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Designing ductile CuZr-based metallic glass matrix composites
Materials Science & Engineering A 682 (2017) 542–549
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Designing ductile CuZr-based metallic glass matrix composites a
Y.J. Liu , H.W. Yao Z.H. Wangc
a,b
c
, T.W. Zhang , Z. Wang
a,b
a,b
, Y.S. Wang
a,b,⁎
, J.W. Qiao
a,b,⁎
, H.J. Yang
,
crossmark
a Laboratory of Applied Physics and Mechanics of Advanced Materials, College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, China b Key Laboratory of Interface Science and Engineering in Advanced Materials, Ministry of Education, Taiyuan University of Technology, Taiyuan 030024, China c Institute of Applied Mechanicsand Biomedical Engineering, Taiyuan University of Technology, Taiyuan 030024, China
A R T I C L E I N F O
A BS T RAC T
Keywords: Metallic glasses Composites Work-hardening Microstructure Martensitic phase transformation
The CuZr-based metallic glass matrix composites (MGMCs) with different volume fractions of crystalline phases were designed by doping of nickel. Minor addition of nickel element can change the glass-forming ability of the resultant composites. Large compressive plasticity accompanied by strong work-hardening capacity was achieved in the Cu47Zr48Al4Ni1 composite with a volume fraction of crystalline phases of 33%. The excellent compressive properties were mainly attributed to the inhibition for the propagation of shear bands by the ductile crystals and deformation induced martensitic transformation of the B2-CuZr phase. However, no obvious global tensile ductility was obtained, due to the mode I fracture toughness and small plastic-zone size of glass matrix. To uncover the shear-band evolution during deformation, finite element simulation was conducted, revealing that the ductile B2 phase can tune the shear-stress distribution and consequently initiate and retard shear bands, which stimulates the multiplication of shear bands. Accordingly, the spacing of B2 CuZr particles is a vital factor dominating the plasticity of CuZr-based MGMCs, especially upon tension.
1. Introduction Bulk metallic glasses (BMGs) have attracted large attention because of their unique properties, such as high strength, large elastic limit, excellent corrosion and wear resistances, etc [1]. However, they are known for being extremely brittle, failing in a catastrophic fracture at room temperature, due to formation of localized shear bands [2,3]. To enhance macroscopic plasticity, metallic glass matrix composites (MGMCs) have been developed by in situ or ex situ introduction of secondary phases, which can hinder the rapid propagation of shear bands [4]. In recent years, a series of in-situ dendrite-reinforced MGMCs have been developed, such as Ti-based and Zr-based MGMCs [5–7]. Nevertheless, these MGMCs exhibit a macroscopic strain softening with localized necking due to lack of work hardening in the glass matrix at room temperature, especially under tensile conditions, limiting the applications significantly. Recently, these shortcomings have been successfully addressed in some CuZr-based MGMCs reinforced by ductile B2 CuZr phases [8– 15]. It is worth noting that the B2-type CuZr phases are prone to deform plastically compared to other intermetallic compounds with
complex crystalline structure [16–18], resisting fast propagation of highly localized shear bands effectively and promoting the generation of multiple shear bands in the glass matrices. Moreover, as crystalline phases with a shape memory effect, B2-type CuZr phases can undergo a stress-induced martensitic transformation from B2 phase to a monoclinic B19ʹ phase, which imparts an appreciable work-hardening capability and then compensates the softening of the matrix significantly during plastic flows [9,12–14,19]. In this way, the effect of “transformation-induced plasticity” (TRIP) [11,14,20] can be introduced into the ductile MGMCs. Depending on the TRIP effect, the Zr48Cu47.5Al4Co0.5 composite, containing 25% volume fractioned spherical B2-CuZr phases, presents strong work-hardening capability and considerable tensile ductility [14]. Subsequently, Wu et al. [21] have demonstrated that a minor 0.5 at% addition of Co in the Zr48Cu47.5Al4Co0.5 composites can promote deformation twinning and martensite transformation of B2-CuZr phases by reducing stacking fault energy (SFE). Meanwhile, the SFE of B2 phases is predicted to be strongly reduced by the addition of Ni theoretically. In addition, since the glass-forming ability (GFA) of these BMGs and MGMCs is sensitive to minor additions [9,22,23], an effective adjustment of microstruc-
⁎ Corresponding authors at: Laboratory of Applied Physics and Mechanics of Advanced Materials, College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, China. E-mail addresses: [email protected] (J.W. Qiao), [email protected] (H.J. Yang).
http://dx.doi.org/10.1016/j.msea.2016.11.079 Received 4 August 2016; Received in revised form 19 October 2016; Accepted 22 November 2016 Available online 24 November 2016 0921-5093/ © 2016 Elsevier B.V. All rights reserved.
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tures can be obtained via appropriate trace additions. It is noted that microstructures plays a most effective role in improving the mechanical properties of such alloys [9,13,24]. As such, an improvement of mechanical properties is expected by a minor addition of Ni. However, the effect of Ni addition on the microstructures and mechanical properties of such kinds of MGMCs have been rarely investigated. In this study, a series of CuZr-based MGMCs with different volume fractions of crystalline phases are systematically designed via minor addition of Ni. Compressive and tensile experiments are carried out to investigate the deformation and fracture mechanisms. Moreover, the correlation between the evolution of shear bands and plastic flow can be revealed vividly, combined with finite element simulation (FES). Additionally, dependences of the mechanical properties on microstructures are explored as well. 2. Experimental Fig. 1. XRD spectra of the as-cast Cu48−xZr48Al4Nix alloys with different Ni additions.
Master alloys with normal compositions of Zr48Cu48−xAl4Nix (x=0, 0.5, 0.8, 1, 1.5, and 5 at%) were prepared by arc melting the mixture of constituent elements with purity greater than 99.9% (wt%) under a Tigettered argon atmosphere. The ingots were melted four times to ensure compositional homogeneity. Plate-shape samples with 1.5 mm in thickness and 10 mm in width were prepared by the copper-mouldcasting method. The tensile specimens with gauge dimensions of 10 mm (length)×1.5 mm (width)×1.5 mm (thickness) and compressive specimens with gauge dimensions of 3 mm (length)×1.5 mm (width)×1.5 mm (thickness) were prepared by the electric spark method, respectively. Both compressive and tensile tests were conducted at a strain rate of 5×10−4 s−1 at room temperature. Optical microscope (OM) was used to examine microstructures of as-cast samples. The volume fraction of B2-CuZr phase was determined by an Image-Pro Plus 6.0 software based on the OM images. The phases of the plates were identified by X-ray diffraction (XRD), and the thermal properties were analyzed by differential scanning calorimetry (DSC) at a rate of 20 K/min. Scanning electron microscopy (SEM) was employed to observe the microstructures and fractographs. Nanoindentation was performed on Nano-Indenter II tester (MTS Systems, USA) at room temperature to measure the hardness of the glass matrix and B2 crystalline phases of the as-cast and deformed samples. A Berkovich diamond tip was used. The nominal strain rate was 0.002 s−1 and the maximum indentation depth was 1000 nm. The cross-section of the samples for the nanoindentation tests was carefully polished.
Fig. 2. DSC curves of the as-cast Cu48−xZr48Al4Nix alloys with different Ni additions.
indicates that the volume fraction of the crystalline phase increases. In contrast, with additions of 1.5% and 5% Ni, the exothermic peaks increase, which means that the volume fraction of the amorphous phase increases. The DSC results are consistent with XRD results. Fig. 3 shows the optical cross-section metallographs of as-cast Cu48−xZr48Al4Nix (x=0, 0.5, 1, 1.5, and 5) alloys. For the Ni-free sample, only limited numbers of B2 CuZr particles can be observed, and their sizes are less than 50 µm, as shown in Fig. 3(a). With a 0.5% Ni addition, in addition to some large B2 particles (Fig. 3b), a number of small B2 particles with a diameter of about 5–10 µm can be found, as presents in the inset of Fig. 4(b). The bimodal sized phenomenon of B2 particles has been reported in Ti-based MGMCs as well [19]. For the 1% Ni-added sample, the crystalline volume fraction further increases, as shown in Fig. 4(c). And a large number of B2 CuZr particles can be found throughout the whole section in Fig. 4(f). With further increase of Ni additions, the volume fraction of B2 crystalline phases decreases, specially, for 5% Ni added sample, only a small amount of B2 phase can be found at the out surface of samples. Uniaxial compression tests were carried on the as-cast Cu48−xZr48Al4Nix (x=0, 0.5, 0.8, 1, and 1.5) samples at room temperature, and the corresponding engineering stress-strain curves for these samples are displayed in Fig. 4. The compressive plasticity of the current composites is enhanced with increasing the B2 CuZr crystalline phase. The ternary Cu48Zr48Al4 alloy without Ni addition exhibits a distinct yielding at ~1840 MPa, and has a fracture with the plasticity of ~0.5%. With minor additions of 0.5% and 0.8% Ni, the resultant
3. Results 3.1. Microstructural evolution with Ni additions Fig. 1 shows XRD spectra of as-cast plates of the alloys with different Ni additions. For the ternary Cu48Zr48Al4 alloy, Bragg diffraction peaks corresponding to the body-centered cubic B2 CuZr phases are superimposed on a broad hump. With addition of 0.5–1% Ni, the peak intensity of the crystalline phase increases gradually, indicating an increasing amount of the B2 phase. Specially, this scenario is obvious appreciably for the 1% Ni-added sample, revealing that 1% Ni addition greatly deteriorates the GFA of Cu48−xZr48Al4NiX alloys. However, the volume fractions of B2 CuZr phases decrease indicated by the declined relative intensity of crystalline peaks for the 1.5% Ni-added sample. With further Ni additions (5%), the crystalline peaks become indistinctive visibly, suggesting that excessive Ni additions could enhance the GFA of Cu48−xZr48Al4Ni1 alloys. Fig. 2 presents the DSC curves of these MGMCs. The features of the DSC curves are all similar, namely, an endothermic platform caused by the glass transition and an exothermic heat peak caused for crystallization. The heat release from the crystallization process of these composites decreases gradually with the Ni added from 0% to 1%. This 543
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Fig. 3. cross-section optical metallographs of the as-cast Cu48−xZr48Al4Nix alloys with different Ni additions, (a) x=0, (b) x=0.5, (c) x=1.0, (d) x=1.5, (e) x=5, and (f) low-magnitude image of the 1% Ni added alloys.
[27,28]. With further addition of 1.5% Ni, the composite exhibits a yielding strength of ~1640 MPa, and an uniform plastic strain of ~2%.Fig. 5 presents fractographs of the Cu48Zr48Al4 and Cu47Zr48Al4Ni1 composites upon compression. As for the Cu48Zr48Al4 composite with less B2 phases, it exhibits a shear fracture along a shear angle of 45 degrees with respect to the compressive loading direction in the inset of Fig. 5(a). Normally, the shearing angle of metallic glasses is less than 45 degrees upon compression [29,30]. Only a few primary shear bands are found on the magnified image for the lateral surface of the fractured Cu48Zr48Al4 composite in Fig. 5(a). Fig. 5(b) shows the sideview micrographs of deformed Cu47Zr48Al4Ni1 composite. The fracture plane of the current composite inclines about 45 degrees with the loading direction as well, and some detour shear bands and cracks can be found even at low magnitude in the inset of Fig. 5(b). Highdensity shear bands and a number of martensite phases with slat structures, denoted by the bright arrows, can be observed from the high-magnitude image in Fig. 5(b). In addition, shear bands propagate in a wave-like path as they intersect with the reinforced crystals embedded in the glassy matrix, as shown in the rectangular frame (Fig. 5b), indicating that the propagation of shear bands is hindered by B2 particles effectively. From the detail of the lateral surface for the deformed Cu47Zr48Al4Ni1 composite, the “blocking effect” [13,19] can be clearly observed, as displayed in Fig. 5(c). Profuse shear bands are formed at the interface between B2 phases and the glassy matrices. Afterwards, the fracture surface of the current composite presents dense vein patterns in Fig. 5(d). It is noted that vein patterns are intersected by the torn and elongated crystalline phases, pointed by the bright arrows, suggesting that B2 phases impede the catastrophic fracture, even at the final stage. Such fractographs which contribute to excellent mechanical properties have been found in Ti-based MGMCs [31]. As depicted above, there are martensite phases on the side surface and fracture surface of the Cu47Zr48Al4Ni1 composite, which induces the plasticity improvement. In particular, the pronounced work-hardening capability is attributed to the TRIP effect. In order to further explore the deformation micromechanisms, XRD measurements and nanoindentation tests of the as-cast and deformed Cu47Zr48Al4Ni1
Fig. 4. Compressive engineering stress-strain curves of the as-cast Cu48−xZr48Al4Nix alloys with different Ni additions.
MGMCs exhibit pronounced work-hardening behavior after yielding at ~1710 and ~1720 MPa, respectively, and the uniform plastic strains are ~1.5 and ~2.5%, respectively. For the 1% Ni-added composite, it is clearly shown that the composite yields at ~1580 MPa and undergoes large plasticity of ~11%, together with distinct work hardening until final fracture at ~2090 MPa. The serrated flow behavior is displayed obviously during plastic deformation after yielding for Cu47Zr48Al4Ni1. The fluctuations was selected randomly due to the similarity of the jerky flow, and the amplitudes of the stress drops were counted, as presented in the inset of Fig. 4. The power-law distribution of the shear avalanches will occur spatiotemporally, suggesting a dynamic transition to the self-organized critical state (SOC), which leads to excellent ductility [25,26]. The stress drop generates the characteristic of a power-law relation apparently and the goodness of fit reaches 0.98, shown in the inset of Fig. 4. With the increase of stress, strains are accommodated elastically until the stress level reaches the value that it can activate a serrated flow. The shear band can be hindered by the B2 phase and the energy will transmit to it, giving rise to the change of stress gradient, which leads to the movement of the dislocations
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Fig. 5. The fractographs of (a) Cu48Zr48Al4 and (b)–(d) Cu47Zr48Al4Ni1 composites upon compression.. Table 1 hardness H of the B2 phase and matrix of as-cast and deformed Cu47Zr48Al4Ni1 composite. H (GPa) As-cast
After deformation
B2 crystal
7.64 ± 0.29
8.64 ± 0.23
Matrix
8.52 ± 0.15
8.26 ± 0.16
composites are conducted, and the corresponding results are shown in Figs. 6 and 7, respectively. In Fig. 6, XRD results indicate that the B2 peaks significantly decrease, and different sharp peaks, which can be identified as the martensite B19' phase, appear in the deformed Cu47Zr48Al4Ni1 composite, indicating that an obvious martensitic transformation does occur during plastic deformation. A comparison of the average hardness of the spherical crystalline phases and amorphous matrix between the as-cast and deformed specimens (Fig. 7) provides a further evidence for the occurrence of the stressinduced martensitic transition. It is clear that the B2 phases are softer than the amorphous phase in the as-cast composite from the result of nanoindentation tests. For the deformed samples, the crystalline B2 phase becomes harder due to the bcc B2→monoclinic B19' martensitic transformation, while the amorphous matrix is softened owing to highdensity shear-band formation. Therefore, it is believed that the observed work hardening is due to hardening of the B2 phases. The specific hardness values of the B2 phases and glass matrix of as-cast and deformed Cu47Zr48Al4Ni1 composites are displayed in Table 1. The mechanical properties of the composites was further evaluated by tensile tests. Fig. 8 shows the tensile engineering stress-strain curves. All composites exist a clear deviation from linearity indicated by bright dotted lines, and have detectable ductility accompanied by work hardening. The plastic strain of all samples increases from 0.09 ± 0.05% to 0.3 ± 0.02% with increasing Ni contents gradually. In contrast, the yielding strength exhibits a decreasing tendency. It is known that the yielding strength of MGMCs is strongly influenced by
Fig. 6. XRD patterns of the as-cast and deformed Cu47Zr48Al4Ni1 composites.
Fig. 7. Microhardness of the crystalline phase and amorphous matrix in the as-cast and deformed Cu47Zr48Al4Ni1 composites.
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ing 1% Ni is deteriorated considerably. In addition, in the Cu48−xZr48Al4Nix (x=0, 0.5, 0.8, 1, 1.5, and 5 at %) alloys, compared with the other constitute elements, the Ni element has the smallest atom radius [41]. With further additions of Ni, atomic size mismatches increase, hindering the long-range diffusion of atoms [40,42]. Thus, the GFA of the present alloys increase, and the volume fractions of crystalline phases decrease in the 5% Ni-added sample distinctly. This conclusion conforms to the atomic size criterion of the three empirical rules proposed by Inoue [43]. Moreover, due to the increasing in the number of constituent elements, the entropy raises, which leads to the density of disordered atoms increasing undoubtedly [44,45]. The increased GFA can be explained by these two factors with further Ni additions from 1.5% to 5%. 4.2. Deformation and fracture mechanisms Fig. 8. Tensile engineering stress-strain curves of the as-cast Cu48−xZr48Al4Nix alloys with different Ni additions.
Compared with most dendrite-reinforced MGMCs, the deformation mechanisms of CuZr-based MGMCs are related to two aspects, namely, martensitic transformation (MT) of the B2 CuZr phase and the interaction between B2 CuZr phase and shear bands. It is known that a disparity of the elastic limit, Young's modulus, and yielding strength exists between B2 phases and amorphous phases [11,14], which consequently leads to the distinct misfit of deformation modes between both phases [46,47]. When subjected to a relatively low stress about 500 MPa, the yielding point of the B2 phase appears, and then, the martensitic transformation occurs [11,12,14]. With the loading further increasing, when the total strain of the B2 crystalline phase reaches the elastic limit of the matrix (about 2%), the matrix enters into the plastic range, accompanied by more severe strain-induced MT phenomenon. Shear bands initiate at the interface between both phases simultaneously, and the composites yield and then enters into the plastic range [14,22]. With the stress increasing, a high density of shear bands form and bifurcate, leading to macroscopic softening. Meanwhile, the degree of MT develops quickly, which results in the formation of a large number of harder martensites. According to previous results [12], the result of the competition and interaction between the MT and formation of shear bands will make the composite possess large plasticity and excellent work-hardening ability. It is believed that the origins of the strong work-hardening capacity after yielding in the current composites may be attributed to the high work-hardening rate of the metastable B2 phases [13,19,24]. Subsequently, when numerous shear bands are clustered together, the microcrack forms easily and propagates, making composites entering into the final fracture stage [48]. Combining stress-strain curves with fractographs under different loading conditions, it is reasonable to rationalize that shear bands initiate at the B2-matrix interface and then rapidly propagate, thus resulting in small tensile ductility. In contrast, the surface of the deformed composites has profuse shear bands, resulting from the strong interaction between B2 phase and shear bands, and the “blocking effect” of B2 phases upon compression [Fig. 5(b) and (c)]. Therefore, the interaction between B2 CuZr phases and shear bands could be regarded as a key factor governing further plastic flows after yielding. In order to better understand the evolution of shear bands in glassy matrices, finite element simulation (FES) was conducted. Fig. 10 shows contour maps of the shear stress at the different strains. Fig. 10(a) displays the maximum shear stress (Tresca stress) field at a global strain of 1.5%. It clearly presents that the shear stress concentrations at the interfaces, owing to the strain misfit between the soft crystalline phase and hard amorphous matrix. The direct results of such stress distributions can be mainly two-fold: i.e. to promote the initiation of shear bands at the interfaces and to effectively hinder the propagation of shear bands by significant stress gradient [35]. Fig. 10(b) shows the maximum shear stress field at a global strain of 2.5%, it clearly shows that the B2 crystalline particles can alter the shear stress field. Stress gradient within the sample increases with increasing strain (556 MPa of 1.5% strain and 1282 MPa of 2.5%), and
the volume fractions of the crystalline and amorphous phases, which can be modeled by a simple rule of mixtures [32–34]:
σcomposite = f crystal σcrystal + f glass σglass
(1)
where f and σ are the volume fractions of the constituent phases and the yielding strength, respectively. With the increase of the volume fraction of soft crystals, the yielding strength decreases. This is consistent with above microstructures. Similar to other MGMCs [8,35,36], an asymmetry for the yield strengths under tension and compression can be observed in the current composites. The SEM images for two composites subjected to tensile loading are displayed in Fig. 9. Fig. 9(a) indicates that the fracture occurs along a shear angle of 57 degrees for the Cu48Zr48Al4 composite, and no discernable shear bands other than few cracks are observed on the lateral surface. In the enlarged area, circled by a white box in Fig. 9(a), it can be seen that the cracks are present with a wavy feature, as displayed in Fig. 9(b), indicating that the propagation of cracks can be inhibited and arrested [9]. Dimples can be found inside the whole surfaces, as presented in the inset of Fig. 9(b). Afterwards, further observations show that dimples can be found on the fracture surfaces of the Cu48Zr48Al4 composite in Fig. 9(c) as well, which is in accordance with the detectable plasticity. Fig. 9(d) shows that the shear fracture of the Cu48Zr47Al4Ni1 composite inclines about 52 degrees and some shear bands appear on the lateral surface, denoted by bright arrows. A high-magnification image for the lateral surface of the fractured Cu47Zr48Al4Ni1 composite shows the formation of multiple shear bands near the cracks, as displayed in Fig. 9(e). The propagation of shear bands is retarded by B2 crystalline phases embedded in the glassy matrix, leading to the deflection of shear bands, as circled by an oval. Besides, ductile dimples, appears, as shown in Fig. 9(f). The observed ductile dimples imply that there is a strong interface between the glassy matrix and the precipitate [7,37]. 4. Discussion 4.1. Effects of Ni on the GFA of the Cu48−xZr48Al4Nix alloys The influence of Ni addition on the GFA of Cu-Zr-Al-Ni alloys is investigated. In general, large negative values of heat of mixing between any two constituent elements in the ternary Cu48Zr48Al4 alloy without Ni addition will enhance the interactions among the components and promote chemical short-range ordering [38], which can improve the local packing efficiency and subsequently restrain the long-range rearrangement of atoms required for crystallization [39]. Nevertheless, when Ni is added to replace Cu, since the heat of mixing of binary Cu-Ni is positive, Cu and Zr, the two major constituents in this system have a high negative heat of mixing [40], which favors the formation of B2 CuZr phases. Therefore the GFA of the alloys contain546
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Fig. 9. SEM images of the tensile fracture surface and its lateral surface of (a)–(c) Cu48Zr48Al4 and (d)–(f) Cu47Zr48Al4Ni1 composites.
the distribution of crystalline phases plays an important role in controlling the mechanical properties [9,13]. To uncover the origins of the phenomenon clearly in this respect, the characteristic material parameter, Rp, which presents the glassy matrix-the plastic zone size ahead of a crack tip is proposed by the yield strength and fracture toughness as:
shear stress concentrate in ligament regions between the neighboring particles. Instead of passing on along the 45 degrees plane, the maximum shear stress turning to nearest particles, which indicates the shear bands propagating among B2 particles, unlike that in the monolithic BMGs [49]. In this way, the B2 crystalline particles can act as strong barriers for the rapid propagation of shear bands, and thus, more shear bands are activated and arrested around B2 particles. The similar shear stress field at strains of 1.5% and 2.5% upon tension is exhibited in Fig. 10(c) and (d), respectively. The shear-band evolution is believed to occur under tension. Yet, all composites failed in a premature manner after having reached a detectable ductility of only below 0.5% upon tension. Specific reasons for the limited ductility upon tension is discussed. Since the formation and distribution of shear bands can be altered between crystalline particles effectively, the spacing of the B2 crystalline particles plays a important role in tuning the shear stress field [48]. Normally, the spacing of B2 crystalline particles is related to two factors, the volume fraction and distribution of B2 phases. With different Ni additions, the volume fractions of B2 phases have significantly changed. High volume fraction of B2 phases implies that large stress concentration emerges at the interface, and thus, multiple shear bands can be activated and promoted. As a result, large plasticity will be accompanied. Nevertheless, the volume fraction of B2 CuZr phases has little effect on the tensile properties. It is rationalized that
2 RP = KIC /2πσy2
(2)
where KIC is the plane-strain fracture toughness [5,50]. It is known that the composites exhibit “asymmetric” mechanical properties between compression and tension like monolithic BMGs [30]. During the active phase of the shear band, shear movement occurs on a plane with some waviness, because both normal tensile stress and normal compressive stress exist at the path of shear-band propagation. Cavities are expected to form at locations with high tensile components, and cavities coalesce into microcacks [51]. On fracture surface, normal tensile stress and shear stress are dominant under uniaxial tension. And normal compressive stress and shear stress dominate under uniaxial compression. The normal tensile stress promotes the propagation of shear bands and the normal compressive stress restrains the propagation of shear bands, respectively [30,49]. The normal tensile stress opens microcacks, which is analogous to mode I type crack in the fracture mechanics. The normal compressive stress closes microcacks and thus shear stress dominates on fracture plane [52], similar to mode II type 547
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Fig. 10. Contour maps of maximum shear stress at different loading conditions obtained by FES: (a) and (b) for strain 1.5% and 2.5% in compression; (c) and (d) for strain 1.5% and 2.5% in tension, respectively.
crack, which makes plastic flows more stable. The mode II type has a larger plastic zone size than that of mode I type. Tandaiya et al. [53] studied mixed mode (I and II) loading of BMGs through detailed finite element simulations, drawing a conclusion that the simulated shear bands are straight and extend over a long distance ahead of the notch tip under mode II loading. And this distance is six times larger than the plastic zone size under mode I. Usually, an increased plastic zone size implies larger plastic dissipation and crack shielding [50]. The shear bands and microcracks evolving from shear bands have almost avoided the B2 phase, thus, the plasticity have not been significantly improved upon tension. In contrast, B2 particles can prevent shear bands from rapid propagation, owing to larger plastic zone size under compression. Therefore, the intrinsic reasons for the “asymmetric” mechanical properties of the present composites between compression and tension are: the critical value of Rp is smaller than the interparticle spacing upon tension, which is attributed to that the B2 phase is easy to grow up quickly and tend to form patch-like framework. Therefore, how to control nucleation and growth of the B2 phase is vital on fabricating MGMCs with uniformly distributed B2 phases. Many methods have been investigated, such as cooling rates (i.e., casting size) [13], superheating (i.e., melt current) [11], and microalloying, among which, the alloying strategy has been proved a most effective approach to homogenize the distribution of B2 crystalline particles [9,54]. As an element with high melting point, minor addition of Ta or Nb [9,54] can be used as the core of heterogeneous nucleation for the B2 crystalline phases. In this study, although the addition of the low melting point element Ni can change the GFA of the Cu48−xZr48Al4Nix markedly, how to uniform the crystalline phase is still a challenge. If the low melting point element and major component elements, such as Cu and Zr to form high-temperature intermetallic compound as heterogeneous nucleation sites and thus may be a feasible way, which requires further study in depth.
alloys were investigated systematically. The microstructural results show that the volume fraction of B2 crystalline phases precipitated within the glassy matrix can be adjusted by suitably adding the Ni contents. Combined experimental results with FES, deformation mechanisms contributing to the improved mechanical properties for the present MGMCs, especially in compression, related to the stressinduced martensitic transformation of the crystalline phase and the strong interaction between the shear bands, cracks and the ductile B2 crystalline phase is stated. while upon tension, the intrinsic reason for the lack of plasticity is that the interparticle spacing and the plastic zone size of the glassy matrix do not satisfy the matching relationship, leading to a rapid propagation of shear bands around the B2 crystalline particles. Consequently, the good tensile plasticity can not be obtained in these composites. Acknowledgement J.W.Q. would like to acknowledge the financial support of National Natural Science Foundation of China (No. 51371122), and the Youth Science Foundation of Shanxi Province, China (No. 2015021005). H.J.Y. would like to acknowledge the financial support from the National Natural Science Foundation of China (No. 51401141) and the Youth Science Foundation of Shanxi Province, China (No. 2014021017-3). Z.H.W. would like to acknowledge the National Natural Science Foundation of China (No. 11390362). References [1] M.K. Miller, P.K. Liaw, Bulk Metallic Glasses, Springer, New York, 2007. [2] C.A. Pampillo, J. Mater. Sci. 10 (1975) 1194–1227. [3] C.T. Liu, L. Heatherly, J.A. Horton, D.S. Easton, C.A. Carmichael, J. Wright, J.L. Schneibel, M.H. Yoo, C.H. Chen, A. Inoue, Metall. Mater. Trans. A 29 (1998) 1811–1820. [4] J.W. Qiao, H.L. Jia, P.K. Liaw, Mater. Sci. Eng. R 100 (2016) 1–69. [5] D.C. Hofmann, J.Y. Suh, A. Wiest, G. Duan, M.L. Lind, M.D. Demetriou, W.L. Johnson, Nature 451 (2008) 1085–1089. [6] F. Szuecs, C.P. Kim, W.L. Johnson, Acta Mater. 49 (2001) 1507–1513. [7] J.W. Qiao, A.C. Sun, E.W. Huang, Y. Zhang, P.K. Liaw, C.P. Chuang, Acta Mater. 59 (2011) 4126–4137. [8] D.C. Hofmann, Science 329 (2010) 1294–1295.
5. Conclusion In summary, effects of Ni additions on the microstructures and mechanical properties of Cu48−xZr48Al4Nix (x=0, 0.5, 0.8, 1, 1.5, and 5) 548
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Designing ductile CuZr-based metallic glass matrix composites a
Y.J. Liu , H.W. Yao Z.H. Wangc
a,b
c
, T.W. Zhang , Z. Wang
a,b
a,b
, Y.S. Wang
a,b,⁎
, J.W. Qiao
a,b,⁎
, H.J. Yang
,
crossmark
a Laboratory of Applied Physics and Mechanics of Advanced Materials, College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, China b Key Laboratory of Interface Science and Engineering in Advanced Materials, Ministry of Education, Taiyuan University of Technology, Taiyuan 030024, China c Institute of Applied Mechanicsand Biomedical Engineering, Taiyuan University of Technology, Taiyuan 030024, China
A R T I C L E I N F O
A BS T RAC T
Keywords: Metallic glasses Composites Work-hardening Microstructure Martensitic phase transformation
The CuZr-based metallic glass matrix composites (MGMCs) with different volume fractions of crystalline phases were designed by doping of nickel. Minor addition of nickel element can change the glass-forming ability of the resultant composites. Large compressive plasticity accompanied by strong work-hardening capacity was achieved in the Cu47Zr48Al4Ni1 composite with a volume fraction of crystalline phases of 33%. The excellent compressive properties were mainly attributed to the inhibition for the propagation of shear bands by the ductile crystals and deformation induced martensitic transformation of the B2-CuZr phase. However, no obvious global tensile ductility was obtained, due to the mode I fracture toughness and small plastic-zone size of glass matrix. To uncover the shear-band evolution during deformation, finite element simulation was conducted, revealing that the ductile B2 phase can tune the shear-stress distribution and consequently initiate and retard shear bands, which stimulates the multiplication of shear bands. Accordingly, the spacing of B2 CuZr particles is a vital factor dominating the plasticity of CuZr-based MGMCs, especially upon tension.
1. Introduction Bulk metallic glasses (BMGs) have attracted large attention because of their unique properties, such as high strength, large elastic limit, excellent corrosion and wear resistances, etc [1]. However, they are known for being extremely brittle, failing in a catastrophic fracture at room temperature, due to formation of localized shear bands [2,3]. To enhance macroscopic plasticity, metallic glass matrix composites (MGMCs) have been developed by in situ or ex situ introduction of secondary phases, which can hinder the rapid propagation of shear bands [4]. In recent years, a series of in-situ dendrite-reinforced MGMCs have been developed, such as Ti-based and Zr-based MGMCs [5–7]. Nevertheless, these MGMCs exhibit a macroscopic strain softening with localized necking due to lack of work hardening in the glass matrix at room temperature, especially under tensile conditions, limiting the applications significantly. Recently, these shortcomings have been successfully addressed in some CuZr-based MGMCs reinforced by ductile B2 CuZr phases [8– 15]. It is worth noting that the B2-type CuZr phases are prone to deform plastically compared to other intermetallic compounds with
complex crystalline structure [16–18], resisting fast propagation of highly localized shear bands effectively and promoting the generation of multiple shear bands in the glass matrices. Moreover, as crystalline phases with a shape memory effect, B2-type CuZr phases can undergo a stress-induced martensitic transformation from B2 phase to a monoclinic B19ʹ phase, which imparts an appreciable work-hardening capability and then compensates the softening of the matrix significantly during plastic flows [9,12–14,19]. In this way, the effect of “transformation-induced plasticity” (TRIP) [11,14,20] can be introduced into the ductile MGMCs. Depending on the TRIP effect, the Zr48Cu47.5Al4Co0.5 composite, containing 25% volume fractioned spherical B2-CuZr phases, presents strong work-hardening capability and considerable tensile ductility [14]. Subsequently, Wu et al. [21] have demonstrated that a minor 0.5 at% addition of Co in the Zr48Cu47.5Al4Co0.5 composites can promote deformation twinning and martensite transformation of B2-CuZr phases by reducing stacking fault energy (SFE). Meanwhile, the SFE of B2 phases is predicted to be strongly reduced by the addition of Ni theoretically. In addition, since the glass-forming ability (GFA) of these BMGs and MGMCs is sensitive to minor additions [9,22,23], an effective adjustment of microstruc-
⁎ Corresponding authors at: Laboratory of Applied Physics and Mechanics of Advanced Materials, College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, China. E-mail addresses: [email protected] (J.W. Qiao), [email protected] (H.J. Yang).
http://dx.doi.org/10.1016/j.msea.2016.11.079 Received 4 August 2016; Received in revised form 19 October 2016; Accepted 22 November 2016 Available online 24 November 2016 0921-5093/ © 2016 Elsevier B.V. All rights reserved.
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tures can be obtained via appropriate trace additions. It is noted that microstructures plays a most effective role in improving the mechanical properties of such alloys [9,13,24]. As such, an improvement of mechanical properties is expected by a minor addition of Ni. However, the effect of Ni addition on the microstructures and mechanical properties of such kinds of MGMCs have been rarely investigated. In this study, a series of CuZr-based MGMCs with different volume fractions of crystalline phases are systematically designed via minor addition of Ni. Compressive and tensile experiments are carried out to investigate the deformation and fracture mechanisms. Moreover, the correlation between the evolution of shear bands and plastic flow can be revealed vividly, combined with finite element simulation (FES). Additionally, dependences of the mechanical properties on microstructures are explored as well. 2. Experimental Fig. 1. XRD spectra of the as-cast Cu48−xZr48Al4Nix alloys with different Ni additions.
Master alloys with normal compositions of Zr48Cu48−xAl4Nix (x=0, 0.5, 0.8, 1, 1.5, and 5 at%) were prepared by arc melting the mixture of constituent elements with purity greater than 99.9% (wt%) under a Tigettered argon atmosphere. The ingots were melted four times to ensure compositional homogeneity. Plate-shape samples with 1.5 mm in thickness and 10 mm in width were prepared by the copper-mouldcasting method. The tensile specimens with gauge dimensions of 10 mm (length)×1.5 mm (width)×1.5 mm (thickness) and compressive specimens with gauge dimensions of 3 mm (length)×1.5 mm (width)×1.5 mm (thickness) were prepared by the electric spark method, respectively. Both compressive and tensile tests were conducted at a strain rate of 5×10−4 s−1 at room temperature. Optical microscope (OM) was used to examine microstructures of as-cast samples. The volume fraction of B2-CuZr phase was determined by an Image-Pro Plus 6.0 software based on the OM images. The phases of the plates were identified by X-ray diffraction (XRD), and the thermal properties were analyzed by differential scanning calorimetry (DSC) at a rate of 20 K/min. Scanning electron microscopy (SEM) was employed to observe the microstructures and fractographs. Nanoindentation was performed on Nano-Indenter II tester (MTS Systems, USA) at room temperature to measure the hardness of the glass matrix and B2 crystalline phases of the as-cast and deformed samples. A Berkovich diamond tip was used. The nominal strain rate was 0.002 s−1 and the maximum indentation depth was 1000 nm. The cross-section of the samples for the nanoindentation tests was carefully polished.
Fig. 2. DSC curves of the as-cast Cu48−xZr48Al4Nix alloys with different Ni additions.
indicates that the volume fraction of the crystalline phase increases. In contrast, with additions of 1.5% and 5% Ni, the exothermic peaks increase, which means that the volume fraction of the amorphous phase increases. The DSC results are consistent with XRD results. Fig. 3 shows the optical cross-section metallographs of as-cast Cu48−xZr48Al4Nix (x=0, 0.5, 1, 1.5, and 5) alloys. For the Ni-free sample, only limited numbers of B2 CuZr particles can be observed, and their sizes are less than 50 µm, as shown in Fig. 3(a). With a 0.5% Ni addition, in addition to some large B2 particles (Fig. 3b), a number of small B2 particles with a diameter of about 5–10 µm can be found, as presents in the inset of Fig. 4(b). The bimodal sized phenomenon of B2 particles has been reported in Ti-based MGMCs as well [19]. For the 1% Ni-added sample, the crystalline volume fraction further increases, as shown in Fig. 4(c). And a large number of B2 CuZr particles can be found throughout the whole section in Fig. 4(f). With further increase of Ni additions, the volume fraction of B2 crystalline phases decreases, specially, for 5% Ni added sample, only a small amount of B2 phase can be found at the out surface of samples. Uniaxial compression tests were carried on the as-cast Cu48−xZr48Al4Nix (x=0, 0.5, 0.8, 1, and 1.5) samples at room temperature, and the corresponding engineering stress-strain curves for these samples are displayed in Fig. 4. The compressive plasticity of the current composites is enhanced with increasing the B2 CuZr crystalline phase. The ternary Cu48Zr48Al4 alloy without Ni addition exhibits a distinct yielding at ~1840 MPa, and has a fracture with the plasticity of ~0.5%. With minor additions of 0.5% and 0.8% Ni, the resultant
3. Results 3.1. Microstructural evolution with Ni additions Fig. 1 shows XRD spectra of as-cast plates of the alloys with different Ni additions. For the ternary Cu48Zr48Al4 alloy, Bragg diffraction peaks corresponding to the body-centered cubic B2 CuZr phases are superimposed on a broad hump. With addition of 0.5–1% Ni, the peak intensity of the crystalline phase increases gradually, indicating an increasing amount of the B2 phase. Specially, this scenario is obvious appreciably for the 1% Ni-added sample, revealing that 1% Ni addition greatly deteriorates the GFA of Cu48−xZr48Al4NiX alloys. However, the volume fractions of B2 CuZr phases decrease indicated by the declined relative intensity of crystalline peaks for the 1.5% Ni-added sample. With further Ni additions (5%), the crystalline peaks become indistinctive visibly, suggesting that excessive Ni additions could enhance the GFA of Cu48−xZr48Al4Ni1 alloys. Fig. 2 presents the DSC curves of these MGMCs. The features of the DSC curves are all similar, namely, an endothermic platform caused by the glass transition and an exothermic heat peak caused for crystallization. The heat release from the crystallization process of these composites decreases gradually with the Ni added from 0% to 1%. This 543
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Fig. 3. cross-section optical metallographs of the as-cast Cu48−xZr48Al4Nix alloys with different Ni additions, (a) x=0, (b) x=0.5, (c) x=1.0, (d) x=1.5, (e) x=5, and (f) low-magnitude image of the 1% Ni added alloys.
[27,28]. With further addition of 1.5% Ni, the composite exhibits a yielding strength of ~1640 MPa, and an uniform plastic strain of ~2%.Fig. 5 presents fractographs of the Cu48Zr48Al4 and Cu47Zr48Al4Ni1 composites upon compression. As for the Cu48Zr48Al4 composite with less B2 phases, it exhibits a shear fracture along a shear angle of 45 degrees with respect to the compressive loading direction in the inset of Fig. 5(a). Normally, the shearing angle of metallic glasses is less than 45 degrees upon compression [29,30]. Only a few primary shear bands are found on the magnified image for the lateral surface of the fractured Cu48Zr48Al4 composite in Fig. 5(a). Fig. 5(b) shows the sideview micrographs of deformed Cu47Zr48Al4Ni1 composite. The fracture plane of the current composite inclines about 45 degrees with the loading direction as well, and some detour shear bands and cracks can be found even at low magnitude in the inset of Fig. 5(b). Highdensity shear bands and a number of martensite phases with slat structures, denoted by the bright arrows, can be observed from the high-magnitude image in Fig. 5(b). In addition, shear bands propagate in a wave-like path as they intersect with the reinforced crystals embedded in the glassy matrix, as shown in the rectangular frame (Fig. 5b), indicating that the propagation of shear bands is hindered by B2 particles effectively. From the detail of the lateral surface for the deformed Cu47Zr48Al4Ni1 composite, the “blocking effect” [13,19] can be clearly observed, as displayed in Fig. 5(c). Profuse shear bands are formed at the interface between B2 phases and the glassy matrices. Afterwards, the fracture surface of the current composite presents dense vein patterns in Fig. 5(d). It is noted that vein patterns are intersected by the torn and elongated crystalline phases, pointed by the bright arrows, suggesting that B2 phases impede the catastrophic fracture, even at the final stage. Such fractographs which contribute to excellent mechanical properties have been found in Ti-based MGMCs [31]. As depicted above, there are martensite phases on the side surface and fracture surface of the Cu47Zr48Al4Ni1 composite, which induces the plasticity improvement. In particular, the pronounced work-hardening capability is attributed to the TRIP effect. In order to further explore the deformation micromechanisms, XRD measurements and nanoindentation tests of the as-cast and deformed Cu47Zr48Al4Ni1
Fig. 4. Compressive engineering stress-strain curves of the as-cast Cu48−xZr48Al4Nix alloys with different Ni additions.
MGMCs exhibit pronounced work-hardening behavior after yielding at ~1710 and ~1720 MPa, respectively, and the uniform plastic strains are ~1.5 and ~2.5%, respectively. For the 1% Ni-added composite, it is clearly shown that the composite yields at ~1580 MPa and undergoes large plasticity of ~11%, together with distinct work hardening until final fracture at ~2090 MPa. The serrated flow behavior is displayed obviously during plastic deformation after yielding for Cu47Zr48Al4Ni1. The fluctuations was selected randomly due to the similarity of the jerky flow, and the amplitudes of the stress drops were counted, as presented in the inset of Fig. 4. The power-law distribution of the shear avalanches will occur spatiotemporally, suggesting a dynamic transition to the self-organized critical state (SOC), which leads to excellent ductility [25,26]. The stress drop generates the characteristic of a power-law relation apparently and the goodness of fit reaches 0.98, shown in the inset of Fig. 4. With the increase of stress, strains are accommodated elastically until the stress level reaches the value that it can activate a serrated flow. The shear band can be hindered by the B2 phase and the energy will transmit to it, giving rise to the change of stress gradient, which leads to the movement of the dislocations
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Fig. 5. The fractographs of (a) Cu48Zr48Al4 and (b)–(d) Cu47Zr48Al4Ni1 composites upon compression.. Table 1 hardness H of the B2 phase and matrix of as-cast and deformed Cu47Zr48Al4Ni1 composite. H (GPa) As-cast
After deformation
B2 crystal
7.64 ± 0.29
8.64 ± 0.23
Matrix
8.52 ± 0.15
8.26 ± 0.16
composites are conducted, and the corresponding results are shown in Figs. 6 and 7, respectively. In Fig. 6, XRD results indicate that the B2 peaks significantly decrease, and different sharp peaks, which can be identified as the martensite B19' phase, appear in the deformed Cu47Zr48Al4Ni1 composite, indicating that an obvious martensitic transformation does occur during plastic deformation. A comparison of the average hardness of the spherical crystalline phases and amorphous matrix between the as-cast and deformed specimens (Fig. 7) provides a further evidence for the occurrence of the stressinduced martensitic transition. It is clear that the B2 phases are softer than the amorphous phase in the as-cast composite from the result of nanoindentation tests. For the deformed samples, the crystalline B2 phase becomes harder due to the bcc B2→monoclinic B19' martensitic transformation, while the amorphous matrix is softened owing to highdensity shear-band formation. Therefore, it is believed that the observed work hardening is due to hardening of the B2 phases. The specific hardness values of the B2 phases and glass matrix of as-cast and deformed Cu47Zr48Al4Ni1 composites are displayed in Table 1. The mechanical properties of the composites was further evaluated by tensile tests. Fig. 8 shows the tensile engineering stress-strain curves. All composites exist a clear deviation from linearity indicated by bright dotted lines, and have detectable ductility accompanied by work hardening. The plastic strain of all samples increases from 0.09 ± 0.05% to 0.3 ± 0.02% with increasing Ni contents gradually. In contrast, the yielding strength exhibits a decreasing tendency. It is known that the yielding strength of MGMCs is strongly influenced by
Fig. 6. XRD patterns of the as-cast and deformed Cu47Zr48Al4Ni1 composites.
Fig. 7. Microhardness of the crystalline phase and amorphous matrix in the as-cast and deformed Cu47Zr48Al4Ni1 composites.
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ing 1% Ni is deteriorated considerably. In addition, in the Cu48−xZr48Al4Nix (x=0, 0.5, 0.8, 1, 1.5, and 5 at %) alloys, compared with the other constitute elements, the Ni element has the smallest atom radius [41]. With further additions of Ni, atomic size mismatches increase, hindering the long-range diffusion of atoms [40,42]. Thus, the GFA of the present alloys increase, and the volume fractions of crystalline phases decrease in the 5% Ni-added sample distinctly. This conclusion conforms to the atomic size criterion of the three empirical rules proposed by Inoue [43]. Moreover, due to the increasing in the number of constituent elements, the entropy raises, which leads to the density of disordered atoms increasing undoubtedly [44,45]. The increased GFA can be explained by these two factors with further Ni additions from 1.5% to 5%. 4.2. Deformation and fracture mechanisms Fig. 8. Tensile engineering stress-strain curves of the as-cast Cu48−xZr48Al4Nix alloys with different Ni additions.
Compared with most dendrite-reinforced MGMCs, the deformation mechanisms of CuZr-based MGMCs are related to two aspects, namely, martensitic transformation (MT) of the B2 CuZr phase and the interaction between B2 CuZr phase and shear bands. It is known that a disparity of the elastic limit, Young's modulus, and yielding strength exists between B2 phases and amorphous phases [11,14], which consequently leads to the distinct misfit of deformation modes between both phases [46,47]. When subjected to a relatively low stress about 500 MPa, the yielding point of the B2 phase appears, and then, the martensitic transformation occurs [11,12,14]. With the loading further increasing, when the total strain of the B2 crystalline phase reaches the elastic limit of the matrix (about 2%), the matrix enters into the plastic range, accompanied by more severe strain-induced MT phenomenon. Shear bands initiate at the interface between both phases simultaneously, and the composites yield and then enters into the plastic range [14,22]. With the stress increasing, a high density of shear bands form and bifurcate, leading to macroscopic softening. Meanwhile, the degree of MT develops quickly, which results in the formation of a large number of harder martensites. According to previous results [12], the result of the competition and interaction between the MT and formation of shear bands will make the composite possess large plasticity and excellent work-hardening ability. It is believed that the origins of the strong work-hardening capacity after yielding in the current composites may be attributed to the high work-hardening rate of the metastable B2 phases [13,19,24]. Subsequently, when numerous shear bands are clustered together, the microcrack forms easily and propagates, making composites entering into the final fracture stage [48]. Combining stress-strain curves with fractographs under different loading conditions, it is reasonable to rationalize that shear bands initiate at the B2-matrix interface and then rapidly propagate, thus resulting in small tensile ductility. In contrast, the surface of the deformed composites has profuse shear bands, resulting from the strong interaction between B2 phase and shear bands, and the “blocking effect” of B2 phases upon compression [Fig. 5(b) and (c)]. Therefore, the interaction between B2 CuZr phases and shear bands could be regarded as a key factor governing further plastic flows after yielding. In order to better understand the evolution of shear bands in glassy matrices, finite element simulation (FES) was conducted. Fig. 10 shows contour maps of the shear stress at the different strains. Fig. 10(a) displays the maximum shear stress (Tresca stress) field at a global strain of 1.5%. It clearly presents that the shear stress concentrations at the interfaces, owing to the strain misfit between the soft crystalline phase and hard amorphous matrix. The direct results of such stress distributions can be mainly two-fold: i.e. to promote the initiation of shear bands at the interfaces and to effectively hinder the propagation of shear bands by significant stress gradient [35]. Fig. 10(b) shows the maximum shear stress field at a global strain of 2.5%, it clearly shows that the B2 crystalline particles can alter the shear stress field. Stress gradient within the sample increases with increasing strain (556 MPa of 1.5% strain and 1282 MPa of 2.5%), and
the volume fractions of the crystalline and amorphous phases, which can be modeled by a simple rule of mixtures [32–34]:
σcomposite = f crystal σcrystal + f glass σglass
(1)
where f and σ are the volume fractions of the constituent phases and the yielding strength, respectively. With the increase of the volume fraction of soft crystals, the yielding strength decreases. This is consistent with above microstructures. Similar to other MGMCs [8,35,36], an asymmetry for the yield strengths under tension and compression can be observed in the current composites. The SEM images for two composites subjected to tensile loading are displayed in Fig. 9. Fig. 9(a) indicates that the fracture occurs along a shear angle of 57 degrees for the Cu48Zr48Al4 composite, and no discernable shear bands other than few cracks are observed on the lateral surface. In the enlarged area, circled by a white box in Fig. 9(a), it can be seen that the cracks are present with a wavy feature, as displayed in Fig. 9(b), indicating that the propagation of cracks can be inhibited and arrested [9]. Dimples can be found inside the whole surfaces, as presented in the inset of Fig. 9(b). Afterwards, further observations show that dimples can be found on the fracture surfaces of the Cu48Zr48Al4 composite in Fig. 9(c) as well, which is in accordance with the detectable plasticity. Fig. 9(d) shows that the shear fracture of the Cu48Zr47Al4Ni1 composite inclines about 52 degrees and some shear bands appear on the lateral surface, denoted by bright arrows. A high-magnification image for the lateral surface of the fractured Cu47Zr48Al4Ni1 composite shows the formation of multiple shear bands near the cracks, as displayed in Fig. 9(e). The propagation of shear bands is retarded by B2 crystalline phases embedded in the glassy matrix, leading to the deflection of shear bands, as circled by an oval. Besides, ductile dimples, appears, as shown in Fig. 9(f). The observed ductile dimples imply that there is a strong interface between the glassy matrix and the precipitate [7,37]. 4. Discussion 4.1. Effects of Ni on the GFA of the Cu48−xZr48Al4Nix alloys The influence of Ni addition on the GFA of Cu-Zr-Al-Ni alloys is investigated. In general, large negative values of heat of mixing between any two constituent elements in the ternary Cu48Zr48Al4 alloy without Ni addition will enhance the interactions among the components and promote chemical short-range ordering [38], which can improve the local packing efficiency and subsequently restrain the long-range rearrangement of atoms required for crystallization [39]. Nevertheless, when Ni is added to replace Cu, since the heat of mixing of binary Cu-Ni is positive, Cu and Zr, the two major constituents in this system have a high negative heat of mixing [40], which favors the formation of B2 CuZr phases. Therefore the GFA of the alloys contain546
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Fig. 9. SEM images of the tensile fracture surface and its lateral surface of (a)–(c) Cu48Zr48Al4 and (d)–(f) Cu47Zr48Al4Ni1 composites.
the distribution of crystalline phases plays an important role in controlling the mechanical properties [9,13]. To uncover the origins of the phenomenon clearly in this respect, the characteristic material parameter, Rp, which presents the glassy matrix-the plastic zone size ahead of a crack tip is proposed by the yield strength and fracture toughness as:
shear stress concentrate in ligament regions between the neighboring particles. Instead of passing on along the 45 degrees plane, the maximum shear stress turning to nearest particles, which indicates the shear bands propagating among B2 particles, unlike that in the monolithic BMGs [49]. In this way, the B2 crystalline particles can act as strong barriers for the rapid propagation of shear bands, and thus, more shear bands are activated and arrested around B2 particles. The similar shear stress field at strains of 1.5% and 2.5% upon tension is exhibited in Fig. 10(c) and (d), respectively. The shear-band evolution is believed to occur under tension. Yet, all composites failed in a premature manner after having reached a detectable ductility of only below 0.5% upon tension. Specific reasons for the limited ductility upon tension is discussed. Since the formation and distribution of shear bands can be altered between crystalline particles effectively, the spacing of the B2 crystalline particles plays a important role in tuning the shear stress field [48]. Normally, the spacing of B2 crystalline particles is related to two factors, the volume fraction and distribution of B2 phases. With different Ni additions, the volume fractions of B2 phases have significantly changed. High volume fraction of B2 phases implies that large stress concentration emerges at the interface, and thus, multiple shear bands can be activated and promoted. As a result, large plasticity will be accompanied. Nevertheless, the volume fraction of B2 CuZr phases has little effect on the tensile properties. It is rationalized that
2 RP = KIC /2πσy2
(2)
where KIC is the plane-strain fracture toughness [5,50]. It is known that the composites exhibit “asymmetric” mechanical properties between compression and tension like monolithic BMGs [30]. During the active phase of the shear band, shear movement occurs on a plane with some waviness, because both normal tensile stress and normal compressive stress exist at the path of shear-band propagation. Cavities are expected to form at locations with high tensile components, and cavities coalesce into microcacks [51]. On fracture surface, normal tensile stress and shear stress are dominant under uniaxial tension. And normal compressive stress and shear stress dominate under uniaxial compression. The normal tensile stress promotes the propagation of shear bands and the normal compressive stress restrains the propagation of shear bands, respectively [30,49]. The normal tensile stress opens microcacks, which is analogous to mode I type crack in the fracture mechanics. The normal compressive stress closes microcacks and thus shear stress dominates on fracture plane [52], similar to mode II type 547
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Fig. 10. Contour maps of maximum shear stress at different loading conditions obtained by FES: (a) and (b) for strain 1.5% and 2.5% in compression; (c) and (d) for strain 1.5% and 2.5% in tension, respectively.
crack, which makes plastic flows more stable. The mode II type has a larger plastic zone size than that of mode I type. Tandaiya et al. [53] studied mixed mode (I and II) loading of BMGs through detailed finite element simulations, drawing a conclusion that the simulated shear bands are straight and extend over a long distance ahead of the notch tip under mode II loading. And this distance is six times larger than the plastic zone size under mode I. Usually, an increased plastic zone size implies larger plastic dissipation and crack shielding [50]. The shear bands and microcracks evolving from shear bands have almost avoided the B2 phase, thus, the plasticity have not been significantly improved upon tension. In contrast, B2 particles can prevent shear bands from rapid propagation, owing to larger plastic zone size under compression. Therefore, the intrinsic reasons for the “asymmetric” mechanical properties of the present composites between compression and tension are: the critical value of Rp is smaller than the interparticle spacing upon tension, which is attributed to that the B2 phase is easy to grow up quickly and tend to form patch-like framework. Therefore, how to control nucleation and growth of the B2 phase is vital on fabricating MGMCs with uniformly distributed B2 phases. Many methods have been investigated, such as cooling rates (i.e., casting size) [13], superheating (i.e., melt current) [11], and microalloying, among which, the alloying strategy has been proved a most effective approach to homogenize the distribution of B2 crystalline particles [9,54]. As an element with high melting point, minor addition of Ta or Nb [9,54] can be used as the core of heterogeneous nucleation for the B2 crystalline phases. In this study, although the addition of the low melting point element Ni can change the GFA of the Cu48−xZr48Al4Nix markedly, how to uniform the crystalline phase is still a challenge. If the low melting point element and major component elements, such as Cu and Zr to form high-temperature intermetallic compound as heterogeneous nucleation sites and thus may be a feasible way, which requires further study in depth.
alloys were investigated systematically. The microstructural results show that the volume fraction of B2 crystalline phases precipitated within the glassy matrix can be adjusted by suitably adding the Ni contents. Combined experimental results with FES, deformation mechanisms contributing to the improved mechanical properties for the present MGMCs, especially in compression, related to the stressinduced martensitic transformation of the crystalline phase and the strong interaction between the shear bands, cracks and the ductile B2 crystalline phase is stated. while upon tension, the intrinsic reason for the lack of plasticity is that the interparticle spacing and the plastic zone size of the glassy matrix do not satisfy the matching relationship, leading to a rapid propagation of shear bands around the B2 crystalline particles. Consequently, the good tensile plasticity can not be obtained in these composites. Acknowledgement J.W.Q. would like to acknowledge the financial support of National Natural Science Foundation of China (No. 51371122), and the Youth Science Foundation of Shanxi Province, China (No. 2015021005). H.J.Y. would like to acknowledge the financial support from the National Natural Science Foundation of China (No. 51401141) and the Youth Science Foundation of Shanxi Province, China (No. 2014021017-3). Z.H.W. would like to acknowledge the National Natural Science Foundation of China (No. 11390362). References [1] M.K. Miller, P.K. Liaw, Bulk Metallic Glasses, Springer, New York, 2007. [2] C.A. Pampillo, J. Mater. Sci. 10 (1975) 1194–1227. [3] C.T. Liu, L. Heatherly, J.A. Horton, D.S. Easton, C.A. Carmichael, J. Wright, J.L. Schneibel, M.H. Yoo, C.H. Chen, A. Inoue, Metall. Mater. Trans. A 29 (1998) 1811–1820. [4] J.W. Qiao, H.L. Jia, P.K. Liaw, Mater. Sci. Eng. R 100 (2016) 1–69. [5] D.C. Hofmann, J.Y. Suh, A. Wiest, G. Duan, M.L. Lind, M.D. Demetriou, W.L. Johnson, Nature 451 (2008) 1085–1089. [6] F. Szuecs, C.P. Kim, W.L. Johnson, Acta Mater. 49 (2001) 1507–1513. [7] J.W. Qiao, A.C. Sun, E.W. Huang, Y. Zhang, P.K. Liaw, C.P. Chuang, Acta Mater. 59 (2011) 4126–4137. [8] D.C. Hofmann, Science 329 (2010) 1294–1295.
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