- Email: [email protected]
Torsional behaviors in Zr-based bulk metallic glass with high stored energy structure
Materials Science & Engineering A 751 (2019) 128–132
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Torsional behaviors in Zr-based bulk metallic glass with high stored energy structure Yu-Bai Maa, Yan Jiangb, Hua-Ping Dingc, Qi-Dong Zhanga, Xiao-Yun Lia, Fang-Qiu Zua,
T
⁎
a
Liquid/Solid Metal Processing Institute, School of Materials Science & Engineering, Hefei University of Technology, Hefei 230009, China Panzhihua International Research Institute of Vanadium and Titanium, Panzhihua 617000,China c State Key Laboratory of Materials Processing and Die & Mould Technology, Huazhong University of Science and Technology, 1037 Luoyu Road, 430074 Wuhan, Hubei, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Bulk metallic glasses Torsional plasticity Stored energy Heterogeneity
To overcome brittleness of bulk metallic glasses, great efforts have been devoted to optimizing BMGs′ plasticity under compressive, tensional or bending load. Yet, it is hardly found any work done to enhance torsional plasticity of BMGs. Herein, based on the different stored energy of BMGs determining diverse properties, we introduce the energy into Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 BMG by a high rheological rate forming method, leading to the higher microstructural heterogeneity, which is conducive to increase not only the plasticity (from 0.2% up to 1.56%) but also the ultimate strength from (990 MPa up to 1110 MPa) on the torsional deformation effectively. This work provides a new orientation to break through the brittleness of BMGs.
1. Introduction
BMGs can be diverse at different preparation conditions [17], while its internal relation with the torsional plasticity needs to further explore still. Here, in this paper, the high rheological rate forming (HRRF) technique [18] is applied to introduce energy into Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 BMG to illuminate this issue. The result shows that the preserved higher energy in this brittle BMG can cause the higher microstructural heterogeneity, resulting in enhancing both the ductility and ultimate strength on the torsion.
Bulk metallic glasses (BMGs) hold a variety of unique properties due to the amorphous state of the atoms, which have been an area of intense research promptly for several decades [1,2]. Nevertheless, this peculiar microstructure has brought the brittleness to BMGs at room temperature [3]. In fact, the plasticity of this material generates in the narrow shear bands of only 10–20 nm width [4,5]. Thus, when the shear bands take place and cross the sample [6], the whole deformation becomes instability and fractures immediately. To address this challenge, many researchers have investigated the possibilities to improve the plasticity including various intrinsic and extrinsic ways [7–12]. However, most of these studies have paid close attention to the mechanical behaviors on compressive [7,8], bending [9,10] and tensile conditions [11,12]. Up to now, to our knowledge, no researches have been carried out on improving the torsional behaviors of BMGs. Actually, as a candidate material for structural applications, the torsional properties can process great significances, such as torsion bars in scanners, transmission shafts, screws, etc. Moreover, similar to the tension, the torsional deformation also occurs in the unconstrained stress field and the fracture of BMGs is always controlled by a single dominant shear band. Thus, most of BMGs show a limited plastic strain (~0.2%) during the torsional deformation [13–16]. Therefore, how to improve the torsional behaviors of BMGs is a significant issue to be solved. Furthermore, it is understandable that the stored energy of ⁎
2. Experimental Amorphous Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 rods with a diameter of 6 mm, and a length of 120 mm were prepared using copper mold casting. Then cut them into 27 mm long cylinders by a cutting machine to fabricate the treated cylinder with 4.4 mm in diameter and a length of 40 mm through HRRF [18]. A test sample was processed into the hexagonal prisms with a gauge length of L = ~9 mm and diameter of d = ~1.8 mm joining with a radius to the long sample heads. The hexagonal heads were designed to prevent slipping while twisting, and the gauge sections of samples were polished with diamond paste (inset of Fig. 3). Torsional tests were performed on the Instron 8874 testing machine at a constant rotation rate of 0.057°/s with careful alignment between the specimen and the rotation axis. One hexagonal head of the sample was fixed on the free end of the testing machine to avoid
Corresponding author. E-mail address: [email protected] (F.-Q. Zu).
https://doi.org/10.1016/j.msea.2019.02.074 Received 10 January 2019; Received in revised form 21 February 2019; Accepted 22 February 2019 Available online 23 February 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 751 (2019) 128–132
Y.-B. Ma, et al.
channel and subsequently filled into the mold cavity. According to the displacement of plunger recorded by a high-speed camera, the whole experiment only lasted ~25 ms, which was enough to avoid thermal embrittlement. 3. Results Fig. 2(a) shows the X-ray diffraction patterns obtained from Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 BMG in the as-cast and treated states respectively. The two patterns exhibit the same broad diffusion maxima and conform to a characteristic of the full amorphous structure. The DSC curves of them (in Fig. 2(b)) have presented the sequential transition of glass transition and supercooled liquid region. Although the onset crystallization temperature (Tx) has almost no change, it is obvious seen that the glass transition temperature (Tg) increases from about 684–694 K (see Table (1)), and the area of an exothermic peak before reaching the Tg in DSC which represents the structural relaxation enthalpy ΔH rises from ~5.432 to ~11.54 J/g dramatically (see the black shaded areas in inset of Fig. 2(b)) with the change from as-cast to treated state. The high structural relaxation enthalpy could be regarded as proofs for much more multiple free volume concentration in treated BMGs because of a positive proportional relationship between free volume change [19]. Additionally, the onset of structural relaxation temperature of as-cast specimens began at ~518 K, while treated specimens began ~503 K. These phenomena indicate the energy system of the metallic glass has increased into a high-level energy state by HRRF [20]. Fig. 3 displays the typical shear-strain curve of Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 BMG under the torsional deformation for as-cast and treated specimens. From these curves, the as-cast specimen fails in a brittle manner with only a small nonlinear trend (plastic strain ~0.2%) before the final catastrophic fracture. However, for the treated sample, it is noticeable that the curve deviates from linearity at the high stress (~880 MPa) and displays the increase of the stress with a further increase in strain after yielding, which is most probably a consequence of noticeable interactions between the shear bands [21]. The macroscopic plastic strain of the treated sample increases into ~1.56% which is ~seven times than that of the as-cast sample. While the torsional ultimate strength grows up to ~1110 MPa, which is higher than that of as-cast one (~990 MPa), (The calculated yielding strengths σy, ultimate strengths σf, and the plastic strain εp were summarized in Table (1)) indicating that the HRRF method has a prominent effect on improving the torsional property of the metallic glass. In order to reflect the macroscopic plastic deformation intuitively, a suitable ink line which is parallel to the torsion axis has been marked on the lateral surface of the sample. After fracture, the variation of the marking line slope (θ) is showed in Fig. 4(a) and (b). By carefully measuring the slope, the ink lines of the as-cast sample is almost
Fig. 1. Schematic drawing of the HRRF method.
uniaxial stresses. At least three samples were measured to ensure the accuracy of the results. The nanoindentation experiment was conducted by an Agilent G200 Nanoindenter with an indenter of Diamond Berkovich at room temperature. Eighty-one indentations were made in a 9 × 9 matrix with the same depth (300 nm) at a constant loading rate of 0.05 s−1, and the spacing between adjacent indentations was set as 6 µm to avoid overlapping of neighboring strained zones. The glassy nature of BMG samples was confirmed by x-ray diffraction (XRD) using a BRUKER D8 ADVANCE diffractometer with Cu Kα radiation source and differential scanning calorimetry (DSC) performing under a purified argon atmosphere in a NETZSCH 449F3. The field emission scanning electron microscope (FESEM, SU8020) was used to observe the lateral and fracture surface morphology of samples. The thin samples were prepared by Ar ion-milling under liquid nitrogen cooling (Gatan Model 600) at 2.5 keV to investigate the detailed microstructures using the transmission electron microscope (TEM; JEM2100F). As mentioned above, to obtain high stored energy BMGs, we adopt a high rheological rate forming process (HRRF) to treat the as-cast samples. The device schematic diagram of the HRRF method consists of a power supply, a plunger, an arrow channel and a mold cavity (see in Fig. 1). A high force of (~6 kN) was pre-loaded on the plunger ahead of time. Then the as-cast BMG was heated to the supercooled liquid state (~745 k) measured by a nanovoltmeter applying the method of fast Joule heating. Therefore, the liquid BMG was squeezed into the narrow
Fig. 2. (a) The XRD patterns and (b) DSC traces of the as-cast and treated samples. Inserted graph of (b) shows as-cast and treated samples marked by black dashed rectangular. 129
Materials Science & Engineering A 751 (2019) 128–132
Y.-B. Ma, et al.
Table 1 Summary of the thermal parameters and torsional properties of the Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 BMG with as-cast and treated. Alloy
Tg (K)
Tx (K)
ΔH (J/g)
εp (%)
σy (MPa)
σf (MPa)
As-cast Treated
684 ± 3 694 ± 3
761 ± 3 758 ± 3
5.43 ± 0.05 11.54 ± 0.05
0.20 ± 0.05 1.56 ± 0.05
930 ± 13 880 ± 13
990 ± 10 1110 ± 10
which violates the Tresca criterion. It indicates that in addition to the shear stress, there is some normal stress applying on the sample. This normal stress is calculated by σy·sin(π −2δ )= 91 MPa, which is ~8% of the applied shear stress. Such a little normal stress can be explained by the secondary shear bands. Secondary shear bands may result in the significant difference between values of critical strain accommodation along perpendicular and a parallel direction with respect to the shear band planes. This physical separation is very similar to the normal stress [22]. Furthermore, the red-lined rectangular in Fig. 4(a) and (b), which are the side surface morphology near the fracture of as-cast and treated specimens, are magnified in Fig. 4(c) and (d) respectively. As illustrated in Fig. 4(d), multiple and intersecting shear bands including primary and secondary shear-bands (white arrows in Fig. 4(d)) are clearly observed on the surface of the treated sample which is considered one of the main evidence that account for the extent of plasticity in glass metals. The primary shear-bands are almost parallel to the shear stress direction. While the secondary shear-bands, which are initiated from the intersection points [23], propagate in disorder orientations. This indicates that there is not a preferential slip plane for secondary ones in BMGs [24]. In turn, almost no shear band is visible on the surface for the as-cast sample in Fig. 4(c), subjected to only a single dominant shear band cross the whole sample during the instantaneous fracture [25].
Fig. 3. The shear stressstrain curves obtained on as-cast and treated samples under torsional loading. The inset shows a pre-treated BMG rod, a treated BMG sample and a typical torsion sample, respectively.
parallel to the torsional axis and the slope θ is just about 0.2° (manifested by the yellow arrows) indicating that only a little plastic deformation happened during the torsion. However, the treated sample presents the larger slope (θ = 1.6°) than as-cast one, which is also obvious with its enhanced plasticity. According to the Tresca criterion, a fracture angle (δ) which is between the torsional fracture surface and the axial axis of the specimen should be 90°. For the as-cast one, its fracture angle is equal to ~90° (marked by the red arrow). While, the treated sample is just ~87°,
Fig. 4. (a) and (b) are the global views of the side surface fractured samples of the as-cast and treated with an ink line, respectively. (c) and (d) High-magnification detailed micrographs corresponding to the part marked by the red-lined rectangular in (a) and (b), respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.) 130
Materials Science & Engineering A 751 (2019) 128–132
Y.-B. Ma, et al.
Fig. 5. (a) Spatial distribution of the nano-hardness values of the Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 BMGs before and after HRRF. (b) and (c) the typical bright-field TEM images with insert diffraction pattern images of the as-cast and treated samples respectively.
7.66 GPa. Further to verify the HRRF imposing the microstructural heterogeneity of BMGs, TEM was carried out to clarify the differences of microstructures between as-cast and treated BMGs in Fig. 5(b) and (c). All images show a maze-like structural configuration characteristic of the amorphous structures, and the selected-area electron diffraction pattern (SAED) (inserted at the upper right corner) consists of a halo ring indicative of the glassy structure formation, which are consistent with XRD results. Fig. 5(c) shows a typical bright-field TEM image recorded from the treated BMG, demonstrating a high density of dark contrast regions distributed uniformly throughout the sample. The bright/dark contrast variation indicates a microstructural heterogeneity in the material [26,27]. However, no such inhomogeneity can be found in the as-cast BMG in Fig. 5(b), which mostly displays a featureless glassy phase throughout the whole sample [28]. All the phenomena demonstrate higher stored energy of BMGs can hold more heterogeneous microstructure, displaying softening of soft regions and hardening the hard regions. In essence, combining with DSC results (see in Fig. 2(b)), it can be considered that the high energy states result in the enhancement of microstructural heterogeneity in BMGs. In the process of HRRF, a large amount of energy could be accumulated inside the BMG breaking the initial state of BMG system at the same time [29]. According to the Lennard-Jones-like potential function, when the interatomic distance is larger or smaller than the original distance, the potential energy will increase (in Fig. 6). Thus, the atoms of BMGs will rearrange in short range, some of the atoms pack more densely as hard regions (from a to c)), while the others arranged more loosely as soft regions (from b to d) forcing the system of BMGs to take place in a inhomogeneous-to-more
Fig. 6. Schematic illustration of the Lennard-Jones-like potential, a and c, b and d are the average interatomic distances in hard regions and soft regions with ascast and treated states respectively. Inset: Sketch of the atomic structures for ascast and treated BMGs, respectively.
4. Discussion To reveal the originations, the as-cast and treated BMGs were tested the spatial nanohardeness distributions as demonstrated on the contour maps by nanoindentation over 48 × 48 µm square areas in Fig. 5(a). The as-cast BMG displays only a little fluctuation of nano-hardness values (7.865–7.471 GPa). While, the treated material exhibits a wide range from 8.164 to 7.228 GPa. Both of their averages are about 131
Materials Science & Engineering A 751 (2019) 128–132
Y.-B. Ma, et al.
inhomogeneous transition regime, so as to stay a high-level energy state (see inset of Fig. 6). It is noteworthy that the differences of torsional behaviors in BMGs are confirmed to be associated with their stored energy. During the stress of torsion applying to this material, at first, the STZs initiate in the soft regions preferentially and gradually evolves into shear bands which appear on the surface [30]. As the deformation going on, the local yielding on the surface begins to extend towards the interior of the sample through the propagation of the stable shear bands. The defects of BMGs can give rise to the formation of cavitation and microcracks resulting in the transition of shear bands from stabilization to unstabilization [14]. Ultimately, with the cracks propagating, the sample fails in the catastrophic fracture. Consequently, for the treated metallic glass, the lower potential energy barrier of atomic moving, due to the much energy content for the looser atomic arrangement in the soft regions, makes STZs activations and operations easier than the as-cast glass, which explains the lower yield strength with respect to the as-cast material [31]. By contrast, in the hard regions, a further densification of atomic packing which has been obtained by the high stored energy, results in strengthening the capacity to impede the expansion of the primary shear bands and promoting the formation of second shear bands through changing the direction of dilatation process and branching, thus leading to the global ductility effectively. The intersection of shear bands decreases their sharpness, hinders their rapid propagation, and increases the flow stress. These make the higher ultimate strength behavior of treated samples and compensate the shear band–induced softening [22,32]. So, the treated BMGs can sustain a high amount of structural reconfiguration avoiding the occurrence of catastrophic failure along a single shear band as observed for the as-cast BMGs.
glasses, Nat. Commun. 3 (2012) 6. [2] J.R. Greer, J.T.M. De Hosson, Plasticity in small-sized metallic systems: intrinsic versus extrinsic size effect, Prog. Mater. Sci. 56 (2011) 654–724. [3] M.M. Khan, A. Nemati, Z.U. Rahman, Recent advancements in bulk metallic glasses and their applications: a review, Crit. Rev. Solid State Mater. Sci. 43 (2018) 233–268. [4] D.B. Miracle, A. Concustell, Y. Zhang, Shear bands in metallic glasses: size effects on thermal profiles, Acta Mater. 59 (2011) 2831–2840. [5] A. Inoue, A. Takeuchi, Recent development and application products of bulk glassy alloys, Acta Mater. 59 (2011) 2243–2267. [6] L.H. Liu, Z.Y. Liu, Y. Huan, Effect of structural heterogeneity on serrated flow behavior of Zr-based metallic glass, J. Alloy. Compd. 766 (2018) 908–917. [7] H.B. Yu, J. Hu, X.X. Xia, Stress-induced structural inhomogeneity and plasticity of bulk metallic glasses, Scr. Mater. 61 (2009) 640–643. [8] N. Chen, D. Pan, D.V. Louzguine-Luzgin, Improved thermal stability and ductility of flux-treated Pd40Ni40Si4P16 BMG, Scr. Mater. 62 (2010) 17–20. [9] L. Zhang, F. Jiang, Y. Zhao, Shear band multiplication aided by free volume under three-point bending, J. Mater. Res. 25 (2010) 283–291. [10] J. Fu, Y. Zhu, C. Zheng, R. Liu, Effect of laser shock peening on mechanical properties of Zr-based bulk metallic glass, Appl. Surf. Sci. 313 (2014) 692–697. [11] Y. Cheng, S. Pang, C. Chen, T. Zhang, Tensile plasticity in monolithic bulk metallic glass with sandwiched structure, J. Alloy. Compd. 688 (2016) 724–728. [12] S.H. Joo, D.H. Pi, A.D.H. Setyawan, Work-hardening induced tensile ductility of bulk metallic glasses via high-pressure torsion, Sci. Rep. 5 (2015) 9. [13] Z.-Q. Song, E. Ma, J. Xu, Failure of Zr61Ti2Cu25Al12 bulk metallic glass under torsional loading, Intermetallics 86 (2017) 25–32. [14] X. Zhang, L. Huang, X. Chen, Mechanical behavior of Fe75Mo5P10C7.5B2.5 bulkmetallic glass under torsional loading, Mat. Sci. Eng. A 527 (2010) 7801–7807. [15] C. Jiang, H. Lu, K. Cao, In situ SEM torsion test of metallic glass microwires based on micro robotic manipulation, Scanning 2017 (2017) 1–7. [16] Y.J. Dai, Y. Huan, M. Gao, Development of a high-resolution micro-torsion tester for measuring the shear modulus of metallic glass fibers, Meas. Sci. Technol. 26 (2015) 025902. [17] M.D. Ediger, P. Harrowell, Perspective, supercooled liquids and glasses, J. Chem. Phys. 137 (2012) 080901. [18] Q.D. Zhang, L.F. Wang, Y. Zhao, Remarkable improving plasticity of a brittle Zrbased bulk metallic glass by a high rheological rate forming method in centesimal seconds, Mater. Lett. 164 (2016) 348–352. [19] S.V. Ketov, Y.H. Sun, S. Nachum, Rejuvenation of metallic glasses by non-affine thermal strain, Nature 524 (2015) 200–203. [20] C. Wang, Z.Z. Yang, T. Ma, High stored energy of metallic glasses induced by high pressure, Appl. Phys. Lett. 110 (2017) 111901. [21] M. Ezzeldien, Z. Kovacs, J. Lendvai, F. Chmelik, K. Mathis, Stages in room temperature torsional deformation of a Vitreloy bulk metallic glass, J. Alloy. Compd. 577 (2013) 25–29. [22] J.M. Park, G. Wang, S. Pauly, N. Mattern, Ductile Ti-based bulk metallic glasses with high specific strength, Meter. Mater. Trans. A 42 (2011) 1456–1462. [23] J.W. Liu, Q.P. Cao, L.Y. Chen, Shear band evolution and hardness change in coldrolled bulk metallic glasses, Acta Mater. 58 (2010) 4827–4840. [24] Y. Yang, C.T. Liu, Size effect on stability of shear-band propagation in bulk metallic glasses: an overview, J. Mater. Sci. 47 (2012) 55–67. [25] B.A. Sun, W.H. Wang, The fracture of bulk metallic glasses, Prog. Mater. Sci. 74 (2015) 211–307. [26] Y.B. Wang, D.D. Qu, X.H. Wang, Introducing a strain-hardening capability to improve the ductility of bulk metallic glasses via severe plastic deformation, Acta Mater. 60 (2012) 253–260. [27] J.G. Wang, D.Q. Zhao, M.X. Pan, Mechanical heterogeneity and mechanism of plasticity in metallic glasses, Appl. Phys. Lett. 94 (2009) (031904 − 031904-3). [28] M.H. Lee, K.S. Lee, J. Das, Improved plasticity of bulk metallic glasses upon cold rolling, Scr. Mater. 62 (2010) 678–681. [29] Y.Q. Cheng, E. Ma, Atomic-level structure and structure–property relationship in metallic glasses, Prog. Mater. Sci. 56 (2011) 379–473. [30] Y.Q. Cheng, E. Ma, Atomic-level structure and structure–property relationship in metallic glasses, Prog. Mater. Sci. 56 (2011) 379–473. [31] L.Y. Chen, Z.D. Fu, G.Q. Zhang, New class of plastic bulk metallic glass, Phys. Rev. Lett. 100 (2008) 075501. [32] J. Pan, Y.X. Wang, Q. Guo, Extreme rejuvenation and softening in a bulk metallic glass, Nat. Commun. 9 (2018).
5. Conclusions In summary, ductile and high-strength BMGs were obtained by HRRF under torsion condition. During this process, high energy was introduced into brittle BMG precursors to endow an extremely inhomogeneous microstructure of the BMGs without any change in the local chemical composition and nanocrystallization, which promotes STZ operation and retards the formation of detrimental runaway shear bands. These microscopic behaviors lead to a dramatic increase in macroscopic plasticity and ultimate strength of the BMGs. This study provides an effective strategy for brittle BMGs to enhance their torsional plasticity, thus optimizing their practical application as an excellent engineering material. Acknowledgements This work was supported by the National Natural Science Foundation of China (Grant nos. 50971053 and 11304073), and the National Basic Research Program of China (Grant no. 2012CB825702). The authors thank X. P. Wang and H. J. Wang for their technical assistance. References [1] L. Tian, Y.Q. Cheng, Z.W. Shan, J. Li, Approaching the ideal elastic limit of metallic
132
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Torsional behaviors in Zr-based bulk metallic glass with high stored energy structure Yu-Bai Maa, Yan Jiangb, Hua-Ping Dingc, Qi-Dong Zhanga, Xiao-Yun Lia, Fang-Qiu Zua,
T
⁎
a
Liquid/Solid Metal Processing Institute, School of Materials Science & Engineering, Hefei University of Technology, Hefei 230009, China Panzhihua International Research Institute of Vanadium and Titanium, Panzhihua 617000,China c State Key Laboratory of Materials Processing and Die & Mould Technology, Huazhong University of Science and Technology, 1037 Luoyu Road, 430074 Wuhan, Hubei, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Bulk metallic glasses Torsional plasticity Stored energy Heterogeneity
To overcome brittleness of bulk metallic glasses, great efforts have been devoted to optimizing BMGs′ plasticity under compressive, tensional or bending load. Yet, it is hardly found any work done to enhance torsional plasticity of BMGs. Herein, based on the different stored energy of BMGs determining diverse properties, we introduce the energy into Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 BMG by a high rheological rate forming method, leading to the higher microstructural heterogeneity, which is conducive to increase not only the plasticity (from 0.2% up to 1.56%) but also the ultimate strength from (990 MPa up to 1110 MPa) on the torsional deformation effectively. This work provides a new orientation to break through the brittleness of BMGs.
1. Introduction
BMGs can be diverse at different preparation conditions [17], while its internal relation with the torsional plasticity needs to further explore still. Here, in this paper, the high rheological rate forming (HRRF) technique [18] is applied to introduce energy into Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 BMG to illuminate this issue. The result shows that the preserved higher energy in this brittle BMG can cause the higher microstructural heterogeneity, resulting in enhancing both the ductility and ultimate strength on the torsion.
Bulk metallic glasses (BMGs) hold a variety of unique properties due to the amorphous state of the atoms, which have been an area of intense research promptly for several decades [1,2]. Nevertheless, this peculiar microstructure has brought the brittleness to BMGs at room temperature [3]. In fact, the plasticity of this material generates in the narrow shear bands of only 10–20 nm width [4,5]. Thus, when the shear bands take place and cross the sample [6], the whole deformation becomes instability and fractures immediately. To address this challenge, many researchers have investigated the possibilities to improve the plasticity including various intrinsic and extrinsic ways [7–12]. However, most of these studies have paid close attention to the mechanical behaviors on compressive [7,8], bending [9,10] and tensile conditions [11,12]. Up to now, to our knowledge, no researches have been carried out on improving the torsional behaviors of BMGs. Actually, as a candidate material for structural applications, the torsional properties can process great significances, such as torsion bars in scanners, transmission shafts, screws, etc. Moreover, similar to the tension, the torsional deformation also occurs in the unconstrained stress field and the fracture of BMGs is always controlled by a single dominant shear band. Thus, most of BMGs show a limited plastic strain (~0.2%) during the torsional deformation [13–16]. Therefore, how to improve the torsional behaviors of BMGs is a significant issue to be solved. Furthermore, it is understandable that the stored energy of ⁎
2. Experimental Amorphous Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 rods with a diameter of 6 mm, and a length of 120 mm were prepared using copper mold casting. Then cut them into 27 mm long cylinders by a cutting machine to fabricate the treated cylinder with 4.4 mm in diameter and a length of 40 mm through HRRF [18]. A test sample was processed into the hexagonal prisms with a gauge length of L = ~9 mm and diameter of d = ~1.8 mm joining with a radius to the long sample heads. The hexagonal heads were designed to prevent slipping while twisting, and the gauge sections of samples were polished with diamond paste (inset of Fig. 3). Torsional tests were performed on the Instron 8874 testing machine at a constant rotation rate of 0.057°/s with careful alignment between the specimen and the rotation axis. One hexagonal head of the sample was fixed on the free end of the testing machine to avoid
Corresponding author. E-mail address: [email protected] (F.-Q. Zu).
https://doi.org/10.1016/j.msea.2019.02.074 Received 10 January 2019; Received in revised form 21 February 2019; Accepted 22 February 2019 Available online 23 February 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 751 (2019) 128–132
Y.-B. Ma, et al.
channel and subsequently filled into the mold cavity. According to the displacement of plunger recorded by a high-speed camera, the whole experiment only lasted ~25 ms, which was enough to avoid thermal embrittlement. 3. Results Fig. 2(a) shows the X-ray diffraction patterns obtained from Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 BMG in the as-cast and treated states respectively. The two patterns exhibit the same broad diffusion maxima and conform to a characteristic of the full amorphous structure. The DSC curves of them (in Fig. 2(b)) have presented the sequential transition of glass transition and supercooled liquid region. Although the onset crystallization temperature (Tx) has almost no change, it is obvious seen that the glass transition temperature (Tg) increases from about 684–694 K (see Table (1)), and the area of an exothermic peak before reaching the Tg in DSC which represents the structural relaxation enthalpy ΔH rises from ~5.432 to ~11.54 J/g dramatically (see the black shaded areas in inset of Fig. 2(b)) with the change from as-cast to treated state. The high structural relaxation enthalpy could be regarded as proofs for much more multiple free volume concentration in treated BMGs because of a positive proportional relationship between free volume change [19]. Additionally, the onset of structural relaxation temperature of as-cast specimens began at ~518 K, while treated specimens began ~503 K. These phenomena indicate the energy system of the metallic glass has increased into a high-level energy state by HRRF [20]. Fig. 3 displays the typical shear-strain curve of Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 BMG under the torsional deformation for as-cast and treated specimens. From these curves, the as-cast specimen fails in a brittle manner with only a small nonlinear trend (plastic strain ~0.2%) before the final catastrophic fracture. However, for the treated sample, it is noticeable that the curve deviates from linearity at the high stress (~880 MPa) and displays the increase of the stress with a further increase in strain after yielding, which is most probably a consequence of noticeable interactions between the shear bands [21]. The macroscopic plastic strain of the treated sample increases into ~1.56% which is ~seven times than that of the as-cast sample. While the torsional ultimate strength grows up to ~1110 MPa, which is higher than that of as-cast one (~990 MPa), (The calculated yielding strengths σy, ultimate strengths σf, and the plastic strain εp were summarized in Table (1)) indicating that the HRRF method has a prominent effect on improving the torsional property of the metallic glass. In order to reflect the macroscopic plastic deformation intuitively, a suitable ink line which is parallel to the torsion axis has been marked on the lateral surface of the sample. After fracture, the variation of the marking line slope (θ) is showed in Fig. 4(a) and (b). By carefully measuring the slope, the ink lines of the as-cast sample is almost
Fig. 1. Schematic drawing of the HRRF method.
uniaxial stresses. At least three samples were measured to ensure the accuracy of the results. The nanoindentation experiment was conducted by an Agilent G200 Nanoindenter with an indenter of Diamond Berkovich at room temperature. Eighty-one indentations were made in a 9 × 9 matrix with the same depth (300 nm) at a constant loading rate of 0.05 s−1, and the spacing between adjacent indentations was set as 6 µm to avoid overlapping of neighboring strained zones. The glassy nature of BMG samples was confirmed by x-ray diffraction (XRD) using a BRUKER D8 ADVANCE diffractometer with Cu Kα radiation source and differential scanning calorimetry (DSC) performing under a purified argon atmosphere in a NETZSCH 449F3. The field emission scanning electron microscope (FESEM, SU8020) was used to observe the lateral and fracture surface morphology of samples. The thin samples were prepared by Ar ion-milling under liquid nitrogen cooling (Gatan Model 600) at 2.5 keV to investigate the detailed microstructures using the transmission electron microscope (TEM; JEM2100F). As mentioned above, to obtain high stored energy BMGs, we adopt a high rheological rate forming process (HRRF) to treat the as-cast samples. The device schematic diagram of the HRRF method consists of a power supply, a plunger, an arrow channel and a mold cavity (see in Fig. 1). A high force of (~6 kN) was pre-loaded on the plunger ahead of time. Then the as-cast BMG was heated to the supercooled liquid state (~745 k) measured by a nanovoltmeter applying the method of fast Joule heating. Therefore, the liquid BMG was squeezed into the narrow
Fig. 2. (a) The XRD patterns and (b) DSC traces of the as-cast and treated samples. Inserted graph of (b) shows as-cast and treated samples marked by black dashed rectangular. 129
Materials Science & Engineering A 751 (2019) 128–132
Y.-B. Ma, et al.
Table 1 Summary of the thermal parameters and torsional properties of the Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 BMG with as-cast and treated. Alloy
Tg (K)
Tx (K)
ΔH (J/g)
εp (%)
σy (MPa)
σf (MPa)
As-cast Treated
684 ± 3 694 ± 3
761 ± 3 758 ± 3
5.43 ± 0.05 11.54 ± 0.05
0.20 ± 0.05 1.56 ± 0.05
930 ± 13 880 ± 13
990 ± 10 1110 ± 10
which violates the Tresca criterion. It indicates that in addition to the shear stress, there is some normal stress applying on the sample. This normal stress is calculated by σy·sin(π −2δ )= 91 MPa, which is ~8% of the applied shear stress. Such a little normal stress can be explained by the secondary shear bands. Secondary shear bands may result in the significant difference between values of critical strain accommodation along perpendicular and a parallel direction with respect to the shear band planes. This physical separation is very similar to the normal stress [22]. Furthermore, the red-lined rectangular in Fig. 4(a) and (b), which are the side surface morphology near the fracture of as-cast and treated specimens, are magnified in Fig. 4(c) and (d) respectively. As illustrated in Fig. 4(d), multiple and intersecting shear bands including primary and secondary shear-bands (white arrows in Fig. 4(d)) are clearly observed on the surface of the treated sample which is considered one of the main evidence that account for the extent of plasticity in glass metals. The primary shear-bands are almost parallel to the shear stress direction. While the secondary shear-bands, which are initiated from the intersection points [23], propagate in disorder orientations. This indicates that there is not a preferential slip plane for secondary ones in BMGs [24]. In turn, almost no shear band is visible on the surface for the as-cast sample in Fig. 4(c), subjected to only a single dominant shear band cross the whole sample during the instantaneous fracture [25].
Fig. 3. The shear stressstrain curves obtained on as-cast and treated samples under torsional loading. The inset shows a pre-treated BMG rod, a treated BMG sample and a typical torsion sample, respectively.
parallel to the torsional axis and the slope θ is just about 0.2° (manifested by the yellow arrows) indicating that only a little plastic deformation happened during the torsion. However, the treated sample presents the larger slope (θ = 1.6°) than as-cast one, which is also obvious with its enhanced plasticity. According to the Tresca criterion, a fracture angle (δ) which is between the torsional fracture surface and the axial axis of the specimen should be 90°. For the as-cast one, its fracture angle is equal to ~90° (marked by the red arrow). While, the treated sample is just ~87°,
Fig. 4. (a) and (b) are the global views of the side surface fractured samples of the as-cast and treated with an ink line, respectively. (c) and (d) High-magnification detailed micrographs corresponding to the part marked by the red-lined rectangular in (a) and (b), respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.) 130
Materials Science & Engineering A 751 (2019) 128–132
Y.-B. Ma, et al.
Fig. 5. (a) Spatial distribution of the nano-hardness values of the Zr54.46Al9.9Ni4.95Cu29.7Pd0.99 BMGs before and after HRRF. (b) and (c) the typical bright-field TEM images with insert diffraction pattern images of the as-cast and treated samples respectively.
7.66 GPa. Further to verify the HRRF imposing the microstructural heterogeneity of BMGs, TEM was carried out to clarify the differences of microstructures between as-cast and treated BMGs in Fig. 5(b) and (c). All images show a maze-like structural configuration characteristic of the amorphous structures, and the selected-area electron diffraction pattern (SAED) (inserted at the upper right corner) consists of a halo ring indicative of the glassy structure formation, which are consistent with XRD results. Fig. 5(c) shows a typical bright-field TEM image recorded from the treated BMG, demonstrating a high density of dark contrast regions distributed uniformly throughout the sample. The bright/dark contrast variation indicates a microstructural heterogeneity in the material [26,27]. However, no such inhomogeneity can be found in the as-cast BMG in Fig. 5(b), which mostly displays a featureless glassy phase throughout the whole sample [28]. All the phenomena demonstrate higher stored energy of BMGs can hold more heterogeneous microstructure, displaying softening of soft regions and hardening the hard regions. In essence, combining with DSC results (see in Fig. 2(b)), it can be considered that the high energy states result in the enhancement of microstructural heterogeneity in BMGs. In the process of HRRF, a large amount of energy could be accumulated inside the BMG breaking the initial state of BMG system at the same time [29]. According to the Lennard-Jones-like potential function, when the interatomic distance is larger or smaller than the original distance, the potential energy will increase (in Fig. 6). Thus, the atoms of BMGs will rearrange in short range, some of the atoms pack more densely as hard regions (from a to c)), while the others arranged more loosely as soft regions (from b to d) forcing the system of BMGs to take place in a inhomogeneous-to-more
Fig. 6. Schematic illustration of the Lennard-Jones-like potential, a and c, b and d are the average interatomic distances in hard regions and soft regions with ascast and treated states respectively. Inset: Sketch of the atomic structures for ascast and treated BMGs, respectively.
4. Discussion To reveal the originations, the as-cast and treated BMGs were tested the spatial nanohardeness distributions as demonstrated on the contour maps by nanoindentation over 48 × 48 µm square areas in Fig. 5(a). The as-cast BMG displays only a little fluctuation of nano-hardness values (7.865–7.471 GPa). While, the treated material exhibits a wide range from 8.164 to 7.228 GPa. Both of their averages are about 131
Materials Science & Engineering A 751 (2019) 128–132
Y.-B. Ma, et al.
inhomogeneous transition regime, so as to stay a high-level energy state (see inset of Fig. 6). It is noteworthy that the differences of torsional behaviors in BMGs are confirmed to be associated with their stored energy. During the stress of torsion applying to this material, at first, the STZs initiate in the soft regions preferentially and gradually evolves into shear bands which appear on the surface [30]. As the deformation going on, the local yielding on the surface begins to extend towards the interior of the sample through the propagation of the stable shear bands. The defects of BMGs can give rise to the formation of cavitation and microcracks resulting in the transition of shear bands from stabilization to unstabilization [14]. Ultimately, with the cracks propagating, the sample fails in the catastrophic fracture. Consequently, for the treated metallic glass, the lower potential energy barrier of atomic moving, due to the much energy content for the looser atomic arrangement in the soft regions, makes STZs activations and operations easier than the as-cast glass, which explains the lower yield strength with respect to the as-cast material [31]. By contrast, in the hard regions, a further densification of atomic packing which has been obtained by the high stored energy, results in strengthening the capacity to impede the expansion of the primary shear bands and promoting the formation of second shear bands through changing the direction of dilatation process and branching, thus leading to the global ductility effectively. The intersection of shear bands decreases their sharpness, hinders their rapid propagation, and increases the flow stress. These make the higher ultimate strength behavior of treated samples and compensate the shear band–induced softening [22,32]. So, the treated BMGs can sustain a high amount of structural reconfiguration avoiding the occurrence of catastrophic failure along a single shear band as observed for the as-cast BMGs.
glasses, Nat. Commun. 3 (2012) 6. [2] J.R. Greer, J.T.M. De Hosson, Plasticity in small-sized metallic systems: intrinsic versus extrinsic size effect, Prog. Mater. Sci. 56 (2011) 654–724. [3] M.M. Khan, A. Nemati, Z.U. Rahman, Recent advancements in bulk metallic glasses and their applications: a review, Crit. Rev. Solid State Mater. Sci. 43 (2018) 233–268. [4] D.B. Miracle, A. Concustell, Y. Zhang, Shear bands in metallic glasses: size effects on thermal profiles, Acta Mater. 59 (2011) 2831–2840. [5] A. Inoue, A. Takeuchi, Recent development and application products of bulk glassy alloys, Acta Mater. 59 (2011) 2243–2267. [6] L.H. Liu, Z.Y. Liu, Y. Huan, Effect of structural heterogeneity on serrated flow behavior of Zr-based metallic glass, J. Alloy. Compd. 766 (2018) 908–917. [7] H.B. Yu, J. Hu, X.X. Xia, Stress-induced structural inhomogeneity and plasticity of bulk metallic glasses, Scr. Mater. 61 (2009) 640–643. [8] N. Chen, D. Pan, D.V. Louzguine-Luzgin, Improved thermal stability and ductility of flux-treated Pd40Ni40Si4P16 BMG, Scr. Mater. 62 (2010) 17–20. [9] L. Zhang, F. Jiang, Y. Zhao, Shear band multiplication aided by free volume under three-point bending, J. Mater. Res. 25 (2010) 283–291. [10] J. Fu, Y. Zhu, C. Zheng, R. Liu, Effect of laser shock peening on mechanical properties of Zr-based bulk metallic glass, Appl. Surf. Sci. 313 (2014) 692–697. [11] Y. Cheng, S. Pang, C. Chen, T. Zhang, Tensile plasticity in monolithic bulk metallic glass with sandwiched structure, J. Alloy. Compd. 688 (2016) 724–728. [12] S.H. Joo, D.H. Pi, A.D.H. Setyawan, Work-hardening induced tensile ductility of bulk metallic glasses via high-pressure torsion, Sci. Rep. 5 (2015) 9. [13] Z.-Q. Song, E. Ma, J. Xu, Failure of Zr61Ti2Cu25Al12 bulk metallic glass under torsional loading, Intermetallics 86 (2017) 25–32. [14] X. Zhang, L. Huang, X. Chen, Mechanical behavior of Fe75Mo5P10C7.5B2.5 bulkmetallic glass under torsional loading, Mat. Sci. Eng. A 527 (2010) 7801–7807. [15] C. Jiang, H. Lu, K. Cao, In situ SEM torsion test of metallic glass microwires based on micro robotic manipulation, Scanning 2017 (2017) 1–7. [16] Y.J. Dai, Y. Huan, M. Gao, Development of a high-resolution micro-torsion tester for measuring the shear modulus of metallic glass fibers, Meas. Sci. Technol. 26 (2015) 025902. [17] M.D. Ediger, P. Harrowell, Perspective, supercooled liquids and glasses, J. Chem. Phys. 137 (2012) 080901. [18] Q.D. Zhang, L.F. Wang, Y. Zhao, Remarkable improving plasticity of a brittle Zrbased bulk metallic glass by a high rheological rate forming method in centesimal seconds, Mater. Lett. 164 (2016) 348–352. [19] S.V. Ketov, Y.H. Sun, S. Nachum, Rejuvenation of metallic glasses by non-affine thermal strain, Nature 524 (2015) 200–203. [20] C. Wang, Z.Z. Yang, T. Ma, High stored energy of metallic glasses induced by high pressure, Appl. Phys. Lett. 110 (2017) 111901. [21] M. Ezzeldien, Z. Kovacs, J. Lendvai, F. Chmelik, K. Mathis, Stages in room temperature torsional deformation of a Vitreloy bulk metallic glass, J. Alloy. Compd. 577 (2013) 25–29. [22] J.M. Park, G. Wang, S. Pauly, N. Mattern, Ductile Ti-based bulk metallic glasses with high specific strength, Meter. Mater. Trans. A 42 (2011) 1456–1462. [23] J.W. Liu, Q.P. Cao, L.Y. Chen, Shear band evolution and hardness change in coldrolled bulk metallic glasses, Acta Mater. 58 (2010) 4827–4840. [24] Y. Yang, C.T. Liu, Size effect on stability of shear-band propagation in bulk metallic glasses: an overview, J. Mater. Sci. 47 (2012) 55–67. [25] B.A. Sun, W.H. Wang, The fracture of bulk metallic glasses, Prog. Mater. Sci. 74 (2015) 211–307. [26] Y.B. Wang, D.D. Qu, X.H. Wang, Introducing a strain-hardening capability to improve the ductility of bulk metallic glasses via severe plastic deformation, Acta Mater. 60 (2012) 253–260. [27] J.G. Wang, D.Q. Zhao, M.X. Pan, Mechanical heterogeneity and mechanism of plasticity in metallic glasses, Appl. Phys. Lett. 94 (2009) (031904 − 031904-3). [28] M.H. Lee, K.S. Lee, J. Das, Improved plasticity of bulk metallic glasses upon cold rolling, Scr. Mater. 62 (2010) 678–681. [29] Y.Q. Cheng, E. Ma, Atomic-level structure and structure–property relationship in metallic glasses, Prog. Mater. Sci. 56 (2011) 379–473. [30] Y.Q. Cheng, E. Ma, Atomic-level structure and structure–property relationship in metallic glasses, Prog. Mater. Sci. 56 (2011) 379–473. [31] L.Y. Chen, Z.D. Fu, G.Q. Zhang, New class of plastic bulk metallic glass, Phys. Rev. Lett. 100 (2008) 075501. [32] J. Pan, Y.X. Wang, Q. Guo, Extreme rejuvenation and softening in a bulk metallic glass, Nat. Commun. 9 (2018).
5. Conclusions In summary, ductile and high-strength BMGs were obtained by HRRF under torsion condition. During this process, high energy was introduced into brittle BMG precursors to endow an extremely inhomogeneous microstructure of the BMGs without any change in the local chemical composition and nanocrystallization, which promotes STZ operation and retards the formation of detrimental runaway shear bands. These microscopic behaviors lead to a dramatic increase in macroscopic plasticity and ultimate strength of the BMGs. This study provides an effective strategy for brittle BMGs to enhance their torsional plasticity, thus optimizing their practical application as an excellent engineering material. Acknowledgements This work was supported by the National Natural Science Foundation of China (Grant nos. 50971053 and 11304073), and the National Basic Research Program of China (Grant no. 2012CB825702). The authors thank X. P. Wang and H. J. Wang for their technical assistance. References [1] L. Tian, Y.Q. Cheng, Z.W. Shan, J. Li, Approaching the ideal elastic limit of metallic
132