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Cryogenic treatment induced hardening of Cu45Zr45Ag7Al3 bulk metallic glass
Physica B 433 (2014) 84–88
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Physica B journal homepage: www.elsevier.com/locate/physb
Cryogenic treatment induced hardening of Cu45Zr45Ag7Al3 bulk metallic glass L.K. Zhang, Z.H. Chen, D. Chen n, Q. Zheng College of Materials Science and Engineering, Hunan University, Changsha 410082, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 27 August 2013 Received in revised form 24 September 2013 Accepted 1 October 2013 Available online 18 October 2013
Little is known about the mechanical properties of bulk metallic glassy alloys (BMGs) after cryogenic treatment (CT). Cu45Zr45Ag7Al3 BMGs were treated in cryogenic temperatures for four chosen holding times. Compared with the as-cast sample, the compressive fracture strength of this BMG increased with increasing CT time. Furthermore, the CT changes the fracture mode from the ductile to the brittle. And CT brought about the precipitation of the AlCu2Zr and Cu5Zr phases which were accompanying with the morphology modification and contributed to the mechanical properties improvement of this BMG. & 2013 Elsevier B.V. All rights reserved.
Keywords: Bulk metallic glass Cryogenic treatment Mechanical properties Harding Fracture
1. Introduction
2. Experiments
Bulk metallic glasses (BMGs), which exhibit short-range structural periodicity, possess many unique physical, mechanical and chemical properties [1–5]. Among those most promising bulk metallic glasses, Cu–Zr–Ag–Al alloy system has exhibited very special not only for its high GFA, but also for its unique mechanical properties, such as extremely high strength and good toughness [6,7]. Thus, the Cu45 Zr45Ag7Al3 was chosen for the current investigation. Although many current studies are directed toward researching the microstructure of BMGs or bulk metallic glass matrix composites (BMGCs) during the annealing process [8,9], little is known about microstructure and mechanical properties of BMGs and BMGCs after cryogenic temperatures. As we know, CT was widely used for high precision parts, and has been shown to result in significant increase in mechanical properties of steels, which was caused by the reduction in the amount of retained austenite and partly transforming of the structure induced by the internal stress of CT [1,10,11]. Recently, we have reported that a Cu46Zr46Al8 matrix BMGs exhibit transformation of microstructure and mechanical behaviors after CT [12]. The objective of the present work was to clarify the microstructure of the Cu45Zr45Ag7Al3 bulk specimens after CT. Furthermore, the effects of the CT on the micro-hardness and the compression fracture strength of the samples were also investigated.
The material used in the present study has a composition of Cu45Zr45Ag7Al3 (atomic percent). Alloy ingots were prepared by arc melting the mixture of metals (4 99.9 mass%) in an argon atmosphere, followed by suction casting into a copper mold. From these master alloys, cast cylinders with 4 mm diameter and 50 mm length were obtained by suction-casting into a copper mold under a purified argon atmosphere. The as-cast samples were cooled to 80 K with a rate of 20 K/h in a controlled box and kept at this temperature for different periods of time (96, 192 and 240 h). The phase constitute was examined by X-ray diffractometry (XRD, D/Max-2550) using a monochromatic Cu Kα radiation. The microstructure was investigated by using a Tecnai F220 highresolution transmission electron microscopy (HRTEM), operating at 300 kV. The mechanical testing was performed by compression at room temperature. The strain rate was 5 10 4s 1. The test samples were 4 mm in diameter and 8 mm in length. Fracture surface morphology was examined by JSM-6700F scanning electron microscopy (SEM).
3. Results 3.1. Microstructure
n
Corresponding author. Tel./fax: þ 86 731 88821648. E-mail address: [email protected] (D. Chen).
0921-4526/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.physb.2013.10.003
Fig. 1 shows the XRD patterns of the Cu45Zr45Ag7Al3 amorphous alloys treated under different conditions. The as-cast sample
L.K. Zhang et al. / Physica B 433 (2014) 84–88
exhibits a broad diffraction peak, which is the characteristic of an amorphous structure. After 192 h CT, several sharp diffraction peaks emerged, which can be indexed as AlCu2Zr and Cu5Zr phases. It is clear that the width of the peaks became narrower, suggesting that the residual internal stresses decreased and the crystalline structures of phases were complete. In order to clarify the structure of the Cu–Zr–Ag–Al BMGs after CT, the microstructures for BMG samples were investigated using
85
TEM and HRTEM. Fig. 2(a) shows a TEM bright-field image with inset corresponding selected area diffraction from the as-cast Cu45Zr45Ag7Al3 bulk metallic glass. The TEM bright-field image and the corresponding selected area diffraction pattern exhibit no visible crystals and only halo diffraction ring also indicating the formation of an amorphous phase, shown in Fig. 2(a). This feature of image is almost the same to the literature reported by Zeng et al. [13]. Fig. 2(b) exhibits bright-field TEM image for the sample after CT 96 h. The SEAD pattern (inset in Fig. 2(b)) presents not only diffuse diffraction intensity but also Debye spots. The HRTEM image of the same specimen in Fig. 2(c) shows that the size of nanoparticle has been precipitated and is less than 10 nm. This is to say that there are some nanocrystalline structure phases occurred in the glassy matrix. Fig. 3(a) and (b) show that these TEM images and SAED patterns of the Cu45Zr45Ag7Al3 alloys after CT 120 h and 240 h, respectively. Some sub-micro particles are seen in the image. The electron diffraction patterns of them suggest that the crystalline phases respectively correspond to the ClCs structure AlCu2Zr (Fig. 3(a)), and the face-centered cubic (fcc) Cu5Zr (Fig. 3(b)), which has been frequently reported [14]. And it is clear that the structure of the phases has not changed and the amount of the particles did not change obviously, comparing 192 h with 240 h. 3.2. Compressive behaviors
Fig. 1. X-ray diffractions of the Cu45Zr45Ag7Al3 amorphous alloys with different treatments.
Fig. 4 shows compressive stress–strain curves of the cylindrical Cu45Zr45Ag7Al3 glass alloy rods in as-cast and CT states, and at an initial strain rate of 5 10 4 s 1. The slopes of the stress–strain curves were calibrated by the Young's modulus obtained by an
Fig. 2. TEM and HREM images as well as corresponding selected-area electron diffraction patterns of the Cu45Zr45Ag7Al3 as-cast bulk metallic glass (a) and alloy after CT 96 h (b) and (c).
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Fig. 3. Bright field TEM images of the Cu45Zr45Ag7Al3 alloys CT after 192 h (a) and 240 h (b) and corresponding selected-area electron diffraction patterns: (a) AlCu2Zr and (b) Cu5Zr.
Fig. 4. Compressive strain–stress curves of Cu45Zr45Ag7Al3 BMG and its CT alloys: as-cast (a), CT after 96 h (b), 192 h (c) and 240 h (d).
ultrasonic pulse method [15]. The values of the Young's modulus of the as-cast and CT (96 h, 192 h and 240 h) were 108, 109, 107 and 108 GPa, respectively. Because the difference of the Young's modulus is too small, it is difficult to see the difference of the slopes of the stress–strain curves, as shown in Fig. 4. And the fracture strength of them were 1510, 1790, 1910 and 1892 MPa. The as-cast and CT 192 h samples show initial elastic strain, followed by final fracture immediately after the onset of yielding. The CT 240 h sample shows that final fracture immediately and the maximum of the compression stress is 1910 MPa respectively after CT 192 h. It is clear that the compressive strength increases with increasing CT time and plastic strain reaches the maximum (0.5%) after CT 96 h. 3.3. Fracture morphology Fig. 5 shows SEM micrographs exhibiting the compressive fracture feature of the as-cast and Cu45Zr45Ag7Al3 BMGs with different CT times, which were tested at a strain rate of 5 10 4 s 1. It is shown in Fig. 5(a), that the fracture surface is similar to many previous studies [16,17], in which the vein-like patterns in the fracture surface are clearly observed. Furthermore, the arrangement of the veins was in agreement with the flow direction in the shear plane, which was consistent with the previous literatures [16,18]. The partly nanocrystallized alloy which exhibits high strength of 1790 MPa (Fig. 4(b)), which is higher than those of the as-cast alloy. Furthermore, distinct plastic
deformation of about 0.5% for the partly nanocrystallized alloy after CT 96 h is also observed presumably, because the nanoparticles can suppress the deformation of shear bands and shear band initiation in glassy alloys which is related to free volume coalescence [19,20]. In addition, the fracture surface of the alloy after CT 96 h is still in the vein-like pattern type (Fig. 5(b)). As the CT time rises to 192 h, some cracks can be seen on the fracture surface, but some vein-like patterns in the surface are still clearly observed, shown in Fig. 5(c), and the compressive strength and micro-hardness reach the maximum. In some regions, localized remelting was observed (Fig. 5(c)), indicating the concentration of the existence of phases with low melting temperatures, e.g. glassy phase. With further increasing CT time to 240 h not only the typical brittle fracture mode but also vein-like pattern can be seen, comparing with CT 192 h, as shown in Fig. 4(d) and Fig. 5(d). It can be seen that Cu45Zr45Ag7Al3 bulk metallic glass has a tendency of CT induced hardening especially in the long time CT.
4. Discussion It is thus concluded that glassy alloys become stiff and rigid at CT [21]. This means that the effective atomic distance of the alloys decreases with increasing CT time, resulting in squeezing the free volumes. According to the literatures [22,23], the free volume coalescence plays the key role in the shear band initiation in glassy alloys and the shear bands form just in the higher free volume fraction. However, it is well established that the diffusion of the free volume is related to the temperature [22,23]. It has currently been accepted that shear band initiation in glass alloys is related to free volume coalescence [19,24]. Since the mobility of the free volume at cryogenic temperature is much lower than that at room temperature, stiffness of atomic bond increases through the increase in the difficulty of free volume coalescence. The coalescence of free volume requires a higher applied load, resulting in an increase of the maximum compressive stress. In addition, amorphous state is thermodynamically unstable and has a larger excessive volume than the counterpart of its crystalline. Crystallization is a process of increasing density. Pressure is therefore favorable to crystallization. Some researchers [24–26] have confirmed this conclusion. In general, crystallization includes crystal nucleation and subsequent growth. The thermodynamic potential barrier governs the initial stage of nucleation. According to the classical theory of nucleation, the following expression is generally used to predict the
L.K. Zhang et al. / Physica B 433 (2014) 84–88
87
Fig. 5. Fracture morphologies of the Cu45Zr45Ag7Al3 bulk metallic glass and its CT samples: as-cast (a), CT 96 h (b), CT 192 h (c), and CT 240 h (d).
homogeneous nucleation [27]. ΔG ¼
16π s3 16πs3 ¼ 3 ΔG2v 3ðGc Gam Þ2
ð1Þ
r¼
2s ΔGv
ð2Þ
where ΔGv is the difference between the Gibbs free energy of amorphous alloy (Gam) and crystallization phase (Gc) per unit volume, ΔG is the critical energy, s is the nucleus–matrix specific interface free energy, and r is the critical radius of nucleation. Taking pressure (P) into account and assuming that the interfacial energy (s) is pressureindependent, the change in ΔG with pressure (P) at a fixed temperature is obtained from Eq. (1) ∂ðΔGÞ 32πs3 1 ∂ΔGv ¼ ð3Þ 3 ∂P 3 ðΔGv Þ ∂P
∂ðr Þ 2s ∂ΔGv ¼ ∂P ∂P ðΔGv Þ2
ð4Þ
Since Gam 4Gc, assuming that the interfacial energy (s) is pressure-independent, the first term in Eq. (3) is always negative, implying that pressure can promote crystallization. Eq. (3) indicates that pressure may promote crystallization when long-range diffusion of the atoms is not needed in amorphous alloy. That is to say, if the term on the right side of Eq. (3) may be neglected, the nucleation activation energy (G) decreases with increasing the pressure. And Eq. (4) indicate that the critical nucleation radius decreases with increasing pressure in the process of the crystallization of the BMGs. Therefore, the crystal nucleus, that is smaller than the critical nucleation radius in the absence of pressure, exceeds the critical value under pressure and become stable phase. In the present study, CT-induced phase precipitation from the glassy matrix may be due to the compressive stress during the CT
process [28,29]. The phase precipitation makes the metallic glass unstable, and crystallization becomes more easily. At the same time, with the further CT, the volume of amorphous alloys decreases under the pressure, which will promote the atoms to recompose in a short-ranged area [30,31]. The atoms of amorphous matrix could “short-ranged remove” directionally under the pressure. Furthermore, the pressure can reduce the critical nucleation radius. Crystal nucleus can therefore change into crystal particles and the nucleation occurs. In addition, under pressure the constituent atoms in amorphous matrix are “pressed transported” into the interface of the grains, which promotes the growth of grains. This might be the reason why sub-micron phases (AlCu2Zr and Cu5Zr) may precipitate in the Cu45Zr45Ag7Al3 amorphous alloy after a long-time CT. As mentioned above, nanocrystalline structure and crystalline phases are formed in the Cu45Zr45Ag7Al3 BMG subjected to an optimum CT. The crystallized structure seriously changes the mechanical properties and fracture morphology. That is, the deformation behavior is associated with the nature of crystallites precipitated in the glassy matrix. CT 96 h, i.e., results in the formation of nanoparticles with size smaller than 10 nm in the glass matrix. To the best of our knowledge, one of the mechanisms of deformation-induced crystallization in BMGs is stress [25,30,32,33], although other factors, such as large local shear strain by deformation and local temperature rise in the shear bands resulting from adiabatic heating, may also play important roles in the deformation-induced crystallization [34,35]. In the current study, the cryogenic surrounding may cause the generation of internal stress in glass matrix. Under pressure, the densities of the CT samples are larger than that of the as-cast alloy, which will promote the atoms to rearrangement in short-range [31,36]. So, the atoms of amorphous matrix could perform short-ranged removal directionally under the pressure. In this condition, some nano-size phases may be formed in the glass matrix, which is shown in Fig. 3. With further CT, the amorphous matrix will contract further, resulting
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that the density of the crystalline phases are larger than the amorphous phase. Then some atoms are “pressed transported” into the surface of nanocrystalline, which will promote the growth of the phases (AlCu2Zr and Cu5Zr). As a result of this change in the microstructure, the mechanical properties of CT samples increased in comparison to the as-cast samples. So far the phenomenon that CT induced crystalline has not been reported yet. Further studies are in progress. 5. Conclusion Cu45Zr45Ag7Al3 bulk metallic glass shows apparent hardening due to long-time CT. After CT 96 h, nano-size composites are formed in alloys, and higher strength of 1790 MPa and distinct plastic deformation of 0.5% are obtained in the alloy. AlCu2Zr and Cu5Zr phases are formed in the alloys that exhibited the maximum compressive strength of 1910MPa after CT 196 h. The compressive strength increases for the alloys with increasing CT time. The as-cast bulk glassy alloys and the alloy CT 96 h show vein-like patterns. On the other hand, the fracture morphology of the alloys CT 192 h and 240 h changes to brittle fracture and some cracks can be seen. Acknowledgment Authors greatly acknowledge the financial support by Program for New Century Excellent Talents in University (NCET-10-03600). References [1] A. Inoue, Mater. Trans JIM 36 (1995) 866. [2] A.L. Greer, Science 267 (1995) 1947.
[3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25] [26] [27] [28] [29] [30] [31] [32] [33] [34] [35] [36]
W.L. Johnson, MRS Bull. 24 (1999) 42. A. Inoue, Acta Mater. 48 (2000) 279. A. Inoue, B.L. Shen, A. Takeuchi, Mater. Trans. 47 (2006) 1275. A. Inoue, Acta Mater. 48 (2000) 279. F.R. Niessen, Cohesion in Metals, Elsevier Science Publishers, Amsterdam, 1988. W. Dmowski, C. Fan, M.L. Morrison, P.K. Liaw, T. Egami, Mater. Sci. Eng. A 471 (2007) 125. K.B. Kim, P.J. Warren, B. Cantor, J. Non-Cryst. Solids 353 (2007) 3338. A. Kawashima, Y.Q. Zeng, M. Fukuhara, Mater. Sci. Eng A 498 (2008) 475. K.M. Asl, A. Tari, F. Khomamizadeh, Mater. Sci Eng. A 523 (2009) 1. G.Z. Ma, D. Chen, Y. Jiang, W. Li, Intermetallics 18 (2010) 1254. Y.Q. Zeng, N. Nishiyama, A. Inoue., Mater Trans. 48 (2007) 1355. W. Heine, U. Zwicker, Naturwissenschaften 49 (1962) 391. M. Fukuhara, A. Kawashima, W. Zhang, A. Inoue, F. Yin., J. Appl. Phys. 103 (2008) 013503. Z.F. Zhang, J. Eckert, L. Schultz, Acta Mater. 51 (2003) 1167. Z. Bian, G.L. Chen, G. He, X.D. Hui, Mater. Sci. Eng. A 316 (2001) 135. Z. Bian, G. He, G.L. Chen, Trans. Nonferrous Metals Soc. China 10 (2000) 345. A.S. Argon, Acta Metall. 27 (1979) 47. K.M. Flores, R.H. Dauskardt, Acta Mater. 49 (2001) 2527. A. Kawashima, T. Okuno, H. Kurishita, W. Zhang, H. Kimura, A. Inoue, Mater. Trans. 48 (2007) 2787. A. Kawashima, Y.Q. Zeng, M. Fukuhara, Mater. Sci. Eng. A 498 (2008) 475. Y.J. Huang, J. Shen, J.F. Sun, Mater. Sci Eng. A 498 (2008) 203. V.F. Degtyareva, F. Porsch, E.G. Ponyatovskiian, W.B. Holzapfel, Phys. Rev. B 53 (1996) 8337. L.L. Sun, W.K. Wang, D.W. He, et al., Appl. Phys. Lett. 76 (2000) 2874. X.Y. Zhang, J.W. Zhang, W.K. Wang., J. Appl. Phys. 89 (2001) 477. D.A. Porter, K.E. Easterling, Phase Transformations in Metal and Alloys, Van Nostrand-Reinhold, New York, 1981, p. 263. S. Pauly, G. Liu, G. Wang, J. Eckert., Acta Mater. 57 (2009) 5445. J. Das, S. Paulya, M. Bostrom, et al., J. Alloys Compd. 483 (2009) 97. K. Lu, J.T. Wang., Acta Metall. Sin. 26 (1990) 316. K. Lu, J.T. Wang, L. Dong., Acta Metall. Sin. 27 (1991) 108. X.Y. Zhang, J.W. Zhang, W.K. Wang, J. Appl. Phys. 89 (2001) 477. A. Inoue, A. TaKeuchi, Mater. Trans. JIM 36 (1995) 963. B. Lohwongwatana, J. Schroers, W.L. Johnson, Phys. Rev. Lett 96 (2006) 075503. L.Y. Chen, A.D.H. Setyawan, H. Kato, A. Inoue, G.Q. Zhang, J. Saida, et al., Scr. Mater. 59 (2008) 75. M. Sutton, Y.S. Yang, J. Mainville, SJL. Jordan, K.F. Ludwig Jr., G.B. Stephenson, Phys. Rev. Lett. 62 (1989) 288.
Contents lists available at ScienceDirect
Physica B journal homepage: www.elsevier.com/locate/physb
Cryogenic treatment induced hardening of Cu45Zr45Ag7Al3 bulk metallic glass L.K. Zhang, Z.H. Chen, D. Chen n, Q. Zheng College of Materials Science and Engineering, Hunan University, Changsha 410082, PR China
art ic l e i nf o
a b s t r a c t
Article history: Received 27 August 2013 Received in revised form 24 September 2013 Accepted 1 October 2013 Available online 18 October 2013
Little is known about the mechanical properties of bulk metallic glassy alloys (BMGs) after cryogenic treatment (CT). Cu45Zr45Ag7Al3 BMGs were treated in cryogenic temperatures for four chosen holding times. Compared with the as-cast sample, the compressive fracture strength of this BMG increased with increasing CT time. Furthermore, the CT changes the fracture mode from the ductile to the brittle. And CT brought about the precipitation of the AlCu2Zr and Cu5Zr phases which were accompanying with the morphology modification and contributed to the mechanical properties improvement of this BMG. & 2013 Elsevier B.V. All rights reserved.
Keywords: Bulk metallic glass Cryogenic treatment Mechanical properties Harding Fracture
1. Introduction
2. Experiments
Bulk metallic glasses (BMGs), which exhibit short-range structural periodicity, possess many unique physical, mechanical and chemical properties [1–5]. Among those most promising bulk metallic glasses, Cu–Zr–Ag–Al alloy system has exhibited very special not only for its high GFA, but also for its unique mechanical properties, such as extremely high strength and good toughness [6,7]. Thus, the Cu45 Zr45Ag7Al3 was chosen for the current investigation. Although many current studies are directed toward researching the microstructure of BMGs or bulk metallic glass matrix composites (BMGCs) during the annealing process [8,9], little is known about microstructure and mechanical properties of BMGs and BMGCs after cryogenic temperatures. As we know, CT was widely used for high precision parts, and has been shown to result in significant increase in mechanical properties of steels, which was caused by the reduction in the amount of retained austenite and partly transforming of the structure induced by the internal stress of CT [1,10,11]. Recently, we have reported that a Cu46Zr46Al8 matrix BMGs exhibit transformation of microstructure and mechanical behaviors after CT [12]. The objective of the present work was to clarify the microstructure of the Cu45Zr45Ag7Al3 bulk specimens after CT. Furthermore, the effects of the CT on the micro-hardness and the compression fracture strength of the samples were also investigated.
The material used in the present study has a composition of Cu45Zr45Ag7Al3 (atomic percent). Alloy ingots were prepared by arc melting the mixture of metals (4 99.9 mass%) in an argon atmosphere, followed by suction casting into a copper mold. From these master alloys, cast cylinders with 4 mm diameter and 50 mm length were obtained by suction-casting into a copper mold under a purified argon atmosphere. The as-cast samples were cooled to 80 K with a rate of 20 K/h in a controlled box and kept at this temperature for different periods of time (96, 192 and 240 h). The phase constitute was examined by X-ray diffractometry (XRD, D/Max-2550) using a monochromatic Cu Kα radiation. The microstructure was investigated by using a Tecnai F220 highresolution transmission electron microscopy (HRTEM), operating at 300 kV. The mechanical testing was performed by compression at room temperature. The strain rate was 5 10 4s 1. The test samples were 4 mm in diameter and 8 mm in length. Fracture surface morphology was examined by JSM-6700F scanning electron microscopy (SEM).
3. Results 3.1. Microstructure
n
Corresponding author. Tel./fax: þ 86 731 88821648. E-mail address: [email protected] (D. Chen).
0921-4526/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.physb.2013.10.003
Fig. 1 shows the XRD patterns of the Cu45Zr45Ag7Al3 amorphous alloys treated under different conditions. The as-cast sample
L.K. Zhang et al. / Physica B 433 (2014) 84–88
exhibits a broad diffraction peak, which is the characteristic of an amorphous structure. After 192 h CT, several sharp diffraction peaks emerged, which can be indexed as AlCu2Zr and Cu5Zr phases. It is clear that the width of the peaks became narrower, suggesting that the residual internal stresses decreased and the crystalline structures of phases were complete. In order to clarify the structure of the Cu–Zr–Ag–Al BMGs after CT, the microstructures for BMG samples were investigated using
85
TEM and HRTEM. Fig. 2(a) shows a TEM bright-field image with inset corresponding selected area diffraction from the as-cast Cu45Zr45Ag7Al3 bulk metallic glass. The TEM bright-field image and the corresponding selected area diffraction pattern exhibit no visible crystals and only halo diffraction ring also indicating the formation of an amorphous phase, shown in Fig. 2(a). This feature of image is almost the same to the literature reported by Zeng et al. [13]. Fig. 2(b) exhibits bright-field TEM image for the sample after CT 96 h. The SEAD pattern (inset in Fig. 2(b)) presents not only diffuse diffraction intensity but also Debye spots. The HRTEM image of the same specimen in Fig. 2(c) shows that the size of nanoparticle has been precipitated and is less than 10 nm. This is to say that there are some nanocrystalline structure phases occurred in the glassy matrix. Fig. 3(a) and (b) show that these TEM images and SAED patterns of the Cu45Zr45Ag7Al3 alloys after CT 120 h and 240 h, respectively. Some sub-micro particles are seen in the image. The electron diffraction patterns of them suggest that the crystalline phases respectively correspond to the ClCs structure AlCu2Zr (Fig. 3(a)), and the face-centered cubic (fcc) Cu5Zr (Fig. 3(b)), which has been frequently reported [14]. And it is clear that the structure of the phases has not changed and the amount of the particles did not change obviously, comparing 192 h with 240 h. 3.2. Compressive behaviors
Fig. 1. X-ray diffractions of the Cu45Zr45Ag7Al3 amorphous alloys with different treatments.
Fig. 4 shows compressive stress–strain curves of the cylindrical Cu45Zr45Ag7Al3 glass alloy rods in as-cast and CT states, and at an initial strain rate of 5 10 4 s 1. The slopes of the stress–strain curves were calibrated by the Young's modulus obtained by an
Fig. 2. TEM and HREM images as well as corresponding selected-area electron diffraction patterns of the Cu45Zr45Ag7Al3 as-cast bulk metallic glass (a) and alloy after CT 96 h (b) and (c).
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L.K. Zhang et al. / Physica B 433 (2014) 84–88
Fig. 3. Bright field TEM images of the Cu45Zr45Ag7Al3 alloys CT after 192 h (a) and 240 h (b) and corresponding selected-area electron diffraction patterns: (a) AlCu2Zr and (b) Cu5Zr.
Fig. 4. Compressive strain–stress curves of Cu45Zr45Ag7Al3 BMG and its CT alloys: as-cast (a), CT after 96 h (b), 192 h (c) and 240 h (d).
ultrasonic pulse method [15]. The values of the Young's modulus of the as-cast and CT (96 h, 192 h and 240 h) were 108, 109, 107 and 108 GPa, respectively. Because the difference of the Young's modulus is too small, it is difficult to see the difference of the slopes of the stress–strain curves, as shown in Fig. 4. And the fracture strength of them were 1510, 1790, 1910 and 1892 MPa. The as-cast and CT 192 h samples show initial elastic strain, followed by final fracture immediately after the onset of yielding. The CT 240 h sample shows that final fracture immediately and the maximum of the compression stress is 1910 MPa respectively after CT 192 h. It is clear that the compressive strength increases with increasing CT time and plastic strain reaches the maximum (0.5%) after CT 96 h. 3.3. Fracture morphology Fig. 5 shows SEM micrographs exhibiting the compressive fracture feature of the as-cast and Cu45Zr45Ag7Al3 BMGs with different CT times, which were tested at a strain rate of 5 10 4 s 1. It is shown in Fig. 5(a), that the fracture surface is similar to many previous studies [16,17], in which the vein-like patterns in the fracture surface are clearly observed. Furthermore, the arrangement of the veins was in agreement with the flow direction in the shear plane, which was consistent with the previous literatures [16,18]. The partly nanocrystallized alloy which exhibits high strength of 1790 MPa (Fig. 4(b)), which is higher than those of the as-cast alloy. Furthermore, distinct plastic
deformation of about 0.5% for the partly nanocrystallized alloy after CT 96 h is also observed presumably, because the nanoparticles can suppress the deformation of shear bands and shear band initiation in glassy alloys which is related to free volume coalescence [19,20]. In addition, the fracture surface of the alloy after CT 96 h is still in the vein-like pattern type (Fig. 5(b)). As the CT time rises to 192 h, some cracks can be seen on the fracture surface, but some vein-like patterns in the surface are still clearly observed, shown in Fig. 5(c), and the compressive strength and micro-hardness reach the maximum. In some regions, localized remelting was observed (Fig. 5(c)), indicating the concentration of the existence of phases with low melting temperatures, e.g. glassy phase. With further increasing CT time to 240 h not only the typical brittle fracture mode but also vein-like pattern can be seen, comparing with CT 192 h, as shown in Fig. 4(d) and Fig. 5(d). It can be seen that Cu45Zr45Ag7Al3 bulk metallic glass has a tendency of CT induced hardening especially in the long time CT.
4. Discussion It is thus concluded that glassy alloys become stiff and rigid at CT [21]. This means that the effective atomic distance of the alloys decreases with increasing CT time, resulting in squeezing the free volumes. According to the literatures [22,23], the free volume coalescence plays the key role in the shear band initiation in glassy alloys and the shear bands form just in the higher free volume fraction. However, it is well established that the diffusion of the free volume is related to the temperature [22,23]. It has currently been accepted that shear band initiation in glass alloys is related to free volume coalescence [19,24]. Since the mobility of the free volume at cryogenic temperature is much lower than that at room temperature, stiffness of atomic bond increases through the increase in the difficulty of free volume coalescence. The coalescence of free volume requires a higher applied load, resulting in an increase of the maximum compressive stress. In addition, amorphous state is thermodynamically unstable and has a larger excessive volume than the counterpart of its crystalline. Crystallization is a process of increasing density. Pressure is therefore favorable to crystallization. Some researchers [24–26] have confirmed this conclusion. In general, crystallization includes crystal nucleation and subsequent growth. The thermodynamic potential barrier governs the initial stage of nucleation. According to the classical theory of nucleation, the following expression is generally used to predict the
L.K. Zhang et al. / Physica B 433 (2014) 84–88
87
Fig. 5. Fracture morphologies of the Cu45Zr45Ag7Al3 bulk metallic glass and its CT samples: as-cast (a), CT 96 h (b), CT 192 h (c), and CT 240 h (d).
homogeneous nucleation [27]. ΔG ¼
16π s3 16πs3 ¼ 3 ΔG2v 3ðGc Gam Þ2
ð1Þ
r¼
2s ΔGv
ð2Þ
where ΔGv is the difference between the Gibbs free energy of amorphous alloy (Gam) and crystallization phase (Gc) per unit volume, ΔG is the critical energy, s is the nucleus–matrix specific interface free energy, and r is the critical radius of nucleation. Taking pressure (P) into account and assuming that the interfacial energy (s) is pressureindependent, the change in ΔG with pressure (P) at a fixed temperature is obtained from Eq. (1) ∂ðΔGÞ 32πs3 1 ∂ΔGv ¼ ð3Þ 3 ∂P 3 ðΔGv Þ ∂P
∂ðr Þ 2s ∂ΔGv ¼ ∂P ∂P ðΔGv Þ2
ð4Þ
Since Gam 4Gc, assuming that the interfacial energy (s) is pressure-independent, the first term in Eq. (3) is always negative, implying that pressure can promote crystallization. Eq. (3) indicates that pressure may promote crystallization when long-range diffusion of the atoms is not needed in amorphous alloy. That is to say, if the term on the right side of Eq. (3) may be neglected, the nucleation activation energy (G) decreases with increasing the pressure. And Eq. (4) indicate that the critical nucleation radius decreases with increasing pressure in the process of the crystallization of the BMGs. Therefore, the crystal nucleus, that is smaller than the critical nucleation radius in the absence of pressure, exceeds the critical value under pressure and become stable phase. In the present study, CT-induced phase precipitation from the glassy matrix may be due to the compressive stress during the CT
process [28,29]. The phase precipitation makes the metallic glass unstable, and crystallization becomes more easily. At the same time, with the further CT, the volume of amorphous alloys decreases under the pressure, which will promote the atoms to recompose in a short-ranged area [30,31]. The atoms of amorphous matrix could “short-ranged remove” directionally under the pressure. Furthermore, the pressure can reduce the critical nucleation radius. Crystal nucleus can therefore change into crystal particles and the nucleation occurs. In addition, under pressure the constituent atoms in amorphous matrix are “pressed transported” into the interface of the grains, which promotes the growth of grains. This might be the reason why sub-micron phases (AlCu2Zr and Cu5Zr) may precipitate in the Cu45Zr45Ag7Al3 amorphous alloy after a long-time CT. As mentioned above, nanocrystalline structure and crystalline phases are formed in the Cu45Zr45Ag7Al3 BMG subjected to an optimum CT. The crystallized structure seriously changes the mechanical properties and fracture morphology. That is, the deformation behavior is associated with the nature of crystallites precipitated in the glassy matrix. CT 96 h, i.e., results in the formation of nanoparticles with size smaller than 10 nm in the glass matrix. To the best of our knowledge, one of the mechanisms of deformation-induced crystallization in BMGs is stress [25,30,32,33], although other factors, such as large local shear strain by deformation and local temperature rise in the shear bands resulting from adiabatic heating, may also play important roles in the deformation-induced crystallization [34,35]. In the current study, the cryogenic surrounding may cause the generation of internal stress in glass matrix. Under pressure, the densities of the CT samples are larger than that of the as-cast alloy, which will promote the atoms to rearrangement in short-range [31,36]. So, the atoms of amorphous matrix could perform short-ranged removal directionally under the pressure. In this condition, some nano-size phases may be formed in the glass matrix, which is shown in Fig. 3. With further CT, the amorphous matrix will contract further, resulting
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that the density of the crystalline phases are larger than the amorphous phase. Then some atoms are “pressed transported” into the surface of nanocrystalline, which will promote the growth of the phases (AlCu2Zr and Cu5Zr). As a result of this change in the microstructure, the mechanical properties of CT samples increased in comparison to the as-cast samples. So far the phenomenon that CT induced crystalline has not been reported yet. Further studies are in progress. 5. Conclusion Cu45Zr45Ag7Al3 bulk metallic glass shows apparent hardening due to long-time CT. After CT 96 h, nano-size composites are formed in alloys, and higher strength of 1790 MPa and distinct plastic deformation of 0.5% are obtained in the alloy. AlCu2Zr and Cu5Zr phases are formed in the alloys that exhibited the maximum compressive strength of 1910MPa after CT 196 h. The compressive strength increases for the alloys with increasing CT time. The as-cast bulk glassy alloys and the alloy CT 96 h show vein-like patterns. On the other hand, the fracture morphology of the alloys CT 192 h and 240 h changes to brittle fracture and some cracks can be seen. Acknowledgment Authors greatly acknowledge the financial support by Program for New Century Excellent Talents in University (NCET-10-03600). References [1] A. Inoue, Mater. Trans JIM 36 (1995) 866. [2] A.L. Greer, Science 267 (1995) 1947.
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