key: cord-1006778-p1kl9y2w
authors: Sabban, Rushikesh; Dash, K.; Suwas, S.; Murty, B. S.
title: Strength–Ductility Synergy in High Entropy Alloys by Tuning the Thermo-Mechanical Process Parameters: A Comprehensive Review
date: 2022-03-24
journal: J Indian Inst Sci
DOI: 10.1007/s41745-022-00299-9
sha: 51505f574084fbe1341adfa3ff912e41506118cd
doc_id: 1006778
cord_uid: p1kl9y2w

The strength–ductility trade-off is an eminent factor in deciding the mechanical performance of a material with regard to specific applications. The strength–ductility synergy is generally inadequate in as-synthesized high entropy alloys (HEAs); however, it can be tailored owing to its tunable microstructure and phase stability. Thermo-mechanical processing (TMP) allows the microstructure to be tailored to achieve desired strength–ductility combination. The additional attribute is evolution of texture, which also significantly influences the mechanical properties. This review presents a critical insight into the role of TMP to achieve superior strength–ductility symbiosis at room temperature in single-phase (FCC, BCC) and multiphase HEA. The role of overall processing strategy of HEAs encompassing rolling and subsequent annealing in relation to the evolution of microstructure and texture in have been discussed. Recently practiced severe plastic deformation processes have also shown promise in improving the strength–ductility combination. The relevance of these processes in the processing of HEAs has also been analysed. At the end, futuristic approaches have been elaborated to enable efficient as well as hassle-free process towards achieving the proficiency of strength–ductility in HEAs.

strength of many HEA, in as-synthesized condition, restricts their use in structural applications such as automobile sector 16 . Hence, enhancing the strength while maintaining ductility of HEA is a crucial factor to become potential structural material to compete with conventionally used materials in automobile and other sectors. The HEA are synthesized conventionally through various routes such as melting and casting 17 , powder metallurgy by mechanical alloying (MA) and sintering 18-21 , surface deposition 22 , etc. Techniques like additive manufacturing (AM) 23 and carbothermal shock synthesis (CTS) 24 have also been utilized widely in recent times. The HEA manufactured through these routes have inherent limitations including gas porosity, 1 3 J. Indian Inst. Sci. | VOL xxx:x | xxx-xxx 2022 | journal.iisc.ernet.in shrinkage porosity, un-melted particles and detrimental microstructural features, etc. [25] [26] [27] . These defects in the specimens create inhomogeneous stress distribution which limit the optimum strength-ductility combination. The strengthductility synergy is quantified as a product of strength and ductility (PSD in GPa × %) defined elsewhere 28 . Since the strength and ductility specifications for many room-temperature structural applications are critical, the synthesized HEA need to be engineered further to imbibe the desirable properties.

Thermo-mechanical processing (TMP) constitutes of a series of plastic deformation and thermal operations to tune the microstructure consequentially enhancing the strength-ductility synergy of as-synthesized components. The plastic deformation processes, like rolling, combined with a post-deformation annealing constitute the most commonly used sequence of a TMP schedule. A combination of rolling and heat treatment process (RHP) has been effectively utilized for aluminium, steel, and titanium alloys for enhancing room-temperature mechanical properties 29-31 . The rolling process, apart from changing the shape, is known to strengthen the as-cast or PM processed materials through various mechanisms such as dislocation strengthening, grain and twin boundary strengthening 32,33 , etc. A subsequently optimized heat treatment facilitates the tailoring of the strength-ductility combination 34 , leading to mechanical properties superior to ascast or PM processed specimens.

The schematic Fig. 1 depicts three major modes of HEA processing along with key factors in individual stages. First step of HEA synthesis routes render characteristic microstructural features, e.g., type and fraction of phases, grain size, etc. All these parameters play a vital role in deciding post RHP properties 35, 36 . During secondary processing, i.e., deformation through rolling or HPT, the working temperature along with recrystallization temperature governs the microstructural evolution. Rolling at various temperatures exhibits different strengthening mechanisms and hence influences the properties of HEA 37-39 . The rolling processes can be differentiated based on operating temperature which includes cryo-rolling (CR), room-temperature rolling (RTR), hot rolling above the recrystallization temperature (HR), and warm rolling (WR) (around average temperature of hot rolling and RTR). In addition to temperature, the parameters like amount of total strain, strain path, symmetricity of rolling (symmetric/ asymmetric rolling), etc. also play an important role in ascertaining the final properties of an HEA [40] [41] [42] . During the heat treatment of the HEA, the annealing temperature primarily dictates microstructural evolution 43 . Additionally, the duration of annealing 44 and heating rate to achieve temperature of interest 45 also can be used to tune the strength-ductility synergy. Taken together, there are various parameters associated with different stages of TMP which decides the enhancement in tensile room-temperature mechanical properties of HEA.

There are a substantial number of journal papers on RHP of HEA starting from the year 2009. Figure 2 highlights the rapid increasing trends in number of publications in TMP starting from its inception. Review on evaluation of mechanical properties of HEA through tensile and compressive behaviour has also been undertaken 46, 47 . However, any of the reviews does not deal with the generic aspects of thermo-mechanical processing involving the rolling and annealing process. In summary, the focus of this review is to establish the role of variety of parameters in RHP process.

The rolling process in TMP has been effectively employed in the strengthening of as-synthesized HEA [48] [49] [50] [51] [52] [53] . Figure 3 clearly indicates significant strengthening of HEA due to rolling. The compositions of HEA, rolling parameters, and percentage increase in the strength values after rolling have been given in Table 1 . The primary strengthening mechanisms in HEA operative during rolling are associated with increased dislocation density 54 , twin boundaries 37 , kink bands 40 , shear bands 55 , grain refinement 56 , transformation-induced plasticity 57,58 , back stress hardening due to accumulation of dislocations 59-61 , etc. In the following section, the tuning of various factors such as pre-rolling microstructural features, temperature of rolling, total strain, strain path, symmetricity, etc. and their corresponding strengthening mechanisms will be discussed.

Based on Phases The various types of as-synthesized alloys with different phases can be classified as: (i) singlephase alloys: face-centered cubic (FCC) 50 , body centered cubic (BCC) 40 , B2 phase 62 , etc.; (ii) multiphase alloys: FCC + BCC 63 , FCC + hexagonal closed packed (HCP) 64 , L1 2 + B2 65 , etc. The deformation mechanisms active in these two types of alloys leading to strengthening will be elaborated in this section.

During the initial stages, the single-phase FCC HEA undergo strengthening through dislocation activities leading to cell formation 66 -68 , deformation twins 69 , microbands 70,71 , shear bands due to heavy deformation 72,73 , etc. At constant deformation temperature, the deformation of FCC HEA is primarily controlled by stacking fault energy (SFE) 74 . The nucleation of deformation twins is directly proportional to SFE 75,76 . Hence, the HEA with lower SFE exhibit deformation twinning as their deformation mechanism 77 . Additionally, with reduction in the SFE by tuning of composition of alloy, the deformation twin thickness can also be reduced 76 . This reduction of twin thickness will provide higher strengthening in FCC HEA 78 . In contrast, the increase in SFE (lowering the distance between Shockley partials) leads to easier cross-slip. Hence, at higher SFE values, microband formation was found to be dominant mechanism for deformation 79 . The deformation mechanism changes from twinning to microband formation with increasing nitrogen content 70 . The SFE increases by alloying elements like nitrogen leading to change in the deformation mechanism 80 . To sum up, in the initial and intermediate stages of deformation, the microstructural evolution and strengthening in FCC HEA is influenced by SFE. J. Indian Inst. Sci. | VOL xxx:x | xxx-xxx 2022 | journal.iisc.ernet.in Table 1 : Strengthening of as-synthesized high entropy alloys after rolling. In BCC HEA, the deformation is driven by mechanisms such as increase in dislocation density 40 , kink bands 81 , deformation bands 82,83 , microbands 84 , shear bands 40 , etc. The misorientation of boundaries was found to be related to kink bands 40,81,85 after rolling which was lower compared to the misorientation of twin boundary in BCC HEA 40,81,86 . Along with BCC alloys, the partially ordered single-phase B2 also deforms primarily by dislocation microbands formation 62 . The formation of microbands in B2 has been attributed to the preference of deformation via planar slip over other mechanisms 71 , wherein the leading Shockley partially destroys the barrier for ordering of B2 phase making it easy for trailing Shockley partial to continue planar slip leading to formation of microbands 62,87 . At larger plastic strains, the shear band formation is the dominant mechanism for most of the HEAs irrespective of initial phase 40,72 . The shear band formation represents plastic instability 88 and occurs irrespective of crystallography of sample at sufficiently larger rolling deformation 89 . Taken together, there are variations in deformation mechanisms possible in single phases based on type of phase and SFE.

The deformation mechanisms of multiphase alloys include most of the mechanisms applicable to single-phase alloys, such as dislocation cell formation 43,60 , deformation twinning 73,90 , shear banding 65,73 , etc. In addition, the multiphase alloys have additional strengthening mechanisms such as back stress strengthening due to accumulation of dislocations at phase boundaries 59-61 , strain-induced martensitic transformation 73,90 , etc. The major strain partitions to soft FCC phase compared to hard phases like carbides, ordered phase B2, etc. For maintaining continuity, the geometrically necessary dislocations (GNDs) participate in the deformation, which provides additional strengthening in dual-phase alloys compared to single-phase alloys 59 . The mechanism of strain-induced FCC-to-HCP martensite transformation in the dual-phase alloy containing FCC and thermally stabilized HCP phase has also been reported 64 . The stacking fault generated in the initial stages of deformation in FCC phase acts as nuclei for transformation of FCC phase to HCP martensite 91 . Hence, in general, the strengthening effect is higher in dual-phase HEA compared to single-phase HEA.

As described in previous section, the active strengthening mechanisms in HEA are dislocation activity 50 , microbands 71 , stacking fault (SF) 92 , deformation twinning 70 , shear bands 93 , etc. depending on pre-rolling phases and SFE. The selection and evolution of any of the abovementioned mechanisms, however, depends on the rolling temperature 37 . During cryo-rolling (CR) and room-temperature rolling (RTR), the evolution of microstructure is governed by dislocation activity and twinning depending on SFE of HEA. The SFE of the conventional alloys and HEA is directly proportional to temperature 92,94 . Hence, during CR and RTR, the microstructural evolution will change according to SFE. However, at higher temperature regime (near or above the recrystallization temperature of alloy), the microstructural changes will get affected by thermal activation. The effect of rolling temperature at elevated temperatures on properties can be explained through change in Zener-Hollomon (Z-H) parameter 95 . At constant strain rate, lowering the rolling temperature increases the Z-H parameter and the strength thereof. In this section, the difference in mechanisms of strengthening and their effect on properties after rolling at different temperatures will be discussed.

The difference in the mechanical properties with change in the rolling temperature in the temperature regime spanning from room temperature to cryogenic temperature is shown in Fig. 4 . Figure 4a describes the difference in hardening at two rolling temperatures of CoCrFeNiMo 0.15 HEA 71 and Fig. 4b highlights the effect of rolling temperature on strengthening in CoCrFeNiMn alloy 37 . The cryo-rolled (CR) specimens exhibited more strengthening than room-temperature rolled (RTR) specimens. The difference between RTR and CR is that the microstructural evolution kinetics is faster in the latter 37,71,92,96 . The initial stages of deformation in RTR are dominated by dislocation activity and SFs, and with increasing strain, the deformation twins become dominant. The SFE of the studied alloy is low, and hence, deformation twining is favorable over microbands at room temperature 75 . Compared to RTR specimens the activation of deformation twinning for CR specimens happened at earlier stages. This is explained with the fact that the SFE at liquid nitrogen temperature is further less compared to SFE at room temperature 92 . With increasing strain in CR, the multiple twin systems get activated and distortion of nanotwins occurs later. The shear bands were also found post CR after imparting high strain 37,71,96 . The enhanced dislocation density was also observed in the conventional FCC alloys after deforming at cryo-temperature in comparison with room temperature 97 . The deformation-driven FCC-to-HCP transformation-induced plasticity (TRIP) was observed after CR. The SFs present in the initial stages acted as nucleating sites for FCCto-HCP transformation 64 . The texture after CR and RTR was brass-type texture which is typical for low SFE material 93 . Hence, there is little role of texture in additional strengthening after CR than RTR. Hence, the HEA exhibits the enhanced strengthening at cryo-temperature compared to room temperature (Fig. 4a , b) due to higher dislocation density, intersection of twins from nonparallel systems, TRIP, and more shear banding.

The influence of rolling in high-temperature regime at intermediate as well as at high temperatures, namely warm rolling (WR) 98 and hot rolling 99-102 , has been investigated for HEAs. In general, hot working refers to rolling at temperatures above the recrystallization temperature 103 , while warm rolling is performed at intermediate temperatures of cold and hot rolling 55 . For the HEAs that are less workable, hot working is preferred 35 . Another associated attribute of deformation at high temperature is associate control of microstructures 55,100 . The hot rolling of HEAs involves dynamic recovery, dynamic recrystallization, grain growth, etc. as other micro-mechanisms like phase evolution and transformation depending on rolling temperature 100,104 . Mostly recovery and partial recrystallization is prevalent in the FCC HEA rolled at temperatures below recrystallization temperature of HEA 100 . This recovery-to-recrystallization ratio is strongly dependent on temperature and SFE of the material. Recovery is dominant at lower rolling temperatures and in high SFE alloys, while at higher rolling temperatures and for low SFE alloys, recrystallization is more prevalent 105 . For HEAs, increased degree of recrystallization is observed at higher rolling temperatures 100 . For much higher rolling temperatures, grain growth has also been observed 100 due to increased mobility of grain boundary at higher temperature 105 . The microstructural transformation through any of these routes led to specific strength-ductility synergy in the HEA. The deformation mechanisms during the warm rolling of AlCoCrFeNi 2.1 eutectic HEA were found to be function of rolling temperature 55 . The shear banding and disordering of L1 2 phase was influenced by rolling temperature. In the dual-phase (FCC + BCC) Al 0.5 CoCrFeMn HEA, the FCC phase underwent deformation, while the harder BCC phase underwent grain fragmentation instead of deformation 35 . Another influence of temperature was noticed in terms of propensity of twinning. Twinning was reported to be suppressed with increase in temperature in warm working regime due to increase in SFE in the HEAs that exhibit twinning induced plasticity (TWIP) 106 . In general, the strength-ductility combination has been tailored by a combination of deformation in multiple temperature domains, for example, a combination of cryo-rolling and warm rolling. This combination results in generation of heterogeneous microstructure to enhance the strength-ductility synergy 53 . Such a processing is sometimes referred to as hybrid processing or hybrid rolling. To summarize, a broad range of micro-mechanisms act to play, depending on the working temperature. Hence, by varying the rolling temperature, hence by tailoring the microstructure optimally, the strength-ductility combination can be fine-tuned.

The evolution of microstructure in various HEA with increasing rolling reductions (plastic strains) and its effect on mechanical properties thereof has been evaluated extensively 107-109 . The evolution of strength after cold rolling in FCC HEA Al 0.25 CoCrFeNi with increasing total plastic strain Fig. 5 48 . The increase in % thickness reduction led to increase in strength at the cost of ductility (Fig. 5a) . The strain hardening analysis has been carried out on results presented in Fig. 5a . The strain hardening regime of stress-strain curve (between yield strength and ultimate tensile strength) was considered for the calculations of specimens rolled to different reductions. The Hollomon analysis plot (ln (true uniform plastic stress) vs. ln (true uniform plastic strain)) after various rolling reductions is presented in Fig. 5b . The details of the Hollomon calculation can be found elsewhere 110 . The slope of the curve after applying linear regression (R 2 > 0.8) is the strain hardening exponent (n). Figure 5b highlights the decrease in 'n' (early necking) after increasing the rolling reduction. Similar behavior was observed for some HEA 111,112 and conventional aluminium alloys 110 . The decrease in strain hardening capacity in Al 0.25 CoCrFeNi HEA with increasing rolling reduction can be explained based on deformation mechanisms. The mechanisms of deformation in various phases are different at different temperatures, and dislocation activity and cell formation are the dominant mechanism of deformation in the initial stages 43,60 . As the deformation proceeds, the multiplication of dislocations and sub-grain size refinement take place which makes it difficult for further strain hardening and nonuniform deformation leads to necking 110 . Taken together, the strength evolution as a function of plastic strain has been observed and increased strengthening can be achieved with increasing plastic strain.

The effect of varying strain path during rolling on microstructure evolution and final properties in the conventional alloys containing different phases has been well studied [113] [114] [115] . The elongated grain structure is developed by virtue of unidirectional rolling (UDR), whereas the lamellar structure fragmentation occurs by cross rolling. There are a few reports on difference in microstructure due to strain path change in HEA 41,117,118 . The variation in deformation microstructure gives rise to difference in the mechanical properties post-annealing. In the case of the HEA AlCoCrFeNi 2.1 , the UDR sample resulted in heterogeneous microstructure with lamellar and recrystallized grains in contrast to cross-rolled specimens which exhibited duplex recrystallized grains. Subgrains are well developed in UDR specimen, whereas it gets distorted during cross rolling. The destabilization of dislocation structure during cross rolling is also reported in the case of cold rolling of CoCrFeMnNi FCC alloy 117 . This destabilization effect has also been reported in the conventional FCC alloys 114, 119 , and the distortion in development of misorientation provides lesser nucleating sites for recrystallization leading to coarse grain size in cross-rolled specimens in comparison to UDR 117 . The difference in the grain size affects the strength-ductility synergy 120 . The increase in the volume fraction of intersecting twins in low SFE HEA after cross rolling and intermittent annealing due to destabilization of substructure has also been reported 118 . Therefore, it is clear that changing the strain path during rolling could enable the microstructure and strength-ductility combination to be tuned.

The symmetricity of rolling can be varied by varying the (a) diameter of the rolls, (b) friction conditions at the roll and sample surface, and (c) the speed of the roll 121, 122 . Along with plain strain compressive stress (as in conventional rolling), asymmetric rolling imparts additional shear stress during deformation 122 . The equivalent strains in asymmetric rolling were also found to be slightly higher compared to the conventional rolling 123 . Researchers have utilized asymmetric rolling for refining the microstructure and enhancing the mechanical properties of conventional aluminium, magnesium alloys [123] [124] [125] , etc. The process of asymmetric rolling (ASR) at room temperature using different roll speeds (speed ratio 1.5) for FCC CoCrFeMnNi alloy was also performed 42 . Asymmetrical rolling strengthened the HEA by 96% more in comparison with symmetric rolling (SR) (Fig. 6) . As discussed earlier, the deformation mechanisms dominant in FCC HEA are dislocation cell formation, deformation twinning, shear bands, etc. 43,65,90 . The CoCrF-eMnNi alloy exhibited a higher number of dislocation cell formation in ASR compared with SR specimens. The asymmetrically rolled FCC metals like aluminium exhibited higher percentage of low-angle grain boundaries (LAGBs) converted to high-angle grain boundaries (HAGBs) due to higher dislocation activity. This conversion was promoted by additional shear stress 123 . In addition to dislocation activity, SR specimens showed parallel set of deformation twins in contrast to intersecting twins in ASR specimens. These intersecting twins exhibit additional hardening in ASR specimens 126 . Higher volume fraction of shear bands formation was seen in ASR as compared to conventionally rolled specimens. The complex strain distribution in ASR is reported to be the plausible reason for higher fraction of intersecting twins and shear bands 124, 126 . Owing to the above-mentioned discussion, ASR displays additional strengthening in comparison with SR specimens. Similar kind of additional strengthening in FCC HEA has been reported with different roll speeds 127 . The dislocation density is found to be twice in ASR specimens compared to SR. Asymmetric rolling is proven to be effective way to attain the additional strengthening compared to symmetric rolling for HEA.

The generation of texture during TMP of HEA significantly influences the strength of the material. Orientation of grains in HEA post TMP is decided by the strain path, working temperature, and the recrystallization parameters. Texture evolution in HEA post-deformation in HEA is discussed here based on phases present in the alloy. Table. 2 shows the deformation and annealing texture components of common HEA.

HEA like Cantor alloy contain typical rolling components post-room-temperature rolling (RTR) such as Bs, Cu, cube, and S components. The S component strengthens up to 80% RTR reduction, beyond which it decreases 129 . The brass component strengthens beyond 80% and keeps on increasing till 90% reduction. Slip planes at low RTR reduction and fine lamellae with deformation bands coexisting at high RTR reduction are responsible for the texture development. The Goss and Bs components show up in the FCC phase of FeCrCuMnNi in 90% RTR alloy 109 . Twins present in the Cantor alloy promote the transition of Cu to brass type of texture to during RTR 126 . Al 0.5 CoCrFeNi HEA shows {110} < 112 > and {111} < 110 > components on RTR possessing FCC with trace BCC phase, which on recrystallization become weak 43 . Cryo-rolling of Cantor alloy shows a similar texture as after cold rolling 93 . Multistage cross cold rolling of Cantor alloy shows stronger Bs component than unidirectional rolled specimen 117 .

The L1 2 phase in the eutectic HEA AlFeCrCoNi 2.1 possesses Bs type of texture post-warm rolling along with α-fibre (Goss, Bs, G/B); whereas the B2 phase shows {112} < 110 > type of texture along with RD||110 and ND||111 fibres post-warm rolling 130 . Different strain paths imparted at cryotemperature on EHEA generate Bs along with Goss, Rt-Goss in L1 2 phase, and {001} < 110 > in B2 phase after multistep cross rolling 41 . HfZ-rTiTaNb HEA with BCC phase, when cold rolled, exhibits strong ND fibre (ND// < 111 >), and RD fibre (RD// < 110 >) along with cube and Rt-cube components 131 .

Rolling leads to strengthening of HEAs, compromising on the ductility factor 132, 133 . The strengthductility combination has been optimized by designing different heat treatment regimens post-rolling, which is plotted and highlighted in Fig. 7 128, 134, 135 . The composition of the alloys with respective RHP and properties are presented in Table 3 . The quantification of strength-ductility synergy which is realised by the parameter PSD (GPa × %) displays that RHP has significantly enhanced the strength-ductility combination of as-synthesized HEA (Table 3) . This improvement is dependent on the cascade of events taking place (microstructural evolution) during annealing heat treatments, which depends on various parameters such as annealing temperature 56 , annealing time 136 , heating rate during annealing 45 , etc. In this section, the role of various annealing parameters in microstructure evolution will be discussed in detail.

The different mechanisms of microstructure evolution during heat treatments reported in literature are as follows: recovery 126, 137 , precipitation 138, 139 , recrystallization 140 -142 , grain growth [143] [144] [145] , and annealing twins 146, 147 , etc. The mechanism in play triggering the microstructural change for HEA is dictated by annealing temperature 148 . The strong dependence on annealing temperature can be attributed to microstructural changes occurring, which in turn is strongly influenced by diffusion 149 which is exponentially related (Arrhenius dependence) to temperature 150 . Before annealing, rolling enhances the yield strength of FCC HEA manifold with considerable reduction in ductility 148 . Hence, the product of strength and ductility value reduced marginally, as shown in Fig. 8a 148 . However, annealing at different temperatures increases the PSD parameter significantly in comparison to both pre-and post-rolled HEA. Annealing renders higher yield strength values compared to pre-rolled status of HEA (Fig. 8a) . The plausible reasons behind the enhanced mechanical properties post-annealing at different temperatures will be discussed vis-à-vis microstructural changes.

The driving force for recovery during annealing post-deformation (static recovery) is the stored energy in the rolled specimens 151 . The static recovery in HEA involves steps such as dislocation interaction leading to formation of dislocation cells 68 , reduction of dislocation density inside cell [152] [153] [154] , intensification of texture 126 , formation of subgrains 81 , etc. The recovery temperature regime is dependent on compositional complexity and SFE of the material 155 , etc. The higher is the SFE, more is the recovery as discussed earlier in hot rolling section. Static recovery is found to be dominant for FCC 72,98,137 , BCC 81 , and dualphase 153, 156 HEA after annealing below recrystallization temperature. The activation energy required for static recovery in conventional alloys is same as that of dislocation annihilation by climb and cross-slip, which is lower compared to that of recrystallization 157, 158 . In terms of mechanical property evolution during recovery, the decrease in hardness of rolled FCC HEA is insignificant 72,148,156 , whereas increase in ductility is notable 137 . The strength decrease during recovery was reported to be logarithmic and not as drastic as recrystallization 159 . This is the probable reason for increase in PSD after annealing below recrystallization temperatures of HEA (Fig. 8a ).

The mechanism of recrystallization involves the migration of HAGBs 105 . Hence, recrystallization needs higher thermal energy compared to recovery 157, 158 . Above the recrystallization temperature, the recrystallization starts to dominate recovery. Similarly, the recrystallized fraction increased with increase in temperature in different HEA having various phases 62,72,81,160 possibly due to increase in HAGBs' mobility 161 . The HAGBs' mobility impeding elements such as carbon increase the recrystallization onset temperature 162 . The recrystallization activation energy in FCC HEA (549 kJ/mol) 163 is significantly higher compared to high manganese steel (230 kJ/ mol) 164 and TWIP steel (229 kJ/mol) 165 due to precipitates pinning the grain boundaries in addition to solute drag effect. Hence, the nucleation of recrystallization in HEA requires higher annealing temperature compared to the conventional high-performance alloys. The recrystallization nucleation sources in HEA are deformation bands 52 , grain boundaries 43 , shear bands 98 , second-phase particles 160 , etc. The regions such as deformation bands, shear bands, etc. exhibited recrystallization initialization due to large driving force of stored energy 52,98 . The particle stimulated nucleation (PSN) is also prevalent in multiphase HEA. The harder phase deform less compared to softer phase and dislocation pile-up at the phase interface causes formation of deformation zone leading to nucleation of recrystallized grains 160 . The ductility is enhanced after recrystallization, and hence, the PSD value also gets enhanced with annealing for various HEA 140,166-169 which is shown in Fig. 8a 148 .

As the HEA gets fully recrystallized, the grain growth starts dominating significantly at higher annealing temperatures compared to recrystallization temperature 152 . The excessive grain growth can deteriorate the strength of HEA 162 . The strategy of controlled precipitation in single-phase HEA is effectively employed in HEA to inhibit the excessive grain growth significantly 166 167 , and volume fraction of precipitates 139, 172 . With increase in the annealing temperature, the coarsening of precipitates occurs by Oswald ripening mechanism 166 . Along with coarsening of precipitates, the volume fraction of precipitates reduces with increase in the annealing temperature 172 . Hence, annealing temperature plays a vital role in optimizing the precipitation and controlling grain growth thereof in singlephase HEA. Similar effect of grain boundary pinning is observed in multiphase alloys where harder phase inhibits the grain growth 35, 166, 173 . The volume fraction of harder second phase decreases with increase in annealing temperature. Hence, grain growth dominates with increasing annealing temperature 35, 173 . In addition to this, the annealing twin fraction varies proportionally to grain size 143, 146 , and hence, annealing twinning increases as annealing temperature rises 142, 173 . These microstructural changes result in higher activation energy for grain growth in HEA compared to the conventional alloys 163 . The yield strength and PSD are also higher than pre-rolling HEA even after annealing at temperatures significantly higher than recrystallization temperature (Fig. 8a) . The excessive grain growth has been successfully inhibited with precipitates in singlephase HEA or with harder phase in multiphase alloys and the strength ductility synergy has been enhanced compared with pre-rolling HEA.

After the microstructural evolution mechanism is ascertained by the annealing temperature, the microstructure can be tailored by varying the annealing time to enhance the PSD further 149, 172, 174, 175 . The increasing trend in strength-ductility combination with increasing annealing time with respect to pre-rolled specimens was observed in Al 0.5 CoCrCuFeNi HEA 176, 177 (Fig. 8b) . The specimens heat treated for different times exhibit higher yield strength values in comparison with pre-rolled HEA (Fig. 8b) . The enhancement in the properties in relation to change in the microstructural features will be elaborated here. The recrystallization kinetics in HEA is formulated with Johnson-Mehl-Avrami-Kolmogorov (JMAK) equation 160 . The recrystallization kinetics in HEA is slower 90 compared to other conventional high-performance alloys such as high manganese steel 164 , TWIP steel 178 , microalloyed steel 179 etc. The sluggish diffusion, severe lattice distortion effect and precipitates in HEA, results in delaying the recrystallization 6,90,163 . The sluggish diffusion renders restrictions on HAGBs' movement for recrystallization 90 . The severe lattice distortion leads to generation of local concentration fluctuation (LCF) regions which restricts the dislocation motion during the softening 90 . The precipitates formed during annealing in FCC HEA impede the mobility of HAGBs, and hence, the recrystallization requires higher thermal energy to occur 163 .

The evolution of precipitates in HEA with increase in annealing time at particular temperature has been studied extensively 134, 175 . The precipitate growth exponent (n) reported in various J. Indian Inst. Sci. | VOL xxx:x | xxx-xxx 2022 | journal.iisc.ernet.in HEA is 3 148, 169, 180 . The significance of n = 3 lies in the fact that the coarsening of precipitates is via Oswald ripening mechanism 166 and volume diffusion 82, 148 . The volume fraction of precipitates remains constant up to a specific annealing time and reduces drastically with further increase in the annealing time 172 . The grain boundary precipitates are stable, but the precipitates inside the grain start to dissolve in the matrix leading to overall decrease in volume fraction of precipitates 169 .

The precipitates also slow down the kinetics of grain growth in FCC 136, 169 , BCC 81,145 , and multiphase HEA 166, 180 . The grain growth is described with modified Zener-Smith model (Z-S) model describing the pinning pressure exerted by the precipitates to restrict the grain growth 169 . Abnormal grain growth is also observed in FCC HEA due to heterogeneous distribution of precipitates in the initial stages of annealing 181 . Hence, the discussed microstructural evolution leads to achieving higher PSD values compared to pre-rolled HEA.

The heating rate to reach annealing temperature also decides the microstructural evolution during annealing 182 and mechanical properties of conventional alloys 183 . The difference in grain size after annealing due to variation of heating rate in FCC HEA is represented in Fig. 9 45 . At high heating rates, HEA specimens exhibited lower grain size at two different annealing times than a lower heating rate (0.013 °C/s) of HEA specimens (Fig. 9) . In lower heating rate specimens, the early nucleation and growth of recrystallized grains at preferred sites takes place 45, 184 . Hence, larger grain size is achieved in low heating rate compared to high heating rate. This difference in final grain size leads to difference in the yield strength according to Hall-Petch relationship 162 . Hence, high heating rate achieves better yield strength and PSD compared to low heating rate 183 .

The annealing of ternary medium entropy alloy leads to a strong cube recrystallization texture, similar to high SFE alloys. Quaternary and quinary medium entropy alloys retain some deformed texture components indicating delayed recovery phase 128 . Sluggish diffusion of grain boundaries in quaternary and quinary alloys hinders preferential texture growth, leading to randomisation of the texture. First-order annealing twins give rise to new annealing texture components such as K and M component apart from retained CR components 129 . No significant change in texture was observed with the annealing temperature. Bs, Goss, and S texture components appear post-cold rolling-annealing at 800 °C in CoCrFeMnNi (FCC) alloy with 1 atom % of carbon content elevating the strength 185 . Annealing twinning in Cantor alloy aids strong texture modification post-recrystallisation 126 . Randomisation of texture post-annealing was corroborated with Cellular Automata simulation results. Another study on annealing texture analysis of Cantor alloy showed S component dominant compared to brass and Goss components, and became stronger with annealing temperature 142 .

Annealing of EHEA post-warm rolling at 800, 1000, and 1200 °C renders the FCC phase with retained deformed texture components. {112} < 110 > components are present post-annealing at 800 °C in B2 phase, and {111} < 110 > component shows up at 1200 °C 130 . {111} < 110 > component predominates along with ND fibre post-annealing of the HfZrTiTaNb HEA 131 .

HEA have been processed and engineered using high-pressure torsion (HPT) to improve the strength, hardness, ductility, and superplasticity. The studies performed on HEA, out of which some significant cases will be discussed here. HPT combined with thermal annealing imparts 400% increase in hardness to Al 0.3 CoCrFeNi HEA. Formation of ordered BCC phase at high-temperature (500-700 °C) annealing as well as heterostructure promotes the elevation in hardness 186 . A grain size of 25 nm was achieved by HPT performed on AlNbTiVZr 0.5 alloy which had an initial coarse-grained structure with B2 matrix embedded with C14 Laves phase (rich in aluminium and Zirconium). Increase in nanohardness (550-665 HV) was observed in the B2 phase, whereas the C14 Laves phase becomes softer post-HPT 187 . Betterment in hardness in HPT processed HfNbTiZr BCC alloy from 2600 to 4450 MPa was realised by the aid of friction stress, possessing dislocation density of the order of 10 16 m −2188 . Chromium oxide precipitates of size 7-10 nm in a matrix of CoCrFeMnNi alloy consisting of FCC + BCC solid solutions show hardness of 6700 MPa which is a staggering improvement in this kind of alloys 189 . Hardness improvement of 910 HV by forming a multiphase nanostructured microstructure obtained after long time (100 h) annealing of HPT processed Cantor alloy is reported by Schuh et al. 190 . Cyclic HPT (changes in strain path) creates unstable dislocation structure and fine grains which is responsible for high hardness of CoCuFeMnNi alloy 191 . Room temperature and cryo-HPT led to high hardness and fine grain morphology in Cantor alloy 192 .

The synergy of tensile strength and ductility was demonstrated in case of HPT followed by annealing in V 10 Cr 15 Mn 5 Fe 35 Co 10 Ni 2 alloy, possessing 1.54 GPa UTS and 6% of ductility 194 . V 10 Cr 15 Mn 5 Fe 35 Co 10 Ni 25 alloy showcased a hardness of 505 MPa and tensile strength of 2 GPa with elongation failure of ~ 6%, post-HPT of coarse-and fine-grained starting material 195 . This was assisted by dislocation substructure formation along with twinning in the HEA. The Ashby plot for HPT processed and rolled HEAs is compared and shown in Fig. 10 . The properties of fine and nanometer size grains post-HPT as well as rolling have been depicted in the plot. Nanometer size grain formation in HPT processed CoCrFeMnNi alloy improved the superplasticity behaviour (> 600% total elongation) at high temperatures; grain boundary sliding being instrumental for the former behaviour 198, 199 . Addition of 2 atom % titanium in CoCrFeMnNi alloy followed by HPT showed 830% total elongation at 700 °C defining a new benchmark of superplasticity in HEA. This is possible due to grain size of 30 nm and retention of equiaxed nature of grains as titanium triggers sluggish diffusion 200 .

with High-PerformanceMaterials Post-RHP The comparison of tensile mechanical properties at room temperature of HEA with highperformance materials after RHP is presented in Fig. 10 [201] [202] [203] [204] [205] [206] [207] [208] [209] . The best strength ductility combination (PSD) of RHPed HEA (maximum PSD of 60 GPa × % FeCoCrNiMn-1 at % C) is in the same range of TWIP steel (PSD: 58 GPa × %) 204 , TRIP steel (PSD: 51 GPa × %) 208 The strength-ductility combination for many RHP HEA is also higher compared to microalloyed steel (PSD: 18.4 GPa × %) 201 , dual-phase steel (PSD: 11.2 GPa × %) 209 , and ferritic stainless steel (PSD: 16.8 GPa × %) 203 (Fig. 11) . The microstructural evolution in the above-mentioned alloys include precipitations inhibiting grain growth 203 , twinning 204 , grain refinement 207 , etc. The role of various parameters during rolling and heat treatment processes is significant in these high-performance materials also in deciding final strength-ductility combination 201 . Figure 10 shows compilation of various yield strength and PSD values of the RHPed HEA and their comparison with conventional RHP high-performance materials 201 -209 .

In this review, the role of processing has been emphasized on room-temperature mechanical properties of a wide range of HEAs. The key variables of rolling and heat treatment process (RHP) that can be tuned to enhance the performance of synthesized HEA have been elaborated and debated. The strengthening mechanisms in HEA which operate during the deformation are explained in terms of metamorphosis of the microstructure. These changes are dependent on the nature of phases, working temperature along with SFE of the HEA. Tuning the total strain, strain path, and symmetricity of rolling optimally could impart additional strengthening. Further in the processing, the annealing temperature primarily dictates the microstructure evolution. The optimization of annealing time, heating rate, etc. enhances the strength-ductility synergy further. Following the evaluation of the role of possible parameters in influencing the microstructure evolution and room-temperature tensile properties during RHP in HEA, some futuristic ideas are enlisted below:

1. The strength-ductility combination can be enhanced to a desirable magnitude by combination of the advantages of multiple RHP domains in terms of temperature and, e.g., combination of cryo-rolling and warm rolling as well as asymmetric and symmetric rolling (hybrid rolling). This combination will enable the generation of hierarchical heterostructures with gradient microstructures which will help to promote the abovementioned synergy. 2. Integrated Computational Materials Engineering (ICME) approach could find a solution for predicting processing-property correlation of HEA, by conducting lesser number of experiments. These simula-tion models can be used for better design of selective experiments to achieve optimal performance of HEA. The better design of experiments will promote energy efficiency and hassle-free methods to scale up the methods. 3. Microstructure and crystallographic texture simulation studies on HEA can be carried out to enable better maneuvering of microstructure-property correlation. 4. SPD techniques such as accumulative roll bonding can be performed to develop multilayers which inherently develop hierarchical microstructure (responsible for strengthductility alliance) owing to the non-uniform strain-induced during the process. Along with developing the above properties and microstructures, the scale up of these processes is highly possible which could be undertaken to compete with the conventional materials in the market.

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| VOL xxx:x | xxx-xxx 2022 | journal.iisc.ernet.in on the microstructure and mechanical properties of the Al5Ti5Co35Ni35Fe20 high-entropy alloy

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| VOL xxx:x | xxx-xxx 2022 | journal.iisc.ernet.in nanocrystalline CoCrFeMnNi high-entropy alloy after severe plastic deformation

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Effect of cryorolling on the microstructure and tensile properties of bulk nanoaustenitic stainless steel

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Microstructural evolution and mechanical properties of Ni-containing light-weight medium-Mn TRIP steel processed by intercritical annealing

Effect of continuous annealing process on various structure parameters of martensite of dual-phase steels

The authors are thankful to the online resources provided by Indian Institute of Technology, Madras in remote mode in the COVID-19 pandemic situation.

Authors declare that they have no potential conflict of interest. 

His specialization includes materials processing, crystallographic texture, and mechanical behaviour of materials. He is the author of more than 300 research papers and has co-authored/co-edited three books. He is a Humboldt fellow and has also been conferred with the Friedrich Wilhelm Bessel Award by the Humboldt foundation, Germany.