Introduction: Refractory high-entropy alloys (RHEAs) have excellent softening resistance and thermal stability at high temperatures, but their brittleness at room temperature often limits their practical application. In this paper, room-temperature plasticization was successfully achieved by forming a heterogeneous structure (HS) in situ and doping Zr in brittle (TaMoTi)92Al8 RHEA. Different from the concept of "soft solid solution matrix and hard intermetal phase" proposed in the mainstream literature, the newly developed TaMoZrTiAl RHEA has the characteristic of embedding the hard disordered BCC phase into the soft metal inter-B2 matrix. This HS results in a significant strength-plastic synergy at room temperature, with high plasticity, greater than 20% plasticity, and greater than 2380 MPa. It is found that the heterogeneous deformation strengthening caused by solution strengthening and phase boundary dislocation accumulation is the main reason for the increase of yield strength, and the frequent operation of strain distribution and dislocation cross-slip caused by deformation greatly improves the work hardening ability of the alloy, so that the alloy has high strength and good plasticity. In conclusion, this study not only reveals the microscopic mechanism of the effect of heterogeneous duplex structure on the plasticity behavior of brittle RHEAs, but also provides a useful strategy for the plasticization of brittle RHEAs.
Refractory-entropy alloys (RHEAs) are new high-entropy alloys composed of metal elements with high melting points (>1600°C), which have attracted much attention in recent years due to their great potential as high-temperature structural materials. However, despite the recent successful design and development of some tensile ductility RHEAs based on IVB and VB transition metal elements, the vast majority of reported RHEAs, especially those rich in W/Cr/Mo elements, still face the embarrassing dilemma of room temperature (RT) brittleness, which is severely weakened compared to traditional superalloys.
A number of new strategies have been proposed, including modifying the microstructure, doping small elements (such as B or C), maximizing lattice distortion, and promoting transformation-induced plasticity (TRIP). The results show that the uniform tensile strain of the alloy is increased to 27%, which is about 7 times that of TaHfZrTiRHEA without TRIP effect. However, in the face of a large number of Titan systems, these recently proposed strategies still seem to be a drop in the bucket. On the one hand, RHEA systems capable of achieving the TRIP effect are still very limited because there is a lack of a suitable metastable phase that can undergo a martensite phase transition after loading.
Xiaoming Wang's team at the University of Chinese Academy of Sciences successfully obtained the super-RT plasticity of ZrCuNiAl bulk metallic glass by using the HS strategy. However, since the engineering of metal HS is often inseparable from the complex thermomechanical processing route, i.e., mechanical deformation and subsequent partial recrystallization, there are few relevant studies reported so far. However, the effect of this HS on the mechanical properties of brittle RHEAs at room temperature and its underlying mechanism are not well understood.
A heterogeneous duplex structure of TaMoZrTiAlRHEA was successfully prepared in situ, and its thermal compression behavior was systematically studied, focusing on the effects of thermal compression on the mechanical properties, damage mechanism and microstructure evolution of the material. The solidification process was analyzed, which provided a new way to improve the RT brittleness of RHEAs.
相关研究成果以“Significantly amelioratingroom-temperature brittleness of refractory high-entropy alloys via in situheterogeneous structure”发表在Journal of Materials Science &Technology上
链接：Significantly ameliorating room-temperature brittleness of refractory high-entropy alloys via in situ heterogeneous structure - ScienceDirect
Figure 1. Design scheme of heterogeneous duplex refractory high-entropy alloy.
Figure 2(a) pseudo-binary phase diagram. (b) a radar chart of the criteria used in the selection of higher education institutions.
Fig. 3 XRD spectra of 0Zr and 23Zr RHEAs.
Microstructure of 40Zr and 23Zr alloys. Backscatter images of (a) 0Zr and (c) 23Zr alloys. Energy spectra of (b) 0Zr and (d) 23Zr alloys. (e) EBSD phase plots of 0Zr and (f) 23Zr alloys.
(g) 0Zr和(h) 23Zr合金的EBSD IPF图。 It's a matter of opinion, GB.
BCC and B2 phase distributions were identified using the TruPhase calibration mode, which collects both structural and elemental information. Obviously, the proportion of BCC and B2 phases in the 23Zr alloy is about 55.1% and 44.9%, respectively, while only one BCC phase exists in the 0Zr alloy. In addition, the EBSD inverse pole diagram (IPF) shows that both the 0Zr and 23Zr alloys are composed of near-equiaxed grains with random orientation. Using the linear intercept method, the average grain sizes of 0Zr and 23Zr alloys are about 56 μm and 44 μm, respectively, indicating that the influence of grain size on the mechanical properties of the alloy can be excluded due to the similar grain size.
Table 1 Distribution of elements in 0Zr and 23Zr alloys.
Fig.5 TEM and SAED analysis of 23Zr alloy. (a) HAADF-STEM imagery. (b, c) (a) SAED mode corresponding to BCC and B2 in (a).
(d) HRTEM image of alloy 23Zr. (e) FFT image, corresponding to the rectangle in (d). (f) GPA graph, corresponding to (e).
Fig. 6(a) Nanoindentation morphology and hardness of 23Zr alloy.
(b) Typical load-displacement curves for BCC and B2 phases. (c) Average nanohardness and modulus values for BCC and B2 phases.
Figure 7. Mechanical properties of 0Zr and 23Zr alloys. (a) Engineering stress-strain curves. Note that the error bars represent the standard errors of yield strain and breaking strain, respectively.
(b) True stress-strain curve and corresponding work hardening rate curve. (c) Fracture morphology and corresponding energy spectrum.
Figure 7(a) shows a comparison of typical engineering compressive stress-strain curves for 0Zr and 23Zr alloys at room temperature. As expected, unlike the traditional paradigm of rhea, which typically has high strength but very limited plasticity at room temperature, the 23Zr alloy with HS effectively circumvents the trade-off between strength and plasticity.
Among them, the plasticity of 23Zr alloy containing HS is significantly increased from <8% of 0Zr matrix to >20%. The breaking strength has also been increased from 1408 MPa to 2388 MPa of 0Zr alloy. Fig. 7(b) shows the work hardening rate (θ = dσtrue/dεtrue) curves for the 0Zr and 23Zr alloys, and different work hardening behaviors can be observed. The work hardening rate curve of 23Zr alloy shows a multi-stage work hardening behavior. In any case, this multi-stage work hardening behavior is thought to give the 23Zr alloy a strong continuous work hardening capacity, resulting in a good strength-plastic match at the rt. The 0Zr alloy only leads to a sharp reduction in the work hardening rate, which leads to the early appearance of plastic instability.
Fig. 8(a) LUR true stress-strain curves for 0Zr and 23Zr alloys. (b) Relationship between HDI stress and plastic true strain in 0Zr alloy. (c) Relationship between HDI stress and plastic true strain in 23Zr alloy.
(d) Relationship between HDI stress/flow stress and plastic true strain of 0Zr and 23Zr alloys.
Figure 90Zr and 23Zr alloy surface deformation and damage characteristics. SEM images of (a) 0Zr and (b) 23Zr alloys. (c) EBSD KAM plots of the 0Zr and (d) 23Zr alloys.
(b) (d1) 0Zr和(d2) 23Zr合金的KAM值统计分布。 (g) 0Zr和(h) 23Zr合金的EBSD相图。
The TEM analysis of the deformed 0Zr and 23Zr alloys in Figure 10 shows the evolution of the defect as the strain increases. (a) 0Zr alloy deformed at 5% strain. (b, c) 0Zr alloy deformed at 8% strain. (d) 23Zr alloy deformed at 10% strain. (e, g) 23Zr alloy deformed at 20% strain. (f) Tamo-rich BCC phase dislocation in 23Zr alloy. (h) Schematic diagram of the formation mechanism of dislocation loops.
(i) 23Zr合金 化 意图。
Apparently, in order to accommodate the strain gradient between the elastic BCC phase and the plastic B2 phase, a number of dislocations are accumulated at the B2/BCC interface (Figure 10(d)). When the strain increases to 20%, both the BCC phase and the B2 phase produce obvious plastic deformation. This is evidenced by the role of the median dislocation in the hard BCC phase.
The dislocation slip of the B2 phase becomes more frequent and active compared to the early stage of plastic deformation, resulting in a high density of dislocation tangles and the formation of dislocation cells (Fig. 10(g)). In addition, some dislocation loops were found in the local region (Fig. 10(g)), which may be attributed to the double-cross slip mechanism of the screw dislocation. As shown in Fig. 10(h), when the screw dislocation slip is blocked, the local chemical fluctuation characteristics of the HEA promote the screw dislocation to transfer to another slip surface with a low energy barrier. The transfer occurs through cross-slip, which results in further slippage.
Table 2 Calculated values of empirical parameters (ΔSmix, ΔHmix, δ, VEC, Ω, Λ) of BCC and B2 phases in 23Zr alloy.
Fig. 11(a) Equilibrium phase diagram of 23Zr alloy. (b) Temperature vs. molar fraction of the liquid under non-equilibrium conditions, showing the segregation/depletion behavior of individual elements.
(c) Schematic diagram of the solidification process of 23Zr alloy.
Fig. 12(a) Strain contour plots of 0Zr and 23Zr alloys after fracture. (b) Enlarged view of the pink borderline area in (a). (c) Statistical distribution of the maximum orientation deviation of 0Zr and 23Zr alloys.
(b) Statistical distribution of the maximum orientation deviation of the BCC phase and the B2 phase in the 23Zr alloy.
Table 3: Relative modulus difference (underlined) ηij and atomic size difference (bold) δij for alloying element pairs.
Figure 13. Calculated and experimental values of σy for 0Zr and 23Zr alloys
。
Enhancement contributions of σ0, σss, σ0, and σgb can be obtained. After subtracting the above strengthening contribution from the σsyn value, the strengthening contribution of 0Zr and 23Zr alloys is about 12 and 290 MPa, respectively. In summary, the contribution of each reinforcement mechanism to σy is shown in Figure 13. Compared with the 0Zr alloy, the main components that are significantly enhanced in the σy of the 23Zr alloy are σss caused by the addition of Zr and σsyn caused by HS.
Fig. 14 Microstructure study of the B2 phase in a compressed 23Zr alloy near the axis in the [1 0 0] region under the two-beam condition (g = (1 0 0) (a), (0 1 3) (b), and (0 3 2) (c) TEM brightfield images taken near the [1 1 1 1] band axis at g = (1 0 1) (d), (0 1 1 1) (e) and (1 1 2) (f) show a typical B2 phase dislocation structure.
IN THIS STUDY, A METHOD WITH IN-SITU FORMING WAS SUCCESSFULLY PREPARED BY USING THE PHASE SEPARATION CHARACTERISTICS AND THE CALPHAD METHOD
A new type of HS
The HS is composed of a ductile (Zr,Al) B2 matrix and a hard (Ta,Mo) BCC phase. The main achievements are as follows:
(1) It effectively solves the problem that it is difficult to obtain HS in brittle RHEAs in the traditional thermomechanical processing route.
(2) The obtained HS makes the strength and plasticity of 23Zr alloy well matched. The plasticity at room temperature was significantly increased from <8% to >20%, an increase of nearly 150%, indicating that the HS strategy was effective in improving the brittleness of RHEAs at room temperature.
(3) Due to the incompatibility of the plastic strain between the B2 phase and the BCC phase during the deformation process, a large number of gds are formed. Continuous, strong HDI hardening is obtained, which improves the work hardening capacity and thus improves the RT plasticity of the 23Zr alloy.
(4) It provides a basic understanding of the deformation mechanism of heterogeneous two-phase RHEAs, and provides a useful strategy for the plasticization of brittle RHEAs.
Reprint: Materials Science Network