Effect of annealing heat treatment on microstructural evolution and

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Jeong Min Parka, Jongun Moona, Jae Wung Baea, Jaimyun Junga, Sunghak Leea,b, .... duction ratio: 78.6%) to break up the initial coarse-grained micro-.
Materials Science & Engineering A 728 (2018) 251–258

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Effect of annealing heat treatment on microstructural evolution and tensile behavior of Al0.5CoCrFeMnNi high-entropy alloy

T

Jeong Min Parka, Jongun Moona, Jae Wung Baea, Jaimyun Junga, Sunghak Leea,b, ⁎ Hyoung Seop Kima,b, a b

Department of Materials Science and Engineering, POSTECH (Pohang University of Science and Technology), Pohang 790-784, South Korea Center for High Entropy Alloys, POSTECH (Pohang University of Science and Technology), Pohang 790-794, South Korea

A R T I C LE I N FO

A B S T R A C T

Keywords: High-entropy alloy (HEA) Heat treatment Microstructures Mechanical properties Deformation twinning

In this work, the mechanical characteristics and microstructural evolution of Al0.5CoCrFeMnNi high-entropy alloy (HEA) were studied after annealing at various temperatures (1000, 1100, and 1200 °C). X-ray diffraction, scanning electron microscopy, and energy dispersive spectroscopy analyses were performed to reveal the phase and microstructural variations. The mechanical properties related to different microstructures of the alloy were characterized using tensile testing with digital image correlation. Annealing at lower temperatures led to a higher fraction of B2 phase and finer grain size of FCC (face-centered cubic) phase. A good combination of strength and ductility in this alloy was attributed to the ductile FCC matrix and hard secondary B2 phase. The alloy showed the active evolution of deformation twinning due to the low stacking fault energy when Al was added to CoCrFeMnNi to make the HEA. However, for alloy annealed at lower temperatures, twinning activity was suppressed by the smaller size of grains and depletion of Al content in the FCC matrix. The correlation between the microstructure and mechanical properties was also explored using a simple composite model.

1. Introduction High-entropy alloys (HEAs) are defined as multi-component alloys of which most contain five or more principal elements in equiatomic or near-equiatomic composition with high configurational entropy. In contrast, most conventional (low-entropy) metallic alloys are composed of one or two principal elements and some minor elements [1,2]. The high configurational entropy of HEAs promotes the formation of simple single-phase crystalline materials of face-centered cubic (FCC) [3], body-centered cubic (BCC) [4,5], or sometimes hexagonal close-packed (HCP) structures [6]. These multicomponent alloys have been attracting considerable attention all over the world [2–6]. This is because many HEAs reportedly exhibit good mechanical properties, excellent corrosion resistance, high thermal stability, and resistance to fatigue [7–10]. In general, the FCC-HEAs are ductile and soft, while the BCC-HEAs have high strength but relatively low tensile ductility [11]. From this point of view, the FCC-HEAs need improved yield strength and the BCCHEAs need more plasticity. Therefore, there have been several attempts to enhance the mechanical properties of HEAs using secondary phases to reach a reasonable balance between strength and ductility to broaden the applications for HEAs [12–15].



Recent studies showed that addition of Al to FCC-HEAs could have strong effects on the formation of secondary phases resulting in the variation of mechanical properties [16,17]. He et al. [18] studied the alloying effect of Al on the structure and tensile properties with respect to an FCC CoCrFeMnNi HEA. They found that the crystalline structure changed from the initial single FCC structure to an FCC + BCC structure in the range 8–11 at% of Al. This dual phase alloy showed composite behavior with sharp increases in strength and hardness due to the hard BCC phase. However, a certain problem arose as the Al content was increased in the alloy: a large fraction of disordered and ordered (A2 and B2, respectively) BCC phases could lead to poor ductility. Moon et al. [19] investigated cracking phenomenon in Al0.5CoCrFeMnNi HEA during cold-rolling. They reported, from the results of microstructural analysis and thermodynamic calculation, that the cracking behavior was induced by the formation of an AlNi-rich B2 phase. They proposed that homogenization heat treatment at high temperature might minimize the poor workability of this alloy by reducing the B2 phase which is known to be brittle in the ingot [11]. However, no studies on the microstructural and mechanical properties of the Al0.5CoCrFeMnNi alloy after annealing heat treatment and recrystallization have been reported. In this study, we explored the microstructural evolution of the

Corresponding author at: Department of Materials Science and Engineering, POSTECH (Pohang University of Science and Technology), Pohang 790-784, South Korea. E-mail address: [email protected] (H.S. Kim).

https://doi.org/10.1016/j.msea.2018.05.041 Received 23 December 2017; Received in revised form 11 April 2018; Accepted 11 May 2018 0921-5093/ © 2018 Elsevier B.V. All rights reserved.

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Al0.5CoCrFeMnNi HEA after cold rolling followed by annealing heat treatment. The characterization of tensile behaviors of the annealed alloy was also investigated in parallel. In addition, the strengthening behavior associated with the microstructure was described using a simple composite model to quantify the effect of the annealing heat treatment in the present alloys.

VBCC =

FCC phase: {111} FCC, {200} FCC, {220} FCC, {311} FCC, and {222} FCC. BCC phase: {110} BCC, {200} BCC, {211} BCC, and {220} BCC. Also, based on the Bragg's law (nλ = 2 dhklsinθ ), the interplanar spacing (dhkl ) of each plane was obtained from the diffraction angles (θ ) of each XRD peak position of FCC and BCC phases. Using these values, the lattice parameters (a) of the FCC and BCC phases could be estimated using the following relationship; dhkl = a/(h2 + k 2 + l 2)1/2 for cubic systems. The estimated lattice parameters were ~3.61 and ~2.89 Å for the FCC and BCC phases, respectively. The XRD analysis was repeated three times to ensure the reliability of the data for calculating the phase volume fractions of all the samples.

2.1. Sample preparation Al0.5CoCrFeMnNi HEA ingots were fabricated using vacuum induction melting (VIM) of the pure elements (purity above 99.9%). The homogenization heat treatment for the cast ingots was carried out at 1200 °C for 6 h to reduce the inhomogeneity of the chemical compositions. This heat treatment also contributed to dissolving the secondary phase and improving their workability for cold rolling in accordance with recent publication [19]. After the homogenization treatment, the ingot was cold rolled for thickness reduction from 7.0 to 1.5 mm (reduction ratio: 78.6%) to break up the initial coarse-grained microstructure. To control the phase fraction of the Al0.5CoCrFeMnNi, annealing heat treatments were performed under various conditions. Annealing heat treatments for the as-rolled samples were conducted at 1000 °C for 15 min, 1100 °C for 10 min, and 1200 °C for 5 min with subsequent water quenching. An equilibrium phase diagram of the present alloy can be utilized to determine annealing temperatures for controlling the BCC phase fraction. This was calculated using Thermo-Calc software, the thermodynamic database TCFE2000, and its upgraded version [19]. The annealing times were varied to ensure sufficient nucleation and growth processes of the secondary phase because lower temperature led to slower diffusion of the elements [20]. In this paper, the alloys annealed at 1000 °C, 1100 °C, and 1200 °C are labelled H10, H11, and H12.

2.3. Microstructural analysis Microstructural analysis was performed using scanning electron microscopy (SEM) and electron backscatter diffraction (EBSD). All the specimens were finely polished to a surface roughness of 1 µm and then polished more using a colloidal silica suspension. SEM analysis was conducted using a high-resolution field emission SEM (FE-SEM, JSM7800F PRIME, JEOL Ltd., Japan) equipped with back-scattered electron (BSE) and dual energy-dispersive X-ray spectroscopy (Dual EDS) detectors. After the SEM analysis, EBSD analysis was performed using FE-SEM (XL-30S FEG, FEI Company, USA). The EBSD data were then interpreted using orientation imaging microscopy (OIM) analysis software (TSL OIM Analysis 5.2). The inverse pole figure (IPF), grain boundary (GB), and phase maps were extracted from the EBSD data. Furthermore, the transmission electron microscope (TEM, JEM2200FS, JEOL Ltd., Japan) analysis was carried out to clarify the secondary phase. The TEM specimen was also prepared by fine mechanical polishing (1 µm) followed by colloidal silica polishing. The site-specific specimen involving secondary phase was extracted using a focused ion beam (FIB, FEI Company, USA) technique [23].

2.2. Phase identification of the Al0.5CoCrFeMnNi HEA Phase identification of the Al0.5CoCrFeMnNi HEA was performed by X-ray diffraction (XRD) using a RIGAKU D/MAX-2500 × -ray diffractometer with an incident beam of Cu Kα radiation (wavelength λ = 1.5418 Å). All samples were polished with silicon carbide papers (400–1200 grit) and 1 µm diamond powder for the XRD analysis. The scans were performed from 30 to 100° of 2θ with a step size of 0.02° and scan speed of 2°/min. For quantitative identification of the FCC and BCC phases, the volume fractions of each phase were determined using the relative intensities of the diffraction peaks. The integrated intensity (I) of a diffraction line can be expressed as [19,21]:

2.4. Mechanical testing and characterization The tensile tests were performed at a strain rate of 10−3 s−1 using a universal testing machine (Instron 1361, USA) with the digital image correlation (DIC) technique using an optical 3-D deformation analysis system (ARAMIS 5 M, GOM Co., Germany) to measure precise strains. Dog-bone shaped specimens with a gauge length of 5.0 mm, a width of 2.5 mm, and a thickness of 1.5 mm were used in the tensile tests. All the tests were performed at room temperature (RT) and repeated at least three times. After the tensile tests, the EBSD analysis was carried out on the planes with normal directions of the specimens for interpretation of the microstructural behavior of the alloys using FE-SEM (XL-30S FEG, FEI Company, USA).

(1)

where K is a constant for the given experimental conditions, R is the material scattering factor, and V is the volume of a phase exposed to the X-ray beam [22]. The scattering factor R value depends on the diffraction angles, diffraction planes, and materials, as follows [19]:

R=

m (LP ) e−2mF 2, ν2

3. Results and discussion (2) 3.1. Phase evolution behaviors of the Al0.5CoCrFeMnNi HEAs after annealing

where m is the multiplicity, ν is the volume of the unit cell, LP is the Lorentz polarization factor, e−2m is the Debye-Waller temperature factor, and F is the structure factor. The R values for each peak position were from Ref. [19]. Using the relationship of intensities of the diffraction peaks (I) and exposed volume fraction (V) from the Eq. (1), the relative volume fractions of FCC (VFCC ) and BCC (VBCC ) phases can be estimated using the following equations:

VFCC =

IFCC / RFCC , [(IFCC /RFCC ) + (IBCC /RBCC )]

(4)

From the XRD data, the diffraction peak positions used for determining the volume fractions of each phase are:

2. Experimental procedure

I = KRV ,

IBCC / RBCC . [(IFCC /RFCC ) + (IBCC /RBCC )]

Fig. 1(a) indicates the XRD patterns for the samples, which reveal the presence of two phases designated as FCC and BCC in the H10 and H11 alloys, while the H12 alloy indicates a single FCC crystal structure. The measured volume fraction of the BCC phase (shown in Fig. 1b) was compared with the volume fraction calculated using the thermodynamic method [19]. In the H10 and H11 alloys, the volume fraction of the BCC phase measured, was ~10.56 and ~4.07%, respectively.

(3) 252

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Fig. 1. (a) X-ray diffraction patterns of the annealed samples under different conditions, and (b) comparison of measured phase volume fractions (BCC) by XRD with thermodynamic calculations [19].

Fig. 2. Microstructure of the Al0.5CoCrFeMnNi alloy for H12 (a, d), H11 (b, e), and H10 (c, f): (a–c) SEM-BSE images, (d–f) EBSD phase maps. The FCC and BCC phases are indicated by red and green colors, respectively, in the EBSD phase maps.

grain size of the FCC phase in the H12 alloy is much larger than that in the H11 alloy. It could be supposed that the secondary BCC phase sufficiently limited grain growth of the matrix FCC phase due to the Zener pinning effect [27]. For the elemental analysis of the secondary phase, EDS analysis was carried out, and these results are summarized in Table 1. Fig. 3 shows the EDS mapping result for the BCC phase and surrounding region in the H11 alloy. There are apparent Al-Ni segregations in the BCC phase area, while the concentrations of Cr and Fe were

Fig. 2 shows the SEM-BSE images (Fig. 2a–c) and EBSD phase maps (Fig. 2d–f) for the H12, H11, and H10 alloys. In Fig. 2(d–f), the FCC and BCC phases are indicated in red and green colors, respectively, and the high-angle grain boundaries (GBs, misorientation angles > 15°) and annealing twin boundaries (TBs, ∑ 3 = 60° @ < 111 >) were indicated by black and blue lines, respectively. The microstructures clearly reveal that all of the alloys are fully recrystallized to form equiaxed grains. Abundant annealing twins can easily be found in their microstructure, which were also highlighted in blue lines in the EBSD phase maps. The result of the profuse annealing twins on the microstructure of the alloys could indicate that the alloy has low stacking-fault energy (SFE) [24,25]. The other HEAs having low SFE (e.g., CoCrFeMnNi or CoCrFeNi HEAs), were also reported to have abundant annealing twins in recrystallized microstructures [26,27]. The secondary B2 phase was clearly observed mostly in grain boundaries as a spherical shape in the H10 and H11 alloys, whereas there was no observable secondary phase in H12 alloy except for a number of oxides. These results corresponded to the minor peak intensity observed in the XRD analysis (Fig. 1a). The average grain sizes of the FCC phase in H12, H11, and H10 alloys were (100, 8.5, and 2) μm, respectively. The sizes of the secondary BCC phase for H11 and H10 alloys were 1 and 0.5 µm, respectively. Although there is a difference in annealing temperature between the H11 and H12 alloys, the

Table 1 Chemical compositions of the different regions in the H12, H11, and H10 alloys by EDS. Specimens

H12 H11 H10

253

Regions

Overall Al-oxide Overall BCC phase Overall BCC phase

Chemical compositions, at% Al

Co

Cr

Fe

Mn

Ni

O

9.19 48.64 9.44 19.92 8.97 16.26

17.94 3.56 18.24 15.19 18.04 15.5

17.57 6.09 17.86 9.21 17.86 11.36

18.32 3.47 18.33 10.03 17.64 11.21

18.46 5.11 17.03 17.52 18.55 18.39

18.52 3.95 19.11 28.13 18.93 27.29

– 29.17 – – – –

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Fig. 3. EDS mapping result around a secondary phase in the H11 alloy: SEM micrographs (left side), and EDS mapping images of Al, Ni, Cr, Mn, Fe, and Co (right side).

relatively low compared to their occurrence in the surrounding region. These results are in good agreement with the results in precedent publications [19]. It is well known that the aluminum generally acts as a strong AlNi-rich ordered BCC (B2) phase stabilizer in the Co-Cr-FeMn-Ni based HEA system [16,17]. Many similar observations have been extensively reported for Al-containing HEAs (e.g., AlxCoCrFeNi [28], AlxCoCrCuFeNi [16], and AlxCoCrFeNiTi [29]). It was attributed to the strong negative enthalpy of formation of the B2 structure in Al-Ni system [30]. To closely examine the B2 phase, TEM analysis was conducted for the H11 alloy. Fig. 4 shows the results of TEM micrographs; the EDS layered image indicating Al and Cr contents to distinguish each phase (Fig. 4a) and the magnified image of the interested area (Fig. 4b) with selected area diffraction pattern (SADP) of the B2 phase region (Fig. 4c). The SADP in Fig. 4(c) showed clear superlattice reflections in diffraction patterns of the BCC phase with respect to the [001] zone axis [30–32]. However, the high magnification image (Fig. 4b) shows the presence of nano-sized precipitates inside the B2 phase. Reddy et al. [32] also investigated the phase evolution behavior of Al0.5CoCrFeMnNi HEA having dual phases (FCC+B2). They also founded that nano-precipitates in B2 phase. They have suggested that the origin of precipitates in B2 phase could be the spinodal decomposition, which were extensively reported in Al-containing HEAs [18,28–32], but not fully clarified yet.

Fig. 5. Tensile properties of the Al0.5CoCrFeMnNi HEAs annealed under different conditions (H10–H12). The inset table shows the yield strength (YS), tensile strength (TS), and total elongation (T. EL) for each alloy.

3.2. Tensile behaviors of the Al0.5CoCrFeMnNi HEAs Fig. 5 shows the stress-strain curves for the current HEA annealed at different temperatures. The H12 alloy shows yield and tensile strengths of 278 and 619 MPa, respectively, along with a total elongation of 60.4%. The strengths increased in the H11 alloy with a small loss of elongation (43.9%). In the H10 alloy, the yield strength, tensile strength, and elongation are 730 MPa, 968 MPa, and 29%, respectively.

Fig. 4. TEM micrographs of the H11 alloy: (a) EDS layered images with respect to Al and Cr contents, (b) the magnified image, and (c) the SADP of the B2 phase region. 254

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The trade-off aspects on strength-ductility of the alloys can be explained using their initial microstructures (shown in Fig. 2). Generally, BCC phase is much harder than FCC phase due to the difference in the number of available slip systems under the plastic deformation [18,31]. Moreover, the Al-Ni rich B2 phase in Al-containing HEAs was reported that it has exceptional high hardness values (e.g., for the Al2.8CoCrCuFeNi HEA having a single B2 phase, the Vickers hardness reached 655 HV) [11,33,34]. This was attributed to high Peierls friction and the extended strain-field of the highly ordered B2 phase [33,34]. He et al. investigated the tensile properties of dual phase Alx(CoCrFeMnNi)100-x with change of Al concentrations (8–11 at%) [18]. They showed that addition of the Al contents resulted in considerable increase of both BCC fraction and strength of the alloys. In the present study, lower annealing temperature promoted the formation of high fraction of hard BCC phase in soft FCC matrix. It also contributed to sharp increase of strength of the annealed alloys like the traditional composites. However, increasing the fraction of the hard phase not only increases the strength but also causes loss of ductility. The hard secondary phase in a very soft matrix mostly behaves as initiation sites of microvoids by localization of stress around them, which promotes fracture of the alloys [35,36]. Thus, it is very important to control the volume fraction, as well as the distribution and morphology of the hard secondary phase. In the present results, all of the engineering stress-strain curves exhibited good work hardening behavior and a good combination of strength and ductility. In particular, the yield strength of the H10 alloy was much higher (730 MPa) than those in previously reported FCC based-HEAs (200–600 MPa), while it had moderate tensile ductility (29%) compared to BCC based-HEAs [11]. The mechanical properties were also improved in comparison with the as-cast Alx(CoCrFeMnNi)100-x (x = 9, at%) HEA reported previously [18]. This is because the cold rolling and subsequent annealing process efficiently broke up the casting-microstructure and reconstructed the secondary phase distribution more finely via recrystallization. Fig. 6 presents DIC images of the microstructures after tensile deformation for the H12 (Fig. 6 a–b), H11 (Fig. 6c–d), and H10 (Fig. 6e–f) alloys. The extracted locations on the EBSD IPF maps are marked with blue boxes in the local strain distribution maps. The local strains for the regions marked on the DIC images were (59, 63, and 44) % for the H12,

Fig. 7. Work hardening rate (d σ /d ε ) as a function of the true plastic strain for the annealed Al0.5CoCrFeMnNi alloys (H10–H12).

H11, and H10 alloys, respectively. For the DIC images (Fig. 6a, c, and e), it can be seen that the fracture surfaces were cup-and-cone shapes (usual in ductile fracture), and that the specimens annealed at lower temperature have clearly shorter necking regions. In the IPF maps (Fig. 6b, d, and f), most of the grains of the FCC matrix were elongated along the tensile directions, while the B2 phase retained shapes close to their initial ones. It should be noted that there are profuse deformation twins in the microstructure of the H12 alloy. However, in the H11 alloy, there are fewer observations of deformation twinning than in the H12 alloy. Furthermore, for the H10 alloy, deformation twins were scarce (Fig. 6f). 3.3. Work hardening behaviors in the tensile tests Fig. 7 shows the work-hardening rate (∂σ/ ∂ε , WHR) as a function of the true plastic strain (εp,true) curves for the annealed alloys. From these results, it can be seen that the initial WHRs of the H10 and H11 alloys are much higher than that of the H12 alloy. This can be attributed to the large difference in grain size of the FCC matrix (Fig. 2: mean grain size of FCC phase is ~100 µm for H12, ~8.5 µm for H11, and ~2 µm for H10), mainly due to dislocations and GBs interactions [37,38]. Moreover, the WHR of the H11 alloy is slightly higher than that of the H10 alloy, and has a large volume fraction of B2. These results are in good

Fig. 6. Microstructures of the present HEAs after tensile deformation; (a, c, and e) show local strain maps of the fractured tensile samples extracted by DIC methods, and (b, d, and f) indicate the IPF maps corresponding to the areas highlighted by a deep-blue box on the specimen images for (a–b) H12, (c–d) H11, and (e–f) H10 alloys. The GBs, TBs, and phase boundaries (PBs) are also indicated by black, blue, and thick red lines, respectively. Deformation twins are highlighted by white arrows in the EBSD maps. 255

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quasi-static conditions at RT [39,42]. This could be due primarily to the smaller grain size (compared with H12 alloy). Wu et al. investigated the effect of grain size on mechanical twinning of Al.1CoCrFeNi HEAs [38]. They showed that a decrease in the grain size could result in low twinning activity, which degraded the work-hardening ability and tensile ductility. Furthermore, the WHR values of the H11 alloy were relatively higher than that of the H10 alloy. Tian et al. [52] also explored the grain-size effect on WHR behaviors of Cu-Al (1:15 at%) alloys with different grain sizes (~47, 7, and 0.6 µm), and reported WHR behaviors similar to the present trends. In addition, comparing the H11 and H10 HEAs, the increased B2 phase fraction in the latter decreased the Al-concentration in the FCC phase, as shown in Fig. 3. This might also contribute to suppression of deformation twinning activity due to increased SFE by decline of Al content in the FCC matrix [49].

agreement with the results in a previous study [18]. Noticeably, the WHR of the H12 alloy became higher than those of the H11 and H10 alloys at high strains. Three different regimes could be noted on the WHR curves for the H12 alloy. In detail, the WHR for the H12 alloy rapidly decreased with strain in the initial stage of deformation (Stage I, εp,true < 4.3%), which was governed by dislocation glide [38]. After Stage I, the WHR increased with the plastic strain (Stage II, 4.3% < εp,true < 15.6%), followed by decrease of WHR with further increase in strain (Stage III, εp,true > 15.6%). The increase in the WHR in Stage II is rare in HEAs, but is very common in TWIP steels [38–40]. The literature suggests that the onset of increasing WHR with the plastic strain was primarily due to a change in the deformation mechanism with respect to slip and twinning [37,41]. The obstacles to dislocation motion are primarily associated with grain boundaries. When deformation twins are generated during plastic deformation, dislocations are blocked by the twin boundaries and dislocation mean free path decreases, especially for dislocations on slip planes intersecting the twinning plane [40]. The decreasing mean free path has the same effect as the dynamically decreased grain size, leading to acceleration WHR (dynamic Hall-Petch effect) upon Stage II [40–43]. Thereafter, twinning evolution gradually decreases with further increase in strain, and the deformation mechanism is mainly governed by dislocation-slip with decrease of WHR again in Stage III [37,38,41]. In the case of the most popular FCC-structured HEAs, the CoCrFeNi and CoCrFeMnNi HEAs have been reported as having low SFE (20 and 30 mJ/m2, respectively) [44]. Compare this with Al (SFE ~ 86 mJ/m2) or Ni (SFE ~ 120–130 mJ/m2) [45,46]. Thus, deformation twins were often observed in these HEAs in various deformation modes, particularly under cryogenic conditions [39,47,48]. Laplanche et al. conducted uniaxial tension tests for the CoCrFeMnNi HEA, and also investigated microstructural evolution under different strains [39]. They observed that deformation twins could appear at sufficiently high strain (~20%) when the flow-stress exceeded the critical twinning stress (720 ± 30 MPa at 293 K). However, the Co-Cr-Fe-Mn-Ni based HEAs [39,42] did not show increasing WHR with strain as shown in Fig. 5, even though Al-containing HEAs often exhibited this behavior as deformation proceeded [38,41,46]. The reason for this difference might be attributable to the difference in SFE. Several recent papers [46,49] reported that the addition of aluminum contributed to reduction of the SFE for a Co-Cr-Fe-Mn-Ni HEA system due to the large difference in atomic size between Al and the other elements. This leads to greater lattice distortion, compared to other elements that might be added to the HEAs, and probably caused a decrease in SFE [49]. Generally, the SFE and grain size have strong influence on the critical twinning stress of materials; the low SFE or large grain size usually reduces the critical twinning stress [50]. In the case of the CoCrFeMnNi HEAs having mean grain sizes of 17 and 53 µm, reported critical stress values for twinning were ~720 and ~620 MPa, in the uniaxial tensile and compressive tests, respectively [41,42]. Kireeva et al. conducted tensile tests for CoCrFeMnNi HEA single crystals, and they determined the critical resolved shear stress for twinning (τT), 110–140 MPa, at RT [51]. From these data, the critical tensile stress for the onset of twinning can be estimated (~ 380 MPa), assuming a Taylor factor of M = 3.06. For the present results on the WHR curve of the H12 alloy, the slope changed from Stage I to Stage II for the H12 alloy after plastic strain of 4.3 ± 0.4%. This corresponds to applied true stress of 379 ± 7 MPa. This value is similar to the critical twinning stress (~380 MPa) for a single crystal of CoCrFeMnNi HEA, which supports the notion that Al addition to the CoCrFeMnNi HEA effectively lowered SFE, leading to decreased critical twinning stress. Therefore, the H12 alloy showed increasing tendencies of the WHR curves (Stage II), and it could be initiated at earlier plastic strains than the CoCrFeMnNi HEA did [41,42,51]. Meanwhile, the H10 and H11 alloys did not show an increase of WHR, while they did show gradually decreasing trends of WHR with straining. This is WHR behavior typical of CoCrFeMnNi HEAs under

3.4. Relationship between microstructure and strength of the Al0.5CoCrFeMnNi HEAs In accordance with Figs. 2 and 5, the tensile properties of the present HEAs were strongly affected by their microstructures. The lower annealing temperature produced finer grain-size in the FCC matrix and higher volume fraction of the B2 phase in these HEAs. This results primarily from the grain size and fraction of hard B2 phase in the present alloys. He et al. [18] described the strengthening behavior of the Alx(CoCrFeMnNi)100-x (x = 8–11, at%) HEAs having a duplex FCC + BCC structure with increasing Al content as a simple composite model, (5)

σ = σFCC *VFCC + σBCC *VBCC ,

where σFCC and σBCC are the strengths of the FCC and BCC phases, respectively, and VFCC and VBCC are the volume fractions of the two phases. However, they did not consider grain boundary hardening. It is known that grain refinement definitely improves the strength of an alloy via the Hall-Petch relationship. Sun et al. [53] investigated the tensile properties of the CoCrFeMnNi HEAs over a wide range of grain sizes, and they suggested the following Hall-Petch relationship of the alloy: σYS = 194 + 490*d−1/2 MPa. To apply these effects readily to the present HEAs, several assumptions were employed. The variation in composition of each phase caused by the heat treatment was neglected. The B2 phase was regarded as having interfaces with FCC phase only. The boundaries of mutual (adjacent) B2 phases were ignored. Finally, Eq. (5) was modified to produce the following equation:

σYS, overall = (σ0,

FCC

−1/2 + KFCC *dFCC )*VFCC + (σB2 + KFCC / B2*dB2 −1/2)*VB2,

(6) where σ0, FCC is the friction stress of the FCC phase and KFCC is the strengthening coefficient of the grain refinement for the FCC matrix, which were derived from CoCrFeMnNi HEAs [53]. Here, σB2 is the strength of a B2 phase. The size effect of the B2 phase was also embodied as KFCC / B2 in Eq. (6), which was an experimental parameter. The volume fraction of the secondary phase (VB2) was measured at ~10.56% and ~4.07% for the H10 and H11 alloy, respectively. Then, to calculate the σ0, FCC of the Al0.5CoCrFeMnNi HEAs, a substitutional solution strengthening effect was applied to the CoCrFeMnNi HEAs using the following model [18]:

∆σsolid solution = M*

G

*εs3/2 * 700

c

1/2

,

(7)

where G is the shear modulus of the CoCrFeMnNi HEA (estimated to be ~81 GPa [42]), c is the molar ratio of Al, and the M is the Taylor factor. Here, εs is the interaction parameter involving elastic and atomic size mismatch of Al with the CoCrFeMnNi HEA. The value was derived from [18]. Assuming that M is 3.06, the ∆σsolid solution is calculated to be ~21 MPa. 256

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FCC matrix, with increasing volume-fraction of B2 phase. (5) A simple composite model employing the Hall-Petch relationship was used to describe the strengthening behaviors of the annealed HEA alloys. This model offers reasonable interpretation of the strengthening effect in the annealed Al0.5CoCrFeMnNi HEAs. Acknowledgement This work was supported by the Future Material Discovery Project of the National Research Foundation of Korea (NRF) funded by the Ministry of Science, ICT and Future Planning (MSIP) of Korea (NRF2016M3D1A1023383). Fig. 8. Experimental results vs. composite model for yield strengths of the H10, H11, and H12 alloys. The contribution of the FCC and B2 phase to yield strength is indicated by a yellow bar (FCC) and a blue bar (B2), respectively. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)

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For σB2 , a Vickers hardness of a single B2 phase (Hv ≈ 642 MPa) of Al2.8CoCrCuFeNi HEA [34] was converted to yield strength (~2141 MPa) using the general relationship between hardness and yield strength, in regards to a material having weak WHR behavior (Hv ≈ 3 σYS ) [54,55]. By substituting these parameters and experimental data into Eq. (6), the strengthening model for the present HEAs was obtained as follows: −1/2 σYS, overall = (215 + 490*dFCC )*VFCC + (2141 + 61*dB2 −1/2)*VB2.

(8)

The predicted strengths from Eq. (8) were plotted with the experimental results in Fig. 8. Also presented was the strengthening contribution of FCC (yellow bars) and B2 phases (blue bars) to the composite structure of the present HEAs. This could provide a good interpretation for strength increment in relation to the microstructure of the Al0.5CoCrFeMnNi HEAs (Fig. 8). 4. Conclusions In this study, an Al0.5CoCrFeMnNi HEA, after cold rolling and subsequent annealing heat treatment at different temperatures, was investigated to explore the effects of microstructural evolution on tensile deformation behavior at RT. (1) After the annealing heat treatment, the secondary B2 phase was clearly observed in the H10 and H11 alloys, while no secondary phase was observed in the H12 alloy. The lower annealing temperature caused a greater B2 phase fraction and smaller grain size of FCC phase. The measured volume fractions of the B2 phase using XRD peaks were determined to be ~10.56% (H10) and ~4.07% (H11), respectively. (2) The EDS mapping shows that Al and Ni were enriched in the B2 phase, while the Cr and Fe were relatively lean compared to the FCC matrix region. In addition, profuse annealing twins were observed in the microstructures of all the alloys (H10–H12). It is proposed that these alloys have low stacking fault energy (SFE). (3) All of the tensile properties exhibited fine work-hardening behaviors and provided an improved combination of strength and ductility in comparison with the as-cast state. This good balance was attributed to the notion that the alloy could behave as a composite structure consisting of a fine combination of a ductile FCC matrix and hard B2 phase. (4) The early onset of the deformation twinning was observed in the H12 alloy than for other Co-Cr-Fe-Mn-Ni based HEAs. This could be derived from the decrease of SFE caused by addition of Al to the CoCrFeMnNi HEAs. On the other hand, weak deformation twinning behaviors were observed in the H10 and H11 alloys. This could be due to decrease in the grain size and depletion of Al content in the 257

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