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Intermetallics 17 (2009) 607–613
Contents lists available at ScienceDirect
Intermetallics
journal homepage: www.elsevier.com/locate/intermet
Phase-separated microstructures and shear-banding behavior in a designed
Zr-based glass-forming alloy
X.H. Du a, b, J.C. Huang a, *, H.M. Chen a, H.S. Chou a, Y.H. Lai a, K.C. Hsieh a, J.S.C. Jang c, P.K. Liaw d
a
Institute of Materials Science and Engineering, Center for Nanoscience and Nanotechnology, National Sun Yat-Sen University, Kaohsiung 804, Taiwan, ROC
Department of Materials Engineering, Shenyang Institute of Aeronautical Engineering, Shenyang 110034, PR China
c
Department of Materials Science and Engineering, I-Shou University, Kaohsiung 840, Taiwan, ROC
d
Department of Materials Science and Engineering, The University of Tennessee, Knoxville, TN 37996, USA
b
a r t i c l e i n f o
a b s t r a c t
Article history:
Received 11 August 2008
Received in revised form
12 January 2009
Accepted 30 January 2009
Available online 28 February 2009
We have employed a thermodynamic-computation approach to identify the composition of the
Zr–Cu–Ni–Al alloy system exhibiting a two-liquid miscibility phase equilibrium in the liquid-temperature region, which tends to favor the occurrence of the liquid-phase separation. Guided by these
calculated diagrams, a Zr-based alloy with a 10 at.% Al is designed, and its bulk-metallic glass (BMG) is
prepared successfully by the copper-mould suction casting. A heterogeneous microstructure characterized by the existence of phase-separated regions with several to decades micrometers in size forms in
the BMG. Under uniaxial compressive loading, the designed Zr-based BMG demonstrates the continuous
‘‘work hardening’’ and remarkable macroscopic plastic strain at room temperature. The improvement of
mechanical properties is attributed to the unique glassy structure correlated with both the heterogeneous microstructure and the micro-scaled phase separation, leading to the extensive shear-band
formation, interaction, and multiplication.
Ó 2009 Elsevier Ltd. All rights reserved.
Keywords:
B. Glasses, metallic
B. Mechanical properties at ambient
temperature
E. Phase diagram, prediction
F. Electron microscopy, transmission
1. Introduction
According to Inoue’s third empirical rule [1], the formation of
monolithic metallic glasses with the high glass-forming ability
(GFA) requires a high negative heat of mixing among the main
constitute elements. However, the resulting monolithic bulkmetallic glasses (BMGs) usually demonstrate highly inhomogeneous deformation below the glass-transition temperature even
under compressive loading with or without confinement [2–5]. To
solve the problem of low plastic strains, the development of bulkmetallic glass composites (BMGCs) has been proven to be a promising way. Two main approaches have been explored so far, one is to
in situ precipitate crystalline phases in the BMG matrix, and the
other is to introduce foreign particles or micrometer-sized pores
into the BMG matrix [6–10].
Apart from forming BMGCs, one can develop a hierarchical
glassy microstructure, which is of potential interest in terms of the
nucleation and propagation of shear bands. Some studies have
demonstrated that for some glass-forming systems, if there exist
repulsive interactions among atoms, its plasticity can be enhanced
by the introduction of the nano-scale chemical inhomogeneity in
* Corresponding author. Tel.: þ886 7 5252000; fax: þ886 7 525 4099.
E-mail address: [email protected] (J.C. Huang).
0966-9795/$ – see front matter Ó 2009 Elsevier Ltd. All rights reserved.
doi:10.1016/j.intermet.2009.01.019
the amorphous matrix [11–14]. Recently, the influences of different
length-scale ductile constituent phases on the mechanical properties of Zr-based BMG composites have been investigated,
showing that the deformation is controlled by the types of the
constituent phases and their morphology [15]. Following this line
of thought, the development of a microstructure with micro-scale
chemical heterogeneities is expected to be of potential interest in
terms of the deformation mechanics of BMGs.
Although the concept of a phase-separating system is contrary
to Inoue’s rule, many valuable efforts have been devoted to the
topic of phase-separating metallic glasses, and impressive progresses have been made so far [16–20]. For the glass-forming system,
which consists of a strong repulsive tendency of components, it is
expected that a liquid-state phase separation into two glasses can
occur during cooling from the liquid state due to the existence of
the two-liquid miscibility phase equilibrium in the liquid-temperature region, and two glassy phases can form during rapid cooling
from the liquid state. This trend has been verified by the recently
observed multi-component La–Zr–Cu–Ni–Al alloys (La–Zr: þ74 kJ/
mol [16]) [17], a quaternary Y–Ti–Al–Co (Ti–Y: þ58 kJ/mol [16])
system [18], a ternary Ni–Nb–Y (Nb–Y: þ127 kJ/mol [16]) alloy [19],
and a Cu–Zr–Al–Y (Zr–Y: þ35 kJ/mol [16]) [20] system. In these
materials, the decomposition and structure formation take place
already in the melt prior to the solidification. The heterogeneity
size in this new group of amorphous-metallic alloys ranges from
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a nanometer to micrometer dimension with the features of selfsimilarity. However, due to the inferior GFA originated from the
existence of a strong de-mixing tendency of components, all these
samples are prepared as the ribbons. Thus, no investigation has
been made so far concerning the effect of micro-scaled phaseseparated regions on the ductility.
To overcome the dilemma, recently, based on thermodynamic
calculations, a micro-scale phase-separated Zr-based BMG
(Zr63.8Ni16.2Cu15Al5) with remarkable plasticity has been successfully developed by the same authors [21]. A unique mechanism
operative in the phase-separated BMG, i.e., the effective hindrance
of a hard glassy phase on the propagation of shear bands, has been
suggested. This trend indicates that by designing the compositions
based on the thermodynamic calculation, some monolithic BMGs
can deform plastically through the formation of phase-separated
microstructures. In this paper, a newly designed Zr-based glassforming alloy with a 10 at.% Al based on the calculated diagrams, as
well as its microstructure and mechanical response is reported.
Especially, the physical process underlying the formation of microscale-separated microstructures and the resulting shear-banding
behavior are systemically described.
2. Alloy design and experimental procedures
The thermodynamic-computation approach to identify the
potential compositions of Zr–Ni–Cu–Al alloys exhibiting the twoliquid miscibility phase equilibrium in the liquid temperature
region has been described in the previous paper for 5 at.% Al
(Zr63.2Ni16.2Cu15Al5) [21]. For the quaternary system, the complete
isotherm is a tetragonal volume under a constant pressure. Fig. 1 is
a two-dimensional section obtained from cutting with a constant
10 at.% Al content from this isothermal tetragonal. The isopleth A-A
cut (a temperature versus composition section through the multidimensional phase diagram) marked in Fig. 1(a) is calculated and
shown in Fig. 1(b), in which a two-liquid miscibility region can be
seen. According to the diagram, the ‘‘X’’ alloys with the compositions locating in the two-liquid-phase region are anticipated to
produce the two-phase BMGs by a liquid-phase separation. Thus,
the alloy composition in the present study is designed.
Using a subregular solution model [22,23], the free energy of
a liquid can be obtained from the expression:
G ¼ RT
X
i
xi lnxi þ
n
n
X
X
i ¼ 1 j ¼ iþ1
xi xj
X
uvij xi xj
v
(1)
v¼0
where uij represents the interaction parameter between the i and j
components. According to the solution thermodynamic principle,
there are positive uij values or positive enthalpy mixing value to
induce the two-liquid miscibility region. Then, we can calculate the
spinodal lines based on the conditions that the second derivate of
Gibb’s free energy is equal to zero within the two-liquid miscibility
region. In our parallel study using thermodynamic computations
[24], such spinodal boundaries are carefully established. Fig. 2
shows the schematic diagram including the spinodal boundary as
a dashed-dotted line. The liquid composition within the spinodal
region may decompose into two liquid phases assigned as a Ni-rich
liquid and Cu-rich liquid and form the two liquid metallic glass. The
designed alloy composition is selected, namely, Zr65.8Ni15.8Cu8.4Al10
in at.%. It should be noted that the current alloy has more Ni than
Cu, which is different from most Zr-based BMGs containing more
Cu than Ni. The typical example is the one reported by Zhang et al.
[25], designated as the Inoue’s alloy. The current Zr65.8Ni15.8Cu8.4Al10 alloy lies within the spinodal region.
The Zr65.8Ni15.8Cu8.4Al10 cylindrical ingots, 2 mm in diameter,
were prepared by arc-melting the pure elements under a purified
Fig. 1. (a) Zr-rich isothermal section calculated for the Zr–Ni–Cu–Al system for
a constant 10 at.% Al at 1050 C (X1: Zr65.8Ni15.8Cu8.4Al10), and (b) the temperature, T,
versus composition, X, section for the AA cut, where the composition of the X1 alloy
can be located.
Ar atmosphere and in situ suction casting in a copper mould. The
differential scanning calorimetry (DSC) with a heating rate of
20 K/min, X-ray diffraction (XRD, Diano 8536X, Cu-Ka radiation)
and high-resolution transmission electron microscopy (HRTEM,
Jeol JEM-3010) coupled with the energy-dispersive X-ray spectroscopy analysis (EDS) are used to assure the amorphous nature
and characterize the phase-separating phenomenon of the as-cast
alloys.
The rod specimens with an aspect ratio of 2:1 (height/diameter)
are tested in uniaxial compression under an initial strain rate of
2 104 s1 at room temperature using an Instron 5582 universal
testing machine. Both ends of the specimens are polished to make
them parallel to each other prior to the compression test. The outer
surfaces as well as the fracture surfaces of the deformed specimens
are examined by the scanning electron microscopy (SEM, Jeol JSM6330 TF) with EDS. Foils for TEM observations are thinned by ion
milling to produce a central hole. For specimen thinning, the ionbeam energy was about 3 keV, and the milling angle about 7 –9 .
Under these ion-milling conditions, no heating or radiationinduced crystallization can occur as reported by Sun et al. [26]. All
images and diffractions are recorded within 30 s to prevent the
possible artifacts caused by the electron beam.
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609
Liquid phase region
Tl
Temperature, T
L1
L2
Td
Tg
N domain
Ni-rich
M region
Composition
C domain
Cu-rich
Fig. 2. The schematic diagram includes the spinodal boundary as a dashed-dotted line.
The liquid composition within the spinodal region may decompose into two liquid
phases assigned as a Ni-rich liquid and Cu-rich liquid.
3. Results and discussion
3.1. Microstructural characterizations of the as-cast ingots
The representative XRD and DSC curves of the as-cast alloy are
shown in Fig. 3. The diffuse hump in the XRD curve, coupled with
the clear glass transitions and sharp crystallization events in the
DSC curve, confirms the glassy nature of the BMG. The glass-transition temperature, Tg, crystallization temperature, Tx, supercooled
temperature range, DTx, solidus temperature, Ts, and liquidus
temperature, Tl, are 647, 745, 98, 1105 and 1178 K, respectively. The
GFA parameters, such as Trg (¼Tg/Tl) [27], g [¼Tx/(Tg þ Tl)] [28], and
gm [¼(2Tx Tg)/Tl] [29] are 0.55, 0.41, and 0.72, respectively. The
corresponding critical cooling rate to produce the glassy phase is
around 1 K/s [28]. All of these results prove that the designed alloy
has a good glass-forming ability.
Fig. 4(a) shows the TEM bright-field image of the as-cast alloy,
also presenting the apparent phase separation, one with a darker
contrast and the other with a brighter contrast. The inserted
diffraction pattern was taken from a larger region covering both the
dark and bright domains. Due to the similar sizes of the Ni
(an atomic radius of 1.24 Å) and Cu (an atomic radius of 1.28 Å)
atoms and the diffuse halo nature, it is not able to differentiate the
halos originated from the Ni or Cu domains. From the TEM observations, it seems that the phase separation occurs only in partial
regions. There are still some glassy matrix regions with no phase
separation, exhibiting a simple gray contrast and single halodiffraction pattern, as shown in Fig. 4(b).
The TEM/EDS results measured for the local domain with
a darker contrast in the Zr65.8Ni15.8Cu8.4Al10 (at.%) alloy give
a composition of Zr65Ni20.3Cu5.2Al9.5 (at.%), i.e., the Ni-rich N
domains, whereas the chemistry of the brighter contrast areas is
Zr66.4Ni10.6Cu12.7Al10.3 (at.%), i.e., Cu-rich C domains. The EDS
measurements at different locations would give slightly different
readings Thus, the above two compositions are not fixed, still
showing either Ni- or Cu-rich phases for the darker or brighter
domains, respectively. Fig. 4(c) and (d) displays the high magnitude
and lattice images obtained from the interface between two glassy
phases with the darker and brighter contrasts denoted by ‘‘A’’ in
Fig. 4(a). The round interface areas do not show a sharp lattice
fringes boundary, indicating that the two glassy phases form by the
Fig. 3. XRD (a) and DSC (b) patterns obtained from the 2-mm-diameter ingots of the
Zr63.8Ni16.2Cu15Al5 alloy.
separation in the liquid state during the solidification process. This
phenomenon has also been observed in many liquid phase-separated glass-forming alloys [17–19].
In this case, the special microstructure in the designed Zrbased BMG can be understood from the pseudo-binary phase
diagrams in Figs. 1(b) and 2 of the Ni-rich composition on the left
and Cu-rich composition on the right. During cooling, the melt
crosses the critical temperature and separates into the corresponding compositions along the spinodal boundaries, i.e., the
primary liquid-phase separation occurs. The present results show
that the cooling rate during the copper-mould casting provides
a sufficient time for the Zr-based alloy melt to form two separated
liquid phases at relatively high temperatures, where the L1 and L2
liquid phases are immiscible. In this case, the liquid-phase separation mainly involves the diffusion of Ni and Cu (with a heat of
mixing of þ4 kJ/mol), which will be promoted by the sudden
agitation and homogeneous mixing through the violent melt flow
during the suction–casting process. Since the viscosity is low in
the liquid, there is enough driving force for the separated phase to
achieve the equilibrium shape with the lowest interfacial energy,
i.e., in the shape of a sphere. The size may be micro-scaled due to
the rapid diffusion of atoms in the liquid. Depending on the
concentration fluctuation of local chemical compositions in
the melt, separation can only occur in the regions, where the
compositions are falling into the spinodal region. Following that,
with the rapid cooling rate of the copper-mould suction casting,
these separated liquid phases are frozen into the separated glassy
phases.
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Fig. 4. Microstructures of the 2-mm-diameter as-cast Zr65.8Ni15.8Cu8.4Al10 rods. (a) and (b) TEM bright-field images for different regions, respectively, with their corresponding
selected-area-diffraction patterns; (c) a high-magnitude image of the interface marked by ‘‘A’’ in (a), and (d) a high-resolution lattice image of the interface marked by ‘‘B’’ in (c).
3.2. Compression studies at ambient temperature and observations
of shear-banding behavior
Fig. 5 shows the engineering stress versus engineering strain
curves of the designed BMGs obtained in the uniaxial compression
under the initial strain rate of 2 104 s1 at room temperature.
While the alloy of the Inoue’s composition fails right after the
elastic stage, the two-glassy-phase Zr-based BMG shows an engineering strain up to 4.6% prior to the unstable deformation. In this
engineering stress–strain curve, the yield stress and the maximum
engineering compressive stress are w1750 and 1950 MPa, respectively. The flow stress increment from the yield to maximum stress
is about 200 MPa, or w11.5% increase. The apparent strain-hardening exponent is measured to be about 0.08, which is close to the
low limit of a face-centered cubic metals (0.10–0.50). The inserted
figure presents the amplified compressive stress-displacement
curve for the designed BMG, which shows that plastic deformation
proceeds mainly by numerous serrated flows.
Traditionally, the discussion on the issue of working hardening
or softening is based on the true stress–strain curves in which the
uniform change in the sample cross-sectional area is taken into
account. For the homogeneous plastic deformation under
compression, the cross-sectional area would increase accordingly.
However, because the plastic deformation of the BMG is inhomogeneous and accomplished by the accumulation of the offset from
individual shear bands, the true sample cross-sectional area change
with increasing compressive straining is actually minimum and
cannot account for the 11.5% flow-stress increment. In fact, due to
the shear-banding offset effect, the effective cross-sectional area
that withstands the applied forces is actually decreasing with
increasing compression straining. Thus, the intrinsically mechanical property of the BMG might be overshadowed by the true
Fig. 5. Engineering stress versus strain curves obtained from the uniaxial compression
test of the 2-mm-diameter rod. The inset figure is the enlarged curves, which shows
a serrated flow.
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X.H. Du et al. / Intermetallics 17 (2009) 607–613
stress–strain curve transformed by the assumption of homogeneous and uniform plastic deformation. The minor work hardening-like behavior seen in the engineering stress and strain curve
of the phase-separated BMG might imply the resistance force
encountered by the shear bands when they propagate within the
two periodically-spaced harder and softer glassy phases. When one
propagating shear band faces the resistance, the neighboring shear
band, with the same or deviated angle with respect to the loading
axis, might be induced at a higher flow stress level.
The SEM observations on the outer surface of the fractured
Zr65.8Ni15.8Cu8.4Al10 specimen are presented in Fig. 6. Both the
primary shear bands and the final fracture plane are inclined about
45 with respect to the compression axis. Therefore, the shear
stress is mainly responsible for the deformation and failure of the
BMG. Fig. 6(a) shows that multiple primary shear bands have been
formed prior to the final fracture for the BMG. It is interesting to
observe that the offset amount of an individual primary shear band
is appreciable, as presented in the circle part in Fig. 6(a). The
compression strain results from the accumulation of the offset
amount associated with these shear bands.
In this case, the significant offset for an individual shear band
should originate from the sliding behavior between the two parts of
the BMG divided by the shear band. However, the ‘‘sliding’’
behavior seems to exhibit a very limited development. Thus, the
lateral constraint effect should be significant for the arrest of shear
bands. Fig. 6(b) is a high-magnitude SEM observation of the
deformed sample, showing that the slip steps are jagged and
typically occur as a group of several slip steps close together. We
believe that these features are originated from the branching of
individual shear bands as they propagate through the materials.
Branching can distribute the plastic strains associated with the
main shear band, and this trend makes it more difficult for
a propagating shear band to nucleate a crack. Furthermore, the
observed shear bands can be characterized as being semi-straight,
wavy, or twisting, indicating the existence of the effective resistance to the propagation of shear bands. Based on these observations, it is reasonable to conclude that an arrest mechanism
responsible for the effective prevention of the runaway slip along
a single shear band plays a critical role for the development of the
marked global plasticity for the designed alloy.
3.3. Mechanisms of the formation of multiple shear bands
To investigate the mechanism underlying the arresting of shear
bands, the plastically-deformed sample was characterized by TEM.
A bright-field TEM image (Fig. 7) shows a severe deformation
region in the vicinity of a fracture surface. As shown in Fig. 7,
a bundle of small shear bands with a width of 30–50 nm as marked
by the dotted lines can be observed, indicating that the deformation
611
strain is achieved by multiple shearing behavior. It should be
mentioned that no diffraction contrast generated by the nanocrystals within each shear band has been observed. This trend is
different from many previous observations showing profuse
nanocrystals in a primary shear band [30,31]. The selected-area
electron diffraction pattern taken from the ‘‘A’’ site in Fig. 7a is
shown in the inset plot of Fig. 7b, demonstrating the amorphous
nature in the shearing regions.
On the other hand, the morphology of the developed shear
bands on the fracture surface can provide the useful information.
The examination on the fracture surface of the Zr65.8Ni15.8Cu8.4Al10
compressive specimen reveals the mixed morphologies, consisting
of vein-like patterns and highly rough regions, as exhibited in
Fig. 8(a). In this case, a locally-melted region is observed on the
fracture surface, suggesting that a large amount of strains along the
shear band led to localized melting before fracture. The mixture of
veins and rough regions is not consistent with some investigations,
which show that the vein-like pattern is the main part of the
compressive fracture surfaces for most monolithic BMGs [32,33].
Fig. 8(b) is a part within the highly rough regions, where the shear
bands can be clearly observed. The shear bands are highly
branched, and their movement is rather wavy in nature. This trend
suggests that the propagation of shear bands is effectively hindered
during the compression process. Hence, the catastrophic shear-off
through the whole sample is avoided. According to the state of
shear bands, the fracture surface can be divided into the ‘‘A’’ and ‘‘B’’
regions, as denoted in Fig. 8(b). By SEM/EDS analyses, the compositions of regions ‘‘A’’ and ‘‘B’’ are different, and the composition of
the ‘‘B’’ region resembles the composition of phase-separated
regions observed by TEM (i.e., the Ni-rich region). Fig. 8(c) and (d)
shows the shear-band morphology in regions ‘‘A’’ and ‘‘B’’,
respectively, at a higher magnification. The shear bands in the
region ‘‘A’’ are highly wavy and branched, which confirms the
formation of a high density of shear bands in the region ‘‘A’’. The
limited length of the abundant shear bands surrounding the region
‘‘B’’ indicates the strong interaction between shear bands and
phase-separated regions. Most interestingly, intercrossed step-like
shear bands are also observed in the region ‘‘B’’ of the fracture
surface, as illustrated in Fig. 8(d). The distance between shear bands
is about 1–2 mm and a large number of shear bands appear at about
45 from the stress axis. All these results indicate that the separated
phases can promote the homogeneous deformation in an individual shear band accommodating the applied strain rather than
the accumulation of damage at some particular shear bands.
We know that a direct result of the phase separation is the
emergence of the regions with different chemical compositions,
which will cause the inhomogeneous distribution of hardness in
the BMGs. In this case, in the phase-separated regions, the softer
Cu-rich glassy phase in the two-phase BMG is the glassy matrix
Fig. 6. SEM observations of the outer surface of the 2-mm-diameter fractured sample at room temperature.
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Fig. 7. A bright-field TEM image of a bunch of shear bands nearby the fracture surface (a) and corresponding selected area electron diffraction pattern (b).
Fig. 8. Observations of the sample and fracture surface after the deformation: (a) SEM image of the fracture surface; (b) a shear-band pattern on the fracture surface; (c) shear
banding in the region A; and (d) shear banding in the region B.
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X.H. Du et al. / Intermetallics 17 (2009) 607–613
with a brighter contrast, and the local hardness measured by nanoindentation for the brighter phases is 5.02 0.04 GPa, which is
softer than that of the darker phases (6.5 0.04 GPa). Thus,
a unique microstructure composed of hard regions surrounded by
soft regions can be realized at the phase-separated regions for the
current Zr-based BMGs. As we have known, due to the different
chemical compositions, different areas in the glasses may exhibit
different critical shear stresses for yielding [34,35]. We propose that
upon yielding of the two-glassy-phase BMG, the softer phase
deform first. Simultaneously, the load is transferred to the
surrounding harder glassy phase, causing the nucleation of a shear
band due to the stress concentration. Upon further loading, the
nucleated shear band propagates and interacts with other phaseseparated regions/shear bands. In this case, the growth of an individual shear band is arrested, and the shear band is cooled. Thus,
the crack disappears. This trend hinders a single shear band to
extend critically through the whole sample at the onset of the
plastic deformation and cause an early fracture. Consequently,
abundant shear bands can be developed in the phase-separated
regions, as shown in Fig. 8(b)–(d). Recent investigations have
demonstrated that if the BMGs are divided into compartments on
a micron scale, the fracture mode changes, and their compressive
ductility will be improved significantly [36–38]. Indeed, the small
regions, if standing alone, can deform greatly with a high failure
resistance by limiting the propensity of forming mature shear
bands, as shown in the study [39,40].
However, a single shear band cannot carry more plastic strains
due to the small longitudinal offset. The remarked compression
strain should be equal to the amount of longitudinal offsets
provided by the multiple shear bands. For this reason, the interrelation of the shear-banding behavior with the ‘‘work hardening’’
mechanical behavior is of interest for this phase-separated Zrbased BMG. The interactions between the shear bands and phaseseparated regions [Fig. 8(b)] would compensate the softening
behavior induced by the shear-band formation [4]. With the
propagation of an individual shear band, more phase-separated
regions are swept, making the hindering force increase. On the
other hand, because of the strong interface between the two glassy
phases, the hard phase can also accommodate the stress intensity
within the shear band. This feature results in a decrease of the local
temperature inside the shear band, leading to a ‘‘hardening’’ of the
glassy matrix. This trend is consistent with the fact that the portion
with a vein-like pattern on the fracture surface is very limited in the
current Zr BMGs [Fig. 8(a)]. These factors will result in a ‘‘work
hardening’’ behavior macroscopically, by which the plastic
instability can be controlled, as demonstrating in many ductile
BMGs [41].
4. Conclusions
Based on the thermodynamic computation, a phase-separated
Zr-based bulk-metallic glass with a dramatically-enhanced
compressive strain over 16% at room temperature is successfully
fabricated. During the compression process, no nanocrystalline
phase forms even in the shear bands. Based on the SEM and TEM
observations, the development of plasticity is attributed to phaseseparated regions heterogeneously distributed in the as-cast
samples with the size scale of several to decades micrometers,
resulting in the extensive shear-band formation, branching, and
interaction. The research result presented here suggests that the
613
creation of a microscale two-glassy-phase microstructure through
the liquid-phase separation might be a promising way to enhance
the intrinsic plasticity of bulk-metallic glasses.
Acknowledgements
We would like to greatly acknowledge the sponsorship by the
National Science Council of Taiwan, ROC, under the Project No. NSC
95–2218-E-110–006. PKL very much appreciated the support of the
National Science Foundation International Materials Institutes
(IMI) Program (DMR-0231320) with Dr. C. Huber as the Program
Director.
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