the surfaces b) correctinterpretation of STM c) tailor made optical properties d)circuit elements e) nanosensors.
In order to address these issues, we have studied
geometries of five different classes of SiNWs (referred as
NWn (n=1-5)) differing in orientations and surface
morphologies Fig. 2. The band gap changes with theorientation and thickness of the nanowires. Importantly, italso depends strongly upon the surface morphologies.

Energy (eV)
Γ

NW3

NW2

NW1

1.30 eVX Γ

1.34 eVX Γ

1.63 eV

X Γ

NW4

NW5

1.36 eV

1.83 eVX Γ

surface of NW4 nanowires as they involve only (110) facets.
Removal of one H atom from each of the corner Si atoms of
the NW4 nanowire has an effect of creating unpaired
electrons on the surface, which leads to drastic changes inthe band structure. In contrast to this, we find that SiNWs
with (100) facets can easily offset the effects of H deficit on
the surface by dimerization such as in NW2 and NW3. Thisshows that the electronic properties of nanowires with (100)surfaces cannot be easily modified and therefore are poor
candidate for the sensors. Nanowires with (110) or (111)
surfaces are ideal for sensors as their electronic properties
could be modified drastically by changing H concentration.
Due to significant differences in the electronic properties
of NW1 and NW4 by a change in hydrogen concentration, we
considered them as our prototypes and doped them with B
and P to have p- and n-type nanowires, respectively. Theband structures of the doped nanowires show a shift of the
Fermi level towards the valence and conduction band for thep- and n- type SiNW, respectively. Nanowires grown along
different directions and with different surface morphologies
have different dispersions (Fig. 2) in their conduction andvalence bands and this information can be used to design
doped semiconductors with better conduction properties.
In summary, we show these nanowires could be ideal for
applications such as optoelectronic devices, and lasers.
Further we have shown for the first time that H-defects on
nanowires with (110) facets modify their electronic properties
significantly so that they could be used as sensors and thatthe behavior of n- or p- type doping in hydrogenated SiNWs
is very similar to bulk or oxidized thicker nanowires, therefore
establishing them to be the smallest possible building blocks
for silicon based nanoscale devices [2, 3].
References[1] A. M. Morales and C. M. Lieber, Science 279, 208(1998).
[2] A. K. Singh, V. Kumar, R. Note and Y. Kawazoe, Nano
Lett. 5, 2302 (2005).
[3] A. K. Singh, V. Kumar, R. Note and Y. Kawazoe, Nano
Lett. 6, 920 (2006).
Contact toYoshiyuki Kawazoe (Materials Design by Computer simulation
Division)e-mail: kawazoe@imr.edu

X

Fig. 2 The band structures of SiNWs: NW1, NW2, NW3,
NW4, and NW5. Insets: shows their corresponding atomic
structure.
In order to determine the surface sensitivity of SiNWs, westudy the effects of changing the concentration of H on the

IMR KINKEN Research Highlights 2007

Nano / Micro Structure Controlled Materials

3FTFBSDI

X-ray Standing Wave and Atomic Resolution Holography using
X-ray Excited Optical Luminescence
X-ray excited optical luminescence (XEOL) was applied
to the x-ray standing wave (XSW) method and atom-resolved
x-ray holography (XH). We measured the incident beam
angular dependence of the luminescence intensity from a
ZnO-coated Al2O3, and found that it corresponded to the
XSW/XH pattern of an Al2O3. The present result demonstrates
that hard x-ray XSW/XH studies for light elements, whichhave been difficult by fluorescence detection, can easily be
carried out by XEOL detection.
When x rays are irradiated on solids, scattering x rays,
fluorescent x rays, photoelectrons and Auger electrons are
typically emitted. In addition, some materials emit visible,
infrared or ultraviolet rays; this phenomenon is called x-ray
excited optical luminescence (XEOL). XEOL has been
applied to site-selective XAFS measurements, whereas
XAFS techniques using x-ray fluorescence and transmittancecannot be used for such measurements. On the other hand,
the x-ray standing wave (XSW) method and atomic-resolution
x-ray holography (XH)[1,2], which together enable us todetermine the local structure around a specified atom,
normally use x-ray fluorescence from the specified atom, butthere has been no report on the use of XEOL to date.
Thus, we measured the XEOL intensity variation of a
ZnO-coated sapphire single crystal with a change in theincident beam direction, and we discuss the possibility of
XSW/XH measurements. The XSW/XH experiment was
conducted at undulator beamline BL37XU in the SPring-8
synchrotron radiation facility. Figure 1 (a) shows angularanisotropies of the XEOL intensity. The pattern in Fig. 1 (a)exhibits clear XSW lines, implying that XSW/XH data wererecorded by the XEOL detection. To estimate the validity ofthis calculation more exactly, the 3030 lines in Fig. 1 (a) were

(3 0 3 0)(a)averaged over the ky-direction, as shown in Fig. 1(b). The
position and amplitude of the experimental pattern well agreewith those of the calculated pattern.
Next, we applied the scattering pattern extraction
algorithm using the maximum-entropy method (SPEA-MEM)
to the pattern in Fig. 1(a) to reconstruct the atomic images.
Figure 2 shows the obtained real space image for the (0001)
plane at z =1.1 Å, in which O atoms exist. Circles in Figs. 2
indicate theoretical positions of O atoms. Al occupies two
distinct sites in the (0001) plane, and therefore, two types ofenvironments are overlapped in Fig. 2. Oxygen atoms seemto be reconstructed despite the smallness of their scatteringcross sections. The XSW/XH measurement using x-ray
fluorescence is generally difficult for light elements, such as
Al or Si, because of their small cross sections for inner-shellionization and the absorption of their fluorescence by air. Our
results demonstrate that hard x-ray XSW/XH measurement
is possible for light elements using XEOL.[3,4]References[1] K. Hayashi, Advances in Imaging and Electron Physics
140, 119 (2006) .
[2] S. Hosokawa, T. Ozaki, K. Hayashi, N. Happo, M. Fujiwara,
K. Horii, P. Fons, A. V. Kolobov, J. Tominaga, Appl. Phys. Lett.
90, 1361913 (2007).
[3] K. Hayashi, T. Hayashi, T. Shishido, E. Matsubara, H. Makino
and T. Yao, Bunseki Kagaku 55, 441 (2006). (Japanese).
[4] K. Hayashi, T. Hayashi, T. Shishido, E. Matsubara, H. Makino,
T. Yao and T. Matsushita, Phys. Rev. B 76, 014119 (2007).
Contact toKouichi Hayashi (Chemical Physics of Non-Crystalline Materials
Division)e-mail: khayashi@imr.tohoku.ac.jp

5
4

(b)

3

10

2

(2 1 1 0)

ky

1
0

-1

5

-2

y(Å)

-3
-4
-5
-5

-4

-3

-2

-1

0

1

2

3

4

0

5

(b)

4

χ x 10-4

kx

2

-5

0
-2

-10
-10

-4
-6

1

1.5

2

2.5
kx (Å-1)

3

3.5

4

Fig. 1 XSW/XH patterns produced by luminescence from
ZnO/Al2O3. (a) Pattern on kx-ky plane. (b) XSW line profile
averaged over ky-direction

Nano / Micro Structure Controlled Materials

-5

0

x(Å)

5

10

Fig. 2 Real-space images reconstructed from hologram in
Fig. 1 (a) by SCEA-MEM. The plane is parallel to the {001}lattice plane, and is 1.1 Å above the Al atom. Circles indicate
theoretical O positions.

KINKEN Research Highlights 2007 IMR

3FTFBSDI

Creation of Multi-functional Titanium Alloys
with Low Elastic Moduli for Biomedical Applications
Ti-Nb-Ta-Zr system alloy exhibits the functionalities suchas low Young’s modulus, super elasticity, and shape memory
effect. The Ti-Nb-Ta-Zr system alloy is expected to be highly
competitive biomedical β-type titanium alloy in the world.
Titanium alloys have been receiving much attention as
high biocompatible material with a modulus similar to that ofbone. Thus, a number of titanium alloys have been currently
developed for biomedical applications. Especially, the β-type
Ti-29Nb-13Ta-4.6Zr alloy, which has been designed by DVXα cluster method targeting a low Young’s modulus and highworkability by Niinomi et al., is a titanium alloy composed of
non-toxic and allergy-free elements for biomedical
applications [1]. It has been reported that this alloy possesses
excellent properties that are required for biomaterials.
Recently, it has been reported that there is a high possibility
for Ti-Nb-Ta-Zr system alloy to exhibit super elasticity
because the behavior of stress-strain curve at tensile loadingis nearly equal to that of super elastic alloys such as TiNi
alloy, which has been already put into practical applications.
However, the numbers of the reports on the effects of alloying
elements on the various properties of the Ti-Nb-Ta-Zr systemalloys including the Ti-29Nb-13Ta-4.6Zr alloy are few.
Therefore, the effects of conditions of fabrication and thermomechanical treatments, and alloying elements on themechanical properties, tensile deformation behaviors andbiocompatibility of the Ti-Nb-Ta-Zr system alloys were
investigated with relating to nano/micro-structures in thisstudy.
The cell viability of Ti-Nb-Ta-Zr system alloy, where thehot forged Ti-29Nb-13Ta-4.6Zr was applied, was nearly thesame as that of pure Ti, which has relatively good
biocompatibility, and greater than that of Ti-6Al-4V ELI, which
is a representative α+β titanium alloy, in both as-extracted
and filtrated extracted solutions. Namely, it was confirmedthat the nontoxicity of Ti-29Nb-13Ta-4.6Zr is equivalent tothat of pure Ti and is greater than that of Ti-6Al-4V.
Young’s moduli of the Ti-XNb-10Ta-5Zr alloys subjectedto swaging after fabrication using power metallurgy method
decrease with an increase in the Nb content as shown in Fig
1. However, the Young’s moduli of the Ti-XNb-10Ta-5Zr
alloys with 15 mass%, 20 mass%, and 25 mass% Nb, whichhave ω phase, exhibit a reverse trend. The Ti-25Nb-10Ta-5Zr
alloy exhibits the greatest elongation among the other TiXNb-10Ta-5Zr alloys because multi-deformation mechanisms
act simultaneously. Stress-induced transformation ofmetastable β phase to martensite phase, and its reversion
are recognized in the Ti-XNb-10Ta-5Zr alloys with 20 mass%and 25 mass% Nb at tensile loading-unloading as shown in
Fig. 2. The shape memory effect and the super elastic
property are expected to be achieved in these types of the
Ti-Nb-Ta-Zr system alloys. The elastic deformation behaviorof the Ti-30Nb-10Ta-5Zr alloy disobeys Hooke’s law. In thiscase, the maximum elastic strain is around 2.9%.
The wires of the Ti-29Nb-13Ta-4.6Zr alloy with diametersof 0.3 and 1.0 mm fabricated by thermo-mechanical

IMR KINKEN Research Highlights 2007

Fig. 1 Mechanical properties of Ti-XNb-10Ta-5Zr alloy as afunction of Nb content.

Fig. 2 Tensile loading - unloading stress-strain curves of TiX(= (a) 20Nb, (b) 25Nb, (c) 30Nb, and (d) 35Nb)Nb-10Ta-5Zr
alloys.
processing including cold-drawing and heat treatment also
show an unique elastic behavior with two gradients under
tensile loading. Their maximum elastic strains are around 2.8
and 2.9%, respectively, which are around twice higher thanthat of the hot forged Ti-29Nb-13Ta-4.6Zr alloy subjected toa solution treatment. Therefore, it is highly expected to be
applied to dental and surgical wires.
From these results, in this study, an important guideline
can be obtained for designing new titanium alloys for medical
applications, which have the functionalities such as low
Young’s modulus, shape memory effect and super elasticity.
Reference[1] M. Niinomi, Int. J. Fatigue, 29, 992 (2007).
Contact to
Mitsuo Niinomi (Biomaterials Science Division)e-mail: niinomi@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

3FTFBSDI

Fabrication and Applications of Bulk Metallic Glasses
This article reviews our results of the formation and resultant

materials.

applications of bulk metallic glasses (BMGs). The BMGs haveHis theory has led to the discovery of several hundreds of glassy

become to be fabricated considerably stably since the early 1990's

bulk metal alloy systems of Mg, rare earth metals (lanthanide), Zr,owing to endeavors of Dr. Inoue who has played a leading role inTi, Hf, Fe, Pd, Co, No, Cu, Ca, and Pt. Inoue et al. have also successfully

fusing academia and industry to create superior materials.

fabricated rod-like materials with maximum diameters of 80 mm or
more, which has enabled upsized materials to finally be produced.

Dr. Akihisa Inoue and his colleagues have discovered a novel

The measured quenching rates were extremely low and the slowest

phenomenon that disproved this common theory. Dr. Inoue et al.

rate recorded was 0.02 K/s. In contrast, manufacturing conventionalfound that if a metallic liquid is gently cooled below its melting point at

amorphous alloy systems requires an ultra rapid quenching (106 K/

a slow rate such as approximately 10 K/s, a supercooled metallics), which is difficult to control and limits the widths to several tens ofliquid can maintain its state without crystallization and produces a

micrometers.

stable glassy bulk metal (Fig. 1). Dr. Inoue and coworkers haveNon-equilibrium materials such as bulk metal glasses haveexamined the factors involved in this novel phenomenon and havebeen rapidly industrialized and have already been applied to various

determined the generation mechanism using experiments andproducts. These products include electromagnetic and shielding

computational algorithms. His work has advanced both basic andmaterials, materials for machines, robots, bicycles, wheel-chairs,

applied materials science technologies. In other words, his group has

electrodes, optic mirrors and junction terminals, anti-abrasion

elucidated that a glassy bulk metal is stably fabricated in an alloymaterials, materials for shot-peening balls, glasses and cell-phonesystem if the following three empirical rules (i.e., three components

frames, moldings, accessories, tools, fishing rods, golf clubs, tennisrule) are satisfied [1]. 1) The system must be comprised of multi-

rackets, and baseball bats.

components where each component is ternary or greater. 2) The

Furthermore, the development and commercialization of bulk

atomic radius of each component must differ by at least approximately

metal glass materials have been actively explored for use in

12%. 3) A negative heat must be produced when mixing the elements.

automobile pressure sensors, the world smallest 2.4 mm-diameter

Furthermore, Inoue’s group has experimentally demonstrated thathigh tork gear motor (Fig. 2), sensor pipes for Coriori flowmeters,

these alloy solid systems are structured into dense random packing

York materials for linear actuators, biomaterials, materials forstates that have short-range order atomic configurations compared to

medical devices, portable electronic devices, spring materials,

conventional amorphous materials. They have also revealed thatmaterials for high-speed transporting equipment, materials for fueleven though a metal-metal alloy system is in an icosahedral cluster

cells, impact absorbents, etc. The frontier spirit and endless creativity

state, a metal-quasimetal alloy system is formed with new local atomicof the Inoue group has significantly contributed to the development

configurations based on a transformed trigonal prism network, whichof “supercooled metal science” [3-4].

confers the stability of these new alloy systems.
In addition, Inoue and coworkers have demonstrated that analgorithm based on the three empirical rules accurately predicts thecomposition range for forming metallic glasses and the degree ofcomposition dependency on the stability of the supercooled liquids.
These findings led Inoue to propose a theory for developing new

Fig. 2 Outlook and durability of a micro-geared motor of 2.4 mm indiameter constructed by using two super precision sun-carrier
gears made of Ni-based glassy alloy [2].
References[1] A. Inoue, Acta Mater. 48, 279 (2000).
Fig. 1 Schematic illustration showing fabrications of tape-shaped[2] A. Inoue, B.L. Shen and A. Takeuchi, Mater. Sci. Eng., A 441, 18

conventional amorphous alloy (thickness: ~ 20 μm) and bulk

(2006).

metallic glasses (BMGs) (~ maximum diameter: 72 mm), both[3] A. Inoue, Proc. Jpn. Acad. Ser. B-Phys. Biol. Sci., 81, 156 (2005).

from a liquid phase. Followed by a couple of decades of the regime[4] A. Inoue, Proc. Jpn. Acad. Ser. B-Phys. Biol. Sci., 81, 172 (2005).

of "amorphous alloy" since its first success in 1960, now, one cansteadily fabricate BMGs owing to findings by Dr. Inoue.

Contact toFurthermore, researches on applications and developments

Akira Takeuchi (Non-Equilibrium Materials Division)utilizing BMGs as engineering materials have also opened up

e-mail: nano2000@imr.tohoku.ac.jp

through the success in fabricating BMGs.

Nano / Micro Structure Controlled Materials

KINKEN Research Highlights 2007 IMR

3FTFBSDI

Preparation of TiC-TiB2-SiC Ternary Eutectic Composite
TiC-TiB2-SiC composites were synthesized by arc and floating

spots were painted by yellow in color (Fig. 2 (b) and (c)). The EDzone (FZ) melting using TiC, TiB2 and β-SiC powders as starting

patterns of TiC[011] and TiB2[010] were obtained at the interface

materials. In TiC-TiB2-SiC system, a ternary eutectic compositionof TiC and TiB2, implying these zone axis were parallel to eachof 34TiC-22TiB2-44SiC (mol%) and three binary eutectic

other (Fig. 2 (b)). TiC(200) and TiB2(001) located at the solid line

compositions of 40TiB2-60SiC, 28TiB2-72TiC and 40TiC-60SiC

across the (000), suggests that these planes are parallel to each

(mol%) were found. TiC(022), TiB2(010) and SiC(111) in theother. TiC(111) and TiB2(101) located at the dash line through theeutectic composite were perpendicular to the growth direction.

(000), indicates that these planes are also parallel to each other. At

The TiC-TiB2-SiC ternary eutectic composite had the relationshipthe interface of TiB2 and SiC grains, the relationship of TiB2[010]//among crystal planes: TiC[011]//TiB2[010]//SiC[112], TiC(200)//

SiC[112], TiB2(001)//SiC(402) and TiB2(101)//SiC(220) wereTiB2(001)//SiC(402) and TiC(111)//TiB2(101)//SiC(220).

obtained. Therefore, the ternary eutectic composite prepared by

TiC-TiB2-SiC system composites can be an excellent material

FZ had the parallel relation of directions and planes: TiC[011]//due to several combined properties of TiC, TiB2 and SiC, such as

TiB2[010]//SiC[112], TiC(200)//TiB2(001)//SiC(402) and TiC(111)//

excellent oxidation resistance, high hardness and high electricalTiB2(101)// SiC(220).

conductivity. The composites of carbide and boride have been
prepared by conventional solid state sintering methods. On the

References

other hand, the self-assembly of microstructure during the[1] W. J. Li, R. Tu and T. Goto: Mater. Trans. 47, 1193-1197

solidification of melted ceramic materials is a useful technique to(2006).

develop high-performance composite materials. Particularly, an[2] W. J. Li, R. Tu and T. Goto: Mater. Lett. 60, 839-843 (2006).

eutectic texture may improve the properties of ceramic composites.

[3] R. Tu, W. J. Li and T. Goto: Mater. Sci. Forum 534-536, 1057-

Some promising results have been reported for the B4C-SiC, B4C-

1060 (2007).

TiB2 and B4C-TiB2-SiC eutectic composites by our research group.
However, the study on the TiC-TiB2-SiC eutectic composite has

Contact tonot been reported.

Takashi Goto (Multi-functional Materials Science Division)

TiC, TiB2 and SiC powders were weighed and mixed, and thene-mail: goto@imr.tohoku.ac.jp

pressed into disks with 10 mm in diameter. The specimens weremelted and solidified on a water-cooled copper holder by arcmelting. The mixture powders were also isostatically pressed in alatex tube with 10 mm in diameter. The pressed rods were sinteredat 1873 K for 3.6 ks and then directionally melted by FZ using a Xelamp heater in an Ar atmosphere.
Figure 1 demonstrates the phase diagram of TiC-TiB2-SiC
ternary system and the microstructure of binary and ternary
eutectic composites [1]. The binary eutectic points were abbreviatedas BE1 for 40TiB2-60SiC, BE2 for 28TiB2-72TiC [2] and BE3 for
40TiC-60SiC(mol%).

The

TiC-TiB2-SiC

ternary

eutectic

composition was determined as 34TiC-22TiB2-44SiC (mol%) andabbreviated as TE. The microstructure can be categorized into

Fig. 1 Phase diagram of TiC-TiB2-SiC system and microstructure

TiB2+TE, TiC+TE, TE+BE1, SiC+BE1, TiB2+BE1, TiC+SiC+TE andof binary eutectic 40TiB2-60SiC (BE1), 28TiB2-72TiC (BE2) andTiB2+TiC+TE. BE1 and BE2 showed labyrinthine microstructures,

40TiC-60SiC (BE3) and ternary eutectic 34TiC-22TiB2-44SiC

whereas BE3 and TE showed typical lamellar textures. TiC, TiB2

(mol%) (TE) composites.

and SiC grains in the ternary eutectic composite showed a lamellar
texture of 3, 3 and 2 μm in thickness, respectively, and paralleled
to each other along the growth direction(Fig. 1 (TE)).
In the ternary eutectic composite
prepared by FZ, TiC(220), TiB2(100), andSiC(111) planes were almost perpendicular
to the growth direction since significant
orientation was identified by XRD.
Figure 2 shows the TEM image of
34TiC-22TiB2-44SiC(mol%)

ternary

eutectic composite prepared by FZ at thecross-section perpendicular to the growthdirection (a) and the electron diffraction

Fig. 2 TEM image of 34TiC-22TiB2-44SiC (mol%) ternary eutectic composite perpendicular topatterns at the interface of TiC-TiB2 (b)growth direction (a) and the electron diffraction patterns at the interface of TiC-TiB2 (b) andand TiB2-SiC (c) [3]. The index from TiB2

TiB2-SiC (c). (yellow in color: TiB2)

IMR KINKEN Research Highlights 2007

Nano / Micro Structure Controlled Materials

3FTFBSDI

Ti-Nb-Sn and Ti-V-Sn Alloys with Low Young’s Modulus andHigh Strength
Low Young’s modulus at and around 50 GPa is obtained
for Ti-Nb-Sn and Ti-V-Sn alloys consisting of martensitic
microstructures. Strength increases more than 850 MPa by
the thermo mechanical processing utilizing the stressinduced and reversed martensitic transformation for Ti-NbSn alloy and the rolling texture development for Ti-V-Sn
alloy.
Titanium alloys with low Young’s modulus and highstrength are strongly demanded in various fields. Many Ti-V
base alloys for industrial applications like suspension springand several Ti-Nb base alloys for biomedical applications
have been developed to meet these demands. In this work,
microstructures and mechanical properties for Ti-Nb-Sn and
Ti-V-Sn alloys were investigated.
In metastable β (Ti-35%Nb)-4%Sn alloy, microstructure
undergoes the stress induced martensitic transformation (β/α”) with [010]α” cold rolling texture by cold rolling, and evolves
into the refined β microstructure with grain size less than 1μm
after reversion from stress induced α” to β by 523 K heat
treatment [1]. Figure 1 demonstrates the (a) Young’s modulus,(b) ultimate tensile strength and (c) elongation to fracture in(Ti-35mass%Nb)-4mass%Sn alloy at various processing
conditions. In Fig. 1(a), constituent phases identified by XRD
and TEM are also shown. As compared with the Young’s
modulus of quenched (Ti-35%Nb)-4%Sn alloy, Young’s
modulus decreases by cold rolling and recovers to the initialvalue by heat treatment at 523 K for 7.2 ks. The decrease ofYoung’s modulus will be ascribed to the α” formation with(200)α”[010]α” texture, while the recovery of Young’s modulus
to the value of quenched sample by heat treatment at 523 K
for 7.2 ks will be related to the reversion to β phase. On theother hand, marked increase in Young’s modulus withincreasing the holding time to 93.6 ks of 523 K heat treatment
can be seen. Then, it would be ascribed to the α precipitationand its increase in volume fraction. As shown in Fig. 1(b) and(c), by cold rolling and the subsequent heat treatment at 523
K strength increases, while elongation decreases. After coldrolling and the subsequent heat treatment, refined β structureand α precipitates are formed. Therefore, the increase instrength after this thermo-mechanical processing isreasonably explained by this microstructural evolution.
In martensite α’ Ti-V-Sn alloy, Young’s modulus shows
the value from 52 to 60 GPa by optimizing the composition[2]. Therefore, it is suggested that a new type α’ martensite
Ti-V-Sn alloy with low Young’s modulus is expected to
substitute for metastable β Ti-V base alloy. Furthermore, coldrolling is effective for decreasing Young’s modulus in rolling
direction due to the texture formation that rolling direction isparallel to the orientations in basal plane of h.c.p. structure.
Figure 2 shows the stress-strain curves of quenched andcold rolled α’ martensite (Ti-8%V)-4%Sn alloy. Strength
drastically increases by cold rolling, and shows the ultimate
tensile strength higher than 1000 MPa, while elongation does
not so decrease by cold rolling. After cold rolling, microstructureof acicular α’ structure for quenched condition evolves into
refined cell like dislocation structure. So the drastic increasein strength after cold rolling as shown in Fig. 2 should beconcluded to be due to the refined α’ structure.
Thus, it is found that low Young’s modulus and highstrength can be obtained in Ti-Nb-Sn alloys by utilizing
martensitic transformation and in Ti-V-Sn alloys by rolling
texture formation.
References[1] H. Matsumoto et al., J.Alloys and Comp, 439, 146(2007).
[2] H. Matsumoto et al., Metall. Trans. 37A , 3239 (2006).
Contact toHiroaki Matsumoto (Deformation Processing Division)e-mail: matumoto@imr.tohoku.ac.jp

Fig. 1 (a) Young’s modulus, (b) ultimate tensile strength and (c) elongation
to fracture in (Ti-35mass%Nb)-4mass%Sn alloy at various processing
conditions.

Nano / Micro Structure Controlled Materials

Fig. 2 Nominal stress-nominal strain curves of
α’ martensite (Ti-8mass%V)-4mass%Sn alloyquenched and cold rolled.

KINKEN Research Highlights 2007 IMR

3FTFBSDI

State of Two-dimensional Long-period Superstructure in theMetastable Fe-Nd-B Phase
A new two-dimensional long-period superstructure (LPS) has
been discovered in the Fe-Nd-B alloy system, and examined bytransmission electron microscopy. The LPS appeared during the
crystallization of an Fe-Nd-B amorphous ribbon, and is based on themetastable D03 phase. This is the first account where LPS with D03
has ever been reported.
LPS in binary alloys have been known for decades, and in thepast, IMR contributed significantly to the understanding of theformation of LPS. For example, the electron diffraction pattern of the
CuAu-II phase, in which the tetragonal CuAu-I unit cell with L10(Strukturbericht notation) structure shifts by c/2 at every M tetragonalcells thus forming a one-dimensional (1D) LPS, was reported as early
as 1953. Two-dimensional (2D) LPSs have also been known including
the Cu-Pd system. These reported LPS's exist as thermodynamicallystable phases, and an approach toward the equilibrium state canexperimentally be followed by an appropriate annealing.
Recently, we have discovered 2D LPS in a metastable crystalline
phase, which appeared during the first crystallization of Fe-Nd-B
amorphous ribbon. This newly found LPS is distinguished from thepreviously known LPS's in several aspect. First, unlike the LPS's
reported so far, the present phase is metastable, and it emerges as antransient phase. Secondly, the LPS is based on the D03 structure.,
and include two possible displacement vectors to form the 2D orderedstructure. Thirdly, the LPS appeared during the research anddevelopment effort in pursuit of so-called spring magnets based on
the Fe-Nd-B alloy systems.
Figure 1 shows a schematic model of the two kinds of anti-phaseboundaries (APB's) with two types of the displacement vectors, that is
possible in the D03 structure, having the composition of AB3. Here,closed circles denote the A atoms, corresponding to Fe in the present
case; open circles the B atoms, corresponding to Nd atoms.
Figure 2 (a) is a selected area diffraction pattern of a spherulite
whose [110] zone axis is parallel to the incident electron beam. (Rings
arise from the matrix TbCu7-type phase) It can be seen that hkl spots
with l=4n±2, e.g. 002, split in the [001] direction, while those withl=4n±1, e.g. 111, show an asymmetric two-fold splitting in the [001]and [110] directions. Also notice that these superlattice spots areaccompanied with non-negligible diffuse scattering intensity. These
splittings indicate the presence of LPS, and can be explained byassuming a 2D LPS, where APBs exist perpendicular to the [001] and[110] directions, with displacement vectors of 1/4a0[111] and1/2a0[001], respectively. Figure 2 (b) is a bright-field (BF) image of thespherulite; while Figs.2 (c), (d), and (e) are corresponding dark-field(DF) images taken with 004, 002, and 111 spots, respectively. As canbe seen here, the spherulite embedded in the matrix exhibits overall
rectangular shape of about 300nm. The DF image taken with thefundamental 004 spot (Fig.2(c)) displays bright contrasts arising from

Fig.1 Schematic models of APBs based on fundamental D03 cells,
characterized by two different displacement vectors, τ: (a)1/4a0<111>, (b) 1/2a0<100>

10

IMR KINKEN Research Highlights 2007

tiny grains from all over inside the spherulite, whereas DF images
taken with the superlattice 002 (Fig.2(d)) and 111 (Fig.2(e)) spots
show unusual morphology: bright contrasts arise from small grains in
the 'core' region of the spherulite, as well as from grains distributed
radially towards 'outskirts' of the spherulite. Note also that the brightregions of the spherulite taken with the 002 spot are essentially
identical to these taken with the 111 spot, suggesting that thesesuperlattice spots, which exhibit aforementioned different splittings, infact arise from the same regions of the spherulite.
These observations indicate that the 2D LPS emerged in thefollowing reaction sequence: amorphous -> A2 -> D03 with 2D LPS ->D03, where all the phases are metastable. Also, the 2D LPS here canbe characterized by 1/4a0<111> APBs perpendicular to <100>directions and 1/2a0<100> APBs perpendicular to <110> directions,where a0 is the unit cell parameter of the D03 cell, with average
superperiods of 3.2 and 2.4 nm, respectively.
References[1] T.J. Konno, M. Uehara, S. Hirosawa, K. Sumiyama and K. Suzuki,
J. Alloys Compds. 268, 278 (1998).
[2] T.J. Konno, M. Uehara, S. Hirosawa, K. Sumiyama and K. Suzuki,
Phil. Mag. A 79, 2413 (1999).
[2] T.J. Konno, M. Uehara, and S. Hirosawa, Phil. Mag. 86, 1427
(2006).

Contact toToyohiko Konno (Advanced Analysis of Materials Division)e-mail: tjkonno@imr.tohoku.ac.jp

Fig.2 Diffraction pattern and images of a spherulite observed along
the [110] zone axis. (a) SAD pattern, showing twofold split ofsuperlattice spots, (b) BF, (c) DF(004), (d) DF(002), and (e) DF(111)
images. Note that the spherulite is composed of tiny grains of 2-3 nm,
which possess essentially the same cubic orientation, yet (d) and (e)show that ordering to the D03 phase is incomplete in the outer regionof this spherulite.

Nano / Micro Structure Controlled Materials

3FTFBSDI

LEEM Study of Single-nucleus Polycrystallizationin Thin Film Epitaxial Growth
We have observed, by use of low energy electron microscopy(LEEM), the first direct evidence of self-driven polycrystallization
evolved from a single nucleus in epitaxial growth of pentacene on
the H-Si(111) surface. This finding gives new insight into the
crystallization of complex molecular systems, elucidating theimportance of nano-scale control of the growth conditions.
The thin film growth of molecular materials is being studied
intensely in light of their potential electronic and optical applications.
Growth optimization is desired to improve the performance oforganic devices. Therefore, it is important to account for thelimitations put on the crystal structure, which originate from theintrinsic properties of the material and the growth conditions, andmay critically affect the growth mechanism.
Among organic semiconductors, pentacene (C22H14) is one of
the most widely studied, since it has been successfully employed
in organic thin-film transistors. Upon the deposition of pentaceneon H-Si(111), LEEM images show single-monolayer high Pn
islands with fractal shapes, typical of the DLA-like growth.
Surprisingly, a quite complicated pattern of different contrasts
levels is observed in the individual branches of the Pn islands(Fig.1(a)) when the incident electron beam direction is tuned bytilting slightly away from the surface normal in the LEEM experiment.
The LEED pattern obtained from the island in Fig.1(a) is shown in
Fig.1(b). There are several distinctive features in this pattern.
Firstly, the diffraction spots originating form the bare H-Si(111)-1x1
surface are still visible (marked by yellow circles in Fig.1(b)),
showing that nucleation of the Pn islands starts immediately
without formation of a wetting layer. Secondly, careful analysis of
the LEED pattern shows six pairs of double spots, corresponding
to six different rotational Pn domains grown epitaxially on the HSi(111)-1x1 surface.
The above observation makes immediately clear that thedifferent contrast levels in Fig.1(a) originate from different rotational
Pn domains, and demonstrates a polycrystalline structure within
the single Pn grain, which was observed to develop from a single
nucleus [1].

It is found that each individual contrast level can be assigned
to one of three micro-beam LEED patterns, corresponding to thePn b-axis (longer in-plane axis) aligning almost exactly (rotated byan angle of 2.3 ± 0.3°) to each of three equivalent <110> directionsof the Si substrate surface (Fig.2(a)).
The growth and evolution of the Pn islands into multiple domains
are governed by a well-defined set of rules: (1) Branch growth
begins with the development of the “main” (center) domain, whoseb-axis aligns to the one of the <110> directions of the substrate; (2)“Outer” domains grow at each side of the “main” domain with theirb-axis rotated by ±60o; (3) New (“side”) branches grow from the sidewalls of the “outer” domains replicating procedures (1) and (2), as
shown in the Fig.2(b). Even if domains with different orientations are
created in random secondary nucleation events, such domains aresuppressed by the favored domain (with proper b-axis alignment),
developing a well defined, polycrystalline structure. Under kineticsdominated conditions with a significant dendritic growth mode, theformation of side domains whose b-axis aligns to the density gradientand develops into a new branch, is preferred over the continuous
growth of the center domain.
In conclusion, we have for the first time documented the selfdriven polycrystallization in an anisotropic growth system. The
observed polycrystallization is a result of kinetic growth processes
in conjunction with the intrinsic anisotropy of the crystal (in crystal
structure or bonding energy). The observed mechanism can beapplied to various systems, in particular, crystallization of organic
films, where both kinetic growth conditions and anisotropic effectare present.

Fig.2. (a) Schematic of Pn 2D crystal alignment reflecting thecontrast difference in the polycrystalline domain structure; longer
side of each rectangle corresponds to b-axis of the 2D unit cell ina particular domain; (b) Sketch illustrating the branchingmechanism; the equipotential lines of the density gradient ofdiffusing molecules are shown schematically (arrows indicate thefast growth directions of the branches, which align to the density
gradient).
Fig.1. (a) Tilted-beam LEEM image of the pentacene island grownon the H-Si(111) surface (imaging electron energy of 2.2 eV) – atthese imaging conditions a complicated domain pattern is visible
within the island; (b) LEED diffraction pattern obtained from theisland shown in (a); spots originating from Si(111)-1x1 lattice aremarked by yellow circles and a unit cell of one of the Pn rotational
domains is outlined by a rectangle; magnified image of the pair ofspots defining the “paired” Pn domains, rotated with respect toeach other by an angle of 4.6 ± 0.3°, is shown in the inset.

Nano / Micro Structure Controlled Materials

References[1] J. T. Sadowski, G. Sazaki, S. Nishikata, A. Al-Mahboob, Y.
Fujikawa, K. Nakajima, R. M. Tromp, and T. Sakurai, Phys. Rev.
Lett. 98, 046104 (2007).

Contact toJerzy T. Sadowski (Materials Processing and Characterization
Division)e-mail: sadows-0@imr.tohoku.ac.jp

KINKEN Research Highlights 2007 IMR

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3FTFBSDI

Surface Properties of Bulk Metallic Glass: Zr50Cu40Al10
Atomic force microscope (AFM) makes it possible toimage the morphological evolution of the Zr50Cu40Al10
metallic glass surface. The visualizations of ~200 nm
nanoclusters and viscous flow like surface features involving
crystalline phases are succeeded. The results of Auger
electron spectroscopy (AES) confirm that the observed
nanoclusters correspond to ZrO surface products. Further,
the nanocluster formation does not influence to a
crystallization process but the viscous flow is crucial for thesurface modification. The temperature dependent
observations using AFM combined chemical analysis canbe a new protocol to study fundamental properties of metallic
glass surfaces [1].
The materials of bulk metallic glass (BMG) have received
much attention because they offer extraordinary mechanicalproperties that have been applied for industrial products suchas sub-micro sized gears and liner actuators [2]. There havebeen developing to pattern and fabricate features not only for
micro-electro mechanical system but also for much smaller
sized fabrications toward nano-electro mechanical system.
Most studies of BMG have focused on both growthdevelopment of ingot sizes and improvement for multimechanical properties including static and dynamic
mechanical strength, toughness, and fatigue. However, farless attention paid for surface properties of BMG though theyplay an important role for nano- sized components where
surface effects cannot be ignored.
Here we demonstrate surface morphological evolutions
induced by thermal heating from the amorphous solid into thesupercooled liquid region. Figure 1(a) shows the surface
morphology after the heating at 673 K. Scratches run
diagonally across the image and the width is about few
hundreds nm. They were created during mechanicalpolishing. The average surface roughness that is defined by
the absolute value of surface heights averaged over thescanned area of 10 x 10 μm2 is 12 nm. Figure 1(b) shows thesurface morphology at 698 K. Surprisingly, there are nanosized structures created by thermal heating in the vacuum
below 1 x 10-8 Torr. The image of Fig. 1(c) demonstrates very
different morphology where the nanostructures disappeared
but viscous flow like features appeared at 723 K. Height
analysis shows that the maximum vertical interval between
features is ~30 nm in this image. The crystallization was
confirmed by ex-situ X-ray diffraction at this temperature.
When the sample was heated at 748 K, viscous features
drastically enhanced having chain of large mountains, as
shown in Fig. 1(d). The maximum height interval between
mountains is ~150 nm and mountains are separated about
3-4 μm.
To obtain the relationship between the morphologicalchanges and the surface compositions, we carried out AES.
The AES results reveal that the nanoclusters are composedfrom residual ZrO surface products because of the existenceof surface oxygen even at 748 K. At this temperature, theintensity of O remains in double ration compared with that ofZr, indicating the formation of a ZrO2 layer. Above the glasstransition temperature (~700K), the nanoclusters are
disappeared while the crystallization is confirmed. Thus,
these nano-products do not affect the crystallization
process.

Fig. 2 AES spectra were measured at 573 and 748 K underthe vacuum of ~10-9 Torr. At 573 K, 36 eV: Zr NVV, 59 eV: Cu
MVV, 91 and 115 eV: Zr MNN, 141 eV: ZrO chemical shift,
147 eV: Zr MNN, and 509 eV: O KLL transition peaks wereidentified. At 748K, in addition to the results at 573K, 69 eV:
Al LVV, 180 eV: Zr MNN transition peaks were obtained.
References[1] K.S. Nakayama, Y. Yokoyama, T. Sakurai, and A. Inoue,
Appl. Phys. Lett. 90, 183105 (2007).
[2] N. Nishiyama, K. Amiya, and A. Inoue, Mater. Trans. 45,
1245 (2004).

Fig. 1 AFM images were obtained after the heating at (a) 673
K, (b) 698 K, (c) 723 K, and (d) 748 K in the vacuum below
1x10-8 Torr. Above 723K, the crystallization phases wereconfirmed by XRD.

12

IMR KINKEN Research Highlights 2007

Contact toKoji S. Nakayama (Research and Development Project on
Advanced Metallic Glasses, Inorganic Materials and Joining
Technology)
e-mail: kojisn@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

3FTFBSDI

Electronic Rule for Formation of Glassy Alloys
The glass transition temperature Tg of 100 kinds of glassy
alloys can be expressed as a function of their valence
electron concentration (VEC): Tg = 131VEC +57 for metal/
metal bonding type and Tg = -236VEC +2375 for metal/
metalloid bonding one. The resonance bonding of the glassy
alloys could be caused by spd or spf hybridization.
The concept of "valence electron" has been derived by
Pauling for elements as they are bonded in the crystalline
metallic state. The valence electrons are defined as outerfreely-moving s and p or d electrons that are not bound to any
particular cations. On the other hand, the distinguishingfeature of glassy alloys without crystal structure is therandomness of their potential energy. The potential energy
minimum among atoms in the multi-component metallic glass
is not as rigid as that of crystal alloys, so that the wave
functions of these electrons are irregularly spread out much
more in space than those of the crystalline metallic valence
electrons.
Our interest lies in studying thermal stability of glassy
alloys as the function of their valence electron concentrationin terms of amorphous metallic bonding, using 100 kinds ofthe glassy alloys. The VEC in the glassy alloys (IIA, IIIB, IVC,
VD, VIE, VIIF, VIIIG) and (IIA, IIIB, IVC, VD, VIE, VIIF, VIIIG) (Mx)w(M= B, C, Si, P, Ge, Sn, Sb, Bi) respectively, of group II to VIII
elements has been defined as follows:[2, 3]
VEC=IIA+IIIB+IVC+VD+VIE+VIIF+VIIIG
for metal/metal type, (1)VEC=IIA+IIIB+IVC+VD+VIE+VIIF+VIIIG+ΣMiXW
for metal/metalloid type, (2)glassy alloys. Thus, the stable region for formation of theglassy phase would be 2.6 < VEC < 5.3 and 6.7 < VEC < 8.8
for metallic- and metalloid- type metallic glasses, respectively.
This means that transition metal elements with higher valence(e.g. group VIII and Ib elements) are the most desirable ones
for the formation of glassy alloys with higher Tg. The
resonance bonding of the glassy alloys could be caused by
spd or spf hybridization. As corroborating evidence, X-ray
photoelectron spectroscopy analysis of glassy alloy
Zr55Al10Cu30Ni5 suggested that stabilization in the glass
phase is derived from formation of a pseudo gap below Fermi
energy EF due to spd hybridization.

Fig. 1 Relation between Tg and VEC for 100 kinds of glassy
allows: green and red solids are metal/metal and metal/
metalloid bonding types, respectively.
References[1] M.Fukuhara, M.Takahashi, Y.Kawazoe and A.Inoue,
Appl.Phys.Lett. 90, 073114 (2007).
[2] M.Fukuhara, T.Mitsuda, Y.Katsumura and A.Fukawa,
J.Mater.Sci. 20, 710 (1985).
[3] M.Fukuhara and H. A.McKinstry, Phys.Stat.Sol.(b) 157,
357 (1990).

where A, B, C, D, E, F and G are atomic fractions,
A+B+C+D+E+F+G =1 and∑MiX=1. The VEC for 83 kinds ofternary and 17 kinds of quaternary Pd-, Zr-, Fe-, Ni-, Co-,
Cu-, Mg-, Ti-, Hf-, Au-, Pt-, La- and Be- based glassy alloys
have been calculated using electron valences of 35 kinds oftheir constituent elements. After we assigned the electron
valence from 1 to 15 for s, p, d and f orbits in all elements,
Contact totaking multiple orbital electron hybridization into consideration,
Mikio Fukuhara (Research and Development Project on
based on the standard method of least squares, we repeatedly
Advanced Metallic Glasses, Inorganic Materials and Joining
calculated up till we can obtain the best linearity for theTechnology)
correlation between Tg and VEC. The best result is shown ine-mail: fukuhara@imr.tohoku.ac.jp
Fig.1, for metal/metal and metal/metalloid bonding types.
Both type alloy groups show fairly good linearity, r =0.913
and 0.906, respectively. To certificate
application of VEC for the electronic
rule, we calculated the stability foranionic clusters by an excess amountof electrons transferring from inner
orbital d or f electrons, in comparison
with the stability for neutral one. We
can conclude that the increase in thevalence of magnesium stabilizes thecluster structures and affirms positively
the application of VEC for the electronic
Table 1 The electron valence for elements used in this study. Light yellow, lightrule associated with the stability ofgreen, light pink and light blue present s-, p-, d- and f- electrons.

Nano / Micro Structure Controlled Materials

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3FTFBSDI

Evolution of Size Enlargement of Bulk Metallic Glasses
In order to try enlargement of Zr-based bulk glassy alloy,
the purity of the master alloy was enhanced and supercool
solidification was also analyzed. As a result, the sort ofcrystal phase changes to simple structure by purity
enhancement. Therefore, high-purified metallic glass forming
master alloy makes more than 300K supercooling possible
easily. Furthermore, in casting, the cap-cast method inwhich three-dimensional cooling is possible was developed,
and it succeeded in enlargement of bulk metallic glasses up
to φ30mm.
The solidification morphology and structure of arc-melted
Zr50Cu40Al10 glass-forming alloys were examined [1] todetermine the impurity influences using two grades of Zr
metals: “sponge Zr” purified somewhat but not highly by the
Kroll method, and highly purified “crystal Zr”. When crystal Zr
is used, arc-melted Zr50Cu40Al10 alloy exhibits superior glassforming ability in forming glassy phase, even in a 40g weightmaster alloy. When sponge Zr is used, we can see distinct a
chain reaction of exothermic heat due to crystallization after
vitrification during solidification. Therefore, we can obtain
metallic shinning master alloy meaning complete vitrificationin the topside of master alloy as shown in Fig. 1 (a).
Furthermore, vitrification in front of the solidification interface
of arc-melted Zr50Cu40Al10 alloy with crystal Zr can occur
when the growing crystalline phase is an Al-supersaturated
B2 (B19’) -type ZrCu phase.

was observed at the crystal→glass interface by nanobeam
EDS. No change in the Al concentration was observed
around the interface, however, a significant fluctuation in the
Al concentration in the early stage of eutectic regions asshown in Fig. 1 (b). Only the high purified master alloy exhibit
glassy region with high reliability meaning a superior glassforming ability.
By use of high-purified master alloys, we tried enragementof “Inoue’s alloy” Zr55Cu30Ni5Al10 with modified tilt-cast
method [3] called cap-casting method. As shown in Fig. 2 (a),we succeeded in producing a larger sized glassy alloy rod of
30 mm in diameter by the cap cast technique. Fig. 2 (b)shows the absence of crystalline inclusion over the crosssectional area at the site of 10 mm from the bottom side. The
X-ray diffraction pattern of the cross section shown in Fig. 2(c) is also composed of distinct Bragg peaks, implying theformation of a single glassy phase.
(a)

(b)

3000

(a)

Zr50Cu40Al10

(b)

(c)

2500

Zr55Cu30Ni5Al10

Intensity (a.u.)

2000

glass

1500

crystal

1000

20 mm

500

crystal (eutectic)

(c)

0

20

Liquid (glassy phase)

30

40

50

60

70

80

90

100

2θ

Fig. 2 Outer appearance of (a) cast rod, (b) cross sectional
image at dot line and (c) X-ray diffraction pattern of the
Zr55Cu30Ni5Al10 glassy rod (φ30mm) prepared by the capcast method.

Solid (B2&B19’ phases)

2 nm

Fig. 1 Outer view (a), cross sectional optical micrograph image(b) and high-resolution transmission electron micrographimage of the interface between the glassy and crystalline
phases (c) of arc-melted high-purified Zr50Cu40Al10 ingot.
A HRTEM image of the crystal/glass interface ofsupercooled solidified high-purity Zr50Cu40Al10 alloy is shownin Fig. 1 (c). No nanocrystalline inclusions were seen in frontof the solidified interface, moreover, the solid/liquid interface
is flat. Both of these morphologies suggest that thetemperature gradient at the crystal/glass interface was notvery steep. Careful observation of the HRTEM image shows“onion-like” contrast [2], which suggests the existence of acluster in liquid in front of the crystallized interface.
Additionally, a little Zr-diluted and Cu-enriched compositional
fluctuation, whose absolute values were less than a few at%,

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IMR KINKEN Research Highlights 2007

With the aim of producing a much larger sized bulk glassy
alloy, a cap-cast technique, leading to achieve a higher
cooling rate even in an upper region, was developed bymodification of the tilt casting technique. We succeeded in
the production of a glassy Zr55Cu30Ni5Al10 alloy rod with adiameter of 30 mm.
References[1] Y. Yokoyama, H. Fredriksson, H. Yasuda, M. Nishijima
and A. Inoue, Materials Transactions, 48, No6 1363 (2007).
[2] J. Saida, M. Matsushita and A. Inoue: Appl. Phys. Lett.,
79, 412 (2001).
[3] Y. Yokoyama, K. Inoue and K. Fukaura, Mater. Trans. 43,2316 (2002).
Contact toYoshihiko Yokoyama (Advanced Research Center of Metallic
Glasses)e-mail: yy@imr.tohoku.ac.jp

Nano / Micro Structure Controlled Materials

3FTFBSDI

Nano-sized Deformation and Failure of Ultrahigh Strength Materials
Ultrahigh strength materials, such as bulk metallic glasses,
nanocrystalline metals and high performance ceramics, canwithstand very high stresses prior to failure. Nevertheless, theunderlying micro-mechanisms that control the deformation anddamage remain to be comprehensively clarified. Our recent
investigations reveal that the mechanical behaviour of theseultrahigh strength materials is associated with localized shearing
and deformation-induced phase transitions.
The superior mechanical properties of bulk metallic glasses,
nanocrystalline metals and high performance ceramics, such as
ultrahigh strength and extraordinary elasticity have made them
ideal candidates in micro-/submicro-scale applications, e.g.,
MEMS and NEMS, where both high strength and small part sizes
are required. Thus, the efforts to characterize their deformationand damage mechanisms are important for modeling theirperformances and developing new materials with improved
properties. In recent years, we employed transmission electron
microscopy and Raman spectroscopy to investigate plasticity andfailure of the ultrahigh strength materials. Our preliminary results
suggest that nonlinear deformation of these high strength materials
is associated with localized deformation and phase transitions,mainly within nano-sized shear bands.
Figure 1 shows a narrow shear band in shock-loaded boron
carbide, one of the hardest materials. The plastic deformation ofthis material takes place at high pressures and is mainly carried by
the 2-3 nm wide shear bands [1]. The operation of the local
shearing was found to accomplish by a phase transition that the
crystalline phase turns to be amorphous within the narrow shear
bands. The amorphization induced lattice changes can provide
irreversible plastic strains. Meanwhile, the amorphous bands withlow strength cause the catastrophic failure of the high strength
materials. This stress-induced amorphization has been observedin a number of other hard materials, such as silicon and boron
nitride. However, with certain loading conditions, the high strength
ceramics sometimes exhibit metal-like plasticity. An example ishigh purity alumina subjected to plate impacts. Well-defined
dynamic yielding can be observed, which is proved to be performedby global twinning. These twinning planes act as the preferredpaths for subsequent cleavage and lead to materials failure [2].
Different from the high-strength inorganic materials in whichstress-induced localized phase transitions lead to material failure,
we found that deformation induced microstructural evolution inhigh strength metallic materials, such as bulk metallic glasses andnanocrystalline metals, benefits the ductility. More recently,
extraordinary plasticity and work hardening were observed in Cuand Pt-based monolithic BMGs with high strength. Although anumber of explanations have been suggested, including highPoisson’s ratio and structure inhomogeneity, the underlying
physical mechanism, in particular microscopic explanation,
remains mystery. We employed high-resolution electron
microscopy to characterize the atomic structure of shear bands inthese ductile BMGs. Our observations reveal a novel self-locked
effect in ductile bulk metallic glasses. Strain softening caused by
localized shearing can be effectively prevented by the crystalline
nanoparticles that are in situ produced by plastic flow within theshear bands (Fig. 2), leading to large plasticity and strain hardening.
These atomic-scale observations not only well explain theextraordinary plasticity that was recently observed in the materials,
but also reveal a novel deformation mechanism that can effectively
improve the ductility of monolithic metallic glasses.
Understanding the micromechanisms of the deformation andfailure of untrahigh strength materials is an interesting topic forboth basic and applied research. The well-defined physicalmechanisms will allow for more robust models to predict the theirmechanical performances and lead the way toward processing
improved materials.

Fig. 2 Self-locked shear bands in a ductile Zr50Cu50 BMG with
deformation-induced nanocrystallie particles
References[1] M. W. Chen, J. W. McCauley and K. J. Hemker, Science 299,
1563 (2003).
[2] M. W. Chen, J. W. McCauley, D. Dandekar and Bourne, Nature
Materials 5, 614 (2006).
[3] M. W. Chen, A. Inoue, W. Zhang and T. Sakurai, Phys. Rev.
Lett. 96, 245502 (2006).

Fig. 1 Nano-sized shear band in shock-loaded boron carbide. The
high resolution transmission electron microscopy reveals that
amorphization takes place within the 2-3 nm wide bands.

Nano / Micro Structure Controlled Materials

Contact to
Mingwei Chen (International Frontier Center for Advanced
Materials)
e-mail: mwchen@imr.tohoku.ac.jp

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