Mechanical and Fatigue Behavior of Ca 65 Mg 15 Zn 20 Bulk-Metallic Glass

CO
DOI: 10.1002/adem.200800313
MM

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. Morrison
ON

[*] 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).

ADVANCED ENGINEERING MATERIALS 2009, 11, No. 1--2 � 2009 WILEY-VCH

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

Verlag GmbH & Co. KGaA, Weinheim 27

COM

MUNIC

ATIO

N

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

KGaA, Weinheim ADVANCED ENGINEERING MATERIALS 2009, 11, No. 1--2

COM

MU


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

NIC

ATIO

N

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.

ADVANCED ENGINEERING MATERIALS 2009, 11, No. 1--2 � 2009 WILEY-VCH Ve

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.

rlag GmbH & Co. KGaA, Weinheim http://www.aem-journal.com 29

COM

MUNIC

ATIO

N


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.


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.


COM

MUNIC

ATIO

N


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


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-

rlag GmbH & Co. KGa

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

A, Weinheim http://www.aem-journal.com 31

COM

MUNIC

ATIO

N


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


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


COM

MUNIC

ATIO

N


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,


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


[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,


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.


Mechanical and Fatigue Behavior of Ca 65 Mg 15 Zn 20 Bulk-Metallic Glass

Documents

Transcript of Mechanical and Fatigue Behavior of Ca 65 Mg 15 Zn 20 Bulk-Metallic Glass