1. Introduction
Metallic glasses have been one of the extensive research subjects over the past twenty years. The unique
properties such as high strength and hardness, special
damping property, soft magnetic characteristics, and
oxidation/corrosion resistance enable this new category
of metallic materials to be applied for structural or
functional applications. Before 1990’s the research and
development (R&D) focuses of metallic glasses were
mainly on the Au-Si and Pd based system. After the
breakthrough in early 1990’s, many other systems such
as La, Mg, and Zr based bulk metallic glasses (BMGs)
were developed. So far, the existing and on-going efforts
have developed numerous successful systems, including
Zr, Fe, Mg, Au, Pd, Pt, Cu, Ni, Ti, Co, Al, Ce, Ca, La,
and many other rare-earth based amorphous alloys, as
reviewed in a number of review papers [1-5].
In addition to the metallic glasses in the bulk form,
recently, metallic glasses have been extended to the thin
or thick film area, in an attempt to find applications in
the micro-electro-mechanical systems (MEMS). For such
applications, the amorphous alloys sometimes need to be
fabricated via the sputtering or other thin film processes.
The as-deposited thin film metallic glasses (TFMGs)
usually exhibit a uniform composition without
micro-segregation, avoiding the possible local
crystallization during the fabrication process.
The initial R&D activities of the BMGs and
THMGs in Taiwan start from early 1990’s, in academic
university such as National Tsing-Hua University and
National Taiwan Ocean University, as well as in research
institutes such as Industry Technology Research Institute
(ITRI) and Chung-Shan (Military) Institute of Science
and Technology. The initial R&D was aimed on the Zr,
Mg and Fe based BMGs. Later research was
continuously extended in several universities, such as
National Sun Yat-Sen University, I-Shou University and

National Taiwan Ocean University, together with the
reinforcement from Metal Industries Research &
Development Centre (MIRDC). The major joint efforts
were lunched in 2000’s. An international collaboration
project on metallic glasses was approved in 2005,
covering eight research laboratories in National Sun
Yat-Sen University, National Taiwan Ocean University,
I-Shou University, National Cheng-Kung University and
National Taiwan University of Science and Technology,
as well four USA laboratories in Oak Ridge National
Laboratory, University of Tennessee, and University of
Wisconsin, Madison. MIRDC was also involved to
promote industry applications.

2. Alloy design
The initial alloy design was limited to the Zr based
BMGs [6-13]. Effects of individual micro-alloying with
Si (0~10 at%) or B (0~4 at%) on the thermal properties
of the Zr65Cu17.5Ni10Al7.5 base alloy have been
systematically investigated. It is found that the addition
of either B or Si would posses a positive effect on
thermal stability improvement of the base alloy, while
keeping the same glass forming ability (GFA) as the base
alloy. The largest activation energy of crystallization,
about 360 kJ/mol, occurs at the compositions of
Zr63Cu17.5Ni10Al7.5B2 and Zr61Cu17.5Ni10Al7.5Si4.
Furthermore, the Zr65-x-yAl7.5 Cu17.5Ni10SixBy alloy system,
where x = 0~4 and y = 0~2, was selected for studying the
effect of simultaneously adding B (with a smallest
atomic size) and Si (with a negative heat of mixing) on
its crystallization and thermal stability. The experimental
evidence and the kinetic analyses revealed that, with the
addition of 4 at% Si and 1 at% B, the Trg (=Tg/Tl) and
γ (=Tx/(Tg+Tl), where Tg, Tx, and Tl are the glass
transition, crystallization, and liquidus temperature,
respectively) values increase from 0.57 and 0.40 for the
Zr65Al7.5Cu17.5Ni10 base alloy to 0.62 and 0.43,

suggesting the improved glass forming ability. The
saturation point for nucleation for 4%Si-1%B and
1%Si-2%B amorphous alloys occurs at the 88%
crystallization fraction, much higher than the 65% for the
base alloy [10-12], implying that these metalloid
elements would extend the nucleation stage and result in
smaller crystalline particles. Moreover, the variation of
the incubation time, as shown in Fig. 1 [6], exhibits an
increasing trend with increasing silicon addition, directly
supporting that the Si addition would improve the
thermal stability of the Zr65Al7.5Cu17.5Ni10 base alloy.
Parallel studies have also been done on the
Zr53Cu30Ni9Al8 amorphous alloy [13] which
computationally designed by thermodynamics and deep
eutectic methodology [14]. The improvement of thermal
stability by the Si addition is contributed by the increase
of their atomic packing density as well as their Zr-Si and
Ni-Si strong atomic bonding, promoting the coating of Si
thin layer on the nanocrystalline Zr2Ni phase [6].

Fig. 1 Incubation time as a function of isothermal
temperature for the Zr65-xAl7.5Cu17.5Ni10Six alloys with Si
content from 0 to 10at% [6].

The Mg65Cu25-xGd10Agx alloy system was
systematically investigated to reveal the effect of Ag (in
substituting Cu) on the GFA and crystallization behavior
of the base alloy [15]. Because of a large difference in
atomic size between Ag and the constituting elements;
the atomic radius of Ag is larger than that of Cu by 11%
and is smaller than those of Mg and Gd by 10% and
19.5%, respectively, the atomic packing density of the
Mg65Cu25-xGd10Agx alloy would increase with increasing
Ag content and concurrently increase its GFA. The
highest γ value (0.407) and a relatively high Trg (0.58)
occur at the Mg65Cu22.5Gd10Ag2.5 alloy. More recently,
new optimum alloy designs are made based on the recent
model of optimum composition extension from the
binary eutectic pairs, originally proposed by Lu et al.
[14]. This model suggests that the optimum Mg based
BMGs might possess a composition with a lower amount
of Mg and a higher amount of rare earth (RE) element,
for example, the Mg58Cu31Y11 and Mg58Cu31Gd11[16,17].
A series of Mg based BMGs with 50-60at% Mg and
10-15at% dual RE elements (Nd and Y) are prepared by
atmosphere controlled induction melting and injection
casting methods [18]. The selection of Nd for partially
replacing Y element is inspired by the facts such as the
difference between Nd and Y in covalent atomic radius
(Nd: 0.164 nm; Y: 0.162 nm), electronegativity (Nd: 1.14,
Y: 1.22), electronic configuration (Nd: 4f⁴5d⁰6s², Y:
4d¹5s²), and near-zero heat of mixing of the Y–Nd binary

density, 360 in Hv hardness, and 700 MPa for the
compressive strength. Moreover, the toughness of the
Mg based BMGCs exhibit an increasing trend with
increasing nano-sized ZrO2 content, and can reach a level
of 8.9 MPa.m^0.5 [32,33].

3. Fabrication
The fabricating routes for the BMGs in I-Shou
University and National Sun Yat-Sen University are
either injection or suction casting, with arc or induction
furnace. For producing the Zr based (or CuZr based)
BMGs, an atmosphere-controlled
arc-melting/suction-casting process is the preferable way
to prepare the Zr based alloys. The rods with dimension
of 8 mm φ x 70 mm L and plates with dimension of 3
mm T x 30 mm W x 60 mm L of the Zr based BMGs can
be successfully fabricated by this method. On the other
hand, for making the Mg based BMGs and BMGCs,
two-steps atmosphere control induction melting and
injection-casting process is in use to prepare the highly
reactive and high vapor-pressured Mg based alloys. The
rods with dimension of 10 mm φ x 60 mm L and plate
with dimension of 2 mm T x 20 mm W x 60 mm L of the
Mg based BMGs can be successfully prepared. Some
examples are shown in Fig. 3.

Fig. 3 Showing the Zr- or Mg-based based BMG rods
and plates in various dimensions.

Fig. 4 Morphology of a spray-formed Mg–Cu–Gd
layered composite plate with a Gaussian-shape, and 300
mm in diameter, 12 mm in maximum thickness, and 510
g in mass [40].

Spray forming has also been applied to fabricate Al
and Mg based amorphous plates measuring 1 to 10 mm
in thickness in National Cheng Kung University [36-40].
Spray forming can be considered as an intermediate
processing between casting and deposition, with a
cooling rate in the range of 1-100 K/s. Spray forming
was firstly employed to fabricate Al89La6Ni5 based
BMGs and later Mg-Cu-Y or Mg-Cu-Gd systems, as

shown in Fig. 4 [40].
BMGs and BMGCs have been fabricated using the
mechanical alloying and consolidation methods [41-48].
Examples include the Mg and Ti based amorphous
composites with additives of WC, SiC, carbon nanotube
(CNT) powders. It is found that mechanical alloying may
cause partial dissolutions of the additives, and thus the
thermal stability of the amorphous matrix is altered due
to the deviation in chemical composition of the glassy
matrix. Yet, the supercooled liquid region of resulting
materials remains almost unchanged. Compared with the
pure amorphous structure, the hardness of the BMGCs
increases as much as 30% due to the presence of
additives. The composites suffer partial crystallization
during the consolidation process in the supercooled
liquid region.
The amorphous materials have also been prepared
by accumulative roll bonding (ARB) via room
temperature cold rolling in the solid state for alternating
thin layers of various metals with a special arrangement
of composition [49-56]. Binary Zr-Ti, Zr-Ni, Zr-Cu and
Zr-Al, ternary Zr-Cu-Ni, quaternary Zr-Cu-Ni-Ti and
pentanary Zr-Cu-Ni-Ti-Al systems were systematically
explored. It is found the crystalline foils can be mixed
into homogeneous nanocrystalline and/or amorphous
phase with increasing ARB cycles. After around 80 to
120 cycles, the alloys can be mostly transformed into the
amorphous matrix, as shown in Fig. 5 [49]. Note that the
volume fractions of the remianed nanocrystalline phases
and the interdiffusion-induced amorphous matrix can be
controlled by the applied ARB cycle, as depicted in Fig.
6 [51]. For example, if an alloy is designed to possess
around 50% amorphous phase in Zr-Cu system, then
about 80 ARB cycles are needed. At that time, the
remaining nanocrystalline pure Zr or pure Cu phases are
about 20 nm in size. The diffraction spots and rings in
the TEM diffraction patterns of the ARB specimens are
all originated from the pure elements, meaning that the
nanocrystalline phases are those unmixed hard particles
left from the previous severe deformation and diffusion
processes. A critical size of the nanocrystalline phases
around 3 nm is consistently observed in all binary,
ternary, and pentanary Zr-X based alloys, below the
critical size a sudden transformation from the
nanocrystalline to amorphous state would occur [50-54],
as shown in Fig. 7 [50].
The friction stir process with effective liquid
nitrogen cooling has also been adopted in preparing the
MgAlZn intermetallic compound alloys, with
nanocrystalline icosahedral particles plus a certain
amount of amorphous phases [57,58].

Fig. 5 (a) The stacked Zr and Ti foils after 50 ARB
cycles, (b) typical TEM micrograph and electron
diffraction pattern of the vitrified Zr50Ti50 region by ARB,
and (c) typical XRD patterns of the ARB Zr50Ti50
samples after different cycles at room temperature [49].

Fig. 6 The variations of (a) nanocrystalline grain size
and (b) amorphous volume fraction in various Zr–X
alloys as a function of ARB cycles, based on the X-ray
diffraction results [51].

Fig. 7 TEM bright field images of the interface region
between the nanocrystalline and amorphous phases. The
smallest nc phase size near the interface is around 3 nm
[50].

4. Mechanical behavior
Extensive research has been devoted to the
mechanical properties of BMGs, in an attempt to develop
strong and tougher BMGs, as described below in Sec. 5.
The effect of the compression specimen aspect ratio was
examined in the Mg based BMGs [59].
Room-temperature compression tests are conducted on
specimens with various height-to-diameter ratios (h/d)
from 2:1, 1:1, 1:2, to 1:4. The failure strength,
deformation strain, and the fracture surface morphologies
are seen to vary systematically in accordance with the
specimen h/d ratio. For specimens with h/d of 2:1 or 1:1,
the compression response is similar to those in most
reports. In contrast, for specimens with lower h/d ratios,
especially at h/d = 1:4 (or 0.25), the shear band
propagation appears to be constrained by the top and
bottom platens, resulting in the enhanced ductility.
Micro-pillars with a diameter of 1 and 3.8 μm were
fabricated from the Mg based metallic glasses using
focus ion beam, and then tested in micro-compression at
strain rates ranging from 6x10^-5 to 6x10^-1 s^-1, as shown in
Fig. 8 [60]. The strength of the micropillars varies within
1342-1580 MPa, or 60-100% increment over that of the
bulk specimens. This strength increase can be
rationalized using the Weibull statistics for brittle
materials, and the Weibull modulus (m) of the Mg based
metallic glasses is estimated to be about 35. The surface
examination of the micro-pillar specimens indicates that
the number of shear bands increase with the sample size
and strain rates.
The specimen size dependence was also explored
in the more ductile Zr based metallic glasses (with
phase-separation [21]), covering rod specimens with
diameters of 2-4 mm down to micro-pillars with

diameters of 0.7 to 4 μm [61]. The increment of strength
is also analyzed in terms of the Weibull statistics. The
extracted Weibull modulus is ~60 for this ductile
phase-separated Zr BMGs [61], compared with the
values ~35 and ~40 for the brittle Mg based BMGs [60]
and Au49Ag5.5Pd2.3Cu26.9Si16.3 BMG [62]. These results
are consistent with other reported Weibull modulus
values for the malleable Zr48Cu45Al7 (m=73.4) [63] and
brittle (Zr48Cu45Al7)98Y2 (m=25.5) [63]. It seems that the
size dependence is related to the ductility and cast defect
of the amorphous alloys.

Fig. 8 Micro-pillars prepared by focus ion beam: (a)
before micro-compression and (b) after
micro-compression to about 30% strain [60].

5. Toughening
The current critical issue for metallic glasses appears
to be the toughening of their brittle nature. Based on
previous efforts, the most promising approaches might
be the enhancement of the Poisson ratio by tailoring the
compositions, the employment of in-situ phase
separation to form an amorphous matrix mixed with a
dendrite crystalline phase (forming a dendrite-reinforced
amorphous composite) or a second glassy phase (forming
a dual-phased glass) [21-23]. Tremendous efforts have
been made in terms of the above approaches. One is to
form liquid phase separation at high temperatures during
induction melting, followed by rapid quenching to form
the phase-separated glass. Successful results in Zr based
glasses (e.g., Zr63.8Ni16.2Cu15Al5 or Zr65.8Ni15.8Cu8.4Al10)
have achieved compressive ductility over 30%, as shown
in Fig. 9 [21]. Another approach is to employ the porous
particles into the amorphous matrix to form an
amorphous matrix composite. Successful results in Mg
based glasses added with 5-25 volume percent of porous
Mo particles (~30-50 μm) have achieved compressive
ductility over 10%, as shown in Fig. 10 [64]. Both these
two approaches are designed with the concept in
separating the homogeneous glassy matrix into numerous
individual compartments with smooth but distinctive
interfaces. The complicated interaction of propagating
shear bands with these compartments and interfaces
would promote multiple shear banding and resulting in
improved ductility.
The gain of high compressive strength and
remarkable macroscopic plastic strain of 30% in the
phase-separated Zr63.8Ni16.2Cu15Al5 or
Zr65.8Ni15.8Cu8.4Al10 is attributed to the unique glassy
structure correlated with the chemical inhomogeneity on
the micron scale; the hard phases surrounded by the soft
phases, leading to an inherent capability of extensive
shear-band formation, interactions, and multiplication
[21-23]. The work shows that, by the two-liquid phase
separation in accordance with the thermodynamic

prediction, a good-GFA and ductile BMG can be
achieved from the two-glassy-phase microstructure.
Recent works on the Mg based BMGs added with
micro-scaled Fe [65] or Ta [66] particles are under
investigation. The BMGCs of these systems also show
compression plasticity over 10%. The brittle nature of
the Mg based BMGs can be sufficiently toughened by
the Fe and Ta particles.

Fig. 9 (a) TEM bright-field image, with an inserted
electron diffraction pattern, of the 2 mm Zr-based BMG
rod with micro-scaled phase separation, (b)
room-temperature stress-strain curve for as-cast 2 mm
diameter cylinders. The inset shows the deformed sample
[21].

Fig. 10 (a) SEM micrograph of the homogeneous
distribution of porous Mo particles in the Mg-based bulk
metallic glasses, with the inserted X-ray diffraction
pattern, (b) representative room-temperature compressive
engineering stress-strain curves for the as-cast 2 mm
diameter compression specimens. The curves are shifted
relative to each other for clarity [64].

6. Modeling
Molecular dynamics (MD) simulation was applied
for the structural evolution and mechanical response of
the metallic glasses using the tight-binding potentials for
the elements [53-56]. For the simulation for the ARB
solid-state vitrification process, the molecular dynamics
simulation reveals that it needs only 6 cycles for a
bi-layer structure, with 5 nm in each layer thickness, to
fully transform into a completely amorphous state, as
shown in Fig. 11 [54]. The potential energy between two
distinct atoms increases during the transition between
nanocrystalline and amorphous state due to the severe
structure change. It increases the difficulty for atoms to
maintain the crystalline structures. The values of the
nearest neighboring distance extracted from the
transmission electron microscopy (TEM) observation
and the MD simulation are both smaller than the
theoretically calculated value. A new tighter packing
atomic structure in the Zr50Ni50 amorphous alloy is
formed due to the large negative mixing enthalpy.
The gradual vitrification evolution of atom mixing
and local atomic pairing structure of the binary Zr-Ni
alloy during severe deformation at room temperature is

traced numerically by MD simulation in terms of the HA
index (Fig. 12 [67,55]) developed by Honeycutt and
Anderson [67]. It is found that the icosahedra clusters
will gradually develop with increasing disorder
environment of alloys. Other compound-like transition
structures were also observed as transient in the Zr-Ni
ARB couple during the solid-state amorphization process
under severe plastic deformation. Since the time scale of
the MD simulation is several orders of magnitude shorter
than reality, the simulated atomic evolution can be
viewed as that would occur at ultra-high strain rates.
Simulation of the deformation behavior and shear band
evolution during indentation, compression and fatigue
loading is now undertaken.

Fig. 11 The microstructural evolution and the associated
two-dimensional Fourier transform of the bi-layered
Zr50Ni50 model subjected to various ARB cycles: (a)
initial state with the distinct Fourier transformed
diffraction patterns for crystalline Zr and Ni, and (b)
after 6 ARB cycles with the Fourier transformed
diffraction patterns for an amorphous structure [54].

Fig. 12 (a) The schematic drawing of the related HA
pairs [67]. (b) Variations of the HA indices of the Zr-Ni
alloys during different ARB cycles: 1431, 1541, and
1551 [55].

7. Thin film metallic glasses
The vapor-solid quenching during sputtering
enables the amorphous or nanocrystalline phases to form.
The thin film amorphous alloys are generally called as
the thin film metallic glasses (TFMGs), which can be
prepared by many physical vapor deposition methods
such as sputtering or evaporation. TFMGs might have
tremendous application potentials in the areas of surface
coating and MEMS [68,69]. Co-sputtering by two or
three guns appears to be a promising process in studying
the TFMGs with a wide variation of composition and
microstructure.
It is also the first to report annealing-induced
amorphization in glass-forming Zr47Cu31Al13Ni9 [68] and
Fe65Ti13Co8Ni7B6Nb1 [69] films at a temperature within
the supercooled liquid region (Fig. 13 [69]). The
extensive amorphization is presumably attributed to
sufficient thermal and interfacial energies between
nanocrystallites and glassy matrix that are present in the

as-deposited condition. The formation of comprehensive
amorphous structure gives rise to notable alterations in
the electrical, mechanical and magnetic properties of
annealed films (Fig. 14 [68]). Important feature of these
works is that a prominent strengthening effect is
observed due to the amorphous matrix dispersed with
nanocrystalline phases upon annealing. In addition, the
glass-forming film coating is applied to improve the
fatigue properties of material for the first time [70,71].
The excellent adherence and the smooth film surface
have a noticeable influence on the fatigue life. With the
deposition of the Zr47Cu31Al13Ni9 glass-forming film, the
fatigue life of the 316L stainless steel could be increased
by 30 times, and the fatigue limit could be elevated by
30%, depending on the maximum stress applied to the
steel (Fig. 15 [70]). The high strength and the good
bending ductility might be other important factors for the
improvement.

Fig. 13 Typical TEM bright-field images and
diffraction patterns of the Fe65Ti13Co8Ni7B6Nb1 films in
as-deposited and annealed conditions [69].

Fig. 14 (a) DSC thermogram of an as-deposited
Zr47Cu31Al13Ni9 film; variations of (b) Knoop
ultra-microhardness and (c) electrical resitivity with the
annealing temperature. Approximate location of
supercooled liquid region is marked by the dash lines to
facilitate visual comparison [68].

corrosion mechanism [85].

9. Thermomechanical behavior
The viscous flow behavior of the Mg58Cu31Y11
bulk amorphous rod in the supercooled viscous region is
investigated using differential scanning calorimetry
(DSC) and thermomechanical analyzer (TMA) [87].
Below the glass transition temperature, Tg, a linear
thermal expansion coefficient of 3±1x10^-6 m/m.K was
obtained. In contrast, significant viscous deformation
occurred as a result of a compressive load above Tg. The
onset, semi-steady-state, and finish temperatures for
viscous flow, determined by TMA, are slightly different
from the glass transition and crystallization temperatures
measured by DSC. All these characteristic temperatures
decrease with increasing applied stress, suggesting
accelerated crystallization in the present Mg58Cu31Y11
under stress. The appropriate working temperature for
microforming as determined by the semi-steady-state
viscous flow temperature is about 460-474 K. The
effective viscosity within this temperature range is
estimated to be about 10⁷-10⁹ Pa.s, and it increases with
increasing applied stress, as demonstrated in Fig. 16(a)
[87]. The viscosity behavior of the Mg based BMGs falls
in-between the strong SiO2 glass and the polymer
o-terphenyl fragile glass, as shown in Fig. 16(b) [87].

Fig. 16 (a) Temperature dependence of the effective
viscosity for the indicated applied stress at a heating rate
of 10 K/min.(b) The viscosity data on Mg58Cu31Y11 (or
the Angell plot). The data on the strong SiO2 and fragile
o-terphenyl glasses are compared with Mg58Cu31Y11 [87].

The viscous flow behavior and thermomechanical
properties of Mg65Cu25-xBxGd10 (x = 0 and 3 at%) [88]
and Mg65Cu25-xAgxGd10 (x=0, 3, 10 at%) [89] BMGs in
the supercooled liquid region have also been investigated.
It is found that the fragility of the supercooled liquid is
reduced by the boron or silver addition, thus degrading
the deformability. This conclusion is supported by the
many other extracted parameters. The lowest viscosity
was found in the base alloy, Mg65Cu25Gd10, in the level
of 10⁶~10⁷ Pa.s. Thus, even with much higher hardness,
the B- or Ag-additive Mg based BMG will be more
difficult to be formed, which appears to be a negative
factor in applying in the micro-forming or
nano-imprinting industry.
With the relatively low viscosity between 10⁶~10⁷
Pa.s within the supercooled temperature region of the
Mg58Cu31Y6Nd5 BMG, a high plastic strain can be
obtained easily by compression test within this
temperature region. In addition, the XRD result shows
that the Mg58Cu31Y6Nd5 BMG sample maintains its
amorphous state after compression at 458 K and different
strain rates. In parallel, a strain rate sensitivity exponent

(m) close to 1.0 can be obtained for Mg58Cu31Y6Nd5
deformed at 448 K, 453 K and 458 K, suggesting an
ideal Newtonian fluid superplasticity behavior. A 6 mm
Mg58Cu31Y6Nd5 BMG rod was extruded at 458 K into a
long wire for more than 160 mm in length, as depicted in
Fig. 17 [90].

Fig. 17 A Mg58Cu31Y6Nd5 metallic glass wire with
dimension of 1 mm in diameter and a length more than
160 mm fabricated by extruding the 6 mm diameter
BMG rod at 458 K [90].

10. Forming and Application
Extensive plasticity (or superplasticity) of BMGs
at temperatures within the supercooled liquid region has
been studied extensively [91-94]. The alloys studied
included the Pd, Zr, Cu, and Ce based BMGs. The
superplastic-like behavior confirmed in the supercooled
liquid region can be used to process imprints of nanoand
micro-parts. A typical example is presented in Fig.
18 [95] for nano-grating using a Pd-based BMG. It is
demonstrated that nanostructured gratings with 600-nm
and 1500-nm periods can be replicated with high fidelity
on BMG in air using master Si dies. The grating features
and diffraction properties of the first-generation BMG
gratings are comparable to those of the original Si master
dies, but the quality of replica depends on the feature size
[95]. These replicated BMG gratings can be used as
molding dies to further create the second-generation
replicas on conventional polymethylmethacrylate
(PMMA). This study shed lights on the possibility of
nano-imprinting optical gratings on BMG in air and is a
first step towards fabrication of integrated optical
components such as diffractive optical nanostructured
elements.

Fig. 18 SEM micrographs of the 600-nm (a,b,c) and
1500-nm (d,e,f) gratings in tilted view. The feature
widths of 600-nm and 1500-nm Pd40Ni40P20 BMG
gratings are (b) ~150 nm and (e) ~655 nm, respectively.
Green arrows indicate the polishing marks in BMG [95].

Potential application of the Zr based BMG can be
extended to the medical tool such as the surgical razor or
micro-surgery scissors. Because the razor made by the Zr
BMG presents much smoother edge than the razor made
by martensitic stainless steel. One example is shown in
Fig. 19 [96].

Fig. 19 (a) Appearance of Zr-based BMG razor and
commercial surgical razor. Enlarged knife edge of the
blade: (b) Zr-based BMG blade and (c) commercial
S-15C blade [96].

Fig. 20 The Mg58Cu31Y6Nd5 BMG photographs of (a)
pre-engraved die and imprinted hologram pattern at 458
K, (b) secondary electron image of SEM and surface
profile of the imprinted hologram pattern [90].

Fig. 21 (a) Photograph of imprinted hologram
pattern of the (Zr53Cu30Ni9Al8)99.5Si0.5 BMG, (b) SEM
images of circle A area in (a), (c) enlarges SEM image of
circle B area in (b), and (d) surface profile of the
imprinted hologram pattern [13].

Since both of the Zr based and Mg based BMGs
present superior viscous flow behavior within their
supercooled region. Therefore, these two BMGs can be
applied to duplicate the hologram pattern via a
micro-imprinting process. The replication of a hologram
pattern with 100 nm depth was replicated by the
Mg58Cu31Y6Nd5 BMG plate in the supercooled
temperature region (Fig. 20) [90]. In addition, an
imprinted hologram pattern with 0.6 μm depth also
demonstrates extremely good microforming ability of the
(Zr53Cu30Ni9Al8)99.5Si0.5 BMG, as shown in Fig. 21 [13].
The other attractive aspect in using BMGs for
micro-forming is the relatively low forming temperature.

The Mg [97-100] and Au based BMGs [100] have low
glass transition temperatures in the range of 140-160oC,
making the viscous forming or imprinting easy to operate.
The low forming temperature lies in the same
temperature range for polymer materials such as PMMA,
and the low forming temperature also prevent from the
oxidation problem. The pressing can also be performed
in air due to the low forming temperatures within
140-160oC for the Mg and Au based BMGs. Successful
micro-imprinting of the Mg58Cu31Y11 BMGs have
achieved for making hexagonal micro-lens arrays. The
original convex Ni-Mo mold made by electro-plating is
pressed onto the BMG at 140-160oC for a few minutes at
a low pressure in the range 0.1 to 1 MPa (Fig. 22(a) [98]).
And the printed concave BMG mold is applied onto
PMMA, forming a convex pattern. The results
demonstrate that the imprinting is feasible and promising.
The resulting hexagonal micro-lens arrays are shown in
Fig. 22(b) [98]. Micro-scaled V-grove patterns (Fig. 22(c)
[98]) can also be imprinted on the Mg [97-99] and Au
based BMGs [100]. With the anti-oxidation,
anti-corrosion, and good forming ability, the Au based
BMG may be a material with high potential for micro
electro mechanical systems (MEMS) applications.

Fig. 22 (a) The schematic replication process first on
BMG and then on the PMMA, (b) SEM micrograph
showing the imprinted micro-lens, and (c) SEM
micrograph showing the V-groove pattern [98].

Acknowledgments
The authors gratefully acknowledge the long-term
sponsorship by National Science Council of Taiwan,
ROC, under the projects NSC 94-2218-E-110-009, NSC
95-2218-E-110-006, NSC 96-2218-E-110-001, NSC
95-2221-E-011-224-MY3, NSC 96-2628-E-011-117-
MY3, etc. The authors gratefully acknowledge the
collaboration with the USA team, Dr. C. T. Liu in Oak
Ridge National Laboratory and University of Tennessee,
Prof. Y. A. Chang in University of Wisconsin, Madison,
Prof. T. G. Nieh in University of Tennessee (formerly in
Lawrence Livermore National Laboratory), and Prof. P.
K. Liaw in University of Tennessee. The authors also
expressed their gratitude to their other local Taiwan team
members, Prof. K. C. Hsieh in National Sun Yat-Sen
University, Profs. P. Y. Lee, W. Kai and J. K. Wu in
National Ocean University and Tatung University, Prof.
C. Y. A Tsao in National Cheng Kung University, as well
as many devoted post-doctors and graduate students.
Thanks are due to Dr. C. T. Liu and Prof. K. L. Lin (of
NSC) who strongly promoted this international
collaboration projects.

References
[1] A. Inoue, “Stabilization of Metallic Supercooled Liquid
and Bulk Amorphous Alloys”, Acta Mater., 48 (2000) 279.
[2] W. H. Wang, C. Dong, C. H. Shek, “Bulk Metallic

Review article (可點選連結以下載本文)
Thin Film Metallic Glasses: Preparations, Properties, and Applications
J.P. Chu, J.C. Huang, J.S.C. Jang, Y.C. Wang, and P.K. Liaw

(J.O.M., vol. 62, No. 4, 2010, pp. 19-24.)

論文資料：請依發表時間之先後順序填寫，內容依序包括作者姓名(依原出版順序，主要作者請加註*)、題目、期刊名稱、卷數、起訖頁數及出版年。

SCI/SSCI Rank
Factor
N / M

SCI/SSCI Cited Number

是否為通訊作者?

1

J. C. Huang and A. J. Ardell*, "Crystal Structure and Stability of the T1 Precipitates in Aged Al-Li-Cu Alloys", Mater. Sci. & Tech., vol. 3, 1987, pp. 176-188.

?22 / 71

50

No

2

J. C. Huang and A. J. Ardell*, "Addition Rules and the Contribution of d' Precipitates to Strengthening of Aged Al-Li-Cu Alloys", Acta Metall., vol. 36, 1988, pp. 2995-3006.

1 / 66

61

No

3

J. C. Huang and A. J. Ardell*, "Precipitation Strengthening of Binary Al-Li Alloys by d' Precipitates", Mater. Sci. Eng., vol. A104, 1988, pp. 149-156.?

42 / 189

31

No

4

J. C. Huang* and G. T. Gray, "Microband Formation in Shock-Loaded and Cold-Rolled Metals", Acta Metall., vol. 37, 1989, pp. 3335-3347

1 / 66

59

Yes

5

J. C. Huang and G. T. Gray, "Serrated Flow and Negative Rate Sensitivity in Al-Li Base Alloys", Scripta Metall. Mater., vol. 24, 1990, pp. 85-90.?

2 / 66

43

Yes

6

P. S. Follansbee, J. C. Huang*, and G. T. Gray, "Low Temperature and High Strain Rate Deformation of Ni and Ni-C Alloys and Analysis of the Constitutive Behavior According to an Internal State Variable Model", Acta Metall. Mater., vol. 38, 1990, pp. 1241-1254.???

1 / 66

33

Yes

7

H.-P. Pu, F. C. Liu, and J. C. Huang, "Characterization and Analysis of Low-Temperature Superplasticity in 8090 Al-Li Alloys", Metall. Mater. Trans., vol. 26A, 1995, pp. 1153-1166.?

9 / 66

61

Yes

8

H.-C. Shih, N. J. Ho*, and J. C. Huang, "Precipitation Behavior in Al-Cu-Mg and 2024 Aluminum Alloys", Metall. Mater. Trans., vol. 27A, 1996, pp. 2479-2494.

9 / 66

47

No

9

H. C. Fu, J. C. Huang*, T. D. Wang, and C. C. Bampton, "Superplastic Deformation Mechanisms of Super-a2 Ti3Al Thin Sheets as a Function of Strain", Acta Mater., vol. 46, 1998, pp. 465-479.??

1 / 66

28

Yes

10

I. C. Hsiao and J. C. Huang*, "Development of Low Temperature Superplasticity in Commercial 5083 Al-Mg Alloys", Scripta Mater., vol. 40, 1999, pp. 697-703.

2 / 66

39

Yes

11

I. C. Hsiao, S. W. Su, and J. C. Huang*, "Evolution of Texture and Grain Misorientation in Al-Mg Alloy Exhibiting Low Temperature Superplasticity", Metall. Mater. Trans., vol. 31A, 2000, pp. 2169-2180.??

9 / 66

29

Yes

12

J. C. Huang*, I. C. Hsiao, T. D. Wang, and B. Y. Lou, "EBSD Study on Grain Boundary Chracteristics in Fine-Grained Al Alloys", Scripta Mater., vol. 43, 2000, pp. 213-220.

42/ 66

37

Yes

13

H. K. Lin and J. C. Huang*, “High Strain Rate and Low Temperature Superplasticity in AZ31 Mg Alloy Using Simple High-Extrusion-Ratio Extrusion Method”, Mater. Trans. JIM, vol. 43, 2002, pp. 2424-2432.

11 / 66

32

Yes

14

I. C. Hsiao and J. C. Huang*, "Deformation Mechanisms during Low and High Temperature Superplasticity in 5083 Alloy”, Metall. Mater. Trans, vol. 33A, 2002, pp. 1373-1384. ?

9 / 66

37

Yes

15

S. F. Su, J. C. Huang*, H. K. Lin, and N. J. Ho, “Electron Beam Welding Behavior in Mg-Al-Based Alloys”, Metall. Mater. Trans., vol. 33A, 2002, pp. 1461-1473.

9 / 66

52

Yes

16

Y. N. Wang and J. C. Huang*, “Texture Analysis on Hexagonal Materials”, Materials Chemistry and Physics, Vol. 81, 2003, pp. 11-26.

42 / 189

101

Yes

17

Y. N. Wang and J. C. Huang*, “Comparison of Grain Boundary Sliding in Fine-Grained Mg and Al alloys during Superplastic Deformation”, Scripta Mater., vol. 48, 2003, pp. 1117-1122.

2 / 66

28

Yes

18

C. J. Lee and J. C. Huang*, “Cavitation Characteristics in AZ31 Mg Alloys during LTSP or HSRSP”, Acta Mater., Vol. 52, 2004, 3111-3122.

1 / 66

45

Yes

19

C. I. Chang, C. J. Lee, and J. C. Huang*, “Relationship between Grain Size and Working Strain Rate and Temperature during Friction Stir Processing in AZ31 Mg Alloy”, Scripta Mater., vol. 51, 2004, pp. 509-514.

2 / 66

145*

Yes

20

H. K. Lin, J. C. Huang*, and T. G. Langdon, “Relationship between Texture and Low Temperature Superplasticity in AZ31 Mg Alloy”, Mater. Sci. Eng. A, vol. A402, 2005, pp. 250-257.

40 / 189

129*

Yes

21

M. C. Kuo, J. C. Huang*, and M. Chen, “PEEK Composites Reinforced by Nano-Sized SiO2 and Al2O3 Particulates”, Mater. Chem. Phys, vol. 90, 2005, pp. 185-195.

42 / 189

96*

Yes

22

C. J. Lee, J. C. Huang*, and P. J. Hsieh, “Mg based Nanocomposites Fabricated by Friction Stir Processing”, Scripta Mater., vol. 54, 2006, pp. 1415-1420.

2 / 66

88*

Yes

23

Y. N. Wang, C. I. Chang, J. C. Huang, Y. N. Wang, C. I. Chang, J. C. Huang, “Texture and weak grain size dependence in friction stir processed Mg-Al-Zn alloy”, Scripta Mater., vol. 55, 2006, pp. 637-640.

2 /66

45

Yes

24

M. C. Kuo, J. C. Huang*, and M. Chen, “Non-isothermal Crystallization Kinetic Behavior of Alumina Nanoparticle Filled Poly(ether ether ketone)”, Mater. Chem. Phys, vol. 99/2-3, 2006, pp. 258-268.

42 / 189

32

Yes

25

Y. N. Wang and J. C. Huang*, “The Role of Twinning and Untwinning in Yielding Behavior in Hot-Extruded Mg-Al-Zn Alloy”, Acta Mater., vol. 55, 2007, pp. 897-905.

1 / 66

162*

Yes

26

X. H. Du, J. C. Huang*, C. T. Liu, and S.P. Lu, “New criterion of glass forming ability for bulk metallic glasses”, J. Appl. Phys, vol. 101, 2007, 086108.

20 / 95

58

Yes

27

X. H. Du, J. C. Huang*, K. C. Hsieh, J. S. C. Jang, and P. K. Liaw, Y. H. Lai, and H. M. Chen, “Two-glassy-phase bulk metallic glass with remarkable plasticity”, Appl. Phys. Lett., vol. 91, 2007, p. 131901.

10 / 95

57

Yes

28

C. J. Lee, J. C. Huang, and T. G. Nieh, “Sample size effect and microcompression of Mg65Cu25Gd10 metallic glass”, Appl. Phys. Lett., vol. 91, 2007, p. 161913.

10 / 95

76*

Yes

29

C. I. Chang, X. H. Du, J. C. Huang, “Achieving ultrafine grain size in Mg-Al-Zn alloy by friction stir processing”, Scripta Mater., vol. 57, 2007, pp. 209-212.

2 / 66

68

Yes

30

Y. H. Lai, M. C. Kuo, J. C. Huang*, and M. Chen, “On the PEEK Composites Reinforced by Surface-Modified Nano-Silica”, Mater. Sci. Eng. A, vol. 458, 2007, pp. 158-169.

40 / 189

41

Yes

31

J. S. C. Jang, J. Y. Ciou, T. H. Hung, J. C. Huang*, and X. H. Du, “Enhanced mechanical performance of Mg metallic glass with porous Mo particles”, Appl. Phys. Lett., vol. 92, 2008, p. 011903.

10 / 95

33

Yes

32

Y. H. Lai, C. J. Lee, Y. T. Cheng, H. M. Chen, H. S. Chou, X. H. Du, C. I. Chang, J. C. Huang*, S. R. Jain, J. S. C. Jang, and T. G. Nieh, “Bulk and Microscale Compressive Properties of Zr-Based Metallic Glass”, Scripta Mater., vol. 58, 2008, pp. 890-893.

2 / 66

64*

Yes

33

C. I. Chang, X. H. Du, and J. C. Huang*, “Processing Nano-Grained Microstructure in Mg-Al-Zn Alloy by Two-Pass Friction Stir Processing”, Scripta Mater., vol. 59, 2008, pp. 356-359.

2 / 66

35

Yes

34

J. C. Huang*, J. P. Chu, and J. S. C. Jang, “Recent Progress on Metallic Glasses in Taiwan”, vol. 17, Intermetallics, 2009, pp. 973-987.

4 / 66

32

Yes

35

H. M. Chen, J. C. Huang, S. X. Song, and T. G. Nieh, “Flow serration and shear-band propagation in bulk metallic glasses”, Appl. Phys. Lett., vol. 60, 2009, 141914.

10 / 95

39

Yes

SCI/SSCI Rank

Factor²