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Variations of the Elastic Properties of the CoCrFeMnNi High Entropy Alloy Deformed by Groove Cold Rolling Paul Lohmuller 1,2,3 , Laurent Peltier 1,3 , Alain Hazotte 1,3 , Julien Zollinger 2,3 Pascal Laheurte 1,3 and Eric Fleury 1,3, * 1

2 3

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ID

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Laboratoire d’Etude des Microstructures et de Mécanique des Matériaux, UMR 7239 CNRS, Arts et Métiers ParisTech Campus de Metz, Université de Lorraine, 7 rue Félix Savart, F-57073 Metz, CEDEX 03, France; [email protected] (P.L.); [email protected] (L.P.); [email protected] (A.H.); [email protected] (P.L.) Institut Jean Lamour, UMR 7198 CNRS, Université de Lorraine, Campus ARTEM, F-54011 Nancy, CEDEX 01, France; [email protected] LABoratory of EXcellence Design of Alloy Metals for low-mAss Structures, Université de Lorraine, F-54011 Nancy, CEDEX 01, France Correspondence: [email protected]; Tel.: +33-372-747-772





Received: 29 June 2018; Accepted: 27 July 2018; Published: 2 August 2018

Abstract: The variations of the mechanical properties of the CoCrFeMnNi high entropy alloy (HEA) during groove cold rolling process were investigated with the aim of understanding their correlation relationships with the crystallographic texture. Our study revealed divergences in the variations of the microhardness and yield strength measured from samples deformed by groove cold rolling and conventional cold rolling processes. The crystallographic texture analyzed by electron back scattered diffraction (EBSD) revealed a hybrid texture between those obtained by conventional rolling and drawing processes. Though the groove cold rolling process induced a marked strengthening effect in the CoCrFeMnNi HEA, the mechanical properties were also characterized by an unusual decrease of the Young’s modulus as the applied groove cold rolled deformation increased up to about 0.5 before reaching a stabilized value. This decrease of the Young’s modulus was attributed to the increased density of mobile dislocations induced by work hardening during groove cold rolling processing. Keywords: high entropy alloy; crystallographic texture; groove rolling; elastic properties

1. Introduction High entropy alloys (HEAs), a new class of metallic materials in which the stability of the solid solution is explained by contribution of the configurational entropy, have recently received particular attention, owing to the prospect of developing new systems with tailored properties particularly suitable for a large range of different applications from dies and molding materials to corrosion-resistant coatings in manufacturing, energy, transport, and aeronautical industries [1–4]. The stabilization of the solid solution is also associated to special features such as low diffusion kinematics, lattice distortion and “cocktail effect” whose effects have been discussed in the literature [1,3,5]. The equi-atomic CoCrFeMnNi composition developed by Cantor and colleagues [6] is currently one of the most studied HEA. Though its mechanical properties are not outstanding [7–9], this alloy characterized by a single solid solution of face centered cubic structure has appeared attractive as a model material to gain a better understanding on characteristics particular to solid solution HEA Materials 2018, 11, 1337; doi:10.3390/ma11081337

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alloys against those of conventional solid solution alloys. Some key studies, such as those published by Wu et al. [10] and Laurent-Brocq and co-workers [11,12], already helped in the understanding the formation of these solid solutions, particularly the influence of each constituent element on the stability of the microstructure. The mechanical properties of HEAs, as those of conventional metallic materials, are also controlled by the microstructure. If several studies have revealed the dependency of the mechanical properties with the grain size [7,13,14], a few have recently reported the variation of the crystallographic texture during mechanical processing of HEAs. For example, conventional cold rolling [15–18], rotary swaging [19], or severe plastic deformation [20] have been investigated and the results of these studies show characteristics very similar to those obtained with low stacking fault energy (SFE) materials such as brass alloys [21,22]. Examinations of the deformation mechanisms have demonstrated that deformation occurred mainly by dislocation glide and sub-cell creation [8,18]. However, under particular conditions, such as tests performed under high strain level or at cryogenic temperature [16,18,23,24], the CoCrFeMnNi also presented deformation twinning, which can be seen as a secondary deformation mechanism. Though most of studies are concerned about the mechanical properties, the modifications of the Young’s modulus induced by work hardening have not been intensively studied. However, as seen earlier in the 19500 s, in their theory of work hardening, both Mott [25] and Friedel [26] took into account a network of dislocations in face-centered cubic crystals to explain the changes in the elastic modulus upon work hardening. More recently, others authors have proposed to account for the increase of internal stresses during deformation, the development of crystallographic texture and change in the dislocation arrangement [27] to describe the variation of solid solution alloys. This study was thus undertaken to investigate the evolution of the hardness and tensile properties of CoCrFeMnNi HEA alloy with the increasing area reduction achieved by grooved cold rolling. Though this process resulted in a marked strengthening effect, it was also accompanied by a reduction of the Young’s modulus. The variation of the crystallographic texture with the increasing groove rolling deformation was found to be slightly different from that obtained during conventional cold rolling. This study attempted to provide elements enabling the understanding of the observed changes in the mechanical properties. 2. Materials and Methods 2.1. Samples Production A 5 kg ingot of the equi-atomic CoCrFeMnNi alloy of composition Co20 .1 Cr19 .9 Fe20 .6 Mn19 .5 Ni19 .9 (in at.%) as analyzed by electron dispersive spectroscopy (EDS) was produced by vacuum arc melting at KIST, Korea, using initial elements with purity above 99.9%. Hot forging treatment and hot rolling at 1000 ◦ C were performed to homogenize the microstructure and to reduce the initial ingot thickness down to 12 mm. After another thickness reduction by conventional cold rolling, 4 samples of dimension 9.3 × 9.3 × 80 mm3 were machined perpendicularly to the rolling direction. Finally an annealing heat treatment was then applied during 3600 s at 900 ◦ C under a protective Argon atmosphere. Following this heat treatment, samples were considered, for this study, in a reference state in term of microstructure and mechanicals properties. Room temperature groove cold rolling process was performed on the annealed square bars. In this process, rollers present V-grooves of different dimensions as illustrated in Figure 1A. These grooves result in a bidirectional loading on the square sections of the samples, which prevents transversal displacement. This transversal displacement, though moderate, is inherent to the conventional rolling process (Figure 1B). From this viewpoint, groove cold rolling may be seen as a hybrid process between conventional rolling and swaging or extrusion. The samples were respectively rolled down to 14% and 33% area reduction ratios, i.e., rolling deformation ε = 0.2 and ε = 0.4, respectively. The obtained square wire-type samples were cut off in order to investigate the tensile properties (minimal length about 90 mm). The rests of the

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samples were rolled to successively up to 89% area reduction ratio (i.e., rolling deformation ε = 2.2), corresponding to a final section of about 3.1 × 3.1 mm2 . For each deformation level, a wire was Materials 2018, 11, x FOR PEER REVIEW 3 of 14 extracted with a minimal length of about 120 mm for tensile test and microhardness measurements. In order to determine the tensile properties of the initial state, the last sample was rolled until 38% area In order to determine the tensile properties of the initial state, the last sample was rolled until 38% reduction ratio (i.e., rolling deformation ε = 0.7) and annealed under the same conditions as described area reduction ratio (i.e., rolling deformation ε = 0.7) and annealed under the same conditions as above. Table the thickness of each sample, the associated reduction state calculated by  1 summarizes 1 summarizes described above. Table the thickness of each sample, the associated reduction state 2 2 ϕ = 100 × 1 − e f /ei and the true obtained by ε = 2 × ln(e /e ) [28], where e and ⁄ deformation calculated by = 100 × (1 − ) and the true deformation obtainedi byf = 2 × ln ( i ⁄ ) e[28], correspond to the initial and final thicknesses, respectively. f where ei and ef correspond to the initial and final thicknesses, respectively.

Figure 1. Grooved cold rolling (A) versus conventional rolling (B), the red arrows correspond to the Figure 1. Grooved cold rolling (A) versus conventional rolling (B), the red arrows correspond to the in in plane displacement direction induced by the process. plane displacement direction induced by the process. Table 1. Different grooved cold rolling deformation level produced for microhardness and tensile Table 1. Different grooved cold rolling deformation level produced for microhardness and tensile test, test, grayed background correspond to samples used for electron back scattered diffraction (EBSD) grayed background correspond to samples used for electron back scattered diffraction (EBSD) analyses. analyses. SectionReduction Reduction (%) (%) True True Deformation f (mm) Section i (mm) ef e ei e(mm) (mm) Deformation 9.3 9.39.3 0.0 0.0 9.3 0.0 0.0 8.6 14.5 0.2 8.6 14.5 0.2 33.2 0.4 -7.67.6 33.2 0.4 7.3 38.4 0.5 -7.36.6 38.4 0.5 49.6 0.7 -6.65.4 49.6 0.7 66.3 1.1 79.6 1.6 -5.44.2 66.3 1.1 3.1 88.9 2.2 4.2 79.6 1.6 3.1 88.9 2.2 2.2. Mechanical Testing 2.2. Mechanical Testingtests were carried out on 100 kN universal Zwick tensile test machine, using a Uniaxial tensile

50 kN load cell. Withtests the aim determining the Young’s modulus, were first loaded Uniaxial tensile wereofcarried out on 100 kN universal Zwicksamples tensile test machine, usingtoa1% 50 strain, charge/discharge cycles were performed. Strains were were measured using to an1% Epsylon kN loadthen cell. 5With the aim of determining the Young’s modulus, samples first loaded strain, extensometer with a gage length of performed. 25 mm, andStrains the samples were deformed atEpsylon a constant strain rate then 5 charge/discharge cycles were were measured using an extensometer −3 s−1 . Values of the Young’s modulus were determined as the mean of the initial slope of −3 of 1 × 10 with a gage length of 25 mm, and the samples were deformed at a constant strain rate of 1 × 10 each s−1. discharging of modulus the curves. Values of the yield strength fromofthe conventionally Values of theportion Young’s were determined as the mean were of theobtained initial slope each discharging 0.2% of plastic deformation. obtained shown in thethe right-hand side inset inof Figure 2. portion of the curves. Values A of typical the yield strengthcurve wereis obtained from conventionally 0.2% plastic For microstructural observations, samples polished side (along the direction) using deformation. A typical obtained curve is shown in were the right-hand inset inrolling Figure 2. SiC papers with grit sizeobservations, about 80 in order to reduce the thickness reachdirection) approximately the For microstructural samples were polished (alongand the to rolling using SiC center part of the samples. Then surfacing polishing was performed with grit size from 1000 to 4000. papers with grit size about 80 in order to reduce the thickness and to reach approximately the center Hardness values wereThen measured by micro-indentation Vickers tests carried Zwick part of the samples. surfacing polishing was performed with grit out sizeusing froma 1000 toRowel 4000. machine a load 0.3 gf, leading to indentation size of about 35 to 70out µm. Distance between Hardnesswith values wereofmeasured by micro-indentation Vickers tests carried using a Zwick Rowel adjacent at least 200 to µm. For each sample, hardness value averaged from machine measurements with a load of was 0.3 gf, leading indentation size of the about 35 to 70 µm.was Distance between 15 measurements, and the errors bars correspond to the measured standardvalue deviation. adjacent measurements was at least 200 µm. For each sample, the hardness was averaged from 15 measurements, and the errors bars correspond to the measured standard deviation.

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Figure procedure and and characterization. characterization. Figure 2. 2. Graphical Graphical representation representation of of samples samples production production procedure

2.3. Microstructural Analysis 2.3. Microstructural Analysis Before tensile testing, pieces of rolled materials were cut off for microstructural analyses. Before tensile testing, pieces of rolled materials were cut off for microstructural analyses. Samples were polished using SiC paper with grit size down to 4000 followed by a Al2O3 polishing Samples were polished using SiC paper with grit size down to 4000 followed by a Al2 O3 polishing with final particle size of 0.1 µm. Microstructural analyses were carried out on SEM Jeol 6490, at 20 with final particle size of 0.1 µm. Microstructural analyses were carried out on SEM Jeol 6490, at 20 kV kV with a working distance of 21 mm. Electron back scattered diffraction (EBSD) analyses were with a working distance of 21 mm. Electron back scattered diffraction (EBSD) analyses were performed performed by Oxford instrument detector with a minimal number of detected Kikuchi bands equal by Oxford instrument detector with a minimal number of detected Kikuchi bands equal to 8 and a step to 8 and a step size of 0.4 µm. EBSD were studied in four cases: after annealing and after section size of 0.4 µm. EBSD were studied in four cases: after annealing and after section reduction of ε = 0.5, reduction of ε = 0.5, 1.1 and 2.2. EBSD maps, Figure pole (FP), inverse figure pole (IFP) and orientation 1.1 and 2.2. EBSD maps, Figure pole (FP), inverse figure pole (IFP) and orientation density function density function (ODF) were obtained using ATEX software developed at LEM3 [29]. Texture (ODF) were obtained using ATEX software developed at LEM3 [29]. Texture components were also components were also performed using a spread of 15° from their ideal orientation to compute their performed using a spread of 15◦ from their ideal orientation to compute their volume fractions. All the volume fractions. All the obtained texture components were summed in order to determine the obtained texture components were summed in order to determine the randomly oriented components randomly oriented components by evaluating the difference between this sum and the totality [16,18]. by evaluating the difference between this sum and the totality [16,18]. The different analyzed samples The different analyzed samples are indicated in Table 1. The whole procedure from hot forging to are indicated in Table 1. The whole procedure from hot forging to EBSD analysis is summarized in EBSD analysis is summarized in Figure 2. Figure 2. 3. Results 3. Results 3.1. Microhardness and Tensile Results Tensile Results Figure 3a shows the variations of the microhardness values and associated standard deviations the different differentdeformation deformationstates statesapplied appliedbyby groove cold rolling data obtained by Otto al. for the groove cold rolling andand data obtained by Otto et al.et[15] [15]CoCrFeMnNi on CoCrFeMnNi processed by conventional cold rolling haveadded been added for comparison. on HEAHEA processed by conventional cold rolling have been for comparison. In the In the as-annealed state, theofvalue of the microhardness HVis 0.3,in which is in good as-annealed state, the value the microhardness was aboutwas 165about HV0.3 ,165 which good agreement agreement valuesinobtained in articles published articles underannealing the same annealing The with valueswith obtained published under the same treatment treatment [8,23]. The[8,23]. variation variation of the microhardness can be described by two different linearwith curves with a breakdown at of the microhardness can be described by two different linear curves a breakdown at a level a level of deformation of about 0.5. From the initial annealing state to a condition corresponding to of deformation of about 0.5. From the initial annealing state to a condition corresponding to the the groove rolling deformation ε ~the 0.5,microhardness the microhardness displayed a marked increased from 165 groove rolling deformation ε ~0.5, displayed a marked increased from 165 HV 0.3 2) than those 2 )kgf/mm HVabout 0.3 to 310 about HV 0.3 , characterized by a larger value of the slope (284 to HV310 , characterized by a larger value of the slope (284 kgf/mm than those reported 0.3 2) represented by the green dashed line in the Figure 3a. For reported et kgf/mm al. [8] (922 )kgf/mm by Otto etby al.Otto [8] (92 represented by the green dashed line in the Figure 3a. For values of values of the groove rolling deformation beyond 0.5, the microhardness continued to increase the groove rolling deformation beyond 0.5, the microhardness continued to increase however with however with a reduced slope until reaching a value of 380 HV 0.3 for a groove rolling deformation a reduced slope until reaching a value of 380 HV0.3 for a groove rolling deformation ε = 2.2. It isε = 2.2. It is remarkable that the microhardness valuesthe followed the same groove remarkable that the microhardness values followed same trend undertrend bothunder grooveboth rolled and rolled and conventional rolling processes, though the initial slope was larger and that the breakdown in the slope occurred earlier for samples deformed by groove rolled process.

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the initial slope was larger and that the breakdown in the 5slope of 14 occurred earlier for samples deformed by groove rolled process.

(a)

(b)

Figure 3. (a) Microhardness Vickers (HV0.3) evolution with the surface reduction, for this study Figure 3. (a) Microhardness Vickers (HV0.3 ) evolution with the surface reduction, for this study (orange (orange triangle) and comparison with results published by Otto et al. [15] (green dot), (b) yield triangle) and comparison with results published by Otto et al. [15] (green dot), (b) yield strength strength evolution during grooved cold rolling process. evolution during grooved cold rolling process.

The evolution of the yield strength, as shown in Figure 3b, followed a similar trend to that of the The evolution of the yield strength, as shown in in Figure 3b, followed a similar trend toagreement that of the microhardness. Values of the yield strength obtained the annealed condition are in good microhardness. Values of the yield strength obtained in the annealed condition are in good agreement with those reported in references [7,8]. The high values of the yield strength obtained after large area with thosewere reported in references The high values of the reported yield strength obtained large area reduction also found to be in[7,8]. accordance with the values by Sun et al. forafter CoCrFeMnNi reduction were also found to be in accordance with the values reported by Sun et al. for CoCrFeMnNi HEA samples with fine-grained microstructure [7]. HEAAs samples with fine-grained microstructure mentioned in the introduction, a part of[7]. this study focused on understanding the evolution mentioned in theas introduction, part of thisrolling study deformation. focused on understanding the evolution of theAsYoung’s modulus function of athe groove Figure 4 indicates that the of the Young’s modulus as function of the groove rolling deformation. Figure 4 indicates that the value value obtained in the as-annealed sample, i.e., 210 GPa, is also in agreement with those reported in obtained in the as-annealed sample, i.e., 210 GPa, is also in agreement with those reported in the the literature [9,30]. It also shows that the value decreased down to about 175 GPa, as the groove literature [9,30]. It also shows0.5, thatwhich the value decreased down about 175 GPa, as thedecrease groove rolling rolling deformation reached is corresponding to to approximately a 15% in the deformation reached 0.5, which is corresponding to approximately a 15% decrease in theHEA. Young’s Young’s modulus. As far as we know, no similar results have to be reported on deformed To modulus. As far as we know, no similar results have to be reported on deformed HEA. To the authors’ the authors’ knowledge, only a few studies have attempted to correlate the variations of the Young’s knowledge, a few studies havewith attempted to correlate the plastic variations of the Young’s modulus modulus of only engineering materials the increase of cold deformation [27,31,32]. As of engineering materials with the increase of cold plastic deformation [27,31,32]. As mentioned by mentioned by Benito et al. [27], four possible contributions have been proposed to explain the Benito et in al. Young’s [27], fourmodulus: possible contributions been proposed to explain the decrease Young’s decrease (i) debondinghave between a matrix and a second-phase or in inclusions modulus: (i) debonding between a matrix and a second-phase or inclusions [33,34]; (ii) increase the [33,34]; (ii) increase in the internal stresses during plastic deformation [35]; (iii) introductioninand internal stresses during plastic deformation [35]; (iii) introduction and development of texture during development of texture during deformation [27,31]; and (iv) change in dislocation arrangement and deformation [27,31]; and (iv) change in dislocation arrangement and kinematics [27,31]. kinematics [27,31]. By considering the single-phase single-phase character character of of CoCrFeMnNi CoCrFeMnNi HEA, HEA, the the first first mentioned mentioned reason reason By considering the appeared to to be In their their study, study, Benito Benito et have suggested suggested that that the the internal internal stresses stresses appeared be irrelevant. irrelevant. In et al. al. have generated during the early stage of cold rolling were not sufficient to describe the variation of the the generated during the early stage of cold rolling were not sufficient to describe the variation of Young’s modulus [27]. Consequently, the decrease in the Young’s modulus during the early stage Young’s modulus [27]. Consequently, the decrease in the Young’s modulus during the early stage of of groove cold rolling could result fromeither eithera acrystallographic crystallographictexture textureeffect effector oraa change change in in the the groove cold rolling could result from dislocation arrangement and deformation mechanism. dislocation arrangement and deformation mechanism. The different different mechanisms mechanisms in in the the CoCrFeMnNi CoCrFeMnNi HEA HEA are are now now well well documented documented in in the the literature literature The for different different deformation deformation modes modes and and different different degrees degrees of From these these for of deformation deformation [7,8,18,23,36,37]. [7,8,18,23,36,37]. From studies, it appears that the deformation mechanisms are governed by dislocation glide upon low strain studies, it appears that the deformation mechanisms are governed by dislocation glide upon low strain value, and by dislocation glide combined with twinning upon high level of plastic deformation. In an attempt to elucidate the decrease of the Young’s Modulus, a thorough study of the evolution of the crystalline texture as a function of the section reduction has thus been undertaken.

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value, and by dislocation glide combined with twinning upon high level of plastic deformation. In an attempt to elucidate the decrease of the Young’s Modulus, a thorough study of the evolution of the crystalline texture as PEER a function of the section reduction has thus been undertaken. Materials 2018, 11, x FOR REVIEW 6 of 14

Figure4.4.Young’s Young’s modulus variation the thickness reduction during cold grooved cold rolling Figure modulus variation withwith the thickness reduction during grooved rolling (maximal (maximal incertitude ± 6 GPa). incertitude ± 6 GPa).

3.2. Texture TextureAnalysis Analysis 3.2. Figure 55 shows maps obtained for the state and samples after three different Figure showsthe theEBSD EBSD maps obtained for annealed the annealed statefor and for samples after three values of the area reduction, i.e., ε = 0.5, 1.1, and 2.2. The initial state microstructure consisted of different values of the area reduction, i.e., ε = 0.5, 1.1, and 2.2. The initial state microstructure equiaxed grains with a mean grain diameter about 27 µm (as computed by ATEX). Figure 5b,d show consisted of equiaxed grains with a mean grain diameter about 27 µm (as computed by ATEX). the evolution of the microstructure increasing strain induced by thestrain groove cold rolling process. Figure 5b,d show the evolution of with the microstructure with increasing induced by the grooveIt can rolling be observed that grains were progressively along the elongated rolling direction (RD) and cold process. It can be observed that grainselongated were progressively along the rolling compressed in the normal direction (ND). For samples processed under low deformation level (Figure direction (RD) and compressed in the normal direction (ND). For samples processed under low 5b), some deformation features oriented at 45° to features the rolling direction appeared their fraction ◦ to the and deformation level (Figure 5b), some deformation oriented at 45 rolling direction progressively with the increasing level of groove rolling deformation (Figure 5c,d). Figure 5e appeared andincreased their fraction progressively increased with the increasing level of groove rolling shows a focused area of the Figure 5d obtained with a step size of 0.05 µm and a high indexation rate. deformation (Figure 5c,d). Figure 5e shows a focused area of the Figure 5d obtained with a step size These features from different deformation operating CoCrFeMnNi HEA that of 0.05 µm and aresulted high indexation rate. These featuresmechanisms resulted from differentin deformation mechanisms are dislocation glide and twinning as already described in different studies [7,8,19,36]. operating in CoCrFeMnNi HEA that are dislocation glide and twinning as already described in Figure 6 shows the evolution of the ODF section intensity and IFP intensity along the rolling different studies [7,8,19,36]. direction. map of the sample in showsand a slight recrystallization texture of FigureThe 6 shows the evolution ofthe theas-annealed ODF sectionstate intensity IFP intensity along the rolling the cube on face Cb {001} type as well as a partial α-fiber ( along ND). Such texture has direction. The map of the sample in the as-annealed state shows a slight recrystallization texture been for CoCrFeMnNi conventional cold rolling hotSuch forging prior to of the reported cube on face Cb {001} HEA type deformed as well as abypartial α-fiber ( along or ND). texture has an annealing at 900 °CHEA during 1 h [11,18]. One may also note that recrystallized been reported treatment for CoCrFeMnNi deformed by conventional cold rolling or hot forgingtwinning prior to occurred during the annealing treatment as reported in [16,18]. ◦ an annealing treatment at 900 C during 1 h [11,18]. One may also note that recrystallized twinning The during Figure 6, thetreatment IFP of theas deformed indicates a strong increase of the fiber occurred theshowing annealing reportedstate, in [16,18]. texture along RD as the level of deformation induced by groove cold rolling increased. Thisoffiber The Figure 6, showing the IFP of the deformed state, indicates a strong increase the texture is associated to the fiber, which is a minor component of the texture, as indicated by local maximal fiber texture along RD as the level of deformation induced by groove cold rolling increased. This fiber intensity. Both fiber textures are frequently observed for FCC materials [38,39], and have also been texture is associated to the fiber, which is a minor component of the texture, as indicated reported for CoCrFeMnNi HEA [19] deformed to swaging or drawing deformation. These observations by local maximal intensity. Both fiber textures are frequently observed for FCC materials [38,39], based onalso IFP been may reported be balanced with the analyses of the ODF sections. In the map, the α-fiber and have for CoCrFeMnNi HEA [19] deformed to swaging or ODF drawing deformation. component ( along ND, in the IFP (Figure is represented red sections. dashed line. The These observations based onnot IFPshown may be balanced with6)the analyses of by thethe ODF In the and directions along RD are highlighted by, respectively, the (Copper (Cu), A) and (Cube ODF map, the α-fiber component ( along ND, not shown in the IFP (Figure 6) is represented (Cb), (G)) texture in the ODF directions with φ2 = 45° (referred in highlighted Table 2) [40].by, respectively, by theGoss red dashed line.components The and along RD are

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Figure 5. EBSD-band contrast maps described by inverse figure pole (IFP) maps (upper right corner):

Figure 5. EBSD-band contrast maps described by inverse figure pole (IFP) maps (upper right corner): Figure 5. inverse maps (a) as-annealed; (b)contrast after ε = maps 0.5 by described groove coldby rolling; (c)figure after εpole = 1.1 (IFP) groove cold (upper rolling; right (d) ε =corner): 2.2 Figure 5. EBSD-band EBSD-band contrast maps described by inverse figure (IFP) maps (upper right corner): (a) as-annealed; (b) after εε = 0.5 by groove cold rolling; (c) after εεpole == 1.1 groove cold rolling; (d) εε == 2.2 (a) as-annealed; (b) after = 0.5 by groove cold rolling; (c) after 1.1 groove cold rolling; (d) 2.2 (a) as-annealed; (b) after ε = 0.5 by groove cold rolling; (c) after ε = 1.1 groove cold rolling; (d) εεall== 2.2 groove cold rolling; and (e) focused area of the image (d). Note that magnification is the same for (a) as-annealed; (b) after ε = 0.5 by groove cold rolling; (c) after ε = 1.1 groove cold rolling; (d) 2.2 groove cold rolling; and (e) focused area of the image (d). Note that magnification is the same for all groove cold rolling; and (e) focused area of the image (d). Note that magnification is the same for all maps (maps (d) corresponding to aarea stackof 2 EBSD maps a step of about 0.2is µm). groove cold rolling; and image (d). Note that size magnification the same groove cold rolling; and (e) (e) focused focused area ofof2the the image (d).with Note is the same for for all all maps (maps (d) corresponding to a stack of EBSD maps with athat stepmagnification size of about 0.2 µm). maps maps (maps (maps (d) (d) corresponding corresponding to to aa stack stack of of 22 EBSD EBSD maps maps with with aa step step size size of of about about 0.2 0.2 µm). µm). Table 2. Definition of the texture components in the alloy.

Table 2. Definition of the texture components in the alloy. 2. of components alloy. TextureTable Component Symbol Euler Angle φ1, Φ, in φ2the Table 2. Definition Definition of the the texture texture components in theMiller alloy. Indices

Cb (Cube) Texture Component Symbol Texture Component Symbol Texture Component Symbol  G (Goss) Cb (Cube) CbCb (Cube) (Cube) B (Brass)  G ⨂ GG/B (Goss) G (Goss) (Goss) (Goss/Brass) B (Brass) B (Brass) B (Brass) BCu (Brass) (Copper) N ⨂ G/B (Goss/Brass) G/B (Goss/Brass) ⨂ G/B (Goss/Brass) A (Copper) CuCu (Copper) F Cu (Copper) A A A P A at each analyzed ction (ODF) section at φ2 = 0°, 45° and 63° (right)  F F FCuT ective IFP along rolling direction, red arrows indicate maximal P S P P 001}. Horizontal and vertical bars, respectively, P correspond to ODF

 CuT    ed dashed line represent the α-fiber. The CuT symbols CuT for each texture S , G: , B: , G/B: ⨂, Cu: , A: , F: S, P: S , CuT: , S:

roove rolling deformation ε = 0.5, a slight partial α-fiber appeared nd it increased till ε = 1.1. As the deformation level increased, this her deformation states as indicated by the ODF with φ2 = 45°, in ss (RtG) texture components were not detected. The S component

0, 0, 0φϕ 1 , Φ,φ 2 Euler Angle Euler Angle Euler Angle φ111,,1Φ, Φ, φϕ222 2 0, 45, 0 0, 0, 0, 0,45, 0, 0,000 0 35, 0, 45, 00 0 0, 0, 90, 45,45, 74, 45 35, 45, 0 35, 45, 35,35, 45,45, 00 90, 35, 450 74, 90, 45 74,74, 90,90, 4545 35, 90, 45 90, 35, 45 35, 45 30/90, 55, 45 90,90, 35, 45 35,35, 90,90, 4545 35, 90, 45 30, 90, 45 35, 90, 45 30/90, 45 90, 74,55, 45 30/90, 55, 30/90, 55, 4545 30, 90, 45 59, 37, 63 30, 90, 45 30, 90, 30, 90, 4545 90, 74, 90,90, 74,74,45 4545 59, 37, 63 59,59, 37,37,6363

{001} Miller Indices Miller Indices Miller Indices {110} {001} {001} {001} {110} {110} {110} {110} {110} {110} {110} {110} {110} {112} {110} {110} {110} {110} {112} {112} {111} {112} {110} {110} {110} {011} {110} {111} {552} {111} {111} {011} {123} {011} {011} {011} {552} {552} {552} {123} {123} {123}

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Figure 5. EBSD-band contrast maps described by inverse figure pole (IFP) maps (upper right as-annealed;contrast (b) aftermaps ε = 0.5 by groove rolling; (c)pole after(IFP) ε = 1.1 groove cold rolling; (d Figure(a) 5. EBSD-band described by cold inverse figure maps (upper right corner groove cold (e) area rolling; of the image (d). that(upper magnification is(d) theεsam as-annealed; (b)rolling; aftermaps ε =and 0.5 by focused groove (c)pole after ε Note = 1.1 groove cold rolling; = 2. Figure(a) 5. EBSD-band contrast described by cold inverse figure (IFP) maps right corner): (maps corresponding to arolling; of EBSD maps withmagnification a step of is about 0.2 µm). groovemaps cold (e) area ofstack the image (d).ε Note theεsame for a (a) as-annealed; (b)rolling; after ε(d) =and 0.5 by focused groove cold (c)2 after = 1.1 that groove cold size rolling; (d) = 2.2 (maps (d) corresponding to aofstack of 2 EBSD mapsthat withmagnification a step size of is about 0.2 µm). groovemaps cold rolling; and (e) focused area the image (d). Note the same for all Table 2. Definition of the texture components in the alloy. maps (maps (d) corresponding tofigure a stackpole of 2(IFP) EBSD maps with aright stepcorner): size of about 0.2 µm). Figure 5. EBSD-band contrast maps described by inverse maps (upper Table 2. Definition of the texture components in the alloy. (a) as-annealed; (b) after ε = 0.5 by groove cold rolling; (c) after ε = 1.1 groove cold rolling; (d) ε = 2.2 Figure 5. EBSD-band contrast maps described by inverse figureComponent pole (IFP) mapsSymbol (upper right corner): Texture Euler Angle φ1, Φ, φ2 Miller Indices Table 2. Definition of (upper the texture components in the alloy. groove cold (e) of the image (d). Note that magnification is Angle theεsame for allφ2 Miller Indices Figure(a) 5. EBSD-band contrast described by area inverse figure pole (IFP) maps right corner): as-annealed; (b)rolling; aftermaps ε and = 0.5 byfocused groove cold rolling; (c) after ε = 1.1 groove cold rolling; (d) = 2.2 Texture Component Symbol Euler φ 1 , Φ, Cb (Cube) 0, 0, 0 {001} corresponding to arolling; stack 2 EBSD withmagnification a step of about µm). (a) as-annealed; (b)(maps after ε(d) =and 0.5 by focused groove cold (c) after εmaps = 1.1 groove coldsize rolling; (d) = 2.2 groovemaps cold rolling; (e) area of the of image Note that isAngle theε0.2 same Texture Component Symbol Euler φ 1,for Φ,0all φ245,Miller Indices  Cb(d). (Cube) 0, 0, {001} G (Goss) 0, 0 {110} groovemaps cold rolling; and (e) focused area the image Note that magnification is the same (maps (d) corresponding to aofstack of 2 EBSD maps with a step size of about 0.2 µm). Cb(d). (Cube) 0,for 0, 0all G (Goss) 0, 45, 35, 0 45, 0{001} {110} B (Brass) {110} Table 2. Definition of the texture components in the alloy. maps (maps (d) corresponding to a stack of 2 EBSD maps with a step size of about 0.2 µm). Figure 6. Orientation density function (ODF) section at φ 2 = 0°, 45° and 63° (right) at each analyzed  ◦ ◦ G (Goss) 45, 35, 0 at45, B G/B (Brass) 0 90, {110} ⨂63◦0,(right) (Goss/Brass) 74, 45{110} {110} Figure 6. Orientation density function (ODF) section at ϕand =the 063° , 45 andarrows each analyzed Figure 6. Orientation density (ODF) section atcomponents φalong 2 = 0°, 45° (right) at eachindicate analyzed 2direction, deformation and their IFP rolling red maximal Tablestate 2.function Definition ofrespective the texture in alloy. Texture Component Symbol Euler Angle φ 1, Φ, φ 2 Miller Indices B (Brass) 35, 45, 0 {110} ⨂ G/B (Goss/Brass) 74, 90, 45 {110} Cu (Copper) 90, 35, 45 {112} state and their respective IFP along rolling direction, redrespectively, arrows indicate maximal intensity deformationdeformation state and2. their respective IFP rolling direction, redbars, arrows indicate maximal Table Definition of {111} the texture components in alloy. intensity for orientation andalong {001}. Horizontal andthe vertical correspond to ODF ⨂ G/B (Goss/Brass) 74, 90, 45 {110} Cb (Cube) 0, 0, 0 {001} Cu (Copper) 90, 35, 45 {112} 90, 45IFP {110} The Texture Component Symbol Angle φ 1A , bars, Φ,respectively, φrespectively, 2 Miller Indices for orientation {111} and {001}. Horizontal vertical correspond to 35, ODF and intensity for orientation {111} and {001}. Horizontal andand vertical bars, correspond to each ODF and IFP maximal intensity. The redEuler dashed line represent the α-fiber. symbols for texture  Texture Component Symbol Euler Angle φ 1A , Φ, φ 2 Miller Indices Cu (Copper) 90, 35, 45 {112} G (Goss) 0, 45, 0 {110} 35, 90, 45 {110} F 30/90, 55, 45 {111} Cb (Cube) 0, 0, 0 {001} maximal intensity. The red dashed line represent the α-fiber. The symbols for each texture component component respectively, Cb:line ,represent G: , B: the , G/B: ⨂, Cu:The , A: , F: for , P: , CuT: , S: , and IFP maximal intensity.are, The red dashed α-fiber. symbols N A 35, 90, 45 {110} Cb (Cube) 0, 0, 0 {001} B (Brass) 35, 45, 0 {110} F 30/90, 55, 45 {111} P 30, 90, 45 {011} as defined in Table 2. Cu: , A: , ,F:F: {110} CuT:, , S: S: , as defined G (Goss) Cb: 45,, 0 A: component are, are, respectively, respectively, Cb: ,, G: G:  ,, B: B: , G/B: G/B: ⨂0, ,, Cu: , ,P:P: , ,CuT:   ⨂ 0,F45, 35, 30/90, 55, 0 P 45, 74, {110} 30,45 90,90, 45 74, 45{111} {011} CuT {552} in(Goss) Table 2. (Goss/Brass) BG/B (Brass) 0 90, 45{110} {110} as defined G in Table 2. For the processed to a groove deformation ε = 0.5, a slight {110} 30,partial 90,90, 45α-fiber {011} B G/B (Brass) 35,Prolling 45,74, 0 90, Cusample (Copper) 90, 35, 45{110} {112} CuT 74,59, 45appeared {552} S 37, 63 {123} ⨂ (Goss/Brass) 45 For the sample processed to a groove rolling deformation = 0.5, a slight partial α-fiber appeared between the G and P components, and it increased till ε = 1.1. Asεthe deformation level increased, this  ⨂ G/B (Goss/Brass) 74, 90, 45 {110} CuT 90, 74, 45 {552} A 35, 90, 45 {110} S 59, 37, 63 {123} For the sampleCu processed to a groove rolling ε till = 0.5,=a 1.1. slight partial α-fiber appeared increased, (Copper) 90, 35, 45 {112} between the G and P “broken” components, anddeformation it increased As the fiber seemed to be for higher deformation statesε as indicated by deformation the ODF with level φ2 = 45°, in Cuand (Copper) 90,till 451.1. {112} S35, 59, 37, 63 {123} F and 30/90, 55,deformation 45 {111} it increased between the G P components, ε = As the level increased, this A 35, 90, 45 {110} this fiber to(B) beand “broken” higher states as indicated by the ODF with ϕ2 = 45◦ , whichseemed the Brass Rotatedfor Goss (RtG)deformation texture components were not detected. The S component A 35, 90, 45 {110} P 30, 90, 45 {011} fiber seemed to be “broken” for higher deformation states as indicated by the ODF with φ 2 = 45°, in F to (B) 30/90, 55,texture, 45 also seemed follow same trend the α-fiber with a {111} weaker decrease for The deformation in which the Brass andthe Rotated Gossas (RtG) texture components were not detected. S component and F 30/90, 55, 45 {111} CuTGoss (RtG) 90, 74, 45This {552} which the Brass (B) and Rotated texture components were not detected. The S component P 30, 90, 45 {011} up to 2.2 in comparison to the B RtG components. first features observed during the up also seemed to follow the same trend as the α-fiber texture, with a weaker decrease for deformation P the 30, 90, 45 74,with {011} S trend 59, 37,afound 63 {123} also seemed to follow same as the α-fiber texture, weaker decrease for deformation  CuT 90, 45 {552} conventional rolling process, with the α-fiber was to continuously increased as the to 2.2 in comparison to the B and RtG components. This first features observed during the conventional and CuT S to 90, 74, 45 {552} deformation increased, as it was reported by Sathiaraj et al. [16] and Haase et al. [18]. The fiber up to 2.2 rolling in comparison the B RtG components. This first features observed during the 59,continuously 37, 63 {123} process, with the α-fiber was found to increased as the deformation increased, (CuSand A) texture component displayed the with the grooved rolling deformation, 59,was 37,strongest 63found increase {123} conventional rolling process, with the α-fiber to continuously increased as and theA) texture as it was reported by Sathiaraj et al. [16] and Haase et al. [18]. The fiber (Cu which firstly confirms the existence of this fiber but may also indicate the presence offiber twinning deformation increased, as it was reported by Sathiaraj et al. [16] and Haase et al. [18]. The component displayed the strongest increase with themay grooved rolling deformation, which firstly deformation mechanism. Thisthe deformation by twinning also occur atrolling lower level of deformation (Cu and A)confirms texture component displayed strongest increase with the grooved deformation, the existence of this fiber butHaase mayetalso indicate of twinning in comparison to conventional rolling. al. have shownthe thatpresence for conventional rolling,deformation the Cu which firstly confirms This the deformation existence of by this fiber but may also indicate thelevel presence of twinning mechanism. twinning may also occur at lower of deformation in comparison texture component appeared after a rolling deformation ε = 1.6 [18]. In our study, for a deformation

deformation deformation alsothat occur lower level ofrolling, deformation to mechanism. conventionalThis rolling. Haase by et twinning al. have may shown foratconventional the Cu texture in comparison to conventional rolling. Haase et al. have shown that for conventional rolling, the Cu of about component appeared after a rolling deformation ε = 1.6 [18]. In our study, for a deformation texture component appeared after a rolling deformation ε = 1.6 [18]. In our study, for a deformation 1.1, this component appeared to be the major one in the ODF section and it was found to continuously

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Materials 2018, 11, x FOR PEER REVIEW 9 of 14 increase. However, the two components of the RD appeared to have different behavior: as the deformation by groove cold rolling increased, the Cb component tended to progressively disappear, of about 1.1, this component appeared to be the major one in the ODF section and it was found to while the G component slightly increased without been of affected by theappeared rupture to ofhave the αdifferent fiber. It is continuously increase. However, the two components the RD important to noteasthat G components are both appertaining to along and along behavior: thethe deformation by groove cold rolling increased, the Cb ND component tended toND progressively while the G be component slightly increased without been two affected the fiber texture, so thatdisappear, their variations may correlated with the evolution of these fiberby textures. rupture the αthe fiber. It is important to note that the Gofcomponents both appertaining Figure 7a of shows variation of the volume fraction each textureare component with thetoincreasing along ND and ND fiber texture, so that their variations be correlated thewith deformation by groove coldalong rolled process, the texture components beingmay obtained from thewith ODF evolution of these two fiber textures. ◦ a maximum deviation around the ideal orientation of 10 [18]. In addition to the grain elongation Figure 7a shows the variation of the volume fraction of each texture component with the as a function to the groove rolling strain (Table 3), a global tendency is a decrease of the random increasing deformation by groove cold rolled process, the texture components being obtained from orientation, which could be correlated to the formation of a strong crystallographic texture. It can the ODF with a maximum deviation around the ideal orientation of 10° [18]. In addition to the grain also beelongation observedasthat at leastto4the texture components fraction isofa8% after a a function groove rolling strainreached (Table 3),a avolume global tendency decrease ofgroove the rollingrandom deformation ε = 2.2, which can be directly correlated to the complex texture described above. orientation, which could be correlated to the formation of a strong crystallographic texture. Figure It7bcan shows the different volume fraction associated to fiber textures reported previously. The also be observed that at least 4 texture components reached a volume fraction of 8% after aaim rolling conventional deformation ε =cold 2.2, which can directly correlated to the texture described here isgroove to compare rolling tobe grooved cold rolling by complex summing different texture above.associated Figure 7b shows thefiber different volume fraction associated to fiberThe textures reported previously.was component to each texture at different area reduction. ND component Theby aim here is toG, compare conventional cold rolling to grooved rolling summing differentand obtained summing G/B, and B components, RD withcold Cu and A by texture component texture component associated to each fiber texture at different area reduction. The ND RD with Cb and G. One may firstly note that the divergence from the as-annealed state for the component was obtained by summing G, G/B, and B components, RD with Cu and A texture two deformation processes may be explained by the different annealing temperatures and by slightly component and RD with Cb and G. One may firstly note that the divergence from the asdifferent mechanical post-processing conditions applied to the HEA (see ref. [16]). As expected by annealed state for the two deformation processes may be explained by the different annealing the previous analysis, of themechanical volume fraction of each fiber texture are clearly temperatures andthe by evolutions slightly different post-processing conditions applied to the different HEA for the(see tworef. conventional and groove processes. reported by Sathiaraj al. [16], [16]). As expected by therolling previous analysis, As the previously evolutions of the volume fraction ofeteach ND the main fiber-texture conventional cold rolling, while theprocesses. two other fibers fiberistexture are clearly different by for the two conventional and groove rolling Asstudied previously by Sathiaraj et al.In[16], is the main conventional rolling,relatively while show areported continuous decrease. theND case of groove coldfiber-texture rolling, theby RD fibercold remained the from two other studiedstate fibers a continuous decrease. state. In the case groove rolling, the constant the initial upshow to the high deformation The of two othercold fibers conserved RD fiber remained relatively constant from the initial state up to the high deformation state. the same evolution until intermediate deformation state, with a simultaneous increase. For higher The two other conserved the shown same evolution intermediate deformation with a for deformation level, thefibers rupture of α-fiber in Figureuntil 6 (between ε = 1.1 and 2.2) isstate, responsible simultaneous increase. For higher deformation level, the rupture of α-fiber shown in Figure 6 the observed decrease of this fiber. In contrast, the RD fiber texture continued to evolve and it (between ε = 1.1 and 2.2) is responsible for the observed decrease of this fiber. In contrast, the RD became predominant over the other texture components. These final states are closely corresponding to fiber texture continued to evolve and it became predominant over the other texture components. the textures in FCC deformed bytoswaging or drawing as already published These found final states arematerials closely corresponding the textures found inprocesses, FCC materials deformed by by Otto et al. [15]. swaging or drawing processes, as already published by Otto et al. [15].

(a)

Volume fraction

(b)

Figure 7. (a) Volume fraction of the main principal different detected texture component and (b) fibertexture comparison between groove cold rolling (this work) and conventional cold rolling (ref. [16]).

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Table 3. Values of the length/width aspect ratio of the grains, r, and Young’s modulus, E, measured along the rolling direction as a function of the groove rolling strain, ε. ε

0

0.5

1.1

2.2

r E (GPa)

1.9 ± 0.7 262.2 ± 5.8

2.3 ± 0.7 263.0 ± 13.1

3.4 ± 1.0 262.8 ± 13.1

5.8 ± 2.0 279.0 ± 12.3

To summarize, the texture of CoCrFeMnNi HEA alloy by groove cold rolling may be described as a hybrid texture between the texture obtained by conventional cold rolling and that obtained by swaging or drawing processes. Under high level of deformation, the RD fiber tended to become predominant due to the decrease of the ND fiber. This final state appeared to be very similar to the texture of FCC materials deformed by swaging process. 4. Discussion This study firstly focused on the evolution of mechanical properties of the CrCoFeMnNi HEA during deformation by groove cold rolling. Our results revealed two stages in the evolution of the microhardness, yield strength, and Young’s modulus, with a change of slopes for a groove rolling deformation of about ε = 0.5. Both microhardness and yield strength displayed a strong increase with the strain up to ε = 0.5 by groove cold rolling followed by a moderate increase, to reach values of, respectively, around 400 HV and 1400 MPa for an applied groove rolling deformation ε equal to 2.2. In contrast, the Young’s modulus displayed firstly a decrease of about 15% from its initial value as the groove rolling deformation increased to about ε = 0.5, followed by a slight increase from 175 to 179 GPa as the groove rolling deformation was further increased till ε = 2.2. The texture analysis revealed that the as-annealed and deformed microstructures presented similar features to those reported in previous studies on CoCrFeMnNi HEA processed by conventional cold rolling. From these considerations, it has been assumed that the deformation mechanisms involved during groove cold rolling were the same as those already reported for CoCrFeMnNi HEA alloy, i.e., dislocation glide and twinning. In order to interpret the variation of the mechanical properties, identification of the main texture components and their evolutions tended to confirm the emergence of a specific crystallographic texture responsible for the change in the microhardness and yield strength as described above. It has been shown that these textures could be interpreted as a hybrid texture between conventional rolling and swaging texture. Also, under high deformation level, our analyses suggested that the swaging type texture tended to become predominant. The evolution of hardness and yield strength may be regarded through the evolution of the active deformation mechanism. Otto et al. [15], as well as Hasse et al. [18], had reported that plastic deformation was mainly governed by dislocation glide for reduction less than 50% (i.e., ε ~2), while both twinning and dislocation glide are active deformation mechanisms beyond this limit for CrCoFeMnNi HEA processed by conventional cold rolling. For alloy processed by groove rolling process, the similar variation in the properties suggested that the deformation mechanisms might also be operating, and that the breakdown occurred for lower value of rolling deformation owing to the biaxial state of stress applied during groove cold rolling. It might also indicate that the work hardening, and the change in the density of dislocation, in CoCrFeMnNi HEA may be more important when deformed under groove cold rolling, i.e., bi-axial compression, than during conventional cold rolling. The variation of the Young’s modulus followed a different pattern that might be more complicated and less intuitive at this current stage of our investigation. Owing to the strong anisotropy of FCC alloys, such as austenitic Fe and Ni alloys, values of the Young’s modulus are expected to be strongly dependent upon the crystallographic texture [41]. Since the elastic constants of the CoCrFeMnNi HEA have not been reported yet, the Young’s modulus was calculated from the fourth-order elastic compliance tensor Sijkl of the quaternary CoCrFeNi alloy [42]. Without any texture, these elastic compliances give values of 96.1, 232.9, and 442.8 GPa for the Young’s modulus calculated, respectively,

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along the , , and directions. Taking into account the crystallographic texture expressed by the ODF measured at the different levels of groove rolling deformation, the calculated values of the Young’s modulus have been reported in Table 3. Though the calculated values are larger than those reported in the literature and measured in this study, these results tend to suggest that the Young’s modulus do not change significantly for CoCrFeMnNi HEA with texture corresponding to groove rolling deformation varying between ε = 0 and 1.1. Beyond ε = 1.1, the crystallographic texture oriented toward the orientation , as seen in the IPF (Figure 6) and identified as the RD fiber-texture (Figure 7), induced a marked increase of the Young modulus. Consequently, from this point of view, the crystallographic texture cannot provide a satisfactory explanation for the change in the Young’s modulus detected in our study, which is in agreement with earlier studies [27] undertaken on conventional steels. As it has been advanced in the introduction, changes in the dislocation arrangement could be responsible for the variation of the elastic properties. Otto et al. [8] explained that the deformation was essentially the results of dislocations gliding at low deformation level and they reported the formation of dislocation cell-structures under higher degrees of deformation. Laplanche et al. [23] have also shown that the dislocation density increased continuously as the deformation increased. Joo et al. [36] have shown that twinning had an effect in term of dynamic Hall-Petch relationship but also led to an accumulation of stacking fault. In addition Benito et al. [27] have demonstrated a strong connection between dislocation structure developed during deformation and decrease in the Young’s modulus. By considering all of these features, and assuming that they are also operating in solid solution deformed by groove cold rolling, it seems reasonable to consider that the change in the Young’s modulus could be mainly governed by the structure of dislocation developed under level of deformation and the gliding of mobile (non-pinned) dislocations within the cells. The stabilized value of the Young’s modulus observed for larger value of the rolled deformation could be explained by modification in the dislocation arrangement resulting from the contribution of the twinning deformation. 5. Conclusions In this study, samples of CoCrFeMnNi HEA were successively deformed by groove cold rolling at different strain levels. The mechanical properties and crystallographic texture were evaluated as a function of the increasing cold rolling deformation, leading to the following conclusions: -

-

-

Both the microhardness and the yield strength exhibited similar evolution as the area reduction ratio applied by groove cold rolling increased. A first marked increase until a groove rolling deformation of about 0.5 was followed by a moderate increase under higher strain level. This two-step behavior has been attributed to change in the activated deformation mechanisms, and characterized by a transition from dislocation mechanism and dislocation + twinning deformation mechanism as reported by Otto et al. [8] and Haase et al. [18]. For the groove rolling process, these breakdowns occurred at lower levels of deformation in comparison to conventional cold rolling. The Young’s modulus also displayed a two-step variation as the groove cold rolled deformation increased. An approximately 15% decrease from 210 GPa to 175 GPa was observed between annealed HEA samples and those deformed up to about 0.5 by groove rolled process. For higher level of the groove cold rolling deformation, the Young’s modulus value stabilized around 175 GPa. This evolution of the elastic property, which could not be attributed by a modification of the crystallographic texture, has been explained by the transition of operating deformation mechanisms between gliding of non-pinned dislocations within dislocation cell in the first stage, and contribution of twinning for larger values of the groove rolling deformation. The texture analysis highlighted a complex crystallographic texture induced by groove cold rolling deformation. This texture has been described as a hybrid texture between the texture developed during conventional rolling and that from swaging or drawing process. This texture

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was found to consist in a combination of fiber texture along the RD, fiber texture along ND and ND’ and texture along RD. At higher strain level, the fiber texture tended to disappear for the benefit of fiber texture along the RD. This texture was found to be consistent with the bi-axial deformation induced by groove cold rolling but could not account for the decrease in the Young’s modulus. Author Contributions: Conceptualization, E.F., J.Z., and A.H.; Methodology, L.P. and P.L. (Pascal Laheurte); Investigation, P.L. (Paul Lohmuller); Data Curation, P.L. (Paul Lohmuller); Writing-Original Draft Preparation, P.L. (Paul Lohmuller); Writing-Review & Editing, E.F., J.Z., and A.H.; Funding Acquisition, E.F. and J.Z. Funding: The authors acknowledge the financial support from French National Research Agency ANR (LabEx DAMAS, Grant no.ANR-11-LABX-0008-01). Acknowledgments: The authors are thankful to Jin-Yoo Suh researcher at KIST, Seoul, Korea, who provided the ingot. A special thank is also addressed at Charbonnier and Lecoz, for their assistance and advices. Eric Fleury acknowledges the France-Korea Partenariat Hubert Curien (PHC) STAR program. Conflicts of Interest: The authors declare no conflict of interest.

References 1. 2.

3. 4. 5. 6. 7.

8.

9.

10.

11. 12. 13.

14.

Gao, M.C.; Yeh, J.; Liaw, P.K.; Zhang, Y. High-Entropy Alloys: Fundamentals and Applications; Springer Internationnal Publishing: New York, NY, USA, 2016. Yeh, J.W.; Chen, S.K.; Lin, S.J.; Gan, J.Y.; Chin, T.S.; Shun, T.; Tsau, C.H.; Chang, S.Y. Nanostructured highEntropy alloys with multiple principal elements: Novel alloy design concepts and outcomes. Adv. Eng. Mater. 2004, 6, 299–303. [CrossRef] Yeh, J.W. Alloy design strategies and future trends in high-Entropy alloys. JOM 2013, 65, 1759–1771. [CrossRef] Zhang, Y.; Yang, X.; Liaw, P.K. Alloy design and properties optimization of high-Entropy. JOM 2013, 63, 830–838. [CrossRef] Miracle, D.; Senkov, O. A critical review of high entropy alloys and related. Acta Mater. 2017, 122, 448–511. [CrossRef] Cantor, B.; Chang, I.; Knight, P.; Vincent, A. Microstructural development in equiatomic multicomponent alloys. Mater. Sci. Eng. A 2004, 375, 213–218. [CrossRef] Sun, S.; Tian, Y.; Lin, H.; Dong, X.; Wang, Y.; Zhang, Z.; Zhang, Z. Enhanced strength and ductility of bulk CoCrFeMnNi high entropy alloy having fully recrystallized ultrafine-grained structure. Mater. Des. 2017, 133, 122–127. [CrossRef] Otto, F.; Dlouhy, A.; Somsen, C.; Bei, H.; Eggeler, G.; George, E. The influences of temperature and microstructure on the tensile properties of a CoCrFeMnNi high-entropy alloy. Acta Mater. 2013, 61, 5743–5755. [CrossRef] Laplanche, G.; Gadaut, P.; Horst, O.; Otto, F.; Eggeler, G.; George, E.P. Temperature dependencies of the elastic young moduli and thermal expansion coefficient of an equiatomic, single-phase CoCrFeMnNi high-entropy alloy. J. Alloys Compd. 2015, 623, 348–353. [CrossRef] Wu, Z.; Bei, H.; Otto, F.; Pharr, G.; George, E. Recovery, recrystallization, grain growth and phase stability of a family of FCC-Structured multi-Component equiatomic solid solution alloys. Intermetallics 2014, 46, 131–140. [CrossRef] Bracq, G.; Laurent-Brocq, M.; Perrière, L.; Pirès, R.; Joubert, J. The fcc solid solution stability in the Co-Cr-Fe-Mn-Ni multi-Component high-Entropy alloy. Acta Mater. 2017, 128, 327–336. [CrossRef] Laurent-Brocq, M.; Akhatova, A.; Perrière, L.; Chebini, S.; Sauvage, X.; Leroy, E.; Champion, Y. Insights into the phase diagram of CrMnFeCoNi high entropy alloy. Acta Mater. 2015, 88, 355–365. [CrossRef] Bae, J.; Moon, J.; Jang, M.; Yim, D.; Kim, D.; Lee, S.; Kim, H. Trade-off between tensile property and formability by partial recrystallization of CrMnFeCoNi high-Entropy alloy. Mater. Sci. Eng. A 2017, 703, 324–330. [CrossRef] Pradeep, K.; Tasan, C.; Yao, M.; Deng, Y.; Springer, H.; Raabe, D. Non-Equiatomic high entropy alloys: Approach towards rapid alloy screening and property-Oriented design. Mater. Sci. Eng. A 2015, 648, 183–192. [CrossRef]

Materials 2018, 11, 1337

15.

16. 17.

18. 19. 20.

21.

22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33.

34. 35. 36. 37.

38.

13 of 14

Otto, F.; Hanold, N.; George, E. Microstructural evolution after thermomechanical processing in an equiatomic, single-Phase CoCrFeMnNi high-Entropy alloy with special focus on twin boundaries. Intermetallics 2014, 54, 39–48. [CrossRef] Dan Sathiaraj, G.; Bhattacharjee, P. Effect of cold-rolling strain on the evolution of annealing texture of equiatomic CoCrFeMnNi high entropy alloy. Mater. Charact. 2015, 109, 189–197. [CrossRef] Dan Sathiaraj, G.; Bhattacharjee, P. Effect of starting grain size on the evolution of microstructure and texture during thermo-mechanical processing of CoCrFeMnNi high entropy alloy. J. Alloys Compd. 2015, 647, 92–96. [CrossRef] Haase, C.; Barrales-Mora, L. Influence of deformation and annealing twinning on the microstructure and texture evolution of face-centered cubic high-entropy alloy. Acta Mater. 2018, 150, 88–103. [CrossRef] Laplanche, G.; Horst, O.; Otto, F.; Eggeler, G.; George, E. Microstructural evolution of a CoCrFeMnNi high-Entropy alloy after swaging and annealing. J. Alloys Compd. 2015, 647, 548–557. [CrossRef] Skrotzki, W.; Pukenas, A.; Joni, B.; Odor, E.; Ungar, T.; Hohenwarter, A.; Pippan, R. Microstructure and texture evolution during severe plastic deformation of CrMnFeCoNi high-entropy alloy. In Proceedings of the IOP Conference Series: Materials Science and Engineering, Sydney, Australia, 2–7 July 2017; Volume 194. Dan Sathiaraj, G.; Bhattacharjee, P. Analysis of microstructure and microtexture during grain growth in low stacking fault energy equiatomic CoCrFeMnNi high entropy and Ni–60 wt.%Co alloys. J. Alloys Compd. 2015, 637, 267–276. [CrossRef] Huang, S.; Li, W.; Lu, S.; Tian, F.; Shen, J.; Holmström, E.; Vitos, L. Temperature dependent stacking fault energy of FeCoCrNiMn high entropy alloy. Scr. Mater. 2015, 108, 44. [CrossRef] Laplanche, G.; Kostka, A.; Horst, O.; Eggeler, G.; George, E. Microstructure evolution and critical stress for twinning in the CrMnFeCoNi high-entropy alloy. Acta Mater. 2016, 118, 152–163. [CrossRef] Kang, M.; Won, J.; Kwon, J.; Na, Y. Intermediate strain rate deformation behavior of a CoCrFeMnNi high entropy alloy. Mater. Sci. Eng. A 2017, 707, 16–21. [CrossRef] Mott, N. A theory of work-hardening of metal crystals. Philos. Mag. 1952, 43, 1151–1178. [CrossRef] Friedel, F. Anomaly in the rigidity modulus of copper alloys for small concentrations. Philos. Mag. 1953, 44, 444–448. [CrossRef] Benito, J.; Manero, J.; Jorba, J.; Roca, A. Change of Young’s Modulus of Cold-Deformed Pure Iron in a Tensile Test. Metall. Mater. Trans. A 2005, 36, 3317–3324. [CrossRef] Callister, W.D. Materials Science and Engineering: An Introduction, 8th ed.; John Wiley & Sons: Hoboken, NJ, USA, 2010. Beausir, B.; Fundenberger, J.J. ATEX-Analysis Tools for Electron and X-Ray Diffraction. Available online: http://atex-software.eu/ (accessed on 28 June 2018). Ge, H.; Song, H.; Shen, J.; Tian, F. Effect of alloying on the thermal-elastic properties of 3d high-Entropy alloys. Mater. Chem. Phys. 2018, 210, 302–326. [CrossRef] Ledbetter, H.; Kim, S. Low Temperature Elastic constants of deformed polycrystalline copper. Mater. Sci. Eng. A 1988, 101, 87–92. Morestin, F.; Boivin, M. On the necessity of taking into account the variation in the young modulus with plastic strain in elastic-Plastic software. Nucl. Eng. Des. 1996, 162, 107–116. [CrossRef] Elmay, W.; Prima, F.; Gloriant, T.; Bolle, B.; Zhong, Y.; Patoor, E.; Laheurte, P. Effects of thermomechanical process on the microstructure and mechanical properties of a fully martensitic titanium-Based biomedical alloy. J. Mech. Behav. Biomed. Mater. 2013, 18, 47–56. [CrossRef] [PubMed] Yamaguchi, K.; Adachi, H.; Takakura, N. Effects of plastic strain and strain path on young’s modulus of sheet metals. Met. Mater. 1998, 4, 420–425. Böcker, W.; Bunge, H.J.; Reinert, T. Anomalies of young modulus in Fe-Cu composites after high degree of deformation. Mater. Sci. Forum 1994, 157, 1551–1558. [CrossRef] Joo, S.; Kato, H.; Jang, M.; Moon, J.; Tsai, C.; Yeh, J.; Kim, H. Tensile deformation behavior and deformation twinning of an equimolar CoCrFeMnNi high-Entropy alloy. Mater. Sci. Eng. A 2017, 689, 122–133. [CrossRef] Kireeva, I.; Chumlyakov, Y.; Pobedennaya, Z.; Vyrodova, K.I.; Karaman, I. Orientation dependence of twinning in single crystalline CoCrFeMnNi high-Entropy alloy. Mater. Sci. Eng. A 2017, 705, 176–181. [CrossRef] Kalua, P.; Brandao, L.; Ortiz, F.; Egungwu, O.; Igeb, F. On the texture evolution in swaged Cu-Based wires. Acta Metall. 1998, 38, 1755–1761.

Materials 2018, 11, 1337

39.

40. 41. 42.

14 of 14

Kauffman, A.; Freudenberger, J.; Klauss, H.; Klemm, V.; Schillinger, W.; Subramanya-Sarma, V.; Schultz, L. Properties of cryo-Drawn copper with severely twinned microstructure. Mater. Sci. Eng. A 2013, 588, 132–141. [CrossRef] Bunge, H. Texture Analysis in Materials Science; Mathematical Methods; Butterworth-Heinemann: Oxrord, UK, 1968. Nye, J. Physical Properties of Crystal: Their Representation by Tensors and Matrices; Oxford Science Publications; Oxford University Press: Oxford, UK, 1957. Tian, F.Y.; Varga, L.K.; Chen, N.X.; Delczeg, L.; Vitos, L. Ab initio investigation of high-entropy alloys of 3d elements. Phys. Rev. Lett. B. 2013, 87, 075144. [CrossRef] © 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).