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Strain rate dependence of mechanical behavior in a CuZr-based bulk metallic glass composite containing B2-CuZr phase
Materials Science & Engineering A 606 (2014) 268–275
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Strain rate dependence of mechanical behavior in a CuZr-based bulk metallic glass composite containing B2-CuZr phase R. Wei a, S. Yang a, C.J. Zhang b, L. He a,n a b
State Key Laboratory for Mechanical Behavior of Materials, Xi’an Jiaotong University, Xi'an 710049, China School of Materials Science and Engineering, Chang'an University, Xi'an 710064, China
art ic l e i nf o
a b s t r a c t
Article history: Received 23 January 2014 Received in revised form 18 March 2014 Accepted 25 March 2014 Available online 1 April 2014
Strain rate effect on the mechanical behavior of a Cu48Zr48Al4 bulk metallic glass composite (BMGC) containing a B2-CuZr phase was investigated under quasi-static and dynamic compression. It was found that the CuZr-based BMGC exhibits remarkable strain-rate-softening, although the deformation-induced martensitic transformation of the B2-CuZr phase similarly occurred under both quasi-static and dynamic loading. The significant transformation-induced plasticity (TRIP) and pronounced work-hardening characteristic of the BMGC upon quasi-static loading were found to be considerably deteriorated under dynamic loading. Furthermore, the BMGC alloy was aboratively heat-treated into austenite and martensite states mainly containing the B2-CuZr phase and the B190 -CuZr phase respectively. By comparing the strain-rate responses of the mechanical behaviors in the austenite and martensite states, it was shown that the B2-CuZr phase itself presents dynamic softening due to the deformation-induced martensitic transformation with heat release effect. The dynamic softening characteristic of the B2-CuZr phase limits its contribution to the potential mechanical behavior of the CuZr-based BMGC containing the B2-CuZr phase under dynamic loading. & 2014 Elsevier B.V. All rights reserved.
Keywords: Mechanical characterization Bulk amorphous alloys Composites Dynamic loading Shear bands Martensitic transformation
1. Introduction Bulk metallic glasses (BMGs) have attracted great attention owing to their excellent strength and elasticity, which make them attractive as potential candidates for structural and functional applications [1–4]. However, BMGs can only accommodate plastic strain in a highly localized shear-banding manner due to the lack of dislocations in their long-range disordering atomic structure [5]. The shear-bandbased deformation mode makes BMGs suffering from limited plasticity and work-softening [6]. Extensive efforts on BMGs have been devoted to improving their mechanical behavior. Among many approaches, the fabrication of BMG-based composites (BMGCs) consisting of harder glassy matrixes and softer second crystalline phases has been proved to be effective in enhancing the plasticity and changing the work-softening of BMGs [7]. For e.g., Zr-based BMGCs reinforced with refractory metal phases were firstly fabricated and present better plasticity under compression [8–10]. Zr-based and Ti-based BMGCs reinforced by ductile β-phase dendrites can exhibit remarkable plasticity under compression and tension, and obvious work-hardening under compression [11–14]. The improved mechanical behaviors are fundamentally attributed to the introduction of
n
Corresponding author. Tel.: þ 86 29 82665165; fax: þ 86 29 826634533. E-mail address: [email protected] (L. He).
http://dx.doi.org/10.1016/j.msea.2014.03.092 0921-5093/& 2014 Elsevier B.V. All rights reserved.
dislocation-based deformation mode of the crystalline phases into the BMGCs [15,16]. Recently, a new strategy was proposed to form the BMGCs reinforced with a shape memory B2 phase [17]. New CuZr-based or TiCuNi-based BMGCs containing B2-CuZr phase or B2-Ti(Cu,Ni) phase can display significant plasticity and pronounced workhardening under compression [18–22], and especially exhibit alleviated work-softening behavior under tension testing [23–26]. It has been identified that deformation-induced martensitic transformation of the B2 phase, besides its dislocation-based deformation, is responsible for the unusual mechanical behavior of the BMGCs [16,20], which is similar to the transformation-induced plasticity (TRIP) in some traditional crystalline alloy systems [27]. This opens up a promising and innovative route for developing high-performance BMGs as engineering materials [7]. Another characteristic of BMGs concerns their strain rate dependence of mechanical behavior. Much current research has studied the dynamic uniaxial compressive mechanical behavior of BMGs by means of split-Hopkinson pressure bar (SHPB) apparatus. Although there is still some debate on the strain rate sensitivity (SRS) of BMGs, most experimental results indicate that BMGs performed apparent strain-rate-softening (i.e. negative SRS) as the strain rate beyond a critical range [28–34]. Differing from the strain-rate-hardening in traditional crystalline alloys with dislocation-based deformation mode, the strain-rate-softening of BMGs is also attributed to their
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shear-band-based deformation mode [32–34]. By modeling the deformation behavior of BMGCs [18,19,25], it is expected that the SRS of BMGCs can be affected by the competition of the shearband-based deformation within the glassy matrixes and the dislocation-based deformation within the crystalline phases. This has been demonstrated in some BMGCs. For e.g., the Zr-based BMGCs containing tungsten phase or β-phase dendrites can exhibit positive SRS [35–39]. A Ti-based BMGC containing βphase dendrites can present not only strain-rate-hardening but also distinguished work-hardening capacity upon dynamic loading [40]. For the CuZr-based or TiCuNi-based BMGCs containing a B2 phase and with TRIP effect, it can be reasonably inferred that the SRS may be additionally affected by the martensitic transformation of the B2 phase. However, so far, little attention has been devoted to this issue. In the present work, the strain rate dependence of mechanical behavior in a CuZr-based BMGC alloy containing a B2CuZr phase was investigated. In addition, the mechanical responses to strain rate for the austenitic B2-CuZr phase and the corresponding martensitic B190 -CuZr phase were compared, martensitic transformation effect on the SRS of the BMGC was analyzed.
2. Experimental methods A representative CuZr-based alloy with a nominal composition of Cu48Zr48Al4 was chosen for the current investigation because of its capability of obtaining the as-cast BMGC microstructure with a single B2-CuZr phase embedded in a glassy matrix at a suitable cooling rate [24]. The master alloy ingot was produced by arcmelting Zr, Cu and Al of 99.99% purity in a Ti-gettered argon atmosphere. As-cast cylindrical rods with a diameter of Ф3 mm were prepared via suction casting using a copper mold. In order to explore the martensitic transformation effect of the B2-CuZr phase on the SRS of the Cu48Zr48Al4 BMGC, some of the as-cast rods were aboratively heat-treated to have the microstructures mainly containing the austenitic B2-CuZr phase or the corresponding martensitic B190 -CuZr phase respectively. According to the binary Cu–Zr phase diagram (http://www1.asminterna tional.org/AsmEnterprise/APD), the B2-CuZr is a high-temperature metastable phase and should decompose into the low-temperature equilibrium phases, Cu10Zr7 and CuZr2, via a eutectoid reaction at 988 K if the system is given the time to equilibrate. For the present ternary Cu48Zr48Al4 alloy, the eutectoid reaction was found to occur within a temperature range from 977 K to 996 K [17]. Thus, some of the as-cast rods were annealed at 1073 K (above the eutectoid temperature range of the alloy) for 60 min to form the high-temperature B2-CuZr phase, and then water-quenched to room temperature. The water-quenching could limit the eutectoid
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reaction, the microstructure mainly containing the martensitic B190 -CuZr phase could be obtained because the martensitic transformation start temperature (Ms) from the B2-CuZr phase to the B190 -CuZr phase in the present Cu48Zr48Al4 alloy is higher than room temperature (about 337 K) [24,26]. To obtain the microstructure mainly containing the austenitic B2-CuZr phase at room temperature, the shape memory effect arising in CuZr-based alloys was exploited [19,41]. It had been testified that the Ms temperature of the B2-CuZr phase could be effectively lowered by repeated martensitic transformation of the phase induced by thermal cycling, and the high-temperature austenitic phase could be stabilized to room temperature [19,42]. Thus, some of the annealed rods were further thermally cycled between 673 K and 77 K, the temperatures are much higher than the austenitizing temperature of the B190 -CuZr phase and much lower than the martensitic transformation temperature of the B2-CuZr phase respectively [19,24]. The thermal cycling treatment was performed by heating the annealed rods at 673 K for 10 min and then quenched in liquid nitrogen for another 10 min. The rods were heat-treated that way at least six times. The microstructures of the as-cast and heat-treated rods were examined by X-ray diffraction (XRD) using an X'Pert-Pro diffractometer with Cu-Kα radiation and observed by scanning electron microscopy (SEM) using a JSM-7000F microscope. The crystallization characteristics of the rods were identified by differential scanning calorimetry (DSC) using a SETARAM-LabsysTM TG-DSC at a heating rate of 0.677 K/s with a sample mass of 3071 mg. Cylindrical specimens with the aspect ratios of 2 and 1 were cut from the as-cast and heat-treated rods. Quasi-static compression tests were conducted using the higher-aspect-ratio specimens on a screw-driven SUNS CMT5105 testing machine at a constant crosshead displacement rate of 0.05 mm/min, which corresponds to an initial strain rate of 1.39 10 4 s 1. Dynamic compression tests were performed using the lower-aspect-ratio specimens on a split Hopkinson pressure bar (SHPB) apparatus at a strain rate range of 2–3.3 103 s 1, the detailed process was described elsewhere [43]. In order to observe the deforming characteristic of glassy matrix and B2-CuZr phase in the CuZr-based BMGC, the side surfaces of some ascast specimens were previously polished, and examined by the SEM observation after both quasi-static and dynamic compression tests.
3. Results and discussion The XRD patterns of representative Ф3 mm Cu48Zr48Al4 alloy specimens in as-cast and heat-treated states are shown in Fig. 1a. For the as-cast specimen, crystalline diffraction peaks superimposed on a broad amorphous scattering maximum are displayed,
Fig. 1. (a) XRD patterns of representative ∅3 mm Cu48Zr48Al4 alloy specimens in as-cast, annealed and annealed plus thermally cycled states. (b) DSC curves of the specimens and the corresponding as-spun ribbon. (c) SEM image of the as-cast specimen.
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the Bragg peaks can be allocated to a cubic primitive B2-CuZr (Pm-3m) phase. For the annealed specimen, the XRD pattern consists of stronger diffraction peaks of a B190 -CuZr (P21/m and Cm) phase and relatively weaker diffraction peaks of retained B2CuZr phase. This is because the martensitic transformation finish temperature (Mf) of the B2-CuZr phase in the present Cu48Zr48Al4 alloy is lower than room temperature [24,26]. It is difficult to exactly determine the percentage of the retained B2-CuZr phase in the annealed specimen just based on the XRD pattern. For the annealed plus thermally cycled specimen, the XRD pattern was remarkably changed, consisting of stronger diffraction peaks of a B2-CuZr phase and relatively weaker diffraction peaks of retained B190 -CuZr phase. This is because the shape recovery of the present Cu48Zr48Al4 alloy is not complete [19], a small amount of the martensitic B190 -CuZr phase did not transform back to the austenitic B2-CuZr phase after the thermal cycling. The XRD patterns of the two heat-treated specimens are consistent with the results in Ref. [19]. Fig. 1b shows the DSC curves of the as-cast and heattreated specimens at a heating rate of 0.677 K/s. For the sake of comparison, the DSC curve for the corresponding as-spun glassy ribbon is also presented. It can be seen that the as-cast specimen and the ribbon began glass transition at a very close temperature of Tg 698 K, and exhibit an almost same crystallization temperature of TX 753 K. However, the crystallization peak area of the BMGC specimen distinctly reduces compared with that of the glassy ribbon. Based on the crystallization enthalpy reduction, the B2-CuZr phase volume fraction in the as-cast is estimated to be about 31%. Fig. 1c (inset in Fig. 1b) shows the SEM image of the ascast specimen in a backscattered electron (BSE) mode. It displays that some dark B2-CuZr precipitates with patch-like shape are embedded in a brighter, featureless matrix. The information in Fig. 1a and c indicates that the as-cast specimen is an in-situ BMGC with the B2-CuZr phase embedded in the glassy matrix. For the two heat-treated specimens, no apparent crystallization peak appeared on their DCS curves, which means that very limited amount of glassy phase exists in the specimens because they all underwent long time annealing at the temperature of 1073 K. Although desirable single B190 -CuZr phase in the annealed specimen and single B2-CuZr phase in the annealed plus thermally cycled specimen were not obtained, previous research testified that a small amount of the retained B2-CuZr phase or B190 -CuZr phase did not obviously affect the typical mechanical behaviors of the two heat-treated specimens under quasi-static compression [19,42]. Thus, it can be reasonably expected that mechanical behaviors of the two heat-treated specimens during dynamic compression would be similarly controlled by their predominant crystalline phases respectively. Based on the descriptions mentioned above, the as-cast, annealed and annealed plus thermally
cycled specimens are considered to be in BMGC, martensite and austenite states respectively in the following description. Fig. 2a shows the quasi-static and dynamic compressive engineering stress–strain curves of the specimens in the BMGC, martensite and austenite states. Upon quasi-static compression, it can be seen that the BMGC specimen displays a yield stress of about 1580 MPa (characterized by s0.2) and a fracture strain of about 18.5%. Furthermore, based on the corresponding true stress– strain curve (not shown here), the BMGC specimen also exhibits a pronounced work-hardening behavior and a resultant maximum stress of about 2200 MPa, which are in agreement with former reports [18,19]. For the martensitic and austenitic specimens under quasi-static compression, it is displayed that the yield strength of the austenite specimen is much lower than that of the martensitic specimen. However, the softer austenite specimen exhibits rather larger plasticity than that of the martensitic specimen, accompanying by strong work-hardening and a maximum stress very close to that of the martensitic specimen before finally fractured. The mechanical behaviors of the martensitic and austenitic specimens upon quasi-static compressive loading mentioned above are also consistent with former reports [19,42]. Upon dynamic compression, the representative engineering stress–strain curves in Fig. 2a illustrate that remarkable strainrate responses for the specimens in the BMGC, martensite and austenite states all occurred. Compared with the corresponding quasi-static stress–strain curves, the following phenomena can be observed. For the specimens in BMGC state, the significant plasticity and pronounced work-hardening under quasi-static compression are considerably deteriorated under dynamic compression. The maximum stress under the dynamic compression is even lower than the yield stress under quasi-static compression, indicating remarkable strain-rate-softening of the BMGC. For the specimens in martensite state, remarkably increased maximum stress upon dynamic loading is presented, indicating obvious strain-rate-hardening in the martensite state, like that in traditional crystalline alloys with dislocation-based deformation mode. However, for the specimen in austenite state, apparently decreased maximum stress upon dynamic loading is exhibited, revealing contrary strain-rate-softening in the austenite state. The strain rate dependences of maximum stresses for the BMGC, martensitic and austenitic specimens are displayed in Fig. 2b. At least nine specimens in each state were tested under both quasi-static and dynamic loading to verify the scattering of the compression data. The scattering of the quasi-static compression data at the strain rate of 1.39 10 4 s 1 for each state of specimens is displayed in error bars in Fig. 2b, the dynamic compression data at the strain rate range of 2–3.3 103 s 1 for all the specimens in three states are plotted in the figure. In despite of the
Fig. 2. (a) Quasi-static (denoted as QS) and dynamic (denoted as D) compressive engineering stress–strain curves of the specimens in different-states, the strain rate for the dynamic compression was about 2.8 103 s 1. (b) The dependences of maximum flow stresses on the strain rate for the specimens in different states.
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data scatter under dynamic loading, the general strain-ratesoftening tendency of the BMGC and austenitic specimens as well as the general strain-rate-hardening tendency of the martensitic specimens are evident. As seen in Fig. 2a, serrations with increased amplitude on the quasi-static compressive engineering stress–strain curve of the BMGC specimen (denoted as BMGC-QS in Fig. 2a) can be observed after yielding. It is well known that serrated flow is the characteristic plastic deformation behavior of BMGs and BMGCs at room temperature and at a lower strain rate, which correlates with the shear-band-based deformation mode in their glassy phase [5,6]. The appearance of serrations on the BMGC-QS stress–strain curve in Fig. 2a should be an indicator for the repeated shear-banding in the BMGC specimen during plastic deformation. For the martensitic and austenitic specimens upon the quasi-static compression, however, no serration appeared on their stress–strain curves (denoted as Martensite-QS and Austenite-QS in Fig. 2a), which corresponds with their dislocation-based deformation mode. In addition, serrations seemingly appeared on the three dynamic compressive engineering stress–strain curves (denoted as BMGCD, Martensite-D and Austenite-D in Fig. 2a). It is worth mentioning that the serrations should not be similarly attributed to the occurrence of repeated shear-banding in the three specimens because the serrations started to appear from the very beginning of the stress–strain curves. The seeming serrations might be possibly related with the load fluctuation during the dynamic loading using the SHPB apparatus [28], which did not affect the obtained results for the intrinsic dynamic behaviors of the specimens. Micromechanism for the significant plasticity and pronounced work-hardening of CuZr-based BMGCs containing a B2-CuZr phase under quasi-static loading has been deeply investigated recently [19,20,23–26]. It is expected that the quasi-static mechanical behavior of the CuZr-based BMGCs is derived from the in-situ competition and interaction between deformation-induced martensitic transformation of the B2-CuZr phase and shear-band multiplication in the glassy matrix. When a shear band is generated and propagates through a CuZr-based BMGC specimen, it has a possibility to reach a B2-CuZr precipitate. The crystalline phase will act as a barrier for the intrinsically rapid propagation of the shear band because the phase is quite ductile and can accommodate the strain [19]. The arrest of the shear band propagation by the “blocking effect” will provide an opportunity for the multiplication of shear bands. Furthermore, the martensitic transformation of the metastable B2-CuZr phase during the deformation can make the deformed regions harder than the undeformed regions, which can cause higher stress for further shear-banding. The sustained development of the in-situ interplay between the martensitic transformation and the shear-banding can endow the BMGCs with significant deforming ability and pronounced work-hardening characteristic under quasi-static compression and even tension. As CuZr-based BMGC alloys are heat-treated into an austenite state mainly containing the B2-CuZr phase, rather larger plasticity and higher fracture strength obtained upon quasi-static loading are primarily attributed to the deformation-induced martensitic transformation of the B2-CuZr phase [20,42,44]. Fig. 3 shows the XRD patterns of the specimens in BMGC and austenite states after quasi-static and dynamic compressive fracture. Compared with their original XRD patterns before compression in Fig. 1a, it can be clearly seen that a deformation-induced martensitic transformation from the B2-CuZr to the B190 -CuZr phase occurred in both the BMGC and austenitic specimens under both the quasi-static and dynamic compression. Based on comparing the relative diffraction intensities of the two phases, only the transformation extent under dynamic loading might be lower than that under quasi-static loading. According to the results in Fig. 2, the martensitic transformation under dynamic compression did
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Fig. 3. XRD patterns of the BMGC and austenitic specimens after quasi-static (denoted as QS) and dynamic (denoted as D) compressive fracture.
not cause similar significant plasticity and pronounced workhardening in the BMGC like that under quasi-static compression. On the contrary, strain-rate-softening occurred in the BMGC and austenite states. Effect of strain rate on the in-situ interplay between the B2-CuZr phase martensitic transformation and the glassy matrix shear-banding in the BMGC is shown in Fig. 4. Fig. 4a and b shows the SEM images of lateral surface of a quasi-statically fractured BMGC specimen. The inset in Fig. 4a displays the shear-fracture mode of the specimen. Shear-band blocking effect of the B2-CuZr phase and plenty of shear bands can be clearly observed on the specimen (Fig. 4a). Slats of martensite lath inside a deformed spherical particle of the parent B2-CuZr phase is visible (Fig. 4b), indicating the occurrence of deformation-induced martensitic transformation. As the results of the “blocking effect” and the “martensitic transformation effect”, multiple shear bands containing secondary and even ternary shear bands formed in the surrounding region of the crystalline particle (Fig. 4b). The shear-band multiplication characteristic of the BMGC specimen in Fig. 4a and b corresponds to its significant plasticity and pronounced work-hardening behavior under quasi-static compression. Fig. 4c and d shows the SEM images of lateral surface of a dynamically fractured BMGC specimen. The inset in Fig. 4c displays that the specimen similarly fractured in a shear mode. It can be seen that the shear-band density on the specimen is much lower than that on the quasi-statically fractured specimen (comparing Fig. 4c with a). It appears that the lower density of shear bands on the dynamically fractured specimen is due to a few secondary shear bands were generated in the surrounding region of a B2-CuZr phase particle (Fig. 4d), even though the primary shear bands had been blocked by the B2-CuZr phase particle, and slats of martensite lath inside the parent B2-CuZr phase is still visible (Fig. 4d). The shear-band multiplication characteristic of the BMGC specimen under dynamic compression in Fig. 4c and d corresponds to its deteriorated plasticity and work-hardening characteristic upon dynamic loading. Based on the information in Fig. 4, it appears that the “blocking effect” and the “deformation-induced martensitic transformation effect” of the B2-CuZr phase did not effectively contribute to the shear-band multiplication in the BMGC under dynamic loading. According to Spaepen's free-volume model [45], the initiation and propagation of shear bands in metallic glasses are correlated with the competition between stress-biased creation and diffusional annihilation of excess free-volume via discrete atomic jumps. According to the cooperative shearing model based on sheartransformation-zones (STZs), the generation of shear bands in metallic glasses associates with the spontaneous and cooperative
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Fig. 4. SEM images of lateral surfaces of the fractured BMGC specimens after quasi-static compression (a) and (b), and after dynamic compression (c) and (d). The insets in (a) and (c) showing the corresponding SEM images of the lateral surfaces under low magnification.
Fig. 5. SEM fractographs of the BMGC specimens after quasi-static compression (a), and after dynamic compression (b).
reorganization of STZs in aid of free-volumes when subjected to extrinsic shear stress [2]. Anyway, shear banding in metallic glasses should be a thermally activated and time-dependent process. In contrast, martensitic transformation of the B2-CuZr phase should be time-independent. For the quasi-static loading on the BMGC, there is enough time for the in-situ interplay between the martensitic transformation of the B2-CuZr phase and the shear-banding in the glassy matrix, causing effective multiplication of shear bands in the BMGC under quasi-static compression. However, under dynamic loading, shear banding will not have sufficient time to operate [46,47], the martensitic transformation of the B2-CuZr phase will not perform its expected contribution to shear band multiplication, and thus causing the plasticity and work-hardening deterioration of the BMGC under dynamic compression. Fractography can provide important information for understanding the deformation features in BMGs and BMGCs. Effect of strain rate on the fracture surface morphology of the present BMGC after quasi-static and dynamic compression is shown in Fig. 5. As can be seen in Fig. 5a, the SEM fractograph of a quasistatically fractured BMGC specimen presents a complex morphology consisting of three different distinct patterns: vein-like,
smooth featureless and river-like features. Fig. 5b shows the SEM fractograph of a dynamically fractured BMGC specimen. It can be seen that the vein-like pattern widely spread on the fractured surface, melted belts and liquid droplets even appeared, indicating the increased area with a higher temperature rise on the surface compared with that after quasi-static loading in Fig. 5a. A characteristic vein-like pattern caused by adiabatic heating on shear fracture surfaces is considered a unique feature of the fractographs of monolithic BMGs after quasi-static compression. After dynamic compression, the most striking feature of the fractographs is the co-appearance of melted belts and liquid droplets, implying much higher temperature rises occur on the shear fracture surfaces of monolithic BMGs due to the higher rate of energy input during dynamic loading [32,48,49]. This phenomenon has been correlated with the strain-rate-softening characteristic of monolithic BMGs [50]. For BMGCs after quasi-static compression, the fracture surface morphology usually consists of the vein-like, smooth featureless and river-like patterns [51], like the situation in Fig. 5a. The presence of the complex morphology indicates that the development of the fracture planes in the BMGCs occur in a stepwise mode [52]. For ZrTi-based BMGCs reinforced by ductile β-phase dendrites upon quasi-static and
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dynamic loading, no distinct fractograph difference was found [46,47]. It was reported that the temperature rise on the fracture surface of a ZrTi-based BMGC containing ductile β-phase dendrites after dynamic compression is even lower than that in quasi-static situation, which was associated with the strain-rate-hardening characteristic of the BMGC [39]. Consequently, the contrary information in Fig. 5 corresponds to the strain-rate-softening characteristic of the present BMGC containing the B2-CuZr phase. This should be correlated with the strain-rate response of the B2-CuZr phase itself. Fig. 6 shows the fracture morphology of the specimens in the austenite state after quasi-static and dynamic compression. It appears that the specimens exhibit a shear fracture mode (see the insets in Fig. 6a and c). For the quasi-statically fractured specimen, no obvious vein-like pattern but sliding trail can be observed on the fracture surface at a lower magnification (Fig. 6a). At a higher magnification, it can be seen that the sliding trail presents a viscous flow characteristic (Fig. 6b). The morphology information signifies a temperature rise on the surface. Similar viscous flow morphology was also reported on the fracture surface of a high strength Fe-based crystalline alloy with TRIP effect originating from a deformation-induced martensitic transformation [53]. It was supposed that the viscous flow morphology was related to the martensitic transformation latent heat, which could cause an adiabatic heating process, and facilitate shear-banding deformation [54]. The fracture morphology for the present austenitic specimen might be caused by the similar reason. Furthermore, melted belts and liquid droplets can be clearly observed on the dynamic fracture surface of the austenitic specimen at a lower magnification (Fig. 6c). At a higher magnification, it can be seen that the vein-like pattern even appeared (Fig. 6d), indicating a much higher temperature rise on the surface. The shear-banding deformation facilitated by the adiabatic heating would counteract the conventional strain-rate hardening induced by dislocationbased deformation in the austenitic specimen, resulting in its contrary strain-rate softening. It was reported that strain-rate responses of mechanical behaviors for some traditional TRIP crystalline alloys, e.g., TRIP steels [55], Cu-based superelastic alloys
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and austenitic stainless steels with TRIP effect [54,56], were commonly affected by the adiabatic heating of martensitic transformation. It is worth mentioning that the adiabatic heating extent in the austenitic specimens might not be as high as that in the BMGC specimens (comparing Fig. 6 with Fig. 5), thus the general strain-rate-softening tendency of the austenite is not so pronounced like that of the BMGC (see Fig. 2). Fig. 7 shows the fracture morphology of the specimens in the martensite state after quasi-static and dynamic compression. It appears that the fracture surfaces obviously deviate from their maximum shear stress plane (insets in Fig. 7a and c). For the quasistatically fractured specimen, only a small fraction of the fracture surface was covered by viscous sliding trail (Fig. 7a and b).The local viscous sliding trail had a moderately increasing tendency on the dynamic fracture surface (Fig. 7c and d). According to the XRD pattern in Fig. 1b, retained B2-CuZr phase existed in the martensitic specimens, which might also undergo deformation-induced martensitic transformation, causing the local viscous sliding trail by local adiabatic heating. Nevertheless, by comparing the information in Figs. 6 and 7, it can be reasonably inferred that the temperature rise on the fracture surfaces of the martensitic specimens should be much lower than that on the corresponding fracture surfaces of the austenitic specimens. Under the circumstances, the conventional dislocation-based deformation would dominate the strain-rate response of mechanical behavior, causing the general strain-rate-hardening tendency of the martensitic specimen. It is worth to mention that phase transition between the B2-CuZr and B190 -CuZr phases may be possibly interplayed due to the rise in temperature during dynamic loading. For the BMGC and austenitic specimens containing the B2-CuZr phase upon dynamic loading, higher temperature rise caused by the combined action of martensitic transformation with heat release effect and dynamic loading with higher energy input rate would decrease the transformation extent from the B2-CuZr phase to the B190 -CuZr phase due to the temperature-dependence of martensitic transformation. Furthermore, reverse martensitic transformation (austenitization) from the B190 -CuZr phase to the B2-CuZr phase might
Fig. 6. SEM fractographs of the austenitic specimens after quasi-static compression (a) and (b), and after dynamic compression (c) and (d). The insets in (a) and in (c) showing the corresponding SEM images of the lateral surfaces of the fractured specimens under low magnification.
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Fig. 7. SEM fractographs of the martensite specimen after quasi-static compression (a) and (b), and after dynamic compression (c) and (d). The insets in (a) and in (c) showing the corresponding SEM images of the lateral surfaces of the fractured specimens under low magnification.
occur as the increased temperature was higher than the austenite start temperature (As) of the B190 -CuZr phase. The conceivable decreased extent of martensitic transformation can be verified by comparing the relative diffraction intensity of the deformationinduced martensitic B190 -CuZr phase under dynamic loading with that under quasi-static loading in Fig. 3. The decrease of martensitic transformation extent would reduce the TRIP effect of the B2-CuZr phase, resulting in the strain-rate-softening of the BMGC and austenitic specimens. For the martensitic specimen containing the B190 -CuZr phase upon dynamic loading, however, much lower temperature rise occurred, which would have lesser effect on the constituent phase in the specimen, causing the usual strain-ratehardening like that in traditional crystalline alloys. Based on the descriptions mentioned above, the B2-CuZr phase itself exhibits strain-rate-softening because of the deformationinduced martensitic transformation of the phase with remarkable temperature rising. This characteristic would limit the contribution of the phase to the mechanical behavior of the CuZr-based BMGC under dynamic loading. Thus, compared with the ZrTibased BMGCs reinforced by tungsten phase or β-phase dendrites, the new CuZr-based BMGC containing the B2-CuZr phase presents a different SRS.
there is not enough time for the in-situ interplay between the B2-CuZr phase martensitic transformation and the glassy matrix shear-banding in the BMGC under dynamic loading, which would cause the declined shear band multiplication in the BMGC. Furthermore, the BMGC alloy was aboratively heat-treated into austenite and martensite states mainly containing the B2-CuZr phase and the B190 -CuZr phase respectively. By comparing the quasi-static and dynamic mechanical behaviors in the austenite and martensite states, it was found that the B2-CuZr phase itself presents dynamic softening. The mechanism was supposed to be correlated with the heat release effect of the deformation-induced martensitic transformation. The martensitic transformation could cause an adiabatic heating process, and facilitate shear-banding deformation in the austenite. The shear-banding deformation would counteract the conventional strain-rate hardening induced by dislocation-based deformation in the austenite, resulting in the contrary strain-rate softening. The dynamic softening characteristic of the B2-CuZr phase limits its contribution to the potential mechanical behavior of the CuZr-based BMGC containing the B2-CuZr phase under dynamic loading.
Acknowledgments 4. Conclusions Quasi-static and dynamic mechanical behaviors of a Cu48 Zr48Al4 bulk metallic glass composite (BMGC) containing a B2-CuZr phase was investigated under compression. The results demonstrate that the CuZr-based BMGC displays remarkable strain-rate-softening, although the deformation-induced martensitic transformation from the B2-CuZr phase to the B190 -CuZr phase similarly occurred under both quasi-static and dynamic compression. The significant transformation-induced plasticity (TRIP) and pronounced work-hardening characteristic of the BMGC upon quasi-static loading were found to be considerably deteriorated under dynamic loading. It should be expected that
The authors would like to gratefully acknowledge the financial support by the National Natural Science Foundation of China (NSFC) under Grant nos. 51171137 and 50671076, and the Specialized Research Fund for the Doctoral Program of Higher Education (SRFDP) under Grant no. 20110201110001. References [1] [2] [3] [4]
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Strain rate dependence of mechanical behavior in a CuZr-based bulk metallic glass composite containing B2-CuZr phase R. Wei a, S. Yang a, C.J. Zhang b, L. He a,n a b
State Key Laboratory for Mechanical Behavior of Materials, Xi’an Jiaotong University, Xi'an 710049, China School of Materials Science and Engineering, Chang'an University, Xi'an 710064, China
art ic l e i nf o
a b s t r a c t
Article history: Received 23 January 2014 Received in revised form 18 March 2014 Accepted 25 March 2014 Available online 1 April 2014
Strain rate effect on the mechanical behavior of a Cu48Zr48Al4 bulk metallic glass composite (BMGC) containing a B2-CuZr phase was investigated under quasi-static and dynamic compression. It was found that the CuZr-based BMGC exhibits remarkable strain-rate-softening, although the deformation-induced martensitic transformation of the B2-CuZr phase similarly occurred under both quasi-static and dynamic loading. The significant transformation-induced plasticity (TRIP) and pronounced work-hardening characteristic of the BMGC upon quasi-static loading were found to be considerably deteriorated under dynamic loading. Furthermore, the BMGC alloy was aboratively heat-treated into austenite and martensite states mainly containing the B2-CuZr phase and the B190 -CuZr phase respectively. By comparing the strain-rate responses of the mechanical behaviors in the austenite and martensite states, it was shown that the B2-CuZr phase itself presents dynamic softening due to the deformation-induced martensitic transformation with heat release effect. The dynamic softening characteristic of the B2-CuZr phase limits its contribution to the potential mechanical behavior of the CuZr-based BMGC containing the B2-CuZr phase under dynamic loading. & 2014 Elsevier B.V. All rights reserved.
Keywords: Mechanical characterization Bulk amorphous alloys Composites Dynamic loading Shear bands Martensitic transformation
1. Introduction Bulk metallic glasses (BMGs) have attracted great attention owing to their excellent strength and elasticity, which make them attractive as potential candidates for structural and functional applications [1–4]. However, BMGs can only accommodate plastic strain in a highly localized shear-banding manner due to the lack of dislocations in their long-range disordering atomic structure [5]. The shear-bandbased deformation mode makes BMGs suffering from limited plasticity and work-softening [6]. Extensive efforts on BMGs have been devoted to improving their mechanical behavior. Among many approaches, the fabrication of BMG-based composites (BMGCs) consisting of harder glassy matrixes and softer second crystalline phases has been proved to be effective in enhancing the plasticity and changing the work-softening of BMGs [7]. For e.g., Zr-based BMGCs reinforced with refractory metal phases were firstly fabricated and present better plasticity under compression [8–10]. Zr-based and Ti-based BMGCs reinforced by ductile β-phase dendrites can exhibit remarkable plasticity under compression and tension, and obvious work-hardening under compression [11–14]. The improved mechanical behaviors are fundamentally attributed to the introduction of
n
Corresponding author. Tel.: þ 86 29 82665165; fax: þ 86 29 826634533. E-mail address: [email protected] (L. He).
http://dx.doi.org/10.1016/j.msea.2014.03.092 0921-5093/& 2014 Elsevier B.V. All rights reserved.
dislocation-based deformation mode of the crystalline phases into the BMGCs [15,16]. Recently, a new strategy was proposed to form the BMGCs reinforced with a shape memory B2 phase [17]. New CuZr-based or TiCuNi-based BMGCs containing B2-CuZr phase or B2-Ti(Cu,Ni) phase can display significant plasticity and pronounced workhardening under compression [18–22], and especially exhibit alleviated work-softening behavior under tension testing [23–26]. It has been identified that deformation-induced martensitic transformation of the B2 phase, besides its dislocation-based deformation, is responsible for the unusual mechanical behavior of the BMGCs [16,20], which is similar to the transformation-induced plasticity (TRIP) in some traditional crystalline alloy systems [27]. This opens up a promising and innovative route for developing high-performance BMGs as engineering materials [7]. Another characteristic of BMGs concerns their strain rate dependence of mechanical behavior. Much current research has studied the dynamic uniaxial compressive mechanical behavior of BMGs by means of split-Hopkinson pressure bar (SHPB) apparatus. Although there is still some debate on the strain rate sensitivity (SRS) of BMGs, most experimental results indicate that BMGs performed apparent strain-rate-softening (i.e. negative SRS) as the strain rate beyond a critical range [28–34]. Differing from the strain-rate-hardening in traditional crystalline alloys with dislocation-based deformation mode, the strain-rate-softening of BMGs is also attributed to their
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shear-band-based deformation mode [32–34]. By modeling the deformation behavior of BMGCs [18,19,25], it is expected that the SRS of BMGCs can be affected by the competition of the shearband-based deformation within the glassy matrixes and the dislocation-based deformation within the crystalline phases. This has been demonstrated in some BMGCs. For e.g., the Zr-based BMGCs containing tungsten phase or β-phase dendrites can exhibit positive SRS [35–39]. A Ti-based BMGC containing βphase dendrites can present not only strain-rate-hardening but also distinguished work-hardening capacity upon dynamic loading [40]. For the CuZr-based or TiCuNi-based BMGCs containing a B2 phase and with TRIP effect, it can be reasonably inferred that the SRS may be additionally affected by the martensitic transformation of the B2 phase. However, so far, little attention has been devoted to this issue. In the present work, the strain rate dependence of mechanical behavior in a CuZr-based BMGC alloy containing a B2CuZr phase was investigated. In addition, the mechanical responses to strain rate for the austenitic B2-CuZr phase and the corresponding martensitic B190 -CuZr phase were compared, martensitic transformation effect on the SRS of the BMGC was analyzed.
2. Experimental methods A representative CuZr-based alloy with a nominal composition of Cu48Zr48Al4 was chosen for the current investigation because of its capability of obtaining the as-cast BMGC microstructure with a single B2-CuZr phase embedded in a glassy matrix at a suitable cooling rate [24]. The master alloy ingot was produced by arcmelting Zr, Cu and Al of 99.99% purity in a Ti-gettered argon atmosphere. As-cast cylindrical rods with a diameter of Ф3 mm were prepared via suction casting using a copper mold. In order to explore the martensitic transformation effect of the B2-CuZr phase on the SRS of the Cu48Zr48Al4 BMGC, some of the as-cast rods were aboratively heat-treated to have the microstructures mainly containing the austenitic B2-CuZr phase or the corresponding martensitic B190 -CuZr phase respectively. According to the binary Cu–Zr phase diagram (http://www1.asminterna tional.org/AsmEnterprise/APD), the B2-CuZr is a high-temperature metastable phase and should decompose into the low-temperature equilibrium phases, Cu10Zr7 and CuZr2, via a eutectoid reaction at 988 K if the system is given the time to equilibrate. For the present ternary Cu48Zr48Al4 alloy, the eutectoid reaction was found to occur within a temperature range from 977 K to 996 K [17]. Thus, some of the as-cast rods were annealed at 1073 K (above the eutectoid temperature range of the alloy) for 60 min to form the high-temperature B2-CuZr phase, and then water-quenched to room temperature. The water-quenching could limit the eutectoid
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reaction, the microstructure mainly containing the martensitic B190 -CuZr phase could be obtained because the martensitic transformation start temperature (Ms) from the B2-CuZr phase to the B190 -CuZr phase in the present Cu48Zr48Al4 alloy is higher than room temperature (about 337 K) [24,26]. To obtain the microstructure mainly containing the austenitic B2-CuZr phase at room temperature, the shape memory effect arising in CuZr-based alloys was exploited [19,41]. It had been testified that the Ms temperature of the B2-CuZr phase could be effectively lowered by repeated martensitic transformation of the phase induced by thermal cycling, and the high-temperature austenitic phase could be stabilized to room temperature [19,42]. Thus, some of the annealed rods were further thermally cycled between 673 K and 77 K, the temperatures are much higher than the austenitizing temperature of the B190 -CuZr phase and much lower than the martensitic transformation temperature of the B2-CuZr phase respectively [19,24]. The thermal cycling treatment was performed by heating the annealed rods at 673 K for 10 min and then quenched in liquid nitrogen for another 10 min. The rods were heat-treated that way at least six times. The microstructures of the as-cast and heat-treated rods were examined by X-ray diffraction (XRD) using an X'Pert-Pro diffractometer with Cu-Kα radiation and observed by scanning electron microscopy (SEM) using a JSM-7000F microscope. The crystallization characteristics of the rods were identified by differential scanning calorimetry (DSC) using a SETARAM-LabsysTM TG-DSC at a heating rate of 0.677 K/s with a sample mass of 3071 mg. Cylindrical specimens with the aspect ratios of 2 and 1 were cut from the as-cast and heat-treated rods. Quasi-static compression tests were conducted using the higher-aspect-ratio specimens on a screw-driven SUNS CMT5105 testing machine at a constant crosshead displacement rate of 0.05 mm/min, which corresponds to an initial strain rate of 1.39 10 4 s 1. Dynamic compression tests were performed using the lower-aspect-ratio specimens on a split Hopkinson pressure bar (SHPB) apparatus at a strain rate range of 2–3.3 103 s 1, the detailed process was described elsewhere [43]. In order to observe the deforming characteristic of glassy matrix and B2-CuZr phase in the CuZr-based BMGC, the side surfaces of some ascast specimens were previously polished, and examined by the SEM observation after both quasi-static and dynamic compression tests.
3. Results and discussion The XRD patterns of representative Ф3 mm Cu48Zr48Al4 alloy specimens in as-cast and heat-treated states are shown in Fig. 1a. For the as-cast specimen, crystalline diffraction peaks superimposed on a broad amorphous scattering maximum are displayed,
Fig. 1. (a) XRD patterns of representative ∅3 mm Cu48Zr48Al4 alloy specimens in as-cast, annealed and annealed plus thermally cycled states. (b) DSC curves of the specimens and the corresponding as-spun ribbon. (c) SEM image of the as-cast specimen.
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the Bragg peaks can be allocated to a cubic primitive B2-CuZr (Pm-3m) phase. For the annealed specimen, the XRD pattern consists of stronger diffraction peaks of a B190 -CuZr (P21/m and Cm) phase and relatively weaker diffraction peaks of retained B2CuZr phase. This is because the martensitic transformation finish temperature (Mf) of the B2-CuZr phase in the present Cu48Zr48Al4 alloy is lower than room temperature [24,26]. It is difficult to exactly determine the percentage of the retained B2-CuZr phase in the annealed specimen just based on the XRD pattern. For the annealed plus thermally cycled specimen, the XRD pattern was remarkably changed, consisting of stronger diffraction peaks of a B2-CuZr phase and relatively weaker diffraction peaks of retained B190 -CuZr phase. This is because the shape recovery of the present Cu48Zr48Al4 alloy is not complete [19], a small amount of the martensitic B190 -CuZr phase did not transform back to the austenitic B2-CuZr phase after the thermal cycling. The XRD patterns of the two heat-treated specimens are consistent with the results in Ref. [19]. Fig. 1b shows the DSC curves of the as-cast and heattreated specimens at a heating rate of 0.677 K/s. For the sake of comparison, the DSC curve for the corresponding as-spun glassy ribbon is also presented. It can be seen that the as-cast specimen and the ribbon began glass transition at a very close temperature of Tg 698 K, and exhibit an almost same crystallization temperature of TX 753 K. However, the crystallization peak area of the BMGC specimen distinctly reduces compared with that of the glassy ribbon. Based on the crystallization enthalpy reduction, the B2-CuZr phase volume fraction in the as-cast is estimated to be about 31%. Fig. 1c (inset in Fig. 1b) shows the SEM image of the ascast specimen in a backscattered electron (BSE) mode. It displays that some dark B2-CuZr precipitates with patch-like shape are embedded in a brighter, featureless matrix. The information in Fig. 1a and c indicates that the as-cast specimen is an in-situ BMGC with the B2-CuZr phase embedded in the glassy matrix. For the two heat-treated specimens, no apparent crystallization peak appeared on their DCS curves, which means that very limited amount of glassy phase exists in the specimens because they all underwent long time annealing at the temperature of 1073 K. Although desirable single B190 -CuZr phase in the annealed specimen and single B2-CuZr phase in the annealed plus thermally cycled specimen were not obtained, previous research testified that a small amount of the retained B2-CuZr phase or B190 -CuZr phase did not obviously affect the typical mechanical behaviors of the two heat-treated specimens under quasi-static compression [19,42]. Thus, it can be reasonably expected that mechanical behaviors of the two heat-treated specimens during dynamic compression would be similarly controlled by their predominant crystalline phases respectively. Based on the descriptions mentioned above, the as-cast, annealed and annealed plus thermally
cycled specimens are considered to be in BMGC, martensite and austenite states respectively in the following description. Fig. 2a shows the quasi-static and dynamic compressive engineering stress–strain curves of the specimens in the BMGC, martensite and austenite states. Upon quasi-static compression, it can be seen that the BMGC specimen displays a yield stress of about 1580 MPa (characterized by s0.2) and a fracture strain of about 18.5%. Furthermore, based on the corresponding true stress– strain curve (not shown here), the BMGC specimen also exhibits a pronounced work-hardening behavior and a resultant maximum stress of about 2200 MPa, which are in agreement with former reports [18,19]. For the martensitic and austenitic specimens under quasi-static compression, it is displayed that the yield strength of the austenite specimen is much lower than that of the martensitic specimen. However, the softer austenite specimen exhibits rather larger plasticity than that of the martensitic specimen, accompanying by strong work-hardening and a maximum stress very close to that of the martensitic specimen before finally fractured. The mechanical behaviors of the martensitic and austenitic specimens upon quasi-static compressive loading mentioned above are also consistent with former reports [19,42]. Upon dynamic compression, the representative engineering stress–strain curves in Fig. 2a illustrate that remarkable strainrate responses for the specimens in the BMGC, martensite and austenite states all occurred. Compared with the corresponding quasi-static stress–strain curves, the following phenomena can be observed. For the specimens in BMGC state, the significant plasticity and pronounced work-hardening under quasi-static compression are considerably deteriorated under dynamic compression. The maximum stress under the dynamic compression is even lower than the yield stress under quasi-static compression, indicating remarkable strain-rate-softening of the BMGC. For the specimens in martensite state, remarkably increased maximum stress upon dynamic loading is presented, indicating obvious strain-rate-hardening in the martensite state, like that in traditional crystalline alloys with dislocation-based deformation mode. However, for the specimen in austenite state, apparently decreased maximum stress upon dynamic loading is exhibited, revealing contrary strain-rate-softening in the austenite state. The strain rate dependences of maximum stresses for the BMGC, martensitic and austenitic specimens are displayed in Fig. 2b. At least nine specimens in each state were tested under both quasi-static and dynamic loading to verify the scattering of the compression data. The scattering of the quasi-static compression data at the strain rate of 1.39 10 4 s 1 for each state of specimens is displayed in error bars in Fig. 2b, the dynamic compression data at the strain rate range of 2–3.3 103 s 1 for all the specimens in three states are plotted in the figure. In despite of the
Fig. 2. (a) Quasi-static (denoted as QS) and dynamic (denoted as D) compressive engineering stress–strain curves of the specimens in different-states, the strain rate for the dynamic compression was about 2.8 103 s 1. (b) The dependences of maximum flow stresses on the strain rate for the specimens in different states.
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data scatter under dynamic loading, the general strain-ratesoftening tendency of the BMGC and austenitic specimens as well as the general strain-rate-hardening tendency of the martensitic specimens are evident. As seen in Fig. 2a, serrations with increased amplitude on the quasi-static compressive engineering stress–strain curve of the BMGC specimen (denoted as BMGC-QS in Fig. 2a) can be observed after yielding. It is well known that serrated flow is the characteristic plastic deformation behavior of BMGs and BMGCs at room temperature and at a lower strain rate, which correlates with the shear-band-based deformation mode in their glassy phase [5,6]. The appearance of serrations on the BMGC-QS stress–strain curve in Fig. 2a should be an indicator for the repeated shear-banding in the BMGC specimen during plastic deformation. For the martensitic and austenitic specimens upon the quasi-static compression, however, no serration appeared on their stress–strain curves (denoted as Martensite-QS and Austenite-QS in Fig. 2a), which corresponds with their dislocation-based deformation mode. In addition, serrations seemingly appeared on the three dynamic compressive engineering stress–strain curves (denoted as BMGCD, Martensite-D and Austenite-D in Fig. 2a). It is worth mentioning that the serrations should not be similarly attributed to the occurrence of repeated shear-banding in the three specimens because the serrations started to appear from the very beginning of the stress–strain curves. The seeming serrations might be possibly related with the load fluctuation during the dynamic loading using the SHPB apparatus [28], which did not affect the obtained results for the intrinsic dynamic behaviors of the specimens. Micromechanism for the significant plasticity and pronounced work-hardening of CuZr-based BMGCs containing a B2-CuZr phase under quasi-static loading has been deeply investigated recently [19,20,23–26]. It is expected that the quasi-static mechanical behavior of the CuZr-based BMGCs is derived from the in-situ competition and interaction between deformation-induced martensitic transformation of the B2-CuZr phase and shear-band multiplication in the glassy matrix. When a shear band is generated and propagates through a CuZr-based BMGC specimen, it has a possibility to reach a B2-CuZr precipitate. The crystalline phase will act as a barrier for the intrinsically rapid propagation of the shear band because the phase is quite ductile and can accommodate the strain [19]. The arrest of the shear band propagation by the “blocking effect” will provide an opportunity for the multiplication of shear bands. Furthermore, the martensitic transformation of the metastable B2-CuZr phase during the deformation can make the deformed regions harder than the undeformed regions, which can cause higher stress for further shear-banding. The sustained development of the in-situ interplay between the martensitic transformation and the shear-banding can endow the BMGCs with significant deforming ability and pronounced work-hardening characteristic under quasi-static compression and even tension. As CuZr-based BMGC alloys are heat-treated into an austenite state mainly containing the B2-CuZr phase, rather larger plasticity and higher fracture strength obtained upon quasi-static loading are primarily attributed to the deformation-induced martensitic transformation of the B2-CuZr phase [20,42,44]. Fig. 3 shows the XRD patterns of the specimens in BMGC and austenite states after quasi-static and dynamic compressive fracture. Compared with their original XRD patterns before compression in Fig. 1a, it can be clearly seen that a deformation-induced martensitic transformation from the B2-CuZr to the B190 -CuZr phase occurred in both the BMGC and austenitic specimens under both the quasi-static and dynamic compression. Based on comparing the relative diffraction intensities of the two phases, only the transformation extent under dynamic loading might be lower than that under quasi-static loading. According to the results in Fig. 2, the martensitic transformation under dynamic compression did
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Fig. 3. XRD patterns of the BMGC and austenitic specimens after quasi-static (denoted as QS) and dynamic (denoted as D) compressive fracture.
not cause similar significant plasticity and pronounced workhardening in the BMGC like that under quasi-static compression. On the contrary, strain-rate-softening occurred in the BMGC and austenite states. Effect of strain rate on the in-situ interplay between the B2-CuZr phase martensitic transformation and the glassy matrix shear-banding in the BMGC is shown in Fig. 4. Fig. 4a and b shows the SEM images of lateral surface of a quasi-statically fractured BMGC specimen. The inset in Fig. 4a displays the shear-fracture mode of the specimen. Shear-band blocking effect of the B2-CuZr phase and plenty of shear bands can be clearly observed on the specimen (Fig. 4a). Slats of martensite lath inside a deformed spherical particle of the parent B2-CuZr phase is visible (Fig. 4b), indicating the occurrence of deformation-induced martensitic transformation. As the results of the “blocking effect” and the “martensitic transformation effect”, multiple shear bands containing secondary and even ternary shear bands formed in the surrounding region of the crystalline particle (Fig. 4b). The shear-band multiplication characteristic of the BMGC specimen in Fig. 4a and b corresponds to its significant plasticity and pronounced work-hardening behavior under quasi-static compression. Fig. 4c and d shows the SEM images of lateral surface of a dynamically fractured BMGC specimen. The inset in Fig. 4c displays that the specimen similarly fractured in a shear mode. It can be seen that the shear-band density on the specimen is much lower than that on the quasi-statically fractured specimen (comparing Fig. 4c with a). It appears that the lower density of shear bands on the dynamically fractured specimen is due to a few secondary shear bands were generated in the surrounding region of a B2-CuZr phase particle (Fig. 4d), even though the primary shear bands had been blocked by the B2-CuZr phase particle, and slats of martensite lath inside the parent B2-CuZr phase is still visible (Fig. 4d). The shear-band multiplication characteristic of the BMGC specimen under dynamic compression in Fig. 4c and d corresponds to its deteriorated plasticity and work-hardening characteristic upon dynamic loading. Based on the information in Fig. 4, it appears that the “blocking effect” and the “deformation-induced martensitic transformation effect” of the B2-CuZr phase did not effectively contribute to the shear-band multiplication in the BMGC under dynamic loading. According to Spaepen's free-volume model [45], the initiation and propagation of shear bands in metallic glasses are correlated with the competition between stress-biased creation and diffusional annihilation of excess free-volume via discrete atomic jumps. According to the cooperative shearing model based on sheartransformation-zones (STZs), the generation of shear bands in metallic glasses associates with the spontaneous and cooperative
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Fig. 4. SEM images of lateral surfaces of the fractured BMGC specimens after quasi-static compression (a) and (b), and after dynamic compression (c) and (d). The insets in (a) and (c) showing the corresponding SEM images of the lateral surfaces under low magnification.
Fig. 5. SEM fractographs of the BMGC specimens after quasi-static compression (a), and after dynamic compression (b).
reorganization of STZs in aid of free-volumes when subjected to extrinsic shear stress [2]. Anyway, shear banding in metallic glasses should be a thermally activated and time-dependent process. In contrast, martensitic transformation of the B2-CuZr phase should be time-independent. For the quasi-static loading on the BMGC, there is enough time for the in-situ interplay between the martensitic transformation of the B2-CuZr phase and the shear-banding in the glassy matrix, causing effective multiplication of shear bands in the BMGC under quasi-static compression. However, under dynamic loading, shear banding will not have sufficient time to operate [46,47], the martensitic transformation of the B2-CuZr phase will not perform its expected contribution to shear band multiplication, and thus causing the plasticity and work-hardening deterioration of the BMGC under dynamic compression. Fractography can provide important information for understanding the deformation features in BMGs and BMGCs. Effect of strain rate on the fracture surface morphology of the present BMGC after quasi-static and dynamic compression is shown in Fig. 5. As can be seen in Fig. 5a, the SEM fractograph of a quasistatically fractured BMGC specimen presents a complex morphology consisting of three different distinct patterns: vein-like,
smooth featureless and river-like features. Fig. 5b shows the SEM fractograph of a dynamically fractured BMGC specimen. It can be seen that the vein-like pattern widely spread on the fractured surface, melted belts and liquid droplets even appeared, indicating the increased area with a higher temperature rise on the surface compared with that after quasi-static loading in Fig. 5a. A characteristic vein-like pattern caused by adiabatic heating on shear fracture surfaces is considered a unique feature of the fractographs of monolithic BMGs after quasi-static compression. After dynamic compression, the most striking feature of the fractographs is the co-appearance of melted belts and liquid droplets, implying much higher temperature rises occur on the shear fracture surfaces of monolithic BMGs due to the higher rate of energy input during dynamic loading [32,48,49]. This phenomenon has been correlated with the strain-rate-softening characteristic of monolithic BMGs [50]. For BMGCs after quasi-static compression, the fracture surface morphology usually consists of the vein-like, smooth featureless and river-like patterns [51], like the situation in Fig. 5a. The presence of the complex morphology indicates that the development of the fracture planes in the BMGCs occur in a stepwise mode [52]. For ZrTi-based BMGCs reinforced by ductile β-phase dendrites upon quasi-static and
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dynamic loading, no distinct fractograph difference was found [46,47]. It was reported that the temperature rise on the fracture surface of a ZrTi-based BMGC containing ductile β-phase dendrites after dynamic compression is even lower than that in quasi-static situation, which was associated with the strain-rate-hardening characteristic of the BMGC [39]. Consequently, the contrary information in Fig. 5 corresponds to the strain-rate-softening characteristic of the present BMGC containing the B2-CuZr phase. This should be correlated with the strain-rate response of the B2-CuZr phase itself. Fig. 6 shows the fracture morphology of the specimens in the austenite state after quasi-static and dynamic compression. It appears that the specimens exhibit a shear fracture mode (see the insets in Fig. 6a and c). For the quasi-statically fractured specimen, no obvious vein-like pattern but sliding trail can be observed on the fracture surface at a lower magnification (Fig. 6a). At a higher magnification, it can be seen that the sliding trail presents a viscous flow characteristic (Fig. 6b). The morphology information signifies a temperature rise on the surface. Similar viscous flow morphology was also reported on the fracture surface of a high strength Fe-based crystalline alloy with TRIP effect originating from a deformation-induced martensitic transformation [53]. It was supposed that the viscous flow morphology was related to the martensitic transformation latent heat, which could cause an adiabatic heating process, and facilitate shear-banding deformation [54]. The fracture morphology for the present austenitic specimen might be caused by the similar reason. Furthermore, melted belts and liquid droplets can be clearly observed on the dynamic fracture surface of the austenitic specimen at a lower magnification (Fig. 6c). At a higher magnification, it can be seen that the vein-like pattern even appeared (Fig. 6d), indicating a much higher temperature rise on the surface. The shear-banding deformation facilitated by the adiabatic heating would counteract the conventional strain-rate hardening induced by dislocationbased deformation in the austenitic specimen, resulting in its contrary strain-rate softening. It was reported that strain-rate responses of mechanical behaviors for some traditional TRIP crystalline alloys, e.g., TRIP steels [55], Cu-based superelastic alloys
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and austenitic stainless steels with TRIP effect [54,56], were commonly affected by the adiabatic heating of martensitic transformation. It is worth mentioning that the adiabatic heating extent in the austenitic specimens might not be as high as that in the BMGC specimens (comparing Fig. 6 with Fig. 5), thus the general strain-rate-softening tendency of the austenite is not so pronounced like that of the BMGC (see Fig. 2). Fig. 7 shows the fracture morphology of the specimens in the martensite state after quasi-static and dynamic compression. It appears that the fracture surfaces obviously deviate from their maximum shear stress plane (insets in Fig. 7a and c). For the quasistatically fractured specimen, only a small fraction of the fracture surface was covered by viscous sliding trail (Fig. 7a and b).The local viscous sliding trail had a moderately increasing tendency on the dynamic fracture surface (Fig. 7c and d). According to the XRD pattern in Fig. 1b, retained B2-CuZr phase existed in the martensitic specimens, which might also undergo deformation-induced martensitic transformation, causing the local viscous sliding trail by local adiabatic heating. Nevertheless, by comparing the information in Figs. 6 and 7, it can be reasonably inferred that the temperature rise on the fracture surfaces of the martensitic specimens should be much lower than that on the corresponding fracture surfaces of the austenitic specimens. Under the circumstances, the conventional dislocation-based deformation would dominate the strain-rate response of mechanical behavior, causing the general strain-rate-hardening tendency of the martensitic specimen. It is worth to mention that phase transition between the B2-CuZr and B190 -CuZr phases may be possibly interplayed due to the rise in temperature during dynamic loading. For the BMGC and austenitic specimens containing the B2-CuZr phase upon dynamic loading, higher temperature rise caused by the combined action of martensitic transformation with heat release effect and dynamic loading with higher energy input rate would decrease the transformation extent from the B2-CuZr phase to the B190 -CuZr phase due to the temperature-dependence of martensitic transformation. Furthermore, reverse martensitic transformation (austenitization) from the B190 -CuZr phase to the B2-CuZr phase might
Fig. 6. SEM fractographs of the austenitic specimens after quasi-static compression (a) and (b), and after dynamic compression (c) and (d). The insets in (a) and in (c) showing the corresponding SEM images of the lateral surfaces of the fractured specimens under low magnification.
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Fig. 7. SEM fractographs of the martensite specimen after quasi-static compression (a) and (b), and after dynamic compression (c) and (d). The insets in (a) and in (c) showing the corresponding SEM images of the lateral surfaces of the fractured specimens under low magnification.
occur as the increased temperature was higher than the austenite start temperature (As) of the B190 -CuZr phase. The conceivable decreased extent of martensitic transformation can be verified by comparing the relative diffraction intensity of the deformationinduced martensitic B190 -CuZr phase under dynamic loading with that under quasi-static loading in Fig. 3. The decrease of martensitic transformation extent would reduce the TRIP effect of the B2-CuZr phase, resulting in the strain-rate-softening of the BMGC and austenitic specimens. For the martensitic specimen containing the B190 -CuZr phase upon dynamic loading, however, much lower temperature rise occurred, which would have lesser effect on the constituent phase in the specimen, causing the usual strain-ratehardening like that in traditional crystalline alloys. Based on the descriptions mentioned above, the B2-CuZr phase itself exhibits strain-rate-softening because of the deformationinduced martensitic transformation of the phase with remarkable temperature rising. This characteristic would limit the contribution of the phase to the mechanical behavior of the CuZr-based BMGC under dynamic loading. Thus, compared with the ZrTibased BMGCs reinforced by tungsten phase or β-phase dendrites, the new CuZr-based BMGC containing the B2-CuZr phase presents a different SRS.
there is not enough time for the in-situ interplay between the B2-CuZr phase martensitic transformation and the glassy matrix shear-banding in the BMGC under dynamic loading, which would cause the declined shear band multiplication in the BMGC. Furthermore, the BMGC alloy was aboratively heat-treated into austenite and martensite states mainly containing the B2-CuZr phase and the B190 -CuZr phase respectively. By comparing the quasi-static and dynamic mechanical behaviors in the austenite and martensite states, it was found that the B2-CuZr phase itself presents dynamic softening. The mechanism was supposed to be correlated with the heat release effect of the deformation-induced martensitic transformation. The martensitic transformation could cause an adiabatic heating process, and facilitate shear-banding deformation in the austenite. The shear-banding deformation would counteract the conventional strain-rate hardening induced by dislocation-based deformation in the austenite, resulting in the contrary strain-rate softening. The dynamic softening characteristic of the B2-CuZr phase limits its contribution to the potential mechanical behavior of the CuZr-based BMGC containing the B2-CuZr phase under dynamic loading.
Acknowledgments 4. Conclusions Quasi-static and dynamic mechanical behaviors of a Cu48 Zr48Al4 bulk metallic glass composite (BMGC) containing a B2-CuZr phase was investigated under compression. The results demonstrate that the CuZr-based BMGC displays remarkable strain-rate-softening, although the deformation-induced martensitic transformation from the B2-CuZr phase to the B190 -CuZr phase similarly occurred under both quasi-static and dynamic compression. The significant transformation-induced plasticity (TRIP) and pronounced work-hardening characteristic of the BMGC upon quasi-static loading were found to be considerably deteriorated under dynamic loading. It should be expected that
The authors would like to gratefully acknowledge the financial support by the National Natural Science Foundation of China (NSFC) under Grant nos. 51171137 and 50671076, and the Specialized Research Fund for the Doctoral Program of Higher Education (SRFDP) under Grant no. 20110201110001. References [1] [2] [3] [4]
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