INVITED FEATURE REVIEW Damage-tolerant Zr Cu Al … · · 2014-08-18INVITED FEATURE REVIEW Damage-tolerant Zr–Cu–Al-based bulk metallic glasses with record-breaking fracture

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INVITED FEATURE REVIEW

Damage-tolerant Zr–Cu–Al-based bulk metallic glasses withrecord-breaking fracture toughness

Jian Xua)

Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences,Shenyang 110016, China

Evan Mab)

Department of Materials Science and Engineering, The Johns Hopkins University, Baltimore, Maryland21218, USA

(Received 28 February 2014; accepted 4 June 2014)

Bulk metallic glasses (BMGs) exhibit high yield strength but little tensile ductility. For this classof materials, damage tolerance is a key mechanical design parameter needed for their engineeringuse. Recently we have discovered a correlation between the local structural characteristics in theglass and the propensity for shear transformations. Based on the dependence of glass structure onalloy composition, zirconium (Zr)-rich Zr–titanium (Ti)–copper (Cu)–aluminum (Al) compositionsare predicted to be more prone to spread-out plastic deformation and hence profuse shear banding.This structural perspective has guided us to locate a Zr61Ti2Cu25Al12 (ZT1) BMG that exhibitsa record-breaking fracture toughness, on par with the palladium (Pd)-based BMG recentlydeveloped at Caltech. At the same time, the new BMG consists of common metals and has robustglass-forming ability. Interestingly, the ZT1 BMG derives its high toughness from its highpropensity for crack deflection and local loading-mode change (from mode I to substantially mode II)at the crack tip due to extensive shear band interactions. A crack-resistance curve (R-curve) has beenobtained following American Society for Testing and Materials (ASTM) standards, employing both“single-specimen” and “multiple-specimen” techniques as well as fatigue precracked specimens. Thecombination of high strength and fracture toughness places ZT1 atop all engineering metallic alloys inthe strength–toughness Ashby diagram, pushing the envelop accessible to a structural material interms of its damage tolerance.

I. INTRODUCTION

Bulk metallic glasses (BMGs) are currently investi-gated worldwide, as a new class of promising structuralmaterials. Different from conventional crystalline metal-lic alloys, the internal structures of BMGs are amorphous,and therefore they do not have well-defined dislocationdefects to mediate plastic deformation. As a result, theyexhibit extraordinarily high strength, but lack tensileductility. This brings in an important difference, in termsof the key properties that are needed for engineeringdesign and practical use. For conventional alloys, yieldstrength is often the design parameter used in materialsselection; but for BMGs, fracture toughness is a moretell-tale indicator of mechanical performance, needed indesign to guarantee damage tolerance. The fracturetoughness assesses a material’s resistance to crack prop-agation, measured by the energy needed to cause fracture.

Simultaneous presence of high fracture toughness andstrength imparts high damage tolerance to an alloy.This article discusses a strategy to locate BMGs with

high fracture toughness. In this regard, the first questionto ask is whether a BMG has the potential to be highlydamage-tolerant. The answer is positive because, differ-ent from oxide glasses, the nondirectional metallicbonding and dense-packed structures of BMGs offer theintrinsic mechanism of relatively easy shear transforma-tions (local relocation/shuffling of a group of atomsrelative to one another) to accommodate plastic strain,which shields crack tip and dissipates energy.1 Althoughunder uniaxial tension loading the plastic strains localizeinto thin shear bands, making BMGs highly susceptible tofailure, they are not necessarily brittle under confinedloading conditions such as bending, rolling, compression(especially when the sample has a low aspect ratio), andtriaxially stressed states. The challenge is how to findBMG compositions/structures that are particularly suscep-tible to shear transformations and profuse generation ofshear bands at the crack tip.Thus far, the fracture toughness values reported for

BMGs span a wide range.2–4 Very few monolithic BMGshave been found to be able to exhibit subcritical crack

Address all correspondence to these authors.a)e-mail: [email protected])e-mail: [email protected]: 10.1557/jmr.2014.160

J. Mater. Res., Vol. 29, No. 14, Jul 28, 2014 �Materials Research Society 2014 1489

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growth and substantial crack propagation resistance(in terms of measured R-curve). In general, this is due tothe lack of microstructural variations inside the BMGs(except in crystal-containing composites based onBMGs5,6), such as a high content of “liquid-like”heterogeneities in BMGs, to play the role of dislocationsin crystalline metals. However, an exciting break-through has recently been reported7: A monolithicBMG (Pd79Ag3.5P6Si9.5Ge2) was developed at Caltechand demonstrated to rival the best known crystalline alloysin damage tolerance. This Pd-based BMG has by far thehighest fracture toughness among all BMGs. Unfortunately,this BMG is compositionally based on the expensivepalladium and suffers from a rather limited glass-formingability (GFA). These severely limit this BMG for mass-production/application in bulk form. Also, the guidelinefor finding such high-toughness BMGs remains unset-tled. Demetriou et al.7 postulated that the high fracturetoughness of their Pd-based BMG is due to its highPoisson’s ratio m of 0.42 (or equivalently, a high ratio ofbulk modulus to shear modulus, B/G8). That connectiondrastically reduces the number of possible candidates,restricting the choices to alloy systems based on high-melements, which are few and tend to be of high cost. In thefollowing, we report an alternate route for the develop-ment of highly damage-tolerant BMGs, from a BMGstructural perspective.

II. EXPERIMENTS: ASTM STANDARD TESTS FORFRACTURE TOUGHNESS

In the experimental front, our study emphasizes twoissues regarding quantitative and appropriate fracturetoughness measurements. First, the test samples shouldhave sufficiently large dimensions, such that the fracturetoughness can be evaluated in a reliable way, usingAmerican Society for Testing and Materials (ASTM)standard tests. In comparison, small samples would leaveinadequate dimensions for the ligaments, and shear bandsnear the roller on the other side of the specimen in athree-point bending test would consume energy and couldhence be a source of uncertainties.7 Also, the toughnessmeasured can exhibit sample-size dependence, as planestrain conditions are no longer valid in small-dimensionsamples. Second, fatigue precrack is required for a standardtest following ASTM standard procedures, such that thework associated with crack nucleation does not come intoplay. In other words, we strive to employ adequately sizedspecimens to strictly follow ASTM standards, to improveon previous tests and minimize the possibility of observingand reporting artificially elevated toughness values.

Our BMG plates were larger than 3 mm in thicknessin their as-cast conditions,9 as shown in Fig. 1. Thealloy composition to be focused on in this article isZr61Ti2Cu25Al12 (denoted as ZT1), which is a monolithic

BMG without noticeable heterogeneities such as phaseseparation and nanocrystals. The oxygen content in theas-cast BMG plates was determined to be as low as 90–130weight parts per million (wppm) using inert-gas fusionanalysis. After machining and polishing, the toughnessspecimendimensions are at a ratio of 1:2:8 forB (thickness53 mm):W (width):S (span). A straight, through-thicknessnotchwasmade using a diamondwire saw, with a notch rootradius of 150 lm and a length of ;0.25W. The fatigueprecracking of the specimens was conducted on a 2.5 kNfatigue test machine (MTS Systems Corporation, 858 miniBionix, Eden Prairie, MN), as shown in Fig. 2, at the fre-quency of 30 Hz under a constant load ratio of minimumto maximum of 0.1 with the stress intensity factor DK of30–45 MPaOm. The notch plus the fatigue precrackamounted to a length of 0.5–0.7W after 10,000–100,000fatigue cycles, conforming toASTM standards. Three-pointbend tests of the fatigue precracked specimens werecarried out on a 5 kN Instron 8871 testing machine(Instron Corporation, Norwood, MA) at a constant

FIG. 1. As-cast Zr61Ti2Cu25Al12 BMG plates used for preparingsamples of fracture toughness measurement.

FIG. 2. Three-point-bending setup with single-edge-notchZr61Ti2Cu25Al12 BMG specimen mounted on MTS fatigue testmachine for fatigue precracking.

J. Xu et al.: Damage-tolerant Zr–Cu–Al-based bulk metallic glasses with record-breaking fracture toughness

J. Mater. Res., Vol. 29, No. 14, Jul 28, 20141490



displacement rate of 0.3 mm/min. The crack openingdisplacements (CODs) were monitored across the crackmouth using a clip gage, mounted between knife edges andaffixed to the front of the specimen.

To measure the crack resistance curve (R-curve), boththe “single-” and “multiple-” specimen techniques wereused.10,11 For the single-specimen technique, the fatigueprecracked specimen was first loaded monotonically at therate of 0.3 mm/min. In the nonlinear part of the load versusload line displacement (LLD) curve, observed after LLDof about 0.3–0.55 mm, the specimen was unloaded by20–30% of the current load at the rate of 0.15 mm/min,followed by reloading at 0.3 mm/min. The number ofload–unload sequences depends on the original cracklength (10–30 runs were used in our case). The length ofcrack extension (Da) can be estimated based on themeasured specimen compliance,

u ¼ 1

BWECi

S=4

h i1=2þ 1

; ð2Þ

where ai is the crack length and Ci is the changingcompliance during the sequence of multiple unloading/reloading. The eventual Da was calibrated by comparingthe compliance-derived value with the physical crackgrowth width sandwiched between the fatigue precrackedregion and refatigued region observed on the fracturesurface.

For the multiple-specimen technique, the first specimenwas unloaded at the appropriate displacement to produceDa in a desired position for the R curve. Subsequently,several specimens were unloaded at desired displacements(distributed uniformly), less than the displacement ofthe first specimen. After unloading, the specimens wereheat-tinted at 300 °C for 60–90 min, then refatigued ona 2.5 kN MTS fatigue test machine at the frequency of30 Hz under a constant load ratio (minimum to maximum)of 0.1 until the ultimate fracture. This way, the physicalcrack length in each fractured samples can be visuallydetermined by using a scanning electron microscope(SEM; Quanta model 600, FEI, Eindhoven, Netherlands)and a 3D measuring laser microscope (LEXT OLS4000,Olympus Corporation, Tokyo, Japan).

Using the following standard equations, the J value(nonlinear strain-energy release rate) was calculated fromload versus COD curves,

J ¼ K2 1� m2ð ÞE

þ Jpl ; ð3Þ

Jpl ¼ 1:9Apl

Bb; ð4Þ

K ¼ PS

BW3=2f

a

W

� �; ð5Þ

fa

W

� �¼ 3

ffiffiffiffiffia

W

r 1:99� aW

� �1� a

W

� �2:15� 3:93 a

W þ 2:7 aW

� �2h i

2 1þ 2 aW

� �1� a

W

� �3=2 ;

ð6Þ

where Apl is the area under the load (P) versus CODrecord, a is the instantaneous crack size, and b is theligament length equal toW� a. E is the Young’s modulus

and m the Poisson’s ratio. For the ZT1 BMG, E and m are80 GPa and 0.367, respectively.9

III. SELECTING A SUITABLE ALLOYCOMPOSITION

Our working hypothesis is that the BMG toughnessscales with the population and proliferation of multipleshear bands formed in front of the crack tip. Extensiveplastic events (shear transformations) lead to such profuseshear banding, developing an extensive (e.g., millimeter-sized) process zone and therefore causing plastic shield-ing and possibly deflection of crack path. To encouragespread-out plastic events, the prerequisite is that the BMGinternal structure should contain a large fraction of fertilesites for shear transformations. An illustration of differentlocal structures leading to different propensity for sheartransformations was presented before.12 In a nutshell, it ispreferable for the local structures to be “flexible” andprone to change. Such local structures should containundesirable coordination polyhedra that deviate from thosepreferred at the particular alloy composition. Specifically,geometrically unfavored motifs (GUMs) are more amena-ble to configurational changes than the relatively rigidclusters that possess the favored short-range order. The“GUMs” include those with unusually high or lowcoordination number for the alloy composition andatomic size ratio, and the clusters that have the favoredcoordination number but contain excess disclinations/distortion (e.g., non-Kasper polyhedra). Such relatively

aiW

¼ 0:999748� 3:9504uþ 2:9821u2 � 3:21408u3 þ 51:51564u4 � 113:031u5� �

; ð1Þ


J. Mater. Res., Vol. 29, No. 14, Jul 28, 2014 1491



uncomfortable clusters would be more likely to recon-figure under external stimuli (in this case, stresses).

There are two ways to increase the population of suchfertile sites for shear transformations. One is by usingrapid quenching of the melt, such that the resultingmetallic glass traps inside a higher fraction of these“liquid-like” GUM varieties in the local structures. Thetrend can be illustrated using a molecular dynamicssimulation. As shown in Fig. 3, the “participation ratio”(i.e., the fraction of atoms that undergo large strains)increases, or the degree of strain localization decreases,upon loading samples prepared at increasingly higherquench rates (with more frozen-in GUMs).13 Experimen-tally, one way to achieve more effective cooling is bycasting in a mixture of argon and helium. For a giventhickness, a BMG plate cooled faster using the lattermedium during casting exhibits a higher toughness, asshown in Fig. 4.14

Of course, a more interesting approach is to use routinecasting but change alloy composition, and accordinglythe internal structure, to induce an increased tendency forspread-out (plastic) shear transformations. Experiencesindicate that metallic glasses (MGs) obtained from morefragile supercooled liquids tend to be more plastic. It isalso known that different alloys compositions lead todifferent fragility of the liquids. To illustrate this point forthe range of compositions that form BMGs in the Cu–Zrbased system, Movie I and Movie II in SupplementaryMaterials compare a Cu64Zr36 metallic glass with anotherone at a Zr-rich composition, Cu20Zr80 (Cu40Zr60 behaves

in a very similar way). Clearly, upon loading these twootherwise identical molecular dynamics (MD) models(same dimensions and cooling history), the latter developsmuch more, and wider bands of, shear transformations. Incomparison, the former exhibits more localized strains ina narrower shear band that leads to an obvious shear offset.This comparison is also made in Figs. 5(a) versus 5(b),using snapshots at the end of the two movies. A similardifference was reported earlier at slightly different com-positions.13 Such a difference is also observed when

FIG. 3. Comparison of the propensity for strain localization upon plastic deformation of Cu64Zr36 metallic glass (MG) models containing 10,000atoms quenched usingMD simulations at different cooling rates: (a) 1� 1010 K/s, (b) 1� 1011 K/s, and (c) 1� 1012 K/s. Only atoms with vonMissesstrain larger than 0.3 are shown (in blue), at an overall sample strain of 12%.13

FIG. 4. Relationship between the fracture energy release rate and theglass-forming cooling rate used during fabrication of Cu49Hf42Al9BMG.14





FIG. 5. Different tendency toward shear localization in MG models of different compositions (cooled at the same molecular dynamics simulationquench rate of 1012 K/s) under tensile loading. In (a) and (c) for the Cu64Zr36 MG, the fraction of full icosahedra in Cu-centered polyhedra is about19%, whereas that in (b) and (d) for Cu20Zr80 MG is 2%. Only atoms with local atomic strain larger than 0.35 are shown with blue dots.





a notch was premade in the MD model sample, asdisplayed in Figs. 5(c) versus 5(d), for Cu64Zr36 versusCu20Zr80.

15 The structural origin is that during the BMGpreparation of the former composition, its internal struc-ture develops a dominantly high fraction of full icosahe-dra, which are the geometrically favored coordinationpolyhedra that tend to resist relaxation and deforma-tion.12,16 In comparison, for the latter composition, thefavored local motifs have more varieties and these poly-hedra have a higher tendency to undergo relaxations.

Figure 5 and the movies above give the impression thatfor Cu–Zr based BMGs, the right direction to encouragemore spread-out shear transformations is to move the alloycomposition toward the Zr-rich side. This is also true forCu–Zr–Al based ternary compositions. Adding Alimproves the GFA so that larger BMG samples can bemade, so some Al is necessary. However, adding Al

degrades toughness, as illustrated with experimentalresults in Fig. 6.9 This is because Al tends to increasethe fraction of rigid full icosahedra and introduces somecovalent character to the interatomic bonding (somedegree of polarized electron charge density distributionaround Al).9,15,16 Therefore, our strategy is to use a limitedAl percentage for an acceptable level of GFA, and thenreduce Cu content while increasing Zr content as much aspossible. This, according to our simulations, did reduce thepopulation of atoms involved in full icosahedra andincrease the fraction of atoms in GUMs.14 As a result,atoms that have participated in shear transformationsincrease, see the comparison in Fig. 7.15

While our modeling suggests that the more Zr thebetter, in experiments it is difficult to incorporate veryhigh Zr concentrations because there the metallic glassesdo not have sufficient GFA (incidentally, one could alsoattempt to move to very Cu-rich compositions, where thefraction of full icosahedra would eventually turn around todrop. But there the GFA is also poor). As experimentallydetermined in Fig. 8, only a small group of compositionscan be cast into fully amorphous BMGs sufficiently large(10 mm rods, or 3 mm plates) for our toughness experi-ments following ASTM standards. We have thereforesettled at Zr61Ti2Cu25Al12 (ZT1), where the criticaldiameter for copper-mold casting is Dc 5 10 mm. ThisBMG has a yield strength of ry 5 1600 MPa, a plasticstrain of 4% in compression and zero plastic elongation intension, a Poisson’s ratio of m 5 0.367, and a glasstransition temperature Tg of 652 K.17

Figure 9(a) shows the tensile specimens of the ZT1BMG, before and after failure. This BMG has appre-ciable plasticity in compression test, but nearly zeroductility in uniaxial tension. For the fractured specimenafter tension test, the inclined fracture angle (h) of thefracture plane, with respect to the loading axis, is about53°, as shown in Fig. 9(b). A dominant single shear band

FIG. 7. Atoms that have participated in shear transformation events, after 5% shear strain: (a) Zr45Cu45Al10, versus (b) Zr64Cu26Al10. Color scheme:Blue for Zr atoms, red for Cu atoms, and green for Al atoms.15

FIG. 6. Maximum strain energy release rate Jmax measured using notchedsamples versus Al concentration, in Zr–TM(TM 5 Co,Ni,Cu)–Al basedBMGs.9





leads to quick failure after the onset of yielding. On thefracture surface, a region with smooth features, asshown in Fig. 9(c), suggests the presence of a shear-off

region, which was found before only in the high-toughness Pd79Ag3.5P6Si9.5Ge2.

7 As shown in Fig. 9(d),viewed along the direction of Arrow II in Fig. 9(b), thecentral portion of the fracture surface is mainly charac-terized by numerous dimple-like or flower-like patternsassociated with the resistance against fracture.

IV. R-CURVE OF THE ZT1 BMG

Figure 10 shows a representative load versus CODcurve for the ZT1 BMG. The curve does not follow linear-elastic fracture behavior but instead shows pronouncedbending indicative of crack-tip plastic flow and stablecrack growth. Beyond the maximum load (Pmax), theapplied load gradually decreases, and the sample waseventually unloaded without fracture at COD 5 1.3 mm.J-integral characterization based on nonlinear elasticfracture mechanics, employing the equations earlier, givesa J values of 564 kJ/m2. In terms of the J-K equivalence,KJ ¼ JE

1�m2� �0:5

, the converted fracture toughness KJ is ashigh as 232 MPaOm (higher than crack growth toughnessreported for the Caltech Pd-BMG7).

Figure 11(a) plots the entire R-curve, for KJ versuscrack extension (Da). Here the Da was estimated fromthe increasing compliance during a sequence of multiple

FIG. 8. Composition map for BMG formation, for alloys with highconcentrations of Zr in pseudo-ternary (Zr0.97Ti0.03)–Cu–Al. Differentsymbols indicate the different critical diameters (Dc) of complete glassformation for as-cast rods fabricated using copper mold casting.9

FIG. 9. (a) Zr61Ti2Cu25Al12 BMG samples before and after tensile testing. SEM images of (b) side view of fractured sample, (c) and (d) high-magnification view of areas I and II in (b), respectively.





unloading/reloading for the single specimen technique(the pink data points). Beyond Pmax, the complianceschanges noticeably, indicating real crack advances. BeforePmax, however, there is onlyminor change in compliance sothe bending regime in the load-COD curve arises pre-dominantly from the formation of a crack-blunting region(this region resulting from numerous shear offsets is akin tothe “stretch zone” before stable crack propagation takesplace for ductile crystalline metals) ahead of the crack tip,due to extensive shear band interactions. The total nominalcrack extension from the compliance method is thereforeactually composed of the stretch (shear-off) zone and thetrue crack propagation. Consistent with the single-specimendata are the KJ values obtained from the multiple-specimentechnique where crack extension was visually determinedusing heat-tinting and refatigue. Data from these fivesamples have been included in Fig. 11(a) using coloredsymbols.

Figure 12 displays an SEMmicrograph of the fracturesurface for an unloaded ZT1 sample [yellow point in

Fig. 11(a)]. Four regions with distinct features areobserved from bottom to top ― the fatigue precrackedregion, a shear-off zone, a region with dimple-like mor-phology extending to ;0.5 mm in length, and finally therefatigued region with finer striations (due to lower load)than prefatigued region. The morphological feature of ZoneIII is consistent with that for mixed-mode crack growth (seediscussion later), corresponding to actual crack advance.

In Fig. 13, it is seen that the crack tip is surrounded bya profuse net of shear bands running extensively andcrossing each other, indicative of complex shear bandinteractions. The crack growth was gradual and undercontrol, as it advances to as far as ;0.45 mm in thedirection of the original fatigue crack (the directionperpendicular to the loading direction) without globalsample failure. One origin of the resistance to crackgrowth is the multiple shear bands, forming (olive/fan-shaped) slip lines and an extensive plastic zone. Theplastic zone size of ;1.7 mm can be compared with theestimated zone radius rp 5 1/2p�(KJ/ry)

2 5 3 mm inplane-stress state and rp5 1/6p�(KJ/ry)

25 1mm in plane-strain state, where KJ and ry is taken to be 232 MPaOmand 1600 MPa,17 respectively. The curved crack pathfollowing the shear band trace, instead of cutting straightthrough, is an expected source for increased energyconsumption. In Fig. 13, we see that the crack is obviouslydeflected to take a detoured path, following the path ofleast resistance (shear bands) and causing the crack tip toadvance under mixed mode (Mode I–Mode II) conditions.

V. TOUGHENING MECHANISMS

Shear-band interactions at the crack tip are generallybelieved to be a major factor responsible for toughening.Our current findings emphasize that this leads to crackdeflection/detour, contributing to the crack-growth resistance(the R-curve behavior). Although the far-field loading wasMode-I in our tests, the appearance of a blunting (shear-off)zone and the meandering crack path suggest that the crack tiploading has turned into a mixed mode. To confirm this, we

FIG. 11. Crack-resistance curves (R-curves) in which the fracture toughness KJ (converted from J values) is plotted against the crack extension,comparing (a) Zr61Ti2Cu25Al12 (Ref. 18) with (b) Pd79Ag3.5P6Si9.5Ge2 BMGs.7

FIG. 10. Load versus COD curve from three-point bending test of theZr61Ti2Cu25Al12 BMG. The maximum applied load (Pmax) is marked byan arrow. The a0 represents the initial crack length.18





calculated the contours of the deformation zones at the cracktip, assuming that it would follow the contour of themaximum shear stress, T. Specifically, at the specimensurface with the plane stress, the T drives the shear bandformation and the cracking simply follows the shear-bandingpath. Under this maximum shear stress (driving force)criterion, the crack would thus be deflected to follow thecontour of T. The crack deflection has two consequences, oneis the reduction in the crack-tip driving force and the other isthat the local loadingmode at crack tip becomes amixed one,involving both Mode-II and Mode-I.18

Under a loading mixed with the opening and shearmode, the T at the crack tip is

T ¼ KIffiffiffiffiffiffiffiffi2pr

p RðhÞ ; ð7Þ

where R(h) is an angular function related to the ratio of theMode-II versus Mode-I stress intensity factors. If the shearband develops when the shear stress reaches a criticalvalue, Tc, the contour of the shear band deformation zone(or plastic zone) is given by

rc ¼ K2I RðhÞ2pTc

; ð8Þ

which under a substantially Mode-II loading conditiongives a contour that compares very well with theexperimentally observed plastic zone shape in Fig. 13.18

As such, the crack-tip loading has clearly transformedfrom the pure Mode-I loading at the initial stage of thecrack growth to substantially Mode-II after the multipleshear bands fully develop at the crack tip. Evidently, sucha transformation in the crack-tip loading results in muchof the energy dissipation in resisting the crack advance.In fact, it has been demonstrated in BMGs before that thefracture resistance under Mode-II loading is significantlyhigher than that under Mode-I.19–22 This results inreduced driving force at the tip, and a deflected crack thatdetours with extended crack path length. A significantportion of the resistance to the Mode-II crack growth maycome from the resistance to crack sliding. One source ofsuch resistance stems from the remaining ligamentsbetween individual shear bands as they line up with themain crack, and the other is from the surface asperitiesseen on the crack surfaces. In other words, the hightoughness of ZT1 is due to its high propensity for crackdeflection following the shear band patch and spontaneouschange in loading mode. A prerequisite is of course thesignificant plastic deformation ahead of the crack tip in theform of profuse shear-banding operation and interactionsof numerous shear bands.

VI. THE TOUGHEST BMGs THUS FAR

We now compare ZT1 with the Caltech Pd-basedBMG, see Fig. 11. For a more meaningful comparison,we choose a critical point in the ZT1 R-curve in Fig. 11(a).Our decision was to use the KJ value of ;150 MPaOm asthe critical value, KC. This is the point where the dashedline crosses, corresponding to Pmax. Beyond this point, theblunting/stretch zone ends and the crack truly advances.We also confirmed that the KC is not sensitively dependenton sample size, by using 14 specimens with thicknessvarying from 2 mm to 4 mm, as shown in Fig. 14. Forthe Pd-BMG in Fig. 11(b), a critical KJ value is also

FIG. 12. SEM micrographs of fracture surface of a Zr61Ti2Cu25Al12BMG specimen subjected to unloading well over Pmax without failure,after subsequent refatigue failure, showing four zones with differentfeatures, fatigue precracked region (I), shear-off zone (II), crack growthregion (III), and refatigued region (IV). Boundaries between adjoiningregions are marked with dashed lines.18

FIG. 13. SEM image of side-view of the Zr61Ti2Cu25Al12 BMGspecimen, showing profuse shear banding and the typical olive/fan-shaped plastic deformation zone at the crack tip in Zr61Ti2Cu25Al12BMG.18





;150 MPaOm, if we follow ASTM E813-87 whichdefines JIC at 0.2 mm crack extension (this is when thecrack advance distance remains short with respect to theligament length). Moreover, both ZT1 and the Pd-BMGundergo extensive subcritical crack growth over hundreds ofmicrometers and the resistance to instability eventually led toa crack growth toughness KJ in excess of 200 MPaOm, asindicated by the extensive R-curve. We conclude thatZT1 and the Pd-BMG are quite comparable, in terms ofthe main properties manifested by the R-curve.

Using a KC of;150MPaOm, ZT1 and the Pd-BMG arecompared with previous BMGs, BMG-based compositesand all known engineering alloys, in the Ashby diagramshown in Fig. 15. The two BMGs are the toughest amongall monolithic BMGs so far, and the toughness is in facthigher than the best known alloys of similar yield strength.The combination of high strength and fracture toughnesspushes the two BMGs toward the upper right corner,indicating their superior damage tolerance.

VII. CONCLUDING REMARKS

In terms of materials discovery, our new Zr61Ti2Cu25Al12(ZT1) BMG stands out as an outstanding metallic alloy forengineering applications. It has a remarkable damagetolerance that not only rivals the previously toughest(Pd–Ag–P–Si–Ge) BMG, but also surpasses that ofconventional crystalline alloys. The advantages of ourBMG also include its high yield strength of ;1600 MPa,the absence of costly elements such as noble metals, anda robust GFA for fabrication (Cu-mold casting) of fullyamorphous rods on the centimeter-scale.

As for toughening of BMGs, we have identified themechanism that effectively enhances the resistance tofracture, i.e., a shear-induced spontaneous transformationof the crack tip loading and drastic crack deflection (detourfollowing the shear banding contour of a mm-sized plasticzone). Our finding is significant in opening up newpossibilities to reach record-high toughness, beyond whatis so far considered necessary (a high Poisson’s ratio around0.4, which severely limits the choice of constituent elementsin the BMG).

Finally, the success reported above, i.e., our search andthe eventual landing of ZT1, was guided by a correlationbetween the local atomic packing structure inside BMGsand deformation properties provided by computer simu-lations. In general, BMGs obtained from liquids of highfragility tend to be tougher. From that perspective, it isdesirable to move the alloy composition close to a pureelement (the very Pd-rich Caltech BMG may be anexample). But in the (commonly studied) Zr–Cu–Albased BMGs, both the very Zr-rich and very Cu-rich endsdo not offer sufficiently high GFA to produce BMGs. Overthe composition range amenable to BMG fabrication toreach 10mm dimensions, molecular dynamics simulationsprovided the insight that the Zr-richer side is more prone toshear transformations (and exhibits higher fragility23)when compared with the Cu-rich side.

Before closing, we note that the discovery of ZT1(and the Pd-BMG as well) also poses challengingquestions. From a conventional engineering point ofview, the combination of very high toughness and yetnearly zero tensile ductility is quite unusual. Metallic glassesare rather brittle in tension, but not necessarily so in bendingand other confined conditions. Future work is necessary to

FIG. 15. Ashby diagram in the form of yield strength versus fracturetoughness (KC or KJC) of engineering materials. For comparison, ournew Zr61Ti2Cu25Al12 (ZT1) BMG and the Caltech Pd79Ag3.5P6Si9.5Ge2BMG are included (labeled using stars). Also marked is the estimatedplastic zone radius.

FIG. 14. Plot of the KJmax converted from Jmax against tested specimenthickness (varying form 2 to 4 mm) for the Zr61Ti2Cu25Al12 BMG. Herethe subscript refers to the point of maximum load in the load versusCOD curve.





fully understand these behaviors and use the consequencesto our advantage. Also, in terms of glass structures, alloysystems not based on Cu–Zr would have atomic size ratiosand chemistry that prefer other types of coordination poly-hedra (e.g., full icosahedra may not be among populousclusters)24,25; the effects of composition on favored andunfavored atomic packing motifs and their response todeformation would need to be characterized in a systematicmanner for locating other plastic and tougher BMGs.

ACKNOWLEDGMENTS

The authors are indebted to their former Ph.D. studentsQiang He, Yongqiang Cheng, Li Zhang, Zhen-dong Zhufor their contributions to this study, and to Prof. J.KShang for assistance with the fracture mechanics analysis.This work was supported by the National Basic ResearchProgram of China (973 Program) under contractNo. 2007CB613906 and National Natural Science Foun-dation of China under Grant No. 51171180. EM wassupported at JHU by the US NSF-DMR-0904188.

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Supplementary Material

Movies

Different degree of strain localization in simulated Cu64Zr36 and (b) Cu20Zr80 MGs. The blue dots correspond tononaffine deformation (mostly plastic strains resulting from shear transformations) induced by tensile loading. Themovies record the subsequent deformation processes. While Cu64Zr36 shows an obvious shear offset associated witha dominant shear band, Cu20Zr80 exhibits spread-out shear transformations (Cu40Zr60 behaved in a similar way). Thisdifferent tendency toward strain localization, due to the difference in the structural order between these two MGs, isalso demonstrated in Fig. 5 using snapshots from the two movies.The movies can be found in the supplementary material at http://dx.doi.org/jmr.2014.160.




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INVITED FEATURE REVIEW Damage-tolerant Zr Cu Al … · · 2014-08-18INVITED FEATURE REVIEW Damage-tolerant Zr–Cu–Al-based bulk metallic glasses with record-breaking fracture