A fundamentally important objective within the 昀椀eld of materials science is to resolve the inherent trade-off between the strength and ductility of advanced metallic materials. Typically, metallic materials are subject to strain-rate effects wherein the yield strength increases, and the ductility decreases with an increase in the strain rate [1]. Furthermore, as the strain rate rises, the in昀氀uence of thermal activation on dislocations declines while the effect of viscous drag on dislocation motion becomes more prominent [2]. Both factors would result in an increase in the Peierls-Nabarro stress [3], hence contributing to the overall strengthening of the material. During the process of dynamic deformation, the rapid multiplication of high-density dislocations impedes their ability to glide easily. As a result, a signi昀椀cant accumulation of dislocations occurs, leading to the concentration of stress and the propagation of cracks at the initial stage of deformation [4]. Furthermore, as the strain rate increases, both the shear stress and the adiabatic effect increase signi昀椀cantly, resulting in severe deformation localization and the generation of adiabatic shear bands (ASBs) [5]. Accordingly, the ductility of the material is impaired. Hence, the enduring problem in the 昀椀eld of metallic materials pertains to the trade-off between strength and ductility with increasing strain rate.
In recent years, researchers have introduced and created a new category of alloys known as high-entropy alloys (HEAs). These alloys are formed by combining multiple elements, typically four or more, in nearequiatomic proportions ranging from 5 to 35 atomic percent (at.%) [6]. The HEA strategy has effectively overcome these limits imposed by traditional alloy-design methods, such as limited base elements and complicated microstructures, resulting in a signi昀椀cant increase of alloydesign possibilities [7]. Through the manipulation of the composition, scientists possess the capability to modify the microstructures of HEAs in order to attain speci昀椀c mechanical properties. For instance, Al [8], Mo [9], Ti [10], V [11], Zr [12], Ta [13], Nb [14], and Y [15] have been added to the Cantor alloy as solutes, resulting in a multi-phase structure, which enhanced the strength of the material. The hardening of the alloys by adding elements as solutes can be attributed to a combination of solid-solution strengthening, second-phase strengthening, and 昀椀negrain strengthening.
Findings from a host of studies have shown that HEAs demonstrate exceptional mechanical properties [16], including elevated strength, outstanding ductility [17], dynamic-impact strength [18,19], superior fracture toughness [20], and favorable thermal stability [21]. Until now, investigations [7,22] have predominantly concentrated on the distinctive mechanical behavior exhibited by different HEAs. However, the various mechanical responses of HEAs under severe conditions, such as high strain rates, have garnered growing interest within the domains of materials science and mechanics [23].
Among the existing HEA systems, face-centered-cubic (FCC)-based HEAs have been extensively studied under dynamic loading [24,25]. In the context of FCC-structured HEAs, it has been observed that the stacking-fault energy and the critical stress required for twinning are in昀氀uenced by the intricate chemical environment at the individual level [26]. The occurrence of deformation twins serves to mitigate stress concentrations resulting from dislocation tangling and the accumulation of large numbers of dislocations [27]. Consequently, FCC HEAs often demonstrate an effective combination of elevated strength and adequate ductility under dynamic-loading conditions [28]. Though, the bodycentered-cubic (BCC) structure in RHEAs, which mainly contain refractory elements and exhibit considerable potential for dynamic applications [24,29], hinders twinning deformation due to their high stacking-fault-energy (SFE). It is crucial to note that in BCC materials, the concept of 'high SFE' indirectly re昀氀ects the complex energy barriers to deformation mechanisms like twinning, necessitating a deeper exploration beyond traditional FCC contexts [30].
Furthermore, the reduced thermal conductivity of RHEAs, which is a result of the inherent properties of their constituent atoms, might exacerbate the occurrence of adiabatic shear effects [31]. The reduced thermal conductivity of RHEAs, in turn, poses signi昀椀cant dif昀椀culties and complexities in achieving ductilization under dynamic-loading conditions. Hypothetically, the mitigation of severe deformation localization and the development of ASBs in RHEAs could potentially lead to an increase in their ductility under dynamic-loading conditions [5].
Accordingly, in the present research, room-temperature compression experiments were conducted on a newly designed (TiNbZr) 89 (AlTa) 11 RHEA subjected to dynamic-loading within the range of 1 × 10 3 -to 3.5 × 10 3 s -1 . An integral part of this investigation involves exploring the potential of post-dynamic compression heat treatment to optimize the mechanical properties of RHEAs. This approach aims not only to enhance our knowledge of the fundamental mechanisms in昀氀uencing the balance between strength and ductility in these alloys, but also to develop novel alloy compositions and treatment protocols that can surmount current limitations in ductility while preserving high strength. Furthermore, the goal of the present research is to enhance microstructural properties through exposure to the high strain rates and subsequent heat treatment, thereby contributing to the development of materials designed for performance in harsh environments. Therefore, it is believed that the current research is pivotal in the ongoing advancement of alloy design, offering novel insights and pathways for material innovation in sectors where good mechanical performance is critically demanded.
Ingots of the (TiNbZr) 89 (AlTa) 11 (at. %) RHEA were produced using a vacuum-melting furnace. The constituent metals, sourced commercially, had a purity level exceeding 99.9 %. During arc melting, a mixture of these constituent elements was melted in an atmosphere consisting of argon gas. The speci昀椀c operational procedure for melting involves the amalgamation of the alloy's source components within the crucible. The technique utilizes an arc current ranging from 60 to 70 A. When the arc is initiated, it is initially melted on a crucible holding pure Ti for a duration of 1-2 min. This step serves to eliminate the oxygen present in the furnace, hence preventing oxidation of the HEA during the melting process. Afterwards, the alloy is melted using a smelting current ranging from 200 to 300 A. The alloy samples underwent multiple cycles of melting, with a minimum of 10 repetitions, in order to achieve consistent melting and enhance their chemical uniformity. Following casting, a homogenization treatment was performed at 750 ç C for 3 h and then at 1,050 ç C for 4 h to achieve a consistent microstructure. The selection of these temperatures was guided by a combination of a series of optimization experiments, and the speci昀椀c characteristics of the RHEA [32]. The density of the RHEA was estimated using the Archimedes method [33]. Additionally, the phase structure, microstructure, and chemical composition prior to compression were examined by X-ray diffraction, SEM (Apreo, Thermo Scienti昀椀c, Netherlands), electron backscatter diffraction (EBSD), and an energy-dispersive X-ray spectrometer (EDS), and the results were reported in our previous work [34]. Cylindrical specimens, measuring 4 mm in diameter and 4 mm in length, were prepared for dynamic and quasi-static-compression tests. The quasistatic-compression tests were conducted at room temperature, with the results reported in our previous work [34]. Additionally, dynamiccompression tests with strain rates ranging from 1.0 × 10 3 to 3.5 × 10 3 s -1 were carried out using a split Hopkinson pressure bar (SHPB). The compression tests were performed at least three times for reproducibility. The samples were fractured after compression at a strain rate of 3.5 × 10 3 s -1 . Post-compression microstructure investigation was performed using EBSD techniques. Besides, after dynamic-compression tests, the strains were assessed employing kernel average misorientation (KEM) methods (by the Aztec software) [35]. Further, microstructures were investigated using a transmission electron microscope (TEM; Tecnai G2 F20) equipped with a selected area electron diffraction (SAED) system and operating at 200 kV. The TEM samples were prepared by electro-polishing, as described in previous studies [36]. The examination of the RHEA microstructure by TEM prior to compression is also reported in our previous work [34].
In the 昀椀nal phase of our experiments, the peak yield strength was achieved in the alloy after dynamic compression using a SHPB at a strain rate of 3.0 × 10 3 s -1 . Following this step, a solution treatment was applied to the samples at 1,000 ç C for 30 min, succeeded by water quenching. Subsequently, the alloy underwent ageing at 600 ç C for varying durations: 0 h, 2 h, 4 h, 8 h, 12 h, 16 h, 24 h, 48 h, and 72 h. This heat treatment aimed to determine the maximum hardness for microstructure analysis utilizing an ISO 6507-standard HV-1000Z micro-Vickers hardness tester at 10 kg-force loads for 15 s [37]. The hardness of the alloy was measured to assess the effect of the heat treatment. Twelve indentations were performed and spaced 0.5 cm in which the average hardness value was calculated.
The (TiNbZr) 89 (AlTa) 11 RHEA demonstrates a very low density of 6.0 g/cm 3 , which is due to the notable presence of Ti (density of 4.5 g/ cm 3 ) and Al (density of 2.7 g/cm 3 ) [38,39]. Furthermore, the density of this RHEA is signi昀椀cantly lower than Inconel 718 (8.22 g/cm 3 ), Mar-M247 (7.8 g/cm 3 ), and most other RHEAs with densities exceeding 8.0 g/cm 3 , as reported in an earlier review article [5]. Given the signi昀椀cantly lower density of this RHEA compared to traditional hightemperature alloys, it becomes crucial to conduct mechanical testing at high strain rates to access its suitability for practical applications. Therefore, an examination is conducted on the dynamic-mechanical behavior and microstructural response of the recently developed lightweight RHEA employing mechanical testing at various strain rates. This analysis aims to evaluate the feasibility of using this material in practical applications, such as the mechanical components of automobiles as these components are often subjected to high strain rates due to collisions during driving.
Fig. 1(a) displays the true stress-true strain curves of the (TiNbZr) 89 (AlTa) 11 RHEA for strain rates ranging from 1.0 × 10 3 to 3.5 × 10 3 s -1 . The yield and ultimate compression strengths, as derived from the curves with standard deviations, are summarized in Table 1. It is apparent that the RHEA displays strain-rate sensitivity where strengths are increasing with increasing strain rates from 1.0 × 10 3 to 3.5 × 10 3 s -1 . For example, the dynamic-compressive yield strength and 昀氀ow stress increase with increasing the strain rate up to 3.0 × 10 3 s -1 . The yield strength after compression at 3.0 × 10 3 s -1 reaches a yield strength of 1,450 MPa. The yield strength of the HEA increases by 25 % after compression at 3.0 × 10 3 s -1 by comparison with the HEA after compression at 1.0 × 10 3 s -1 .
The increase in strength may be due to interactions between the moving dislocations and phonons, which are de昀椀ned as elastic vibrations within a crystal lattice or discrete quanta of vibrational mechanical energy [40]. Under strain rates higher than 10 3 s -1 , dislocations encounter a uniform distribution of phonons emanating from the surrounding area, their motion can be impeded due to a viscous drag effect caused by the phonons [41]. Moreover, these swift dislocations create localized temperature changes, leading to an irreversible heat 昀氀ow and energy depletion in the dislocation, thereby reducing its mobility leading to locking of dislocations and increase in strength. Such a viscousdrag effect on the dislocations causes the solid to behave similarly to a Newtonian viscous material during dynamic deformation. The friction force applied to the dislocation by drag effects under dynamic conditions can be expressed as [2]:
where f v is the friction force, B is the viscous coef昀椀cient, v is the dislocation velocity, τ is the shear stress, and b is the Burgers vector. The factor, B, is dependent on the dislocation velocity as follows [2]:
where B o is the viscosity at rest, and C s is the shear-wave speed of materials. Thus, based on Eqs. ( 1) and ( 2), the dynamic-dislocation motions lead to a strong high strain-rate dependence of the 昀氀ow stress [42]. Under high strain rates, the process of dislocation slip in metals is in昀氀uenced by viscous damping, which includes both phonon and electron-drag viscosities. Studies have demonstrated [41,43] that particularly at room temperature, the phonon-drag viscosity is a sig-ni昀椀cant factor in the damping process. Phonons, which are thermoelastic waves moving through the metal lattice, increase in intensity when the lattice is compressed. This intensi昀椀cation of thermal vibrations among the atoms leads to a corresponding increase in the viscosity associated with the phonon drag. The modi昀椀ed Z-A model considering the effect of phonon drag can be expressed as follows [41]:
where σ o is the thermal stress at 0 K, β o and β 1 are material parameters, ϛ is a material parameter related to viscosity, Y is the initial yield strength, T is the temperature, ε is the strain-rate, and σ g is the grain-boundary strengthening stress. This equation is important as it provides a relationship between the initial yield strength and factors, such as the temperature, strain rate, and grain-boundary strengthening stress.
Increasing the compressive strain rate to 3.2 × 10 3 and 3.5 × 10 3 s -1 leads to a drop in the yield strength by 6 %. The drop in the yield strength can be attributed to thermal softening and strain localization [36,44]. Fig. 1(b) illustrates that the current RHEA exhibits an optimal
The yield and ultimate strengths of the RHEA tested at strain rates ranging from 1.0 × 10 3 to 3.5 × 10 3 S -1 .
Strain rate (S -1 ) Yield strength (MPa) Ultimate compression strength (MPa)
1.0 × 10 3 940 ± 12 1,402 ± 10 1.6 × 10 3 1,276 ± 8 1,437 ± 6 2.1 × 10 3 1,095 ± 10 1,445 ± 8 2.5 × 10 3 1,106 ± 10 1,466 ± 5 3.0 × 10 3 1,454 ± 5 1,487 ± 5 3.2 × 10 3 1,329 ± 8 1,485 ± 8 3.5 × 10 3 1,254 ± 6 1,441 ± 10 yield strength under high compressive strain rates, indicating its potential suitability for applications that involve exposure to such conditions. The second-phase particles, which consist of mainly of nano B2 and AlZr rich particles, are dispersed throughout the grains such that they hinder dislocation motion. These phases are homogeneously distributed and were detected on the as-cast RHEA, as presented in our previous study [34].
Fig. 2 presents the hardness pro昀椀le of the RHEA after undergoing the post-dynamic compression heat treatment. According to the hardness chart, the peak hardness of approximately 870 HV was observed after ageing at 600 ç C for 16 h. Consequently, to gain deeper insights into the effects of the heat treatment, the microstructure of the alloy, postdynamic compression, and post-heat treatment at 600 ç C for 16 h, was further examined using both SEM and TEM.
Typical EBSD-orientation maps of the (TiNbZr) 89 (AlTa) 11 RHEA after compression at strain rates of 3.0 × 10 3 s -1 and 3.5 × 10 3 s -1 are shown in Fig. 3. These maps were recorded on the transverse sections perpendicular to the compression direction. Fig. 3(a) reveals that the RHEA compressed at 3.0 × 10 3 s -1 had mostly equiaxed grains and a few 昀椀ne small grains with average sizes of ~ 50 μm (Fig. S1), which were signi昀椀cantly smaller than those found in the as-cast sample (193 μm) investigated in our recent study [34]. Fig. 3(b) shows the sample after compression at a strain rate of 3.5 × 10 3 s -1 , which indicates the presence of equiaxed grains similar to that observed in Fig. 3(a), although they had a smaller average grain size of 24 μm (see Fig. S1). The dominant orientation texture was along {1 0 1} with the presence of {0 0 1} and {1 1 1} after compression at 3.0 × 10 3 s -1 . The dominant orientation owes to the recrystallization that occurred during deformation. It was found that after compression at 3.5 × 10 3 s -1 , the dominant orientation was still along {0 0 1}.
The kernel average misorientation (KAM) of the alloy after compression at strain rates of 3.0 × 10 3 s -1 and 3.5 × 10 3 s -1 are presented in Fig. 3(c)-(d), respectively. The color gradient, ranging from blue to green, indicates the degree of angular misorientation between neighboring grains. Here, blue signi昀椀es areas of low misorientation that is typical of uniform grain structures, while green to yellow hues represent higher misorientation that correspond to regions of internal strain or grain-boundary distortion. In contrast, Fig. 3(c) shows a higher degree of localized orientation. However, as depicted in Fig. 3(d), an increase in strain rate leads to a decrease in the localized orientation, which can be attributed to thermal softening [5]. The regions highlighted with arrows correspond to areas with high internal strains caused by the presence of micro shear bands. It is suggested that shear bands were only found at a very large, imposed shear strain. The work imposed upon the material during deformation can be used to calculate a (adiabatic) temperature rise inside the shear band using the following equation [2]:
where ρ = 6.0 g/cm 3 is the density, τ is the shear stress, and C p is the speci昀椀c heat capacity. By using a weight averaging method, C p = ω i 3 5 i=1 C pi (ω i is the weight percent, and C pi is the speci昀椀c heat capacity for each element of the alloy) [56]. An increasing strain rate raises the imposed shear strain and leads to thermal heating and further fracture after high strain-rate compression. The presence of thermal heating and high-strain-rate compression leads to dynamic recrystallization and recovery after compression, leading to the grain re昀椀nement that is not present in the as-cast RHEA [34]. Fig. 3(c) features a large-strain area in the KAM map whereas for Fig. 3(a), 昀椀ne grains can be observed within these large-strain regions in addition to very 昀椀ne grains within the region suggested to be shear bands. It is important to clarify that shear bands are a particular kind of strain localization that involves strong shear deformation whereas strain localization is a broader term that includes any form of localized deformation in a material. Similar structures have been recently reported in α-Ti and identi昀椀ed as kink bands [57]. Increasing the strain-rate compression to 3.5 × 10 3 s -1 results in thermal softening that leads to dynamic recovery and slight grain growth accompanied by reduced local strains, as con昀椀rmed through Fig. 3 (b) and (d).
Typical TEM micrographs of the RHEA subjected to high-strain-rate compression at 3.0 × 10 3 s -1 and 3.5 × 10 3 s -1 are displayed in Fig. 4 (a)-(b) and Fig. 4(c)-(d), respectively. Inspection of the bright-昀椀eld (BF) micrograph in Fig. 4(a) reveals the presence of micro shear bands at different orientations in which these bands have great dislocation tangles. Weak beam-dark 昀椀eld (WBDF) TEM micrographs are displayed in Fig. 4(b), which features dislocations that are primarily entangled or in loops (white contrast within the shear band). Further analysis shows that the dislocations inside the shear bands consist mainly of screw dislocations with a g vector = [0 1 1] BCC , which is selected such that |g AE A⋅ bp | = 1 where bp is a Burgers vector. Fig. 4(c) indicates that the shear bands with dislocation tangles are located within the shear bands, which is con昀椀rmed by the WBDF TEM image presented in Fig. 4(d). In this case, small dislocation loops and entanglements are heterogeneously distributed. However, as shown in Fig. 4(b), these features decreased with an increase in the strain rate. As mentioned earlier, this reduction is in昀氀uenced by thermal heating and leads to fracture following high-strain-rate compression at 3.5 × 10 3 s -1 . In the early phases of straining, high densities of dislocations emerge and reorganize into cellular substructures as strain increases. This process results in the subdivision of the original grains into smaller areas with signi昀椀cant misorientation [58]. The boundaries around these small areas consist of stored dislocations, which transform into 昀椀ner grains with further straining [58,59].
High-magni昀椀cation TEM images are displayed in Fig. 5. Fine, homogeneous precipitates can be observed in Fig. 5(a). These precipitates predominantly consist of the B2 and AlZr-rich phases, as con昀椀rmed by the SAED pattern along the 81 1 29 BCC direction. Fig. 5(b) illustrates the presence of very 昀椀ne precipitates along the grain boundaries. Following high strain-rate compression at 3.0 × 10 3 s -1 , the average size of these precipitate is ~ 20 nm, which is a 40 % reduction compared to those in the HEA before compression. This reduction is caused by dislocation shear occurring during the high strain-rate compression. Fig. 5(c) and 5 (d) show BF-TEM under the two-beam-condition images of dislocation networks after compression. Inspection of the two images indicates that some dislocations disappeared under a g vector, [1 1 0] BCC , by comparison with dislocations observed under a g vector, [0 1 1] BCC . It is suggested that the disappeared dislocations on the g vector of [1 1 0] BCC is screw dislocations where the g vector is perpendicular to the dislocation lines and/or edge dislocations, g, where the g vector parallel to the dislocation lines [60,61]. This feature shows the evidence of a mixed type of dislocations observed after high-strain-rate compression. Furthermore, sessile dislocations, constituted by two 60 ç full dislocations, are detected along the g vector, [0 1 1] BCC , as indicated by the arrow in Fig. 5(d). Fig. 5. The high-magni昀椀cation BF-TEM images of the RHEA after high-strain-rate compression at 3.0 × 10 3 s -1 showing (a) a homogeneous distribution of 昀椀ne precipitates, (b) 昀椀ne precipitates along the grain boundaries (indicated by white arrows), Two-beam BF-TEM images of dislocations (indicated by red circles) in the deformed alloy imaged at the [1 1 1] BCC zone axis with (c) a g vector along [1 1 0] BCC , and (d) a g vector along [0 1 1] BCC . Arrows inserted in (Figs. c and d) the SAED images stand for the g-vector direction. 3.0 × 10 3 s -1 and subsequent peak aging treatment at 600 ç C for 16 h. Fig. 6(a) displays equiaxed grains where the average grain size in the HEA increased to 130 μm after high-strain-rate compression. Fig. S2 shows the grain-size distribution of the RHEA after dynamic compression at 3.0 × 10 3 s -1 , followed by the peak-aging treatment at 600 ç C for 16 h. The phase map in Fig. 6(b) reveals the presence of agglomerated coarse second-phase particles that are heterogeneously distributed. Fig. 6(c) displays the results of the KAM analysis, which indicates that heat treatment led to a reduction in the internal strains of the sample by comparison with the sample before the heat treatment, as depicted in Fig. 3(c). Furthermore, the micro shear bands (denoted by the white arrows in the 昀椀gure) became markedly smaller after the heat treatment [62].
The TEM micrographs shown in Fig. 7 illustrate the microstructural changes in the RHEA following the peak-aged heat treatment. Specifically, Fig. 7(a) highlights the existence of large, coarse particles intertwined with networks of dislocations, known as dislocation tangles. Conversely, Fig. 7(b) shows plate-like and spherical second-phase particles located within and along grain boundaries, where the volume fraction of precipitates is slightly higher (an average size of 20 nm) after the heat treatment. Furthermore, the insets of Fig. 7(b)-(c) feature a SAED pattern along 81 1 09 BCC indicates that B2 and Zr 5 Al 3 phases [34,63] were present in the sample. Further inspection of the SAED pattern in Fig. 7(c) con昀椀rms the presence of the B2 phase adjacent to the grain boundaries along 81 1 09 BCC , which are coarse particles formed after the heat treatment. The WBDF-TEM image shown in Fig. 7(d) reveals a signi昀椀cant volume of 昀椀ne second-phase particles that serve as pinning sites for dislocations, chosen with a g vector of [0 1 1] BCC to ensure |g⋅ bp | = 1, which are visible in white contrast where the dislocations are identi昀椀ed.
Fig. 8 exhibits an EDS map of the RHEA after dynamic compression at 3.0 × 10 3 s -1 and subsequent peak-ageing heat treatment, which shows precipitates along a grain boundary. Fig. 8(a)-(b) show, respectively, BF-TEM and high-angle annular dark-昀椀eld scanning-TEM (HAADF-STEM) micrographs of the grain-boundary precipitates. Fig. 8 (c) displays the solute variation of these precipitates, as determined by the EDS mapping. It was found that the precipitates predominately consisted of Al 3 Zr 5 . Also, another group of precipitates was detected that contained a percentage of Ta and Nb, suggesting that they were B2 second-phase particles.
From these results, it is hypothesized that during aging, the equilibrium segregation of solutes to grain boundaries hastens the precipitation process, as compared to the bulk material. This hypothesis is substantiated by the presence of coarse particles observed along the grain boundaries after the aging treatment, as evidenced in Fig. 7(c) and 8(c). The presence of even a minimal quantity of this segregated impurity can signi昀椀cantly change both the microscopic dynamics of the defects and the overall behavior of the material [64,65]. The solute segregation isotherms based on the Langmuir adsorption, which was adopted for grain-boundary segregation by McLean [66], is given by:
where X 1 GB is the atomic concentration of the solute I at the grain boundary, X 1 is the atomic concentration of the solute, I, in the matrix, E 1 is the internal energy of segregation of a solute atom, T is the temperature, and R is the universal gas constant. Based on Eq. ( 5), the solute segregation along grain boundaries is increasing with the peak-aged heat treatment, suggesting faster kinetics along the grain boundaries by comparison with the interior grains. Such a scenario supports the increase in the number of second-phase particles. This result was con昀椀rmed by a recent 昀椀nding, which provides direct evidence of the importance of entropy in grain-boundary segregation [67].
As previously stated, subjecting the samples to dynamic compression at a strain rate of 3.0 × 10 3 s -1 resulted in the acceleration of agehardening processes within them. This 昀椀nding was con昀椀rmed through hardness measurements conducted after the samples were subjected to heat treatment at 600 ç C for 16 h. The peak hardness (Fig. 1) can be attributed to the 昀椀ne precipitates coupled with dislocations, as observed in Fig. 6(b) and 7(b). It is proposed that the solute atoms (Ti, Nb, Zr, Al, and Ta) at substitutional sites cause lattice distortion, which can be mitigated to some extent when these atoms cluster around line defects and form what is known as a Cottrell atmosphere [68]. In this alloy, larger solute atoms tend to gather in the expanded areas of dislocations, especially near the dislocation core beneath the dislocation line. In contrast, smaller solute atoms are more likely to move to the compressed areas above the slip plane. Such migration leads to a concentration of solute atoms along the dislocation lines, forming solute-rich areas [69]. This movement towards the dislocation can occur through two primary mechanisms. First, the stress 昀椀eld generated by the dislocations can facilitate the migration of solute atoms. Second, the formation of Cottrell atmospheres is temperature-dependent: the temperature must be suf昀椀ciently high to enable defect migration, but not so high that it causes the atmosphere to disperse back into the solution matrix due to the impact of entropy on the free energy [70]. During the solution treatment, solutes tend to accumulate along dislocations. Subsequently, during the aging treatment, these solutes evolve into precipitates. Notably, after the heat treatment at 600 ç C for 16 h, a multitude of 昀椀ne, nanosized precipitates are formed. This result is clearly illustrated in Fig. 7(e), where these particles are homogeneously distributed within the grains. In summary, given their unique properties and the ongoing advancements in our understanding and production capabilities, HEAs hold promise in revolutionizing the performance and durability of mechanical components in the automotive industry.
The present research shows a comprehensive analysis of the (TiNbZr) 89 (AlTa) 11 RHEA, demonstrating its unique dynamicdeformation behavior and offering insights into its microstructural evolution under various strain rates. The key 昀椀ndings are described as follows:
1. The (TiNbZr) 89 (AlTa) 11 RHEA exhibits pronounced strain-rate sensitivity, with notable increases in the yield and ultimate compression strengths at strain rates up to 3.0 × 10 3 s -1 . This enhancement in mechanical properties is attributed to the intrinsic resistance of the material to deformation and the effective activation of dislocation mechanisms. Besides, the formation of the B2 phase and Zr 5 Al 3 precipitates was found to be crucial in boosting the alloy strengths at high strain rates. However, at strain rates beyond 3.0 × 10 3 s -1 , a decline in yield strength is observed, likely due to thermalsoftening and strain-localization phenomena. This behavior underscores the complex response of the alloy to dynamic stress, providing valuable insights for applications where variable strain rates are encountered. 2. Under dynamic compression, the RHEA undergoes signi昀椀cant microstructural changes, including grain re昀椀nement, formation of micro shear bands, and development of dislocation tangles. These transformations are indicative of dynamic recrystallization and recovery processes, which play a pivotal role in enhancing the strength and ductility of the material. The observed microstructural adaptability is crucial for the application of this alloy in environments subjected to extreme dynamic stresses. 3. The application of the post-dynamic compression heat treatment, particularly aging at 600 ç C for 16 h, further re昀椀nes the microstructure of the RHEA. This treatment leads to the development of second-phase particles and an increase in hardness, contributing signi昀椀cantly to the overall mechanical properties of the material. The effectiveness of this post-deformation heat treatment suggests a promising approach for tailoring the microstructure and enhancing the performance of RHEAs in practical applications.
Materials & Design 243 (2024) 113062