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Designing Bulk Metallic Glass Composites with Enhanced Formability and Plasticity
Accepted Manuscript Designing Bulk Metallic Glass Composites with Enhanced Formability and Plasticity Y. Wu , H. Wang , X.J. Liu , X.H. Chen , X.D. Hui , Y. Zhang , Z.P. Lu , Ph.D
PII:
S1005-0302(14)00091-7
DOI:
10.1016/j.jmst.2014.03.028
Reference:
JMST 353
To appear in:
Journal of Materials Science & Technology
Received Date: 14 February 2014 Accepted Date: 10 March 2014
Please cite this article as: Y. Wu, H. Wang, X.J. Liu, X.H. Chen, X.D. Hui, Y. Zhang, Z.P. Lu, Designing Bulk Metallic Glass Composites with Enhanced Formability and Plasticity, Journal of Materials Science & Technology (2014), doi: 10.1016/j.jmst.2014.03.028. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Designing Bulk Metallic Glass Composites with Enhanced Formability and Plasticity Y. Wu, H. Wang, X.J. Liu, X.H. Chen, X.D. Hui, Y. Zhang, Z.P. Lu*
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State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, China [Manuscript received February 14, 2014, in revised form March 10, 2014] Ph.D.; Tel.: +86
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82375387; E-mail address:
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*Corresponding author. Prof., [email protected] (Z.P. Lu).
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To address the main stumbling-block of bulk metallic glasses (BMGs), i.e., room temperature brittleness, designing BMG matrix composites has been attracted extensive attention. Up to date, BMG composites in various alloy systems have been successfully developed by forming crystalline phases embedded in the amorphous matrix through either ex-situ or in-situ methods. In this paper, a brief review of our recent work in this topic will be presented and the novel approaches to improving composite formability and mechanical properties will also be highlighted. The main purpose of this manuscript is not to offer a comprehensive review of all the BMG composites, but instead focuses will be placed on illustrating recently developed advanced BMG composites including Fe-based BMG composite with no metalloids, Al-based BMG composite and BMG composites reinforced by the TRIP (transformation-induced plasticity) effects. The basic ideas and related mechanisms underlying the development of these novel BMG composites will be discussed.
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KEY WORDS: Bulk metallic glasses; Designing; Formability; Plasticity
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1. Introduction
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Bulk metallic glasses (BMGs) have shown a unique combination of mechanical, chemical and physical properties[1–5], but their room-temperature brittleness and strain softening behavior has been one of the stumbling blocks to real structural applications[6,7]. To answer this challenge, the concept of developing glassy composite microstructures by forming crystalline phases embedded in the glassy matrix through either in-situ or ex-situ methods have been developed[8–21]. Initially, BMG composites were formed mainly due to partial crystallization of fully glassy precursors via annealing experiments. In the late 1990s, it was recognized that the brittle failure of BMGs could be mitigated with additions of crystalline phases in the BMG matrix. Gradually, formation of appropriate crystalline phases in BMG matrices has become a toughening strategy. By adjusting processing parameters and/or alloy compositions, various crystalline phases could be either in-situ or ex-situ formed inside BMGs, and large compressive plasticity was successfully achieved[18–21]. The basic idea underlying this strategy is to manipulate formation of shear bands (i.e., nucleation and propagation of shear bands) during deformation by utilizing the inclusions as reinforcing media. Compared with the amorphous matrix, most reinforced crystalline inclusions were ductile with a lower elastic modulus and hardness. As a result, these reinforcing phases were designed to yield before the amorphous matrix and thus bear more plastic deformation[18–21]. Nevertheless, a few BMG composites with brittle intermetallic compounds were also found to be capable of improving the plasticity[22]. Early work was mostly focused on enhancing compressive plasticity of BMG materials[20–22]. In recent years, by properly controlling the preparation routes and alloy compositions, the microstructural length scale (i.e., the dendrite spacing of the primary phase) was adjusted to match the mechanical length scale (i.e., the plastic shielding of an opening crack tip), large tensile ductility and superior fracture toughness have been achieved in a series of Zr- and Ti-based BMG composites[23–26]. In this approach, the precipitated dendrites are usually thermodynamic stable phase, which endows the composite technically more processable, e.g., the characteristics of the crystalline inclusions can be controlled by processing techniques such as “semi-solid” formation and Bridgman solidification[26–28]. In-situ formed BMG composites usually have good crystal-amorphous matrix interface and a relative simple fabrication process, but compromises among glass-formation ability (GFA), composition and cooling rates have to be made. In contrast, features of crystalline inclusions in ex-situ formed BMG composites, such as distribution and volume fraction, can be easily adjusted. However, the fabrication processes are complex and crystal-amorphous matrix interface is always the weak zone during deformation. In the late 1990s, various attempts were conducted to develop ex-situ formed BMG composites by infiltration casting with crystalline wires or powders to create two-phase mixtures[10,11]. For instance, additions of W fibers/balls, porous Ti or particles of Ti, Mo and SiC into Zr-based and Mg-based BMGs could lead to remarkable increment in the compressive plasticity[44–47], and the improvement was attributed to the slowdown of shear band propagation due to the ex-situ added reinforcements. In these reported approaches, nevertheless, the added second phases are usually distributed discontinuously with significant inhomogeneity in the BMG matrix,
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2. In-situ Formed BMG Composites
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which greatly reduce the reinforcing effectiveness[48]. It is important to point out that tensile ductility in these ex-situ BMG composites was still absent. The concept of “transformation-induced plasticity” (TRIP) has been proven to be effective in enhancing tensile ductility and toughness in crystalline materials, such as steels and ceramics[29,30]. In recent years, to further improve tensile ductility and work-hardening capability of BMGs, TRIP effect was also introduced and was confirmed to be a novel way to improve both tensile ductility and work-hardening ability in BMGs[31–40]. Up to date, several TRIP-reinforced BMG composites systems have been developed, mostly in CuZr-based BMG[35–40] and Ti-based BMG[41–43]. Formation, mechanical behavior and microstructure evolution during deformation in this type of BMG composites have been preliminarily studied[35–43], but the deformation mechanism, especially the details about contribution of martensitic transformation to the macroscopic properties, remain unclear. Nevertheless, this TRIP concept did prompt extensive attention in developing novel BMG composites with advanced properties, such as Ti- and Fe-based BMGs[41–43]. In this paper, several special approaches developed recently in our group for designing advanced BMG composites with either enhanced formability or mechanical properties were summarized. The intention is to highlight basic ideas and related mechanism underlying development of these novel BMG composites, including Fe-based BMG composite with no metalloids, Al-based BMG composite and TRIP reinforced BMG composites.
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2.1. Large-sized BMG composites fabricated by phase separation Phase separation has been found in metallic glasses and could be used to improve the deformation plasticity[49–51]. Recently, we used the concept of phase separation to design BMG composites[52]. Firstly, a known BMG with a good GFA, e.g., element A-based AxByCzDuEv (x + y + z + u + v = 100 at.%), was selected as the starting composition, and then an element M having a positive enthalpy of mixing and a miscibility gap with the main component A, was chosen and added. As a result, two liquids, i.e. the M-rich region and M depleted region, will form at high temperatures. Due to the phase separation, the M-depleted region would have a composition close to the starting A-based alloy, which has a good GFA. During solidification, it can be thus anticipated that crystalline phases will be precipitated out of the M-rich liquid, while the M-depleted liquid will vitrify. As such, an M-based BMG composite consisting of the M-rich crystalline phases embedded in the A-based BMG matrix can be fabricated. The total volume fraction of the BMG matrix can be adjusted by the amount of the M addition[52]. This approach can be widely used to design special BMG composites based on currently known BMG systems and thus optimize overall properties of BMG materials. 1) Development of centimeter sized BMG composites with high Fe concentration It was widely recognized that the metalloid elements (i.e. B, C, Si and P) are key glass-forming constituents in Fe-based alloy systems[53–55], and a good Fe-based glass-former often contains a metalloid content of 20 at.% or above[53–55]. Nevertheless, it was found that the high content of the metalloid elements usually bring in side effects on the soft magnetic
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properties and plastic deformation at room temperature in these Fe-based BMGs[56,57]. Based on the concept of phase separation, we have successfully developed a series of BMG matrix composites with high Fe concentration but without any metalloid elements. The Fe element which has a positive mixing enthalpy with La and Ce, was added in the BMG composition La32.5Ce32.5Co15Al10Cu10, as shown in Fig. 1(a). Due to the repulsion between the Fe and (La, Cu), phase separation occurred and the melt initially separated into two liquids, i.e, the Fe-rich and the Fe-depleted. During cooling, Fe-rich phases, i.e., CeFe2 and Ce(Fe, Co)2 were firstly formed and the residual composition was adjusted near to our starting composition, i.e. the LaCe-based bulk glass former, and finally solidified into glass phase. Therefore, a composite structure with a large percentage of Fe was fabricated (Fig. 1(b)). The current BMG matrix composites show no significant improvement in plastic deformation, which is probably attributable to the brittle primary phases and the weak interfaces between the primary phases and the glassy matrix. Nevertheless, this work provides a useful basis for developing large sized BMG composites with a large formability and deformability. 2) Centimeter sized BMG composites with a high Al content Unlike that of the other metal-based alloy systems such as Zr-, Ti-, and Cu-based BMGs, the maximal attainable thickness of Al-based metallic glasses still cannot exceed 1 mm[58–61]. The formation of Al-based metallic glasses or the corresponding glass matrix composites is limited to thin ribbons which seriously restrict the utilization of their high specific strength and low density[58–61]. Based on the aforementioned concept, we developed a series of Al–(Co,Cu)–(La,Ce)–Fe BMG matrix composites which contain a high Al content from 35 to 60 at.%[62]. It was found that additions of Fe, which has a positive heat of mixing with the main constitute elements La and Ce, enhanced the volume fraction of glass matrix by suppressing the precipitation of the Al(La,Ce) phase. The strength of these glassy composites reaches up to 1.2 GPa which is twice that of conventional Al-based alloys. These composites possessing extraordinarily large formability with a critical size over 15 mm were produced by common copper mold casting (as shown in Fig. 2). For the alloy containing 30 at.% Al, the BMG composites mainly consisted of the Al2(La,Ce) phase and glass matrix, and the volume fraction of the glass matrix is about 80%. For the alloy containing 30–50 at.%Al, the BMG composites mainly consisted of the Al2(La,Ce) phase, the Al(La,Ce) phase and glass matrix, and the volume fraction of glass matrix is about 30%–60%. Proper addition of the Fe element effectively suppressed the precipitation of the Al(La,Ce) phase and consequently enhanced the volume fraction of glass matrix due to the occurrence of phase separation. The optimum content of the Fe addition varies from 5 to 0.5 at.% as the Al content changed from 50 to 55 at.%. For the other elements such as Ti and Zr, which also have a positive heat of mixing with La and Ce, their addition induced formation of new ternary AlTiCo and AlZrCo phases and consumed most of the glass-forming element Co, thus severely deteriorating formation of the glass matrix. 2.2. BMG composites reinforced by transformation-induced plasticity To make BMG composites viable for engineering applications, both tensile ductility and
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work-hardening capability are needed. By applying the concept of transformation-induced plasticity (TRIP), BMG composites with both large tensile ductility and significant work-hardening capability have been developed[31–43]. This section will give a brief review of formation, mechanical behavior and mechanical properties optimum of TRIP-reinforced BMG composites.
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1) Formation of TRIP-reinforced BMG composites Up to date, the phase transformation used in the TRIP-reinforced BMG composites includes three types: 1) from the B2-CuZr (PM3M lattice structure) to the B19’-CuZr phase (P21/M) or the supercell B33-CuZr (CM) in CuZr-based BMG composites[63,64], 2) from the bcc metastable β-phase to body-centered tetragonal α’ and/or hexagonal close-packed (hcp) α’ martensite phases[41] in TiZr-based BMG composites, and 3) from B2-Ti(CuNi) to B19’-Ti(CuNi) in Ti(CuNi)-based BMG composites[42,43]. The main challenge of forming TRIP-reinforced BMG composites is that most of the austenitic phases utilized in BMG composites are metastable at room temperature. For example, the B2-CuZr phase is stable at high temperatures above 988 K but tends to decompose into Cu10Zr7 and Zr2Cu at low temperatures[65]. Many studies have been conducted to investigate formation of TRIP-reinforced BMG composites from thermodynamic and kinetic point of view[66–68]. The key point is to fully understand phase completion and formation kinetics during solidification, which can be elucidated by the model CuZrAl system[38]. It was found that Al additions not only alter the GFA, but also promote the austenite transformation and destabilize the martensite CuZr phase by decreasing the martensitic transformation temperature Ms, as shown Fig. 3(a). As such, the combined effects of cooling rates and Al addition can be schematically analyzed from the framework of time-temperature-transformation (TTT) diagram based on the competitive formation mechanism[69], as demonstrated in Fig. 3(b). For the alloys < 2% Al with lower thermodynamic stability, the main competing crystalline phase is B2-CuZr which is stable at temperatures above 988 K[65]. To form a fully glassy structure, the cooling rate has to be faster enough to bypass the nose of the B2-CuZr TTT curve (i.e., the critical cooling rate Rc). Once the applied cooling rate is slower than Rc, the B2-CuZr phase would precipitate out of the liquid at high temperatures. Due to the Ms value is relatively close to the glass transition temperature Tg[38,70], the supercooled parent B2-CuZr phase can easily transform into the martensitic CuZr phase via the martensitic transformation, resulting in a composite structure consisting of the martensitic CuZr phases embedded in the amorphous matrix. For the alloys with appropriate amount of Al (i.e., 3%–8% Al, as indicated by solid lines in Fig. 3(b)), formation of the B2-CuZr phase is retarded because of necessary redistribution of the Al atoms during solidification[69]. As a result, the TTT curve of the B2 phase moves toward the longer time, leading to reduction in the critical cooling rate and hence high GFA. Meanwhile, formation of the Al2Zr phase at high temperatures becomes possible due to high amount of Al, although B2-CuZr is still the primary competing phase. Moreover, the martensitic transformation is also suppressed because of the stabilization of the parent B2-CuZr phase, i.e., a decreased Ms value. Therefore, at a suitable cooling rate, the supercooled B2-CuZr phase can be retained in the amorphous matrix without occurrence of
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the matensitic transformation. With the further increase of the Al content to above 8%, the formation of the Al2Zr phase tends to become predominant and precipitation of the B2-CuZr phase is further suppressed (see dashed lines in Fig. 3(b)). Consequently, glass formation is demoted due to the strong forming tendency of the Al2Zr phase. Upon slow cooling, Al2Zr always precipitates first and BMG composites containing the single B2-CuZr phase as reinforcement cannot be obtained anymore.
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2) Effects of features of reinforcing crystalline phase on mechanical properties of BMG composites Morphology characteristics (e.g., shape and distribution) of the crystalline phases were found to have significant effects on mechanical properties of TRIP-reinforced BMG composites[71–73]. If brittle intermetallics phase such as Al2Zr precipitated along with the ductile austenitic phase, the mechanical properties of the composite were deteriorated[38]. If the austenitic phase is the single reinforcement, macroscopic yield strength of the composite was mainly determined by its volume fraction. However, the plasticity and work-hardening behavior were governed by not only its volume fraction, but also its distribution[71–73]. With the increase of the austenite, yield strength decreased and for the composites with a crystalline volume fraction higher than 50%, characteristic of the resultant stress-strain curve was similar to that of the conventional shape memory alloy[72,73]. Moreover, even for the BMG composites with a similar crystalline volume fraction, the fracture strain may vary tremendously due to different distribution of the reinforcing phase[71]. As an example, Fig. 4 shows the stress-strain curves of three BMG composites with ~10% B2 phase for the Cu48Zr48Al4 BMG composite. Surprisingly, the fracture strain changes dramatically from 6% to 22%. SEM (scanning electron microscopy) characterization confirmed that these specimens had a similar volume fraction of the crystalline phase (typical examples are given in Fig. 4), nevertheless, specimen A has the fewest crystals with the largest size, whilst B and C have more tiny crystals. In addition, the tiny crystalline particles are distributed more homogeneously in specimen C than in specimen B. As can be seen, homogeneous distribution of the B2 phase gave rise to a large fracture strain, manifesting that distribution of the reinforced crystals also play important roles in achieving macroscopic plasticity of the current type of BMG composites. To fabricate BMG composites with the homogeneously distributed austenitic phase, controlling nucleation and growth of the B2 phase was the key. Many factors influencing its distribution have been investigated, such as cooling rates (i.e., casting size), superheating (i.e., melt current)[74,75], and microalloying of Ta, Sn and Nb to form high-temperature phases as heterogeneous nucleation sites for the austenitic crystalline phases[76–78]. It might be possible to use annealing or semi-solid processing to partially crystallize a monolithic BMG or homogenize the crystallite distribution, but these processes are difficult to control due to the metastability of the austenitic phase[79]. Among all the proposed approaches, alloying strategy may be a feasible way to develop tough TRIP-BMG composites. 3) Reinforcing mechanism Origins of the enhanced mechanical properties in the TRIP-reinforced BMG composites
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can be attributed to the following two aspects. The first one is the “blocking effect” from heterogeneities at multiple length scales[80]. As indicated in Fig. 5, shear bands during deformation were deflected at the interface between the crystals and glassy matrix, and then divided into many tiny secondary shear bands around the spherical phases. The crystalline spheres acted as strong barriers for the rapid propagation of shear bands and thus effectively enhanced plasticity of the composites. Moreover, the “blocking effect” was also observed for not only the micro-meter sized B2-CuZr phases, but also the nano-meter sized B2-CuZr ones (Fig. 5(c)). Therefore, the “blocking effect” from both micro-meter and nano-meter spherical phases would enhance the plasticity of the composite. The other important aspect is the contribution from the martensitic transformation of B2-CuZr. As shown in Fig. 6, experimental results evidenced that the phase transformation did occurred during deformation of the TRIP-reinforced BMG composite[33,39]. As the deformation proceeds, the crystalline B2 phase starts to transform into B19’ (Fig. (6)) and become harder and harder (Fig. 6(b)) while the amorphous matrix is softened due to extensive shear-band formation. It is reasonable to rationalize that the hardening of B2 will compensate softening of the amorphous matrix and contributes to the global mechanical properties of the composite. According to the literature[81], phase transformation of the austenite B2 spherical crystals during deformation could release the stress concentration around them and restrict free volume accumulation. As a result, the rapid propagation of shear bands can be hindered, which requires a further stress to move the shear bands and consequently inhibits the early necking and work-softening.
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4) Optimizing BMG composites by tailoring stacking fault energy Due to the important role of martensitic transformation in the global deformation behavior of the TRIP-BMG composites, ductilizing BMG composites by tailoring the martensitic transformation via stacking fault energy (SFE) becomes possible[40,82]. Generally, a low SFE tends to facilitate deformation twinning to initiate and thus stimulates the martensitic transformation of the austenite B2 structure[83,84]. By systematically investigating the alloying effects in the Zr48Cu48–xAl4Mx (x=0–2 at.%, M=Co, Ti, Fe, Ni, Ta, Cr, Ga, Zn, Hf, Nb, Ta and Ag) system, it was found that stacking fault energy of the crystalline phase, which usually governs plastic deformation carriers (i.e., dislocation and twinning) in crystalline materials, also can be tailored to ductilize the TRIP-BMG composites. For the bcc B2-CuZr phase in the ZrCuAlM system, the minimum SFE occurred along the (011)[100] slip system among all the possible lattice planes investigated. Some alloying additions, such as Co, dramatically reduce SFE, but some other elements, such as Ti, significantly increase it (Fig. 7(a)). The onset temperature (Ms) was also found to be correlated with SFE. The smaller the SFE, the higher the Ms, confirming that proper alloying additions can facilitate the martensitic transformation[40,82]. In the case of a similar crystal volume fraction, it was found that the TRIP-BMG composite with a smaller SFE yields at a relatively low stress but undergoes a pronounced work-hardening after yielding and a larger plastic deformation, and vice versa for the BMG composite with a higher SFE (Fig. 7(b)). This trend demonstrates that the minor additions reducing SFE of B2-CuZr can remarkably improve the tensile ductility and work-hardening capability of the TRIP-BMG composite.
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Variation of the SFE usually has its origins in the electronic structures, and particularly relate to the charge density change during the stacking fault formation[85,86]. From Fig. 8(a), it can be seen that the doped element that increases the electron charge density redistribution also increases the SFE, which is consistent with the tendency in other alloy systems[85,86]. Furthermore, it is found that both the electronegativity and the atomic radius difference between the doped element and substituted element correlate closely with the SFE of the doped B2 structure, as demonstrated in Fig. 8(b). Clearly, a smaller electronegativity and atomic radius difference between the doping element and Cu lead to a lower SFE of the doped B2-CuZr phase. As a result, the martensitic transformation of the B2 crystals doped with suitable elements that reduce the SFE could occur much easier, which eventually results in an earlier yielding, a larger plasticity and more pronounced work-hardening capability.
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Comparing with those of in-situ formed BMG composite, volume fraction and distribution of foreign inclusions in ex-situ formed BMG composites can be easily adjusted, and consequently preparation of this type of BMG composites is also an effective way to ductilize glassy alloys. In this section, two approaches recently developed were introduced.
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3.1. Continuous formation of metallic glass coated wires and their bulk composites The size of BMG composites prepared by conventional casting processes, such as melt quenching, pressure injection and infiltration casting[9,87] is usually restricted to centimeter to decimeter scale. Recently, a continuous processing method to prepare MG coated composite wires was developed in our group[88,89]. The sketch of the continuous infiltration system for producing BMG coated wires is shown in Fig. 9(a). The system consists of a vacuum, heat, cooling, and motor drive unit. Alloy ingots of BMGs were placed in a crucible and heated to the melt temperature. Tungsten wires or other metallic wires with different diameters were polished and cleaned, and then was fed continuously from a series of supply wheels. It is to be noted that multiple wires can be fed simultaneously. These metallic wires passed through preheating unit, and then immersed into molten metal for infiltration at a tunable drawing velocity. As the wires were drawn out the molten pool, it was cooled forcibly by argon. Eventually, BMG composite rods with a diameter around 0.5 mm were prepared continuously, as shown in Fig. 9(b). SEM images (Figs. 9(b) and (c)) indicted that no contrast corresponding to the crystalline phase can be found at the interface between the amorphous matrix and wire (Figs. 9(d) and (e)), indicating that no reactant was formed during the infiltration process[89]. Furthermore, the composite wires can be weaved at the supercooled liquid region to form bulk form for structural applications, as shown in Fig. 9(f)[88]. Fig. 10(a) shows the uniaxial tension curves of the BMG composite rod with 61.4 vol.% tungsten wires, in comparison with that of the monolithic BMG sample and tungsten wire. It can be seen that the composite have a larger fracture strength and improved plastic deformation. Multiple shear bands can be seen on the lateral surface near the fracture section (Fig. 10(b)). It is inevitable that global mechanical properties of the BMG composite would be affected by properties of the metallic wires used. By properly choosing type, volume
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fraction and diameter of metallic wires used in this approach, mechanical behavior of the BMG composites can effectively be tuned.
4. Summary
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3.2. BMG composites reinforced by three dimensional metallic network Although ex-situ formed BMG composites have been widely studied[44–47], the added second inclusions are usually distributed discontinuously with significant inhomogeneity in the BMG matrix, which greatly reduce the reinforcing effectiveness[48]. Reinforcements with a continuous three-dimensional networked structure are expected to suppress the rapid propagation of shear bands more effectively. To test this hypothesis, a Cu foam with a porosity content of 95.8% and a pore radius ranging from 200 to 250 µm was chosen. The dimension of networked ligament is approximately 12 mm in diameter. The Ti40Zr25Cu12Ni3Be20 amorphous alloy has a melting point of 983 K[90], which is far below the melting point of Cu, was chosen as the amorphous matrix. The Cu foam was initially preloaded in the Cu mold, melted master alloy was subsequently sucked into the Cu mold under argon atmosphere and then solidified to form the final composite. In spite of the small volume fraction of Cu foam (~4.2%), yield strength of the composite slightly decreased while plasticity dramatically increased comparing with the monolithic BMG, as shown in Fig. 11[91]. By analyzing elastic behavior and fracture energy dissipation, effects of the ductile Cu network on the deformation of the composite can be summarized below. When shear bands are formed and propagate in the amorphous phase, they are confined by the individual compartment, which is surrounded by the Cu cell-wall. In other words, long-range shear band propagation is prohibited. The ductile Cu can temporarily block the shear bands and plastically deform to dissipate the kinetic energy. Additionally, the elastic mismatch at the interface between the amorphous matrix and cell-wall can lead to stress concentration and initiate secondary shear bands. The bifurcation of shear bands would certainly contribute to the improvement of plasticity.
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To make BMGs viable for practical applications as engineering materials, their damage tolerance and plastic deformation capability under tension (i.e., toughness, tensile ductility and work-hardening) have to be improved. As elaborated above, our recent research work on BMG composites was briefly highlighted. To fabricate in-situ formed BMG composites, we have applied different alloy design methodologies, i.e., phase separation and transformation-induced plasticity. To develop ex-situ formed BMG composites, we not only developed a new technique for continuously producing BMG coated wires but also fabricated ex-situ formed BMG composites with 3D metallic network inclusions. Our results indicate that development of advanced BMG composites with optimized overall properties provide a promising route to utilize BMG materials for practical engineering applications. Acknowledgements This research was supported in part by the National Natural Science Foundation of China (Nos. 51010001, 51371003, 51001009 and 50901006), 111 Project (No. B07003) and Program for Changjiang Scholars and Innovative Research Team in University. Y. Wu and X.J. Liu acknowledge the financial support from “the Fundamental Research Funds for the Central Universities”.
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References [1] W. L. Johnson, MRS Bull. 24 (1999) 42–56. [2] A. Inoue, Acta Mater. 48 (2000) 279–306. [3] C.A. Schuh, T.C. Hufnagel, U. Ramamurty, Acta Mater. 55 (2007) 4067–4109. [4] M.F. Ashby, A.L. Greer, Scripta Mater. 54 (2006) 321–326. [5] M.W. Chen, Annu. Rev. Mater. Res. 38 (2008) 445–469. [6] Z.F. Zhang, J. Eckert, L. Schultz, Acta Mater. 51 (2003) 1167–1179. [7] F.F. Wu, Z.F. Zhang, S.X. Mao, A. Peker, J. Eckert, Phys. Rev. B 75 (2007) 134201. [8] H.A. Bruck, A.J. Rosakis, W.L. Johnson, J. Mater. Res. 11 (1996) 503–511. [9] R.B. Dandliker, R.D. Conner, W.L. Johnson, J. Mater. Res. 13 (1998) 2896–2901. [10] H. Choi-Yim, R. Busch, W.L. Johnson, J. Appl. Phys. 83 (1998) 7993–7997. [11] H. Choi-Yim, R.D. Conner, F. Szuecs, W.L. Johnson, Acta Mater. 50 (2002) 2737–2745. [12] Y.F. Xue, H.N. Cai, L. Wang, F.C. Wang, H.F. Zhang, Z.Q. Hu, Compos. Sci. Technol. 68 (2008) 3396–3400. [13] Y.C. Kim, E. Fleury, J.C. Lee, D.H. Kim, J. Mater. Res. 20 (2005) 2474–2479. [14] F. Szuecs, C.P. Kim, W.L. Johnson, Acta Mater. 49 (2001) 1507–1513. [15] M.L. Lee, Y. Li, C.A. Schuh, Acta Mater. 52 (2004) 4121–4131. [16] H. Tan, Y. Zhang, Y. Li, Intermetallics 10 (2002) 1203–1205. [17] U. Kuhn, J. Eckert, N. Mattern, L. Schultz, Mater. Sci. Eng. A 375–377 (2004) 322–326. [18] C. Fan, R.T. Ott, T.C. Hufnagel, Appl. Phys. Lett. 81 (2002) 1020–1022. [19] U. Kühn, J. Eckert, N. Mattern, L. Schultz, Appl. Phys. Lett. 80 (2002) 2478–2480. [20] C.C. Hays, C.P. Kim, W.L. Johnson, Phys. Rev. Lett. 84 (2000) 2901. [21] C. Fan, H.Q. Li, L.J. Kecskes, K. X. Tao, H. Choo, P. K. Liaw, C.T. Liu, Phys. Rev. Lett. 96 (2006) 145506. [22] G. Chen, H. Bei, Y. Cao, A. Gali, C.T. Liu, E.P. George, Appl. Phys. Lett. 95 (2009) 081908. [23] D.C. Hofmann, J.Y. Suh, A. Wiest, G. Duan, M.L. Lind, M.D. Demetrious, W.L. Johnson, Natrue 451 (2008) 1085–1089. [24] D.C. Hofmann, J.Y. Suh, A. Wiest, M.L. Lind, M.D. Demetriou, W.L. Johnson, Proc. Natl. Acad. Sci. U.S.A. 105 (2008) 20136. [25] Z. Zhu, H. Zhang, Z. Hu, W. Zhang, A. Inoue, Scripta Mater. 62 (2010) 278–281. [26] J.W. Qiao, S. Wang, Y. Zhang, P.K. Liaw, G.L. Chen, Appl. Phys. Lett. 94 (2009) 151905. [27] D.C. Hofmann, J.Y. Suh, A. Wiest, W. Johnson, Scripta Mater. 59 (2008) 684–687. [28] D.C. Hofmann, H. Kozachkov, H.E. Khalifa, J.P. Schramm, M.D. Demetriou, K.S. Vecchio, W.L. Johnson, JOM 61 (2009) 11–17. [29] V. F. Zackay, E.R. Parker, D. Fahr, R. Busch, Trans. ASM 60 (1967) 252–259. [30] P.J. Jacques, Q. Fumemont, F. Lani, T. Pardoen, F. Delannay, Acta Mater. 55 (2007) 3681–3693. [31] J. Das, M.B. Tang, K.B. Kim, R. Theissmann, F. Baier, W.H. Wang, J. Eckert, Phys. Rev. Lett. 94 (2005) 205501. [32] K.B. Kim, J. Das, S. Venkataraman, S. Yi, J. Eckert, Appl. Phys. Lett. 89 (2006) 071908. [33] S. Pauly, J. Das, J. Bednarcik, N. Mattern, K.B. Kim, D.H. Kim, J. Eckert, Scripta Mater. 60 (2009) 431–434.
ACCEPTED MANUSCRIPT
AC C
EP
TE D
M AN U
SC
RI PT
[34] Y.F. Sun, B.C. Wei, Y.R. Wang, W.H. Li, T.L. Cheung, C.K. Shek, Appl. Phys. Lett. 87 (2005) 051905. [35] S. Pauly, J. Das, C. Duhamel, J. Eckert, Adv. Eng. Mater. 9 (2007) 487–491. [36] S. Pauly, G. Liu, G. Wang, U. Kuhn, N. Mattern, J. Eckert, Acta Mater. 57 (2009) 5445–5453. [37] S. Pauly, G. Liu, S. Gorantla, G. Wang, U. Kuhn, D.H. Kim, J. Eckert, Acta Mater. 58 (2010) 4883–4890. [38] Y. Wu, H. Wang, H.H. Wu, Z.Y. Zhang, X.D. Hui, G.L. Chen, D. Ma, X.L. Wang, Z.P. Lu, Acta Mater. 59 (2011) 2928–2936. [39] Y. Wu, Y.H. Xiao, G.L. Chen, C.T. Liu, Z.P. Lu, Adv. Mater. 22 (2010) 2770–2773. [40] Y. Wu, D.Q. Zhou, W.L. Song, H. Wang, Z.Y. Zhang, D. Ma, X.L. Wang, Z.P. Lu, Phys. Rev. Lett. 109 (2012) 245506. [41] C.P. Kim, Y.S. Oh, S.H. Lee, N.J. Kim, Scripta Mater. 65 (2011) 304–307. [42] P. Gargarella, S. Pauly, K.K. Song, J. Hu, N.S. Barekar, M.S. Khoshkhoo, A. Teresiak, H. Wendrock, U. Kuhn, C. Ruffing, E. Kerscher, J. Eckert, Acta Mater. 61 (2013) 151–162. [43] Y.S. Oh, C.P. Kim, S.H. Lee, N.J. Kim, Acta Mater. 59 (2011) 7277–7286. [44] J.A. Hines, K.S. Vecchio, Acta Mater. 45 (1997) 635–649. [45] B. Zhang, H.M. Fu, Z.W. Zhu, A.M. Wang, H. Li, C. Dong, Mater. Sci. Eng. A 504 (2012) 207–211. [46] H.F. Zhang, H. Li, A.M. Wang, H.M. Fu, B.Z. Ding, Z.Q. Hu, Intermetallics 17 (2009) 1070–1077. [47] Y. Sun, H.F. Zhang, A.M. Wang, H.M. Fu, Z.Q. Hu, C.E. Wen, Appl. Phys. Lett. 95 (2009) 171910. [48] J.S.C. Jang, T.H. Li, S.R. Jian, J.C. Huang, T.G. Nieh, Intermetallics 19 (2011) 738–743. [49] B.J. Park, H.J. Chang, D.H. Kim, W.T. Kim, K. Chattopadhyay, T.A. Abinandanan, S. Bhattacharyya, Phys. Rev. Lett. 96 (2006) 245503. [50] E.S. Park, D.H. Kim, Acta Mater. 54 (2006) 2597–2604. [51] D.H. Kim, W.T. Kim, E.S. Park, N. Mattern, J. Eckert, Prog. Mater. Sci. 58 (2013) 1103–1172. [52] J.E. Gao, Z.P. Chen, Q. Du, H.X. Li, Y. Wu, H. Wang, X.J. Liu, Z.P. Lu, Acta Mater. 61 (2013) 3214–3223. [53] Z.P. Lu, C.T. Liu, J.R. Thomepson, W.D. Porter, Phys. Rev. Lett. 92 (2004) 245503. [54] V. Ponnambalam, S.J. Poon, G.J. Shiflet, J. Mater. Res. 19 (2004) 1320–1323. [55] A. Inoue, T. Zhang, A. Takeuchi, Appl. Phys. Lett. 71 (1997) 464. [56] A.R. Yavari, Mater. Sci. Eng. 98 (1988) 491–493. [57] X.J. Gu, S.J. Poon, G.J. Shiflet, W. Widom, Appl. Phys. Lett. 92 (2008) 161910. [58] S.J. Poon, G.J. Shiflet, Science 241 (1988) 1640–1642. [59] Z.P. Chen, J.E. Gao, Y. Wu, H.X. Li, H. Wang, Z.P. Lu, J. Alloy. Compd. 513 (2012) 387–392. [60] H. Yang, J.Q. Wang, Y. Li, J. Non-Cryst. Solids 354 (2008) 3473–3479. [61] F.Q. Guo, S.J. Enouf, S.J. Poon, G.J. Shiflet, Philos. Mag. Lett. 81 (2001) 203–211. [62] Z.P. Chen, J.E. Gao, Y. Wu, H. Wang, X.J. Liu, Z.P. Lu, Sci. Rep. 3 (2013) 3353. [63] D. Schryvers, G.S. Firstov, J.W. Seo, J.V. Humbeeck, Y.N. Koval, Scripta Mater. 36 (1997) 1119–1125.
ACCEPTED MANUSCRIPT
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SC
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[64] J.W. Seo, D. Schryvers, Acta Mater. 46 (1998) 1165–1175. [65] T. Abe, M. Shimono, M. Ode, H. Onodera, Acta Mater. 54 (2006) 909–915. [66] K.K. Song, S. Pauly, Y. Zhang, P. Gargarella, R. Li, N.S. Barekar, U. Kuhn, M. Stoica, J. Eckert, Acta Mater. 17 (2011) 6620–6630 [67] K.K. Song, S. Pauly, B.A. Sun, Y. Zhang, J. Tan, U. Kuhn, M. Stoica, J. Eckert, Intermetallics 30 (2012) 132–138. [68] K.K. Song, S. Pauly, Y. Zhang, B.A. Sun, J. He, G.Z. Ma, U. Kuhn, J. Eckert, Mater. Sci. Eng. A 559 (2013) 711–718. [69] X.P. Lu, D. Ma, C.T. Liu, Y.A. Chang, Intermetallics 15 (2007) 253–259. [70] Y.N. Koval, G.S. Firstov, A.V. Kotko, Scripta Metall. Mater. 27 (1992) 1611–1616. [71] Y. Wu, W.L. Song, Z.Y. Zhang, X.D. Hui, D. Ma, X.L. Wang, X.C. Shang, Z.P. Lu, Chin. Sci. Bull. 56 (2011) 3960–3964. [72] K.K. Song, S. Pauly, Y. Zhang, R. Li, S. Gorantla, N. Narayanan, U. Kuhn, T. Gemming, J. Eckert, Acta Mater. 60 (2012) 6000–6012. [73] S. Pauly, G. Liu, G. Wang, J. Das, K.B. Kim, U. Kuhn, D.H. Kim, J. Eckert, Appl. Phys. Lett. 95 (2009) 101906. [74] S. Pauly, G. Liu, G. Wang, U. Kuhn, N. Mattern, J. Eekert, Acta Mater. 57 (2009) 5445–5453. [75] Z.Q. Liu, R. Li, G. Liu, K.K. Song, S. Pauly, T. Zhang, J. Eckert, AIP Adv. 2 (2012) 032176. [76] Z.Q. Liu, R. Li, G. Liu, W.H. Su, H. Wang, Y. Li, M.J. Shi, X.K. Luo, G.J. Wu, T. Zhang, Acta Mater. 60 (2012) 3128–3139. [77] Z.Y. Zhang, Y. Wu, J. Zhou, W.L. Song, D. Cao, H. Wang, X.J. Liu, Z.P. Lu, Intermetallics 42 (2013) 68–76. [78] S. Pauly, J. Das, J. Bednarcik, N. Mattern, K.B. Kim, D.H. Kim, J. Eckert, Scripta Mater. 60 (2009) 431–434. [79] H. Kozachkov, J. Kolodziejska, W.L. Johnson, D.C. Hofmann, Intermetallics 39 (2013) 89–93. [80] D.V. Louzguine-Luzgin, A. Vinogradov, G.Q. Xie, S. Li, A. Lazarev, S. Hashimoto, A. Inoue, Philos. Mag. 89 (2009) 2887–2901. [81] J. Das, K.B. Kim, W. Xu, B.C. Wei, Z.F. Zhang, W.H. Wang, S. Yi, J. Eckert, Mater. Trans. 47 (2006) 2606–2609. [82] D.Q. Zhou, Y. Wu, H. Wang, X.D. Hui, X.J. Liu, Z.P. Lu, Comput. Mater. Sci. 79 (2013) 187–192. [83] D.C. Hoffman, Science 329 (2010) 1294–1295. [84] X.Q. Wang, Phys. Rev. B 78 (2008) 092103. [85] S. Ogata, J. Li, S. Yip, Science 298 (2002) 807–811. [86] Y. Qi, R.K. Mishra, Phys. Rev. B 75 (2007) 224105. [87] H. Choi-Yim, W.L. Johnson, Appl. Phys. Lett. 71 (1997) 3808–3810. [88] X.H. Chen, B.Y. Zhang, G.L. Chen, Y. Zhang, X.D. Hui, Z.P. Lu, X.J. Liu, Y. Xu, X.R. Xing, Intermetallics 18 (2010) 2034–2038. [89] B.Y. Zhang, X.H. Chen, S.S. Wang, D.Y. Lin, X.D. Hui, Mater. Lett. 93 (2013) 210–214. [90] G. Wang, J. Shen, J.F. Sun, Y.J. Huang, J. Zou, Z.P. Lu, Z.H. Stachurski, B.D. Zhou, J. Non-Cryst. Solids 351 (2005) 209–217.
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Fig. 1 (a) Critical thickness for the BMG matrix composites as a function of alloy compositions in the Fe–Co–La–Ce–Al–Cu system; (b) schematic illustration of the solidification process of the Fe–Co–La–Ce–Al–Cu alloys[52].
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Fig. 2 (a) Critical thickness for BMG matrix composites as a function of alloy compositions for the Al–(CoCu)–(La,Ce)–Fe system, and (b) outer appearance of the as-cast Al40(CoCu)15(LaCe)45 BMG matrix composites with a diameter of 15 mm, and (c) DSC curves of the rods with a diameter of 15 mm for the Al35(CoCu)15(LaCe)50, Al40(CoCu)15(LaCe)45, Al40(CoCu)10(LaCe)50, Al45(CoCu)10(LaCe)45 and Al45(CoCu)5(LaCe)50 alloys[62].
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Fig. 3 DSC heating and cooling curves of the (Cu0.5Zr0.5)100–xAlx (x=0, 4, 6) alloy ingots showing the martensitic and austenitic transformation (a), and schematic time-temperature-transformation diagram of the (Cu0.5Zr0.5)100–xAlx alloys with a different Al content (b)[38]. Fig. 4 Compressive true stress-strain curves of the Cu48Zr48Al4 alloy with about 10% B2 phase, and SEM images of the three representative samples with a similar crystal volume fraction of B2 but a different distribution pattern[71].
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Fig. 5 (a) SEM image of the lateral surface, (b) blowup of the lateral surface, and (c) TEM image of the tensile fractured sample of the Cu48Zr48Al4 BMG composite[38]. Fig. 6 XRD patterns of the as-cast and fractured samples of the Cu48Zr48Al4Co0.5 BMG composite (a), micro-hardness of the crystalline phase and amorphous matrix in the as-cast and tensioned samples (b)[39].
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Fig. 7 Stacking fault energy on the (011)[100] slip system of the B2-CuZr phase and the ones with half of Cu on the slip plane substituted by Co or Ti (a), and true tensile stress-strain curves of the Zr48Cu48Al4, Zr48Cu47.5Al4Co0.5 and Zr48Cu47.5Al4Ti0.5 BMG composites (b)[40]. Fig. 8 SFE values of the B2-CuZr phase doped with different alloying elements as a function of the electronic density charge redistribution on the (011)[100] slip system (a), and dependence of the stacking fault energy of the doped B2 phase on the eletronegativity and radius difference between the doping and substituted elements (b)[40]. Fig. 9 (a) Schematic illustration of the newly developed continuous infiltration process; (b) a bunch of BMG composite wires prepared, (c) and (d) SEM images of the cross-section of BMG composites containing 61.4 vol% tungsten wires, (e) TEM image of the interface between the tungsten and the amorphous matrix shown in d, and (f) bulk sample weaved with BMG-coated tungsten wire composite[88,89].
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Fig. 10 Quasi-static tensile stress–strain curves of the BMG composite wires containing 61.4 vol.% tungsten wires in comparison with that of the tungsten wire and monolithic BMG Vit 1 (a), and side image near fracture surface[89].
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Fig. 11 Compressive stress-strain curves of the monolithic Ti-based BMG and the fabricated BMG composite (a), as compared with that of the Cu foam (b). Insets are the enlargements of the plastic deformation curves and appearances of the pre-strained Cu foams, respectively[91].
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PII:
S1005-0302(14)00091-7
DOI:
10.1016/j.jmst.2014.03.028
Reference:
JMST 353
To appear in:
Journal of Materials Science & Technology
Received Date: 14 February 2014 Accepted Date: 10 March 2014
Please cite this article as: Y. Wu, H. Wang, X.J. Liu, X.H. Chen, X.D. Hui, Y. Zhang, Z.P. Lu, Designing Bulk Metallic Glass Composites with Enhanced Formability and Plasticity, Journal of Materials Science & Technology (2014), doi: 10.1016/j.jmst.2014.03.028. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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Designing Bulk Metallic Glass Composites with Enhanced Formability and Plasticity Y. Wu, H. Wang, X.J. Liu, X.H. Chen, X.D. Hui, Y. Zhang, Z.P. Lu*
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State Key Laboratory for Advanced Metals and Materials, University of Science and Technology Beijing, Beijing 100083, China [Manuscript received February 14, 2014, in revised form March 10, 2014] Ph.D.; Tel.: +86
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82375387; E-mail address:
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*Corresponding author. Prof., [email protected] (Z.P. Lu).
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To address the main stumbling-block of bulk metallic glasses (BMGs), i.e., room temperature brittleness, designing BMG matrix composites has been attracted extensive attention. Up to date, BMG composites in various alloy systems have been successfully developed by forming crystalline phases embedded in the amorphous matrix through either ex-situ or in-situ methods. In this paper, a brief review of our recent work in this topic will be presented and the novel approaches to improving composite formability and mechanical properties will also be highlighted. The main purpose of this manuscript is not to offer a comprehensive review of all the BMG composites, but instead focuses will be placed on illustrating recently developed advanced BMG composites including Fe-based BMG composite with no metalloids, Al-based BMG composite and BMG composites reinforced by the TRIP (transformation-induced plasticity) effects. The basic ideas and related mechanisms underlying the development of these novel BMG composites will be discussed.
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KEY WORDS: Bulk metallic glasses; Designing; Formability; Plasticity
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1. Introduction
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Bulk metallic glasses (BMGs) have shown a unique combination of mechanical, chemical and physical properties[1–5], but their room-temperature brittleness and strain softening behavior has been one of the stumbling blocks to real structural applications[6,7]. To answer this challenge, the concept of developing glassy composite microstructures by forming crystalline phases embedded in the glassy matrix through either in-situ or ex-situ methods have been developed[8–21]. Initially, BMG composites were formed mainly due to partial crystallization of fully glassy precursors via annealing experiments. In the late 1990s, it was recognized that the brittle failure of BMGs could be mitigated with additions of crystalline phases in the BMG matrix. Gradually, formation of appropriate crystalline phases in BMG matrices has become a toughening strategy. By adjusting processing parameters and/or alloy compositions, various crystalline phases could be either in-situ or ex-situ formed inside BMGs, and large compressive plasticity was successfully achieved[18–21]. The basic idea underlying this strategy is to manipulate formation of shear bands (i.e., nucleation and propagation of shear bands) during deformation by utilizing the inclusions as reinforcing media. Compared with the amorphous matrix, most reinforced crystalline inclusions were ductile with a lower elastic modulus and hardness. As a result, these reinforcing phases were designed to yield before the amorphous matrix and thus bear more plastic deformation[18–21]. Nevertheless, a few BMG composites with brittle intermetallic compounds were also found to be capable of improving the plasticity[22]. Early work was mostly focused on enhancing compressive plasticity of BMG materials[20–22]. In recent years, by properly controlling the preparation routes and alloy compositions, the microstructural length scale (i.e., the dendrite spacing of the primary phase) was adjusted to match the mechanical length scale (i.e., the plastic shielding of an opening crack tip), large tensile ductility and superior fracture toughness have been achieved in a series of Zr- and Ti-based BMG composites[23–26]. In this approach, the precipitated dendrites are usually thermodynamic stable phase, which endows the composite technically more processable, e.g., the characteristics of the crystalline inclusions can be controlled by processing techniques such as “semi-solid” formation and Bridgman solidification[26–28]. In-situ formed BMG composites usually have good crystal-amorphous matrix interface and a relative simple fabrication process, but compromises among glass-formation ability (GFA), composition and cooling rates have to be made. In contrast, features of crystalline inclusions in ex-situ formed BMG composites, such as distribution and volume fraction, can be easily adjusted. However, the fabrication processes are complex and crystal-amorphous matrix interface is always the weak zone during deformation. In the late 1990s, various attempts were conducted to develop ex-situ formed BMG composites by infiltration casting with crystalline wires or powders to create two-phase mixtures[10,11]. For instance, additions of W fibers/balls, porous Ti or particles of Ti, Mo and SiC into Zr-based and Mg-based BMGs could lead to remarkable increment in the compressive plasticity[44–47], and the improvement was attributed to the slowdown of shear band propagation due to the ex-situ added reinforcements. In these reported approaches, nevertheless, the added second phases are usually distributed discontinuously with significant inhomogeneity in the BMG matrix,
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2. In-situ Formed BMG Composites
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which greatly reduce the reinforcing effectiveness[48]. It is important to point out that tensile ductility in these ex-situ BMG composites was still absent. The concept of “transformation-induced plasticity” (TRIP) has been proven to be effective in enhancing tensile ductility and toughness in crystalline materials, such as steels and ceramics[29,30]. In recent years, to further improve tensile ductility and work-hardening capability of BMGs, TRIP effect was also introduced and was confirmed to be a novel way to improve both tensile ductility and work-hardening ability in BMGs[31–40]. Up to date, several TRIP-reinforced BMG composites systems have been developed, mostly in CuZr-based BMG[35–40] and Ti-based BMG[41–43]. Formation, mechanical behavior and microstructure evolution during deformation in this type of BMG composites have been preliminarily studied[35–43], but the deformation mechanism, especially the details about contribution of martensitic transformation to the macroscopic properties, remain unclear. Nevertheless, this TRIP concept did prompt extensive attention in developing novel BMG composites with advanced properties, such as Ti- and Fe-based BMGs[41–43]. In this paper, several special approaches developed recently in our group for designing advanced BMG composites with either enhanced formability or mechanical properties were summarized. The intention is to highlight basic ideas and related mechanism underlying development of these novel BMG composites, including Fe-based BMG composite with no metalloids, Al-based BMG composite and TRIP reinforced BMG composites.
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2.1. Large-sized BMG composites fabricated by phase separation Phase separation has been found in metallic glasses and could be used to improve the deformation plasticity[49–51]. Recently, we used the concept of phase separation to design BMG composites[52]. Firstly, a known BMG with a good GFA, e.g., element A-based AxByCzDuEv (x + y + z + u + v = 100 at.%), was selected as the starting composition, and then an element M having a positive enthalpy of mixing and a miscibility gap with the main component A, was chosen and added. As a result, two liquids, i.e. the M-rich region and M depleted region, will form at high temperatures. Due to the phase separation, the M-depleted region would have a composition close to the starting A-based alloy, which has a good GFA. During solidification, it can be thus anticipated that crystalline phases will be precipitated out of the M-rich liquid, while the M-depleted liquid will vitrify. As such, an M-based BMG composite consisting of the M-rich crystalline phases embedded in the A-based BMG matrix can be fabricated. The total volume fraction of the BMG matrix can be adjusted by the amount of the M addition[52]. This approach can be widely used to design special BMG composites based on currently known BMG systems and thus optimize overall properties of BMG materials. 1) Development of centimeter sized BMG composites with high Fe concentration It was widely recognized that the metalloid elements (i.e. B, C, Si and P) are key glass-forming constituents in Fe-based alloy systems[53–55], and a good Fe-based glass-former often contains a metalloid content of 20 at.% or above[53–55]. Nevertheless, it was found that the high content of the metalloid elements usually bring in side effects on the soft magnetic
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properties and plastic deformation at room temperature in these Fe-based BMGs[56,57]. Based on the concept of phase separation, we have successfully developed a series of BMG matrix composites with high Fe concentration but without any metalloid elements. The Fe element which has a positive mixing enthalpy with La and Ce, was added in the BMG composition La32.5Ce32.5Co15Al10Cu10, as shown in Fig. 1(a). Due to the repulsion between the Fe and (La, Cu), phase separation occurred and the melt initially separated into two liquids, i.e, the Fe-rich and the Fe-depleted. During cooling, Fe-rich phases, i.e., CeFe2 and Ce(Fe, Co)2 were firstly formed and the residual composition was adjusted near to our starting composition, i.e. the LaCe-based bulk glass former, and finally solidified into glass phase. Therefore, a composite structure with a large percentage of Fe was fabricated (Fig. 1(b)). The current BMG matrix composites show no significant improvement in plastic deformation, which is probably attributable to the brittle primary phases and the weak interfaces between the primary phases and the glassy matrix. Nevertheless, this work provides a useful basis for developing large sized BMG composites with a large formability and deformability. 2) Centimeter sized BMG composites with a high Al content Unlike that of the other metal-based alloy systems such as Zr-, Ti-, and Cu-based BMGs, the maximal attainable thickness of Al-based metallic glasses still cannot exceed 1 mm[58–61]. The formation of Al-based metallic glasses or the corresponding glass matrix composites is limited to thin ribbons which seriously restrict the utilization of their high specific strength and low density[58–61]. Based on the aforementioned concept, we developed a series of Al–(Co,Cu)–(La,Ce)–Fe BMG matrix composites which contain a high Al content from 35 to 60 at.%[62]. It was found that additions of Fe, which has a positive heat of mixing with the main constitute elements La and Ce, enhanced the volume fraction of glass matrix by suppressing the precipitation of the Al(La,Ce) phase. The strength of these glassy composites reaches up to 1.2 GPa which is twice that of conventional Al-based alloys. These composites possessing extraordinarily large formability with a critical size over 15 mm were produced by common copper mold casting (as shown in Fig. 2). For the alloy containing 30 at.% Al, the BMG composites mainly consisted of the Al2(La,Ce) phase and glass matrix, and the volume fraction of the glass matrix is about 80%. For the alloy containing 30–50 at.%Al, the BMG composites mainly consisted of the Al2(La,Ce) phase, the Al(La,Ce) phase and glass matrix, and the volume fraction of glass matrix is about 30%–60%. Proper addition of the Fe element effectively suppressed the precipitation of the Al(La,Ce) phase and consequently enhanced the volume fraction of glass matrix due to the occurrence of phase separation. The optimum content of the Fe addition varies from 5 to 0.5 at.% as the Al content changed from 50 to 55 at.%. For the other elements such as Ti and Zr, which also have a positive heat of mixing with La and Ce, their addition induced formation of new ternary AlTiCo and AlZrCo phases and consumed most of the glass-forming element Co, thus severely deteriorating formation of the glass matrix. 2.2. BMG composites reinforced by transformation-induced plasticity To make BMG composites viable for engineering applications, both tensile ductility and
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work-hardening capability are needed. By applying the concept of transformation-induced plasticity (TRIP), BMG composites with both large tensile ductility and significant work-hardening capability have been developed[31–43]. This section will give a brief review of formation, mechanical behavior and mechanical properties optimum of TRIP-reinforced BMG composites.
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1) Formation of TRIP-reinforced BMG composites Up to date, the phase transformation used in the TRIP-reinforced BMG composites includes three types: 1) from the B2-CuZr (PM3M lattice structure) to the B19’-CuZr phase (P21/M) or the supercell B33-CuZr (CM) in CuZr-based BMG composites[63,64], 2) from the bcc metastable β-phase to body-centered tetragonal α’ and/or hexagonal close-packed (hcp) α’ martensite phases[41] in TiZr-based BMG composites, and 3) from B2-Ti(CuNi) to B19’-Ti(CuNi) in Ti(CuNi)-based BMG composites[42,43]. The main challenge of forming TRIP-reinforced BMG composites is that most of the austenitic phases utilized in BMG composites are metastable at room temperature. For example, the B2-CuZr phase is stable at high temperatures above 988 K but tends to decompose into Cu10Zr7 and Zr2Cu at low temperatures[65]. Many studies have been conducted to investigate formation of TRIP-reinforced BMG composites from thermodynamic and kinetic point of view[66–68]. The key point is to fully understand phase completion and formation kinetics during solidification, which can be elucidated by the model CuZrAl system[38]. It was found that Al additions not only alter the GFA, but also promote the austenite transformation and destabilize the martensite CuZr phase by decreasing the martensitic transformation temperature Ms, as shown Fig. 3(a). As such, the combined effects of cooling rates and Al addition can be schematically analyzed from the framework of time-temperature-transformation (TTT) diagram based on the competitive formation mechanism[69], as demonstrated in Fig. 3(b). For the alloys < 2% Al with lower thermodynamic stability, the main competing crystalline phase is B2-CuZr which is stable at temperatures above 988 K[65]. To form a fully glassy structure, the cooling rate has to be faster enough to bypass the nose of the B2-CuZr TTT curve (i.e., the critical cooling rate Rc). Once the applied cooling rate is slower than Rc, the B2-CuZr phase would precipitate out of the liquid at high temperatures. Due to the Ms value is relatively close to the glass transition temperature Tg[38,70], the supercooled parent B2-CuZr phase can easily transform into the martensitic CuZr phase via the martensitic transformation, resulting in a composite structure consisting of the martensitic CuZr phases embedded in the amorphous matrix. For the alloys with appropriate amount of Al (i.e., 3%–8% Al, as indicated by solid lines in Fig. 3(b)), formation of the B2-CuZr phase is retarded because of necessary redistribution of the Al atoms during solidification[69]. As a result, the TTT curve of the B2 phase moves toward the longer time, leading to reduction in the critical cooling rate and hence high GFA. Meanwhile, formation of the Al2Zr phase at high temperatures becomes possible due to high amount of Al, although B2-CuZr is still the primary competing phase. Moreover, the martensitic transformation is also suppressed because of the stabilization of the parent B2-CuZr phase, i.e., a decreased Ms value. Therefore, at a suitable cooling rate, the supercooled B2-CuZr phase can be retained in the amorphous matrix without occurrence of
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the matensitic transformation. With the further increase of the Al content to above 8%, the formation of the Al2Zr phase tends to become predominant and precipitation of the B2-CuZr phase is further suppressed (see dashed lines in Fig. 3(b)). Consequently, glass formation is demoted due to the strong forming tendency of the Al2Zr phase. Upon slow cooling, Al2Zr always precipitates first and BMG composites containing the single B2-CuZr phase as reinforcement cannot be obtained anymore.
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2) Effects of features of reinforcing crystalline phase on mechanical properties of BMG composites Morphology characteristics (e.g., shape and distribution) of the crystalline phases were found to have significant effects on mechanical properties of TRIP-reinforced BMG composites[71–73]. If brittle intermetallics phase such as Al2Zr precipitated along with the ductile austenitic phase, the mechanical properties of the composite were deteriorated[38]. If the austenitic phase is the single reinforcement, macroscopic yield strength of the composite was mainly determined by its volume fraction. However, the plasticity and work-hardening behavior were governed by not only its volume fraction, but also its distribution[71–73]. With the increase of the austenite, yield strength decreased and for the composites with a crystalline volume fraction higher than 50%, characteristic of the resultant stress-strain curve was similar to that of the conventional shape memory alloy[72,73]. Moreover, even for the BMG composites with a similar crystalline volume fraction, the fracture strain may vary tremendously due to different distribution of the reinforcing phase[71]. As an example, Fig. 4 shows the stress-strain curves of three BMG composites with ~10% B2 phase for the Cu48Zr48Al4 BMG composite. Surprisingly, the fracture strain changes dramatically from 6% to 22%. SEM (scanning electron microscopy) characterization confirmed that these specimens had a similar volume fraction of the crystalline phase (typical examples are given in Fig. 4), nevertheless, specimen A has the fewest crystals with the largest size, whilst B and C have more tiny crystals. In addition, the tiny crystalline particles are distributed more homogeneously in specimen C than in specimen B. As can be seen, homogeneous distribution of the B2 phase gave rise to a large fracture strain, manifesting that distribution of the reinforced crystals also play important roles in achieving macroscopic plasticity of the current type of BMG composites. To fabricate BMG composites with the homogeneously distributed austenitic phase, controlling nucleation and growth of the B2 phase was the key. Many factors influencing its distribution have been investigated, such as cooling rates (i.e., casting size), superheating (i.e., melt current)[74,75], and microalloying of Ta, Sn and Nb to form high-temperature phases as heterogeneous nucleation sites for the austenitic crystalline phases[76–78]. It might be possible to use annealing or semi-solid processing to partially crystallize a monolithic BMG or homogenize the crystallite distribution, but these processes are difficult to control due to the metastability of the austenitic phase[79]. Among all the proposed approaches, alloying strategy may be a feasible way to develop tough TRIP-BMG composites. 3) Reinforcing mechanism Origins of the enhanced mechanical properties in the TRIP-reinforced BMG composites
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can be attributed to the following two aspects. The first one is the “blocking effect” from heterogeneities at multiple length scales[80]. As indicated in Fig. 5, shear bands during deformation were deflected at the interface between the crystals and glassy matrix, and then divided into many tiny secondary shear bands around the spherical phases. The crystalline spheres acted as strong barriers for the rapid propagation of shear bands and thus effectively enhanced plasticity of the composites. Moreover, the “blocking effect” was also observed for not only the micro-meter sized B2-CuZr phases, but also the nano-meter sized B2-CuZr ones (Fig. 5(c)). Therefore, the “blocking effect” from both micro-meter and nano-meter spherical phases would enhance the plasticity of the composite. The other important aspect is the contribution from the martensitic transformation of B2-CuZr. As shown in Fig. 6, experimental results evidenced that the phase transformation did occurred during deformation of the TRIP-reinforced BMG composite[33,39]. As the deformation proceeds, the crystalline B2 phase starts to transform into B19’ (Fig. (6)) and become harder and harder (Fig. 6(b)) while the amorphous matrix is softened due to extensive shear-band formation. It is reasonable to rationalize that the hardening of B2 will compensate softening of the amorphous matrix and contributes to the global mechanical properties of the composite. According to the literature[81], phase transformation of the austenite B2 spherical crystals during deformation could release the stress concentration around them and restrict free volume accumulation. As a result, the rapid propagation of shear bands can be hindered, which requires a further stress to move the shear bands and consequently inhibits the early necking and work-softening.
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4) Optimizing BMG composites by tailoring stacking fault energy Due to the important role of martensitic transformation in the global deformation behavior of the TRIP-BMG composites, ductilizing BMG composites by tailoring the martensitic transformation via stacking fault energy (SFE) becomes possible[40,82]. Generally, a low SFE tends to facilitate deformation twinning to initiate and thus stimulates the martensitic transformation of the austenite B2 structure[83,84]. By systematically investigating the alloying effects in the Zr48Cu48–xAl4Mx (x=0–2 at.%, M=Co, Ti, Fe, Ni, Ta, Cr, Ga, Zn, Hf, Nb, Ta and Ag) system, it was found that stacking fault energy of the crystalline phase, which usually governs plastic deformation carriers (i.e., dislocation and twinning) in crystalline materials, also can be tailored to ductilize the TRIP-BMG composites. For the bcc B2-CuZr phase in the ZrCuAlM system, the minimum SFE occurred along the (011)[100] slip system among all the possible lattice planes investigated. Some alloying additions, such as Co, dramatically reduce SFE, but some other elements, such as Ti, significantly increase it (Fig. 7(a)). The onset temperature (Ms) was also found to be correlated with SFE. The smaller the SFE, the higher the Ms, confirming that proper alloying additions can facilitate the martensitic transformation[40,82]. In the case of a similar crystal volume fraction, it was found that the TRIP-BMG composite with a smaller SFE yields at a relatively low stress but undergoes a pronounced work-hardening after yielding and a larger plastic deformation, and vice versa for the BMG composite with a higher SFE (Fig. 7(b)). This trend demonstrates that the minor additions reducing SFE of B2-CuZr can remarkably improve the tensile ductility and work-hardening capability of the TRIP-BMG composite.
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Variation of the SFE usually has its origins in the electronic structures, and particularly relate to the charge density change during the stacking fault formation[85,86]. From Fig. 8(a), it can be seen that the doped element that increases the electron charge density redistribution also increases the SFE, which is consistent with the tendency in other alloy systems[85,86]. Furthermore, it is found that both the electronegativity and the atomic radius difference between the doped element and substituted element correlate closely with the SFE of the doped B2 structure, as demonstrated in Fig. 8(b). Clearly, a smaller electronegativity and atomic radius difference between the doping element and Cu lead to a lower SFE of the doped B2-CuZr phase. As a result, the martensitic transformation of the B2 crystals doped with suitable elements that reduce the SFE could occur much easier, which eventually results in an earlier yielding, a larger plasticity and more pronounced work-hardening capability.
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3. Ex-situ Formed BMG Composite
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Comparing with those of in-situ formed BMG composite, volume fraction and distribution of foreign inclusions in ex-situ formed BMG composites can be easily adjusted, and consequently preparation of this type of BMG composites is also an effective way to ductilize glassy alloys. In this section, two approaches recently developed were introduced.
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3.1. Continuous formation of metallic glass coated wires and their bulk composites The size of BMG composites prepared by conventional casting processes, such as melt quenching, pressure injection and infiltration casting[9,87] is usually restricted to centimeter to decimeter scale. Recently, a continuous processing method to prepare MG coated composite wires was developed in our group[88,89]. The sketch of the continuous infiltration system for producing BMG coated wires is shown in Fig. 9(a). The system consists of a vacuum, heat, cooling, and motor drive unit. Alloy ingots of BMGs were placed in a crucible and heated to the melt temperature. Tungsten wires or other metallic wires with different diameters were polished and cleaned, and then was fed continuously from a series of supply wheels. It is to be noted that multiple wires can be fed simultaneously. These metallic wires passed through preheating unit, and then immersed into molten metal for infiltration at a tunable drawing velocity. As the wires were drawn out the molten pool, it was cooled forcibly by argon. Eventually, BMG composite rods with a diameter around 0.5 mm were prepared continuously, as shown in Fig. 9(b). SEM images (Figs. 9(b) and (c)) indicted that no contrast corresponding to the crystalline phase can be found at the interface between the amorphous matrix and wire (Figs. 9(d) and (e)), indicating that no reactant was formed during the infiltration process[89]. Furthermore, the composite wires can be weaved at the supercooled liquid region to form bulk form for structural applications, as shown in Fig. 9(f)[88]. Fig. 10(a) shows the uniaxial tension curves of the BMG composite rod with 61.4 vol.% tungsten wires, in comparison with that of the monolithic BMG sample and tungsten wire. It can be seen that the composite have a larger fracture strength and improved plastic deformation. Multiple shear bands can be seen on the lateral surface near the fracture section (Fig. 10(b)). It is inevitable that global mechanical properties of the BMG composite would be affected by properties of the metallic wires used. By properly choosing type, volume
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fraction and diameter of metallic wires used in this approach, mechanical behavior of the BMG composites can effectively be tuned.
4. Summary
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3.2. BMG composites reinforced by three dimensional metallic network Although ex-situ formed BMG composites have been widely studied[44–47], the added second inclusions are usually distributed discontinuously with significant inhomogeneity in the BMG matrix, which greatly reduce the reinforcing effectiveness[48]. Reinforcements with a continuous three-dimensional networked structure are expected to suppress the rapid propagation of shear bands more effectively. To test this hypothesis, a Cu foam with a porosity content of 95.8% and a pore radius ranging from 200 to 250 µm was chosen. The dimension of networked ligament is approximately 12 mm in diameter. The Ti40Zr25Cu12Ni3Be20 amorphous alloy has a melting point of 983 K[90], which is far below the melting point of Cu, was chosen as the amorphous matrix. The Cu foam was initially preloaded in the Cu mold, melted master alloy was subsequently sucked into the Cu mold under argon atmosphere and then solidified to form the final composite. In spite of the small volume fraction of Cu foam (~4.2%), yield strength of the composite slightly decreased while plasticity dramatically increased comparing with the monolithic BMG, as shown in Fig. 11[91]. By analyzing elastic behavior and fracture energy dissipation, effects of the ductile Cu network on the deformation of the composite can be summarized below. When shear bands are formed and propagate in the amorphous phase, they are confined by the individual compartment, which is surrounded by the Cu cell-wall. In other words, long-range shear band propagation is prohibited. The ductile Cu can temporarily block the shear bands and plastically deform to dissipate the kinetic energy. Additionally, the elastic mismatch at the interface between the amorphous matrix and cell-wall can lead to stress concentration and initiate secondary shear bands. The bifurcation of shear bands would certainly contribute to the improvement of plasticity.
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To make BMGs viable for practical applications as engineering materials, their damage tolerance and plastic deformation capability under tension (i.e., toughness, tensile ductility and work-hardening) have to be improved. As elaborated above, our recent research work on BMG composites was briefly highlighted. To fabricate in-situ formed BMG composites, we have applied different alloy design methodologies, i.e., phase separation and transformation-induced plasticity. To develop ex-situ formed BMG composites, we not only developed a new technique for continuously producing BMG coated wires but also fabricated ex-situ formed BMG composites with 3D metallic network inclusions. Our results indicate that development of advanced BMG composites with optimized overall properties provide a promising route to utilize BMG materials for practical engineering applications. Acknowledgements This research was supported in part by the National Natural Science Foundation of China (Nos. 51010001, 51371003, 51001009 and 50901006), 111 Project (No. B07003) and Program for Changjiang Scholars and Innovative Research Team in University. Y. Wu and X.J. Liu acknowledge the financial support from “the Fundamental Research Funds for the Central Universities”.
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References [1] W. L. Johnson, MRS Bull. 24 (1999) 42–56. [2] A. Inoue, Acta Mater. 48 (2000) 279–306. [3] C.A. Schuh, T.C. Hufnagel, U. Ramamurty, Acta Mater. 55 (2007) 4067–4109. [4] M.F. Ashby, A.L. Greer, Scripta Mater. 54 (2006) 321–326. [5] M.W. Chen, Annu. Rev. Mater. Res. 38 (2008) 445–469. [6] Z.F. Zhang, J. Eckert, L. Schultz, Acta Mater. 51 (2003) 1167–1179. [7] F.F. Wu, Z.F. Zhang, S.X. Mao, A. Peker, J. Eckert, Phys. Rev. B 75 (2007) 134201. [8] H.A. Bruck, A.J. Rosakis, W.L. Johnson, J. Mater. Res. 11 (1996) 503–511. [9] R.B. Dandliker, R.D. Conner, W.L. Johnson, J. Mater. Res. 13 (1998) 2896–2901. [10] H. Choi-Yim, R. Busch, W.L. Johnson, J. Appl. Phys. 83 (1998) 7993–7997. [11] H. Choi-Yim, R.D. Conner, F. Szuecs, W.L. Johnson, Acta Mater. 50 (2002) 2737–2745. [12] Y.F. Xue, H.N. Cai, L. Wang, F.C. Wang, H.F. Zhang, Z.Q. Hu, Compos. Sci. Technol. 68 (2008) 3396–3400. [13] Y.C. Kim, E. Fleury, J.C. Lee, D.H. Kim, J. Mater. Res. 20 (2005) 2474–2479. [14] F. Szuecs, C.P. Kim, W.L. Johnson, Acta Mater. 49 (2001) 1507–1513. [15] M.L. Lee, Y. Li, C.A. Schuh, Acta Mater. 52 (2004) 4121–4131. [16] H. Tan, Y. Zhang, Y. Li, Intermetallics 10 (2002) 1203–1205. [17] U. Kuhn, J. Eckert, N. Mattern, L. Schultz, Mater. Sci. Eng. A 375–377 (2004) 322–326. [18] C. Fan, R.T. Ott, T.C. Hufnagel, Appl. Phys. Lett. 81 (2002) 1020–1022. [19] U. Kühn, J. Eckert, N. Mattern, L. Schultz, Appl. Phys. Lett. 80 (2002) 2478–2480. [20] C.C. Hays, C.P. Kim, W.L. Johnson, Phys. Rev. Lett. 84 (2000) 2901. [21] C. Fan, H.Q. Li, L.J. Kecskes, K. X. Tao, H. Choo, P. K. Liaw, C.T. Liu, Phys. Rev. Lett. 96 (2006) 145506. [22] G. Chen, H. Bei, Y. Cao, A. Gali, C.T. Liu, E.P. George, Appl. Phys. Lett. 95 (2009) 081908. [23] D.C. Hofmann, J.Y. Suh, A. Wiest, G. Duan, M.L. Lind, M.D. Demetrious, W.L. Johnson, Natrue 451 (2008) 1085–1089. [24] D.C. Hofmann, J.Y. Suh, A. Wiest, M.L. Lind, M.D. Demetriou, W.L. Johnson, Proc. Natl. Acad. Sci. U.S.A. 105 (2008) 20136. [25] Z. Zhu, H. Zhang, Z. Hu, W. Zhang, A. Inoue, Scripta Mater. 62 (2010) 278–281. [26] J.W. Qiao, S. Wang, Y. Zhang, P.K. Liaw, G.L. Chen, Appl. Phys. Lett. 94 (2009) 151905. [27] D.C. Hofmann, J.Y. Suh, A. Wiest, W. Johnson, Scripta Mater. 59 (2008) 684–687. [28] D.C. Hofmann, H. Kozachkov, H.E. Khalifa, J.P. Schramm, M.D. Demetriou, K.S. Vecchio, W.L. Johnson, JOM 61 (2009) 11–17. [29] V. F. Zackay, E.R. Parker, D. Fahr, R. Busch, Trans. ASM 60 (1967) 252–259. [30] P.J. Jacques, Q. Fumemont, F. Lani, T. Pardoen, F. Delannay, Acta Mater. 55 (2007) 3681–3693. [31] J. Das, M.B. Tang, K.B. Kim, R. Theissmann, F. Baier, W.H. Wang, J. Eckert, Phys. Rev. Lett. 94 (2005) 205501. [32] K.B. Kim, J. Das, S. Venkataraman, S. Yi, J. Eckert, Appl. Phys. Lett. 89 (2006) 071908. [33] S. Pauly, J. Das, J. Bednarcik, N. Mattern, K.B. Kim, D.H. Kim, J. Eckert, Scripta Mater. 60 (2009) 431–434.
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AC C
EP
TE D
M AN U
SC
RI PT
[34] Y.F. Sun, B.C. Wei, Y.R. Wang, W.H. Li, T.L. Cheung, C.K. Shek, Appl. Phys. Lett. 87 (2005) 051905. [35] S. Pauly, J. Das, C. Duhamel, J. Eckert, Adv. Eng. Mater. 9 (2007) 487–491. [36] S. Pauly, G. Liu, G. Wang, U. Kuhn, N. Mattern, J. Eckert, Acta Mater. 57 (2009) 5445–5453. [37] S. Pauly, G. Liu, S. Gorantla, G. Wang, U. Kuhn, D.H. Kim, J. Eckert, Acta Mater. 58 (2010) 4883–4890. [38] Y. Wu, H. Wang, H.H. Wu, Z.Y. Zhang, X.D. Hui, G.L. Chen, D. Ma, X.L. Wang, Z.P. Lu, Acta Mater. 59 (2011) 2928–2936. [39] Y. Wu, Y.H. Xiao, G.L. Chen, C.T. Liu, Z.P. Lu, Adv. Mater. 22 (2010) 2770–2773. [40] Y. Wu, D.Q. Zhou, W.L. Song, H. Wang, Z.Y. Zhang, D. Ma, X.L. Wang, Z.P. Lu, Phys. Rev. Lett. 109 (2012) 245506. [41] C.P. Kim, Y.S. Oh, S.H. Lee, N.J. Kim, Scripta Mater. 65 (2011) 304–307. [42] P. Gargarella, S. Pauly, K.K. Song, J. Hu, N.S. Barekar, M.S. Khoshkhoo, A. Teresiak, H. Wendrock, U. Kuhn, C. Ruffing, E. Kerscher, J. Eckert, Acta Mater. 61 (2013) 151–162. [43] Y.S. Oh, C.P. Kim, S.H. Lee, N.J. Kim, Acta Mater. 59 (2011) 7277–7286. [44] J.A. Hines, K.S. Vecchio, Acta Mater. 45 (1997) 635–649. [45] B. Zhang, H.M. Fu, Z.W. Zhu, A.M. Wang, H. Li, C. Dong, Mater. Sci. Eng. A 504 (2012) 207–211. [46] H.F. Zhang, H. Li, A.M. Wang, H.M. Fu, B.Z. Ding, Z.Q. Hu, Intermetallics 17 (2009) 1070–1077. [47] Y. Sun, H.F. Zhang, A.M. Wang, H.M. Fu, Z.Q. Hu, C.E. Wen, Appl. Phys. Lett. 95 (2009) 171910. [48] J.S.C. Jang, T.H. Li, S.R. Jian, J.C. Huang, T.G. Nieh, Intermetallics 19 (2011) 738–743. [49] B.J. Park, H.J. Chang, D.H. Kim, W.T. Kim, K. Chattopadhyay, T.A. Abinandanan, S. Bhattacharyya, Phys. Rev. Lett. 96 (2006) 245503. [50] E.S. Park, D.H. Kim, Acta Mater. 54 (2006) 2597–2604. [51] D.H. Kim, W.T. Kim, E.S. Park, N. Mattern, J. Eckert, Prog. Mater. Sci. 58 (2013) 1103–1172. [52] J.E. Gao, Z.P. Chen, Q. Du, H.X. Li, Y. Wu, H. Wang, X.J. Liu, Z.P. Lu, Acta Mater. 61 (2013) 3214–3223. [53] Z.P. Lu, C.T. Liu, J.R. Thomepson, W.D. Porter, Phys. Rev. Lett. 92 (2004) 245503. [54] V. Ponnambalam, S.J. Poon, G.J. Shiflet, J. Mater. Res. 19 (2004) 1320–1323. [55] A. Inoue, T. Zhang, A. Takeuchi, Appl. Phys. Lett. 71 (1997) 464. [56] A.R. Yavari, Mater. Sci. Eng. 98 (1988) 491–493. [57] X.J. Gu, S.J. Poon, G.J. Shiflet, W. Widom, Appl. Phys. Lett. 92 (2008) 161910. [58] S.J. Poon, G.J. Shiflet, Science 241 (1988) 1640–1642. [59] Z.P. Chen, J.E. Gao, Y. Wu, H.X. Li, H. Wang, Z.P. Lu, J. Alloy. Compd. 513 (2012) 387–392. [60] H. Yang, J.Q. Wang, Y. Li, J. Non-Cryst. Solids 354 (2008) 3473–3479. [61] F.Q. Guo, S.J. Enouf, S.J. Poon, G.J. Shiflet, Philos. Mag. Lett. 81 (2001) 203–211. [62] Z.P. Chen, J.E. Gao, Y. Wu, H. Wang, X.J. Liu, Z.P. Lu, Sci. Rep. 3 (2013) 3353. [63] D. Schryvers, G.S. Firstov, J.W. Seo, J.V. Humbeeck, Y.N. Koval, Scripta Mater. 36 (1997) 1119–1125.
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[64] J.W. Seo, D. Schryvers, Acta Mater. 46 (1998) 1165–1175. [65] T. Abe, M. Shimono, M. Ode, H. Onodera, Acta Mater. 54 (2006) 909–915. [66] K.K. Song, S. Pauly, Y. Zhang, P. Gargarella, R. Li, N.S. Barekar, U. Kuhn, M. Stoica, J. Eckert, Acta Mater. 17 (2011) 6620–6630 [67] K.K. Song, S. Pauly, B.A. Sun, Y. Zhang, J. Tan, U. Kuhn, M. Stoica, J. Eckert, Intermetallics 30 (2012) 132–138. [68] K.K. Song, S. Pauly, Y. Zhang, B.A. Sun, J. He, G.Z. Ma, U. Kuhn, J. Eckert, Mater. Sci. Eng. A 559 (2013) 711–718. [69] X.P. Lu, D. Ma, C.T. Liu, Y.A. Chang, Intermetallics 15 (2007) 253–259. [70] Y.N. Koval, G.S. Firstov, A.V. Kotko, Scripta Metall. Mater. 27 (1992) 1611–1616. [71] Y. Wu, W.L. Song, Z.Y. Zhang, X.D. Hui, D. Ma, X.L. Wang, X.C. Shang, Z.P. Lu, Chin. Sci. Bull. 56 (2011) 3960–3964. [72] K.K. Song, S. Pauly, Y. Zhang, R. Li, S. Gorantla, N. Narayanan, U. Kuhn, T. Gemming, J. Eckert, Acta Mater. 60 (2012) 6000–6012. [73] S. Pauly, G. Liu, G. Wang, J. Das, K.B. Kim, U. Kuhn, D.H. Kim, J. Eckert, Appl. Phys. Lett. 95 (2009) 101906. [74] S. Pauly, G. Liu, G. Wang, U. Kuhn, N. Mattern, J. Eekert, Acta Mater. 57 (2009) 5445–5453. [75] Z.Q. Liu, R. Li, G. Liu, K.K. Song, S. Pauly, T. Zhang, J. Eckert, AIP Adv. 2 (2012) 032176. [76] Z.Q. Liu, R. Li, G. Liu, W.H. Su, H. Wang, Y. Li, M.J. Shi, X.K. Luo, G.J. Wu, T. Zhang, Acta Mater. 60 (2012) 3128–3139. [77] Z.Y. Zhang, Y. Wu, J. Zhou, W.L. Song, D. Cao, H. Wang, X.J. Liu, Z.P. Lu, Intermetallics 42 (2013) 68–76. [78] S. Pauly, J. Das, J. Bednarcik, N. Mattern, K.B. Kim, D.H. Kim, J. Eckert, Scripta Mater. 60 (2009) 431–434. [79] H. Kozachkov, J. Kolodziejska, W.L. Johnson, D.C. Hofmann, Intermetallics 39 (2013) 89–93. [80] D.V. Louzguine-Luzgin, A. Vinogradov, G.Q. Xie, S. Li, A. Lazarev, S. Hashimoto, A. Inoue, Philos. Mag. 89 (2009) 2887–2901. [81] J. Das, K.B. Kim, W. Xu, B.C. Wei, Z.F. Zhang, W.H. Wang, S. Yi, J. Eckert, Mater. Trans. 47 (2006) 2606–2609. [82] D.Q. Zhou, Y. Wu, H. Wang, X.D. Hui, X.J. Liu, Z.P. Lu, Comput. Mater. Sci. 79 (2013) 187–192. [83] D.C. Hoffman, Science 329 (2010) 1294–1295. [84] X.Q. Wang, Phys. Rev. B 78 (2008) 092103. [85] S. Ogata, J. Li, S. Yip, Science 298 (2002) 807–811. [86] Y. Qi, R.K. Mishra, Phys. Rev. B 75 (2007) 224105. [87] H. Choi-Yim, W.L. Johnson, Appl. Phys. Lett. 71 (1997) 3808–3810. [88] X.H. Chen, B.Y. Zhang, G.L. Chen, Y. Zhang, X.D. Hui, Z.P. Lu, X.J. Liu, Y. Xu, X.R. Xing, Intermetallics 18 (2010) 2034–2038. [89] B.Y. Zhang, X.H. Chen, S.S. Wang, D.Y. Lin, X.D. Hui, Mater. Lett. 93 (2013) 210–214. [90] G. Wang, J. Shen, J.F. Sun, Y.J. Huang, J. Zou, Z.P. Lu, Z.H. Stachurski, B.D. Zhou, J. Non-Cryst. Solids 351 (2005) 209–217.
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Fig. 1 (a) Critical thickness for the BMG matrix composites as a function of alloy compositions in the Fe–Co–La–Ce–Al–Cu system; (b) schematic illustration of the solidification process of the Fe–Co–La–Ce–Al–Cu alloys[52].
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Fig. 2 (a) Critical thickness for BMG matrix composites as a function of alloy compositions for the Al–(CoCu)–(La,Ce)–Fe system, and (b) outer appearance of the as-cast Al40(CoCu)15(LaCe)45 BMG matrix composites with a diameter of 15 mm, and (c) DSC curves of the rods with a diameter of 15 mm for the Al35(CoCu)15(LaCe)50, Al40(CoCu)15(LaCe)45, Al40(CoCu)10(LaCe)50, Al45(CoCu)10(LaCe)45 and Al45(CoCu)5(LaCe)50 alloys[62].
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Fig. 3 DSC heating and cooling curves of the (Cu0.5Zr0.5)100–xAlx (x=0, 4, 6) alloy ingots showing the martensitic and austenitic transformation (a), and schematic time-temperature-transformation diagram of the (Cu0.5Zr0.5)100–xAlx alloys with a different Al content (b)[38]. Fig. 4 Compressive true stress-strain curves of the Cu48Zr48Al4 alloy with about 10% B2 phase, and SEM images of the three representative samples with a similar crystal volume fraction of B2 but a different distribution pattern[71].
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Fig. 5 (a) SEM image of the lateral surface, (b) blowup of the lateral surface, and (c) TEM image of the tensile fractured sample of the Cu48Zr48Al4 BMG composite[38]. Fig. 6 XRD patterns of the as-cast and fractured samples of the Cu48Zr48Al4Co0.5 BMG composite (a), micro-hardness of the crystalline phase and amorphous matrix in the as-cast and tensioned samples (b)[39].
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Fig. 7 Stacking fault energy on the (011)[100] slip system of the B2-CuZr phase and the ones with half of Cu on the slip plane substituted by Co or Ti (a), and true tensile stress-strain curves of the Zr48Cu48Al4, Zr48Cu47.5Al4Co0.5 and Zr48Cu47.5Al4Ti0.5 BMG composites (b)[40]. Fig. 8 SFE values of the B2-CuZr phase doped with different alloying elements as a function of the electronic density charge redistribution on the (011)[100] slip system (a), and dependence of the stacking fault energy of the doped B2 phase on the eletronegativity and radius difference between the doping and substituted elements (b)[40]. Fig. 9 (a) Schematic illustration of the newly developed continuous infiltration process; (b) a bunch of BMG composite wires prepared, (c) and (d) SEM images of the cross-section of BMG composites containing 61.4 vol% tungsten wires, (e) TEM image of the interface between the tungsten and the amorphous matrix shown in d, and (f) bulk sample weaved with BMG-coated tungsten wire composite[88,89].
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Fig. 10 Quasi-static tensile stress–strain curves of the BMG composite wires containing 61.4 vol.% tungsten wires in comparison with that of the tungsten wire and monolithic BMG Vit 1 (a), and side image near fracture surface[89].
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Fig. 11 Compressive stress-strain curves of the monolithic Ti-based BMG and the fabricated BMG composite (a), as compared with that of the Cu foam (b). Insets are the enlargements of the plastic deformation curves and appearances of the pre-strained Cu foams, respectively[91].
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