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Microstructure and stored energy evolutions during rolling of Cu60Zr20Ti20 bulk metallic glass
Journal of Non-Crystalline Solids 354 (2008) 5353–5362
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Microstructure and stored energy evolutions during rolling of Cu60Zr20Ti20 bulk metallic glass Q.P. Cao a,b, J.F. Li a,*, J.Z. Jiang b, Y.H. Zhou a a
State Key Laboratory of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai 200240, PR China International Center for New-Structured Materials (ICNSM), Zhejiang University and Laboratory of New-Structured Materials, Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, PR China b
a r t i c l e
i n f o
Article history: Received 2 April 2008 Received in revised form 9 September 2008 Available online 24 October 2008 PACS: 81.05.Kf 81.40.Ef 81.70.Pg 68.37.Lp Keywords: Amorphous metals, metallic glasses Crystallization
a b s t r a c t The microstructure and stored energy of Cu60Zr20Ti20 bulk metallic glass rolled at cryogenic temperature in a wide strain rate range 1.0 104 5.0 101 s1 have been investigated. As the specimen is rolled to be thinner, the stored energy first increases linearly, and then saturates above a critical thickness reduction at lower strain rates, or decreases at high strain rates. At the initial stage of rolling, no phase transformation except shear bands appears in the glass. Phase transformation occurs only when the specimen is severely deformed at strain rates higher than 1.0 104 s1. As strain rate increases, the critical strain for the stored energy to saturate increases, but the critical strain for phase separation to occur decreases, and meanwhile the type of the phase transformation changes from phase separation to nanocrystallization. The stored energy does not change with the occurrence of phase separation, but decreases due to nanocrystallization. It is proposed that coalescence of more free volume in shear bands into nanovoids should be principally responsible for the saturation of the stored energy, which balances the results from the increase in shear band number at higher strains. Ó 2008 Elsevier B.V. All rights reserved.
1. Introduction In comparison with the ideal metallic glass in which the atoms have the densest arrangement while the amorphous state is still held, an actual metallic glass with the same composition always contains an amount of excess free volume [1–3]. Spaepen proposed a free-volume model to account for the deformation behavior of metallic glasses at various temperatures [4], according to which, macroscopic flow of metallic glasses occurs as a result of biased jumps of atoms into their neighboring sites when the external stress is applied. A certain amount of free volume can be created as the atoms are squeezed into the neighboring sites with a smaller volume than the atoms. Van den Buekel and Sietsma interpreted enthalpy relaxation in terms of free volume, i.e. the change in the free-volume content during differential scanning calorimetry (DSC) experiment of a constant heating rate is proportional to the energy released by the structural relaxation of metallic glasses [5]. It is worth noting that the enthalpy relaxation is not fully determined by free volume and may reflect other relevant structural change as well, such as relaxation of residual stress. Casting and rapid cooling of metallic glasses can lead to residual stress
* Corresponding author. Tel./fax: +86 21 5474 8530. E-mail address: jfl[email protected] (J.F. Li). 0022-3093/$ - see front matter Ó 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jnoncrysol.2008.09.024
due to thermal tempering [6]. So can inhomogeneous deformation of metallic glasses [7]. During cold rolling of metallic glasses, very thin bands with intense shear deformation are commonly observed [8], and it definitely alters the amount and distribution of residual stresses. Therefore, for the cold-rolled metallic glasses it may be inexact to directly link the relaxation enthalpy (stored energy) determined by DSC with free volume. An earlier work on the strain energy storage in a cold-rolled Pd77.5Cu6Si16.5 glass ribbon has revealed that the stored energy monotonously increases as the thickness reduction increases to the highest value 36% [9]. However, traditional metallic glasses were fabricated by rapid solidification, and their thicknesses were seldom larger than 100 lm. The highest plastic strain realized in them could not be large, and experimental investigation of the stored energy evolution at larger plastic strains was severely confined. Recently this situation was changed as novel multicomponent bulk metallic glasses (BMGs) were discovered [10–15]. Their thicknesses more than 1 mm make it feasible to obtain large plastic strains by compression or rolling at temperatures far below the glass transition temperature. On the other hand, metallic glasses are metastable in thermodynamics. When they are undergoing plastic deformation, the atomic displacement is significantly activated by the shear stress. While part of the atoms in flow becomes more disordered as a result of the mechanical shear [8,16–18], the other atoms may get a chance
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to rearrange into another structure with a lower free energy, such as crystalline phases or new glassy phases. Mechanically induced crystallization has been observed in some metallic glasses [19,20]. The variation of the microstructure of metallic glasses with plastic deformation has been an interesting subject of materials science [21–30]. In our previous work [31,32], the Cu60Zr20Ti20 BMG was ever rolled at a fixed strain rate at room temperature (RT) as well as cryogenic temperature (CT), and it was found that the increase of rolling temperature redounded to precipitation of nanocrystals. Considering that the structural variation in metallic glasses depends on not only rolling temperature and plastic strain but also the applied strain rate, in the present work we are going to carry out a series of deformation experiments at CT on the Cu60Zr20Ti20 BMG in a particularly wide range of strain rate up to high macroscopic strains, so as to reveal the effect of strain rate and total strain on the microstructure and stored energy of the deformed metallic glass. 2. Experimental Ternary Cu60Zr20Ti20 master alloy ingots were prepared by arc melting the mixture of high purity Cu (99.99%), Zr (99.9%) and Ti (99.9%) under a Ti-gettered argon atmosphere. The ingots were inverted and remelted six times to ensure compositional homogeneity and then suck-cast into a water-cooled copper mold to produce 40 mm long cylindrical rods with a diameter of 2 mm. The amorphous structure was identified by X-ray diffraction (XRD) with monochromatic Cu Ka radiation and high resolution transmission electron microscopy (HRTEM). The rods were cut into short cylinders with a thickness of 1.5 mm for rolling. Both ends of the cylinders were mechanically polished to make them parallel to each other prior to the rolling experiment. Before the rolling deformation, a continuous liquid nitrogen stream was used to cool the specimen, by which the temperature of the specimen was measured by a K-type thermocouple to be about 150 K. During the subsequent rolling process, the specimen was incessantly cooled with the liquid nitrogen stream. Details of the deformation procedure were described elsewhere [32]. The degree of deformation was denoted by the reduction in thickness e = (h0 h)/h0, where h0 and h represented the specimen thicknesses before and after rolling, respectively. Many small deformation passes were used with progressively narrowing the gap between two rollers, and the decrease of the gap during deformation was carefully controlled. The strain rate e_ adopted in the experiment was about 1.0 104, 5.0 104, 5.0 103 and 5.0 101 s1, respectively. At a higher strain rate (e.g., e_ ¼ 1:0 s1 ), the specimen fractured abruptly at about 45to the direction of rolling pressure in the initial rolling stage. The microstructures of the as-rolled specimens were examined by JEOL JEM-3010 HRTEM apparatus operating at 300 kV. The difficulty in preparing a good TEM specimen of the Cu–Zr–Ti alloy is well known just as the previous studies on this alloy system have pointed out that a nanocrystalline microstructure might form when the specimen was prepared without special attention [33]. The thin foils for TEM were prepared by twin-jet electropolishing using a solution of 6% nitric acid in methanol at a temperature of 238 K. The specimens were observed in TEM immediately after the preparation since the Cu–Zr–Ti thin foil readily oxidizes upon exposure in air atmosphere. Selected area electron diffraction (SAED) patterns were taken from an area about 1 lm in diameter. Thermal analyses of the specimens subjected to different degrees of deformation were performed in a Perkin–Elmer Pyris Diamond DSC under a flow of purified argon. The specimens were heated at a heating rate of 20 K/min to 853 K, cooled to RT and then
reheated at the same rate to establish a baseline. The baseline was subtracted from the first scan to better describe the thermal behavior of the glass. Five samples at each strain were measured. The sample and reference pans were made of aluminium. The temperature and the heat flow were calibrated by measuring the melting temperatures and the heats of fusion of pure In, Sn and Zn. 3. Results 3.1. DSC curves The as-cast Cu60Zr20Ti20 rods have been verified by X-ray diffractometry (XRD) and HRTEM to be fully amorphous (the results are not shown). In the strain rate range from 1.0 104 to 5.0 101 s1, the Cu60Zr20Ti20 BMG can be rolled at CT up to 97% reduction in thickness without fracture. Fig. 1(a)–(d) shows the typical DSC curves of the specimens rolled at various strain rates. The data of the DSC curves are the average of five measured DSC curves at each strain. The as-cast alloy exhibits a glass transition with the onset temperature Tg = 706 K, followed by two exothermic events, characterized by the onset temperature for crystallization Tx = 736 K, the peak temperature for the first crystallization event Tp1 = 748 K and the peak temperature for the second crystallization event Tp2 = 786 K. It has been known that the first exothermic event corresponds to the amorphous-to-Cu51Zr14 phase transition while the second one is due to the crystallization of the residual amorphous phase [34,35]. The peak temperature and enthalpy of the second exothermic event do not change with the deformation no matter how high the e_ is. Here, we only present the measured strain dependences of the peak temperature and enthalpy of the first exothermic event at various strain rates in Fig. 1(e) and (f), respectively. The error bars of the peak temperature and enthalpy of each strain are standard deviations of five measured samples. At the lowest strain rate e_ ¼ 1:0 104 s1 , the first peak temperature keeps unchanged up to the highest e of 97%. As e_ increases from 1.0 104 to 5.0 104 s1, the first peak temperature still remains constant at e 6 93%, but decreases by about 1.0 K when e increases up to 97%. When e_ further increases from 5.0 104 to 5.0 103 s1, the starting e for peak temperature to decline varies from 93% to 89%, and the decrease amount at e = 97% enlarges from 1.0 to 1.5 K. However, the enthalpy of the first exothermic event remains unchanged. It is clear that the rolling at a higher e_ leads to a heavier decrease of the glass in thermal stability when e exceeds the critical value. In the experiment we raised the strain rate abruptly to 5.0 101 s1, the critical e for the first peak temperature to decline drops to 67%, and the decrease of the first peak temperature reaches as large as 4.7 K at e = 97%. An important finding is that the enthalpy of the first exothermic event at this strain rate declines with increasing e beyond the critical strain 67%, and a change of about 4.7% is reached at e = 97%. 3.2. Microstructure evolution with rolling It has been observed that when the Cu60Zr20Ti20 BMG is rolled at CT and e_ ¼ 5:0 103 s1 , mechanically driven phase separation occurs with e > 89% [32]. The occurrence of phase separation does not change the enthalpy of the first exothermic event, but decreases the peak temperature. It is therefore speculated that at the lowest e_ ¼ 1:0 104 s1 , no phase transformation occurs in the Cu60Zr20Ti20 specimen during the whole rolling process since both the peak temperature and enthalpy of the first exothermic event are unchanged. For e_ ¼ 5:0 104 and 5.0 103 s1, the deformation should be accompanied with phase separation when e exceeds 93% and 89%, respectively. In view of the obvious
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Fig. 1. (a)–(d) DSC curves of the as-rolled Cu60Zr20Ti20 specimens with different e at various strain rates of 1.0 104, 5.0 104, 5.0 103 and 5.0 101 s1, respectively. (e) Peak temperature and (f) enthalpy of the first exothermic event for the as-rolled specimens. Part of data for 5.0 103 s1 are taken from Ref. [32].
decrease of the enthalpy and peak temperature of the first exothermic event, it is evident that crystallization has been involved in the as-rolled specimen with e > 67% when e_ reaches 5.0 101 s1. In order to confirm such arguments, TEM observation of the asrolled specimens with typical e and e_ was performed. Fig. 2(a) and (b) shows the bright-field TEM image of the as-rolled specimen with e = 53% and 97% at e_ ¼ 1:0 104 s1 , respectively. No obvious contrast was observed except for the shear band. The insets in Fig. 2(a) and (b) are the corresponding SAED patterns, and each of them consists of two broad diffraction halos, indicating the amorphous nature of the structure. The HRTEM images of the areas containing both the shear band and the undeformed matrix are shown in Fig. 2(c) and (d), respectively. No lattice fringes are observed in both regions. They still maintain the amorphous state. Moreover, it is obvious that the shear band appears bright and the undeformed matrix is dark, which is likely a result of both thickness contrast and dilatation of the interatomic distance in shear bands. It is well known that shear bands have a more disor-
dered structure and higher energy than the undeformed matrix, as the chemical and topological short-range order is destroyed more or less by deformation. Shear-band material is more readily etched during chemical and/or electrochemical thinning as compared with the amorphous matrix, and then the shear band in the thin foil is thinner than the matrix. Meanwhile, it is also well accepted that there is some mechanical dilatation among the shear band atoms due to the higher free-volume content in shear bands as compared with the amorphous matrix, which also makes shear bands show a brighter contrast. Fig. 3(a) shows the bright-field TEM image of the as-rolled specimen with e = 93% at e_ ¼ 5:0 104 s1 . Similar to Fig. 2(a) and (b), no obvious contrast was observed except for the shear band. The corresponding SAED pattern is shown in the inset, and only diffraction halos are observed. The HRTEM image of the area containing both the shear band and the undeformed matrix is shown in Fig. 3(b). No lattice fringes are observed in both regions, indicating the amorphous state of the structure. As e exceeds 93%, some gray
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Fig. 2. (a) and (b) TEM image of the as-rolled specimen with e = 53% and 97% at e_ ¼ 1:0 104 s1 , respectively. The insets are the corresponding SAED patterns. (c) and (d) HRTEM images of the region containing the shear band and amorphous matrix in (a) and (b), respectively. The lines in (c) and (d) indicate the boundary between the shear band and the undeformed matrix.
regions near the shear bands begin to appear in the microstructure and their number and size increase with e. A TEM image of the specimen deformed by e = 97% is shown in Fig. 3(c), where the size of the gray regions are in the range of 80–100 nm. The inset in Fig. 3(c) is the corresponding SAED pattern, and it consists of only two broad diffraction halos, without any visible diffraction spots from crystallites. The HRTEM image of the area containing both the shear band and the undeformed matrix outside the gray regions is shown in Fig. 3(d). No lattice fringes are observed in both regions. A TEM image of the gray regions at higher magnification is shown in Fig. 3(e). The microstructure consists of a brighter matrix and darker substructures with an average size of about 5–10 nm. Fig. 3(f) shows a HRTEM image of the area containing the brighter matrix and darker substructure. No obvious lattice fringes are observed in both regions. Hence, it is clear that mechanically driven phase separation occurs in the as-rolled specimen at e_ ¼ 5:0 104 s1 when e exceeds 93%. The microstructure of the metallic glass deformed at e_ ¼ 5:0 103 s1 has been investigated before [32], where it has been revealed that phase separation occurs when e > 89%. Fig. 4(a) shows a bright-field TEM image of the as-rolled specimen with e = 97% at e_ ¼ 5:0 101 s1 . Some gray regions with an average size of 100– 200 nm precipitate in the specimen. The inset in Fig. 4(a) is the corresponding SAED pattern, consisting of only two broad diffraction halos. A TEM image of the gray regions at higher magnification is shown in Fig. 4(b). The microstructure also consists of a brighter matrix and darker substructures with an average size of about 5– 20 nm. Fig. 4(c) shows the HRTEM image of an area containing both brighter matrix and darker substructure, where no obvious lattice fringes are observed in both regions. However, some of the darker substructures are also seen to be composed of precipitated nanocrystals, as shown in Fig. 4(d). In combination with the DSC results,
it is concluded that mechanically driven phase separation and nanocrystallization have occurred in the as-rolled specimen at e_ ¼ 5:0 101 s1 when e exceeds 67%. 3.3. Stored energy evolution during rolling When a metallic glass is heated at a constant rate, the stored energy is released due to the increasing atomic mobility, which is observed as a structural relaxation exothermic peak in the DSC curve. Van den Beukel et al. has assumed that the change in enthalpy of structural relaxation is due to the change in free volume [5], which is probably appropriate for the as-cast metallic glasses rather than the deformed metallic glasses since inhomogeneous deformation may result in the development of residual stress. It is better to directly relate the relaxation enthalpy of the deformed metallic glass with the stored energy. In order to elucidate the stored energy evolution during rolling the Cu60Zr20Ti20 specimens and correlate it with the strain rate, the enthalpy released by structural relaxation was calculated from the isochronal DSC curves shown in Fig. 5(a–d). The relative change of the apparent specific heat DCp = Cp(T) Cp(323 K), where Cp(323 K) and Cp(T) are the apparent specific heat of the specimen at 323 K and temperature T, respectively, is calculated from the DSC curves. The results of the relaxation enthalpy per unit mass of specimen, Er, are shown in Fig. 5(e). It is found that in the initial stage of rolling, Er monotonously increases with e at all strain rates, and the variation rate of Er with e increases when e_ rises. For each e_ there is a critical e above which Er no longer increases and even decreases. From the change of Er with e at e_ ¼ 1:0 104 s1 , it is clear that the stored energy in the asrolled specimens gets to a plateau value when e exceeds 53%. Considering that neither phase separation nor crystallization oc-
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Fig. 3. Microstructure of specimens rolled at e_ ¼ 5:0 104 s1 . (a) TEM image with e = 93% at. (b) HRTEM image of the region containing the shear band and amorphous matrix in (a). (c) TEM image with e = 97%. (d) HRTEM image of the shear band and matrix apart from the gray regions in (c). (e) Magnified TEM image of the gray region in (c). (f) HRTEM images of the brighter matrix and darker substructure in (e). The insets in (a) and (c) are the corresponding SAED patterns.
curs in the deformation at such a low e_ , it can be undoubtedly concluded that a saturation of stored energy driven by plastic deformation is achieved in the monolithic metallic glass. As e_ rises to 5.0 104 s1, the critical e beyond which the stored energy is saturated increases to 77%, and the saturated stored energy also augments as compared with that at e_ ¼ 1:0 104 s1 . It should be pointed out that phase separation takes place at e_ ¼ 5:0 104 s1 at a higher e = 93%, rather than simultaneously with the saturation of the stored energy. But the occurrence of phase separation does not change the value of the saturated stored energy. If e_ further increases to 5.0 103 s1, the plateau of the stored energy is reached at e = 89%, and the amount of saturated stored energy also increases. At this strain rate, the saturation of the stored energy and phase separation concur at the same critical strain. At the highest e_ ¼ 5:0 101 s1 , although the stored energy can increase up to a very large value as deformation proceeds, an immediate decrease is observed in it, following the occurrence of phase separation plus nanocrystallization.
4. Discussion 4.1. Saturation of stored energy In the free-volume model [4], plastic deformation of metallic glasses occurs as a result of a number of individual atomic jumps in the so-called flow defects with a size larger than a critical size. When the deformation temperature remains unchanged, the generation rate of free volume is dependent on the free-volume content and strain rate, while the annihilation rate of free volume correlates with the free-volume content only [36]. At the beginning of the rolling deformation, the generation rate is generally higher than the annihilation rate, and the free-volume content increases with e. With the increase of the free-volume content, which prompts the annihilation of free volume, the net increase rate of free volume gradually decreases. Theoretical analyses have pointed out that whether the deformation is homogeneous or inhomogeneous, the free-volume content in a metallic glass
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Fig. 4. (a) TEM image of the as-rolled specimen with e = 97% at e_ ¼ 5:0 101 s1 . The inset is the corresponding SAED pattern. (b) Magnified TEM image of the gray region in (a). (c) HRTEM images including partial brighter matrix and darker substructure in (b). (d) HRTEM image of a darker substructure in (b). Note that clear lattice fringes have appeared.
increases as the plastic deformation proceeds, and saturates when the strain exceeds a certain value that depends on both the temperature and strain rate [37–40]. This phenomenon has been observed in the homogeneous tensile deformation of Pd40Ni40P20 glass ribbon at high temperature, where the ribbon remained amorphous after the free volume content got saturated [41]. When there is no phase transformation to take place upon deformation, whether the balance between generation and annihilation of free volume can be eventually reached in an inhomogeneous deformation of metallic glasses is still undiscovered experimentally. Huang et al. has pointed out that if no phase transformation occurs, the saturation of free-volume content is achieved in an inhomogeneous deformation only when shear bands extend to the entire volume of the specimen and the distribution of free volume becomes uniform [38]. That is to say, the saturation of free volume is impossible to be achieved in an actual inhomogeneous deformation provided that no phase transformation occurs. However, in the theoretical analysis of Huang et al., formation of nano-voids in shear bands through the coalescence of free volume has not been taken into account. Wright et al. found that excess free volume in shear bands results in excess free energy relative to amorphous matrix, which provides a driving force for nano-voids nucleation [42]. It indicates that free volume in shear bands is thermodynamically unstable during deformation, and whether or not nano-voids form in practice is a question of kinetics. Following the method developed by Miller and Gibson and later extended by Li et al. [43,44], the nanoscale structural defects in the shear band were investigated in the present work. The defocused images at the same locations as Fig. 2(c) and (d) at a defocus value of 200 nm were obtained. The Fourier transform of the defocused images was filtered by passing 0.5 < k < 1.5 nm1 and excluding all
other spatial frequencies in order to image the defects giving rise to the prominent difference in the Fourier transform amplitude between the shear band and the undeformed amorphous matrix in this small-angle scattering range. The reverse Fourier transform was then calculated to obtain filtered images, which show the projected atomic density, and the resulting images are shown in Fig. 6(a) and (b), respectively. To highlight the defects in the shear bands, the threshold-filtered and inverted images were obtained, as shown in Fig. 6(c) and (d), respectively. The small spots indicate the location of low-density defects, and they have been referred to as nano-voids [24,44]. The shear band contains a uniform distribution of nano-voids, while in the neighboring amorphous matrix there are almost no nano-voids existing. It becomes clear that the density of nano-voids in the shear band of Fig. 6(d) is higher than that in Fig. 6(c). On the other hand, as the rolling proceeds, the number of shear bands increases and the heterogeneity of microstructure enhances, resulting in the augmentation of residual stress to some extent. Our present work indicates that in the initial stage of rolling with e 6 53% at e_ ¼ 1:0 104 s1 , although there are a few nano-voids forming in the shear band as shown in Fig. 6(c), it cannot counteract the increased amount of free-volume content and residual stress, which results in a net increase of stored energy in this stage. As e exceeds 53%, the decrease of free-volume content in the shear bands, which results from the coalescence of free volume into many nano-voids, succeeds in compensating for the increase of free-volume content from the increasing density of shear band and the development of residual stress. Thus, saturation of the stored energy is observed in the monolithic Cu60Zr20Ti20 BMG, although it is still dubious whether or not the free-volume content in the monolithic metallic glass gets saturated during the rolling deformation.
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Fig. 5. (a)–(d) Specific heat data of the as-rolled Cu60Zr20Ti20 specimens with different e at various strain rates of 1.0 104, 5.0 104, 5.0 103 and 5.0 101 s1, respectively. (e) The released energy per unit mass of amorphous phase, Er, during structural relaxation of the rolled specimens. Part of data for 5.0 103 s1 is taken from Ref. [31].
4.2. Effect of strain rate on the stored energy and microstructure According to the free-volume model, the increase rate of free volume and the saturated free-volume content increase with e_ [4]. At a higher strain rate, more free volume is generated due to the higher generation rate of free volume as compared with the lower strain–rate deformation [36]. Although the increase in free-volume content also enhances the annihilation rate of free volume, it cannot compensate for the effect of the higher e_ on the increase of the generation rate of free volume. As a result, a higher e_ leads to the storage of more free volume as compared with the low e_ when e is the same. It can also be confirmed by a lot of earlier work, where multiple shear bands were observed at high strain rates but isolated shear bands operated at low strain rates [40,45]. The stored energy has a similar variation tendency, as shown in Fig. 5(e). At the same time, with the increase of e_ , the critical e above which the stored energy in the monolithic metallic
glass saturates shifts to higher values, which is consistent with that of free volume according to the free-volume model. However, at higher e_ phase transformation occurs at a lower critical e. The competition between these two critical strains determines whether the saturation of stored energy in the monolithic metallic glass can be obtained. When the Cu60Zr20Ti20 BMG is rolled at e_ ¼ 5:0 103 s1 , a plateau is observed on the curve of the stored energy as a function of e. However, it concurs with phase separation at the same critical e = 89% [32]. Only by means of the experimental results at this strain rate, it cannot be determined whether the saturation of the stored energy in the monolithic glass is reached at e = 89%, and the effect of phase separation on the stored energy also cannot be clearly understood. That is to say, phase separation would not affect the stored energy if the saturation of stored energy in the monolithic Cu60Zr20Ti20 glass were obtained at e = 89%, or phase separation would reduce the stored energy if the saturation had
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Fig. 6. (a) and (b) HRTEM images of the as-rolled specimen with e = 53% and 97% at e_ ¼ 1:0 104 s1 , respectively, at the same locations as Fig. 2 (c) and (d), but defocus 200 nm, and Fourier filtered. (c) and (d) Images at the same locations as (a) and (b), respectively, but threshold filtered and inverted.
not been obtained. This issue is settled down when the applied strain rate was decreased to 5.0 104 s1 in the experiment, where we have found that the saturation of stored energy takes place before the occurrence of phase separation, which provides a chance to investigate the effect of phase separation on the stored energy. As e exceeds 93%, a value considerably high than the critical strain for the saturation of stored energy, phase separation occurs, but the stored energy does not almost vary and keeps the saturation value up to the highest e. So it is concluded that phase separation does not change the stored energy, and it is very possible that at e_ ¼ 5:0 103 s1 the stored energy is already saturated when e = 89%. Phase separation is such a phenomenon that a homogeneous liquid or amorphous system develops into two or more parts whose compositions are different each other, while the liquid or amorphous structure is still maintained. A miscibility gap, which opens up between different thermodynamically favored compositions, promotes the occurrence of phase separation [46]. Using the subregular solution model, Abe et al. illustrated that among most of the reported ternary BMG system, the glass-forming range overlaps with the miscibility gap only for the Cu–Zr–Ti system, indicating that phase separation is possible for this ternary system [47]. Their further investigation shows that the temperature of miscibility gap for the Cu–Zr–Ti system is at about 600 K, while the glass transition temperature of the Cu–Zr–Ti alloy is about 710 K, indicating that phase separation during solidification and upon annealing is unlikely to happen if the glassy state is obtained in the solidified specimens. Although the temperature of miscibility gap for the Cu–Zr–Ti alloy is much higher than the rolling temperature, the enhanced atomic diffusion by shear deformation and the possible temperature rise within shear bands by local adiabatic heating may destabilize the glass and promote the phase separation and even nanocrystallization.
The fully amorphous Cu60Zr20Ti20 alloy has a tendency for phase separation. At temperature far below the glass transition temperature, however, phase separation cannot spontaneously occur since atomic rearrangement is quite difficult. When the amorphous alloy is subjected to cold rolling, the deformation exhibits an extreme inhomogeneity, i.e. the plastic strain is mainly localized in shear bands. The material near shear bands undergoes a severer deformation compared with the region far away from the shear bands, so the deformation-induced structural evolution should preferentially occur in it. As a result, the phase separation is observed only in localized regions near the shear bands, as shown in Figs. 3(c) and 4(a). In our previous work, the microstructure of the metallic glass deformed at e_ ¼ 5:0 103 s1 has been investigated in detail [32], where the phase-separation-induced brighter and darker substructures in the gray regions when e > 89% were Cu-poor and Cu-rich, respectively. These gray regions have been determined to possess the same average composition as that of the amorphous matrix, but they still show dark contrast compared to the amorphous matrix owing to the contribution of darker substructures. It should be emphasized that the microstructure of the as-rolled specimen is not only associated with e, but also strongly dependent on e_ . At the highest strain rate e = 5.0 101 s1, the DSC and TEM results clearly show evidences for phase separation and nanocrystallization as e exceeds 67%. The flow stress and especially the elastic limit of metallic glasses are known to be very high compared to conventional polycrystalline alloys. Deformation therefore involves significant work done on the specimens, leading to some temperature rise within shear bands [7,48]. At the same time, the amount of the temperature rise is closely related to the applied macroscopic strain rate [49]. It is worthy to note that the local strain rate within the shear band can exceed the macroscopic strain rate by orders of magnitude. Wright et al. found that the local strain rate within shear bands is about 103 s1 even under
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quasistatic compression with the macroscopic strain rate of 104 s1 considering that the shear band is very thin (a few tens of nanometers) [50]. The local strain rate within shear bands is estimated to be as high as 109 s1 under dynamic loading tests [49]. So, it is clear that as the applied strain rate of rolling increases from 1.0 104 to 5.0 101 s1, the local strain rate within shear bands also increases drastically, leading to a larger temperature rise. As the shear band moves fast (close to the sound speed) and is short-lived (order of microseconds or less) [51], it is difficult to measure the temperature rise directly. The calculation of Chen et al. show that the local temperature in shear bands rises up to about 2500 K during bending Al-based metallic glasses at RT, and then drops to RT in one nanosecond, dissipated by conduction [19]. Since there is no driving force for crystallization above the melting point, and the subsequent high cooling rate is also sufficient to bypass crystallization, no crystallization should occur in this case. But they indeed observed the Al-nanocrystal precipitation within shear bands of the Al-based glassy ribbons. They then attributed the nanocrystallization to the enhanced topological and chemical short-range order of the amorphous state by bending deformation. Meanwhile, Jiang et al. observed that nanocrystallization occurred only in the compressive region of amorphous Al90Fe5Gd5 bent at RT, but not in the tensile region [24]. This indicates that the temperature rise due to adiabatic heating is probably not the main cause of crystallization in shear bands since a similar temperature increase is expected to be in both the compressive and tensile regions. Interestingly, for the plastic deformation occurring in the nanoindentation of Zr52.5Cu17.9Ni14.6Al10Ti5 BMG, the temperature rise was only 0.05 K [20], but the deformed glass around the nanoindenter was crystallized, indicating that crystallization can be triggered merely by mechanical force. In our case, therefore, phase separation and nanocrystallization is likely due to the fact that atoms of the metallic glass are activated by mechanical stress, rather than local heating, since the heat can be dissipated over a nanosecond timescale even if it indeed exists. The rise of stored energy in the as-rolled metallic glasses puts the amorphous atoms to a higher energy state, as shown in Fig. 5(e), which enhances atomic mobility and increases the opportunity for the metallic glass to transform into a more stable state with lower free energy. The more the stored energy is, the more active the atomic mobility. As a result, no phase transformation takes place at e_ ¼ 1:0 104 s1 due to the low stored energy, and only phase separation occurs at e_ ¼ 5:0 104 and 5.0 103 s1, while phase separation plus nanocrystallization occurs at e_ ¼ 5:0 101 s1 due to the highest stored energy. Meanwhile, it is quite reasonable that the accumulative strain needed for phase transformation will be considerably decreased at high e_ due to the increased stored energy. Although the effect of the local temperature rise within shear bands on phase transformation is less significant than the enhancement of atomic mobility by rolling deformation, it may need to be taken into account in analyzing the stored energy evolution in the as-rolled specimen, especially at high strain rates. Lewandowski and Greer observed that the tin coating was remelted into spherical beads at shear bands during bending the tin-coated Zr41.2Ti13.8Cu12.5Ni10Be22.5 BMG [7], providing the evidence for local heating and melting of the coating. They further suggested that the local temperature rise can affect shear band evolution, especially at high strain rates where a larger temperature change is expected. It is interesting that the stored energy in the Cu60Zr20Ti20 BMG decreases with strain at the highest strain rate e_ ¼ 5:0 101 s1 when e exceeds 67%. It might partially be a result of temperature rise under these conditions, since the rise in temperature can reduce the residual stress and enhance the annihilation of free volume due to the increased atomic mobility. On the other hand, as the nanocrystals do not contain the so-called free volume, the
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stored energy in the as-rolled specimen will continuously decrease as nanocrystallization develops. At the same time, the appearance of the amorphous/crystal boundary is helpful for annihilating free volume, leading to the decrease of the stored energy. 5. Conclusions (1) By reducing e_ to 1.0 104 s1, the saturation of stored energy in the monolithic metallic glass without phase transformation has been observed in the Cu60Zr20Ti20 BMG as e exceeds 53% up to 97%. It is suggested that formation of nano-voids in shear bands decreases the free-volume content, and eventually leads to saturation of stored energy as a result of balance to the increase in shear band density and residual stress. (2) As e_ increases, the increase rate of stored energy in the initial deformation and the saturated amount at higher strains in the as-rolled Cu60Zr20Ti20 BMG increases, and meanwhile the critical e above which the stored energy in the monolithic metallic glass begins to saturate shifts to a higher value. However, the critical e for phase transformation tends to decrease due to the larger stored energy at higher e_ . (3) When the Cu60Zr20Ti20 BMG is rolled at strain rates of 5.0 104 s1 and 5.0 103 s1, phase separation occurs at high e. If rolling is performed at e_ ¼ 5:0 101 s1 , nanocrystallization takes place. The stored energy does not change with the occurrence of phase separation, but it decreases accompanied by nanocrystallization. It is proposed that both local temperature rise within shear bands and nanocrystallization account for the drop of stored energy at high e_ , because the elevated temperature can significantly reduce the residual stress and enhance the annihilation of free volume, and the amorphous/crystal boundary provides a easy way for free volume to annihilate.
Acknowledgements Financial supports from the National Natural Science Foundation of China under Grant No. 50771064, the Scientific Research Foundation for ROCS and SEM, and the DANIDA Fellowship Center are gratefully acknowledged. Center for analysis at Zhejiang SciTech University is gratefully acknowledged for the use of XRD and DSC. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18]
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Journal of Non-Crystalline Solids journal homepage: www.elsevier.com/locate/jnoncrysol
Microstructure and stored energy evolutions during rolling of Cu60Zr20Ti20 bulk metallic glass Q.P. Cao a,b, J.F. Li a,*, J.Z. Jiang b, Y.H. Zhou a a
State Key Laboratory of Metal Matrix Composites, School of Materials Science and Engineering, Shanghai Jiao Tong University, 800 Dongchuan Road, Shanghai 200240, PR China International Center for New-Structured Materials (ICNSM), Zhejiang University and Laboratory of New-Structured Materials, Department of Materials Science and Engineering, Zhejiang University, Hangzhou 310027, PR China b
a r t i c l e
i n f o
Article history: Received 2 April 2008 Received in revised form 9 September 2008 Available online 24 October 2008 PACS: 81.05.Kf 81.40.Ef 81.70.Pg 68.37.Lp Keywords: Amorphous metals, metallic glasses Crystallization
a b s t r a c t The microstructure and stored energy of Cu60Zr20Ti20 bulk metallic glass rolled at cryogenic temperature in a wide strain rate range 1.0 104 5.0 101 s1 have been investigated. As the specimen is rolled to be thinner, the stored energy first increases linearly, and then saturates above a critical thickness reduction at lower strain rates, or decreases at high strain rates. At the initial stage of rolling, no phase transformation except shear bands appears in the glass. Phase transformation occurs only when the specimen is severely deformed at strain rates higher than 1.0 104 s1. As strain rate increases, the critical strain for the stored energy to saturate increases, but the critical strain for phase separation to occur decreases, and meanwhile the type of the phase transformation changes from phase separation to nanocrystallization. The stored energy does not change with the occurrence of phase separation, but decreases due to nanocrystallization. It is proposed that coalescence of more free volume in shear bands into nanovoids should be principally responsible for the saturation of the stored energy, which balances the results from the increase in shear band number at higher strains. Ó 2008 Elsevier B.V. All rights reserved.
1. Introduction In comparison with the ideal metallic glass in which the atoms have the densest arrangement while the amorphous state is still held, an actual metallic glass with the same composition always contains an amount of excess free volume [1–3]. Spaepen proposed a free-volume model to account for the deformation behavior of metallic glasses at various temperatures [4], according to which, macroscopic flow of metallic glasses occurs as a result of biased jumps of atoms into their neighboring sites when the external stress is applied. A certain amount of free volume can be created as the atoms are squeezed into the neighboring sites with a smaller volume than the atoms. Van den Buekel and Sietsma interpreted enthalpy relaxation in terms of free volume, i.e. the change in the free-volume content during differential scanning calorimetry (DSC) experiment of a constant heating rate is proportional to the energy released by the structural relaxation of metallic glasses [5]. It is worth noting that the enthalpy relaxation is not fully determined by free volume and may reflect other relevant structural change as well, such as relaxation of residual stress. Casting and rapid cooling of metallic glasses can lead to residual stress
* Corresponding author. Tel./fax: +86 21 5474 8530. E-mail address: jfl[email protected] (J.F. Li). 0022-3093/$ - see front matter Ó 2008 Elsevier B.V. All rights reserved. doi:10.1016/j.jnoncrysol.2008.09.024
due to thermal tempering [6]. So can inhomogeneous deformation of metallic glasses [7]. During cold rolling of metallic glasses, very thin bands with intense shear deformation are commonly observed [8], and it definitely alters the amount and distribution of residual stresses. Therefore, for the cold-rolled metallic glasses it may be inexact to directly link the relaxation enthalpy (stored energy) determined by DSC with free volume. An earlier work on the strain energy storage in a cold-rolled Pd77.5Cu6Si16.5 glass ribbon has revealed that the stored energy monotonously increases as the thickness reduction increases to the highest value 36% [9]. However, traditional metallic glasses were fabricated by rapid solidification, and their thicknesses were seldom larger than 100 lm. The highest plastic strain realized in them could not be large, and experimental investigation of the stored energy evolution at larger plastic strains was severely confined. Recently this situation was changed as novel multicomponent bulk metallic glasses (BMGs) were discovered [10–15]. Their thicknesses more than 1 mm make it feasible to obtain large plastic strains by compression or rolling at temperatures far below the glass transition temperature. On the other hand, metallic glasses are metastable in thermodynamics. When they are undergoing plastic deformation, the atomic displacement is significantly activated by the shear stress. While part of the atoms in flow becomes more disordered as a result of the mechanical shear [8,16–18], the other atoms may get a chance
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to rearrange into another structure with a lower free energy, such as crystalline phases or new glassy phases. Mechanically induced crystallization has been observed in some metallic glasses [19,20]. The variation of the microstructure of metallic glasses with plastic deformation has been an interesting subject of materials science [21–30]. In our previous work [31,32], the Cu60Zr20Ti20 BMG was ever rolled at a fixed strain rate at room temperature (RT) as well as cryogenic temperature (CT), and it was found that the increase of rolling temperature redounded to precipitation of nanocrystals. Considering that the structural variation in metallic glasses depends on not only rolling temperature and plastic strain but also the applied strain rate, in the present work we are going to carry out a series of deformation experiments at CT on the Cu60Zr20Ti20 BMG in a particularly wide range of strain rate up to high macroscopic strains, so as to reveal the effect of strain rate and total strain on the microstructure and stored energy of the deformed metallic glass. 2. Experimental Ternary Cu60Zr20Ti20 master alloy ingots were prepared by arc melting the mixture of high purity Cu (99.99%), Zr (99.9%) and Ti (99.9%) under a Ti-gettered argon atmosphere. The ingots were inverted and remelted six times to ensure compositional homogeneity and then suck-cast into a water-cooled copper mold to produce 40 mm long cylindrical rods with a diameter of 2 mm. The amorphous structure was identified by X-ray diffraction (XRD) with monochromatic Cu Ka radiation and high resolution transmission electron microscopy (HRTEM). The rods were cut into short cylinders with a thickness of 1.5 mm for rolling. Both ends of the cylinders were mechanically polished to make them parallel to each other prior to the rolling experiment. Before the rolling deformation, a continuous liquid nitrogen stream was used to cool the specimen, by which the temperature of the specimen was measured by a K-type thermocouple to be about 150 K. During the subsequent rolling process, the specimen was incessantly cooled with the liquid nitrogen stream. Details of the deformation procedure were described elsewhere [32]. The degree of deformation was denoted by the reduction in thickness e = (h0 h)/h0, where h0 and h represented the specimen thicknesses before and after rolling, respectively. Many small deformation passes were used with progressively narrowing the gap between two rollers, and the decrease of the gap during deformation was carefully controlled. The strain rate e_ adopted in the experiment was about 1.0 104, 5.0 104, 5.0 103 and 5.0 101 s1, respectively. At a higher strain rate (e.g., e_ ¼ 1:0 s1 ), the specimen fractured abruptly at about 45to the direction of rolling pressure in the initial rolling stage. The microstructures of the as-rolled specimens were examined by JEOL JEM-3010 HRTEM apparatus operating at 300 kV. The difficulty in preparing a good TEM specimen of the Cu–Zr–Ti alloy is well known just as the previous studies on this alloy system have pointed out that a nanocrystalline microstructure might form when the specimen was prepared without special attention [33]. The thin foils for TEM were prepared by twin-jet electropolishing using a solution of 6% nitric acid in methanol at a temperature of 238 K. The specimens were observed in TEM immediately after the preparation since the Cu–Zr–Ti thin foil readily oxidizes upon exposure in air atmosphere. Selected area electron diffraction (SAED) patterns were taken from an area about 1 lm in diameter. Thermal analyses of the specimens subjected to different degrees of deformation were performed in a Perkin–Elmer Pyris Diamond DSC under a flow of purified argon. The specimens were heated at a heating rate of 20 K/min to 853 K, cooled to RT and then
reheated at the same rate to establish a baseline. The baseline was subtracted from the first scan to better describe the thermal behavior of the glass. Five samples at each strain were measured. The sample and reference pans were made of aluminium. The temperature and the heat flow were calibrated by measuring the melting temperatures and the heats of fusion of pure In, Sn and Zn. 3. Results 3.1. DSC curves The as-cast Cu60Zr20Ti20 rods have been verified by X-ray diffractometry (XRD) and HRTEM to be fully amorphous (the results are not shown). In the strain rate range from 1.0 104 to 5.0 101 s1, the Cu60Zr20Ti20 BMG can be rolled at CT up to 97% reduction in thickness without fracture. Fig. 1(a)–(d) shows the typical DSC curves of the specimens rolled at various strain rates. The data of the DSC curves are the average of five measured DSC curves at each strain. The as-cast alloy exhibits a glass transition with the onset temperature Tg = 706 K, followed by two exothermic events, characterized by the onset temperature for crystallization Tx = 736 K, the peak temperature for the first crystallization event Tp1 = 748 K and the peak temperature for the second crystallization event Tp2 = 786 K. It has been known that the first exothermic event corresponds to the amorphous-to-Cu51Zr14 phase transition while the second one is due to the crystallization of the residual amorphous phase [34,35]. The peak temperature and enthalpy of the second exothermic event do not change with the deformation no matter how high the e_ is. Here, we only present the measured strain dependences of the peak temperature and enthalpy of the first exothermic event at various strain rates in Fig. 1(e) and (f), respectively. The error bars of the peak temperature and enthalpy of each strain are standard deviations of five measured samples. At the lowest strain rate e_ ¼ 1:0 104 s1 , the first peak temperature keeps unchanged up to the highest e of 97%. As e_ increases from 1.0 104 to 5.0 104 s1, the first peak temperature still remains constant at e 6 93%, but decreases by about 1.0 K when e increases up to 97%. When e_ further increases from 5.0 104 to 5.0 103 s1, the starting e for peak temperature to decline varies from 93% to 89%, and the decrease amount at e = 97% enlarges from 1.0 to 1.5 K. However, the enthalpy of the first exothermic event remains unchanged. It is clear that the rolling at a higher e_ leads to a heavier decrease of the glass in thermal stability when e exceeds the critical value. In the experiment we raised the strain rate abruptly to 5.0 101 s1, the critical e for the first peak temperature to decline drops to 67%, and the decrease of the first peak temperature reaches as large as 4.7 K at e = 97%. An important finding is that the enthalpy of the first exothermic event at this strain rate declines with increasing e beyond the critical strain 67%, and a change of about 4.7% is reached at e = 97%. 3.2. Microstructure evolution with rolling It has been observed that when the Cu60Zr20Ti20 BMG is rolled at CT and e_ ¼ 5:0 103 s1 , mechanically driven phase separation occurs with e > 89% [32]. The occurrence of phase separation does not change the enthalpy of the first exothermic event, but decreases the peak temperature. It is therefore speculated that at the lowest e_ ¼ 1:0 104 s1 , no phase transformation occurs in the Cu60Zr20Ti20 specimen during the whole rolling process since both the peak temperature and enthalpy of the first exothermic event are unchanged. For e_ ¼ 5:0 104 and 5.0 103 s1, the deformation should be accompanied with phase separation when e exceeds 93% and 89%, respectively. In view of the obvious
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Fig. 1. (a)–(d) DSC curves of the as-rolled Cu60Zr20Ti20 specimens with different e at various strain rates of 1.0 104, 5.0 104, 5.0 103 and 5.0 101 s1, respectively. (e) Peak temperature and (f) enthalpy of the first exothermic event for the as-rolled specimens. Part of data for 5.0 103 s1 are taken from Ref. [32].
decrease of the enthalpy and peak temperature of the first exothermic event, it is evident that crystallization has been involved in the as-rolled specimen with e > 67% when e_ reaches 5.0 101 s1. In order to confirm such arguments, TEM observation of the asrolled specimens with typical e and e_ was performed. Fig. 2(a) and (b) shows the bright-field TEM image of the as-rolled specimen with e = 53% and 97% at e_ ¼ 1:0 104 s1 , respectively. No obvious contrast was observed except for the shear band. The insets in Fig. 2(a) and (b) are the corresponding SAED patterns, and each of them consists of two broad diffraction halos, indicating the amorphous nature of the structure. The HRTEM images of the areas containing both the shear band and the undeformed matrix are shown in Fig. 2(c) and (d), respectively. No lattice fringes are observed in both regions. They still maintain the amorphous state. Moreover, it is obvious that the shear band appears bright and the undeformed matrix is dark, which is likely a result of both thickness contrast and dilatation of the interatomic distance in shear bands. It is well known that shear bands have a more disor-
dered structure and higher energy than the undeformed matrix, as the chemical and topological short-range order is destroyed more or less by deformation. Shear-band material is more readily etched during chemical and/or electrochemical thinning as compared with the amorphous matrix, and then the shear band in the thin foil is thinner than the matrix. Meanwhile, it is also well accepted that there is some mechanical dilatation among the shear band atoms due to the higher free-volume content in shear bands as compared with the amorphous matrix, which also makes shear bands show a brighter contrast. Fig. 3(a) shows the bright-field TEM image of the as-rolled specimen with e = 93% at e_ ¼ 5:0 104 s1 . Similar to Fig. 2(a) and (b), no obvious contrast was observed except for the shear band. The corresponding SAED pattern is shown in the inset, and only diffraction halos are observed. The HRTEM image of the area containing both the shear band and the undeformed matrix is shown in Fig. 3(b). No lattice fringes are observed in both regions, indicating the amorphous state of the structure. As e exceeds 93%, some gray
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Fig. 2. (a) and (b) TEM image of the as-rolled specimen with e = 53% and 97% at e_ ¼ 1:0 104 s1 , respectively. The insets are the corresponding SAED patterns. (c) and (d) HRTEM images of the region containing the shear band and amorphous matrix in (a) and (b), respectively. The lines in (c) and (d) indicate the boundary between the shear band and the undeformed matrix.
regions near the shear bands begin to appear in the microstructure and their number and size increase with e. A TEM image of the specimen deformed by e = 97% is shown in Fig. 3(c), where the size of the gray regions are in the range of 80–100 nm. The inset in Fig. 3(c) is the corresponding SAED pattern, and it consists of only two broad diffraction halos, without any visible diffraction spots from crystallites. The HRTEM image of the area containing both the shear band and the undeformed matrix outside the gray regions is shown in Fig. 3(d). No lattice fringes are observed in both regions. A TEM image of the gray regions at higher magnification is shown in Fig. 3(e). The microstructure consists of a brighter matrix and darker substructures with an average size of about 5–10 nm. Fig. 3(f) shows a HRTEM image of the area containing the brighter matrix and darker substructure. No obvious lattice fringes are observed in both regions. Hence, it is clear that mechanically driven phase separation occurs in the as-rolled specimen at e_ ¼ 5:0 104 s1 when e exceeds 93%. The microstructure of the metallic glass deformed at e_ ¼ 5:0 103 s1 has been investigated before [32], where it has been revealed that phase separation occurs when e > 89%. Fig. 4(a) shows a bright-field TEM image of the as-rolled specimen with e = 97% at e_ ¼ 5:0 101 s1 . Some gray regions with an average size of 100– 200 nm precipitate in the specimen. The inset in Fig. 4(a) is the corresponding SAED pattern, consisting of only two broad diffraction halos. A TEM image of the gray regions at higher magnification is shown in Fig. 4(b). The microstructure also consists of a brighter matrix and darker substructures with an average size of about 5– 20 nm. Fig. 4(c) shows the HRTEM image of an area containing both brighter matrix and darker substructure, where no obvious lattice fringes are observed in both regions. However, some of the darker substructures are also seen to be composed of precipitated nanocrystals, as shown in Fig. 4(d). In combination with the DSC results,
it is concluded that mechanically driven phase separation and nanocrystallization have occurred in the as-rolled specimen at e_ ¼ 5:0 101 s1 when e exceeds 67%. 3.3. Stored energy evolution during rolling When a metallic glass is heated at a constant rate, the stored energy is released due to the increasing atomic mobility, which is observed as a structural relaxation exothermic peak in the DSC curve. Van den Beukel et al. has assumed that the change in enthalpy of structural relaxation is due to the change in free volume [5], which is probably appropriate for the as-cast metallic glasses rather than the deformed metallic glasses since inhomogeneous deformation may result in the development of residual stress. It is better to directly relate the relaxation enthalpy of the deformed metallic glass with the stored energy. In order to elucidate the stored energy evolution during rolling the Cu60Zr20Ti20 specimens and correlate it with the strain rate, the enthalpy released by structural relaxation was calculated from the isochronal DSC curves shown in Fig. 5(a–d). The relative change of the apparent specific heat DCp = Cp(T) Cp(323 K), where Cp(323 K) and Cp(T) are the apparent specific heat of the specimen at 323 K and temperature T, respectively, is calculated from the DSC curves. The results of the relaxation enthalpy per unit mass of specimen, Er, are shown in Fig. 5(e). It is found that in the initial stage of rolling, Er monotonously increases with e at all strain rates, and the variation rate of Er with e increases when e_ rises. For each e_ there is a critical e above which Er no longer increases and even decreases. From the change of Er with e at e_ ¼ 1:0 104 s1 , it is clear that the stored energy in the asrolled specimens gets to a plateau value when e exceeds 53%. Considering that neither phase separation nor crystallization oc-
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Fig. 3. Microstructure of specimens rolled at e_ ¼ 5:0 104 s1 . (a) TEM image with e = 93% at. (b) HRTEM image of the region containing the shear band and amorphous matrix in (a). (c) TEM image with e = 97%. (d) HRTEM image of the shear band and matrix apart from the gray regions in (c). (e) Magnified TEM image of the gray region in (c). (f) HRTEM images of the brighter matrix and darker substructure in (e). The insets in (a) and (c) are the corresponding SAED patterns.
curs in the deformation at such a low e_ , it can be undoubtedly concluded that a saturation of stored energy driven by plastic deformation is achieved in the monolithic metallic glass. As e_ rises to 5.0 104 s1, the critical e beyond which the stored energy is saturated increases to 77%, and the saturated stored energy also augments as compared with that at e_ ¼ 1:0 104 s1 . It should be pointed out that phase separation takes place at e_ ¼ 5:0 104 s1 at a higher e = 93%, rather than simultaneously with the saturation of the stored energy. But the occurrence of phase separation does not change the value of the saturated stored energy. If e_ further increases to 5.0 103 s1, the plateau of the stored energy is reached at e = 89%, and the amount of saturated stored energy also increases. At this strain rate, the saturation of the stored energy and phase separation concur at the same critical strain. At the highest e_ ¼ 5:0 101 s1 , although the stored energy can increase up to a very large value as deformation proceeds, an immediate decrease is observed in it, following the occurrence of phase separation plus nanocrystallization.
4. Discussion 4.1. Saturation of stored energy In the free-volume model [4], plastic deformation of metallic glasses occurs as a result of a number of individual atomic jumps in the so-called flow defects with a size larger than a critical size. When the deformation temperature remains unchanged, the generation rate of free volume is dependent on the free-volume content and strain rate, while the annihilation rate of free volume correlates with the free-volume content only [36]. At the beginning of the rolling deformation, the generation rate is generally higher than the annihilation rate, and the free-volume content increases with e. With the increase of the free-volume content, which prompts the annihilation of free volume, the net increase rate of free volume gradually decreases. Theoretical analyses have pointed out that whether the deformation is homogeneous or inhomogeneous, the free-volume content in a metallic glass
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Fig. 4. (a) TEM image of the as-rolled specimen with e = 97% at e_ ¼ 5:0 101 s1 . The inset is the corresponding SAED pattern. (b) Magnified TEM image of the gray region in (a). (c) HRTEM images including partial brighter matrix and darker substructure in (b). (d) HRTEM image of a darker substructure in (b). Note that clear lattice fringes have appeared.
increases as the plastic deformation proceeds, and saturates when the strain exceeds a certain value that depends on both the temperature and strain rate [37–40]. This phenomenon has been observed in the homogeneous tensile deformation of Pd40Ni40P20 glass ribbon at high temperature, where the ribbon remained amorphous after the free volume content got saturated [41]. When there is no phase transformation to take place upon deformation, whether the balance between generation and annihilation of free volume can be eventually reached in an inhomogeneous deformation of metallic glasses is still undiscovered experimentally. Huang et al. has pointed out that if no phase transformation occurs, the saturation of free-volume content is achieved in an inhomogeneous deformation only when shear bands extend to the entire volume of the specimen and the distribution of free volume becomes uniform [38]. That is to say, the saturation of free volume is impossible to be achieved in an actual inhomogeneous deformation provided that no phase transformation occurs. However, in the theoretical analysis of Huang et al., formation of nano-voids in shear bands through the coalescence of free volume has not been taken into account. Wright et al. found that excess free volume in shear bands results in excess free energy relative to amorphous matrix, which provides a driving force for nano-voids nucleation [42]. It indicates that free volume in shear bands is thermodynamically unstable during deformation, and whether or not nano-voids form in practice is a question of kinetics. Following the method developed by Miller and Gibson and later extended by Li et al. [43,44], the nanoscale structural defects in the shear band were investigated in the present work. The defocused images at the same locations as Fig. 2(c) and (d) at a defocus value of 200 nm were obtained. The Fourier transform of the defocused images was filtered by passing 0.5 < k < 1.5 nm1 and excluding all
other spatial frequencies in order to image the defects giving rise to the prominent difference in the Fourier transform amplitude between the shear band and the undeformed amorphous matrix in this small-angle scattering range. The reverse Fourier transform was then calculated to obtain filtered images, which show the projected atomic density, and the resulting images are shown in Fig. 6(a) and (b), respectively. To highlight the defects in the shear bands, the threshold-filtered and inverted images were obtained, as shown in Fig. 6(c) and (d), respectively. The small spots indicate the location of low-density defects, and they have been referred to as nano-voids [24,44]. The shear band contains a uniform distribution of nano-voids, while in the neighboring amorphous matrix there are almost no nano-voids existing. It becomes clear that the density of nano-voids in the shear band of Fig. 6(d) is higher than that in Fig. 6(c). On the other hand, as the rolling proceeds, the number of shear bands increases and the heterogeneity of microstructure enhances, resulting in the augmentation of residual stress to some extent. Our present work indicates that in the initial stage of rolling with e 6 53% at e_ ¼ 1:0 104 s1 , although there are a few nano-voids forming in the shear band as shown in Fig. 6(c), it cannot counteract the increased amount of free-volume content and residual stress, which results in a net increase of stored energy in this stage. As e exceeds 53%, the decrease of free-volume content in the shear bands, which results from the coalescence of free volume into many nano-voids, succeeds in compensating for the increase of free-volume content from the increasing density of shear band and the development of residual stress. Thus, saturation of the stored energy is observed in the monolithic Cu60Zr20Ti20 BMG, although it is still dubious whether or not the free-volume content in the monolithic metallic glass gets saturated during the rolling deformation.
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Fig. 5. (a)–(d) Specific heat data of the as-rolled Cu60Zr20Ti20 specimens with different e at various strain rates of 1.0 104, 5.0 104, 5.0 103 and 5.0 101 s1, respectively. (e) The released energy per unit mass of amorphous phase, Er, during structural relaxation of the rolled specimens. Part of data for 5.0 103 s1 is taken from Ref. [31].
4.2. Effect of strain rate on the stored energy and microstructure According to the free-volume model, the increase rate of free volume and the saturated free-volume content increase with e_ [4]. At a higher strain rate, more free volume is generated due to the higher generation rate of free volume as compared with the lower strain–rate deformation [36]. Although the increase in free-volume content also enhances the annihilation rate of free volume, it cannot compensate for the effect of the higher e_ on the increase of the generation rate of free volume. As a result, a higher e_ leads to the storage of more free volume as compared with the low e_ when e is the same. It can also be confirmed by a lot of earlier work, where multiple shear bands were observed at high strain rates but isolated shear bands operated at low strain rates [40,45]. The stored energy has a similar variation tendency, as shown in Fig. 5(e). At the same time, with the increase of e_ , the critical e above which the stored energy in the monolithic metallic
glass saturates shifts to higher values, which is consistent with that of free volume according to the free-volume model. However, at higher e_ phase transformation occurs at a lower critical e. The competition between these two critical strains determines whether the saturation of stored energy in the monolithic metallic glass can be obtained. When the Cu60Zr20Ti20 BMG is rolled at e_ ¼ 5:0 103 s1 , a plateau is observed on the curve of the stored energy as a function of e. However, it concurs with phase separation at the same critical e = 89% [32]. Only by means of the experimental results at this strain rate, it cannot be determined whether the saturation of the stored energy in the monolithic glass is reached at e = 89%, and the effect of phase separation on the stored energy also cannot be clearly understood. That is to say, phase separation would not affect the stored energy if the saturation of stored energy in the monolithic Cu60Zr20Ti20 glass were obtained at e = 89%, or phase separation would reduce the stored energy if the saturation had
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Fig. 6. (a) and (b) HRTEM images of the as-rolled specimen with e = 53% and 97% at e_ ¼ 1:0 104 s1 , respectively, at the same locations as Fig. 2 (c) and (d), but defocus 200 nm, and Fourier filtered. (c) and (d) Images at the same locations as (a) and (b), respectively, but threshold filtered and inverted.
not been obtained. This issue is settled down when the applied strain rate was decreased to 5.0 104 s1 in the experiment, where we have found that the saturation of stored energy takes place before the occurrence of phase separation, which provides a chance to investigate the effect of phase separation on the stored energy. As e exceeds 93%, a value considerably high than the critical strain for the saturation of stored energy, phase separation occurs, but the stored energy does not almost vary and keeps the saturation value up to the highest e. So it is concluded that phase separation does not change the stored energy, and it is very possible that at e_ ¼ 5:0 103 s1 the stored energy is already saturated when e = 89%. Phase separation is such a phenomenon that a homogeneous liquid or amorphous system develops into two or more parts whose compositions are different each other, while the liquid or amorphous structure is still maintained. A miscibility gap, which opens up between different thermodynamically favored compositions, promotes the occurrence of phase separation [46]. Using the subregular solution model, Abe et al. illustrated that among most of the reported ternary BMG system, the glass-forming range overlaps with the miscibility gap only for the Cu–Zr–Ti system, indicating that phase separation is possible for this ternary system [47]. Their further investigation shows that the temperature of miscibility gap for the Cu–Zr–Ti system is at about 600 K, while the glass transition temperature of the Cu–Zr–Ti alloy is about 710 K, indicating that phase separation during solidification and upon annealing is unlikely to happen if the glassy state is obtained in the solidified specimens. Although the temperature of miscibility gap for the Cu–Zr–Ti alloy is much higher than the rolling temperature, the enhanced atomic diffusion by shear deformation and the possible temperature rise within shear bands by local adiabatic heating may destabilize the glass and promote the phase separation and even nanocrystallization.
The fully amorphous Cu60Zr20Ti20 alloy has a tendency for phase separation. At temperature far below the glass transition temperature, however, phase separation cannot spontaneously occur since atomic rearrangement is quite difficult. When the amorphous alloy is subjected to cold rolling, the deformation exhibits an extreme inhomogeneity, i.e. the plastic strain is mainly localized in shear bands. The material near shear bands undergoes a severer deformation compared with the region far away from the shear bands, so the deformation-induced structural evolution should preferentially occur in it. As a result, the phase separation is observed only in localized regions near the shear bands, as shown in Figs. 3(c) and 4(a). In our previous work, the microstructure of the metallic glass deformed at e_ ¼ 5:0 103 s1 has been investigated in detail [32], where the phase-separation-induced brighter and darker substructures in the gray regions when e > 89% were Cu-poor and Cu-rich, respectively. These gray regions have been determined to possess the same average composition as that of the amorphous matrix, but they still show dark contrast compared to the amorphous matrix owing to the contribution of darker substructures. It should be emphasized that the microstructure of the as-rolled specimen is not only associated with e, but also strongly dependent on e_ . At the highest strain rate e = 5.0 101 s1, the DSC and TEM results clearly show evidences for phase separation and nanocrystallization as e exceeds 67%. The flow stress and especially the elastic limit of metallic glasses are known to be very high compared to conventional polycrystalline alloys. Deformation therefore involves significant work done on the specimens, leading to some temperature rise within shear bands [7,48]. At the same time, the amount of the temperature rise is closely related to the applied macroscopic strain rate [49]. It is worthy to note that the local strain rate within the shear band can exceed the macroscopic strain rate by orders of magnitude. Wright et al. found that the local strain rate within shear bands is about 103 s1 even under
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quasistatic compression with the macroscopic strain rate of 104 s1 considering that the shear band is very thin (a few tens of nanometers) [50]. The local strain rate within shear bands is estimated to be as high as 109 s1 under dynamic loading tests [49]. So, it is clear that as the applied strain rate of rolling increases from 1.0 104 to 5.0 101 s1, the local strain rate within shear bands also increases drastically, leading to a larger temperature rise. As the shear band moves fast (close to the sound speed) and is short-lived (order of microseconds or less) [51], it is difficult to measure the temperature rise directly. The calculation of Chen et al. show that the local temperature in shear bands rises up to about 2500 K during bending Al-based metallic glasses at RT, and then drops to RT in one nanosecond, dissipated by conduction [19]. Since there is no driving force for crystallization above the melting point, and the subsequent high cooling rate is also sufficient to bypass crystallization, no crystallization should occur in this case. But they indeed observed the Al-nanocrystal precipitation within shear bands of the Al-based glassy ribbons. They then attributed the nanocrystallization to the enhanced topological and chemical short-range order of the amorphous state by bending deformation. Meanwhile, Jiang et al. observed that nanocrystallization occurred only in the compressive region of amorphous Al90Fe5Gd5 bent at RT, but not in the tensile region [24]. This indicates that the temperature rise due to adiabatic heating is probably not the main cause of crystallization in shear bands since a similar temperature increase is expected to be in both the compressive and tensile regions. Interestingly, for the plastic deformation occurring in the nanoindentation of Zr52.5Cu17.9Ni14.6Al10Ti5 BMG, the temperature rise was only 0.05 K [20], but the deformed glass around the nanoindenter was crystallized, indicating that crystallization can be triggered merely by mechanical force. In our case, therefore, phase separation and nanocrystallization is likely due to the fact that atoms of the metallic glass are activated by mechanical stress, rather than local heating, since the heat can be dissipated over a nanosecond timescale even if it indeed exists. The rise of stored energy in the as-rolled metallic glasses puts the amorphous atoms to a higher energy state, as shown in Fig. 5(e), which enhances atomic mobility and increases the opportunity for the metallic glass to transform into a more stable state with lower free energy. The more the stored energy is, the more active the atomic mobility. As a result, no phase transformation takes place at e_ ¼ 1:0 104 s1 due to the low stored energy, and only phase separation occurs at e_ ¼ 5:0 104 and 5.0 103 s1, while phase separation plus nanocrystallization occurs at e_ ¼ 5:0 101 s1 due to the highest stored energy. Meanwhile, it is quite reasonable that the accumulative strain needed for phase transformation will be considerably decreased at high e_ due to the increased stored energy. Although the effect of the local temperature rise within shear bands on phase transformation is less significant than the enhancement of atomic mobility by rolling deformation, it may need to be taken into account in analyzing the stored energy evolution in the as-rolled specimen, especially at high strain rates. Lewandowski and Greer observed that the tin coating was remelted into spherical beads at shear bands during bending the tin-coated Zr41.2Ti13.8Cu12.5Ni10Be22.5 BMG [7], providing the evidence for local heating and melting of the coating. They further suggested that the local temperature rise can affect shear band evolution, especially at high strain rates where a larger temperature change is expected. It is interesting that the stored energy in the Cu60Zr20Ti20 BMG decreases with strain at the highest strain rate e_ ¼ 5:0 101 s1 when e exceeds 67%. It might partially be a result of temperature rise under these conditions, since the rise in temperature can reduce the residual stress and enhance the annihilation of free volume due to the increased atomic mobility. On the other hand, as the nanocrystals do not contain the so-called free volume, the
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stored energy in the as-rolled specimen will continuously decrease as nanocrystallization develops. At the same time, the appearance of the amorphous/crystal boundary is helpful for annihilating free volume, leading to the decrease of the stored energy. 5. Conclusions (1) By reducing e_ to 1.0 104 s1, the saturation of stored energy in the monolithic metallic glass without phase transformation has been observed in the Cu60Zr20Ti20 BMG as e exceeds 53% up to 97%. It is suggested that formation of nano-voids in shear bands decreases the free-volume content, and eventually leads to saturation of stored energy as a result of balance to the increase in shear band density and residual stress. (2) As e_ increases, the increase rate of stored energy in the initial deformation and the saturated amount at higher strains in the as-rolled Cu60Zr20Ti20 BMG increases, and meanwhile the critical e above which the stored energy in the monolithic metallic glass begins to saturate shifts to a higher value. However, the critical e for phase transformation tends to decrease due to the larger stored energy at higher e_ . (3) When the Cu60Zr20Ti20 BMG is rolled at strain rates of 5.0 104 s1 and 5.0 103 s1, phase separation occurs at high e. If rolling is performed at e_ ¼ 5:0 101 s1 , nanocrystallization takes place. The stored energy does not change with the occurrence of phase separation, but it decreases accompanied by nanocrystallization. It is proposed that both local temperature rise within shear bands and nanocrystallization account for the drop of stored energy at high e_ , because the elevated temperature can significantly reduce the residual stress and enhance the annihilation of free volume, and the amorphous/crystal boundary provides a easy way for free volume to annihilate.
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