**3.1. Zr-based BMGMCs**

β-Zr phase is one of the in-situ reinforcements in Zr-based BMGMCs that has been studied a lot. Hays et al. have reported the development of β phase reinforced Zr-Ti-Cu-Ni-Be-Nb BMGMCs [47]. Both XRD patterns and SEM images suggest the precipitation of β phase during rapid quenching. The β phase is in the dendritic and distributed in the glassy matrix homogeneously, as shown in **Figure 9a**. The volume fraction of β phase is estimated to be ~25%. The dendritic structures are characterized by primary dendrite axes with lengths of 50–150 μm and radius of about 1.5–2 μm. Regular patterns of secondary dendrite arms with spacing of 6–7 μm can be observed from SEM images. This composite shows about 5% plastic strain under three point bending. **Figure 9b** shows that the shear bands propagate preferentially through many successive dendrite arms, occasionally initiate or terminate within the arms, and clearly propagate as localized bands through the β-phase arms. This composite also shows good plasticity during compression, as shown in **Figure 9c**. It yields at 1.3 GPa when the β phase yields and deforms, and shear band patterns develop, as the glassy matrix is locally loaded beyond its critical shear stress. The plastic strain is over 6%. The composite even shows about 5% plastic strain during tension. Clear necking and deformation can be observed. The dendritic microstructure of the β phase acts to seed the initiation of organized shear band patterns, confines the propagation of individual shear bands to domains having a spatial scale of the order of the primary dendritic axes length, and lead to shear band spacing which is related to the dendrite arm spacing.

Another important in-situ secondary phase in Zr-based BMGMCs is refractory metal phase. Fan et al. have introduced in-situ Ta-rich precipitates in Zr-Cu-Al-Ni-Ta BMGMCs [48]. For the high melting temperature of Ta, there are two steps arc-melting during master alloy preparation. Firstly, Zr-Ta ingot is fabricated which forms solid solutions of Zr-Ta, and remaining elements are subsequently mixed with Zr-Ta ingots. The microstructure of the as-cast composite samples consists of both glassy matrix and Ta-rich particles with an average size of 10–30 μm, as shown in **Figure 10a**. The particles are oblong in shape and do not appear to possess a dendritic structure, they distribute homogeneously among the matrix and the volume faction is about 4%. This composite yields at 1.7 GPa and exhibits apparent work hardening and significant plastic strain, as shown in **Figure 10b**. The elastic incompatibility between the particles and the matrix introduces stress concentrations which may promote shear band initiation. The particles may also impede shear band propagation. Guo et al. have applied dealloying in metallic melt method to further optimize the microstructure and mechanical properties of Ta-rich phase reinforced BMGMCs [49]. The dealloying in metallic melt phenomenon occurs when immerging Zr-Ta solid solution precursor in Cu-Al-Ni melt. For the negative enthalpy of mixing between Zr and Cu-Al-Ni and positive enthalpy of mixing between Ta and Cu-Al-Ni, Zr is gradually selectively leached from precursor ingot to

more interfaces, which makes the yield strength follows the load-bearing mode even with low volume fraction of particles. This composite also shows very obvious work hardening behavior, which is considered to originate from the stress-induced martensitic transformation of B2-NiTi phase, which is very attractive and different from those conventional metal or ceramic particles.

**Figure 8.** (a) Porous NiTi powder by top-down process; (b) SEM images of BMGMCs containing 20 vol% porous particles; (c) compressive true stress-strain curves for monolithic base BMG and BMGMCs with various volume fraction

**Figure 7.** (a) SEM observation of the porous Mo particles in the BMG matrix, with the inserted XRD pattern; (b) an enlarged image of a single porous Mo particle; (c) representative room-temperature compressive engineering stress-

For in-situ BMGMCs, the fabrication process is very important and difficult to design, such as how to induce in-situ precipitated phase, how to control the volume fraction and size of secondary phase. For different alloy systems, the process and reinforcing phases are quite dif-

β-Zr phase is one of the in-situ reinforcements in Zr-based BMGMCs that has been studied a lot. Hays et al. have reported the development of β phase reinforced Zr-Ti-Cu-Ni-Be-Nb BMGMCs [47]. Both XRD patterns and SEM images suggest the precipitation of β phase during rapid quenching. The β phase is in the dendritic and distributed in the glassy matrix homogeneously, as shown in **Figure 9a**. The volume fraction of β phase is estimated to be ~25%. The dendritic

ferent, thus, this part will be divided by alloy systems, not by reinforcing phases.

**3. In-situ BMGMCs**

of porous NiTi addition.

strain curves for the BMGMCs.

60 Metallic Glasses - Properties and Processing

**3.1. Zr-based BMGMCs**

**Figure 9.** (a) SEM image of in-situ β-Zr reinforced BMGMCs (inset: XRD patterns); (b) shear band patterns array from compressive failure region of bend test sample; (c) compressive stress strain curve for the composite.

**Figure 10.** (a) SEM image of in-situ Ta reinforced BMGMCs; (b) compressive stress-strain curves for as-cast in-situ Ta reinforced BMGMCs; (c) SEM images of the sample by conventional arc-melting; (d) SEM images of the sample by novel dealloying method; (e) true compressive stress-strain curves and (f) tensile stress-strain curves of monolithic BMG, BMGMC by arc-melting (ZT-A) and BMGMC by dealloying (ZT-D).

the melt and forming a glass-former liquid, while at the same time, remaining Ta breaks into small particles and distribute in the melt. By finally quenching, the BMGMCs can be fabricated. Interestingly, the size of Ta-rich particles by dealloying method is much smaller than that by conventional arc-melting method, 40 μm of arc-melting sample and 10 μm of dealloying sample, as shown in **Figure 10c** and **d**.

> spheroidal crystal phases of B2-CuZr (confirmed by XRD, not shown here) are embedded in the amorphous matrix. The volume fraction of B2-CuZr is estimated to be about 10%. The composite exhibits obvious plastic deformation under compression with various strain rates, as shown in **Figure 11b**. Moreover, the flow stress increases with the increasing strain after yielding, exhibiting obvious work-hardening behavior. Even though the B2-CuZr reinforced Zr-based BMGMCs have shown good plasticity and work-hardening, but the inhomogeneous distribution of B2-CuZr limits the further improvement of the mechanical properties. During cooling, the B2-CuZr phase tends to precipitate in the center of the sample where the cooling rate is lower than that near the edge. Guo and Saida have successfully homogenized the distribution of B2-CuZr phase by minor doping Ta [52]. During cooling, the primary precipitated Ta-rich phase acts as effective nucleants that promoted copious nucleation of the B2-CuZr phase. As shown in **Figure 11c**, the dark phase of shape memory phase forms around the gray phase of Ta-rich particle. The total volume fraction of crystalline secondary phase, including both Ta-rich particles and B2-CuZr (some transforms to B19'-CuZr because of the residual heat at the interface), is estimated to be ~80% by comparing the heat of crystallization of both the composites and monolithic BMG. The volume fraction of Ta-rich particles is estimated to be ~10% based on the image analysis. This composite shows a superior plasticity of 8.4% plastic strain during tension, as well as obvious work-hardening and a unique triple yielding phenomenon, as shown

> **Figure 11.** (a) Typical OM images of the B2-CuZr reinforced BMGMCs (inset: enlarged parts of rectangular area); (b) true stress-strain curves of the composites with various strain rates; (c) SEM image of ta-doped sample; (d) true tensile stress-

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strain curves of both monolithic base alloy (Ta0) and Ta-doped multiphase reinforced BMGMC (Ta5).

in **Figure 11d**.

The reason is considered to be that during conventional arc-melting, the Ta-rich particles precipitate during cooling and its size can be hardly controlled, while during dealloying, for the low melting temperature, the Ta-rich particles remain in solid state and finally fine particle can be obtained. With these finer particles, the composite by dealloying method shows better plasticity than that by arc-melting method under both compression and tension, as shown in **Figure 10e** and **f**. The plastic strain for arc-melting sample is 3% under compression and 0% under tension. However, the value for dealloying sample is 14% under compression and 1.8% under tension. As discussed before, the stress concentration at the interfaces between Ta-rich particles and glassy matrix contributes a lot to the overall plasticity. For the finer size of particles in dealloying samples, the number density of particle and equivalent interface area per mm<sup>3</sup> are quite larger than arc-melting sample, 34.4 × 10<sup>4</sup> , 108 mm2 for dealloying sample and 0.4 × 10<sup>4</sup> and 22.5 mm2 for arc-melting sample. Thus, the five times larger area of interfaces in dealloying sample can plasticize the sample more.

Recently, B2-CuZr shape memory secondary phase in Zr-based BMGMCs have attracted a lot of interests. Unlike conventional metal or ceramic reinforcing phases, the shape memory phase can undergo the stress-induced martensitic transformation during deformation, which can both further improve the plasticity and give the sample obvious work hardening behavior [50]. Xu et al. have induced B2-CuZr phase in Zr-Cu-Al-Co system, as shown in **Figure 11a** [51]. The

**Figure 11.** (a) Typical OM images of the B2-CuZr reinforced BMGMCs (inset: enlarged parts of rectangular area); (b) true stress-strain curves of the composites with various strain rates; (c) SEM image of ta-doped sample; (d) true tensile stressstrain curves of both monolithic base alloy (Ta0) and Ta-doped multiphase reinforced BMGMC (Ta5).

the melt and forming a glass-former liquid, while at the same time, remaining Ta breaks into small particles and distribute in the melt. By finally quenching, the BMGMCs can be fabricated. Interestingly, the size of Ta-rich particles by dealloying method is much smaller than that by conventional arc-melting method, 40 μm of arc-melting sample and 10 μm of dealloy-

**Figure 10.** (a) SEM image of in-situ Ta reinforced BMGMCs; (b) compressive stress-strain curves for as-cast in-situ Ta reinforced BMGMCs; (c) SEM images of the sample by conventional arc-melting; (d) SEM images of the sample by novel dealloying method; (e) true compressive stress-strain curves and (f) tensile stress-strain curves of monolithic BMG,

The reason is considered to be that during conventional arc-melting, the Ta-rich particles precipitate during cooling and its size can be hardly controlled, while during dealloying, for the low melting temperature, the Ta-rich particles remain in solid state and finally fine particle can be obtained. With these finer particles, the composite by dealloying method shows better plasticity than that by arc-melting method under both compression and tension, as shown in **Figure 10e** and **f**. The plastic strain for arc-melting sample is 3% under compression and 0% under tension. However, the value for dealloying sample is 14% under compression and 1.8% under tension. As discussed before, the stress concentration at the interfaces between Ta-rich particles and glassy matrix contributes a lot to the overall plasticity. For the finer size of particles in dealloying samples, the number density of particle and equivalent interface area per mm<sup>3</sup>

Recently, B2-CuZr shape memory secondary phase in Zr-based BMGMCs have attracted a lot of interests. Unlike conventional metal or ceramic reinforcing phases, the shape memory phase can undergo the stress-induced martensitic transformation during deformation, which can both further improve the plasticity and give the sample obvious work hardening behavior [50]. Xu et al. have induced B2-CuZr phase in Zr-Cu-Al-Co system, as shown in **Figure 11a** [51]. The

, 108 mm2

for arc-melting sample. Thus, the five times larger area of interfaces in

for dealloying sample and

ing sample, as shown in **Figure 10c** and **d**.

62 Metallic Glasses - Properties and Processing

BMGMC by arc-melting (ZT-A) and BMGMC by dealloying (ZT-D).

are quite larger than arc-melting sample, 34.4 × 10<sup>4</sup>

dealloying sample can plasticize the sample more.

0.4 × 10<sup>4</sup>

and 22.5 mm2

spheroidal crystal phases of B2-CuZr (confirmed by XRD, not shown here) are embedded in the amorphous matrix. The volume fraction of B2-CuZr is estimated to be about 10%. The composite exhibits obvious plastic deformation under compression with various strain rates, as shown in **Figure 11b**. Moreover, the flow stress increases with the increasing strain after yielding, exhibiting obvious work-hardening behavior. Even though the B2-CuZr reinforced Zr-based BMGMCs have shown good plasticity and work-hardening, but the inhomogeneous distribution of B2-CuZr limits the further improvement of the mechanical properties. During cooling, the B2-CuZr phase tends to precipitate in the center of the sample where the cooling rate is lower than that near the edge. Guo and Saida have successfully homogenized the distribution of B2-CuZr phase by minor doping Ta [52]. During cooling, the primary precipitated Ta-rich phase acts as effective nucleants that promoted copious nucleation of the B2-CuZr phase. As shown in **Figure 11c**, the dark phase of shape memory phase forms around the gray phase of Ta-rich particle. The total volume fraction of crystalline secondary phase, including both Ta-rich particles and B2-CuZr (some transforms to B19'-CuZr because of the residual heat at the interface), is estimated to be ~80% by comparing the heat of crystallization of both the composites and monolithic BMG. The volume fraction of Ta-rich particles is estimated to be ~10% based on the image analysis. This composite shows a superior plasticity of 8.4% plastic strain during tension, as well as obvious work-hardening and a unique triple yielding phenomenon, as shown in **Figure 11d**.

#### **3.2. Ti-based BMGMCs**

The shape memory phase has also been introduced into Ti-based BMGMCs. Hong et al. have successfully induced B2 phase in Ti-Cu-Ni-Zr-Sn-Si system [53]. The size and distribution of B2 phase can be also tailored by varying the composition, as shown in **Figure 12a** and **b**. The size of B2 phase varies from 2–5 to 70–150 μm. The volume fraction varies from 10 to 33%. The composites also shows good plasticity under compression, the largest plastic strain is about 12.7%, as shown in **Figure 12c**. Furthermore, the yield strength exhibits a tendency to decrease with the increase of volume fraction of B2 phase, which originates from the early deformation on softer B2 phase. However, the plastic strain increases with more and larger B2 phase. The large B2 phase is found to be effective in dissipating the localization of shear stress, thus causing branching and multiplication of the shear bands. It is also observable that the severe deformation in the B2 phase, formation of wrinkles suggesting the possible deformation-induced phase transformation, as shown in **Figure 12d**.

Another important reinforcement in Ti-based BMGMCs is the β phase. Very good plasticity can be even observed for such composites under tension [54]. However, most of these composites contains Be, which is toxic and should be avoided when used as biomaterials. To induce β phase in Be-free alloy system, the β phase stabilizers, such as Ta, V and etc. Yamamoto et al. have successfully induced β-Ti phase in Ti-Cu-Ni-Sn-Ta system [55]. As shown in **Figure 13a**, for the low glass-forming ability of the matrix, Ti2 Ni also forms besides β-Ti phase. Thus, the plasticity of this composite is not very good, the plastic strain is about 1.6%, as shown in **Figure 13b**. To further improve the mechanical property, Guo and Kato have chosen a better glass former, Ti-Zr-Cu-Pd-Sn alloy, and used Mo as the β-Ti phase stabilizer element [56]. By doping 2 at%

Mo, homogeneously distributed oblong-like β-Ti phase can be observed from **Figure 14a**. The volume fraction and average size of β-Ti are estimated to be about 25 μm and 25%. This composite also shows a good plasticity under compression, fracture strength of 2160 MPa and plastic strain of 13.4%, as shown in **Figure 14b**. The nanoindentation test has shown that the β-Ti phase is softer than the matrix, indicating the propagation of the main shear band is hindered by the interfaces between the softer β-Ti and glassy matrix. The shear band is deflected, branched, or multiplied. Furthermore, after elastic deformation of both β-Ti and matrix to the yielding point, the β-Ti appears to yield and deform, contributing to the work-hardening behavior. The composite also shows about 3% plastic strain under three point bending test,

**Figure 13.** (a) XRD patterns and (b) stress-strain curves under compression for β phase reinforced Ti-Cu-Ni-Sn-Ta

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The researches on in-situ Mg-based BMGMCs are not as much as those in Zr-based or Ti-based systems for the difficulty to design proper fabrication process. However, recently, the application of novel dealloying in metallic melt method or selective phase leaching method in fabri-

**Figure 14.** (a) OM image of Mo-doped β-Ti based BMGMCs; (b) compressive true stress-strain curves for both monolithic base alloy and Mo-doped BMGMCs; (c) bending stress-strain curves for both base alloy and BMGMCs (inset: XRD

cation of in-situ Mg-based BMGMCs have attracted a lot of attentions.

see **Figure 14c**.

BMGMCs.

**3.3. Mg-based BMGMCs**

patterns for bending samples).

**Figure 12.** (a) and (b) SEM images of B2 phase reinforced Ti-based BMGMCs with different composition; (c) stress-strain curves of B2 phase reinforced Ti-based BMGMCs; (d) SEM image from the lateral surface of fractured sample.

**Figure 13.** (a) XRD patterns and (b) stress-strain curves under compression for β phase reinforced Ti-Cu-Ni-Sn-Ta BMGMCs.

Mo, homogeneously distributed oblong-like β-Ti phase can be observed from **Figure 14a**. The volume fraction and average size of β-Ti are estimated to be about 25 μm and 25%. This composite also shows a good plasticity under compression, fracture strength of 2160 MPa and plastic strain of 13.4%, as shown in **Figure 14b**. The nanoindentation test has shown that the β-Ti phase is softer than the matrix, indicating the propagation of the main shear band is hindered by the interfaces between the softer β-Ti and glassy matrix. The shear band is deflected, branched, or multiplied. Furthermore, after elastic deformation of both β-Ti and matrix to the yielding point, the β-Ti appears to yield and deform, contributing to the work-hardening behavior. The composite also shows about 3% plastic strain under three point bending test, see **Figure 14c**.

#### **3.3. Mg-based BMGMCs**

**Figure 12.** (a) and (b) SEM images of B2 phase reinforced Ti-based BMGMCs with different composition; (c) stress-strain curves of B2 phase reinforced Ti-based BMGMCs; (d) SEM image from the lateral surface of fractured sample.

The shape memory phase has also been introduced into Ti-based BMGMCs. Hong et al. have successfully induced B2 phase in Ti-Cu-Ni-Zr-Sn-Si system [53]. The size and distribution of B2 phase can be also tailored by varying the composition, as shown in **Figure 12a** and **b**. The size of B2 phase varies from 2–5 to 70–150 μm. The volume fraction varies from 10 to 33%. The composites also shows good plasticity under compression, the largest plastic strain is about 12.7%, as shown in **Figure 12c**. Furthermore, the yield strength exhibits a tendency to decrease with the increase of volume fraction of B2 phase, which originates from the early deformation on softer B2 phase. However, the plastic strain increases with more and larger B2 phase. The large B2 phase is found to be effective in dissipating the localization of shear stress, thus causing branching and multiplication of the shear bands. It is also observable that the severe deformation in the B2 phase, formation of wrinkles suggesting the possible

Another important reinforcement in Ti-based BMGMCs is the β phase. Very good plasticity can be even observed for such composites under tension [54]. However, most of these composites contains Be, which is toxic and should be avoided when used as biomaterials. To induce β phase in Be-free alloy system, the β phase stabilizers, such as Ta, V and etc. Yamamoto et al. have successfully induced β-Ti phase in Ti-Cu-Ni-Sn-Ta system [55]. As shown in **Figure 13a**, for the

ity of this composite is not very good, the plastic strain is about 1.6%, as shown in **Figure 13b**. To further improve the mechanical property, Guo and Kato have chosen a better glass former, Ti-Zr-Cu-Pd-Sn alloy, and used Mo as the β-Ti phase stabilizer element [56]. By doping 2 at%

Ni also forms besides β-Ti phase. Thus, the plastic-

deformation-induced phase transformation, as shown in **Figure 12d**.

low glass-forming ability of the matrix, Ti2

**3.2. Ti-based BMGMCs**

64 Metallic Glasses - Properties and Processing

The researches on in-situ Mg-based BMGMCs are not as much as those in Zr-based or Ti-based systems for the difficulty to design proper fabrication process. However, recently, the application of novel dealloying in metallic melt method or selective phase leaching method in fabrication of in-situ Mg-based BMGMCs have attracted a lot of attentions.

**Figure 14.** (a) OM image of Mo-doped β-Ti based BMGMCs; (b) compressive true stress-strain curves for both monolithic base alloy and Mo-doped BMGMCs; (c) bending stress-strain curves for both base alloy and BMGMCs (inset: XRD patterns for bending samples).

Oka et al. have successfully introduced α-Ti phase in Mg-Cu-Gd BMG system by using novel dealloying in metallic melt method [57]. The schematic of this method is shown in **Figure 15a**. For the negative value of heat of mixing (miscible) between Ti-Cu and Gd-Cu while positive value (immiscible) of Ti-Gd, Ti-Cu and Gd-Cu phases are expected to form in the prealloy. Similarly, for the negative value of heat of mixing between Mg-Gd and Mg-Cu while positive value of Mg-Ti. When the Ti-Cu-Gd prealloy consisting of Cu-Gd and Ti-Cu phases is immersed in the Mg-melt, the Cu-Gd phase and Cu dealloyed from the Ti-Cu phase are expected to dissolve. They will form the Mg-Cu-Gd BMG formable liquid if the Mg, Cu and Gd proportions are correctly balanced. The remaining elemental Ti from Ti-Cu phase is thought to form the porous structure by a surface diffusion mechanism in the Mg-Cu-Gd alloy liquid. Rapid cooling of the semi-solid mother alloy yields Mg-Cu-Gd BMG with in-situ Ti dispersoids, as shown in **Figure 15b**. By using this strategy, the in-situ Ti dispersoids have been successfully introduced. Furthermore, for the Ti dispersoids are directly dealloyed from Ti-Cu phase, thus, the size of them can be reduced by decreasing the size of Ti-Cu phase. By increasing the cooling rate of Ti-Cu-Gd prealloy, the Ti-Cu phase is refined and subsequently the size of Ti dispersoids also decreased. As shown in **Figure 15c**, during four point bending test, the composite with large and fine Ti dispersoids shows a fracture strength of 230 and 387 MPa, 6 and 78% higher than the monolithic BMG, respectively. With fine pore Ti dispersoids, the size and inter-particle spacing of Ti phase is estimated to be ~500 nm, which is very close to the characteristic plastic processing zone size of reported Mg-based BMG, 100–1000 nm. Thus, optimum condition for the composite effect has locally achieved within and surrounding the porous Ti dispersoids. Therefore, these regions and surrounded area of glassy near porous Ti could deform plastically. However, for the low volume fraction of Ti dispersoids (~2%), macroscopic plasticity is not obtained in this system. Subsequently, Guo et al. have applied similar dealloying reaction in Mg-Cu-Gd-Ag system with better glass-forming ability [49]. As shown in **Figure 16a**, homogenous distributed α-Ti phase among the glassy matrix can be observed. The average size and volume fraction are estimated to be about 6 μm and 13%, respectively. This composite shows improved mechanical properties compared with its monolithic counterpart,

that is, ~6.1% of plastic strain and 920 MPa of fracture strength during compression test, see **Figure 16b**. The stress concentration at the interfaces between Ti dispersoids and surrounding matrix is in favor of initiating multiple shear bands. Furthermore, when the stress exceeds the yield strength of Ti, it can release the stress concentration condition. The suppression of propagation of the single main shear band is enhanced by such yielding, which causes branching, blocking or multiplying the shear bands. It is therefore, the sample is deformed with a significant plasticity. The composite also shows improved fracture stress and strain during four point bending test, 331 MPa of fracture stress and 0.42% of fracture strain, as shown in **Figure 16c**. The calculated fracture toughness of this composite is ~1.73 MPa m1/2, 45% higher than its monolithic counterpart. However, it is still low and cannot lead to plasticity during bending. As stated above, the shape memory phase has attracted a lot of interest recently for its unique stress-induced martensitic transformation behavior. However, the research on shape memory phase reinforced BMGMCs mainly focus on Zr-based or Ti-based BMGMCs. Guo and Kato have successfully induced in-situ B2-NiTi shape memory phase in Mg-Ni-Gd-Ag BMGMCs by using novel selective phase leaching in metallic melt method [58]. A schematic of the novel designed process is shown in **Figure 17a**, which contains roughly three steps: Ni-Ti-Gd precursor preparation by arc-melting, master alloy preparation by induction melting and composite preparation by copper mold casting. From the Ni-Ti, Ti-Gd and Ni-Gd phase diagrams, it is possible to prepare a Ni-Ti-Gd ternary precursor consisting of only NiTi and Ni-Gd phases if the proportions of Ti, Ni, and Gd are properly balanced. Then, the temperature of master alloy preparation is kept low enough for NiTi dispersoids not to melt or dissolve, but high enough for the Ni-Gd phase to dissolve into the Mg-Ag melt, owing to their different reactivities. Moreover, the amounts of Mg-Ag melt and dissolved Ni-Gd was properly balanced to form the glass-forming matrix. Finally, Mg-based BMGMCs, an Mg-Ni-Gd-Ag BMG matrix with in-situ NiTi dispersoids, is fabricated by casting this semi-solid melt into a copper mold. By using such strategy, the in-situ B2-NiTi phase have been successfully induced, as shown in **Figure 17b**, the average size is ~8 μm and the volume fraction is ~15%. As shown in **Figure 17c**, the composite shows a higher fracture stress (~906 MPa), plastic strain (~7%), and work hardening than its monolithic counterpart. The shear bands were considered to be obstructed by the in-situ ductile NiTi dispersoids, which deflected their propagation and caused branching or multiplying of the shear bands, typically observed in ductile metal reinforced BMGMCs and is known as the "blocking effect". Moreover, a stress-induced phase transformation from

**Figure 16.** (a) SEM images of in-situ Ti reinforced Mg-based BMGMC by dealloying method (inset: XRD patterns of both BMGMC and its monolithic counterpart); (b) true compressive stress-strain curves of both BMGMC and its base alloy

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(inset: SEM images of fractured BMGMC); (c) stress-strain curves by four point banding test.

**Figure 15.** (a) Schematic showing porous Ti formation by dealloying in metallic melt; (b) schematic showing the preparation of the mother alloy; (c) stress-strain curves of BMGMCs using rapid cooling prealloy and slow cooling prealloy, under four point bending test.

Oka et al. have successfully introduced α-Ti phase in Mg-Cu-Gd BMG system by using novel dealloying in metallic melt method [57]. The schematic of this method is shown in **Figure 15a**. For the negative value of heat of mixing (miscible) between Ti-Cu and Gd-Cu while positive value (immiscible) of Ti-Gd, Ti-Cu and Gd-Cu phases are expected to form in the prealloy. Similarly, for the negative value of heat of mixing between Mg-Gd and Mg-Cu while positive value of Mg-Ti. When the Ti-Cu-Gd prealloy consisting of Cu-Gd and Ti-Cu phases is immersed in the Mg-melt, the Cu-Gd phase and Cu dealloyed from the Ti-Cu phase are expected to dissolve. They will form the Mg-Cu-Gd BMG formable liquid if the Mg, Cu and Gd proportions are correctly balanced. The remaining elemental Ti from Ti-Cu phase is thought to form the porous structure by a surface diffusion mechanism in the Mg-Cu-Gd alloy liquid. Rapid cooling of the semi-solid mother alloy yields Mg-Cu-Gd BMG with in-situ Ti dispersoids, as shown in **Figure 15b**. By using this strategy, the in-situ Ti dispersoids have been successfully introduced. Furthermore, for the Ti dispersoids are directly dealloyed from Ti-Cu phase, thus, the size of them can be reduced by decreasing the size of Ti-Cu phase. By increasing the cooling rate of Ti-Cu-Gd prealloy, the Ti-Cu phase is refined and subsequently the size of Ti dispersoids also decreased. As shown in **Figure 15c**, during four point bending test, the composite with large and fine Ti dispersoids shows a fracture strength of 230 and 387 MPa, 6 and 78% higher than the monolithic BMG, respectively. With fine pore Ti dispersoids, the size and inter-particle spacing of Ti phase is estimated to be ~500 nm, which is very close to the characteristic plastic processing zone size of reported Mg-based BMG, 100–1000 nm. Thus, optimum condition for the composite effect has locally achieved within and surrounding the porous Ti dispersoids. Therefore, these regions and surrounded area of glassy near porous Ti could deform plastically. However, for the low volume fraction of Ti dispersoids (~2%), macroscopic plasticity is not obtained in this system. Subsequently, Guo et al. have applied similar dealloying reaction in Mg-Cu-Gd-Ag system with better glass-forming ability [49]. As shown in **Figure 16a**, homogenous distributed α-Ti phase among the glassy matrix can be observed. The average size and volume fraction are estimated to be about 6 μm and 13%, respectively. This composite shows improved mechanical properties compared with its monolithic counterpart,

**Figure 15.** (a) Schematic showing porous Ti formation by dealloying in metallic melt; (b) schematic showing the preparation of the mother alloy; (c) stress-strain curves of BMGMCs using rapid cooling prealloy and slow cooling

prealloy, under four point bending test.

66 Metallic Glasses - Properties and Processing

**Figure 16.** (a) SEM images of in-situ Ti reinforced Mg-based BMGMC by dealloying method (inset: XRD patterns of both BMGMC and its monolithic counterpart); (b) true compressive stress-strain curves of both BMGMC and its base alloy (inset: SEM images of fractured BMGMC); (c) stress-strain curves by four point banding test.

that is, ~6.1% of plastic strain and 920 MPa of fracture strength during compression test, see **Figure 16b**. The stress concentration at the interfaces between Ti dispersoids and surrounding matrix is in favor of initiating multiple shear bands. Furthermore, when the stress exceeds the yield strength of Ti, it can release the stress concentration condition. The suppression of propagation of the single main shear band is enhanced by such yielding, which causes branching, blocking or multiplying the shear bands. It is therefore, the sample is deformed with a significant plasticity. The composite also shows improved fracture stress and strain during four point bending test, 331 MPa of fracture stress and 0.42% of fracture strain, as shown in **Figure 16c**. The calculated fracture toughness of this composite is ~1.73 MPa m1/2, 45% higher than its monolithic counterpart. However, it is still low and cannot lead to plasticity during bending.

As stated above, the shape memory phase has attracted a lot of interest recently for its unique stress-induced martensitic transformation behavior. However, the research on shape memory phase reinforced BMGMCs mainly focus on Zr-based or Ti-based BMGMCs. Guo and Kato have successfully induced in-situ B2-NiTi shape memory phase in Mg-Ni-Gd-Ag BMGMCs by using novel selective phase leaching in metallic melt method [58]. A schematic of the novel designed process is shown in **Figure 17a**, which contains roughly three steps: Ni-Ti-Gd precursor preparation by arc-melting, master alloy preparation by induction melting and composite preparation by copper mold casting. From the Ni-Ti, Ti-Gd and Ni-Gd phase diagrams, it is possible to prepare a Ni-Ti-Gd ternary precursor consisting of only NiTi and Ni-Gd phases if the proportions of Ti, Ni, and Gd are properly balanced. Then, the temperature of master alloy preparation is kept low enough for NiTi dispersoids not to melt or dissolve, but high enough for the Ni-Gd phase to dissolve into the Mg-Ag melt, owing to their different reactivities. Moreover, the amounts of Mg-Ag melt and dissolved Ni-Gd was properly balanced to form the glass-forming matrix. Finally, Mg-based BMGMCs, an Mg-Ni-Gd-Ag BMG matrix with in-situ NiTi dispersoids, is fabricated by casting this semi-solid melt into a copper mold. By using such strategy, the in-situ B2-NiTi phase have been successfully induced, as shown in **Figure 17b**, the average size is ~8 μm and the volume fraction is ~15%. As shown in **Figure 17c**, the composite shows a higher fracture stress (~906 MPa), plastic strain (~7%), and work hardening than its monolithic counterpart. The shear bands were considered to be obstructed by the in-situ ductile NiTi dispersoids, which deflected their propagation and caused branching or multiplying of the shear bands, typically observed in ductile metal reinforced BMGMCs and is known as the "blocking effect". Moreover, a stress-induced phase transformation from

**Figure 17.** (a) Schematic of the fabrication process; (b) SEM image of B2-NiTi reinforced mg-based BMGMC; (c) compressive true stress-strain curves for both base alloy and BMGMC.

**4. Conclusion and outlook**

versus fracture strain for various in-situ Mg-based BMGMCs to date.

3 mm.

ing materials.

In this chapter, both ex-situ and in-situ BMGMCs developed in Zr-based, Ti-based, Mg-based systems have been introduced, such as ceramic particle, metal particle, porous particle reinforced ex-situ BMGMCs and B2-phase, β-phase reinforced in-situ BMGMCs. The microstructures, mechanical properties as well as deformation mechanisms are discussed for each kind of BMGMCs. Compared with nearly zero plasticity of monolithic BMGs, the BMGMCs reinforced by secondary phases show significant improvement in plasticity, e.g., β-phase reinforced Ti-based BMGMCs show over 10% plastic strain under tension, B2-NiTi reinforced Mg-based BMGMCs show over 20% plastic strain under compression, etc. For the limitation of the chapter, more works on various reinforcements and alloy systems cannot be covered. The detailed deformation mechanisms of BMGMCs are not discussed fully either. For further development of BMGMCs, more works should be done on more complex composite structure, the deformation mechanisms, designing novel processing methods, tailoring the microstructures and mechanical properties of the existed BMGMCs. The research on BMGMCs will greatly extend the application potentials of amorphous materials as engineer-

**Figure 19.** Compressive true stress-strain curves of the optimized B2-NiTi reinforced BMGMCs; (b) fracture strength

**Figure 18.** SEM images of the composites using rapid cooling precursor: (a) 8 mm rod precursor; (b) 3 mm rod precursor; (c) compressive true stress-strain curves of the composites using rapid cooling precursor with various diameter, 8, 5 and

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B2-NiTi to B19'-NiTi during deformation releases the stress concentration around the dispersoids to restrict free volume accumulation. This process hindered the rapid propagation of shear bands, meaning that additional stress was required to move the shear bands, called the "TRIP effect". Since the B2-NiTi phase precipitates during Ni-Ti-Gd precursor preparation and has no change from then on, the size of it can be refined by increasing the cooling rate of the precursor. Thus, the precursor rods by tilt casting technique are produced with various diameters, i.e., various cooling rate. With a higher cooling rate, the size of B2-NiTi is smaller, from 8 to 2 μm, as shown in **Figure 18a** and **b**. With finer B2-NiTi phase, the composite shows higher fracture strength and larger plastic strain, as shown in **Figure 18c**. The sample with finest particle size exhibits the best mechanical properties, i.e., 1096 MPa fracture strength and 15.5% plastic strain. The volume fraction of the B2-NiTi phase has also been improved by adjusting the composition with higher Ti amount. The volume fraction increases from 15 to 32% and the optimized composite shows superior plasticity during compression, as shown in **Figure 19a**, a fracture stress of 1212 MPa and a fracture strain of 25.3%. **Figure 19b** summarizes the compressive property data of various in-situ Mg-based BMGMCs, including Fe, long-period stacking ordered structure (LPSO), NiZr, AgMg, and quasicrystal reinforced composites [59–63]. Both the fracture strength and fracture strain of the optimized B2-NiTi reinforced BMGMC are the highest among all in-situ Mg-based BMGMCs reported to date.

**Figure 18.** SEM images of the composites using rapid cooling precursor: (a) 8 mm rod precursor; (b) 3 mm rod precursor; (c) compressive true stress-strain curves of the composites using rapid cooling precursor with various diameter, 8, 5 and 3 mm.

**Figure 19.** Compressive true stress-strain curves of the optimized B2-NiTi reinforced BMGMCs; (b) fracture strength versus fracture strain for various in-situ Mg-based BMGMCs to date.
