A Ti-based amorphous composite material for overcoming relaxation-induced embrittlement and a preparation method thereof

Abstract

The present invention belongs to the field of amorphous alloy composites, and specifically relates to a Ti-based amorphous composite material that overcomes relaxation-induced embrittlement and a preparation method thereof. The amorphous composite material is an in-situ shape memory phase Ti-based amorphous in-situ composite material. The shape memory phase Ti-based amorphous in-situ composite material means that the in-situ β-Ti phase in the amorphous composite material is a metastable phase and can undergo deformation-induced α″ martensitic transformation. First, an alloy ingot is prepared by arc melting in a high-purity argon environment, and the alloy ingot is remelted at least four times to ensure the uniformity of the components. Then, in the high-purity argon environment of the arc furnace, the alloy ingot is remelted, and an alloy rod is obtained by copper mold flipping casting. The present invention avoids the intrinsic problem of relaxation-induced embrittlement of amorphous alloys or amorphous composite materials under the preparation conditions of a lower cooling rate, helps to break through the bottleneck of relaxation-induced embrittlement of the amorphous phase, and has important value for developing large-size high-toughness amorphous composite materials and realizing their applications.

Description

Technical Field

[0001] The present invention belongs to the field of amorphous alloy composites, and specifically relates to a Ti-based amorphous composite material that overcomes relaxation-induced embrittlement and a preparation method thereof. Background Art

[0002] Bulk amorphous alloys do not have the periodic characteristics of atomic packing in crystalline materials, and their unique amorphous structure endows them with excellent mechanical properties, such as high strength, high hardness, and large elasticity, and are regarded as structural materials with broad prospects. However, under uniaxial tension and conditions far below their glass transition temperature, amorphous alloys usually exhibit zero tensile ductility because the plastic strain in amorphous alloys is highly concentrated in narrow shear bands, and the propagation of shear bands is rapid and unstable, resulting in catastrophic failure. To break this bottleneck, amorphous composites containing in-situ formed crystals have been developed to promote the proliferation of shear bands and prevent their propagation. Among all amorphous composites, Ti / Zr-based amorphous composites with β-Ti / Zr dendrites have attracted great attention because their microstructures and the metastability of dendrites can be adjusted to achieve large tensile ductility and high strain hardening ability. However, with the decrease of the cooling rate during preparation, obvious structural relaxation occurs in the amorphous phase, which leads to a significant decrease in the ductility and fracture toughness of amorphous alloys and amorphous composites, and even the phenomenon of ductile-brittle transition occurs. This greatly limits the casting size of amorphous alloys and amorphous composites, and further restricts their applications in large-scale structural materials. Therefore, it is urgent to systematically study the mechanism of relaxation-induced embrittlement in the amorphous phase and explore methods to break through this limitation.

[0003] Recent studies have shown that if the in-situ formed β-Ti phase is a metastable phase and can undergo deformation-induced martensitic transformation during deformation and cause the shear band passivation effect, then Ti-based amorphous in-situ composites can exhibit excellent tensile work hardening ability. In addition, this shear band passivation effect depends on the dendrite size, and the dendrite size is negatively correlated with the cooling rate. Therefore, regulating the microstructure of in-situ shape memory phase Ti-based amorphous in-situ composites provides a solution to overcome relaxation-induced embrittlement of the amorphous phase. Summary of the Invention

[0004] The purpose of the present invention is to provide a large-size Ti-based amorphous composite material that overcomes relaxation-induced embrittlement and a preparation method thereof, and to study the variation law of the fracture toughness of amorphous alloys and amorphous composites by controlling the cooling rate during preparation.

[0005] To achieve the above purpose, the technical solutions adopted by the present invention are as follows:

[0006] A Ti-based amorphous composite material that overcomes relaxation-induced embrittlement, and the chemical composition of the amorphous composite material is: Ti 47.2Zr 33.8 Cu 5.8 Co1Be 12.2 (at. %), the cooling rate was changed by preparing alloy rods with different diameters (8 mm and 20 mm), which were labeled as BT50-8 and BT50-20 respectively. The same method as above was used to prepare amorphous alloys for comparing the influence of the cooling rate. The chemical composition of the amorphous alloy was: Ti 32 Zr 30.2 Cu9Co 4.8 Be 24 (at. %), which were labeled as BT0-8 and BT0-20 respectively.

[0007] In the shape memory phase Ti-based amorphous in-situ composite, the amorphous matrix has a high glass-forming ability, and the alloy cast in a copper mold is the structure of the amorphous in-situ composite. The technical index for characterizing the glass-forming ability of the amorphous matrix is: for alloy rods with diameters of 8 mm and 20 mm cast in a copper mold, there is no crystallization phenomenon in the amorphous matrix. Such shape memory phase Ti-based amorphous in-situ composites refer to: the β-Ti phase in-situ in the amorphous composite is a metastable phase and can undergo deformation-induced α″ martensitic transformation. Therefore, such amorphous composites are large-size shape memory phase Ti-based amorphous in-situ composites. In the bulk Ti-based amorphous alloy, it has a high glass-forming ability, and the alloy cast in a copper mold is an amorphous alloy. The technical index for characterizing its glass-forming ability is: the diameters of the alloy rods obtained by copper mold casting are 8 mm and 20 mm respectively, and there is no crystallization in the center and edge parts.

[0008] In such large-size Ti-based amorphous alloys, with the decrease of the cooling rate, its tensile strength slightly decreases from 1764 ± 15 MPa of BT0-8 to 1727 ± 18 MPa of BT0-20, but the notch toughness of the amorphous alloy significantly decreases: from 147 ± 14 MPa m 1 / 2 decreases to 118 ± 10 MPa m of BT0-20 1 / 2 .

[0009] In such large-size shape memory phase Ti-based amorphous in-situ composites, with the decrease of the cooling rate, its yield strength increases from 1084 ± 15 MPa of BT50-8 to 1170 ± 18 MPa of BT50-20, the tensile strength increases from 1378 ± 15 MPa of BT50-8 to 1393 ± 18 MPa of BT50-20, while the tensile plasticity decreases from 8.0 ± 0.3% of BT50-8 to 8.6 ± 0.4% of BT50-20. Moreover, the notch toughness of the shape memory phase amorphous Ti-based composite increases with the decrease of the cooling rate: from 106 ± 10 MPa m of BT50-8 1 / 2Slightly increased to 108±12 MPa m of BT50-20 1 / 2 。

[0010] The principle by which the shape memory phase Ti-based amorphous in-situ composite of this kind in the present invention overcomes relaxation-induced embrittlement is as follows:

[0011] (1) In such large-sized Ti-based amorphous alloys, the reasons for the lower toughness at a lower cooling rate: The structure and energy state of the amorphous alloy are closely related to the cooling rate. At a lower cooling rate, the supercooled liquid evolves towards a more equilibrium state, resulting in a more relaxed and lower-energy-state amorphous alloy. Since the cooling rate of BT0-8 is more than 6 times higher than that of BT0-20, BT0-20 has a more relaxed structure and a lower energy state, which makes it more difficult to induce shear bands and results in a smaller plastic zone. Therefore, the toughness of the amorphous alloy is extremely sensitive to the cooling rate, and reducing the cooling rate is likely to induce ductile-brittle transition.

[0012] (2) In such large-sized shape memory phase Ti-based amorphous in-situ composites, the reasons for the larger fracture toughness at a lower cooling rate: The size of β-Ti grains in the Ti-based amorphous composite increases with the decrease of the cooling rate. Larger-sized β-Ti crystals are more likely to induce a thicker martensitic transformation by deformation and a higher degree of shear band blunting effect. This toughening effect can overcompensate for the relaxation brittleness of the amorphous matrix, thus enabling the Ti-based amorphous composite to overcome the relaxation embrittlement phenomenon.

[0013] The advantages and beneficial effects of the present invention are:

[0014] 1. For the Ti-Zr-Cu-Co-Be amorphous in-situ composite described in the present invention, its amorphous matrix has a high glass-forming ability and is easy to obtain a larger size. The diameters of the alloy rods obtained by copper mold casting are 8 mm and 20 mm respectively, and their microstructures are β-Ti uniformly distributed in the amorphous matrix.

[0015] 2. By regulating the cooling rate, the present invention studies the fracture toughness of amorphous alloys and shape memory phase amorphous composites, and finds that in amorphous alloys, the fracture toughness decreases significantly with the decrease of the cooling rate, which is called relaxation-induced brittleness.

[0016] 3. The present invention finds that in such large-sized shape memory phase Ti-based amorphous in-situ composites, a larger fracture toughness is obtained at a lower cooling rate, which shows that the higher degree of phase transformation toughening effect and shear band blunting effect of large-sized β-Ti can overcompensate for the relaxation brittleness of the amorphous matrix.

[0017] In summary, the present invention discloses a new strategy to overcome relaxation-induced embrittlement in the amorphous phase. It is found that reducing the cooling rate during the preparation of amorphous alloys or amorphous composites will seriously lead to the structural relaxation of the amorphous phase. Amorphous alloys in the low-energy state undergo brittle fracture due to the difficulty in initiating shear bands. However, for shape-memory-phase amorphous composites, a decrease in the cooling rate is beneficial for dendrite coarsening, which leads to a more likely occurrence of deformation-induced martensitic transformation. In such shape-memory-phase amorphous composites, the martensitic transformation induces a shear-band passivation effect, which is highly sensitive to the dendrite size. And this shape-memory-phase amorphous composite exhibits a higher degree of shear-band passivation at a lower cooling rate, compensating for the defect of fewer shear bands in its low-energy-state amorphous phase. The enhanced shear-band passivation effect is not only beneficial for plasticization but also toughens the material by greatly hindering crack propagation. The higher toughness of this shape-memory-phase amorphous composite at a lower cooling rate is attributed to the significant crack deflection and bifurcation caused by the coarser dendrites undergoing phase transformation, especially the newly emerged numerous crack bridges that greatly promote crack passivation. BRIEF DESCRIPTION OF THE DRAWINGS

[0018] Figure 1 It is the microstructural diagram of the Ti-based amorphous alloy and amorphous in-situ composite of the present invention; wherein: (a) and (b) are the scanning electron micrographs of BT0-8 and BT0-20 respectively, and the insets are the corresponding XRD patterns; (c) and (d) are the high-resolution transmission electron micrographs of BT0-8 and BT0-20 respectively; (e) and (f) are the scanning electron micrographs of BT50-8 and BT50-20 respectively, and the insets are the corresponding XRD patterns; (g) and (h) are the transmission electron micrographs of BT50-8 and BT50-20 respectively.

[0019] Figures 2 - 3 It is the mechanical property diagram of the BT0 amorphous alloy and the BT50 amorphous composite. Among them: Figure 2 They are the typical engineering tensile stress-strain curves of the BT0 amorphous alloy and the BT50 amorphous composite with diameters of 8 mm and 20 mm, and the inset shows the size parameters of the tensile samples. Figure 3 It is the comparison diagram of the notch toughness and relaxation enthalpy of the BT0 amorphous alloy and the BT50 amorphous composite with diameters of 8 mm and 20 mm.

[0020] Figure 4Microstructures of the BT0 amorphous alloy and BT50 amorphous composite after fracture. Among them: (a) and (b) are scanning electron micrographs of BT0-8 and BT0-20 after tensile fracture, respectively; (c) and (d) are scanning electron micrographs of BT50-8 and BT50-20 after tensile fracture, respectively, and the insets are the corresponding XRD patterns; (e) and (f) are scanning electron micrographs of the notch tip regions of BT0-8 and BT0-20 after fracture in Mode I toughness test, respectively, and the insets are magnified plastic deformation images of the corresponding regions; (g) and (h) are scanning electron micrographs of the notch tip regions of BT50-8 and BT50-20 after fracture in Mode I toughness test, respectively.

[0021] Figure 5 Microstructures of the BT50 amorphous composite after the Mode I fracture toughness test. Among them: (a) and (b) are scanning electron micrographs of the plastic zones about 400 μm in front of the notches of BT50-8 and BT0-20 after fracture, respectively; (c) and (d) are scanning electron micrographs of the plastic zones in front of the microcracks of BT50-8 and BT50-20 after fracture, respectively.

[0022] Figure 6 Transmission electron micrographs of the Mode I fracture of the BT50-8 and BT50-20 amorphous composites. Among them: (a) is the transmission electron micrograph of the BT50-8 amorphous composite after fracture; (b) and (c) are the transmission electron micrographs of the BT50-20 amorphous composite after fracture at different positions, respectively; (d) is the selected area diffraction pattern of the α″-Ti martensite laths; (e) and (f) are the dark field transmission electron micrographs of two α″-Ti variants, respectively. Detailed implementation manners

[0023] In the specific implementation process, the present invention regulates the fracture toughness by changing the cooling rate of the amorphous alloy and the amorphous composite. The specific preparation method is as follows:

[0024] Ti-Zr-Cu-Co-Be amorphous alloy and amorphous in-situ composite are used, and their naming method is to be marked according to the dendritic mole fraction and the preparation size. The chemical composition of this amorphous alloy is: Ti 32 Zr 30.2 Cu9Co 4.8 Be 24 (at.%), marked as BT0, and the cooling rate is changed by preparing alloy rods with different diameters (8 mm and 20 mm), which are marked as BT0-8 and BT0-20 respectively; the chemical composition of this amorphous composite is: Ti 47.2 Zr 33.8 Cu 5.8 Co1Be 12.2(at.%), the cooling rate was changed by preparing alloy rods with different diameters (8 mm and 20 mm), labeled as BT50-8 and BT50-20, respectively. Among them, the cooling rate for preparing the 8-mm-diameter alloy rod was ~63 K / s -1 , and the cooling rate for preparing the 20-mm-diameter alloy rod was ~10 K / s -1 .

[0025] The raw materials used were sponge Ti and sponge Zr of industrial purity, and the purity of the remaining elements was higher than 99.8 wt%. Alloy ingots were prepared by arc melting in an environment of high-purity argon (volume purity 99.999%, 0.01-0.1 MPa). The alloy ingots were remelted at least four times to ensure the elemental homogeneity of the alloy. Then, in the environment of high-purity argon in an arc furnace, the alloy ingots were remelted, and alloy rods with diameters of 8 mm and 20 mm were obtained by copper mold flipping casting. First, slices with a thickness of approximately 1 mm were cut from the as-cast alloy rods, and the as-cast microstructure was characterized. Then, fracture toughness samples with dimensions of 2.5 mm × 5 mm × 25 mm were cut from the alloy rods, and the microstructure after fracture was characterized.

[0026] The mechanical properties of the Ti-based amorphous composite material of the present invention are shown in Table 1 below:

[0027] Table 1

[0028]

[0029] The present invention will be described in detail below with reference to the accompanying drawings and embodiments.

[0030] Example 1

[0031] According to the designed chemical composition, Ti 32 Zr 30.2 Cu9Co 4.8 Be 24 (at.%), labeled as BT0, amorphous alloys with different diameters (8 mm and 20 mm) were prepared, labeled as BT0-8 and BT0-20, respectively. The raw materials used were sponge Ti and sponge Zr of industrial purity, and the purity of the remaining elements was higher than 99.8 wt%. Alloy ingots were prepared by arc melting in an environment of high-purity argon (volume purity 99.999%, 0.01-0.1 MPa). The alloy ingots were remelted at least four times to ensure the homogeneity of the components. Then, in the environment of high-purity argon in an arc furnace, the alloy ingots were remelted, and alloy rods with diameters of 8 mm and 20 mm were obtained by copper mold flipping casting. First, slices with a thickness of approximately 1 mm were cut from the as-cast alloy rods, and the as-cast microstructure was characterized. Then, fracture toughness samples with dimensions of 2.5 mm × 5 mm × 25 mm were cut from the alloy rods, and the microstructure after fracture was characterized.

[0032] Figure 1 (a) and (b) show that the SEM micrographs of as-cast BT0-8 and BT0-20 are homogeneous and no second phase is observed. Figure 1 The XRD patterns shown in the insets of (a) and (b) consist of broad peaks, indicating that the samples are amorphous. In addition, the compositions of the two amorphous alloys do not change at different cooling rates. Figure 1 The high-resolution transmission electron micrographs of the two amorphous alloys in (c) and (d) also show typical maze-like patterns without any crystallization traces, and the selected area electron diffraction patterns are also halos (see the insets in (c) and (d)). Figure 1 (c) and (d) in the illustration).

[0033] Figure 2 Shows the typical tensile stress-strain curves of BT0 amorphous alloy and BT50 amorphous composite prepared at different cooling rates. Figure 2 The inset shows the geometric schematic diagram of the small-sized tensile specimens and dimensional parameters used in this study. For the BT0 amorphous alloy, all specimens only exhibit elastic deformation behavior and catastrophic fracture under tensile load without yielding. BT0-8 shows an ultimate tensile strength of 1764 ± 15 MPa at a fracture strain of 1.9 ± 0.3%, while BT0-20 shows a slightly lower ultimate tensile strength of 1727 ± 18 MPa and a fracture strain of 1.6 ± 0.2%. The tensile properties of BT0 are listed in Table 1 for easy reading. However, the BT50 amorphous composite exhibits large tensile plasticity and work-hardening ability, which is attributed to the interaction between the shear bands of the amorphous matrix and the dendrites. The yield strength of BT50-8 is 1084 ± 15 MPa, and after yielding, BT50-8 exhibits large tensile ductility and extraordinary strain-hardening ability until it reaches its ultimate tensile strength of 1378 ± 16 MPa. BT50-8 finally fractures at a strain of 8.6 ± 0.4% and experiences a negligible short softening stage. In contrast, BT50-20 yields at a higher strength of 1084 ± 15 MPa, and after yielding, it undergoes a slightly shorter work-hardening stage until it reaches its ultimate tensile strength of 1378 ± 16 MPa. Then BT50-20 plastically deforms beyond the ultimate tensile strength in a work-softening manner and finally fails at a strain of 8.0 ± 0.3%. For clearer and easier reading, the tensile properties of the BT50 amorphous composite are also summarized in Table 1. Figure 3 Shows that the fracture toughness of the amorphous alloy decreases significantly from 147 ± 14 MPa m of BT0-8 1 / 2 to 118 ± 10 MPa m of BT0-20 1 / 2 , indicating that the lower cooling rate severely deteriorates the fracture toughness of the amorphous alloy.

[0034] Figure 4(a) and (b) respectively show the scanning electron micrographs of the regions near the tensile fracture surfaces of the tensile-fractured BT0-8 and BT0-20 amorphous alloys. Although the BT0 amorphous alloy lacks macroscopic tensile ductility, shear bands and microcracks evolved from the shear bands can be observed near the main crack. However, the number density of shear bands in the BT0-8 tensile specimen is higher than that in BT0-20, significantly decreasing from ∼0.2 μm in BT0-8 -1 to ∼0.05 μm in BT0-20 -1 . Figure 4 (c) and (d) respectively show the scanning electron images of the notch tip regions of the BT0-8 and BT0-20 amorphous alloys after the mode-I fracture toughness tests. Macroscopically, the fractured samples of both amorphous alloys show tortuous crack paths, and a large number of curved shear bands can be observed at the notch tip. Although the initial directions of the main cracks of both amorphous alloys are ∼30°, the number density of shear bands in BT0-8 is twice that in BT0-20. In addition, the plastic zone of BT0-8 is about 300 μm, while that of BT0-20 is about 210 μm, indicating that the plastic deformation ability of BT0-8 is better than that of BT0-20. Compared with BT0-20, the large number of primary shear bands, secondary shear bands and shear band branches formed in BT0-8 lead to more crack deflection and bifurcation, thus increasing the crack propagation path in BT0-8. Therefore, the toughness of the amorphous alloy is extremely sensitive to the cooling rate, and reducing the cooling rate is very likely to induce the ductile-brittle transition.

[0035] Example 2

[0036] According to the principle of two-phase quasi-equilibrium solidification, a β-Ti amorphous in-situ composite with a molar fraction of 50% was designed, in which β-Ti can undergo deformation-induced martensitic transformation. The chemical composition of this amorphous composite is: Ti 47.2 Zr 33.8 Cu 5.8 Co1Be 12.2 (at.%). By preparing alloy rods with different diameters (8 mm and 20 mm) to change the cooling rate, they are respectively labeled as BT50-8 and BT50-20. The raw materials used are industrial purity sponge Ti and sponge Zr, and the purity of the remaining elements is higher than 99.8 wt%. Alloy ingots are prepared by arc melting in an environment of high-purity argon (volume purity 99.999%, 0.01 - 0.1 MPa), and the alloy ingots are remelted at least four times to ensure the uniformity of the components. Then, in the environment of high-purity argon in the arc furnace, the alloy ingots are remelted, and alloy rods with diameters of 8 mm and 20 mm are obtained by the method of copper mold flipping casting. First, slices with a thickness of about 1 mm are cut from the as-cast alloy rods, and the as-cast microstructure is characterized. Then, fracture toughness samples with dimensions of 2.5 mm × 5 mm × 25 mm are cut from the alloy rods, and the microstructure after fracture is characterized.

[0037] Figure 1 (e) and (f) show the microstructures of as-cast BT50 alloys with diameters of 8 mm and 20 mm, which are composed of β-Ti dendrites embedded in a continuous amorphous matrix, as confirmed by the diffraction peaks in the XRD patterns (see Figure 1 the insets in (e) and (f)). The measured volume fractions of β-Ti dendrites in the two amorphous composites are similar at 56 ± 2%, but the average dendrite size (secondary dendrite arm) varies from 4.3 ± 1.1 μm in BT50-8 to 9.8 ± 1.6 μm in BT50-20. As Figure 1 (g) and (h) show, similar as-cast microstructures of the two amorphous composites are further observed by TEM. The β-Ti dendrites in BT50-8 are nearly perfect body-centered cubic (BCC) crystals with some particularly weak diffraction spots at the 1 / 3 and 2 / 3 {112} β positions, which can be indexed as two of the four adiabatic ω-Ti variants, as shown by the selected area electron diffraction in Figure 1 (g). In contrast, in addition to ω-Ti observed at the same positions under the

[110] β zone axis, α″-Ti diffraction spots can also be found, as shown in Figure 1 (h). The formation of α″-Ti martensite laths in BT50-20 implies that the phase stability of β-Ti in BT50-20 is much lower than that in BT50-8.

[0038] The BT50 amorphous composites exhibit large tensile plasticity and work-hardening ability, which is attributed to the interaction between the shear bands of the amorphous matrix and the dendrites. The yield strength of BT50-8 is 1084 ± 15 MPa, and after yielding, BT50-8 exhibits large tensile ductility and extraordinary strain-hardening ability until reaching its ultimate tensile strength of 1378 ± 16 MPa. BT50-8 finally fractures at a strain of 8.6 ± 0.4% and experiences a negligible short softening stage. In contrast, BT50-20 yields at a higher strength of 1084 ± 15 MPa, and after yielding, it undergoes a slightly shorter work-hardening stage until reaching its ultimate tensile strength of 1378 ± 16 MPa. Then BT50-20 plastically deforms in a work-softening manner beyond the ultimate tensile strength and finally fails at a strain of 8.0 ± 0.3%. For clarity and readability, the tensile properties of the BT50 amorphous composites are also summarized in Table 1. Figure 3 The fracture toughness of the BT50 amorphous composites prepared at different cooling rates is shown. It can be clearly seen that the fracture toughness values of the BT50 amorphous composites are different from the dependence of the BT0 amorphous alloy on the cooling rate, with the latter ranging from 106 ± 10 MPa m for BT50-8 1 / 2Slightly increased to 108 ± 12 MPa m of BT50-20 1 / 2 。

[0039] Compared with the amorphous alloy only having shear bands after tensile fracture, the post-fracture microstructures of BT50-8 and BT50-20 amorphous composites show many surface undulations caused by martensite plates within β-Ti dendrites, as well as high-density shear bands in the amorphous matrix, as shown in Figure 4 (e) and (f). As shown in the XRD pattern, the martensite phase is α″-Ti (see the insets in Figure 4 (e) and (f)), which means that under the confinement of the amorphous matrix, the β-Ti dendrites still undergo deformation-induced martensite transformation to increase plasticity.

[0040] Figure 4 (g) and (h) respectively show the scanning electron micrographs of the notch tip regions of the post-fracture BT50-8 and BT50-20 amorphous composites. Similar to the case of the BT0 amorphous alloy, a curved crack path is observed in both amorphous alloys, but it is more obvious in BT50-20 than in BT50-8. However, the main crack initiation direction of BT50-8 is ~67°, which is more than twice that of BT50-20 at ~30°. In addition, they also generate secondary cracks at different angles at the notch tip. Contrary to the long and continuous curved shear bands formed at the notch tip of the BT0 amorphous alloy under Mode I, many shear bands in the two BT50 amorphous composites are confined in the amorphous ligaments between β-Ti crystals with high-density surface protrusions, as shown in the insets. Moreover, since β-Ti can undergo martensitic phase transformation under low stress, the BT50 amorphous composites have a larger plastic zone, and surface relief and shear bands can still be observed at a distance of ~400 μm from the notch, as shown in Figure 5 (a) and (b). It can be clearly observed that this large amount of plastic deformation behavior leads to crack deflection, bifurcation, etc., especially compared with the crack directly passing through the fine grains in BT50-8, while the large-sized dendrites in BT50-20 can interrupt crack propagation (see Figure 5 (c) and (d)). In addition, due to the effective hindrance of the larger dendrites in the BT50-20 amorphous composite, a large number of crack bridges that do not appear in BT50-8 can be observed.

[0041] Figure 6(a) and (b) show the transmission electron micrographs of the BT50-8 and BT50-20 amorphous composites after the fracture toughness tests. Parallel martensite laths can be seen in both amorphous composites. Apparently, the α″-Ti martensite laths in BT50-20 are much thicker than those in BT50-8, and the thicker laths shear cooperatively with the shear bands, resulting in obvious shear steps. α″-Ti variants with different plate orientations are observed in BT50-8, while there are more variants in BT50-20, as shown in Figure 6 (c). Figure 6 (d) shows the selected area diffraction spot pattern from the α″-Ti martensite, and two variants satisfy specific crystallographic orientations. Figure 6 (e) and (f) show the morphologies of the two variants. The two variants shear each other to coordinate deformation, which is beneficial to plastic deformation. In this shape memory phase amorphous composite, the martensitic transformation induces a shear band blunting effect, and the larger dendrite size has a higher degree of shear band blunting. Therefore, this shape memory phase amorphous composite shows a higher degree of shear band blunting at a lower cooling rate, compensating for the defect of fewer shear bands in its low-energy amorphous phase. Moreover, this shape memory phase amorphous composite has higher toughness at a lower cooling rate. On the one hand, the shear band blunting is beneficial to plastic deformation, increasing the plastic deformation work. On the other hand, significant crack deflection and bifurcation are caused by the coarser dendrites undergoing phase transformation, especially the newly emerging large number of crack bridges greatly promote crack blunting. Therefore, this large-size shape memory phase amorphous composite shows the characteristic of overcoming relaxation-induced brittleness with the decrease of the cooling rate.

[0042] The implementation results show that the present invention discovers, by controlling the cooling rate during preparation, that at a low cooling rate, brittle fracture occurs due to structural relaxation. However, the in-situ shape memory phase Ti-based amorphous in-situ composite undergoes dendrite coarsening at a low cooling rate, which enhances the degree of shear band blunting caused by the crystal martensitic transformation, resulting in a higher degree of shear band blunting effect in the large-size shape memory phase amorphous composite prepared at a low cooling rate. This not only increases the plastic deformation ability but also enhances the toughening mechanisms of crack deflection, bifurcation, and bridging, thus avoiding the inherent problem of relaxation-induced embrittlement in amorphous alloys or amorphous composites under the preparation conditions of a lower cooling rate. The present invention helps to break through the bottleneck of relaxation-induced embrittlement of the amorphous phase and has important value for developing large-size high-toughness amorphous composites and realizing their applications.

Claims

1. A Ti-based amorphous composite material for overcoming relaxation embrittlement, characterized in that, The amorphous composite material is an in-situ shape memory phase Ti-based amorphous in-situ composite material. The in-situ shape memory phase Ti-based amorphous in-situ composite material means that the in-situ β-Ti phase in the amorphous composite material is a metastable phase and can undergo deformation-induced α″ martensitic transformation. The chemical composition of the amorphous composite material is in atomic percentage: Ti 47.2 Zr 33.8 Cu 5.8 Co1Be 12.2 .

2. The Ti-based amorphous composite material for overcoming relaxation embrittlement according to claim 1, wherein, The in-situ shape memory phase Ti-based amorphous in-situ composite material obtains alloy rods of different sizes by copper mold casting. The diameters of the alloy rods are 8 mm and 20 mm respectively. The amorphous matrix has high glass-forming ability and no crystallization. The alloy cast by copper mold is the microstructure of the amorphous in-situ composite material.

3. The preparation method of the Ti-based amorphous composite material for overcoming relaxation embrittlement according to any one of claims 1 to 2, characterized in that, This method first prepares an alloy ingot by arc melting in a high-purity argon environment. The alloy ingot is remelted at least four times to ensure the uniformity of the components. Then, in the high-purity argon environment of the arc furnace, the alloy ingot is remelted, and alloy rods are obtained by the method of copper mold inverted casting.

4. The preparation method of the Ti-based amorphous composite material for overcoming relaxation embrittlement according to claim 3, characterized in that When using copper mold inverted casting, reducing the cooling rate enables this amorphous composite material to have higher fracture toughness.

Images