Micromechanisms of serrated flow in a Ni50Pd30P20 bulk metallic glass with a large compression plasticity

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Acta Materialia 56 (2008) 2834–2842 www.elsevier.com/locate/actamat

Micromechanisms of serrated flow in a Ni50Pd30P20 bulk metallic glass with a large compression plasticity K. Wang a, T. Fujita a, Y.Q. Zeng b, N. Nishiyama b, A. Inoue b, M.W. Chen a,* a

International Frontier Center for Advanced Materials, Tohoku University, Sendai 980-8577, Japan b Institute for Materials Research, Tohoku University, Sendai 980-8577, Japan

Received 1 September 2007; received in revised form 4 February 2008; accepted 8 February 2008 Available online 11 March 2008

Abstract Serrated flow is a characteristic feature of plastic deformation of bulk metallic glasses (BMGs) with a large compression strain. However, the underlying mechanisms of the discrete plasticity in the disordered solids have been debated for many years. Here, we report mechanical behavior and microstructural evolution of a Ni50Pd30P20 BMG subjected to uniaxial compression testing. Extensive nanocrystallization within shear bands and in the vicinity of fracture surfaces was observed and various crystal defects, including dislocations, twins and kink bands, were detected in the resultant nanocrystals. These observations suggest a microscopic mechanism of the serrated flow of the BMG, i.e. the stress drop is caused by local strain-softening and the arrest of shear bands is associated with in situ nanocrystallization. Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Bulk metallic glasses; Mechanical properties; Shear bands; Microstructure

1. Introduction Mechanical properties of metallic glasses strongly depend on the nucleation and propagation of shear bands in which shear flow takes place within nano-scaled zones [1–4]. This localized shearing caused by significant strain-softening generally leads to sample failure along a single shear band with limited global plasticity. Unlike conventional metallic glasses, large compression plasticity was observed more recently in a number of monolithic bulk metallic glasses (BMGs) [5–7]. These BMGs have two characteristic deformation features that are serrated flow, i.e. elastically loading and plastically unloading with sharp stress drop, in stress– strain curves and multiple shear bands on deformed sample surfaces. The formation of multiple shear bands during the serrated flow provides a visible plastic strain, but uniquely this localized shearing can automatically cease to avoid the further strain-softening that generally leads the runaway *

Corresponding author. Tel.: +81 22 215 2143; fax: +81 22 215 2196. E-mail address: [email protected] (M.W. Chen).

failure along a single shear band in conventional BMGs. The factors that cause both shear-softening and the arrest of the shear bands have been debated for many years since the serrated flow was first observed in Pd-based BMGs with a detectable plastic strain [8,9]. The strain-softening has been a topic of intense discussion and is believed to result from the decrease of material viscosity within shear bands by either increased free volume or temperature rise [2,8,10–12]. Nevertheless, the arrest of the shear bands has not been well understood. Over the years, several intrinsic and extrinsic factors have been suggested to cause the arrest of the shear bands, such as inhomogeneous microstructure and possible kinetic and instrumentation effects [13]. More recently, deformation-induced nanocrystallization has been observed in a number of BMGs with large plastic deformation [14–19]. It has been suggested that in situ nanocrystallization within shear bands during shear flow can offer a ‘‘self-locking effect” to prevent the further propagation of shear bands [14]. However, the crystallization reported so far was mainly observed in the BMGs deformed with either complex stress states, such as indentation and cold-rolling, or uncontrolla-

1359-6454/$34.00 Ó 2008 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2008.02.015

K. Wang et al. / Acta Materialia 56 (2008) 2834–2842

ble loading conditions, such as manual bending. Therefore, the definite evidence on the deformation-induced nanocrystallization along individual shear bands with known load– displacement history has not been achieved before and the influence of the nanocrystallization on the mechanical performances of BMGs remains to be comprehensively clarified. 2. Material and experimental methods A Ni50Pd30P20 alloy used in this study was prepared by arc-melting a mixture of pure Ni, Pd metals and prealloyed Pd–P ingots in argon atmosphere. The bulk glass samples in a cylindrical form with the diameter of 2 mm were produced by a copper mold casting method [20]. Specimens with a diameter of 2 mm and a length of 4 mm were tested under a quasi-static compressive condition at a nominal strain rate of 104 s1. All the tested samples were carefully prepared to have well finished surfaces and parallel ends. The plastic strains reported in this paper were measured by precise strain gauges and confirmed by measuring the height changes of the deformed samples without detectable bending caused by misalignment during compression tests. The microstructure before and after the compression tests was investigated by a transmission electron microscope (TEM, JEM-3000F) operated at 300 kV with a field emission gun (FEG). Thin TEM foils were prepared by electrochemical polishing with a solution of 20% HClO4 and 80% C2H5OH at 30 °C. Subsequently, a low-angle ion milling and polishing system (Fischione Inst. 1010) was employed to remove the chemical contamination layers on the TEM sample surfaces by short-time gentle milling (5–10 min) at the liquid-nitrogen temperature. For the deformed samples, thin slices were

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cut along a plane parallel to the loading direction and perpendicular to the visible shear planes identified from the traces of slip bands on the sample surfaces. The fractographs of the failed samples were investigated by a FEG scanning electron microscope (SEM, JSM-7000F). The microstructure in the vicinity of the fracture surfaces was characterized by TEM and the samples were prepared by a one-side thinning technique that allows the fracture surface to be retained. 3. Experimental results 3.1. Microstructural characterization of as-cast Ni50Pd30P20 BMG The as-cast specimens were first inspected by an X-ray diffractometer (XRD). Visible crystalline phase cannot be found from the XRD spectrum and only broad amorphous peaks can be seen. The microstructure of the as-cast Ni50Pd30P20 BMG was further characterized by TEM. The dark-field TEM micrograph (Fig. 1a) imaged by the first halo ring of a selected area electron diffraction (SAED) pattern shows uniform contrast of the as-cast alloy and visible nanocrystals cannot be observed. The halo SAED pattern inserted in Fig. 1a confirms that the as-cast alloy is fully amorphous. High-resolution electron microscope (HREM) observations further reveal the amorphous nature of the as-cast alloy (Fig. 1b) and nano-sized crystallites with a periodic lattice contrast cannot be found. 3.2. Compression tests of as-cast Ni50Pd30P20 BMG A typical compression stress–strain curve of the Nibased BMG is shown in Fig. 2a. The yield stress is mea-

Fig. 1. Microstructure of the as-prepared Ni50Pd30P20 BMG. (a) Dark-field TEM micrograph imaged by the first halo ring of the inserted SAED pattern taken by using a 200 nm selected area aperture and (b) the representative HREM image of the metal-metalloid BMG.

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3.3. Characterization of shear bands in deformed Ni50Pd30P20 BMG

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The deformed sample is shown in Fig. 2b. Visible sample bending caused by misalignment between the loading axis and the perpendicularity of sample ends cannot be seen. Corresponding to the serrated flow, multiple shear bands can be observed on pre-polished sample surfaces as shown in Fig. 2b. Both the primary shear bands and the final fracture plane are inclined nearly 45° with respect to the compression axis, suggesting that the shear stresses play a dominant role in the deformation and the failure of the BMG. The relatively uniform distribution of parallel shear bands on the sample surface demonstrates that the large plastic deformation is macroscopically homogeneous and performed under a uniaxial loading condition. Evidently, the multiple shear bands on the sample surface result from the serrated flow and each serration corresponds to an event that a shear band starts at a critical stress and automatically stops with a certain amount of shear strain. This process may be repeated in a single shear band or by the formation of a new shear band. The accumulation of these quasi-periodic birth and arrest of local shearing produces large global plasticity. Apparently, the arrest mechanism that effectively prevents the runaway slip along a single shear band plays an overriding role in the plasticity of the monolithic BMG.

500µm

Fig. 2. (a) A representative compression stress–strain curve of the Ni50Pd30P20 BMG. (b) An optical micrograph of a compressed Ni50Pd30P20 sample with pre-polished surface showing the multiple shear bands and the fracture plane. The insert shows the whole morphology of the deformed sample and visible sample bending cannot be seen.

sured to be 1780 MPa and the maximum plastic strain is 11%. Obvious strain-hardening can be observed in the stress–strain curve and the peak strength of 1870 MPa is obtained at the 9% plastic deformation. After that, a significant strain-softening occurs with the increased amplitude of serrations until failure. The overall plastic deformation is accomplished by serrated flow that has been highlighted by the insert in Fig. 2a. The amplitude of the stress drops gradually increases with strain and the slopes of both loading and unloading of the serrations keep almost constant from the yield point to the final failure. The loading slope is approximately parallel to the elastic part of the stress–strain curve, indicating that the loading process is elastic. The slope of the unloading sections is associated with the inelastic displacements caused by local shear deformation and opposite elastic expansion with rapid stress drops. Thus, each serration includes two processes, i.e. elastic loading to a critical value and inelastic destabilization by local shearing.

To understand the micromechanisms of the serrated flow, in particular the arrest of shear bands, the plastically deformed samples were characterized by TEM. The compressed samples with pre-polished surfaces and the traces of shear bands were carefully cut along a plane perpendicular to the primary shear bands. TEM foils were prepared by mainly thinning the sheets from the original internal sides to retain the traces of the slip on the specimen surfaces for quickly locating the shear bands during TEM observations. The dark-field TEM images (Fig. 3a and d) show a nano-sized band along the shearing direction (parallel to the slip bands shown in pre-polished surface (Fig. 2b). This band contains a large number of nanocrystals that show a bright contrast in the dark-field image. The nanocrystals with a size of 5 nm are only visible within the shear band and cannot be found in the regions out of the band, suggesting that their formation is associated with the localized shear deformation. The nanocrystallization within shear bands can be further confirmed by nano-beam electron diffraction. Fig. 3b is taken from a region within the shear band by using a 15  15 nm square selected area aperture and Fig. 3c is taken from a region out of the shear band (the marked areas in Fig. 3a). The sharp diffraction spots in Fig. 3b demonstrate the formation of nanocrystallites in the shear band and the diffusing ring pattern in Fig. 3c suggests the fully amorphous structure outside the shear band. The density of the nanocrystals in the shear band, measured to be 3  1023/m3 (Fig. 3d), is about two orders of magnitude higher than that in nanocrystallized BMGs by annealing [21].

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Fig. 3. TEM micrograph and SAED patterns of a shear band in the uniaxially deformed Ni50Pd30P20 BMG. (a) Dark-field TEM image of a 100 nm wide shear band with a high density of nanocrystals. (b) and (c) The nanodiffraction patterns taken from the regions within the shear band and out of the shear band, respectively (marked areas in (a)). The nanodiffraction experiments were performed by using a 15  15 nm square selected area aperture. (d) An enlarged dark-field TEM micrograph of the shear band shown in (a).

In addition to the nanocrystallization, shear bands illuminated by TEM with a nanoscale resolution can reveal more details on the structural evolution of the localized shearing, which cannot be detected by optical microscopy and SEM. A bright-field TEM image (Fig. 4) illustrates a deformation region in the vicinity of a fracture surface with a big surface relief that shows a dark contrast. In this region, a bundle of small shear bands with a width of 20 nm as marked by dashed lines can be observed, indicating that the shear deformation in this region is accomplished by multiple individual shearing events.

Fig. 4. A bright-field TEM image of a bunch of narrow shear bands nearby the fracture surface and the corresponding SAED pattern taken by using a 200 nm selected area aperture.

Importantly, the 20 nm wide bands are highly nanocrystallized and the contrast shown in the bright-field TEM image mainly results from the diffraction contrast generated by the high density of nanocrystals within each slip band. The SAED pattern of the shear bands is shown in

Fig. 5. (a) HREM image of the selected region in Fig. 2 showing a high density of nanocrystals mixed with a residual amorphous phase. (b) The enlarged view of the marked region in (a) and (c) the Fourier-filtered image of (b). Both (b) and (c) show a kink band along with dislocations in the deformation-induced nanocrystals.

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the insert of Fig. 4. The sharp diffraction rings confirm the formation of nanocrystals. The index of the ring patterns suggests that deformation-induced nanocrystals have a simple face-centered cubic (fcc) structure with a lattice constant of 0.375 nm, consistent with a palladium–nickel solid solution phase [22]. Moreover, the diffuse rings are still visible in the SAED pattern, indicating that the matrix of the shear bands is a residual amorphous phase. The fcc nanocrystalline phase mixed with the residual amorphous phase in the shear bands was further characterized by HREM (Fig. 5a). Interestingly, a high density of crystal defects that generally appear in severely deformed nanocrystals [23] can be readily found, suggesting that the plastic deformation of the in situ nanocrystallized phase takes place during the propagation and the arrest of the shear bands. In addition to dislocations and deformation twins, kink bands that are rarely seen in cubic crystals are found in some nanoparticles (Fig. 5b and c). The rotated lattices along the shearing direction can be clearly observed and the rotation angle does not follow any twin relationship in cubic crystals. The space of the distorted lattice fringes is 0.22 nm, well consistent with that of the {1 1 1} plane of the palladium–nickel solid solution phase. A number of dislocation pairs are present at the boundaries of the kink band and are characterized to be h0 1 1i edge dislocations. Kinking, a mechanism of plastic deformation, is quite rare in cubic crystals though long known in minerals and hexagonal close packed (hcp) crystals [24]. The formation of kink bands generally requires a high stress where the shearing direction is nearly normal to the slip planes

and conventional deformation modes, such as dislocation sliding and twinning, cannot be activated [25]. 3.4. Microstructural evolution in the vicinity of fracture surfaces The stress–strain curve shown in Fig. 2a reveals that the final failure of the BMG sample is accompanied by the increased amplitude of serrated flow and significant strain-softening. Combining with the cross-sectional observations of the multiple shear bands in the vicinity of the fracture surface (Fig. 4), the fracture of the BMG sample is most likely to be associated with the intensive shear deformation within a narrow region of about hundreds of nanometers in thickness. Fig. 6 shows SEM images of the fractograph of a Ni50Pd30P20 BMG sample failed by compression testing. The typical vein pattern along with re-solidified droplets can be clearly observed (Fig. 6a–c) [26]. It is interesting to note that the ridges of the vein pattern in some regions exhibit the features that are similar to those of crystal fractographs (Fig. 6d), implying possible crystallization of the fracture surface during the intensive shearing and the final fracture. Planar TEM observations of the fracture surface reveal two distinct regions: an amorphous domain surrounded by a partially nanocrystallized zone. Bright-field TEM image (Fig. 7a) shows that the nanocrystallized zones are of 2 lm in width, which is consistent with the width of the vein ridges in the SEM fractographs (Fig. 6d and the insert in Fig. 7a). The size and distribution of the nanocrystals are fairly consistent with

Fig. 6. SEM fractographs of a Ni50Pd30P20 BMG sample failed by compressive testing: (a) Typical vein pattern on the fracture surface; (b) enlarged view of the vein-like morphology showing the nanoscale feature; (c) the re-solidified droplet (as marked by arrow) on the fracture surface; and (d) the typical fracture morphology with tearing ridge-like feature mixed with the vein pattern.

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b A

nanocrystalline

B amorphous

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amorphous

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nanocrystalline

amorphous 100nm Fig. 7. TEM micrographs of the fracture surface of a compressed Ni50Pd30P20 BMG specimen. (a) Bright-field TEM image of the fracture surface showing a mixture structure: dark nanocrystallized region and bright amorphous region (the insert is an enlarged view of Fig. 6d). (b) and (c) the magnified micrograph of area A and B in (a), respectively. (d) HREM image of the nanocrystallized region (marked area in (c)). Various crystal defects can be detected in the nanocrystals. Deformation twins are marked in the micrograph and the squared region includes a high density of stacking faults.

those in cross-sectional TEM images (Fig. 4). The SAED pattern taken from the nanocrystallized zone demonstrates that the nanoparticles are the fcc Pd–Ni solid solution, which is the same as the crystal phase in the shear bands (Figs. 3 and 4). HREM characterization confirms the fcc structure of the nanocrystals and suggests that the nanoparticles are mixed with the residual amorphous phase (Fig. 7d). Again, a high density of crystal defects, such as the stacking faults and deformation twins (as marked in Fig. 7d), can be observed in these nanoparticles. Moreover, the obvious composition difference between the crystalline zone and the glassy region has not been detected by both SEM and TEM energy dispersive X-ray spectrometers. 4. Discussion 4.1. Mechanisms of in situ nanocrystallization within shear bands

deformation. However, the Ni50Pd30P20 BMG has a very good glass-forming ability and a wide supercooled liquid region [20]. Even at a temperature above the crystallization temperature, the incubation period of crystallization is still about tens of seconds. Thus, only the temperature rise is apparently not enough to cause the nanocrystallization within a very short-time scale (<0.02 s) determined by rapid stress drop in serrations (Fig. 2a and Ref. [30]). Other factors, such as high shear strains (5000%), ultra-high shear strain rates (104–105 s1) [27,31] and excess free volumes produced by deformation [2,3,28] along the shear bands, may also play important roles in the in situ nanocrystallization. Moreover, the unique atomic configuration of the metal-metalloid BMG in which some of the metalloid-centered quasi-equivalent clusters or MROs [32,33] are shared by cubic crystal structures, may promote the rapid crystallization during shear flow. 4.2. Microscopic explanations on the serrated flow

A number of suggestions for the mechanisms of deformation-induced nanocrystallization in metallic glasses have been proposed [16,27–29], but a convincing explanation with definite evidence is still missing. Because both crystallization at high temperatures and deformation-induced temperature rise within shear bands are well-proven phenomena in nonequilibrium metallic glasses, one may think that the extensive crystallization along shear bands may result from high temperatures caused by the energy dissipation of shear

Deformation-induced nanocrystallization in metallic glasses has been discussed for many years and received considerable attention in recent years [14–19,27–29]. The previous investigations have been focused on finding crystalline phases in deformed metallic glasses but have not directly tied the nanocrystallization with mechanical properties, in particular plasticity. The extensive nanocrystallization along shear bands can dramatically increase the strength

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of the shearing regions because the viscosity of semi-solid BMG slurries depends exponentially on the fraction of solid nanocrystals [17,34]. This strengthening caused by in situ nanocrystallization can compensate the strain-softening and thus prevents the runaway failure along the shear bands. Apparently, the nanocrystallization is the function of shear deformation along shear bands and the fraction of the nanocrystalline phase spontaneously increases with shear strains and deformation time. The serrations with quasi-periodic birth and arrest of local shearing are probably associated with a ‘‘critical shear strain” at which the density or volume fraction of the nanocrystalline phases reaches a critical value at which the increased viscosity by the nanocrystallization is high enough to prevent the further shearing along a band. The observed deformation defects in the nanoparticles further confirm this assumption. The coincidence of deformation direction of kink bands, twinning and dislocation sliding with the shearing direction rules out the possibility that the crystal defects were produced by TEM sample preparation and non-equilibrium crystal growth processes. During shear flow, the solid nanoparticles will interact with the applied shear stresses that may be high enough to cause the plastic deformation of the nanocrystals. On the other hand, the severe deformation of the nanocrystals gives rise to strain-hardening, which also assists the arrest of the local shearing. Consequently, the in situ nanocrystallization can stop the continuous propagation of a shear band and turn the deformation into elastic state until the onset of a new shear band with further loading. In addition to the microscopic explanation on the serrated flow of the BMG, instrumental effects arising from the lateral motion resistance and longitudinal constraint may play a certain role in the arrest of shear bands. When large shearing takes place along a single shear band, the opposite lateral movement at the two ends of a tested sample will be required to coordinate the overall compression

a

deformation. Such lateral movement may be prevented by the friction between the sample top/bottom and the platens of the mechanical test apparatus, resulting in the lateral force to restrict the further slip along the shear band. However, although the shear strain is very large (5000%) along each shear band, the lateral offset caused by the shear displacement is very small (0.7–3.5 lm). Thus, the lateral constraint effect should not be significant for the arrest of shear bands. In addition to the lateral constraint, one may think that the fast slip along a shear band with an ultra-high shear strain rate (104–105 s1) may not be followed up by the slow platen motion operated in a constant displacement or constant strain rate mode (in the present study, we used the constant displacement rate mode) with a quasi-static compression strain rate (104– 101 s1). However, the longitudinal strain rate corresponding to the fast shearing along a shear band is only 103–100 s1. In our study, the nominal compression strain rates were changed from 105 to 102 s1 and dramatic changes in the serrated flow behavior of the BMG, in particular the disappearance of the serrations at high strain rates, were not observed. Moreover, if the grip constraint for sample motion is the reason for the observed shear band arrest, this effect should be more significant in tensile test mode where both ends of the sample are tightly held by the grips. However, obvious serrated flow along with a detectable plastic strain has not been observed in BMGs subjected to tensile testing. Furthermore, with the exactly same measurement conditions we tested a number of BMGs, including Zr-based, Cu-based and Fe-based BMGs, and found that only the Ni50Pd30P20 BMG shows remarkable serrated flow with multiple shear bands. These observations suggest that the serrations in the stress–strain curves are mainly associated with the intrinsic material properties, although the instrumentation constraints in both longitudinal and transverse directions certainly benefit the arrest of shear bands and thereby plasticity.

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Strain (%) Fig. 8. (a) A representative compression stress–strain curve of the Ni50Pd30P20 BMG at the situation of misalignment. The insert shows the bending shape of the deformed sample. (b) A SEM image showing a high density of shear bands in the intensively deformed region marked in the insert. The complex stress state caused by misalignment results in a disorderly distribution of small shear bands.

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Another possible extrinsic effect that may affect the plasticity of BMGs is the sample alignment between the loading axis and the perpendicularity of the sample ends. We have noticed that slight sample misalignment can lead to an ultra-large plastic strain. However, in this case, the plastic deformation is not macroscopically uniform and mainly takes place in the regions near by the sample ends, which results in obvious sample bending (see the insert of Fig. 8a). Moreover, the serration in the corresponding stress–strain curve is not as conspicuous as that in wellaligned samples (Fig. 8a). The complex stress states caused by the misalignment obviously suppress the extensive propagation of individual shear bands and lead to the discrete displacement, insignificant by a large number of small shear events (Fig. 8b). Thus, this extrinsic effect can be readily found from the shape of deformed samples and the strain vs. stress behavior. In comparison with Fig. 2, again, we can conclude that the large plasticity of the Ni50Pd30P20 BMG reported in this paper is mainly due to an intrinsic material effect. Although shear bands can be halted by complex stress states, the arrest of shear bands in the well-aligned samples is mainly due to the intrinsic effect, i.e. deformation-induced nanocrystallization.

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strain-hardening can be well explained by an exhaustion model proposed in our earlier study [14]. According to the observations, the deformation-induced nanocrystals prevent the shear-softening, leading to the arrest of shear bands. To sustain a plastic deformation rate in a specimen during mechanical testing, new shear bands need to be formed when an active shear band is blocked due to the in situ nanocrystallization. Because the nucleation of shear bands prefers to initiate at the weakest sites, such as casting defects (porosity and inclusions) and surface flaws, where high concentrated stresses are easily generated [36,37]. The critical stresses to drive the formation of new shear bands will gradually increase from easy to difficult nucleation sites. During this deformation process, the applied force is required to progressively increase in order to keep continuous plastic deformation, which results in ‘‘workhardening” and the formation of multiple shear bands as being observed in Fig. 2. Apparently, this type of ‘‘workhardening” is different from the classical one in crystalline metals and alloys which results from the interaction among defects, such as dislocations, deformation twins and grain boundaries. 5. Conclusions

4.3. Strain-hardening of the BMG Because of the lack of crystal slip, the strain-hardening caused by the increase of dislocation density is not expected in glassy materials. However, a number of studies [5–7] have reported that the obvious strain-hardening takes place in BMGs with large compression strains. This unusual mechanical behavior has been explained by inhomogeneous microstructure and the possible interaction between the shear bands [6,35]. As shown in Fig. 2a, the noticeable strain-hardening also takes place in our Ni50Pd50P20 BMG. The strain-hardening exponent is measured to be about 0.06, which is close to the smaller end as compared with that of fcc metals (0.10–0.50). Traditionally, the discussion on strain-hardening and softening is based on true stress– strain curves in which the uniform change in the sample cross-sections is accounted. Because the plastic deformation of BMGs is inhomogeneous and accomplished by individual shear bands, true sample cross-section change cannot be achieved solely from stress–strain curves. However, the effective cross-section that withstands applied forces should not change too much with about 11% deformation and, in fact, becomes smaller with compression strain. Thus, the hardening showed in the engineering stress–strain curve (Fig. 2a) still reflects the intrinsic material behavior. Regarding the underlying mechanism of the observed strain-hardening, we have demonstrated that the as-cast alloy is homogeneous at sub-nano-scale (Fig. 1) and the crossover of primary shear bands is rarely seen from the deformed sample surfaces (for example, Fig. 2b). Thus, the inhomogeneous microstructure and the interaction of shear bands suggested by previous studies cannot be applied to the Ni50Pd30P20 BMG. Instead, the

1. A systematic study has been carried out to characterize the mechanical behavior of a Ni50Pd30P20 BMG by uniaxial compression testing. The material exhibits a large compression strain of 11%. Detectable strain-hardening takes place with the plastic strain up to 9% at the peak strength of 1870 MPa and further deformation leads to significant strain-softening and eventual sample failure. 2. TEM characterization reveals that a high density of nanocrystals precipitate exclusively within 20–100 nm wide shear bands produced by uniaxial compression. Electron diffraction analysis and HREM characterization demonstrate that the nanocrystalline phase is an fcc palladium–nickel solid solution with a lattice constant of 0.375 nm. 3. A high density of crystal defects were observed in the in situ formed nanocrystals, suggesting that they have experienced severely plastic deformation during the propagation and the arrest of shear bands. The deformability of the precipitated fcc phase appears to be important to the large plasticity of the BMG because it can release the applied stress by plastic deformation rather than cracking. 4. The crystallization in the shear bands within a very shorttime scale implies that the deformation-induced temperature rise may not be the only reason leading to the formation of extensive nanocrystals. The assistance of the high shear strain rate and the large shear strain along the shear bands may be important for the in situ nanocrystallization. 5. The obvious strain-hardening observed in the homogeneous BMG can be well explained by an exhaustion model in which the critical stress to drive the

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nucleation of shear bands will gradually increase from easy to difficult nucleation sites during deformation. The applied force is thus required to progressively increase with plastic strain to keep further plastic deformation. 6. This study provides compelling evidence on the micromechanism of serrated flow of the Ni50Pd30P20 BMG, i.e. the stress drop caused by strain-softening and the arrest of shear bands by in situ nanocrystallization.

Acknowledgements We thank Prof. A.L. Greer for helpful discussions. This work was sponsored by the ‘‘Grant-in-Aid for Scientific Research in Priority Areas: Materials Science of Bulk Metallic Glasses”, ‘‘Global COE for Materials Science”, and ‘‘World Premier International Research Center (WPI) initiative for Atoms, Molecules and Materials”, the Ministry of Education, Culture, Sports, and Science, Japan. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]

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