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DOI: 10.1002/adem.200800313MM
UNI
Mechanical and Fatigue Behavior of Ca65Mg15Zn20Bulk-Metallic Glass**
CATI
By Gongyao Wang*, Peter K. Liaw, Oleg N. Senkov, Daniel B. Miracle and Mark L. MorrisonON
[*] Dr. G. Wang, Prof. P. K. Liaw, Dr. M. L. MorrisonDepartment of Materials Science and Engineering, TheUniversity of Tennessee Knoxville, TN 37996, USAE-mail: [email protected]
Dr. O. N. SenkovUES, Inc., DaytonOH 45432-1894, USA
Dr. D. B. MiracleAir Force Research Laboratory, Materials and ManufacturingDirectorateWright-Patterson AFB, OH 45433, USA
[**] Wewould like to acknowledge thefinancial support of theNationalScienceFoundation: theDivisionof theDesign,Manufacture, andIndustrial Innovation Program, under GrantNo. DMI-9724476;the Combined Research-Curriculum Development (CRCD) Pro-grams, under EEC-9527527 and EEC-0203415; the IntegrativeGraduate Education and Research Training (IGERT) Program,under DGE-9987548; the International Materials Institutes(IMI) Program, under DMR-0231320; and the Major ResearchInstrumentation (MRI) Program, under DMR-0421219, to theUniversity of Tennessee, Knoxville, withDr.D.Durham,Ms.M.Poats,Dr.C. J.VanHartesveldt,Dr. J.Giordan,Dr.D.Dutta,Dr.W. Jennings,Dr.L.Goldberg,Dr.C.Huber, andDr.C.R.Bouldinas Program Directors, respectively. Work at the Air ForceResearch Laboratory (AFRL) was conducted through the AFRLon-site contract No. FA8650-04-D-5233 and through anAFOSRTask (01ML05-COR, Dr. J. Fuller, Program Manager).
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A number of BMGs, such as Zr-, Fe-, Pd-, Al-, and Ni-based
alloys, have been discovered after the rapid development of
glass-forming alloys during the early 1990s.[1–4] Ca-Mg-Cu
and Ca-Mg-Cu-Ag BMGs were successfully fabricated by
Amiya and Inoue in 2002.[5,6] Following these reports,
numerous Ca-based BMG systems have been produced and
studied.[5–21] Ca-based BMG alloys are of interest because of
their unique properties, such as low density (�2.0 g cc�1), low
Young’s modulus (�17–20 GPa) that is comparable to
the modulus of human bones, low glass-transition tempera-
ture (Tg� 100 8C) and a wide super-cooled liquid temperature
range (DTxg¼Tx�Tg� 30–80 8C).[15,18] Elements such as Ca,
Mg, and Zn are biocompatible, which makes the
Ca-Mg-Zn-based alloys attractive for use in biomedical
applications.[18]
The amorphous structure gives unique properties to
BMGs, including high elastic strain, high fracture strength,
and high fatigue resistance. Although the mechanical
behavior of BMGs is studied widely,[1–4,22,23] there is no
fatigue data for Ca-based BMGs. A comprehensive under-
standing based on the compression, hardness, and fatigue
behavior is critically important for the application of the
Ca-based BMGs. In the current paper, the compression
behavior, Vickers hardness, and fatigue characteristics
of Ca65Mg15Zn20 BMGs were investigated at room tempera-
ture in air. A mechanistic understanding of the
fatigue and fracture mechanisms of Ca-based BMGs is
proposed.
Experimental
The Ca65Mg15Zn20 (atomic percent, at%) BMG alloy was
fabricated by induction melting pure elements (99.9 wt%)
with a water-cooled copper crucible in an argon atmosphere.
The prepared alloy was subsequently placed in a quartz
crucible with a 2 mm diameter hole at the bottom, induction
melted in an argon atmosphere, and injected into a water-
cooled copper mold with a 15� 15� 4 mm3 cavity. Previous
studies demonstrated that the critical thickness of this alloy,
below which it is fully amorphous, is 6 mm [9,15,18]. X-ray
diffraction and differential scanning calorimetry (DSC)
analyses indeed confirmed the fully amorphous state of
the produced 4-mm-thick plates. Thermal properties of the
cast alloys were determined using a DSC Q1000 differen-
tial-scanning calorimeter (TA Instruments Inc., New Castle,
DE) at a heating rate of 20 K min�1. The weight of the DSC
samples was in the range of 8–15 mg. The DSC results
exhibited that this Ca65Mg15Zn20 BMG had a very low Tg of
91 -C.[18]
The ingots were cut into 4� 4� 4 mm3 samples for
compression and fatigue experiments. All samples were
polished to avoid surface effects. Each side of these samples
was polished to a 600-SiC-grit-surface finish parallel to the
longitudinal axis of the specimens using a polishing fixture
(South Bay Technologies, San Clemente, CA) to keep the sides
parallel and perpendicular.
A computer-controlled Material Test System (MTS Systems
Corporation, Eden Praire, MN) servohydraulic-testing
machine was employed to study these samples. The load
frame was aligned prior to use. The compression experiments
were performed at room temperature with strain rates of 10�4,
10�3, and 10�2 s�1. Four to eight specimens were compression
tested at each of the three strain rates. Load-controlled fatigue
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G. Y. Wang et al./Mechanical and Fatigue Behavior of Ca65Mg15Zn20 Bulk-Metallic Glass
Fig. 2. Powder-like fractured Ca65Mg15Zn20 sample after compression experiments.
Fig. 3. Representative compression engineering stress–strain curves of theCa65Mg15Zn20 BMG at strain rates of 10�4, 10�3, and 10�2 s�1.
tests were conducted at various stress ranges with an R ratio
(R¼ smin/smax, where smin and smax are the applied
minimum and maximum stresses, respectively) of 0.1 using
a sinusoidal waveform at a frequency of 10 Hz. Tungsten-
carbide spacers were employed above and below the speci-
men to prevent the deformation of the pushrods during the
compression and fatigue experiments. Each fatigue test was
continued until the sample failed or achieved ‘‘runout’’ at 106
cycles.
Vickers hardness was measured using a Buehler MMT-3
digital microhardness tester (Buehler Ltd., Lake Bluff, IL) and
a diamond pyramidal Vickers indenter with various loads
(100, 300, and 500 g). The sample surfaces for the hardness
tests were prepared with 1200 grit SiC paper. The indents were
examined after unloading by SEM. A Leo 1526 scanning-
electron microscopy (SEM) (LEO Electron Microscopy Ltd.,
Cambridge, UK) with the energy-dispersive spectroscopy was
used to examine the surfaces of selected specimens to provide
a mechanistic understanding of the fatigue and fracture
behavior of the Ca-based BMG.
Results
Compression Behavior
Many thin pieces were shed progressively from specimen-
free surfaces during the compression tests. Significant
spallation typically occurred after the applied load reached
about 60% of the fracture load. Macroscopically, the planes of
these fracture surfaces were roughly parallel to the loading
axis (Fig. 1). This splitting fracture reduced the remaining
cross-sectional area significantly, so that the strengths
reported here are nominal values of applied load divided
by the initial cross-sectional area. The actual strengths are
expected to be higher than those reported here. In the final
catastrophic failure event, the samples fractured into numer-
ous small pieces (Fig. 2).
Figure 3 shows typical compression stress–strain curves of
amorphous Ca65Mg15Zn20 at initial strain rates of 10�4, 10�3,
Fig. 1. (a)Cracks initiating fromapore-likeflaw in the splitting-fracturemode and (b) thecracks propagating and resulting in a splitting failure under a compression loading.
28 http://www.aem-journal.com � 2009 WILEY-VCH Verlag GmbH & Co.
and 10�2 s�1. The compression test results are summarized in
Table 1. All the samples exhibited elastic deformation behavior
and catastrophic fracture without plasticity. The nominal
fracture strength, sf, increased from 300 to 409 MPa with an
increase in the applied strain rate from 10�4 to 10�2 s�1. The
average nominal fracture strength was 364 MPa, which is
similar to the compressive fracture strength reported by
Senkov et al.[15] for Ca47Mg19Zn7Cu27 at a strain rate of 10�4 s�1.
The total elastic strain before failure increased from 1.4 to 2.1%
when the applied strain rate increased from 10�4 to 10�2 s�1.
Table 1. Results from compression experiments of the Ca65Mg15Zn20 BMG alloy atvarious strain rates.
Strainrate [s�1]
Young’smodulus[GPa]
Fracturestrength[MPa]
Fracturestrain[%]
10�4 22.3 300 1.36
10�3 18.2 383 2.09
10�2 19.6 409 2.10
Average 20.0� 2.3 364� 64 1.85� 0.49
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The Young’s modulus ranged from 18.2 to 22.3 GPa at the
applied strain rates with a mean elastic modulus (E) of
20.0 GPa, which is comparable to the measured value of
26.4 GPa by resonant ultrasound spectroscopy.[19] This value is
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Fig. 4. Fracture-surface morphology of the Ca65Mg15Zn20 BMG after a compressionexperiment, including (a) a splitting-fracture surface, (b) a conchoidal pattern, (c) therock-layer pattern and (d) vein patterns. The solid arrows indicate the crack-growthdirections.
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also in agreement with the Young’s modulus reported for
Ca65Ag35.[5] The increased in the slope of the stress–strain
curves above about 60% of the maximum load achieved is due
to the reduction in cross-sectional area that results from
spallation.
In order to characterize the fracture behavior, a controlled
test was stopped immediately after major cracks were
observed but the sample did not totally fail. The spalled
fragments displayed fracture planes that locally appear to be
inclined by about 0–208 to the direction of the applied load.
SEM observation of the splitting-fracture surfaces reveals a
mixture of rough and flat fracture zones, as shown in
Figure 4a. A conchoidal fracture consists of tear lines and
mirror surfaces (Fig. 4b), and is generally observed in brittle
materials, including Fe-based BMGs.[24–26] A rough fracture
zone exhibits rock-layer patterns, as seen in Figure 4c, which
are similar to compression-fracture surfaces of a Zr60Ti5Cu15-
Ni10Al10 BMG composite.[26] The mixture of rough and flat
fracture zones is presumably due to the easy initiation of
fracture at many flaws.
In addition to splitting, shear fracture was observed locally
on the fracture surfaces. The shear-fracture surface was
relatively flat and displayed typical shear-fracture features,
such as a vein pattern (Fig. 4d). The vein structure has been
widely observed and is generally attributed to the significant
increase in the temperature in shear bands during the
deformation of metallic glasses.[27–32] The shear fracture
surface had a large angle of approximately 358 with respect
to the stress axis. This result is consistent with the reported
results for many other metallic glasses, which demonstrate
that the compressive fracture of metallic glasses does not
occur along the plane of the maximum shear stress, and the
compressive fracture angle is less than 458.[27–30]
Hardness
Figure 5 shows the dependence of the Vickers hardness, Hv,
on the indentation load, P. Under applied loads of 100, 300,
and 500 g, Hv values were 1.47, 1.41, and 1.38 GPa, respec-
tively. The trend for Hv to decrease with increasing P is
Fig. 5. A plot of the Vickers hardness as a function of the applied load for theCa65Mg15Zn20 BMG.
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Fig. 6. SEM images of a Vickers indentation in Ca65Mg15Zn20 demonstrating (a) theindent morphology at a load of 300 g, and (b) shear bands (indicated by arrows) withinand around the indent. The white spots are oxides due to polishing.
consistent with results for other BMGs.[33–35] The mean
hardness of the Ca65Mg15Zn20 BMG for the three applied
loads in this study was 1.42 GPa.
Figure 6 shows a representative Vickers indentation
obtained on the Ca65Mg15Zn20 BMG using a 300 g load.
Pile-up of material near the edges of the indenter and
semi-circular and wavy shear bands that emanate from the tip
of the indentation are shown in Figure 6b. These features
confirm a noticeable local plasticity of the alloy under highly
constrained compressive loading, in spite its extreme macro-
scopic brittleness. Similar indent morphologies are observed
in other BMGs.[35–42] The discrete slip steps in the pile-up
around the indent are generally attributed to the inhomoge-
neous nature of the plastic deformation. Shear-band propaga-
tion in metallic glasses is inhomogeneous, which results in
discrete deformation bands within the deformed zone.[39] The
deformation in elastic-perfectly plastic solids generally occurs
by the pile-up of the material against the faces of the
indenter.[35]
Fatigue Behavior
The fatigue lifetime increased from 104 to 106 cycles when
the maximum applied stress decreased from about 240 to
140 MPa (Fig. 7). The fatigue data, which exhibited a fatigue
Fig. 7. A plot of the maximum stress as a function of the fatigue lifetime for theCa65Mg15Zn20 BMG alloy under compression–compression fatigue loading.
30 http://www.aem-journal.com � 2009 WILEY-VCH Verlag GmbH & Co.
limit of 140 MPa at 106 cycles, appears to be highly variable.
Similar to the monotonic-compression testing, many small
pieces broke off during fatigue experiments. Upon the final
fracture, the samples shattered into very small fragments.
Fatigue–fracture surfaces demonstrated a mixture of fracture
features similar to compression experiments (Fig. 8). Vein
patterns (Fig. 8a) have been seen previously after both
monotonic[27–32] and cyclic loading.[43,44] These results suggest
that the fracture behavior under the compression–
compression fatigue loading is analogous to the monotonic
compressive fracture. A similar conclusion was made earlier
for Zr-based BMGs.[44,45] On the other hand, the fatigue–
fracture behavior of the Ca65Mg15Zn20 BMG under compres-
sion–compression loading is more complex than that of
Fig. 8. Fracture-surfaces of Ca65Mg15Zn20 subjected to compression-compressionfatigue loading, including (a) vein patterns, (b) conchoidal patterns, and (c) rock-layerpatterns. The solid arrows indicate crack-growth directions.
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Zr-based BMGs where the fracture occurs only along specific
shear planes.[44–46]
Discussion
Brittleness is a common feature of many metallic glasses,
including BMGs such as Zr-, Cu-, Mg-, and Pd-based BMGs, at
temperatures well below the glass-transition temperature.[47]
Their brittleness is generally explained by limited amounts of
deformation carriers, such as linear and planar defects, to
accommodate the loading conditions, as well as the absence of
strain hardening in the amorphous structures. Due to these
constraints, deformation by tension is generally localized in a
single shear band, leading to a premature fracture along this
band. To slow down the premature sample failure and be able
to analyze the deformation behavior, soft modes of deforma-
tion are applied to amorphous metals, such as bending of thin
samples (ribbons)[47] or compression of samples with a low
aspect ratio of �1,[47,48] which is smaller than an aspect ratio of
2–3 typically used for the compression study of crystalline
materials. Although in this work we also used compression
loading of samples with the aspect ratio of 1, the samples
showed extreme brittleness, and they exploded into very
small fragments upon failure in the monotonic loading and
high-cycle fatigue compression tests, vividly illustrating the
inherent brittleness of Ca65Mg15Zn20. The splitting fracture
observed in the present work has not been reported
previously for metallic glasses, and it may be influenced by
the possible presence of multi-axial stresses.
Fracture surfaces exhibited a dominant splitting-fracture
mode that is similar to the fracture of brittle nonmetallic
materials (ceramics, rocks, and glasses), and a secondary
shear-fracture mode that is similar to the fracture behavior of
other BMGs, such as Zr-, Cu-, and Pd-based BMGs.[27–30]
Splitting fracture has not been observed previously in
metallic glasses, although it is common in brittle nonmetallic
solids under uniaxial compression.[49–52] Splitting failure
normally involves a sequence of progressive microfracturing
events,[50,51] where microcracks originate at inhomogeneities
such as flaws and inclusions (Fig. 1a) at stresses much lower
than the fracture stress, and propagate with increasing load
to form a macroscopic splitting failure (Fig. 1b). In the present
Ca-based BMG, pore-like flaws were easily observed on
ig. 9. SEM fractographs demonstrating (a) a pore-like flaw, microcrack, and crack on the fracture surface, and) a pore-like flaw, which resulted in macroscopic splitting failure under compression.
F(b
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fracture surfaces (Fig. 9). In Figure 9a, a microcrack can be seen
that initiated from a pore and propagated a short distance.
These cracks typically propagate roughly parallel to the
loading direction. Thus, macroscopic splitting failure under
compression likely resulted from flaws (Fig. 9b). As a result,
cracks near free surfaces can easily result in a buckling failure
near the specimen surfaces.[50,51] Another possible explana-
tion of the failure follows.
Zhang et al.[53] suggested that during compression
deformation, the critical shear stress, tC, responsible for
development of a shear band can be expressed as:
tC¼t0 þ mCsC (1)
where t0 is the critical shear stress in the shear plane in the case
of a pure shear condition, sC the normal stress on the shear
plane, andmC is a constant. Using Equation (1) the compressive
fracture strength sCF can be deduces as:[53]
sCF ¼ 2t0ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi
1 þ ðmCÞ2q
� mC
(2)
and the compressive shear-fracture angle, uC, can be expressed
as:
uC ¼ arctan
ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi1 þ ðmCÞ2
q�mC
� �(3)
It can be seen from Equation (1) and (3) that when mC¼ 0
(i.e., when there is no effect of sC on tC), then uC¼ 458. For
most BMG materials, mC is above zero and thus uC< 458. For
example, uC is approximately 428 for a Zr55Cu30Al10Ni5BMG,[53] giving mC¼ 0.106. For the Ca65Mg15Zn20 BMG
(current work), uC is about 358 for the shear-fracture mode,
and about 108 for the splitting-fracture mode, which gives
mC� 0.364 and �2.75, respectively. Unfortunately, the phy-
sical meaning of mC and why this parameter should be
different for different metallic glasses were not discussed in
ref.[53], where this parameter was introduced.
Our interpretation of the compositional dependence of mC
follows. From definition of mC (Eq. 1), this parameter is
equivalent to the coefficient of friction along the shear plane.
For the amorphous materials, which do not have long-range
order and crystal planes, the shear plane can be micro-
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scopically wavy. This microscopic roughness
can be different in different glasses leading to
a different friction coefficient mC. Higher
atomic size differences between the alloying
elements and/or larger atomic clusters
formed in the glassy material are expected
to provide higher shear plane roughness and,
therefore, higher value of mC. If this inter-
pretation is correct, then mC should be higher
and uC should be smaller in amorphous
materials with microscopically rougher shear
plane surfaces, i.e., in BMGs with larger
atomic size differences and/or larger atomic
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G. Y. Wang et al./Mechanical and Fatigue Behavior of Ca65Mg15Zn20 Bulk-Metallic Glass
Fig. 10. (a) Fracture strength versus applied strain rate and (b) fracture strength vs. thefracture time for the Ca65Mg15Zn20 BMG.
Table 2. The mechanical properties of the Ca65Mg15Zn20 BMG.
Hv [GPa] Hv/sY E/sY E/Hv
1.42 3.90 55 14
clusters. When the plane roughness is extremely high and mC
becomes higher than 2 (uC< 138), then no shear but splitting
should occur at a very low angle to the load axis.
The nominal fracture strength, sC, increased from 300 to
383 MPa when the applied strain rate increased from 10�4 to
10�3 s�1. However, when the applied strain rate changed from
10�3 to 10�2 s�1, the nominal fracture strength only increased
slightly, as seen in Figure 10(a). In addition, the nominal
fracture strength versus fracture time, tf, is plotted in
Figure 10b. This plot demonstrates that the nominal fracture
strength of the Ca65Mg15Zn20 BMG decreases with increase in
fracture time of the samples. The relationship between the
nominal fracture strength and the fracture time can be fitted
by the following equation (R2¼ 0.99):
sf ¼ 405� 0:78tfðMPaÞ (4)
This trend of the Ca65Mg15Zn20 BMG is much different
from other BMGs.[52–55] For example, Hufnagel et al.[54] found
that the compressive fracture strength of a Zr57Ti5Cu20Ni8Al10
BMG decreased with increase in the strain rate. Similarly,
Mukai et al.[55] showed a reduction in the fracture strength
when the strain rate increased during the compression
experiments with the Pd40Ni40P20 metallic glass. In contrast,
the compression fracture strength of the bulk Zr41.25Ti13.75-
Cu12.5Ni10Be22.5 glass seems to be independent of the strain
rate.[56,57] The above differences in the fracture strength versus
32 http://www.aem-journal.com � 2009 WILEY-VCH Verlag GmbH & Co.
strain rate behavior between the Ca65Mg15Zn20 BMG and
other BMGs could be attributed to the differences in the
fracture mechanisms.
For transition metal-based BMG alloys, such as the
Zr-based and Pd-based systems, the inhomogeneous defor-
mation at room temperature usually occurs by the formation
of shear bands that operate rapidly and accommodate
displacements apparently up to nearly the millimeter scale.[58]
Thus, based on the shear-band model, the temperature and
strain-rate dependence of the strength can be modeled by an
equation of the form:[47]
t ¼ ts � ts D � ln gs
g
� �T
Tg
� �1=2(5)
where ts is the athermal stress required to initiate the
critical-shear event, D a dimensionless constant, gs a
characteristic-strain rate, g a shear–strain rate, and T is the
temperature. Based on Equation (5), it is found that the
inhomogeneous strength of metallic glasses exhibits a very
weak scaling with the strain rate. Therefore, Schuh et al.[47]
theorized that metallic glasses exhibited an essentially rate-
independent strength and some softening phenomena due to
the fact that the strain rates were associated with the adiabatic
heat generation. Nevertheless, for the current simple-metal
Ca-based BMG, the deformation included not only a shear
mode but also a splitting mode. Moreover, we found that the
splitting mode is a dominant failure mode in this Ca-based
BMG. In fact, the splitting process appears to account for the
crack initiation, crack propagation, crack linkage, and final
failure. This process is closely related to the test time.
Furthermore, the interaction of cracks with the surfaces causes
them to grow rapidly and results in a buckling failure near the
specimen surfaces.[50,51] This fact is supported by the test
behaviors observed in this study. During the compression and
fatigue experiments of the Ca-based BMG, small pieces of the
samples fractured from the surfaces prior to the final failure,
which could result in a reduction of the cross-sectional area.
As a result, the specimen failed at a higher engineering stress.
Likewise, this process is directly proportional to the test time.
Therefore, the apparent fracture strength of the Ca-based
BMG was inversely proportional to the fracture time. If this
explanation is correct, one may expect a true nominal fracture
strength above 409 MPa (the nominal strength at a strain rate
of 10�2 s�1) for this BMG.
From the slip-line-field theory for rigid-perfectly plastic
materials, the hardness, H, is directly related to the yield
stress, sy, through the simple linear relationship, H¼ 3sy.[59]
This equation has been validated for metallic glasses from the
nature of the glass plasticity and mechanics-based treatments
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Fig. 11. The ratio of the maximum applied stress during fatigue to the compressivefracture strength versus fatigue–lifetime plot of Ca65Mg15Zn20 BMGs,Zr50Cu37Al10Pd3 [44], and Zr41.2Cu12.5Ni10Ti13.8Be22.5
[46] under compression–compression fatigue loading.
of the indentation.[60] For the current Ca65Mg15Zn20 BMG, the
proportionality constant K (¼H/sF) is equal to 3.90 (Table 2).
This is consistent with other metallic glasses, where
3�K� 4.5.[60,61] Similarly, E/sF and E/Hv for the current
Ca65Mg15Zn20 BMG are 55 and 14, respectively (Table 2), and
these fall within the ranges 45�E/sF� 70 and 13�E/Hv � 20
that are typical for BMGs.[62]
The fatigue fracture did not occur only along a specific
shear plane in the Ca65Mg15Zn20 BMG under the compres-
sion–compression loading, which is different from the
previous fatigue results for Zr-based BMGs with a shear
fracture angle to the loading axis close to 458. This trend
results from the brittle nature of the Ca65Mg15Zn20 BMG and
may also be influenced by the small aspect ratio of the test
samples. It is evident that the splitting mode is a dominant
failure mode in the present Ca-based BMG during the fatigue
experiments. Thus, crack initiation from flaws, crack propa-
gation that is roughly parallel to the loading direction, crack
linkage, and final failure form the splitting fracture process in
the Ca65Mg15Zn20 BMG. In Figure 11 the ratio of the maximum
applied stress during fatigue to the monotonic compressive
fracture strength versus fatigue lifetime is plotted for the
current Ca-based BMG, Zr50Cu37Al10Pd3,[44] and Zr41.2Cu12.5-
Ni10Ti13.8Be22.5[46] under compression–compression fatigue
loading. In general, the Ca65Mg15Zn20 BMG exhibits much
lower lifetimes than those of the Zr-based BMGs.
Conclusions
The deformation behavior was characterized for a
Ca65Mg15Zn20 bulk metallic glass during monotonic compres-
sion, Vickers indentation hardness and compression–
compression fatigue behavior. The alloy was macroscopically
brittle, and demonstrated a dominant splitting failure with
fracture planes from 0 to 208 relative to the loading axis.
Secondary shear fracture was also observed in local regions,
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with a fracture plane roughly 358 relative to the loading
direction.Thenominalcompressive fracturestrength increased
with an increase in the applied strain rate and was inversely
proportional to the testing time. The indentation hardness
decreased with an increase in the indentation load. The typical
indentation morphology showed pile-ups near the indenter
edges with associated shear bands. Although the fatigue data
were somewhat scattered, a fatigue limit of 140 MPa was found
at 106 cycles. Compared to Zr-based BMGs, the Ca65Mg15Zn20
BMG demonstrated much shorter fatigue lives and a lower
fatigue endurance limit.
Received: September 20, 2008
Revised: October 29, 2008
[1] W. L. Johnson, MRS Bull. 1999, 24, 42.
[2] A. Inoue, Acta Mater. 2000, 48, 279.
[3] W. L. Johnson, JOM 2002, 54, 40.
[4] A. Inoue, A. Takeuchi, Mater. Sci. Eng. A 2004, 375–377,
16.
[5] K. Amiya, A. Inoue, Mater. Trans. JIM 2002, 43, 81.
[6] K. Amiya, A. Inoue, Mater. Trans. JIM 2002, 43, 2578.
[7] O. N. Senkov, J. M. Scott, MRS Proc. Mater. Res. Soc.
Warrendale, PA 2003, 806, 145.
[8] O. N. Senkov, J. M. Scott, Mater. Lett. 2004, 58, 1375.
[9] O. N. Senkov, J. M. Scott, J. Non-Cryst. Solids 2005, 351,
3087.
[10] E. S. Park, W. T. Kim, D. H. Kim, Mater. Sci. Forum 2005,
475–479, 3415.
[11] O. N. Senkov, J. M. Scott, Scr. Mater. 2004, 50, 449.
[12] F. Q. Guo, S. J. Poon, G. J. Shiflet, Appl. Phys. Lett. 2004,
84, 37.
[13] O. N. Senkov, J. M. Scott, D. B. Miracle, J. Alloys Compd.
2006, 424, 394.
[14] S. Gorsse, G. Orveillon, O. N. Senkov, D. B. Miracle,
Phys. Rev. B 2006, 73, 224202.
[15] O. N. Senkov, D. B. Miracle, J. M. Scott, Intermetallics
2006, 14, 1055.
[16] E. S. Park, D. H. Kim, Appl. Phys. Lett. 2005, 86, 201912.
[17] E. S. Park, D. H. Kim, J. Mater. Res. 2004, 19, 685.
[18] M. L. Morrison, R. A. Buchanan, O. N. Senkov, D. B.
Miracle, P. K. Liaw, Metall. Mater. Trans. A 2006, 37, 1239.
[19] Z. Y. Zhang, V. Keppens, O. N. Senkov, D. B. Miracle,
Mater. Sci. Eng. A 2007, 471, 151.
[20] O. N. Senkov, J. M. Scott, D. B. Miracle, Mater. Trans.
2007, 48, 1610.
[21] O. N. Senkov, D. B. Miracle, V. Keppens, P. K. Liaw,
Metall. Mater. Trans. A 2008, 39, 1888.
[22] H. Li, C. Fan, K. Tao, H. Choo, P. K. Liaw, Adv. Mater.
2006, 18, 752.
[23] C. Fan, H. Li, L. J. Kecskes, K. Tao, H. Choo, P. K. Liaw,
C. T. Liu, Phys. Rev. Lett. 2006, 96, 145506.
[24] X. J. Gu, S. J. Poon, G. J. Shiflet, J. Mater. Res. 2007, 22, 344.
rlag GmbH & Co. KGaA, Weinheim http://www.aem-journal.com 33
COM
MUNIC
ATIO
N
G. Y. Wang et al./Mechanical and Fatigue Behavior of Ca65Mg15Zn20 Bulk-Metallic Glass
[25] D. C. Qiao, G. Y. Wang, P. K. Liaw, V. Ponnambalam, S.
J. Poon, G. J. Shiflet, J. Mater. Res. 2007, 22, 544.
[26] G. He, W. Loser, J. Eckert, L. Schultz, Mater. Sci. Eng. A
2003, 352, 179.
[27] Z. F. Zhang, J. Eckert, L. Schultz, Acta Mater. 2003, 51,
1167.
[28] W. H. Jiang, G. J. Fan, H. Choo, P. K. Liaw, Mater. Lett.
2006, 60, 3537.
[29] D. Xu, B. Lohwongwatana, G. Duan, W. L. Johnson, C.
Garland, Acta Mater. 2004, 52, 2621.
[30] P. E. Donovan, Mater. Sci. Eng. 1988, 98, 487.
[31] J. J. Lewandowski, A. L. Greer, Nat. Mater. 2006, 5,
15.
[32] B. Yang, C. T. Liu, T. G. Nieh, M. L. Morrison, P. K. Liaw,
R. A. Buchanan, J. Mater. Res. 2006, 21, 915.
[33] H. Zhang, X. Jing, G. Subhash, L. J. Kecskes, R. J.
Dowding, Acta Mater. 2005, 53, 3849.
[34] F. Yang, K. Geng, P. K. Liaw, G. Fan, H. Choo, Acta
Mater. 2007, 55, 321.
[35] U. Ramamurty, S. Jana, Y. Kawamura, K. Chattopad-
hyay, Acta Mater. 2005, 53, 705.
[36] J. J. Kim, Y. Choi, S. Suresh, A. S. Argon, Science 2002,
295, 654.
[37] Y. F. Gao, B. Yang, T. G. Nieh, Acta Mater. 2007, 55,
2319.
[38] W. H. Jiang, F. E. Pinkerton, M. Atzmon, J. Appl. Phys.
2003, 93, 9287.
[39] S. Jana, U. Ramamurty, K. Chattopadhyay, Y. Kawa-
mura, Mater. Sci. Eng. A 2004, 375–377, 1191.
[40] R. Vaidyanathan, M. Dao, G. Ravichandran, S. Suresh,
Acta Mater. 2001, 49, 3781.
[41] V. Keryvin, Acta Mater. 2007, 55, 2565.
[42] N. K. Mukhopadhyay, A. Belger, P. Paufler, D. H. Kim,
Mater. Sci. Eng. A 2007, 449–451, 954.
[43] G. Y. Wang, P. K. Liaw, W. H. Peter, B. Yang, Y.
Yokoyama, M. L. Benson, B. A. Green, M. J. Kirkham,
34 http://www.aem-journal.com � 2009 WILEY-VCH Verlag GmbH & Co.
S. A. White, T. A. Saleh, R. L. McDaniels, R. V. Steward,
R. A. Buchanan, C. T. Liu, C. R. Brooks, Intermetallics
2004, 12, 885.
[44] D. C. Qiao, G. Y. Wang, W. H. Jiang, Y. Yokoyama, P. K.
Liaw, H. Choo, Mater. Trans. JIM 2007, 48, 1828.
[45] Z. F. Zhang, J. Eckert, L. Schultz, Metall. Mater. Trans. A
2004, 35, 3489.
[46] P. A. Hess, B. C. Menzel, R. H. Dauskardt, Scr. Mater.
2006, 54, 355.
[47] C. A. Schuh, T. C. Hufnagel, U. Ramamurty, Acta Mater.
2007, 55, 4067.
[48] W. H. Jiang, G. J. Fan, H. Choo, P. K. Liaw, Mater. Lett.
2006, 60, 3537.
[49] M. F. Ashby, S. D. Hallam, Acta Metall. 1986, 34, 497.
[50] C. G. Sammis, M. F. Ashby, Acta Metall. 1986, 34, 511.
[51] C. A. Tang, R. H. C. Wong, K. T. Chau, P. Lin, Eng. Fract.
Mech. 2005, 72, 597.
[52] R. H. C. Wong, C. A. Tang, K. T. Chau, P. Lin, Eng. Fract.
Mech. 2002, 69, 1853.
[53] Z. F. Zhang, G. He, J. Eckert, L. Schultz, Phys. Rev. Lett.
2003, 91, 045505.
[54] T. C. Hufnagel, T. Jiao, Y. Li, L. Q. Xing, K. T. Ramesh, J.
Mater. Res. 2002, 17, 1441.
[55] T. Mukai, T. G. Nieh, Y. Kawamura, A. Inoue, K. Higa-
shi, Intermetallics 2002, 10, 1071.
[56] H. A. Bruck, A. J. Rosakis, W. L. Johnson, J. Mater. Res.
1996, 11, 503.
[57] G. Subhash, R. J. Dowding, L. J. Kecskes, Mater. Sci. Eng.
A 2002, 334, 33.
[58] H. Kimura, T. Masumoto, Acta Metall. 1983, 31, 231.
[59] D. Tabor, The Hardness of Metals, Oxford University
Press, London, UK 1951.
[60] C. A. Schuh, T. G. Nieh, J. Mater. Res. 2004, 19, 46.
[61] M. N. M. Patnaik, R. Narasimhan, U. Ramamurty, Acta
Mater. 2004, 52, 3335.
[62] W. H. Wang, J. Non-Cryst. Solids 2005, 351, 1481.
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