- Email: [email protected]
Improved plasticity and fracture toughness in metallic glasses via surface crystallization
Intermetallics 19 (2011) 1420e1427
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Improved plasticity and fracture toughness in metallic glasses via surface crystallization Jitang Fan a, b, *, Aiying Chen b, Juan Wang b, Jun Shen c, Jian Lu b, d, ** a
Faculty of Civil Engineering and Geosciences, Delft University of Technology, P.O. Box 5048, 2600 GA, Delft, The Netherlands Department of Mechanical Engineering, The Hong Kong Polytechnic University, Hong Kong, PR China c School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, PR China d Department of Mechanical Engineering, City University of Hong Kong, Hong Kong, PR China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 3 December 2009 Received in revised form 28 February 2011 Accepted 13 May 2011 Available online 16 June 2011
Mechanical crystallization was induced in the monolithic bulk metallic glasses by surface mechanical attrition treatment (SMAT) to create the isolated crystallite islands in the top surface layer. Inside the isolated crystallite islands, microstructure consists of the crystallites with a gradient grain size evolution and the residual amorphous phase. Moreover, isolated crystallite islands, acting as the obstacles to restrict the highly localized deformation of shear bands/cracks, effectively limit the shear bands extension, suppress the shear bands opening, and avoid the cracks developing, which significantly enhance the overall plasticity and fracture toughness. They were suggested by the secondary shear bands in the glassy matrix and fine-shearing together with micro-cracking inside the isolated crystallite islands. Finally, the improved plasticity and fracture toughness were systematically discussed. Based on the current results, surface crystallization is proposed to optimize the mechanical properties of metallic glasses. Ó 2011 Elsevier Ltd. All rights reserved.
Keywords: B. Glasses, metallic A. Composite, based on the metallic glass matrix intermetallics, miscellaneous B. Deformation map B. Fracture toughness B. Mechanical properties at ambient temperature
1. Introduction Bulk metallic glasses (BMGs) have inspired considerable scientific interest and industrial attention, due to their unique physical, chemical and mechanical properties [1,2]. Particularly, the characteristics of high strength and hardness are regarded as a potential candidate for the applications in structural engineering [3]. However, although computer simulation shows metallic glasses can display the superhigh plastic strain in the individual shear band on the microscopic scale [4,5], they, macroscopically, effectively lack the ductility with an apparently brittle fracture model. This is a fatal weakness to impede their application in the engineering field. Thus, improving general plasticity has been an urgent topic pursued in the metallic glass community for the recent several years. As pointed out by Chen et al. [6], variation in local material
* Corresponding author. Faculty of Civil Engineering and Geosciences, Delft University of Technology, P.O. Box 5048, 2600 GA, Delft, The Netherlands. ** Corresponding author. Department of Mechanical Engineering, The Hong Kong Polytechnic University, Hong Kong, PR China. E-mail addresses: [email protected] (J. Fan), [email protected] (J. Lu). 0966-9795/$ e see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2011.05.012
properties could arrest persistent slip on the individual shear band, in order to activate the multiple shear bands and contribute to the global plasticity. Generally, introducing a soft/plastic second phase is effective to improve the plasticity of metallic glasses [7e14], such as solid solution ductile particles [7e9], crystalline dendrite phases [10e13]. Especially, Zr-based metallic glass matrix composite exhibited the macroscopic room-temperature tensile ductility [14]. Besides, nanocrystallites are also profitable to improve the plasticity of metallic glasses [15,16]. In a word, the ‘composite’ of crystallites and metallic glasses was successful to optimize the mechanical properties. However, preparing the metallic glasses matrix composite usually has an acerbic demand for the glass forming ability and the casting technology. Thus, an interesting question is issued: whether there is a well-controlled, subsequent treating method to enhance the plasticity of BMGs. The correlative studies include shot-peening [17], coating [18], rolling [19] etc. Subsequently, a new strategy was proposed to improve the plasticity of BMGs, i.e. strengthening the surface of metallic glasses [20]. Since surface mechanical attrition treatment (SMAT) was proposed in 1999 [21], it has been proven to be an effective approach to optimize the surface microstructure and properties,
J. Fan et al. / Intermetallics 19 (2011) 1420e1427
Fig. 1. XRD diffraction patterns taken from the top surfaces of Ti40Zr25Ni3Cu12Be20 alloy after SMAT process with different treating time: 0, 5, 15 and 30 min.
enabling significantly to enhance the global performance of metallic materials [22]. In this work, we show a detailed study on the surface crystallization induced by SMAT process as well as the resultant improvement in mechanical properties. 2. Experimental procedures Alloy ingots with nominal chemical composition of Ti40Zr25Ni3Cu12Be20 (at.%) were prepared by arc melting the mixtures of elemental pieces with a purity above 99.9% in a Ti-gettered highly pure argon atmosphere. Remelting for five times was performed to ensure the homogeneity of composition. The final ingots were cast
1421
into a rectangular plate with a dimension of 3 30 70 mm3 using the copper mould. Surface mechanical attrition treatment (SMAT) is a subsequent treating approach to realize the specific structure and property requirements in the top surface layer [21]. In the current study, SMAT was performed by using high-hardness balls with a 2 mm diameter under a vibrating frequency of 20 kHz at roomtemperature. In order to avoid the temperature rise, treating process was done in a time interval of 5 min with a break time of 2 min. As well, cooling air flow was applied. Before treated, specimens were electric spark cut, and then were machined carefully into rectangle gauge dimension of 3 3 70 mm3. Four lateral surfaces of every specimen were treated for the same time duration. Treating time duration is 5 min, 15 min and 30 min respectively, which formed the four groups of samples in addition to the as-cast specimens. One typical specimen from each group was analyzed using Philips X-ray diffractometer (XRD) with Cu-Ka radiation as a source. And, microstructure analysis was performed for a typical specimen treated for 30 min by transmission electron microscopy (TEM JEM 2010) operated at 200 kV. TEM specimens were taken from the regions with a depth ranging to 90 mm below the top surface. The plane-view TEM foils were slowly ion-thinned at low temperature. Besides, statistical distribution of the grain size was quantitatively measured from the bright and dark-field TEM micrographs. Room-temperature compression tests were performed using a standard MTS Alliance RT/50 Materials Testing System, at a strain rate of 1 104 s1. For each group of samples, three specimens or more were tested with the dimension of 3 3 6 mm3, and the average data was used as the final results. Fracture morphologies, as well as the top surfaces of the treated samples, were observed by scanning electron microscopy (SEM Hitachi S-4200 field emission).
Fig. 2. SEM images taken from the top surface of the treated (for 30 min) sample: (a) homogeneous distribution of isolated crystallite islands on the top surface; (b) melting veinlike structure together with some crystalline particles; (c) amplified vein-like structure of region _ showing the interior nanocrystallites; (d) amplified crystallized particles of region II showing the homogeneous crystalline particles with the grain size of from about 200e500 nm.
1422
J. Fan et al. / Intermetallics 19 (2011) 1420e1427
Fig. 3. TEM and HRTEM images characterizing the crystallization behavior of the treated (for 30 min) sample: (a) bright-field image showing the isolated crystallite island taken from about 60 mm below the top surface; (b) bright-field image showing the homogeneous crystallites in the glassy matrix, and the corresponding SAED patterns in the inset; (c) corresponding dark-field pair; (d) HRTEM image showing the clear lattice fringes in the glassy matrix with the long-range ordering; (e) HRTEM image taken from about 80 mm below the top surface, showing the amorphous nature and the corresponding SAED pattern in the inset.
J. Fan et al. / Intermetallics 19 (2011) 1420e1427
3. Experimental results 3.1. Isolated crystallite islands formation Crystallization behavior in metallic glasses has been studied to be induced by annealing, plastic deformation, room-temperature fracture, and ball milling [23e27]. For the current work, mechanical crystallization was characterized to be induced by the SMAT process. XRD diffraction patterns recorded from the top surfaces are shown in Fig. 1. For the as-cast specimen, i.e. treated for 0 min, the pattern consists of only a broad diffraction maximum without any detectable sharp Bragg peaks, indicating the fully amorphous microstructure. And, for the treated specimens, it can be clearly seen that XRD patterns are composed of a broad diffusion background and a set of several sharp crystalline peaks, which suggest the mixtures of glass matrix and crystalline phases. Crystalline phases are identified as the Cu2TiZr ternary compound, and here maybe exists some incomplete nuclei [28,29]. In addition, the morphology of the treated (for 30 min) samples was revealed from the top surface by high-resolution SEM, as shown in Fig. 2. A typical surface roughness was induced with a well distribution in the top surface (Fig. 2(a)). Wherein, it is interesting to find that some crystalline phases congregate, named as isolated crystallite island, as pointed out by the arrows. Isolated crystallite islands disperse in the rough surface with the size of about 10e100 mm. Local amplified image displays the coexistence of bugle-like patterns and crystalline particles inside the isolated crystallite island, as shown in regions I and II respectively (Fig. 2(b)). This is very similar to the nanocrystallization behavior of Co- and Fe-based BMGs, induced by the quasi-static compression fracture at room-temperature [26]. Moreover, HRSEM image of region I clearly shows the bugle-like patterns, which consist of the elongated viscous liquid and interior nanocrystallites (Fig. 2(c)), implying the local plastic deformation and crystallization behavior during SMAT process. The same, HRSEM image of region II, seemingly overlaying the region I, shows the crystallites which are uniform with the grain size of about 200e500 nm, recognized as submicron crystals (Fig. 2(d)). The formation of the viscous liquid and crystalline particles in the top surface should be attributed to the severe plastic deformation during the SMAT process. To further shed light on the microstructural evolution below the top surface, transmission electron microscopy (TEM) were employed to examine the treated (for 30 min) samples, as shown in Fig. 3. A typical bright-field TEM image is shown in Fig. 3(a), taken from about 60 mm below the top surface. It clearly exhibits that an isolated crystallite island emerged in the glassy matrix, inside which many crystalline particles distribute homogeneously, shown as the dark dots. Furthermore, bright-field TEM image and corresponding dark-field pair obviously display the homogeneous distribution of the nanocrystallites inside the isolated crystallite islands (Fig. 3(b, c)). Herein, grain sizes are about 5e30 nm. Meanwhile, the corresponding selected-area electron diffraction (SAED) pattern consists of the spotty diffraction rings with a diffuse background, as shown in the inset. It further confirms the partial crystallization behavior of the metallic glasses. And, Fig. 3(d) shows a typical HRTEM image with the clear lattice fringes in the glassy matrix. Fig. 3(e) shows the microstructural characteristics taken from about 80 mm below the top surface. It is essentially featureless and no change is discernible in the microstructure. SAED pattern is also included as the inset. Nearly, no local crystalline region could be found. Therefore, crystallization has been induced in the top surface of the metallic glasses by SMAT process. It created the isolated crystallite islands, dispersing in the glassy matrix. The crystallites inside the isolated crystallite islands display a decreasing trend in grain
1423
size along the depth until about 80 mm below the top surface. Grain size range is about 500e0 nm, showing a gradient change. Thus, microstructural characteristics provide the information about the structural evolution in the surface layer of the treated samples. 3.2. Mechanical properties optimization Compression tests reveal the optimized mechanical properties by SMAT process. Fig. 4(a) presents the room-temperature compression engineering stress-strain curves tested at the strain rate of 1 104 s1. Nearly, all the samples display the initial elastic deformation behavior with the equal elastic strain of about 1.86% and Young’s modulus of about 97 GPa, and then begin to yield at the yield strength of about 1.78 GPa. Especially, fracture strength is nearly equal for all the samples. However, plastic strain displays a distinctly incremental trend with the increase in time duration of SMAT process. It is 0.95%, 2.09%, 2.97% and 3.78% respectively for the samples treated for 0, 5, 15 and 30 min. Vividly, for the identical alloy composition, plastic strain is enhanced up to 400% by SMAT process for 30 min, comparing with that of the as-cast sample, nearly without inducing the loss in strength [22]. That is to say, this structural gradient metallic glass matrix composite has an
Fig. 4. Improved mechanical properties with the increase in the time duration of SMAT process: (a) the compressive engineering stress-strain curves tested at the strain rate of 1 10-4 s1, and inset showing the typical macroscopic fracture mode of treated samples with a global plastic deformation; (b) the statistical average results showing the enhanced plasticity and the preserved yield strength with the increase in time duration of SMAT process.
1424
J. Fan et al. / Intermetallics 19 (2011) 1420e1427
improved plasticity plus preserved high strength (Fig. 4(b)). It should be attributed to the suppressive effect of the isolated crystallite islands on the shear bands/cracks propagation, which can stimulate the global plastic deformation. The inset in Fig. 4(a) shows the typical fracture features of the treated samples in a shearing mode with the fracture angle of about 42 , which agrees well with the previous analysis [30,31]. Multiple shear bands are triggered on the top surface, forming a grid pattern, as shown by the arrows. 3.3. Deformation and fracture behaviors Unlike the monolithic metallic glasses, highly localized deformation can be overwhelmed by the sub-formed isolated crystallite islands, so that global plastic deformation is stimulated, which contributes to the obvious macroscopic plasticity. SEM observations show the restraining effect of the isolated crystallite islands on the shear bands/cracks propagation. Fig. 5 shows the fractographies recorded from the top surface of the treated (30 min) samples. Shear bands were observed to be effectively limited by the isolated crystallite islands, as shown by the arrows in Fig. 5(a). Clearly, when the shear band encounters the crystallite island, it is hindered and then deflects. Winding propagation is along the interface between the crystallite island and glassy matrix. Finally, cracking is triggered along the interface, due to the high stress concentration. Meanwhile, some short secondary shear bands nucleate in the neighboring region shown by the arrows in Fig. 5(a). Besides, it also shows that the limited shear band becomes curved and widened distinctly without any observable opening when striking against the crystallite island. The width is measured of about 1700 nm, which is far larger than the prior reports [32e35]. And, shearing deformation is also observed
inside the crystallite island to coordinate the overall shearing deformation. In short, shear bands deflection is induced by the isolated crystallite islands. Fig. 5(b) shows another kind of the fractographies of the treated (30 min) samples. Shear crack extension is observed to be effectively arrested by the isolated crystallite islands. These two cracks do not link together, and are separated by the middle crystallite island. It confirms that former crack is arrested by the crystallite island, and another new crack nucleates afterward. The new crack is along the same direction with the arrested crack, resulted from the effect of shear stress. Meanwhile, the cracking is also triggered inside the crystallite island, which suggests that crystallites occur to deform during arresting the shear crack. In brief, shear cracks arresting is caused by the isolated crystallite islands. Fig. 5(c) shows the third kind of the fractographies of the treated (30 min) samples. Shear band opening is observed to be effectively suppressed by the isolated crystallite islands. Two shear bands initiated from the crack tip, when they propagate inside the crystallite island. It convincingly shows that the crystalline phases have a suppressing effect on the shear band opening. Besides, the inset reveals a local shearing displacement of about 1 mm ahead of the crack tip, due to the restraining effect of the crystallites. It suggests that shear band opening is actually suppressed and local plastic strain is stimulated. Therefore, the suppression for shear band opening and the necessity for shearing deformation should contribute to the resistance to the brittle fracture and the enhancement in plasticity. In a word, shear bands suppression is resulted by the isolated crystallite islands. In addition, Fig. 5(d) shows the numerous of microcracks and fine shear bands inside the isolated crystallite islands. These microcracks are small and winding. And, the high-density shear bands are also found with a nanoscale spacing. These observations
Fig. 5. SEM images showing the deformation and fracture features of the treated (for 30 min) sample: (a) deflected shear band/crack by the isolated crystallite island, together with the initiation of neighboring secondary shear bands; (b) arrested shear crack by the isolated crystallite island, together with the initiation of the internal crack; (c) suppressed shear band opening and formed about 1 mm shearing displacement at the crack tip (see inset); (d) formation of profuse microcracks and shear bands.
J. Fan et al. / Intermetallics 19 (2011) 1420e1427
illuminate that the crystallite islands themselves also undergo the plastic deformation, during restraining the shear bands/cracks propagation. All in all, from the above observations on deformation and fracture behavior, isolated crystallite islands effectively limit shear bands extension, suppress shear bands opening and avoid crack developing. And, these effects not only activate the initiation of secondary shear bands in the neighboring glassy matrix, but also stimulate the shearing deformation and micro-cracking inside the isolated crystallite islands. Thus, five kinds of toughening behaviors are observed on the top surface of the treated samples, i.e. shear band deflection, shear crack arresting, shear band opening suppression, micro-cracking and fine-shearing.
1425
4. Discussion
deformation in the glassy matrix is in the manner of main shear bands and secondary shear bands. Although computer simulation shows the high shear plastic strain within an individual shear band [37], it is limiting for the contribution of main shear bands to the global plasticity. But, numerous secondary shear bands can make the large contribution to the enhancement in global plasticity, which are stimulated by the isolated crystallite islands. Besides, plastic deformation inside the isolated crystallite islands also has a contribution to the global plastic strain. Herein, plastic deformation is represented by the fine-shearing and micro-cracking. The high-density shear bands indicate a large plastic deformation, according to the relationship between plastic strain,3p and shear band spacing, d: d1 ¼ 0.143p [38]. And, microcracks would also contribute to the plastic deformation and micro-hardening, based on the prior study on size dependence of micro-toughness [39,40].
4.1. Improved plasticity
4.2. Fracture toughness
For the current treated samples, since partially crystallized layer (PCL) is much thinner than that of interior glassy matrix, the beginning elastic deformation and following yielding are mainly controlled by the glassy matrix. The slight change in microstructure of the top surface is not expected to induce the great change in glass transition temperature and molar volume. So, strength is not decreased [36]. But, plasticity is distinctly improved. In the current work, fracture observations show that plastic deformation occurs partially through a shear-banding mechanism in the glassy matrix (inset of Fig. 4(a), and Fig. 5(a)), and partially through fine-shearing and micro-cracking mechanism inside the isolated crystallite islands (Fig. 5(c, d)). Thus, plastic strain in the glassy matrix and isolated crystallite islands contributes to the global plasticity. Herein, plastic
In the composite community, there are two possible ways to toughen a brittle matrix: through crack bridging by adding a soft second phase with a strong interface or through crack deflecting by adding a strong and hard phase with a weak interface. Based on the principles, ductile phases and hard particles have been successfully prepared to form the metallic glass matrix composite [41e45]. For our current study, isolated crystallite islands effectively limit shear band extension (Fig. 5(a)), suppress shear band opening (Fig. 5(c)) and avoid crack developing (Fig. 5(b)). These toughening behaviors are in line with the basic toughening methods of metallic glasses [6e16]. Schematic illustrations are shown in Fig. 6(aec). Shear band deflects and widens forwards, and then bypasses the isolated crystallite island, finally opens into a crack, which shows that
Fig. 6. Schematic illumination of the toughening mechanisms induced by the restraining effect of the isolated crystallite islands on the shear bands/cracks propagation: (a) limiting shear band extension; (b) suppressing shear band opening; (c) avoiding crack developing; and corresponding theoretical illumination of the energetics accompanying the shear crack extension and the stress state on the crack tip: (e) elastic energy concentrated on the crack tip in the region with the radius, c, approximately equal to the crack half-length, a; (f) maximum stress occurring adjacent to the crack tip in a linearly elastic solid.
1426
J. Fan et al. / Intermetallics 19 (2011) 1420e1427
isolated crystallite islands limit the shear band extension (Fig. 6(a)). Moreover, non-opening shear bands extend with a micron-scale shearing displacement, initiating from the crack tip inside the crystallite island, which indicates that isolated crystallite islands suppress the shear band opening [Fig. 6(b)]. And, winding shear crack is arrested by the crystallite island forwards and re-nucleates afterward, which suggests that isolated crystallite islands avoid crack developing (Fig. 6(c)). These toughening behaviors contribute to preventing shear bands/cracks from free propagation and the enhancement in plasticity. Wherein, toughening mechanisms include: (1) increase in fracture surface due to the irregular path of shear cracks; (2) local shearing deformation in the glassy matrix around the isolated crystallite islands; (3) shear crack arresting and re-nucleating; and (4) fine-shearing and micro-cracking inside the isolated crystallite islands. As described above, these toughening mechanisms are relative to the effect of the isolated crystallite islands on the shear bands/cracks tip. The energetics accompanying the propagation of a shear crack is illuminated in Fig. 6(e). Two energies are associated with a shear crack extension. One is elastic strain energy (ypa2 ðs2nom =2EÞ per unit sheet thickness. Where, a is the half-length of shear crack; and snom and E are nominally applied stress and Young’s modulus respectively), which are associated with the stress concentration and strain presenting near the crack tip. It represents a driving force for shear crack propagation. Another one is surface-energy (¼4ag per unit sheet thickness, where, g is the energy per unit area (here exist two surfaces)), created as the shear crack propagates. It acts to resist shear crack advance. Surface-energy retarding force and elastic strain energy driving force for shear crack extension depend differently on the crack half-length, a. Thus, due to the restraining effect of the isolated crystallite islands, shear crack half-length, a is effectively decreased. With the increase in local stress, the width of shear cracks (i.e. the radius of curvature of crack tip, r) becomes larger (Fig. 5(b)). So, according to the relationship between maximum stress occurring immediately adjacent to the crack tip, smax and nominally applied stress, snom (¼F/A): smax y2s nom ða=rÞ
1=2
, the decreased crack half-length, a and the increased curvature radius, r will lead to the decrease in maximum stress, smax. It significantly results in the termination of shear crack extension, based on the fracture happening principle: smax s0 . Here, s0 is a constant, the theoretical strength of material. Schematic illumination is shown in Fig. 6(f). Based on the above theoretical analysis, shear bands/cracks would be effectively arrested by the isolated crystallite islands, which makes a toughening effect. Finally, when local stress rises up to high enough, shear bands/cracks have to favor the weak interface to deflect and bypass the isolated crystallite islands. Or, they are arrested forwards, and then re-nucleate afterward. Meanwhile, high local stress stimulates the shearing deformation in the neighboring glassy matrix and inside the crystallite islands, which can improve the global plasticity. Furthermore, since fracture toughness normally refers to the energy absorbed prior to fracture, which is proportional to the area underneath the stress-strain curve, it is also effectively enhanced. Based on elastic-plastic fracture mechanics theory, J-integral can be R H described as follows: J ¼ C F=Adu=dl0 ¼ sd3 [46,47]. Where, F is the force applied at the crack tip; A is the area of crack tip; du=dl0 is the change in energy per unit length; s is the stress; d3 is the change in strain caused by the stress. So, the improved plasticity and preserved high strength can significantly enhance the JIC. And, K-based fracture toughness can be back-calculated from the J-integral pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi using the KeJ equivalence, KIC ¼ EJIC =1 y2 . KIC is the fracture toughness; y is the Poisson’s ratio; E is the Young’s Modulus. Thus, the enhanced JIC indicates the improvement of fracture toughness, KIC. In
pffiffiffiffiffiffi addition, for stress intensity factor, KI, KI ¼ Y s pa. Where, Y is geometrical factor; a is crack half-length. Due to the restraining effect of the isolated crystallite islands on the shear crack developing, the decreased crack half-length, a can reduce the tress intensity factor, KI. While, the improved fracture toughness, KIC and reduced tress intensity factor, KI can induce the discontent with the fracture criteria: KI KIC , according to the Griffith theory. Therefore, isolated crystallite islands actually enhance the resistance to shear bands/ cracks developing and the fracture toughness. 5. Conclusions In this work, isolated crystallite islands are created to optimize the mechanical properties of metallic glasses via surface mechanical crystallization behavior induced by SMAT process. Furthermore, the improved plasticity and toughening mechanism are systematically evaluated. These results are summarized as follows: (1) Mechanical crystallization is induced in the top surface with about 80 mm thickness to form the isolated crystallite islands homogenously distributed in the glassy matrix. Inside the isolated crystallite islands, crystallites are well dispersed into the residual amorphous phase, displaying a microstructural gradient with the grain size range of about 500e0 nm along the depth. (2) Isolated crystallite islands, acting as the obstacles to restrict the highly localized deformation of shear bands/cracks, effectively limit shear band extension, suppress shear band opening and avoid crack developing. They stimulate the secondary shear bands in the neighboring glassy matrix, shearing deformation and micro-cracking inside the isolated crystallite islands, which contribute to the improvement in global plasticity. And, shear band deflection, shear crack pinning, shear band opening suppression, fine-shearing and micro-cracking are observed to contribute to the enhancement in fracture toughness. Acknowledgements The authors acknowledge the financial support by Research Grants Council of the Hong Kong Special Administrative Region of China under the PolyU 5203/08E. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]
Schuh CA, Hufnagel TC, Ramamurty U. Acta Mater 2007;55:4067. Wang WH. Adv Mater 2009;21:1. Zhang ZF, Wu FF, He G, Eckert J. J Mater Sci Tech 2007;23:747. Chen H, He Y, Shiflet GJ, Poon SJ. Nature 1994;367:541. Pampillo CA. J Mater Sci 1975;10:1194. Chen MW, Inoue A, Zhang W, Sakurai T. Phys Rev Lett 2006;96:245502. He G, Zhang ZF, Loser W, Eckert J, Schultz L. Acta Mater 2003;51:2383. Fan C, Ott RT, Hufnagel TC. Appl Phys Lett 2002;81:1020. Pan DG, Zhang HF, Wang AM, Hu ZQ. Appl Phys Lett 2006;89:261904. Fan JT, Zhang ZF, Mao SX, Shen BL, Inoue A. Intermetallics 2009;17:445. Hui X, Dong W, Chen GL, Yao KF. Acta Mater 2007;55:907. Shi LL, Ma H, Liu T, Xu J, Ma E. J Mater Res 2006;21:613. Qiao JW, Zhang Y, Liaw PK, Chen GL. Scr Mater 2009;61:1087. Hofmann DC, Suh JY, Wiest A, Duan G, Lind ML, Demetriou MD, et al. Nature 2008;451:1085. Kim KB, Das J, Baier F, Tang MB, Wang WH, Eckert J. Appl Phys Lett 2006;88: 051911. Inoue A, Zhang W, Tsurui T, Yavari AR, Greer AL. Philos Mag Lett 2005;85:221. Zhang Y, Wang WH, Greer AL. Nat Mater 2006;5:857. Qin FX, Wang XM, Xie GQ, Wada K, Song M, Furuya K, et al. Intermetallics 2009;17:945. Cao QP, Li JF, Zhou YH, Horsewell A, Jiang JZ. Acta Mater 2006;54:4373. Fan JT, Zhang ZF, Mao SX, Eckert J. Adv Eng Mater 2008;10:1117. Lu K, Lu J. J Mater Sci Tech 1999;15:193. Fan JT, Chen AY, Fu MW, Lu J. Scr Mater 2009;61:608. Kim JJ, Choi Y, Suresh S, Argon AS. Science 2002;295:654. He Y, Shiflet GJ, Poon SJ. Acta Metal Mater 1995;43:83.
J. Fan et al. / Intermetallics 19 (2011) 1420e1427 [25] EI-Eskandarany MS, Saida J, Inoue A. Acta Mater 2002;50:2725. [26] Fan JT, Zhang ZF, Mao SX, Shen BL, Inoue A. Philos Mag Lett 2008;88:837. [27] Tian JW, Shaw LL, Wang YD, Yokoyama Y, Liaw PK. Intermetallics 2009;17: 951. [28] Bae DH, Park JM, Na JH, Kim DH, Kim YC, Lee JK. J Mater Res 2004;19:937. [29] Tsai AP, Kamiyama T, Kawamura Y, Inoue A, Masumoto T. Acta Mater 1997; 45:1477. [30] Schuh CA, Lund AC. Nat Mater 2003;2:449. [31] Zhang ZF, Eckert J, Schultz L. Acta Mater 2003;51:1167. [32] Donovan PE, Stobbs WM. Acta Metal Mater 1981;29:1419. [33] Li J, Spaepen F, Hufnagel TC. Philos Mag A 2002;82:2623. [34] Georgarakis K, Aljerf M, Li Y, LeMoulec A, Charlot F, Yavari AR, et al. Appl Phys Lett 2008;93:031907.
[35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47]
1427
Zhang Y, Greer AL. Appl Phys Lett 2006;89:071907. Yang B, Liu CT, Nieh TG. Appl Phys Lett 2006;88:221911. Srolovitz D, Vitek V, Egami T. Acta Metal Mater 1983;31:335. Bei H, Xie S, George EP. Phys Rev Lett 2006;96:105503. Tvergaard V, Hutchinson JW. J Mech Phys Solids 1992;40:1377. Han CS, Gao HJ, Huang YG, Nix WD. J Mech Phys Solids 2005;53:1188. Liu T, Shen P, Qiu F, Yin ZF, Lin QL, Jiang QC, et al. Scr Mater 2009;60:84. Pauly S, Liu G, Wang G, Kuhn U, Mattern N, Eckert J. Acta Mater 2009;57:5445. Choi-Yim H, Schroers J, Johnson WL. Appl Phys Lett 2002;80:1906. Qiao JW, Wang S, Zhang Y, Liaw PK, Chen GL. Appl Phys Lett 2009;15:151905. Lu ZP, Ma D, Liu CT, Chang YA. Intermetallics 2007;15:253. Cherepanov GP. J Appl Math Mech 1967;31:503. Rice JR. J Appl Mech 1968;35:379.
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Improved plasticity and fracture toughness in metallic glasses via surface crystallization Jitang Fan a, b, *, Aiying Chen b, Juan Wang b, Jun Shen c, Jian Lu b, d, ** a
Faculty of Civil Engineering and Geosciences, Delft University of Technology, P.O. Box 5048, 2600 GA, Delft, The Netherlands Department of Mechanical Engineering, The Hong Kong Polytechnic University, Hong Kong, PR China c School of Materials Science and Engineering, Harbin Institute of Technology, Harbin 150001, PR China d Department of Mechanical Engineering, City University of Hong Kong, Hong Kong, PR China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 3 December 2009 Received in revised form 28 February 2011 Accepted 13 May 2011 Available online 16 June 2011
Mechanical crystallization was induced in the monolithic bulk metallic glasses by surface mechanical attrition treatment (SMAT) to create the isolated crystallite islands in the top surface layer. Inside the isolated crystallite islands, microstructure consists of the crystallites with a gradient grain size evolution and the residual amorphous phase. Moreover, isolated crystallite islands, acting as the obstacles to restrict the highly localized deformation of shear bands/cracks, effectively limit the shear bands extension, suppress the shear bands opening, and avoid the cracks developing, which significantly enhance the overall plasticity and fracture toughness. They were suggested by the secondary shear bands in the glassy matrix and fine-shearing together with micro-cracking inside the isolated crystallite islands. Finally, the improved plasticity and fracture toughness were systematically discussed. Based on the current results, surface crystallization is proposed to optimize the mechanical properties of metallic glasses. Ó 2011 Elsevier Ltd. All rights reserved.
Keywords: B. Glasses, metallic A. Composite, based on the metallic glass matrix intermetallics, miscellaneous B. Deformation map B. Fracture toughness B. Mechanical properties at ambient temperature
1. Introduction Bulk metallic glasses (BMGs) have inspired considerable scientific interest and industrial attention, due to their unique physical, chemical and mechanical properties [1,2]. Particularly, the characteristics of high strength and hardness are regarded as a potential candidate for the applications in structural engineering [3]. However, although computer simulation shows metallic glasses can display the superhigh plastic strain in the individual shear band on the microscopic scale [4,5], they, macroscopically, effectively lack the ductility with an apparently brittle fracture model. This is a fatal weakness to impede their application in the engineering field. Thus, improving general plasticity has been an urgent topic pursued in the metallic glass community for the recent several years. As pointed out by Chen et al. [6], variation in local material
* Corresponding author. Faculty of Civil Engineering and Geosciences, Delft University of Technology, P.O. Box 5048, 2600 GA, Delft, The Netherlands. ** Corresponding author. Department of Mechanical Engineering, The Hong Kong Polytechnic University, Hong Kong, PR China. E-mail addresses: [email protected] (J. Fan), [email protected] (J. Lu). 0966-9795/$ e see front matter Ó 2011 Elsevier Ltd. All rights reserved. doi:10.1016/j.intermet.2011.05.012
properties could arrest persistent slip on the individual shear band, in order to activate the multiple shear bands and contribute to the global plasticity. Generally, introducing a soft/plastic second phase is effective to improve the plasticity of metallic glasses [7e14], such as solid solution ductile particles [7e9], crystalline dendrite phases [10e13]. Especially, Zr-based metallic glass matrix composite exhibited the macroscopic room-temperature tensile ductility [14]. Besides, nanocrystallites are also profitable to improve the plasticity of metallic glasses [15,16]. In a word, the ‘composite’ of crystallites and metallic glasses was successful to optimize the mechanical properties. However, preparing the metallic glasses matrix composite usually has an acerbic demand for the glass forming ability and the casting technology. Thus, an interesting question is issued: whether there is a well-controlled, subsequent treating method to enhance the plasticity of BMGs. The correlative studies include shot-peening [17], coating [18], rolling [19] etc. Subsequently, a new strategy was proposed to improve the plasticity of BMGs, i.e. strengthening the surface of metallic glasses [20]. Since surface mechanical attrition treatment (SMAT) was proposed in 1999 [21], it has been proven to be an effective approach to optimize the surface microstructure and properties,
J. Fan et al. / Intermetallics 19 (2011) 1420e1427
Fig. 1. XRD diffraction patterns taken from the top surfaces of Ti40Zr25Ni3Cu12Be20 alloy after SMAT process with different treating time: 0, 5, 15 and 30 min.
enabling significantly to enhance the global performance of metallic materials [22]. In this work, we show a detailed study on the surface crystallization induced by SMAT process as well as the resultant improvement in mechanical properties. 2. Experimental procedures Alloy ingots with nominal chemical composition of Ti40Zr25Ni3Cu12Be20 (at.%) were prepared by arc melting the mixtures of elemental pieces with a purity above 99.9% in a Ti-gettered highly pure argon atmosphere. Remelting for five times was performed to ensure the homogeneity of composition. The final ingots were cast
1421
into a rectangular plate with a dimension of 3 30 70 mm3 using the copper mould. Surface mechanical attrition treatment (SMAT) is a subsequent treating approach to realize the specific structure and property requirements in the top surface layer [21]. In the current study, SMAT was performed by using high-hardness balls with a 2 mm diameter under a vibrating frequency of 20 kHz at roomtemperature. In order to avoid the temperature rise, treating process was done in a time interval of 5 min with a break time of 2 min. As well, cooling air flow was applied. Before treated, specimens were electric spark cut, and then were machined carefully into rectangle gauge dimension of 3 3 70 mm3. Four lateral surfaces of every specimen were treated for the same time duration. Treating time duration is 5 min, 15 min and 30 min respectively, which formed the four groups of samples in addition to the as-cast specimens. One typical specimen from each group was analyzed using Philips X-ray diffractometer (XRD) with Cu-Ka radiation as a source. And, microstructure analysis was performed for a typical specimen treated for 30 min by transmission electron microscopy (TEM JEM 2010) operated at 200 kV. TEM specimens were taken from the regions with a depth ranging to 90 mm below the top surface. The plane-view TEM foils were slowly ion-thinned at low temperature. Besides, statistical distribution of the grain size was quantitatively measured from the bright and dark-field TEM micrographs. Room-temperature compression tests were performed using a standard MTS Alliance RT/50 Materials Testing System, at a strain rate of 1 104 s1. For each group of samples, three specimens or more were tested with the dimension of 3 3 6 mm3, and the average data was used as the final results. Fracture morphologies, as well as the top surfaces of the treated samples, were observed by scanning electron microscopy (SEM Hitachi S-4200 field emission).
Fig. 2. SEM images taken from the top surface of the treated (for 30 min) sample: (a) homogeneous distribution of isolated crystallite islands on the top surface; (b) melting veinlike structure together with some crystalline particles; (c) amplified vein-like structure of region _ showing the interior nanocrystallites; (d) amplified crystallized particles of region II showing the homogeneous crystalline particles with the grain size of from about 200e500 nm.
1422
J. Fan et al. / Intermetallics 19 (2011) 1420e1427
Fig. 3. TEM and HRTEM images characterizing the crystallization behavior of the treated (for 30 min) sample: (a) bright-field image showing the isolated crystallite island taken from about 60 mm below the top surface; (b) bright-field image showing the homogeneous crystallites in the glassy matrix, and the corresponding SAED patterns in the inset; (c) corresponding dark-field pair; (d) HRTEM image showing the clear lattice fringes in the glassy matrix with the long-range ordering; (e) HRTEM image taken from about 80 mm below the top surface, showing the amorphous nature and the corresponding SAED pattern in the inset.
J. Fan et al. / Intermetallics 19 (2011) 1420e1427
3. Experimental results 3.1. Isolated crystallite islands formation Crystallization behavior in metallic glasses has been studied to be induced by annealing, plastic deformation, room-temperature fracture, and ball milling [23e27]. For the current work, mechanical crystallization was characterized to be induced by the SMAT process. XRD diffraction patterns recorded from the top surfaces are shown in Fig. 1. For the as-cast specimen, i.e. treated for 0 min, the pattern consists of only a broad diffraction maximum without any detectable sharp Bragg peaks, indicating the fully amorphous microstructure. And, for the treated specimens, it can be clearly seen that XRD patterns are composed of a broad diffusion background and a set of several sharp crystalline peaks, which suggest the mixtures of glass matrix and crystalline phases. Crystalline phases are identified as the Cu2TiZr ternary compound, and here maybe exists some incomplete nuclei [28,29]. In addition, the morphology of the treated (for 30 min) samples was revealed from the top surface by high-resolution SEM, as shown in Fig. 2. A typical surface roughness was induced with a well distribution in the top surface (Fig. 2(a)). Wherein, it is interesting to find that some crystalline phases congregate, named as isolated crystallite island, as pointed out by the arrows. Isolated crystallite islands disperse in the rough surface with the size of about 10e100 mm. Local amplified image displays the coexistence of bugle-like patterns and crystalline particles inside the isolated crystallite island, as shown in regions I and II respectively (Fig. 2(b)). This is very similar to the nanocrystallization behavior of Co- and Fe-based BMGs, induced by the quasi-static compression fracture at room-temperature [26]. Moreover, HRSEM image of region I clearly shows the bugle-like patterns, which consist of the elongated viscous liquid and interior nanocrystallites (Fig. 2(c)), implying the local plastic deformation and crystallization behavior during SMAT process. The same, HRSEM image of region II, seemingly overlaying the region I, shows the crystallites which are uniform with the grain size of about 200e500 nm, recognized as submicron crystals (Fig. 2(d)). The formation of the viscous liquid and crystalline particles in the top surface should be attributed to the severe plastic deformation during the SMAT process. To further shed light on the microstructural evolution below the top surface, transmission electron microscopy (TEM) were employed to examine the treated (for 30 min) samples, as shown in Fig. 3. A typical bright-field TEM image is shown in Fig. 3(a), taken from about 60 mm below the top surface. It clearly exhibits that an isolated crystallite island emerged in the glassy matrix, inside which many crystalline particles distribute homogeneously, shown as the dark dots. Furthermore, bright-field TEM image and corresponding dark-field pair obviously display the homogeneous distribution of the nanocrystallites inside the isolated crystallite islands (Fig. 3(b, c)). Herein, grain sizes are about 5e30 nm. Meanwhile, the corresponding selected-area electron diffraction (SAED) pattern consists of the spotty diffraction rings with a diffuse background, as shown in the inset. It further confirms the partial crystallization behavior of the metallic glasses. And, Fig. 3(d) shows a typical HRTEM image with the clear lattice fringes in the glassy matrix. Fig. 3(e) shows the microstructural characteristics taken from about 80 mm below the top surface. It is essentially featureless and no change is discernible in the microstructure. SAED pattern is also included as the inset. Nearly, no local crystalline region could be found. Therefore, crystallization has been induced in the top surface of the metallic glasses by SMAT process. It created the isolated crystallite islands, dispersing in the glassy matrix. The crystallites inside the isolated crystallite islands display a decreasing trend in grain
1423
size along the depth until about 80 mm below the top surface. Grain size range is about 500e0 nm, showing a gradient change. Thus, microstructural characteristics provide the information about the structural evolution in the surface layer of the treated samples. 3.2. Mechanical properties optimization Compression tests reveal the optimized mechanical properties by SMAT process. Fig. 4(a) presents the room-temperature compression engineering stress-strain curves tested at the strain rate of 1 104 s1. Nearly, all the samples display the initial elastic deformation behavior with the equal elastic strain of about 1.86% and Young’s modulus of about 97 GPa, and then begin to yield at the yield strength of about 1.78 GPa. Especially, fracture strength is nearly equal for all the samples. However, plastic strain displays a distinctly incremental trend with the increase in time duration of SMAT process. It is 0.95%, 2.09%, 2.97% and 3.78% respectively for the samples treated for 0, 5, 15 and 30 min. Vividly, for the identical alloy composition, plastic strain is enhanced up to 400% by SMAT process for 30 min, comparing with that of the as-cast sample, nearly without inducing the loss in strength [22]. That is to say, this structural gradient metallic glass matrix composite has an
Fig. 4. Improved mechanical properties with the increase in the time duration of SMAT process: (a) the compressive engineering stress-strain curves tested at the strain rate of 1 10-4 s1, and inset showing the typical macroscopic fracture mode of treated samples with a global plastic deformation; (b) the statistical average results showing the enhanced plasticity and the preserved yield strength with the increase in time duration of SMAT process.
1424
J. Fan et al. / Intermetallics 19 (2011) 1420e1427
improved plasticity plus preserved high strength (Fig. 4(b)). It should be attributed to the suppressive effect of the isolated crystallite islands on the shear bands/cracks propagation, which can stimulate the global plastic deformation. The inset in Fig. 4(a) shows the typical fracture features of the treated samples in a shearing mode with the fracture angle of about 42 , which agrees well with the previous analysis [30,31]. Multiple shear bands are triggered on the top surface, forming a grid pattern, as shown by the arrows. 3.3. Deformation and fracture behaviors Unlike the monolithic metallic glasses, highly localized deformation can be overwhelmed by the sub-formed isolated crystallite islands, so that global plastic deformation is stimulated, which contributes to the obvious macroscopic plasticity. SEM observations show the restraining effect of the isolated crystallite islands on the shear bands/cracks propagation. Fig. 5 shows the fractographies recorded from the top surface of the treated (30 min) samples. Shear bands were observed to be effectively limited by the isolated crystallite islands, as shown by the arrows in Fig. 5(a). Clearly, when the shear band encounters the crystallite island, it is hindered and then deflects. Winding propagation is along the interface between the crystallite island and glassy matrix. Finally, cracking is triggered along the interface, due to the high stress concentration. Meanwhile, some short secondary shear bands nucleate in the neighboring region shown by the arrows in Fig. 5(a). Besides, it also shows that the limited shear band becomes curved and widened distinctly without any observable opening when striking against the crystallite island. The width is measured of about 1700 nm, which is far larger than the prior reports [32e35]. And, shearing deformation is also observed
inside the crystallite island to coordinate the overall shearing deformation. In short, shear bands deflection is induced by the isolated crystallite islands. Fig. 5(b) shows another kind of the fractographies of the treated (30 min) samples. Shear crack extension is observed to be effectively arrested by the isolated crystallite islands. These two cracks do not link together, and are separated by the middle crystallite island. It confirms that former crack is arrested by the crystallite island, and another new crack nucleates afterward. The new crack is along the same direction with the arrested crack, resulted from the effect of shear stress. Meanwhile, the cracking is also triggered inside the crystallite island, which suggests that crystallites occur to deform during arresting the shear crack. In brief, shear cracks arresting is caused by the isolated crystallite islands. Fig. 5(c) shows the third kind of the fractographies of the treated (30 min) samples. Shear band opening is observed to be effectively suppressed by the isolated crystallite islands. Two shear bands initiated from the crack tip, when they propagate inside the crystallite island. It convincingly shows that the crystalline phases have a suppressing effect on the shear band opening. Besides, the inset reveals a local shearing displacement of about 1 mm ahead of the crack tip, due to the restraining effect of the crystallites. It suggests that shear band opening is actually suppressed and local plastic strain is stimulated. Therefore, the suppression for shear band opening and the necessity for shearing deformation should contribute to the resistance to the brittle fracture and the enhancement in plasticity. In a word, shear bands suppression is resulted by the isolated crystallite islands. In addition, Fig. 5(d) shows the numerous of microcracks and fine shear bands inside the isolated crystallite islands. These microcracks are small and winding. And, the high-density shear bands are also found with a nanoscale spacing. These observations
Fig. 5. SEM images showing the deformation and fracture features of the treated (for 30 min) sample: (a) deflected shear band/crack by the isolated crystallite island, together with the initiation of neighboring secondary shear bands; (b) arrested shear crack by the isolated crystallite island, together with the initiation of the internal crack; (c) suppressed shear band opening and formed about 1 mm shearing displacement at the crack tip (see inset); (d) formation of profuse microcracks and shear bands.
J. Fan et al. / Intermetallics 19 (2011) 1420e1427
illuminate that the crystallite islands themselves also undergo the plastic deformation, during restraining the shear bands/cracks propagation. All in all, from the above observations on deformation and fracture behavior, isolated crystallite islands effectively limit shear bands extension, suppress shear bands opening and avoid crack developing. And, these effects not only activate the initiation of secondary shear bands in the neighboring glassy matrix, but also stimulate the shearing deformation and micro-cracking inside the isolated crystallite islands. Thus, five kinds of toughening behaviors are observed on the top surface of the treated samples, i.e. shear band deflection, shear crack arresting, shear band opening suppression, micro-cracking and fine-shearing.
1425
4. Discussion
deformation in the glassy matrix is in the manner of main shear bands and secondary shear bands. Although computer simulation shows the high shear plastic strain within an individual shear band [37], it is limiting for the contribution of main shear bands to the global plasticity. But, numerous secondary shear bands can make the large contribution to the enhancement in global plasticity, which are stimulated by the isolated crystallite islands. Besides, plastic deformation inside the isolated crystallite islands also has a contribution to the global plastic strain. Herein, plastic deformation is represented by the fine-shearing and micro-cracking. The high-density shear bands indicate a large plastic deformation, according to the relationship between plastic strain,3p and shear band spacing, d: d1 ¼ 0.143p [38]. And, microcracks would also contribute to the plastic deformation and micro-hardening, based on the prior study on size dependence of micro-toughness [39,40].
4.1. Improved plasticity
4.2. Fracture toughness
For the current treated samples, since partially crystallized layer (PCL) is much thinner than that of interior glassy matrix, the beginning elastic deformation and following yielding are mainly controlled by the glassy matrix. The slight change in microstructure of the top surface is not expected to induce the great change in glass transition temperature and molar volume. So, strength is not decreased [36]. But, plasticity is distinctly improved. In the current work, fracture observations show that plastic deformation occurs partially through a shear-banding mechanism in the glassy matrix (inset of Fig. 4(a), and Fig. 5(a)), and partially through fine-shearing and micro-cracking mechanism inside the isolated crystallite islands (Fig. 5(c, d)). Thus, plastic strain in the glassy matrix and isolated crystallite islands contributes to the global plasticity. Herein, plastic
In the composite community, there are two possible ways to toughen a brittle matrix: through crack bridging by adding a soft second phase with a strong interface or through crack deflecting by adding a strong and hard phase with a weak interface. Based on the principles, ductile phases and hard particles have been successfully prepared to form the metallic glass matrix composite [41e45]. For our current study, isolated crystallite islands effectively limit shear band extension (Fig. 5(a)), suppress shear band opening (Fig. 5(c)) and avoid crack developing (Fig. 5(b)). These toughening behaviors are in line with the basic toughening methods of metallic glasses [6e16]. Schematic illustrations are shown in Fig. 6(aec). Shear band deflects and widens forwards, and then bypasses the isolated crystallite island, finally opens into a crack, which shows that
Fig. 6. Schematic illumination of the toughening mechanisms induced by the restraining effect of the isolated crystallite islands on the shear bands/cracks propagation: (a) limiting shear band extension; (b) suppressing shear band opening; (c) avoiding crack developing; and corresponding theoretical illumination of the energetics accompanying the shear crack extension and the stress state on the crack tip: (e) elastic energy concentrated on the crack tip in the region with the radius, c, approximately equal to the crack half-length, a; (f) maximum stress occurring adjacent to the crack tip in a linearly elastic solid.
1426
J. Fan et al. / Intermetallics 19 (2011) 1420e1427
isolated crystallite islands limit the shear band extension (Fig. 6(a)). Moreover, non-opening shear bands extend with a micron-scale shearing displacement, initiating from the crack tip inside the crystallite island, which indicates that isolated crystallite islands suppress the shear band opening [Fig. 6(b)]. And, winding shear crack is arrested by the crystallite island forwards and re-nucleates afterward, which suggests that isolated crystallite islands avoid crack developing (Fig. 6(c)). These toughening behaviors contribute to preventing shear bands/cracks from free propagation and the enhancement in plasticity. Wherein, toughening mechanisms include: (1) increase in fracture surface due to the irregular path of shear cracks; (2) local shearing deformation in the glassy matrix around the isolated crystallite islands; (3) shear crack arresting and re-nucleating; and (4) fine-shearing and micro-cracking inside the isolated crystallite islands. As described above, these toughening mechanisms are relative to the effect of the isolated crystallite islands on the shear bands/cracks tip. The energetics accompanying the propagation of a shear crack is illuminated in Fig. 6(e). Two energies are associated with a shear crack extension. One is elastic strain energy (ypa2 ðs2nom =2EÞ per unit sheet thickness. Where, a is the half-length of shear crack; and snom and E are nominally applied stress and Young’s modulus respectively), which are associated with the stress concentration and strain presenting near the crack tip. It represents a driving force for shear crack propagation. Another one is surface-energy (¼4ag per unit sheet thickness, where, g is the energy per unit area (here exist two surfaces)), created as the shear crack propagates. It acts to resist shear crack advance. Surface-energy retarding force and elastic strain energy driving force for shear crack extension depend differently on the crack half-length, a. Thus, due to the restraining effect of the isolated crystallite islands, shear crack half-length, a is effectively decreased. With the increase in local stress, the width of shear cracks (i.e. the radius of curvature of crack tip, r) becomes larger (Fig. 5(b)). So, according to the relationship between maximum stress occurring immediately adjacent to the crack tip, smax and nominally applied stress, snom (¼F/A): smax y2s nom ða=rÞ
1=2
, the decreased crack half-length, a and the increased curvature radius, r will lead to the decrease in maximum stress, smax. It significantly results in the termination of shear crack extension, based on the fracture happening principle: smax s0 . Here, s0 is a constant, the theoretical strength of material. Schematic illumination is shown in Fig. 6(f). Based on the above theoretical analysis, shear bands/cracks would be effectively arrested by the isolated crystallite islands, which makes a toughening effect. Finally, when local stress rises up to high enough, shear bands/cracks have to favor the weak interface to deflect and bypass the isolated crystallite islands. Or, they are arrested forwards, and then re-nucleate afterward. Meanwhile, high local stress stimulates the shearing deformation in the neighboring glassy matrix and inside the crystallite islands, which can improve the global plasticity. Furthermore, since fracture toughness normally refers to the energy absorbed prior to fracture, which is proportional to the area underneath the stress-strain curve, it is also effectively enhanced. Based on elastic-plastic fracture mechanics theory, J-integral can be R H described as follows: J ¼ C F=Adu=dl0 ¼ sd3 [46,47]. Where, F is the force applied at the crack tip; A is the area of crack tip; du=dl0 is the change in energy per unit length; s is the stress; d3 is the change in strain caused by the stress. So, the improved plasticity and preserved high strength can significantly enhance the JIC. And, K-based fracture toughness can be back-calculated from the J-integral pffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi using the KeJ equivalence, KIC ¼ EJIC =1 y2 . KIC is the fracture toughness; y is the Poisson’s ratio; E is the Young’s Modulus. Thus, the enhanced JIC indicates the improvement of fracture toughness, KIC. In
pffiffiffiffiffiffi addition, for stress intensity factor, KI, KI ¼ Y s pa. Where, Y is geometrical factor; a is crack half-length. Due to the restraining effect of the isolated crystallite islands on the shear crack developing, the decreased crack half-length, a can reduce the tress intensity factor, KI. While, the improved fracture toughness, KIC and reduced tress intensity factor, KI can induce the discontent with the fracture criteria: KI KIC , according to the Griffith theory. Therefore, isolated crystallite islands actually enhance the resistance to shear bands/ cracks developing and the fracture toughness. 5. Conclusions In this work, isolated crystallite islands are created to optimize the mechanical properties of metallic glasses via surface mechanical crystallization behavior induced by SMAT process. Furthermore, the improved plasticity and toughening mechanism are systematically evaluated. These results are summarized as follows: (1) Mechanical crystallization is induced in the top surface with about 80 mm thickness to form the isolated crystallite islands homogenously distributed in the glassy matrix. Inside the isolated crystallite islands, crystallites are well dispersed into the residual amorphous phase, displaying a microstructural gradient with the grain size range of about 500e0 nm along the depth. (2) Isolated crystallite islands, acting as the obstacles to restrict the highly localized deformation of shear bands/cracks, effectively limit shear band extension, suppress shear band opening and avoid crack developing. They stimulate the secondary shear bands in the neighboring glassy matrix, shearing deformation and micro-cracking inside the isolated crystallite islands, which contribute to the improvement in global plasticity. And, shear band deflection, shear crack pinning, shear band opening suppression, fine-shearing and micro-cracking are observed to contribute to the enhancement in fracture toughness. Acknowledgements The authors acknowledge the financial support by Research Grants Council of the Hong Kong Special Administrative Region of China under the PolyU 5203/08E. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24]
Schuh CA, Hufnagel TC, Ramamurty U. Acta Mater 2007;55:4067. Wang WH. Adv Mater 2009;21:1. Zhang ZF, Wu FF, He G, Eckert J. J Mater Sci Tech 2007;23:747. Chen H, He Y, Shiflet GJ, Poon SJ. Nature 1994;367:541. Pampillo CA. J Mater Sci 1975;10:1194. Chen MW, Inoue A, Zhang W, Sakurai T. Phys Rev Lett 2006;96:245502. He G, Zhang ZF, Loser W, Eckert J, Schultz L. Acta Mater 2003;51:2383. Fan C, Ott RT, Hufnagel TC. Appl Phys Lett 2002;81:1020. Pan DG, Zhang HF, Wang AM, Hu ZQ. Appl Phys Lett 2006;89:261904. Fan JT, Zhang ZF, Mao SX, Shen BL, Inoue A. Intermetallics 2009;17:445. Hui X, Dong W, Chen GL, Yao KF. Acta Mater 2007;55:907. Shi LL, Ma H, Liu T, Xu J, Ma E. J Mater Res 2006;21:613. Qiao JW, Zhang Y, Liaw PK, Chen GL. Scr Mater 2009;61:1087. Hofmann DC, Suh JY, Wiest A, Duan G, Lind ML, Demetriou MD, et al. Nature 2008;451:1085. Kim KB, Das J, Baier F, Tang MB, Wang WH, Eckert J. Appl Phys Lett 2006;88: 051911. Inoue A, Zhang W, Tsurui T, Yavari AR, Greer AL. Philos Mag Lett 2005;85:221. Zhang Y, Wang WH, Greer AL. Nat Mater 2006;5:857. Qin FX, Wang XM, Xie GQ, Wada K, Song M, Furuya K, et al. Intermetallics 2009;17:945. Cao QP, Li JF, Zhou YH, Horsewell A, Jiang JZ. Acta Mater 2006;54:4373. Fan JT, Zhang ZF, Mao SX, Eckert J. Adv Eng Mater 2008;10:1117. Lu K, Lu J. J Mater Sci Tech 1999;15:193. Fan JT, Chen AY, Fu MW, Lu J. Scr Mater 2009;61:608. Kim JJ, Choi Y, Suresh S, Argon AS. Science 2002;295:654. He Y, Shiflet GJ, Poon SJ. Acta Metal Mater 1995;43:83.
J. Fan et al. / Intermetallics 19 (2011) 1420e1427 [25] EI-Eskandarany MS, Saida J, Inoue A. Acta Mater 2002;50:2725. [26] Fan JT, Zhang ZF, Mao SX, Shen BL, Inoue A. Philos Mag Lett 2008;88:837. [27] Tian JW, Shaw LL, Wang YD, Yokoyama Y, Liaw PK. Intermetallics 2009;17: 951. [28] Bae DH, Park JM, Na JH, Kim DH, Kim YC, Lee JK. J Mater Res 2004;19:937. [29] Tsai AP, Kamiyama T, Kawamura Y, Inoue A, Masumoto T. Acta Mater 1997; 45:1477. [30] Schuh CA, Lund AC. Nat Mater 2003;2:449. [31] Zhang ZF, Eckert J, Schultz L. Acta Mater 2003;51:1167. [32] Donovan PE, Stobbs WM. Acta Metal Mater 1981;29:1419. [33] Li J, Spaepen F, Hufnagel TC. Philos Mag A 2002;82:2623. [34] Georgarakis K, Aljerf M, Li Y, LeMoulec A, Charlot F, Yavari AR, et al. Appl Phys Lett 2008;93:031907.
[35] [36] [37] [38] [39] [40] [41] [42] [43] [44] [45] [46] [47]
1427
Zhang Y, Greer AL. Appl Phys Lett 2006;89:071907. Yang B, Liu CT, Nieh TG. Appl Phys Lett 2006;88:221911. Srolovitz D, Vitek V, Egami T. Acta Metal Mater 1983;31:335. Bei H, Xie S, George EP. Phys Rev Lett 2006;96:105503. Tvergaard V, Hutchinson JW. J Mech Phys Solids 1992;40:1377. Han CS, Gao HJ, Huang YG, Nix WD. J Mech Phys Solids 2005;53:1188. Liu T, Shen P, Qiu F, Yin ZF, Lin QL, Jiang QC, et al. Scr Mater 2009;60:84. Pauly S, Liu G, Wang G, Kuhn U, Mattern N, Eckert J. Acta Mater 2009;57:5445. Choi-Yim H, Schroers J, Johnson WL. Appl Phys Lett 2002;80:1906. Qiao JW, Wang S, Zhang Y, Liaw PK, Chen GL. Appl Phys Lett 2009;15:151905. Lu ZP, Ma D, Liu CT, Chang YA. Intermetallics 2007;15:253. Cherepanov GP. J Appl Math Mech 1967;31:503. Rice JR. J Appl Mech 1968;35:379.