Strengthening and fracture of deformation-processed dual fcc-phase CoCrFeCuNi and CoCrFeCu1.71Ni high entropy alloys

Materials Science & Engineering A 781 (2020) 139241

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Strengthening and fracture of deformation-processed dual fcc-phase CoCrFeCuNi and CoCrFeCu1.71Ni high entropy alloys Yong Keun Kim , Byung Ju Lee , Soon-Ku Hong , Sun Ig Hong * Department of Materials Science and Engineering, Chungnam National University, Daejeon, 34134, South Korea

A R T I C L E I N F O

A B S T R A C T

Keywords: High entropy alloy Dual-phase Filament Heterogeneous lamella Strengthening

In this study, the microstructural evolution and strengthening mechanism of deformation-processed CoCrFeCuNi and CoCrFeCu1.71Ni were studied. Segregation and phase separation into dual fcc phases was found to be thermodynamically stable even after homogenizing annealing at 1273K. Dual fcc phase structure with elongated Cu-rich filaments developed by deformation processing and elongated dual-phase structure was found to remain stable even during annealing at1173 K. The high temperature stability of Cu-rich filaments is attributed to the low internal strain energy of filaments due to continuous recrystallization during deformation processing. One interesting consequence of stability of Cu-rich filaments is that the phase boundaries act as barriers to grain growth and the grain of Cr–Fe–Co–Ni matrix and Cu-rich filaments are controlled by the distribution and thickness of Cu-rich phase. The yield strengths of CoCrFeCuNi (506 MPa) and CoCrFeCu1.71Ni (477Mpa) were observed to be greater than that of the original Cantor alloy (220 MPa) with minor reduction of elongation less than 3–7%. Phase boundaries in CoCrFeCuNi or CoCrFeCu1.71Ni HEAs act as effective barriers to grain growth during annealing and to dislocation transmission during deformation. The strengthening with excellent ductility in HEAs of this study is attributed to the grain size refinement and the development of the enhanced back stress by the phase boundaries between Cu-rich filaments and Cu-lean matrix with higher modulus and strength. The lower strength of annealed CoCrFeCu1.71Ni than that of CoCrFeCuNi can be attributed to the presence of higher fraction of soft Cu-rich filaments.

1. Introduction The Cantor alloy, first high entropy alloy with the equiatomic CoCrFeMnNi composition, has been the subject of extensive studies and still attracted the interest of many investigators [1,2]. CoCrFeMnNi is known to have excellent combination of strength and ductility at ambient [3–7] and low temperatures [3–6,8] because of its stabilized solid solution phase structure [1–5]. Solid solution phase in CoCr­ FeMnNi is known to be stabilized by high configurational entropy especially at high temperatures [1–5]. However, phase decomposition and precipitation have been observed after prolonged annealing [9,10] or creep deformation [11] at intermediate temperatures. Many in­ vestigators have developed and studied modified Cantor-type alloys by substituting one major element of Cantor composition with Cu [12–15] or adding new elements such as Al [16], Ti or Nb [17]. The replacement of Ni by Cu in CrMnFeCoNi alloy was observe to cause the phase separation into two face centered cubic (fcc) phases and σ phase (in CrMnFeCoCu) [18]. The substitution of one element by

copper in quinary Cantor alloy composition was found to induce various microstructures with various phases despite the same configurational entropy implying that the configurational entropy is not a predominant factor [18]. Oh and Hong [12] investigated the compressive stress-strain responses of equiatomic CoCrFeCuNi, CuCrFeMnNi, CoCrFeMnCu al­ loys. In CuCrFeMnNi in which Cu substitute Co, σ phase and Cr-rich bcc phase as well as the dual fcc structure was observed [12]. The strength of CuCrFeMnNi increased due to the presence of Cr rich bcc particles and σ phase particles that act as barriers to slip. In CoCrFeCuNi and CoCuFeMnNi in which Mn or Cr was substituted by Cu, dual ductile fcc phase structures without the formation of σ phase were observed [12]. In these alloys, phase separation into Cu-lean phase and Cu-rich phase induces nanoscale or micro-scale microstructure depending on the reduction of enthalpy caused by the phase separation [19]. Excellent compressive deformability was observed in CoCrFeCuNi and CoCuFeMnNi because of the dual ductile fcc phase structures [12]. Recently, Shim et al. [19] observed nanoscale modulated dual fcc phase structure and the rapid work hardening rate in CoCuFeMnNi high

* Corresponding author. E-mail address: [email protected] (S.I. Hong). https://doi.org/10.1016/j.msea.2020.139241 Received 8 November 2019; Received in revised form 27 February 2020; Accepted 10 March 2020 0921-5093/© 2020 Elsevier B.V. All rights reserved.

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Materials Science & Engineering A 781 (2020) 139241

entropy alloys. They observed numerous short segment of dislocations spanned between the interfaces of the modulated structure and attrib­ uted the rapid work hardening to the stable nanoscale modulated dual fcc phase structure. Park et al. [14] studied the recrystallization behavior of CoCrCuFeNi alloy in which Mn was replaced by Cu and observed that their as-cast microstructure comprised of Cu-lean dendrite and Cu-rich inter-dendritic region. The microstructure and the microhardness of the as-cast and rolled dual fcc-phase-structured CoCrFeCuNi in which Cu substitute Mn have been studied by some investigators [13–15], but the tensile mechanical testing and deformation behavior of CoCrFeCuNi have never been studied and reported. In this study, the tensile deformation behaviors and microstructural evolution of deformation-processed CoCrFeCuNi and CoCrFeCu1.71Ni were studied. The elongated dual fcc phase struc­ ture is expected to develop by rolling since CoCrFeCrNi has two immiscible fcc phases and it would be interesting to investigate the ef­ fect of elongated Cu-rich phase on the microstructural evolution and mechanical properties. In this study, the microstructural evolution during heavy deformation processing by rolling and annealing was investigated. The effects of elongated Cu filament on the deformation behavior and mechanical stabilities were also studied.

because of its lower melting temperature and its immiscibility with Cr, Fe or Co [18,19]. In Fig. 1, SEM image (a, d), and elemental mapping images of Cr (b, e) and Cu(c, f) in the homogenizing annealed cast CoCrFeCuNi (a-c) and CoCrFeCu1.71Ni (d-f) are exhibited. Since Cu-rich phase consists of mostly Cu (>83 at. %) and Cu-lean phase consists of Cr, Co, Fe and Ni (each element >20 at.%) as summarized in Table 1, elemental mapping images of only Cu and Cr are exhibited in Fig. 1 to save the space. It is apparent that dendritic arms (the matrix) is Cr-, Fe-, Co- and Ni-rich phase, resulting in the formation of thin Cu-rich inder­ dendritic phase with a lower melting temperature. The spacing between separated Cu-rich phases decreased from 18.5 μm to 12.3 μm as the Cu content increased from 20 at. % to 30 at. %. The separation into two phases was found to be clearer in the as-cast alloys (not shown) [12]. The phase separation was found to be thermodynamically stable even after high temperature homogenizing annealing (1273K for 24 h) because of lower enthalpy of mixing (ΔHmix) and lattice strain energy (ΔHel) of the dual fcc phase separation than those of single fcc phase [19]. The stability of dual fcc phase after homogenizing annealing suggests that the mixing entropy at high temperatures is not large enough to offset the driving forces for the phase separation [18–21]. In Fig. 2, SEM image (a, d) and elemental mapping images of Cr (b, e) and Cu(c, f) of the annealed CoCrFeCuNi (a, b, c) and CoCrFeCu1.71Ni (d, e, f) after rolling are exhibited. As stated, Cu-rich phase consists of mostly Cu and Cu-lean phase consists of Cr, Co, Fe and Ni (Table 1). It is interesting to note that Cu-rich fcc phase developed into elongated Curich filaments during rolling remained stable even after annealing at 1173 K for 1 h. It would be interesting to compare the microstructural morphology of deformation-processed CoCrFeCuNi and CoCrFeCu1.71Ni in Fig. 2 to elongated filamentary structure of Cu-Fe-Co [22], Cu-Cr [23], Cu-Fe [24] and Cu–Nb [25] microcomposites. The microstruc­ tural morphology of deformation-processed CoCrFeCuNi and CoCrFe­ Cu1.71Ni in Fig. 2 appear to be close to the elongated two-phase or filamentary structure in two phase Cu-based microcomposite processed by rolling or drawing [22–25]. It should be noted that the elongated filament in CoCrFeCuNi alloys in the present study is softer Cu-rich phase filament in the harder Cr–Fe–Co–Ni rich matrix in contrast to the harder Cr, Fe or Nb filaments in Cu base microcomposites [22–25]. It would be interesting to examine how these elongated softer Cu filaments contribute to the strength of CoCrFeCuNi and CoCrFeCu1.71Ni in the present study. The structure and phase identifications in the as-cast, homogenizing annealed, as-rolled and annealed CoCrFeCuNi and CoCrFeCu1.71Ni al­ loys were examined by X-ray diffraction (XRD) and XRD patterns are exhibited in Fig. 3(a) and (b). XRD spectra from CoCrFeCuNi HEA (a) exhibit split double peaks from (111), (200) and (220) planes of fcc phases. The splitting of peaks became more apparent after homoge­ nizing annealing, supporting the thermodynamical stability of duplex fcc phase structure. In each double peak, left-hand side lower peak is from Cu-rich phase and right-hand side higher peak is from Cu-lean (Cr–Fe–Co–Ni rich) phase. No other peaks were observed, supporting no intermetallic compound formation. The expanded view of these peaks along the x axis revealed splitting or broadening of the peaks, supporting the presence of separated dual fcc phases. It is interesting to note, in the inset, that the (111) peak splitting became more apparent after homogenizing annealing (1000 � C) of cast ingot and annealing (900 � C) after rolling, implying the phase separation into dual fcc phase structure is thermodynamically stable. The separation of peaks became more apparent after final annealing at 900 � C. In CoCrFeCu1.71Ni, the (111) peak splitting also became more apparent after annealing (900 � C) as shown in the inset of Fig. 3(b). The separation of Cu-rich phase from Cu-lean phase is clearer because of larger fraction of Cu-rich phase in CoCrFeCu1.71Ni in the as-cast structure. In CoCrFeCu1.71Ni (Fig. 3(b)), the (111) peak from Cu-rich phase appears to be suppressed while that from Cu-lean phase sharpened after rolling. The suppression of peaks from Cu-rich phase after rolling could be attributed to more prevalent development of lattice strain and

2. Experimental In this study, equiatomic CoCrFeCuNi and non-equiatomic CoCrFe­ Cu1.71Ni ingots were cast by arc melting a mixture of Co (purity: 99.99%), Cr (purity: 99.97%), Fe (purity: 99.99%), Cu (purity: 99.99%), and Ni (purity: 99.99%) in a pure Ar atmosphere. The ingots were remelted repetitively to enhance chemical and microstructural homo­ geneity [12,19]. The cast ingot was homogenizing annealed at 1273K for 24 h and air-cooled. The complete homogenization with uniform microstructure and chemical compositions cannot be achieved by high temperature homogenizing annealing of cast ingot at 1273K because phase separation is thermodynamically stable at high temperatures. Homogenizing annealed ingots were cold-rolled into plates of 1.0 mm thickness (91% reduction) at room temperature and annealed at 1173K for 1 h to obtain the recrystallized microstructure [19]. Dog-bone sha­ ped tensile specimens with a gauge length of 9 mm were machined from as-rolled and annealed plates, respectively. Tensile testing was per­ formed at an engineering strain rate of 10 3s 1. Microstructural and phase analyses of the alloys were carried out using optical microscopy (OM), X-ray diffraction (XRD), and scanning electron microscopy (SEM) coupled with energy dispersive X-ray spec­ troscopy (EDS). The grain structure and phase distribution were also characterized using an electron backscatter diffraction (EBSD) system (Oxford Instruments, UK) attached to a field-emission scanning electron microscope (FE-SEM; Helios, Pegasus, FEI). Since the grain growth was observed to be limited by the phase boundaries, the average grain size in Cu-rich phase and Cu-lean phase were measured separately in each phase by the linear intercept method. Transmission electron microscope (TEM) specimens were prepared using the dimple grinder (Gatan, model 656, USA) and the precision ion polishing system (Gatan, model 691, USA), with argon gas ions at 3.6 keV. TEM observations were carried out by using a JEOL JEM 3010 operated at 300 keV. Diffraction pattern was obtained by using a smallest selected area diffraction (SED) aperture with a diameter of 190 nm. High resolution TEM (HRTEM) and scanning transmission electron microscope (STEM) observations were carried out using a JEOL JEM-ARM 200F, in which Cs-corrected STEM and EDX (Bruker QUANTAX 400 model) are equipped. 3. Results 3.1. Micro-scale phase separation and deformation processing by rolling In the Cantor variant alloys with copper replacing one element of Cantor composition, copper tends to form the interdendritic fcc phase 2

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Fig. 1. SEM image (a, d) and elemental mapping images of Cr (b, e) and Cu(c, f) of the homogenizing annealed CoCrFeCuNi (a, b, c) and CoCrFeCu1.71Ni (d, e, f). Table 1 Composition of dual fcc phases and their experimental (XRD) and predicted (Vegard’s law) lattice constants in homogenized CoCrFeCuNi and CoCrFeCu1.71Ni. Alloy Type

Phase

CoCrFeCuNi

Cu Cu Cu Cu

CoCrFeCu1.71Ni

lean Phase Rich Phase lean Phase Rich Phase

Composition (at.%)

Lattice Constant (nm)

Co

Cu

Fe

Cr

Ni

XRD

Predicted

22.70 3.61 23.36 3.44

9.34 83.50 8.13 83.46

22.20 3.13 22.33 3.42

21.25 2.69 21.81 2.81

24.51 7.07 24.37 6.87

0.3567 0.3604 0.3561 0.3624

0.3536 0.3605 0.3535 0.3604

Fig. 2. SEM image (a, d) and elemental mapping images of Cr (b, e) and Cu(c, f) of the annealed CoCrFeCuNi (a, b, c) and CoCrFeCu1.71Ni (d, e, f) after cold-rolling.

shearing of softer Cu-rich filaments during deformation processing [25, 26]. The re-separation of dissolved phases after annealing was also observed in CoCrFeCu1.71Ni again driven by the reduction of enthalpy during annealing. The reason why the splitting of (111) peaks is clearer after annealing at 900 � C than after homogenizing annealing at 1000 � C in Fig. 3(a) and (b) is thought to be caused by the decreasing contri­ bution of entropy at 900 � C. Splitting of the peaks supporting the pres­ ence of separated dual fcc phases are more clearly visible in the XRD patterns of CoCrFeCu1.71Ni as shown in Fig. 3(b). The compositions of two separated phases from EDS analyses in the annealed CoCrFeCuNi and CoCrFeCu1.71Ni are summarized in Table 1 along with the lattice constants determined using the XRD data and those obtained by the Vegard’s law using the atomic radii of constituent elements in Table 1. The lattice constants determined using the XRD peaks for two fcc phases in Fig. 3 were 0.3567 nm and 0.3604 nm for CoCrFeCuNi, both of which are close to those (0.3536 nm and 0.3605

nm) calculated by the Vegard’s law. The lattice constants determined using the XRD peaks for CoCrFeCu1.71Ni were 0.3561 nm and 0.3624 and both are close to those (0.3535 and 0.3604 nm) calculated by the Vegard’s law (Table 1). 3.2. Filamentary structure and grain-size refinement in dual fcc phase structure In order to investigate the microstructural evolution during rolling and subsequent annealing of the homogenizing annealed dual phase CoCrFeCuNi and CoCrFeCu1.71Ni alloys, the grain structure and phase distribution were characterized using EBSD and EDS. Since dual fcc phases have the similar lattice constants especially when the lattice strain was induced during deformation processing, two fcc phases became almost indistinguishable in EBSD inverse pole figure (IPF) map. EBSD and EDS images were obtained from the same location at the same 3

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Fig. 3. XRD diffraction patterns of as-cast, homogenizing-annealed, as-rolled and annealed CoCrFeCuNi (a) and CoCrFeCu1.71Ni (b).

time to distinguish two fcc phases more clearly in EDS IPF image. In Fig. 4, Cr (a) and Cu (b) elemental mapping image and EBSD IPF image (c) of cold-rolled CoCrFeCuNi are exhibited. In the EBSD IPF image (c), the phase boundaries between Cu-rich fcc phase and Cu-lean (Cr–Fe–Co–Ni rich) phase are delineated by white dashed lines. It is interesting to note that grain boundaries of equi-axed or slightly elon­ gated ultrafine grains (indicated by red arrows) are clearly visible in Curich phase filaments despite the heavy cold-rolling up to 91% reduction of thickness. Cu-rich phase is thought to undergo more deformation because it is softer than Cu-lean (Cr–Co–Fe–Ni rich) phase. Early studies of severe plastic deformation of metals at low temperatures revealed the presence of equi-axed ultrafine grains [27,28]. The ultrafine-grained structure was suggested to develop by gradual transformation of the deformed cellular dislocation substructure to ul­ trafine grains progressively by generation, glide and deposition of dis­ locations that would increase the misorientation angles of subboundaries [27,28]. Different colors of neighboring grains in Cu-rich phase filaments suggest the high misorientation angles between grains. On the other hand, similar or same colors in Cu-lean (Cr–Co–Fe–Ni rich) phase suggest the textured elongated grains with low or nil misorien­ tation angles. In the Supplementary Fig. S1, EBSD IPF map (a) and image quality (IQ) map (b) with grain boundary misorientation distribution for cold-rolled CoCrFeCuNi are exhibited. This figure clearly shows that the misorientation angles between grains in Cu-rich phase are greater than 15� , supporting the continuous recrystallization in Cu phase during cold-rolling. In Cu-lean (Cr–Co–Fe–Ni rich) phase, the misorientation angles for major fraction of boundaries are lower than 5� and some fraction is with misorientation angles of 5–15� . In the Supplementary

Figs. S2(a) and S2(b), TEM image (a) and HRTEM image (b) near interface regions of Cu-rich filaments/Cr-lean matrix in cold-rolled CoCrFeCuNi are exhibited. This figure clearly shows the presence of recrystallized grains in Cu-rich filament, supporting the recrystallized ultrafine grains (shown in the EBSD IQ map) formed by continuous recrystallization during cold-rolling. HRTEM image in Fig. S2(b) also support that the interface between Cu-rich filaments and Cr-lean matrix is incoherent with high misorientation angle. Fig. 5 exhibits Cr (a) and Cu (b) elemental mapping image and EBSD IPF image of annealed CoCrFeCuNi. Again, the boundaries between Curich fcc phase and Cu-lean (Cr–Fe–Co–Ni rich) phase are delineated by white dashed lines in the EBSD IPF image (c). It is apparent that extensive recrystallization occurred in Cr–Fe–Co–Ni phase after annealing. More annealing twins were observed in Cr–Fe–Mn–Ni phase, suggesting the stacking fault energy of Cr–Fe–Co–Ni phase is lower [29–31]. It is interesting to note that phase boundaries between Cu-rich phase and Cr–Fe–Co–Ni phase act as barriers to grain growth and limit the grain size in each phase. In the Supplementary Fig. S3, EBSD IPF map (a) and image quality (IQ) map (b) for annealed CoCrFeCuNi are exhibited. As expected, the misorientation angles between grains in Cu-rich phase are greater than 15� . The misorientation angles for major fraction of boundaries in Cu-lean (Cr–Co–Fe–Ni rich) phase are greater than 15� with some partial recrystallized region with misorientation angles lower than 15� . In the Supplementary Figs. S4(a) and S4(b), TEM image (a) and HRTEM image (b) near interface regions of Cu-rich fila­ ments/Cr-lean matrix in annealed CoCrFeCuNi are exhibited. After annealing at 1173K, fully recrystallized grain with its width limited by the thickness of Cu-rich filament was observed as shown in Fig. S4(a). In

Fig. 4. Cr (a) and Cu (b) elemental mapping image and EBSD image (c) of cold-rolled CoCrFeCuNi high entropy alloy. 4

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Fig. 5. Cr (a) and Cu (b) elemental mapping image and EBSD image (c) of annealed CoCrFeCuNi high entropy alloy.

Fig. S4(b), the HRTEM lattice image in Cu-lean matrix cannot be ac­ quired simultaneously because of high misorientation angle between Cu-rich filaments and Cr-lean matrix. Fig. 6 shows the Cr (a) and Cu (b) elemental mapping image and EBSD IPF map of cold-rolled CoCrFeCu1.71Ni. Equi-axed or slightly elongated ultrafine grains (indicated by red arrows) with high misori­ entation angle grain boundaries are visible in Cu-rich phase filaments, implying continuous recrystallization also in CoCrFeCu1.71Ni. Fig. 7 exhibits Cr (a) and Cu (b) elemental mapping images and EBSD IPF image (c) of annealed CoCrFeCu1.71Ni. In annealed CoCrFeCu1.71Ni, the phase boundaries between Cu-rich fcc phase and Cu-lean (Cr–Fe–Co–Ni rich) phase act as barriers to grain growth is apparent since extensive recrystallization occurred separately in both Cr–Fe–Co–Ni phase and Curich phase and grains in each phase rarely crossed the phase boundaries. Annealing twins were observed more often again in Cr–Fe–Co–Ni phase [29–31]. Larger areas with more grains than shown in Figs. 5 and 7 are required in order to make a more accurate statistics of grain size. We did observe the larger area at a lower magnification to get the statistically accurate grain size and examples of EBSD and EDS images showing lager areas for CoCrFeCuNi and CoCrFeCu1.71Ni are shown in Supplementary Fig. S5. The grain size in Cu-rich phase and Cu-lean (Cr–Fe–Co–Ni) phase of CoCrFeCuNi was measured to be 1.37 μm and 1.13 μm, respectively and that in Cu-rich phase and Cu-lean phase of CoCrFe­ Cu1.71Ni was measured to be 1.24 μm and 1.02 μm, respectively. The incoherent interfaces between Cu-rich filament and Cu-lean

matrix with high misoientation angles in CoCrFeCoNi HEAs are quite different from conventional dual fcc phase Cu-24 wt % Ag alloy in which the Cu/Ag interfaces are highly coherent. The difference of interface nature between CoCrFeCoNi HEAs and Cu–24Ag alloy is thought to be associated with difference of mutual solubility (nil solubility between Cu and matrix in HEA versus ~10% mutual solubility in Cu–Ag at inter­ mediate temperatures) and difference of second phase filament thick­ ness (0.5–3.0 μm in HEA versus 0.6–10 nm in Cu–24Ag. Supplementary figures Fig. S3 (EBSD IQ map) and S4 (TEM pictures) showing mostly high angle boundaries both in Cu-rich filament and Cu-lean lamella and mostly incoherent interfaces between Cu-rich filament and Cu-lean lamella supports the high misorientation angle phase/grain bound­ aries in annealed HEAs of the present study. Since the grain growth in Cu-rich phase and Cu-lean phase is limited by the thickness and distri­ bution of Cu-rich phase, a good fraction of grain boundaries formed at phase boundaries in Cu-rich phase because the Cu-rich phase is rela­ tively thin (0.5–3.0 μm). The understanding of nature of grain bound­ aries and interphase boundaries are very important in predicting the dislocation behaviors at these boundaries. Grain/phase boundaries with high misorientation angles and high mismatch (as evidenced by EBSD IQ map and HRTEM inages) would act as effective obstacles to dislocation motion. Dislocations tend to pile up or accumulate at these boundaries with high misorientation angles and are likely increase the back stress.

Fig. 6. Cr (a) and Cu (b) elemental mapping image and EBSD image (c) of cold-rolled CoCrFeCu1.71Ni high entropy alloy. 5

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Fig. 7. Cr (a) and Cu (b) elemental mapping image and EBSD image (c) of annealed CoCrFeCu1.71Ni high entropy alloy.

3.3. Mechanical performances of as-rolled and annealed CoCrFeCuNi and CoCrFeCu1.71Ni

increase of ductility by annealing in duplex fcc phase CoCrFeCuNi HEAs of the present study is in sharp contrast to the observation in duplex fcc phase Cu-24 wt % Ag alloy [26]. In Cu–24Ag alloy, a negligible change of ductility (from 7.1% to 6.3%) was observed by annealing that induced the decrease of yield stress from 866 MPa to 442 MPa (the decrease of yield stress by about a half) (Fig. 10 of Ref. 26). The different strength-ductility trade-off behavior between CoCrFeCuNi HEAs and Cu–Ag filamentary composite may be associated with the difference of interfaces and heterogeneous structures of dual fcc phases and will be explained in Discussion. The lower ultimate tensile strength of as-rolled CoCrFeCu1.71Ni alloy than that of CoCrFeCuNi can be attributed to the presence of higher fraction of softer Cu-rich filaments. The yield strengths of annealed CoCrFeCuNi and CoCrFeCu1.71Ni were similar, but the tensile strength of CoCrFeCu1.71Ni was found to be slightly lower than that of CoCrFeCuNi because of the decrease in hardening rate with increasing fraction of softer Cu-rich phase. In Fig. 9, SEM image (a) and the elemental mapping images of Cr (b) and Cu(c) of the longitudinal section close to the fracture surface of asrolled CoCrFeCuNi high entropy alloy deformed until fracture are shown. As stated, the elongated Cu-rich filaments developed parallel to the specimen axis. Cu filaments are not straight but wavy because of the inclined slip along the slip planes in both phases during heavy rolling. Some cracks propagated along the inclined Cu-rich phase filaments (indicated by blue arrows) while some other cracks cut through the aligned Cu-rich phase filaments (indicated by red arrows). In the region enclosed by yellow dotted rectangle, the shape of Cu-rich filaments near fracture surface were not greatly modified by tensile deformation to fracture, reflecting low fracture elongation of as-rolled CoCrFeCuNi. Fig. 10 displays SEM image (a) and the elemental mapping images of Cr

Fig. 8 shows the tensile stress-strain curves of as-rolled (a) and annealed (at 1173 K) (b) CoCrFeCuNi and CoCrFeCu1.71Ni high entropy alloys. Yield strengths of both as-rolled CoCrFeCuNi (1065 MPa) and CoCrFeCu1.71Ni (1020 MPa) high entropy alloys were found to be a little smaller than that of Cantor alloy (1080 MPa) [30–32]. On the other hand, yield strengths of annealed CoCrFeCuNi (502 MPa) and CoCrFe­ Cu1.71Ni (497 MPa) were observed to be greater than that (~220 MPa) of annealed Cantor alloy [30–32]. In Fig. 8(b), the stress-strain curves of Cantor alloy annealed at 1173 K reported by Ko et al. [32] are also plotted for comparison. It should be noted that the yield strength increased dramatically from 220 MPa in CoCrFeMnNi to 480–510 MPa in CoCrFeCuNi and CoCrFeCu1.71Ni by substituting Mn with Cu with minor reduction of elongation less than 3–7%. The ultimate tensile strengths of CoCrFeCuNi and CoCrFeCu1.71Ni are still greater than that of Cantor alloy [9,30,32] despite their lower work hardening rate. The lower strain-hardening rate of HEAs of the present study are thought to be attributed to the presence of softer Cu rich phase. The higher strengths of CoCrFeCuNi alloys compared to the Cantor alloy are attributed to the formation of smaller grain-sized structures limited by the distribution of Cu-rich filaments. It is interesting to note the minor reduction of ductility despite a more than two-fold increase of the yield stress (~500 MPa versus 220 MPa) compared to those of the Cantor alloy as shown in Fig. 8. It should also be noted that the ductility of CoCrFeCuNi alloys increased signifi­ cantly from ~6% to ~50% after annealing that induced the decrease of yield stress by a half (from ~1000 MPa to ~500 MPa). The drastic

Fig. 8. Tensile stress-strain curves of as-rolled (a) and annealed (b) CoCrFeCuNi and CoCrFeCu1.71Ni alloys. 6

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Fig. 9. SEM image (a) and the elemental mapping images of Cr (b) and Cu(c) of the longitudinal sections close to the fracture surfaces of as-rolled CoCrFeCuNi high entropy alloy deformed until fracture.

the range of 0 < ΔHel � 6.9 kJ mol 1 and 10.7 � ΔHmix � 3.9 kJ mol 1 [20]. However, even though the mixing enthalpy (ΔHmix) versus lattice strain energy (ΔHel) criterion for solid solution is satisfied, phase sepa­ ration would occur if the sum of ΔHmix and ΔHel further decreases by the compositional separation or variation. The enthalpy reduction criterion [19] used in the present study within the ranges of solid solutions would add to the accuracy of predictability for single solid solution phase formation or phase separation predicted by the mixing enthalpy versus elastic-strain energy criterion of Andreoli et al. [20]. The mixing enthalpy, ΔHmix and the elastic lattice distortion energy, ΔHel of quinary solid solution alloys can be calculated using the following equations [19,20,34]:

(b) and Cu(c) of the longitudinal section close to the fracture surface of annealed CoCrFeCuNi high entropy alloy deformed to fracture. Note that some thin Cu-rich filaments in annealed CoCrFeCuNi appears to be sheared during extensive deformation to fracture up to 50% (as shown in dotted rectangles in Fig. 10). Some secondary cracks were formed by separation along the aligned softer Cu-rich phase (marked by yellow arrows) in the annealed CoCrFeCuNi, inducing necking of separated columns. The images near fracture surface for CoCrFeCu1.71Ni are similar to those observed in Figs. 9 and 10 and will not be shown. 4. Discussion 4.1. Reduction of ΔHmix and ΔHel by phase separation in CoCrFeCuNi and CoCrFeCu1.71Ni

5 X

ΔHmix ¼

It has been suggested that the strain energy contribution as well as the chemical enthalpy term is important in the stabilities of various phases in multicomponent alloys [19,20]. The discrepancy between experimental observations and thermodynamic predictions in some high entropy alloys [19,20,33,34] can be ascribed to the underestimation of the strain energy in the total enthalpy. The mixing enthalpy (ΔHmix) versus lattice strain energy (ΔHel) criterion was proposed by Andreoli et al. [20] and the regions of fcc solid solution region was found to lie in

(1a)

ci cj Ωij i; j;

i6¼j

5 X

ΔHel ¼

ci Bi i¼1

VÞ2

ðVi 2Vi

(1b)

where Ωij ¼ 4ΔHijmix, ci and cj are the atomic fractions of the element i and j, ΔHijmix is the enthalpy of mixing of a liquid binary alloy, and Bi, Vi, V are the bulk modulus, the atomic volume of the element i and the average volume of atom of the alloy [19,20,34]. If the reduction of sum

Fig. 10. SEM image (a) and the elemental mapping images of Cr (b) and Cu(c) of the longitudinal sections close to the fracture surfaces of annealed CoCrFeCuNi high entropy alloy deformed until fracture. 7

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of mixing enthalpy (ΔHmix) and lattice strain energy (ΔHel) upon the phase separation is greater than the possible increase of elastic misfit strain energy of two separated phases at the interface [19,35], dual fcc-phase separation takes place spontaneously [19,36,37]. The changes of ΔHmix and ΔHel before and after separation into dual fcc-phases was calculated and summarized in Table 2. The average enthalpy of mixing (ΔHmix) and lattice strain energy (ΔHel) of the separated structure were calculated based on the volume fraction of two fcc phases. Since the atomic weight of 5 elements are not much different from each other, the volume percentage of Cu-rich filaments in CoCrFeCuNi and CoCrFe­ Cu1.71Ni alloys were conveniently assumed to be 20% and 30%, respectively. The separation into dual fcc-phases implies that the sum of the mixing enthalpy, ΔHmix and the elastic lattice distortion energy, ΔHel of separated dual phases is lower than that of homogeneously mixed quinary CoCrFeCuNi and CoCrFeCu1.71Ni alloys. The calculated values of ΔHmix and ΔHel for CoCrFeCuNi were 3.52 kJ mol 1 and 1.60 kJ mol 1, respectively. And the values of ΔHmix and ΔHel for CoCrFeCu1.71Ni were 4.71 kJ mol 1 and 2.88 kJ mol 1, respectively. The enthalpy of mixing and lattice strain energy decreased significantly with separation into dual fcc phases in both CoCrFeCuNi and CoCrFeCu1.71Ni, providing the pronounced driving force for dual fcc-phase separation. The large reduction of enthalpy (sum of ΔHmix and ΔHel) after dual fcc phases from homogeneous single fcc phase enables the separated dual fcc phase stable during heavy rolling as shown in Fig. 4 or heat treatments (Figs. 2 and 5). It is apparent that the separation into dual fcc solid solution phases occurs spontaneously by the reduction of enthalpy of mixing and lattice strain energy even if the multicom­ ponent alloys meet the criteria for the complete homogeneous solid solution fcc phase.

Cu-based filamentary composites with different crystalline structures and larger differences of lattice constants. The spheroidization and splitting of Nb and Ag filaments in Cu-based microcomposites are induced by the reduction of mostly cold-work energy stored in deformed Nb or Ag filamnets [25,26,38–40]. In CoCrFeCuNi alloys of the present study, continuous recrystalli­ zation occurred in Cu-rich phase filament during cold-rolling up to 91% reduction of thickness as shown in Figs. 4 and 6 and Supplementary Fig. S1 because of more severe deformation and lower recovery tem­ perature of Cu-rich phase. Because of continuous recrystallization in Curich phase, negligible cold work energy is stored in the Cu-rich phase. Therefore, low cold work energy storage in Cu-rich phase filaments and low interface energies [41] in CoCrFeCuNi and CoCrFeCu1.71Ni alloys would lower the driving force for spheroidization and/or break-up of Cu-rich phase filaments, stabilizing the elongated dual phase fcc struc­ tures. The low solubility of Cu in the matrix and the sluggish diffusion in CrFeCoNi HEA [42,43] matrix may have contributed to the stability of Cu-rich phase filaments. One interesting consequence of stability of Cu-rich filaments is that they act as barriers to grain growth and the grain sizes in both Cr–Fe–Co–Ni matrix and Cu-rich filaments are controlled by the distribution and thickness of Cu-rich phase. 4.3. Strength enhancement and good ductility in CoCrFeCuNi and CoCrFeCu1.71Ni The increase of strength is generally accompanied by the pronounced loss of ductility [44,45]. One interesting observation in this study is that yield strengths of annealed CoCrFeCuNi (502 MPa) and CoCrFeCu1.71Ni (497 MPa) were more than twice greater than that (~220 MPa) of annealed Cantor alloy [9,30,32] with inappreciable loss of ductility (less than 3–7%). The grain sizes of CoCrFeCuNi and CoCrFeCu1.71Ni in the present study were found to be much smaller than that of Cantor alloy. Grain refinement can increase the strength of conventional metals greatly, but with the dramatic loss of ductility [46]. Recently, many measures to overcome the strength ductility trade-off have been devel­ oped [44]. Such measures include engineering heterostructured grains [45], heterogeneous lamella [46–48], gradient structure [49], dual phase structure [50] and grain boundary structure [51]. Wu et al. [46] suggested that the heterogeneous lamella structure characterized with soft micrograined lamellae embedded in hard ultrafine-grained lamella matrix exhibited unusual high strength and high ductility. Wu et al. [44] also suggested that the enhancement of the strength-ductility trade-off resulted from the microstructure comprised with a combination of the non-recrystallized and recrystallized grains arranged in complex het­ erogeneous structures. Those heterogeneous structures were produced by cold-rolling, followed by intermediate-temperature-annealing [44, 46]. The enhancement of the strength-ductility trade-off in dual fcc phase CoCrFeCuNi and CoCrFeCu1.71Ni can be attributed to the char­ acteristic heterogeneous lamella structure [44,46]. CoCrFeCuNi and CoCrFeCu1.71Ni exhibit heterogeneous lamella structure consisting of Cu-rich phase and Cu-lean (Cr–Fe–Co–Ni rich) phase. The heterogeneous lamella structure of the present study is different from those of Wu et al. [44,46] in which the heterogeneity resulted from the different degrees of deformation and/or recrystalli­ zation. The unusual high strength in heterogeneous lamella was sug­ gested to be caused by high back stress and the high ductility was attributed to back-stress hardening and dislocation hardening [46]. As stated, the phase boundaries between Cu-rich phase and Cr–Fe–Co–Ni phase in the present study act as barriers to grain growth and limit the grain size in each phase. The grain size in both phases are very small (1.0–1.3 μm) and the difference of strengths in the heterogeneous lamella structure of the present study is not caused by the differences in grain size or degree of recrystallization, but by the difference of elastic moduli and strengths between two fcc phases. The shear moduli (G) for Cu [52] and Cr–Fe–Co–Ni [53] were reported to be 40 GPa and 84 GPa, respectively. The yield stress (YS) of coarse-grained (24–48 μm)

4.2. Grain/phase boundaries and grain size refinement by Cu-rich phase filaments One of the most interesting observations in rolled filamentary CoCrFeCuNi and CoCrFeCu1.71Ni HEA is the stability of Cu-rich fila­ ments during annealing at 1173K for 1hr. as shown Figs. 2, Figs. 5 and 7. The stability of Cu-rich filaments in dual fcc structure HEAs after annealing at 1173K is in sharp contrast to the instabilities of Cr, Fe and Nb filaments in Cu-based microcmposites [22–26] during annealing at temperatures as low as 400–500 � C. During deformation processing of filamentary Cu-Nb [25,38] and Cu–Ag [26,39] microcomposites, dislo­ cations are stored in the elongated Nb or Ag filaments and dislocations generated in the matrix are heavily deposited at the interfaces because of the differences of lattice structure and lattice constant. Dislocation deposition at the interface of the filaments and dislocation storage in the elongated filaments increases the interface energy and cold work energy of filaments. In HEAs of the present study, the deposition rate of inter­ face dislocations is expected to be smaller compared with other Table 2 Calculated enthalpy of mixing (ΔHmix) and elastic lattice distortion energy (ΔHel) before and after separation into dual fcc phases from a single fcc phase with higher enthalpies.ΔHel Alloy Types CocrFeCuNi

CoCrFeCu1.71Ni

Phase Single Phase Cu lean Phase Cu Rich Phase Single Phase Cu lean Phase Cu Rich Phase

Volume fraction

ΔHmix ðKJ ⋅mol

1

3.52

0.8

0.25

0.2

4.10

1

4.71

0.7

0.29

0.3

4.123

1

Þ

ΔHel ðKJ ⋅mol

1

Þ

1.60 Dual fcc phase 1.06 Dual fcc phase 1.04

1.264 0.787 2.88 1.249 0.772

Dual fcc phase 1.17 Dual fcc phase 1.11

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Materials Science & Engineering A 781 (2020) 139241

Cr–Co–Fe–Ni (295 MPa) [53] was also reported to be greater than that of Cu (82 MPa) with the similar grain size (25 μm) [54]. CoCrFeCuNi and CoCrFeCu1.71Ni can be considered as high entropy alloy composites with Cu-rich filaments and their strengths can be predicted [25,26] using the rule of mixture as follow;

increase the back stress at the interface [58]. Inappreciable loss of ductility in ultrafine-grained HEAs in this study can also be attributed to the back stress as suggested by Wu et al. [46]. The high misorientation angles in grain/phase boundaries in the present study have the positive influence on the ductility enhancement [51]. The formation of high angle grain boundaries (HAGBs) in Cu rich filaments (even during cold-rolling) as shown in the supplementary figure (Fig. S1) and Cu lean matrix during high temperature annealing (Fig. S2) at 1173K enhanced the stability of grain boundaries. The formation of ultrafine-grained structure with HAGBs corroborates the view on the effect of HAGBs on the high ductility in this study. Grain boundaries present effective obstacles to dislocation motion, and the elastic interaction between lattice dislocations and grain boundaries is repulsive [51], causing the back stress. Consequently, the dislocations tend to pile up or accumulate at the grain/phase boundaries and increase the back stress. Back stress by pile-up models has been widely used as an explanation for the empirical Hall–Petch relationship [59,60]. The temperature dependence of back stress by dislocation pile-ups was also reported to be close to that of shear moduli [61]. The back stress build-up at Cu-rich filament grain/phase boundaries with high misorientation angles can be enhanced due to the differences of modulus and strength between Cu rich filament and Cu lean matrix. In conclusion, the strengthening with excellent ductility in HEAs of this study is associated with grain size refinement and the development of the enhanced back stress by the phase boundaries between Cu-rich fil­ aments and Cr–Fe–Co–Ni matrix. Phase boundaries in CoCrFeCuNi or CoCrFeCu1.71Ni HEAs act as effective barriers to grain growth during annealing and to dislocation transmission during deformation. Hetero­ geneous lamella structure in CoCrFeCuNi filamentary composites con­ sists of softer Cu-rich filament embedded in harder Cr–Fe–Co–Ni matrix with greater modulus and strength. The filamentary lamella-type HEA composites in the present study may offer a new design strategy toward heterogeneous lamella structure materials with attractive mechanical properties.

(2)

σHEC ¼ fmσm þ ffσf

where σHEC is the yield strength of high entropy alloy composite and σm and σf are yield stress of Cr–Fe–Co–Ni matrix and Cu rich filament, respectively. Yield stresses of Cr–Fe–Co–Ni matrix and Cu rich filament were predicted using the grain size data in Table 3 and Hall-Petch equations (σy ¼ σo þ k d 1/2). fm and ff are volume fractions of Cr–Fe–Co–Ni matrix and Cu-rich filaments. As discussed in 4.1., the volume percentage of Cu-rich filaments in CoCrFeCuNi and CoCrFe­ Cu1.71Ni alloys were conveniently assumed to be 20% and 30%, respectively. The friction stress σo and Hall-Petch coefficient k of Cr–Fe–Co–Ni matrix were assumed to be close to that of typical Cantor ally (σo ¼ 102 MPa, k ¼ 538 MPa mm1/2) [40,55,56]. The friction stress and Hall-Petch coefficient of Cu-rich filament was assumed to be close to those of copper (σo ¼ 40 MPa, k ¼ 110 MPa mm1/2) [57]. The yield stress was calculated using equation (2) and the calculated stress was sum­ marized in Table 3 in comparison with the experimental data. The grain size in Cr–Fe–Co–Ni HEA matrix and Cu-rich filament was observed to be 1.0–1.4 μm as summarized in Table 3. On the other hand, the grain size of Cantor alloy processed the same way and annealed at 1173 K was reported to be ~20 μm [32]. The grain growth was observed to be effectively restricted by the elongated Cu-rich filaments because of their high stability even at high temperatures. The relatively smaller grain size in Cr–Fe–Co–Ni matrix suggests the presence ultramicron-scale Cu-rich filaments that may not be easily identified in the low magnification SEM figures. Cu-rich phase boundaries act as barriers to grain growth in both Cu-rich phase and Cr–Fe–Co–Ni matrix. The predicted yield stresses in CoCrFeCuNi, CoCrFeCu1.71Ni and Cantor alloy were in excellent agreement with the experimentally measured values in Table 3. The excellent agreement between the higher yield strengths of CoCrFeCuNi (506 MPa) and CoCrFeCu1.71Ni (477 MPa) with predictions by Hall-Petch equations supports that the grain size refinements contribute to the strengthening greatly. Wu et al. [46] suggested that the soft lamellae of recrystallized micrograins in the heterogeneous lamella structure would start plastic deformation first under tensile loading. In CoCrFeCuNi or CoCrFeCu1.71Ni HEAs, the initial plastic deformation would start in Cu-rich phase, but the yield stresses of HEAs in the present study are far greater than those of Cu rich phase (127.5–132.8 MPa) predicted for Cu by Hall-Petch equation as in Table 1. The higher yield stresses of HEAs in this study can be explained by the suggestion of Wu et al. [46] that the soft lamellae are strength­ ened by surrounding hard lamellae that exert the significant back stress at the onset of yielding. In Ti [46] and Al0.1CoCrFeNi HEA [44] with heterogeneoue lamella structures, the back stress was suggested to be induced by the hard lamella with smaller grains or more-deformed structure. In the present study, the Cu-lean (Cr–Co–Fe–Ni) hard lamella with greater elastic modulus and strength would amplify the internal back stress and thereby increase the yield strength of softer Cu-rich lamella. It is also well known that the image force of hard second phase particles or films

5. Conclusions We studied the microstructural evolution and strengthening of deformation-processed CoCrFeCuNi and CoCrFeCu1.71Ni high entropy alloys and obtained the following conclusions. 1) In homogenizing-annealed cast CoCrFeCuNi and CoCrFeCu1.71Ni alloys, Cu-rich and Cu-lean (Cr–Fe–Co–Ni rich) dual fcc phase structure with the similar lattice constants were obtained. The separated Cu-rich phase developed into Cu-rich filament during deformation processing and CoCrFeCuNi HEAs can be considered as filamentary microcomposites with Cu-rich filaments embedded in Cr–Fe–Co–Ni high entropy matrix 2) The reduction of total enthalpy (the mixing enthalpy (ΔHmix) and lattice strain energy (ΔHel)) due to compositional fluctuation within solid solution regions can cause phase separation as in CoCrFeCuNi and CoCrFeCu1.71Ni alloys. 3) The high stability of Cu-rich filaments in dual fcc structure HEAs after annealing at 900 � C is in sharp contrast to the instabilities of Cr, Fe and Nb filaments in conventional Cu-based filamentary compos­ ites. The stability of Cu-rich filaments is attributed to the low internal

Table 3 Experimental and predicted yield stresses of annealed CoCrFeCuNi and CoCrFeCu1.71Ni predicted using the rule of mixture and Hall-Petch equation. The yield stress and grain size of Cantor alloys are also summarized for comparison. Alloy type CoCrFeCuNi CoCrFeCu1.71Ni CoCrFeMnNi (Cantor)

Grain size (μm)

Yield stress (MPa)

Matrix (CrFeCoNi)

Filament (Cu-rich)

Matrix (CrFeCoNi)

Filament (Cu-rich)

HEA (Predicted)

HEA (Experimental)

1.13 1.02 20.77

1.37 1.24

608.1 634.2 220 (single phase)

127.5 132.8

512 484 220

506 477 218

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Materials Science & Engineering A 781 (2020) 139241

strain energy of Cu-rich filaments due to continuous recrystallization during deformation processing. 4) Because of the stability of Cu-rich filaments, phase boundaries act as barriers to grain growth and the grain sizes of Cr–Fe–Co–Ni matrix and Cu-rich filaments are controlled by the distribution and thick­ ness of Cu-rich phase. The grain size in Cr–Fe–Co–Ni HEA matrix and Cu-rich filament was observed to be 1.0–1.4 μm. 5) The yield strengths of CoCrFeCuNi (506 MPa) and CoCrFeCu1.71Ni (477Mpa) were observed to be greater than that of the Cantor alloy (220 MPa) with minor reduction of elongation. The strengthening of CoCrFeCuNi HEAs is associated with grain size refinement and the enhanced back stress by the phase boundaries between Cu-rich fil­ aments and Cu-lean matrix with higher strength and modulus.

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Data availability The data in this study are available from the corresponding author upon reasonable request. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. CRediT authorship contribution statement Yong Keun Kim: Formal analysis, Investigation, Writing - original draft. Byung Ju Lee: Formal analysis, Investigation, Resources. SoonKu Hong: Formal analysis, Data curation, Validation. Sun Ig Hong: Conceptualization, Methodology, Writing - review & editing, Supervision. Acknowledgements This research was supported by the Future Material Discovery Pro­ gram of the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT and Future Planning (MSIP) of Korea (2016M3D1A1023532) and High Value-Added Metallic Materials Specialist Training Program through the Korea Institute for Advance­ ment of Technology (KIAT) funded by the Korea Ministry of Trade, In­ dustry and Energy (G02P00720002001). Appendix A. Supplementary data Supplementary data to this article can be found online at https://doi. org/10.1016/j.msea.2020.139241. References [1] B. Cantor, I.T.H. Chang, P. Knight, A.J.B. Vincent, Microstructural development in equiatomic multicomponent alloys, Mater. Sci. Eng. A375–377 (2004) 213–218. [2] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.H. Tsau, S.Y. Chang, Nanostructured high-entropy alloys with multiple principle elements: novel alloy design concepts and outcomes, Adv. Eng. Mater. 6 (2004) 299–303. [3] D.B. Miracle, O.N. Senkov, A critical review of high entropy alloys and related concepts, Acta Mater. 122 (2017) 448–511. [4] Z. Li, S. Zhao, R.O. Ritchie, M.A. Meyers, Mechanical properties of high-entropy alloys with emphasis on face-centered cubic alloys, Prog. Mater. Sci. 102 (2019) 296–345. [5] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Microstructures and properties of high-entropy alloys, Prog. Mater. Sci. 61 (2014) 1–93. [6] G. Laplanche, A. Kostka, O.M. Horst, G. Eggeler, E.P. George, Microstructure evolution and critical stress for twinning in the CrMnFeCoNi high-entropy alloy, Acta Mater. 118 (2016). [7] J. Moon, S.I. Hong, J.B. Seol, J.W. Bae, J.M. Park, H.S. Kim, Strate rate sensitivity of high entropy alloys and its significance in deformation, Mater. Res. Lett. 7 (2019) 503–509.

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