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Novel Co-rich high performance twinning-induced plasticity (TWIP) and transformation-induced plasticity (TRIP) high-entropy alloys
Scripta Materialia 165 (2019) 39–43
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Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Novel Co-rich high performance twinning-induced plasticity (TWIP) and transformation-induced plasticity (TRIP) high-entropy alloys Daixiu Wei a,⁎, Xiaoqing Li b,⁎, Jing Jiang c, Weicheng Heng c, Yuichiro Koizumi d, Won-Mi Choi e, Byeong-Joo Lee e, Hyoung Seop Kim e, Hidemi Kato a, Akihiko Chiba a a
Institute for Materials Research, Tohoku University, 2-1-1 Katahira, Sendai, Miyagi 980-8577, Japan Department of Materials Science and Engineering, KTH - Royal Institute of Technology, 10044 Stockholm, Sweden Graduate School of Engineering, Tohoku University, 6-6-02 Aramaki Aza Aoba, Sendai, Miyagi 980-8579, Japan d Division of Materials and Manufacturing Science, Graduate School of Engineering, Osaka University, 2-1 Yamadaoka, Suita 565-0871, Japan e Department of Materials Science and Engineering, Pohan University of Science and Technology (POSTEC), Pohan 37673, Republic of Korea b c
a r t i c l e
i n f o
Article history: Received 4 December 2018 Received in revised form 10 February 2019 Accepted 12 February 2019 Available online xxxx Keywords: High-entropy alloy Stacking fault energy Phase stability Deformation twinning Martensitic transformation
a b s t r a c t The equiatomic CoCrMnNiFe high-entropy alloy (HEA) has attracted much attention owing to its exceptional mechanical properties. Here, we designed novel face-centered cubic (fcc) phase Co-rich non-equiatomic CoCrMnNiFe HEAs with tensile properties superior to the counterparts, derived from lowering stacking fault energy (SFE) via modifying constituent concentrations. The decrease of Mn, Ni, Fe meanwhile increase of Co, Cr concentrations does reduce the SFE value, based on ab initio and thermodynamics calculations. Hereinto, Co35Cr20Mn15Ni15Fe15 and Co35Cr25Mn15Ni15Fe10 HEAs overcame the strength-ductility trade-off, contributing to twinning-induced plasticity (TWIP) or transformation-induced plasticity (TRIP) effects, respectively. The present study sheds light on developing high performance HEAs. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
High-entropy alloys (HEAs) consisting of multiple principle elements have attracted much academic attention after the concept was firstly proposed in the year of 2004 [1,2]. The CoCrMnNiFe alloy comprised of single face-centered cubic (fcc) phase is one of the most thoroughly studied HEAs. The alloy is usually plastic deformed by dislocation slips assisted with deformation twinning at room temperature or cryogenic temperatures [2–5]. The fcc phase of the alloy is thermodynamically stable over a large temperature range whereas hcp phase tends to become more stable with the decrease of temperature [6–8]. Recently, non-equiatomic Fe-rich FeMnCoCr or Co-rich CoCrFeNi HEAs with improved mechanical properties were developed [9–11], which were derived from reducing the phase stability and lowering the stacking fault energy (SFE), by means of tuning the compositions and concentrations of constituents. The exceptional mechanical properties are contributing to the assistance of twinning-induced plasticity (TWIP) or transformation-induced plasticity (TRIP) effects [9–11], inspired by the plasticity and strengthening mechanisms in many low SFE metals and alloys such as TRIP steels and TWIP steels. On the other hand, Co-based alloys with mechanical properties superior to the abovementioned steels, have been widely used in manufacturing vanes, gas turbines, and metallic orthopedic implants ⁎ Corresponding authors. E-mail addresses: [email protected] (D. Wei), [email protected] (X. Li).
https://doi.org/10.1016/j.scriptamat.2019.02.018 1359-6462/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
[12–14]. The alloys generally exhibit fcc crystal structure at elevated temperatures but hexagonal close packed (hcp) crystal structure at low temperatures, whereas metastable fcc phase can be retained at room temperature via fast cooling from the fcc phase stable temperatures [15]. In addition, the fcc → hcp strain induced martensitic transformation (SIMT) is a predominant plastic deformation mode of the metastable fcc phase at ambient temperature, due to the extremely low SFEs [16–20]. The SIMT contributes significantly to a high strain hardening rate and superior strength, where the martensitic hcp phase act as secondary hardening effect owing to the lack of plasticity of the hcp crystal structure [11,16–20]. Furthermore, elements of Mn, Ni and Fe are the fcc phase stabilizers but Co and Cr are the hcp phase stabilizers of the Co-based alloys [11,21]. It means that the addition of Mn, Ni or Fe will stabilize the fcc phase meanwhile increase the SFE, whereas the addition of Co or Cr will stabilize the hcp phase and lower the SFE of the Co-based alloys, which are comprised of metastable fcc phase at ambient temperature. Therefore, it is feasible to develop novel Co-rich HEAs exhibiting SIMT behavior by tuning the compositions. In the present study, we aimed to reveal the influence of the concentration of constituents on the SFE and fcc phase stability of the CoCrMnNiFe series HEA, using ab initio and thermodynamics calculations. Furthermore, we tried to design novel Co-rich non-equiatomic HEAs with exceptional mechanical properties, via manipulating the
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D. Wei et al. / Scripta Materialia 165 (2019) 39–43
deformation behaviors by lowering the SFE of the equiatomic CoCrMnNiFe HEA. We designed quinary HEAs with nominal compositions of Co20Cr20Mn20Ni20Fe20, Co35Cr20Mn15Ni15Fe15, and Co35Cr25Mn15Ni15Fe10, hereinafter denoted by Ni20Fe20, Ni15Fe15, and Ni15Fe10, respectively. In order to calculate the intrinsic and extrinsic SFEs of the fcc phase HEAs for the {111} 〈112〉 slip systems, the fault-free fcc structure was described by a super cell containing nine fcc {111} layers (one atom per layer) with ! ! ! its lattice vectors a 1, a 2, and a 3 parallel to the [110], [101], and [111] directions, respectively. The two planar faults were introduced by shifting ! ! a 3 along the [112] direction with prescribed displacement vectors 0 ≤ b ! ! ! ! ≤ 2 b p ( b p = 1/6 [112] is the partial Burgers vector). When b = b p, the fcc lattice is transformed into an fcc lattice with an intrinsic stacking fault (ISF) and corresponding energy γisf, whereas an extrinsic stacking fault ! ! (ESF) and corresponding energy γesf forms when b = 2 b p (shearing involves two adjacent layers) as illustrated in Fig. 1a. We found that the chosen super cell yields converged SFEs (at the level of approximately 3 mJ/m2) with respect to a larger tested cell size (12 layers). Our ab initio calculations were based on density functional theory (DFT) method [22], and the Kohn-Sham equations were solved using the exact muffin-tin orbitals method (EMTO) for all the ab initio calculations [23–25]. The Perdew-Burke-Ernzerhof function was adopted for self-consistent determination of the charge density and total energy [26]. The experimental investigations revealed that FeCrNiMnCo HEA is paramagnetic at room temperature, and magnetic phase transitions with decreasing temperature to first a spin glass and second ferromagnetic order occur well below 100 K [27,28]. The other currently investigated HEAs were experimentally determined to be paramagnetic at room temperature as well. For this reason, the calculations in the present study were performed in the paramagnetic state, which was described by the disordered-local moment model [29]. The problem of chemical disorder was treated within the coherent-potential approximation [30], and the basis set included s,
p, d, and f states. Brillouin zone integrations were performed on a 12 × 24 × 3 k-points mesh after careful testing. The Gibbs free energy difference between the hcp and fcc phases (ΔGfcc→hcp, defined as Ghcp − Gfcc) of the alloys at 300 K, were calculated using Thermo-Calc software with the TCFE2000 thermodynamic database and its upgraded version [31,32]. Samples were prepared by arc melting, and the weight of each sample is 50 g. The purity of all the raw materials was higher than 99.9%. Then, the samples were homogenized at 1473 K for 5 h in an Ar atmosphere, and subsequently forged at 1473 K to a 50% reduction in thickness. The forged samples were aged at 1273 K for 10 min with subsequent water quenching, and followed by room temperature rolling to a 40% reduction in thickness. Finally, the samples were aged at 1273 K for 6 min to obtain fully recrystallized grain structures. After that, dog-bone shaped tensile specimens with a gauge geometry of 10 mm × 2 mm × 1 mm were cut by electrical discharge machining, and then grinded, wet polished, and finally mirror-finished. Then, uniaxial tensile tests were conducted at 25 °C with a strain rate of 1.5 × 10−4 s−1. The microstructures of the specimens were characterized by scanning electron microscope (SEM) equipped with an electron backscatter diffraction detector (EBSD), using an acceleration voltage of 15 kV. The EBSD data was analyzed using a TSL OIM data-analysis software. The X-ray diffraction (XRD) with Co Kα radiation was carried out for the phase identification. The generalized stacking fault energy curves in Fig. 1b show the initiation of one whole stacking fault of the Ni20Fe20, Ni15Fe15, and Ni15Fe10 HEAs at 0 K, achieved by shifting along 1/6 ⟨112⟩ at {111} planes of the alloys. Fig. 1c demonstrates the value of intrinsic (γisf) and extrinsic (γesf) SFE of the HEAs, where the value of γisf is relatively lower than that of γesf for all the three HEAs. Both the values of γisf and γesf are lowered by increasing the Co and Cr content at the expense of decreasing Fe, Ni, and Mn content. The γisf of the three alloys at 0 K is −72 mJ/m2 (Ni20Fe20), −75 mJ/m2 (Ni15Fe15), and −87 mJ/m2
Fig. 1. (a) Schematics of special crystal configurations due to {111}[112] shearing of the fcc structure (from left to right, fault-free fcc lattice, fcc lattice with ISF, and fcc lattice with ESF) and tilted super cell geometry for calculating the corresponding planar fault energies. Letters A, B, C, and associated color-coded atoms indicate the stacking sequence, the thin solid lines ! denote the undistorted and sheared nine-layers unit cells, and the thick arrows indicate the slip planes and displacement vectors ( b p = 1/6[112]); (b) Generalized-stacking fault energy of {111}〈112〉 slip and (c) the value of SFE at 0 K calculated by ab initio using DFT method; (d) the value of ΔGfcc→hcp at 300 K calculated by Thermo-Calc of the Co20Cr20Mn20Ni20Fe20, Co35Cr20Mn15Ni15Fe15, and Co35Cr25Mn15Ni15Fe10 high entropy alloys. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
D. Wei et al. / Scripta Materialia 165 (2019) 39–43
(Ni15Fe10), respectively. Fig. 1d shows the value of ΔGfcc→hcp of the alloys at 300 K, where a larger value of the ΔGfcc→hcp yields more stable fcc phase. It can be seen that decreasing the concentrations of Mn, Ni and Fe meanwhile increasing the concentration of Co can reduce the stability of fcc phase, where the value of ΔGfcc→hcp decreases slightly from 87 (Ni20Fe20) to −34 (Ni15Fe15) J/mol. Moreover, increasing the concentration of Cr at the expense of Fe significantly reduces the stability of fcc phase, where the value of ΔGfcc→hcp drops dramatically from −34 (Ni15Fe15) to −276 (Ni15Fe10) J/mol. The result is consistent with the theoretical calculations using Calphad method and new TCHEA1 database in other study [33]. The relationship between the γisf and ΔGfcc→hcp can be expressed as [34]: γisf = 2ρAΔGfcc→hcp + 2σfcc→hcp, where ρA is the planar packing density of close-packed planes, σfcc→hcp is the coherent fcc/hcp interfacial energy. Therefore, the SFE is lowered accompanied by decreasing the stability of fcc phase. Fig. 2 depicts the EBSD IPF (Fig. 2a–c) and phase (Fig. 2d–f) maps of the Ni20Fe20 (Fig. 2a, d), Ni15Fe15 (Fig. 2b, e), and Ni15Fe10 (Fig. 2c, f) alloys. The average grain diameters are indicated in the IPF maps. It can be seen from the IPF maps in Fig. 2a–c that all the three samples are fully recrystallized, because no obvious distortion inside grains is observed. The grains are relatively fine with average diameters ranging from 11.2 μm (Ni15Fe10) to 19.8 μm (Ni20Fe20). The relationship between yieldpstrength and grain size follows the Hall-Petch law [35,36]: σd = σ0 ffiffiffi + k/ d, where d is the average diameter of grains, k is a material dependent constant. Therefore, we can neglect the effect of grain size on the tensile properties among the three samples, because of the insignificant difference in grain size. On the other hand, all the samples are single fcc phase without any hcp phase or precipitate (Fig. 2d–f). Moreover, N98% of the grain boundaries are high angle grain boundaries denoted by dark lines in the phase maps. We can conclude that single fcc phase Corich non-equiatomic CoCrMnNiFe HEAs were produced, and no obvious difference in the microstructures among the alloys was observed. Fig. 3a shows the engineering tensile stress-strain curves of the three HEAs at room temperature. Fig. 3b demonstrates the corresponding yield strengths, ultimate tensile strengths (UTS), and elongations. The yield strength increases from 199 MPa (Ni20Fe20) to 231 MPa (Ni15Fe15), and finally to 305 MPa (Ni15Fe10). Besides, the UTS increases from 565 MPa (Ni20Fe20) to 688 MPa (Ni15Fe15), and finally to 806 MPa (Ni15Fe10). The elongations are 74% (Ni20Fe20), 96% (Ni15Fe15), and 76% (Ni15Fe10), respectively. Therefore, both the Ni15Fe15 and Ni15Fe10 HEAs
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exhibit tensile properties superior to the Ni20Fe20 counterpart, which become stronger and more ductile overcoming the strength-ductility tradeoff. The yield strength and UTS of the Ni15Fe10 alloy are improved by N40% meanwhile the elongation is also increased by 2%, compared with Ni20Fe20 alloy. Fig. 3c shows the strain hardening curves of the three alloys. The hardening rates of the Ni15Fe10 and Ni15Fe15 alloys are obviously higher and decreasing much slower than that of the Ni20Fe20 alloy. It indicates that the plastic deformation and strain hardening mechanisms of the three alloys are different. Fig. 3d shows UTS versus elongation of the alloys in the present study, compared with equiatomic and Fe-rich non-equiatomic CoCrMnNiFe HEAs in other studies [2–5,37–41]. It can be seen that the Ni15Fe10 and Ni15Fe15 alloys exhibit higher strength and/or larger elongation than that of other HEAs. Fig. 4a–b is the XRD patterns of the three alloys before (Fig. 4a) and after (Fig. 4b) tensile tests. All the samples are single fcc phase before tensile, and the Ni20Fe20 and Ni15Fe15 samples remained single fcc phase after tensile. However, hcp phase is observed in the Ni15Fe10 sample after tensile as shown in Fig. 4b, indicating a phase transformation of fcc → hcp occurred during plastic deformation. Fig. 4c is IPF map showing microstructure of the Ni15Fe15 alloy after tensile, and the inserted chart depicts the misorientation angle from point A to B. It can be seen that deformation-induced nanotwins are formed, which the misorienation angle between twin and matrix is approximate 60°. The twinning system is consistent with the common operative deformation twinning system in fcc phase, which the twinning elements are: K1 = (111), η1 = [11 2 ], K2 = (11 1 ), η2 = [112]. Fig. 4d–e is the IPF (Fig. 4d) and phase (Fig. 4e) maps of the Ni15Fe10 alloy after tensile. It can be seen that stripe-shaped hcp phase are produced inside the fcc matrix by the fcc → hcp martensitic transformation, which is analogous to the SIMT mechanism in other Co-based alloys [11,16–20]. As a result, either deformation twinning in the TWIP Ni15Fe15 alloy or fcc → hcp SIMT in the TRIP Ni15Fe10 alloy, contributed to the strain hardening and plasticity of the alloys significantly. The SFE is the main factor governing the plastic deformation behavior of fcc phase [42,43]. The plastic deformation mode switches from dislocation slip to deformation twinning, and finally to fcc → hcp SIMT, accompanied by lowering the SFE from a high value to an extremely low value gradually. In steels, it has been found that a high level of γisf (γisf N 45 mJ/m2) leads to perfect dislocation slip, an intermediate range of
Fig. 2. EBSD IPF (a–c) and phase maps (d–f) showing grain structure and phase constituent of the (a, d) Co20Cr20Mn20Ni20Fe20, (b, e) Co35Cr20Mn15Ni15Fe15, and (c, f) Co35Cr25Mn15Ni15Fe10 high entropy alloys before tensile tests.
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Fig. 3. (a) Room temperature engineering tensile stress-strain curves, (b) the yield strength, ultimate tensile strength, and elongation, (c) strain hardening curves of the Co20Cr20Mn20Ni20Fe20, Co35Cr20Mn15Ni15Fe15, and Co35Cr25Mn15Ni15Fe10 high entropy alloys. (d) The strength versus elongation of the alloys in present work compared with the quinary alloys in other studies [2–5,37–41].
γisf (15 mJ/m2 b γisf b 45 mJ/m2) is associated with deformation twinning, and a low value of γisf (γisf b 15 mJ/m2) results in SIMT [43]. In fcc phase ! alloy with low SFE, a perfect dislocation with Burgers vector b of 1/2〈110〉 ! tends to dissociate into a pair of Shockley partial dislocations with b of 1/ 6〈211〉, which can be described as: 1/2[110] → 1/6[211] + 1/6[121]. Either deformation twinning or SIMT is proceeded by the motion of Shockley partial dislocations, where a twin will generate if the partial
dislocations glide on every two consecutive {111}fcc close packed planes, but the martensitic hcp phase will be produced if the partial dislocations glide on every second {111}fcc planes [44]. It was found that the SFE of each constituent element affected the SFE of CoCrMnNiFe alloy significantly [6,7,11], and decreasing the Ni, Mn, and Fe concentrations independently of increasing the Cr concentration does lower the SFE [21]. The γisf of the equiatomic CoCrMnNiFe
Fig. 4. (a–b) XRD patterns of the Co20Cr20Mn20Ni20Fe20, Co35Cr20Mn15Ni15Fe15, and Co35Cr25Mn15Ni15Fe10 high entropy alloys before (a) and after (b) tensile tests. (c) EBSD IPF map of the Co35Cr20Mn15Ni15Fe15 alloy after tensile fracture. (d–e) EBSD IPF (d) and phase (e) maps of the Co35Cr25Mn15Ni15Fe10 alloy after tensile fracture.
D. Wei et al. / Scripta Materialia 165 (2019) 39–43
alloy is 20–50 mJ/m2 at 298 K [45,46]. In the present study, the SFE of Ni15Fe15 alloy was slightly lowered compared to the equiatomic counterpart, which makes deformation twinning more preferred. As a result, an enhancement in both strength and ductility was achieved owing to the TWIP effect. Besides, the SFE was further lowered in the Ni15Fe10 alloy, leading to a conversion of deformation mode from twinning to SIMT, which enhances the strength appreciably owing to the TRIP effect. In summary, novel fcc phase Co-rich non-equiatomic CoCrMnNiFe HEAs with extraordinarily enhanced tensile properties were developed, by means of reducing SFE and fcc phase stability of the equiatomic counterpart. The ab initio and thermodynamics calculations indicated that decreasing the concentrations of Mn, Ni and Fe or increasing the concentrations of Co and Cr does lower the SFE and fcc phase stability of the HEAs. The tensile strength or ductility of the novel developed TWIP Co35Cr20Mn15Ni15Fe15 and TRIP Co35Cr25Mn15Ni15Fe10 alloys were superior to the equiatomic or Fe-rich non-equiatomic CoCrFeMnNi series HEAs. The UTS of the two HEAs are 688 MPa (Ni15Fe15) and 806 MPa (Ni15Fe10), and the elongations are 96% (Ni15Fe15) and 76% (Ni15Fe10), respectively. The deformation twinning or fcc → hcp martensitic transformation contributed remarkably to the exceptional properties of the two HEAs, respectively. The present study sheds light on the development of novel strong and ductile HEAs. Acknowledgement This work was primarily supported by the ‘Creation of Life Innovation Materials for Interdisciplinary and International Researcher Development’ project, Tohoku University, Japan and the Core Research Cluster of Materials Science Fusion Research project (No. J180002408). This work is a cooperative program (Proposal No.18G0424) of the Cooperative Research and Development Center for Advanced Materials, Institute for Materials Research, Tohoku University. A part of this work was supported by the Swedish Research Council (Grant No. 2016-00236), and the ab initio calculations were performed on resources provided by the Swedish National Infrastructure for Computing (SNIC) at the National Supercomputer Centre. Meanwhile, this work was partially supported by the Future Material Discovery Project of the National Research Foundation of Korea (NRF) funded by the Ministry of Science and ICT of Korea (NRF-2016M3D1A1023383).
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Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Novel Co-rich high performance twinning-induced plasticity (TWIP) and transformation-induced plasticity (TRIP) high-entropy alloys Daixiu Wei a,⁎, Xiaoqing Li b,⁎, Jing Jiang c, Weicheng Heng c, Yuichiro Koizumi d, Won-Mi Choi e, Byeong-Joo Lee e, Hyoung Seop Kim e, Hidemi Kato a, Akihiko Chiba a a
Institute for Materials Research, Tohoku University, 2-1-1 Katahira, Sendai, Miyagi 980-8577, Japan Department of Materials Science and Engineering, KTH - Royal Institute of Technology, 10044 Stockholm, Sweden Graduate School of Engineering, Tohoku University, 6-6-02 Aramaki Aza Aoba, Sendai, Miyagi 980-8579, Japan d Division of Materials and Manufacturing Science, Graduate School of Engineering, Osaka University, 2-1 Yamadaoka, Suita 565-0871, Japan e Department of Materials Science and Engineering, Pohan University of Science and Technology (POSTEC), Pohan 37673, Republic of Korea b c
a r t i c l e
i n f o
Article history: Received 4 December 2018 Received in revised form 10 February 2019 Accepted 12 February 2019 Available online xxxx Keywords: High-entropy alloy Stacking fault energy Phase stability Deformation twinning Martensitic transformation
a b s t r a c t The equiatomic CoCrMnNiFe high-entropy alloy (HEA) has attracted much attention owing to its exceptional mechanical properties. Here, we designed novel face-centered cubic (fcc) phase Co-rich non-equiatomic CoCrMnNiFe HEAs with tensile properties superior to the counterparts, derived from lowering stacking fault energy (SFE) via modifying constituent concentrations. The decrease of Mn, Ni, Fe meanwhile increase of Co, Cr concentrations does reduce the SFE value, based on ab initio and thermodynamics calculations. Hereinto, Co35Cr20Mn15Ni15Fe15 and Co35Cr25Mn15Ni15Fe10 HEAs overcame the strength-ductility trade-off, contributing to twinning-induced plasticity (TWIP) or transformation-induced plasticity (TRIP) effects, respectively. The present study sheds light on developing high performance HEAs. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
High-entropy alloys (HEAs) consisting of multiple principle elements have attracted much academic attention after the concept was firstly proposed in the year of 2004 [1,2]. The CoCrMnNiFe alloy comprised of single face-centered cubic (fcc) phase is one of the most thoroughly studied HEAs. The alloy is usually plastic deformed by dislocation slips assisted with deformation twinning at room temperature or cryogenic temperatures [2–5]. The fcc phase of the alloy is thermodynamically stable over a large temperature range whereas hcp phase tends to become more stable with the decrease of temperature [6–8]. Recently, non-equiatomic Fe-rich FeMnCoCr or Co-rich CoCrFeNi HEAs with improved mechanical properties were developed [9–11], which were derived from reducing the phase stability and lowering the stacking fault energy (SFE), by means of tuning the compositions and concentrations of constituents. The exceptional mechanical properties are contributing to the assistance of twinning-induced plasticity (TWIP) or transformation-induced plasticity (TRIP) effects [9–11], inspired by the plasticity and strengthening mechanisms in many low SFE metals and alloys such as TRIP steels and TWIP steels. On the other hand, Co-based alloys with mechanical properties superior to the abovementioned steels, have been widely used in manufacturing vanes, gas turbines, and metallic orthopedic implants ⁎ Corresponding authors. E-mail addresses: [email protected] (D. Wei), [email protected] (X. Li).
https://doi.org/10.1016/j.scriptamat.2019.02.018 1359-6462/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
[12–14]. The alloys generally exhibit fcc crystal structure at elevated temperatures but hexagonal close packed (hcp) crystal structure at low temperatures, whereas metastable fcc phase can be retained at room temperature via fast cooling from the fcc phase stable temperatures [15]. In addition, the fcc → hcp strain induced martensitic transformation (SIMT) is a predominant plastic deformation mode of the metastable fcc phase at ambient temperature, due to the extremely low SFEs [16–20]. The SIMT contributes significantly to a high strain hardening rate and superior strength, where the martensitic hcp phase act as secondary hardening effect owing to the lack of plasticity of the hcp crystal structure [11,16–20]. Furthermore, elements of Mn, Ni and Fe are the fcc phase stabilizers but Co and Cr are the hcp phase stabilizers of the Co-based alloys [11,21]. It means that the addition of Mn, Ni or Fe will stabilize the fcc phase meanwhile increase the SFE, whereas the addition of Co or Cr will stabilize the hcp phase and lower the SFE of the Co-based alloys, which are comprised of metastable fcc phase at ambient temperature. Therefore, it is feasible to develop novel Co-rich HEAs exhibiting SIMT behavior by tuning the compositions. In the present study, we aimed to reveal the influence of the concentration of constituents on the SFE and fcc phase stability of the CoCrMnNiFe series HEA, using ab initio and thermodynamics calculations. Furthermore, we tried to design novel Co-rich non-equiatomic HEAs with exceptional mechanical properties, via manipulating the
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deformation behaviors by lowering the SFE of the equiatomic CoCrMnNiFe HEA. We designed quinary HEAs with nominal compositions of Co20Cr20Mn20Ni20Fe20, Co35Cr20Mn15Ni15Fe15, and Co35Cr25Mn15Ni15Fe10, hereinafter denoted by Ni20Fe20, Ni15Fe15, and Ni15Fe10, respectively. In order to calculate the intrinsic and extrinsic SFEs of the fcc phase HEAs for the {111} 〈112〉 slip systems, the fault-free fcc structure was described by a super cell containing nine fcc {111} layers (one atom per layer) with ! ! ! its lattice vectors a 1, a 2, and a 3 parallel to the [110], [101], and [111] directions, respectively. The two planar faults were introduced by shifting ! ! a 3 along the [112] direction with prescribed displacement vectors 0 ≤ b ! ! ! ! ≤ 2 b p ( b p = 1/6 [112] is the partial Burgers vector). When b = b p, the fcc lattice is transformed into an fcc lattice with an intrinsic stacking fault (ISF) and corresponding energy γisf, whereas an extrinsic stacking fault ! ! (ESF) and corresponding energy γesf forms when b = 2 b p (shearing involves two adjacent layers) as illustrated in Fig. 1a. We found that the chosen super cell yields converged SFEs (at the level of approximately 3 mJ/m2) with respect to a larger tested cell size (12 layers). Our ab initio calculations were based on density functional theory (DFT) method [22], and the Kohn-Sham equations were solved using the exact muffin-tin orbitals method (EMTO) for all the ab initio calculations [23–25]. The Perdew-Burke-Ernzerhof function was adopted for self-consistent determination of the charge density and total energy [26]. The experimental investigations revealed that FeCrNiMnCo HEA is paramagnetic at room temperature, and magnetic phase transitions with decreasing temperature to first a spin glass and second ferromagnetic order occur well below 100 K [27,28]. The other currently investigated HEAs were experimentally determined to be paramagnetic at room temperature as well. For this reason, the calculations in the present study were performed in the paramagnetic state, which was described by the disordered-local moment model [29]. The problem of chemical disorder was treated within the coherent-potential approximation [30], and the basis set included s,
p, d, and f states. Brillouin zone integrations were performed on a 12 × 24 × 3 k-points mesh after careful testing. The Gibbs free energy difference between the hcp and fcc phases (ΔGfcc→hcp, defined as Ghcp − Gfcc) of the alloys at 300 K, were calculated using Thermo-Calc software with the TCFE2000 thermodynamic database and its upgraded version [31,32]. Samples were prepared by arc melting, and the weight of each sample is 50 g. The purity of all the raw materials was higher than 99.9%. Then, the samples were homogenized at 1473 K for 5 h in an Ar atmosphere, and subsequently forged at 1473 K to a 50% reduction in thickness. The forged samples were aged at 1273 K for 10 min with subsequent water quenching, and followed by room temperature rolling to a 40% reduction in thickness. Finally, the samples were aged at 1273 K for 6 min to obtain fully recrystallized grain structures. After that, dog-bone shaped tensile specimens with a gauge geometry of 10 mm × 2 mm × 1 mm were cut by electrical discharge machining, and then grinded, wet polished, and finally mirror-finished. Then, uniaxial tensile tests were conducted at 25 °C with a strain rate of 1.5 × 10−4 s−1. The microstructures of the specimens were characterized by scanning electron microscope (SEM) equipped with an electron backscatter diffraction detector (EBSD), using an acceleration voltage of 15 kV. The EBSD data was analyzed using a TSL OIM data-analysis software. The X-ray diffraction (XRD) with Co Kα radiation was carried out for the phase identification. The generalized stacking fault energy curves in Fig. 1b show the initiation of one whole stacking fault of the Ni20Fe20, Ni15Fe15, and Ni15Fe10 HEAs at 0 K, achieved by shifting along 1/6 ⟨112⟩ at {111} planes of the alloys. Fig. 1c demonstrates the value of intrinsic (γisf) and extrinsic (γesf) SFE of the HEAs, where the value of γisf is relatively lower than that of γesf for all the three HEAs. Both the values of γisf and γesf are lowered by increasing the Co and Cr content at the expense of decreasing Fe, Ni, and Mn content. The γisf of the three alloys at 0 K is −72 mJ/m2 (Ni20Fe20), −75 mJ/m2 (Ni15Fe15), and −87 mJ/m2
Fig. 1. (a) Schematics of special crystal configurations due to {111}[112] shearing of the fcc structure (from left to right, fault-free fcc lattice, fcc lattice with ISF, and fcc lattice with ESF) and tilted super cell geometry for calculating the corresponding planar fault energies. Letters A, B, C, and associated color-coded atoms indicate the stacking sequence, the thin solid lines ! denote the undistorted and sheared nine-layers unit cells, and the thick arrows indicate the slip planes and displacement vectors ( b p = 1/6[112]); (b) Generalized-stacking fault energy of {111}〈112〉 slip and (c) the value of SFE at 0 K calculated by ab initio using DFT method; (d) the value of ΔGfcc→hcp at 300 K calculated by Thermo-Calc of the Co20Cr20Mn20Ni20Fe20, Co35Cr20Mn15Ni15Fe15, and Co35Cr25Mn15Ni15Fe10 high entropy alloys. (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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(Ni15Fe10), respectively. Fig. 1d shows the value of ΔGfcc→hcp of the alloys at 300 K, where a larger value of the ΔGfcc→hcp yields more stable fcc phase. It can be seen that decreasing the concentrations of Mn, Ni and Fe meanwhile increasing the concentration of Co can reduce the stability of fcc phase, where the value of ΔGfcc→hcp decreases slightly from 87 (Ni20Fe20) to −34 (Ni15Fe15) J/mol. Moreover, increasing the concentration of Cr at the expense of Fe significantly reduces the stability of fcc phase, where the value of ΔGfcc→hcp drops dramatically from −34 (Ni15Fe15) to −276 (Ni15Fe10) J/mol. The result is consistent with the theoretical calculations using Calphad method and new TCHEA1 database in other study [33]. The relationship between the γisf and ΔGfcc→hcp can be expressed as [34]: γisf = 2ρAΔGfcc→hcp + 2σfcc→hcp, where ρA is the planar packing density of close-packed planes, σfcc→hcp is the coherent fcc/hcp interfacial energy. Therefore, the SFE is lowered accompanied by decreasing the stability of fcc phase. Fig. 2 depicts the EBSD IPF (Fig. 2a–c) and phase (Fig. 2d–f) maps of the Ni20Fe20 (Fig. 2a, d), Ni15Fe15 (Fig. 2b, e), and Ni15Fe10 (Fig. 2c, f) alloys. The average grain diameters are indicated in the IPF maps. It can be seen from the IPF maps in Fig. 2a–c that all the three samples are fully recrystallized, because no obvious distortion inside grains is observed. The grains are relatively fine with average diameters ranging from 11.2 μm (Ni15Fe10) to 19.8 μm (Ni20Fe20). The relationship between yieldpstrength and grain size follows the Hall-Petch law [35,36]: σd = σ0 ffiffiffi + k/ d, where d is the average diameter of grains, k is a material dependent constant. Therefore, we can neglect the effect of grain size on the tensile properties among the three samples, because of the insignificant difference in grain size. On the other hand, all the samples are single fcc phase without any hcp phase or precipitate (Fig. 2d–f). Moreover, N98% of the grain boundaries are high angle grain boundaries denoted by dark lines in the phase maps. We can conclude that single fcc phase Corich non-equiatomic CoCrMnNiFe HEAs were produced, and no obvious difference in the microstructures among the alloys was observed. Fig. 3a shows the engineering tensile stress-strain curves of the three HEAs at room temperature. Fig. 3b demonstrates the corresponding yield strengths, ultimate tensile strengths (UTS), and elongations. The yield strength increases from 199 MPa (Ni20Fe20) to 231 MPa (Ni15Fe15), and finally to 305 MPa (Ni15Fe10). Besides, the UTS increases from 565 MPa (Ni20Fe20) to 688 MPa (Ni15Fe15), and finally to 806 MPa (Ni15Fe10). The elongations are 74% (Ni20Fe20), 96% (Ni15Fe15), and 76% (Ni15Fe10), respectively. Therefore, both the Ni15Fe15 and Ni15Fe10 HEAs
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exhibit tensile properties superior to the Ni20Fe20 counterpart, which become stronger and more ductile overcoming the strength-ductility tradeoff. The yield strength and UTS of the Ni15Fe10 alloy are improved by N40% meanwhile the elongation is also increased by 2%, compared with Ni20Fe20 alloy. Fig. 3c shows the strain hardening curves of the three alloys. The hardening rates of the Ni15Fe10 and Ni15Fe15 alloys are obviously higher and decreasing much slower than that of the Ni20Fe20 alloy. It indicates that the plastic deformation and strain hardening mechanisms of the three alloys are different. Fig. 3d shows UTS versus elongation of the alloys in the present study, compared with equiatomic and Fe-rich non-equiatomic CoCrMnNiFe HEAs in other studies [2–5,37–41]. It can be seen that the Ni15Fe10 and Ni15Fe15 alloys exhibit higher strength and/or larger elongation than that of other HEAs. Fig. 4a–b is the XRD patterns of the three alloys before (Fig. 4a) and after (Fig. 4b) tensile tests. All the samples are single fcc phase before tensile, and the Ni20Fe20 and Ni15Fe15 samples remained single fcc phase after tensile. However, hcp phase is observed in the Ni15Fe10 sample after tensile as shown in Fig. 4b, indicating a phase transformation of fcc → hcp occurred during plastic deformation. Fig. 4c is IPF map showing microstructure of the Ni15Fe15 alloy after tensile, and the inserted chart depicts the misorientation angle from point A to B. It can be seen that deformation-induced nanotwins are formed, which the misorienation angle between twin and matrix is approximate 60°. The twinning system is consistent with the common operative deformation twinning system in fcc phase, which the twinning elements are: K1 = (111), η1 = [11 2 ], K2 = (11 1 ), η2 = [112]. Fig. 4d–e is the IPF (Fig. 4d) and phase (Fig. 4e) maps of the Ni15Fe10 alloy after tensile. It can be seen that stripe-shaped hcp phase are produced inside the fcc matrix by the fcc → hcp martensitic transformation, which is analogous to the SIMT mechanism in other Co-based alloys [11,16–20]. As a result, either deformation twinning in the TWIP Ni15Fe15 alloy or fcc → hcp SIMT in the TRIP Ni15Fe10 alloy, contributed to the strain hardening and plasticity of the alloys significantly. The SFE is the main factor governing the plastic deformation behavior of fcc phase [42,43]. The plastic deformation mode switches from dislocation slip to deformation twinning, and finally to fcc → hcp SIMT, accompanied by lowering the SFE from a high value to an extremely low value gradually. In steels, it has been found that a high level of γisf (γisf N 45 mJ/m2) leads to perfect dislocation slip, an intermediate range of
Fig. 2. EBSD IPF (a–c) and phase maps (d–f) showing grain structure and phase constituent of the (a, d) Co20Cr20Mn20Ni20Fe20, (b, e) Co35Cr20Mn15Ni15Fe15, and (c, f) Co35Cr25Mn15Ni15Fe10 high entropy alloys before tensile tests.
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Fig. 3. (a) Room temperature engineering tensile stress-strain curves, (b) the yield strength, ultimate tensile strength, and elongation, (c) strain hardening curves of the Co20Cr20Mn20Ni20Fe20, Co35Cr20Mn15Ni15Fe15, and Co35Cr25Mn15Ni15Fe10 high entropy alloys. (d) The strength versus elongation of the alloys in present work compared with the quinary alloys in other studies [2–5,37–41].
γisf (15 mJ/m2 b γisf b 45 mJ/m2) is associated with deformation twinning, and a low value of γisf (γisf b 15 mJ/m2) results in SIMT [43]. In fcc phase ! alloy with low SFE, a perfect dislocation with Burgers vector b of 1/2〈110〉 ! tends to dissociate into a pair of Shockley partial dislocations with b of 1/ 6〈211〉, which can be described as: 1/2[110] → 1/6[211] + 1/6[121]. Either deformation twinning or SIMT is proceeded by the motion of Shockley partial dislocations, where a twin will generate if the partial
dislocations glide on every two consecutive {111}fcc close packed planes, but the martensitic hcp phase will be produced if the partial dislocations glide on every second {111}fcc planes [44]. It was found that the SFE of each constituent element affected the SFE of CoCrMnNiFe alloy significantly [6,7,11], and decreasing the Ni, Mn, and Fe concentrations independently of increasing the Cr concentration does lower the SFE [21]. The γisf of the equiatomic CoCrMnNiFe
Fig. 4. (a–b) XRD patterns of the Co20Cr20Mn20Ni20Fe20, Co35Cr20Mn15Ni15Fe15, and Co35Cr25Mn15Ni15Fe10 high entropy alloys before (a) and after (b) tensile tests. (c) EBSD IPF map of the Co35Cr20Mn15Ni15Fe15 alloy after tensile fracture. (d–e) EBSD IPF (d) and phase (e) maps of the Co35Cr25Mn15Ni15Fe10 alloy after tensile fracture.
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alloy is 20–50 mJ/m2 at 298 K [45,46]. In the present study, the SFE of Ni15Fe15 alloy was slightly lowered compared to the equiatomic counterpart, which makes deformation twinning more preferred. As a result, an enhancement in both strength and ductility was achieved owing to the TWIP effect. Besides, the SFE was further lowered in the Ni15Fe10 alloy, leading to a conversion of deformation mode from twinning to SIMT, which enhances the strength appreciably owing to the TRIP effect. In summary, novel fcc phase Co-rich non-equiatomic CoCrMnNiFe HEAs with extraordinarily enhanced tensile properties were developed, by means of reducing SFE and fcc phase stability of the equiatomic counterpart. The ab initio and thermodynamics calculations indicated that decreasing the concentrations of Mn, Ni and Fe or increasing the concentrations of Co and Cr does lower the SFE and fcc phase stability of the HEAs. The tensile strength or ductility of the novel developed TWIP Co35Cr20Mn15Ni15Fe15 and TRIP Co35Cr25Mn15Ni15Fe10 alloys were superior to the equiatomic or Fe-rich non-equiatomic CoCrFeMnNi series HEAs. The UTS of the two HEAs are 688 MPa (Ni15Fe15) and 806 MPa (Ni15Fe10), and the elongations are 96% (Ni15Fe15) and 76% (Ni15Fe10), respectively. The deformation twinning or fcc → hcp martensitic transformation contributed remarkably to the exceptional properties of the two HEAs, respectively. The present study sheds light on the development of novel strong and ductile HEAs. Acknowledgement This work was primarily supported by the ‘Creation of Life Innovation Materials for Interdisciplinary and International Researcher Development’ project, Tohoku University, Japan and the Core Research Cluster of Materials Science Fusion Research project (No. J180002408). This work is a cooperative program (Proposal No.18G0424) of the Cooperative Research and Development Center for Advanced Materials, Institute for Materials Research, Tohoku University. A part of this work was supported by the Swedish Research Council (Grant No. 2016-00236), and the ab initio calculations were performed on resources provided by the Swedish National Infrastructure for Computing (SNIC) at the National Supercomputer Centre. Meanwhile, this work was partially supported by the Future Material Discovery Project of the National Research Foundation of Korea (NRF) funded by the Ministry of Science and ICT of Korea (NRF-2016M3D1A1023383).
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