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Revealing the Hall-Petch relationship of Al0.1CoCrFeNi high-entropy alloy and its deformation mechanisms
Journal of Alloys and Compounds 795 (2019) 269e274
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Revealing the Hall-Petch relationship of Al0.1CoCrFeNi high-entropy alloy and its deformation mechanisms J. Yang a, b, J.W. Qiao c, S.G. Ma a, b, *, G.Y. Wu a, b, D. Zhao a, b, Z.H. Wang a, b, ** a
Institute of Applied Mechanics, College of Mechanical and Vehicle Engineering, Taiyuan University of Technology, Taiyuan 030024, China Shanxi Key Laboratory of Material Strength and Structural Impact, Taiyuan University of Technology, Taiyuan 030024, China c Institute of High-entropy Alloys, College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 22 February 2019 Received in revised form 28 April 2019 Accepted 30 April 2019 Available online 4 May 2019
Al0.1CoCrFeNi high-entropy alloys with fully recrystallized single-phase structures and various grain sizes ranging from 4.7 to 59.5 mm are fabricated through cold rolling and subsequent annealing. Tensile tests reveal that the yielding strength has a linear relationship with the inverse square root of grain size, i.e. a Hall-Petch type relationship, shown as sYTS ¼ 83 þ 464d1=2 . Moreover, similar deformation textures, i.e. strong <111> fiber, moderate <100> fiber, and relative stable Rotated cube and Brass components between fine-grained (FG) and coarse-grained (CG) alloys, in combination with TEM analyses, suggest analogous deformation mechanisms for both alloys and more active twinning behavior of the CG alloy. © 2019 Elsevier B.V. All rights reserved.
Keywords: High-entropy alloy Grain size Hall-Petch relationship Deformation mechanism
1. Introduction High-entropy alloys (HEAs) have drawn great interests in the materials field due to amazing properties, such as high strength and good ductility, potential low- and/or high-temperature properties, good wear and fatigue resistances, and excellent fracture toughness [1e10]. Distinct from traditional alloys that are based on one or two major elements, the HEAs usually contain five or more principal elements in equimolar or near-equimolar ratio. The unique alloy design concept often yields the formation of random solid solutions upon solidification, rather than complex phases or intermetallic compounds, probably due to the unique high mixing entropy effect [1e3]. Therefore, the HEAs exhibit a good combination of great academic values and potential application prospects. In fact, a large number of investigations on phase formation, microstructure evolution, mechanical property and deformation mechanisms of the HEAs have been reported [4e11]. Among various HEAs, the alloys with a face-centered cubic (FCC)
* Corresponding author. Institute of Applied Mechanics, College of Mechanical and Vehicle Engineering, Taiyuan University of Technology, Taiyuan 030024, China. ** Corresponding author. Institute of Applied Mechanics, College of Mechanical and Vehicle Engineering, Taiyuan University of Technology, Taiyuan 030024, China. E-mail addresses: [email protected] (S.G. Ma), [email protected] (Z.H. Wang). https://doi.org/10.1016/j.jallcom.2019.04.333 0925-8388/© 2019 Elsevier B.V. All rights reserved.
solid-solution structure, especially for the low stacking-faultenergy (SFE) systems, have been the current research hotspots [11e19]. Plenty studies have shown that the SFE could significantly affect the deformation modes of materials when subjected to external loading. For example, plastic deformation mainly occurs by dislocation slip in the high SFE alloy, however, deformation twins (DTs) tend to be more dominated when the SFE is low [18e20]. What's more, the formation of DTs often causes a good balance of strength and ductility in FCC alloys with low SFE [20e22]. However, the underlying deformation mechanisms originated from DTs have not been well documented especially for the low SFE HEAs, in which abundant DTs and excellent toughness have been reported in various conditions [23e25]. In this paper, tensile behavior and deformation mechanisms of the Al0.1CoCrFeNi HEAs with various grain sizes are studied, in terms of twin thickness, twin spacing, and twin-dislocation interactions. Analyses about mechanical response, Hall-Petch relationship, strain hardening rate and twinning behavior are presented in Results and discussion. 2. Experimental procedures Button ingots with a nominal composition of Al0.1CoCrFeNi (in atomic ratio) are prepared by arc-melting mixtures of the constituent elements (~99.99 wt percent in purity) under a high-purity
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argon atmosphere. Each master ingot is remelted at least five times to ensure its chemical homogeneity, and then cast into a watercooled copper mold with a dimension of 2 10 100 mm3. Then, the 2-mm-thick plates are further cold rolled into 0.5-mm-thick sheets that are subsequently annealed at varying temperatures ranging from 1073 K to 1373 K and followed by water quenching. Tensile specimens with a gauge dimension of 10 2 0.5 mm3, in accordance with the rolling (RD), transverse (TD), and normal directions (ND), respectively, are cut from the cold-rolled and annealed sheets by electrical discharge machine. Tensile tests are performed at a nominal strain rate of 1 103 s1 with an Instron5969 materials testing machine. Phase structures are characterized by X-ray diffraction (XRD) using a PHILIPS APD-10D diffractometer with Cu-Ka radiation. Electron backscattered diffraction (EBSD), with a JEOL JSM-7100F scanning electron microscope (SEM) operated at 20 kV, is applied to microstructural characterizations of the alloys before and after tension. Transmission electron microscope (TEM, JEM-2010F, operated at 200 kV) observation is used to further examine the deformation substructures.
3. Results and discussion Fig. 1(a) shows the XRD patterns for Al0.1CoCrFeNi HEAs at varying annealed statuses. Reflection patterns reveal that only single-phase FCC solid solutions are obtained for the annealed alloys. Fig. 1(b)-1(c) display EBSD maps of fully recrystallized microstructures annealed at 1073 K (1 h) and 1373 K (1 h), respectively, where equiaxed grain structures embedded with many annealing twins are clearly exhibited. Grain sizes are measured through mean-linear-intercept method in which annealing twins are regarded as subgrains. Accordingly, the average grain sizes are obtained as 4.7 mm for 1073 K (1 h), 7.5 mm for 1073 K (3 h), 15.2 mm for 1173 K (1 h), 32.5 mm for 1273 K (1 h), and 59.5 mm
for 1373 K (1 h), respectively. That is to say, with an increase in annealing temperature (or time), an obvious gradient effect on grain size ranging from fine-grained (FG) state to coarse-grained (CG) state occurs in the annealed alloys. Fig. 2(a) shows the tensile true stress - true train curves of Al0.1CoCrFeNi HEAs with various grain sizes. It is noted that with decreasing the grain size, the yielding tensile strength (YTS) and ultimate tensile strength (UTS) increase obviously while the uniform tensile elongation (UTE) decreases modestly. In particular, the 1073 K (1 h)-annealed alloy (~4.7 mm in grain size) realizes a good balance of strength and ductility with a product of ~30 GPa%, revealing pronounced strain hardening and high energy absorption. Fig. 2(b) further exhibits a Hall-Petch type relationship between YTS and grain size that can be expressed as:
sYTS ¼ sYTS;0 þ kYTS d1=2
(1)
where sYTS is the YTS, sYTS;0 is the initial friction stress, kYTS is the strengthening coefficient, and d is the average grain size. Accordingly, the received sYTS;0 and kYTS , extracted from the fitting curve in Fig. 2(b), are 83 MPa and 464 MPa mm1/2, respectively. The kYTS value is very close to the reported 494 MPa mm1/2 for the CoCrFeNiMn HEA also with a single FCC structure [26]. Moreover, strainhardening response extracted from Fig. 2(a), in terms of strainhardening rates (SHR, i.e. ds/dε), as a measure of the ability to strain hardening [27,28], are plotted in Fig. 2(c). Clearly, three representative zones are presented in the SHR diagrams for various grain sizes. The SHR of Zone 1 exhibits a dramatic decrease from over 3 GPa to about 1.5 GPa due to the elastic to plastic transition process. Interestingly, relatively-steady SHR values (i.e. ~1.5e1.7 GPa) can be seen in Zone 2, particularly for the large grain structure, indicating a relatively-stable competitive relationship between strain hardening and strain softening, probably due to the generation of DTs during plastic deformation. Next, Zone 3 again demonstrates a gradually-decreasing trend in SHR that
Fig. 1. (a) XRD patterns for Al0.1CoCrFeNi HEAs at varying annealed statuses; EBSD maps for Al0.1CoCrFeNi HEAs annealed at (b) 1073 K (1 h) and (c) 1373 K (1 h).
Fig. 2. (a) Tensile true stress - true train curves of Al0.1CoCrFeNi HEAs with various grain sizes, (b) yielding tensile strength versus grain size, and (c) strain-hardening rate versus true strain.
J. Yang et al. / Journal of Alloys and Compounds 795 (2019) 269e274
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Fig. 3. EBSD maps of (a) FG (~4.7 mm) and (e) CG (~59.5 mm) HEAs after tensile deformation (the tensile direction//RD), Pole figures (PFs) of (b) FG and (f) CG alloys on the crystal plane of (110) and (111), Inverse pole figures (IPFs) of (c) FG and (g) CG alloys, and orientation distribution function (ODF) sections at F2 ¼ 0 , 45 , and 65 of (d) FG and (h) CG alloys.
corresponds to a dislocation-mediated plastic deformation as reported in twinning-induced plasticity (TWIP) steels [29,30]. It is worth noting that the ultimate SHR value ahead of fracture is up to 1 GPa, revealing excellent strain-hardening capacity for the current alloys. To analyze the texture evolution after tensile deformation,
representative EBSD maps, pole figures (PFs), inverse pole figures (IPFs), and orientation distribution function (ODF) figures obtained from the FG (~4.7 mm) and CG (~59.5 mm) alloys, are presented in Fig. 3. The tensile deformation direction is parallel to RD. It is seen that the grains with modest elongation along tensile direction occur in both alloys, as shown in Fig. 3(a, e). Furthermore, Fig. 3(b) -
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Fig. 4. Volume fractions of the main texture components developed during uniaxial tension.
(d) and (f) - (h) reveal similar deformation textures between FG and CG alloys, and quantitative analyses for the main texture components are presented in Fig. 4. It is noted that in addition to extensive random textures, strong <111> fiber orientations, i.e. Copper ({112}
<111>) and A ({110}<111>) components, occupy the largest volume fraction. The Cube ({001}<100>) and Goss ({110}<100>) components of the <100> fiber orientations, as well as the S component ({123}<634>), exhibit moderate volume percentage. In addition, a relative stable changing trend in Rotated cube ({001}<110>) and Brass ({110}<112>) components is found for both alloys. This feature is similar to the texture evolution of Fee24Mne3Ale2Sie1Nie0.06C steel during uniaxial tension [31]. Nevertheless, as shown in Fig. 4, different from rolling textures [32], the FG and CG alloys exhibit distinct randomization of the texture during tensile testing (the former randomization level is slightly higher than the latter), possibly due to the multiple grain-boundary resistance to dislocation movement upon tensile loading. Moreover, representative bright-field (BF), dark-field (DF), and high resolution TEM (HRTEM) images, as displayed in Fig. 5, further reveal the underlying deformation mechanisms after tensile testing for both FG (~4.7 mm) and CG (~59.5 mm) alloys. As shown in Fig. 5(a, d), high density of dislocation entanglements, in combination with the generation of DT bundles (evidenced by the insets, corresponding selected-area diffraction patterns), indicating that a mixed deformation mode of dislocation slip plus mechanical twinning occurs in the present alloys [20]. In addition, the DF (Fig. 5(b, e)) and HRTEM (Fig. 5(c, f)) images again confirm that nanoscale DTs with several to tens of nanometers are produced upon tensile loading. Furthermore, Fig. 5 (g, h) respectively depict
Fig. 5. (a) Bright-field (BF) image with corresponding selected-area diffraction pattern (SADP), (b) dark-field (DF) image, and (c) high-resolution TEM (HRTEM) image of FG (~4.7 mm) HEA; (d) BF image with corresponding SADP, (e) DF image, and (f) HRTEM image of CG (~59.5 mm) HEA; statistics of twin thickness (g) and twin spacing (h); (i) schematic illustrations of the contribution of twin boundaries (TBs) and grain boundaries (GBs) to the tensile stress of FG and CG HEAs.
J. Yang et al. / Journal of Alloys and Compounds 795 (2019) 269e274
the twin thickness and twin spacing of FG and CG alloys. It is suggested that the latter generally has wider twin thickness and smaller twin spacing relative to the former, indicating more active twin behaviors for the CG alloy [24]. It is widely accepted that a low SFE could facilitate the generation of DTs [20], and for AlxCoCrFeNi HEAs, the SFE parameter is closely related with the Al concentration. For instance, the reported SFE for Al0.1CoCrFeNi is about 6e21 mJ/m2, which is significantly smaller than those of Al0.3CoCrFeNi (~51 mJ/m2) and Al0.6CoCrFeNi (~150 mJ/m2 with various temperatures) alloys [33e35]. This is because Al has an intrinsic SFE of ~86 mJ/m2 and the addition of Al may restrict the formation of partial dislocations due to local chemical ordering and phase separation, thus yielding a high SFE value [33,36]. Therefore, for the latter two alloys, DTs are not easy to be produced upon static loading at room temperature, conversely often observed at extreme conditions such as high-speed impact loading and/or cryogenic temperatures [8,37,38]. In other words, for Al0.1CoCrFeNi, a good combination of strength and ductility, induced by the TWIP effect, can be realized by simple roomtemperature static loading. DTs, acting as barriers to dislocation motion analogous to grain boundaries (GBs), referred to as a dynamic Hall-Petch effect, have been reported in FCC metals and alloys [39e42]. It is believed that twin boundaries (TBs) could reduce the mobility of dislocations and result in greater mechanical forces required to initiate further dislocation movements, thereby rendering a material stronger. Thus, twins may be treated as a compensation for grain boundaries to impede dislocation motion and strengthen the material. Therefore, for the current alloys, the CG sample should have greater compensatory ability than the FG counterpart due to the more active twin behavior. This elucidation is consistent with the strain-hardening response from Zone 2 (a near-platform region) in Fig. 2(c), in which DTs effectively compensate the annihilation of opposite dislocations and build a relative balance between dislocation storage and dislocation recovery [42]. Meanwhile, this balanced work hardening also postpones the onset of necking instability, which contributes to the enhanced ductility for CG alloy [43]. In addition, by carefully comparing the three SHR zones shown in Fig. 2(c), it is easy to be noted that the SHR difference gradually decreases from Zone 1 to Zone 3 with increasing the grain size, indicating that DTs tend to be initiated at large strains, and vice versa, dislocation slip easily appear at small strains for the CG alloy than the FG counterpart. This phenomenon is also observed in the low SFE Cu-Al alloys [20]. Nevertheless, over the entire deformation stage, the FG alloys generally have larger stress and SHR values during stretching (shown in Fig. 2), revealing that GBs still play an important role in blocking dislocation motions and rendering dislocation storage relative to TBs, as illustrated in Fig. 5 (i). 4. Conclusion In conclusion, fully recrystallized single-phase structured Al0.1CoCrFeNi HEAs with different grain sizes are investigated. A good balance of tensile strength and tensile ductility (~30 GPa%) is realized in the FG samples, and the ultimate SHR value ahead of fracture is up to 1 GPa, revealing excellent strain-hardening capacity for the present alloys. Interestingly, a platform region (i.e. Zone 2) occurs in the SHR curves, particularly for the CG alloys, which is in agreement with the more active twinning behavior (wider twin thickness and smaller twin spacing). This balanced work hardening postpones the onset of necking instability and contributes to the enhanced ductility for CG alloys. Combining similar deformation textures, high density of dislocation entanglements and nanoscale DTs, deformation mechanisms of the current alloys are roughly analogous. Nevertheless, the FG alloys with
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more GBs have larger stress values during stretching, revealing that GBs still play an important role in blocking dislocation motions and rendering dislocation storage. Acknowledgements The authors thank the National Natural Science Foundation of China (Nos. 51501123 and 11602158), the Youth Natural Science Foundation of Shanxi (No. 201601D021026), Scientific and Technological Innovation Programs of Higher Education Institutions in Shanxi (No. 2015127), the Top Young Academic Leaders of Shanxi, the “1331 project” fund and Key Innovation Teams of Shanxi Province, the Youth Academic Backbone Cultivation Project from Taiyuan University of Technology, the Sanjin Young Scholars Project of Shanxi Province, and the opening project from the National Key Laboratory for Remanufacturing (No. 61420050204) for financial supports. References [1] J.W. Yeh, S.K. Chen, S.J. Lin, et al., Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes, Adv. Eng. Mater. 6 (2004) 299e303. [2] B. Cantor, I.T.H. Chang, P. Knight, et al., Microstructural development in equiatomic multicomponent alloys, Mater. Sci. Eng. A 375 (2004) 213e218. [3] F. Otto, Y. Yang, H. Bei, et al., Relative effects of enthalpy and entropy on the phase stability of equiatomic high-entropy alloys, Acta Mater. 61 (2013) 2628e2638. [4] T. Yang, Y.L. Zhao, Y. Tong, et al., Multicomponent intermetallic nanoparticles and superb mechanical behaviors of complex alloys, Science 362 (2018) 933e937. [5] O.N. Senkov, G.B. Wilks, D.B. Miracle, et al., Refractory high-entropy alloys, Intermetallics 18 (2010) 1758e1765. [6] C.Y. Hsu, J.W. Yeh, S.K. Chen, et al., Wear resistance and high-temperature compression strength of Fcc CuCoNiCrAl0.5Fe alloy with boron addition, Metall. Mater. Trans. A 35 (2004) 1465e1469. [7] B. Gludovatz, A. Hohenwarter, D. Catoor, et al., A fracture-resistant highentropy alloy for cryogenic applications, Science 345 (2014) 1153e1158. [8] L. Wang, J.W. Qiao, S.G. Ma, et al., Mechanical response and deformation behavior of Al0.6CoCrFeNi high-entropy alloys upon dynamic loading, Mater. Sci. Eng. A 727 (2018) 208e213. [9] D.B. Miracle, O.N. Senkov, A critical review of high entropy alloys (HEAs) and related concepts, Acta Mater. 122 (2017) 448e511. [10] Y. Zhang, T.T. Zuo, Z. Tang, et al., Microstructures and properties of highentropy alloys, Prog. Mater. Sci. 61 (2014) 1e93. [11] J.Y. He, C. Zhu, D.Q. Zhou, et al., Steady state flow of the FeCoNiCrMn high entropy alloy at elevated temperatures, Intermetallics 55 (2014) 9e14. [12] F. Otto, A. Dlouhý, C. Somsen, et al., The influences of temperature and microstructure on the tensile properties of a CoCrFeMnNi high-entropy alloy, Acta Mater. 61 (2013) 5743e5755. [13] N. Stepanov, M. Tikhonovsky, N. Yurchenko, et al., Effect of cryo-deformation on structure and properties of CoCrFeNiMn high-entropy alloy, Intermetallics 59 (2015) 8e17. €lker, et al., Mechanical properties, micro[14] B. Schuh, F. Mendez-Martin, B. Vo structure and thermal stability of a nanocrystalline CoCrFeMnNi high-entropy alloy after severe plastic deformation, Acta Mater. 96 (2015) 258e268. [15] F. Otto, A. Dlouhy, C. Somsen, et al., The influences of temperature and microstructure on the tensile properties of a CoCrFeMnNi high-entropy alloy, Acta Mater. 61 (2013) 5743e5755. [16] G. Chen, L.T. Li, J.W. Qiao, et al., Gradient hierarchical grain structures of Al0.1CoCrFeNi high-entropy alloys through dynamic torsion, Mater. Lett. 238 (2019) 163e166. [17] M. Komarasamy, N. Kumar, Z. Tang, et al., Effect of microstructure on the deformation mechanism of friction stir-processed Al0.1CoCrFeNi high entropy alloy, Mater. Res. Lett. 3 (2015) 30e34. [18] A.J. Zaddach, C. Niu, C.C. Koch, et al., Mechanical properties and stacking fault energies of NiFeCrCoMn high-entropy alloy, JOM 65 (2013) 1780e1789. [19] S. Huang, W. Li, S. Lu, et al., Temperature dependent stacking fault energy of FeCrCoNiMn high entropy alloy, Scripta Mater. 108 (2015) 44e47. [20] Y.Z. Tian, L.J. Zhao, N. Park, et al., Revealing the deformation mechanisms of Cu-Al alloys with high strength and good ductility, Acta Mater. 110 (2016) 61e72. [21] Y. Deng, C.C. Tasan, K.G. Pradeep, et al., Design of a twinning-induced plasticity high entropy alloy, Acta Mater. 94 (2015) 124e133. [22] Z. Wu, C.M. Parish, H. Bei, Nano-twin mediated plasticity in carbon-containing FeNiCoCrMn high entropy alloys, J. Alloys Compd. 647 (2015) 815e822. [23] S.J. Sun, Y.Z. Tian, H.R. Lin, et al., Transition of twinning behavior in CoCrFeMnNi high entropy alloy with grain refinement, Mater. Sci. Eng. A 712 (2018) 603e607.
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[24] S.W. Wu, G. Wang, J. Yi, et al., Strong grain-size effect on deformation twinning of an Al0.1CoCrFeNi high-entropy alloy, Mater. Res. Lett. 5 (2016) 276e283. [25] G. Frommeyer, U. Brux, P. Neumann, Supra-ductile and high-strength manganese-TRIP/TWIP steels for high energy absorption purposes, ISIJ Int. 43 (2003) 438e446. [26] W.H. Liu, Y. Wu, J.Y. He, et al., Grain growth and the Hall-Petch relationship in a high-entropy FeCrNiCoMn alloy, Scripta Mater. 68 (2013) 526e529. [27] U.F. Kocks, H. Mecking, Physics and phenomenology of strain hardening: the FCC case, Prog. Mater. Sci. 48 (2003) 171e273. [28] H. Mecking, U.F. Kocks, Kinetics of flow and strain-hardening, Acta Metall. 29 (1981) 1865e1875. [29] I. Gutierrez-Urrutia, S. Zaefferer, D. Raabe, The effect of grain size and grain orientation on deformation twinning in a Fe-22 wt.% Mn-0.6 wt.% C TWIP steel, Mater. Sci. Eng. A 527 (2010) 3552e3560. [30] I. Gutierrez-Urrutia, D. Raabe, Dislocation and twin substructure evolution during strain hardening of an Fe-22 wt.% Mn-0.6 wt.% C TWIP steel observed by electron channeling contrast imaging, Acta Mater. 59 (2011) 6449e6462. [31] A.A. Saleh, E.V. Pereloma, A.A. Gazder, Microstructure and texture evolution in a twinning-induced-plasticity steel during uniaxial tension, Acta Mater. 61 (2013) 2671e2691. [32] C. Haase, L.A. Barrales-Mora, Influence of deformation and annealing twinning on the microstructure and texture evolution of face-centered cubic highentropy alloys, Acta Mater. 150 (2018) 88e103. [33] X.D. Xu, P. Liu, Z. Tang, et al., Transmission electron microscopy characterization of dislocation structure in a face-centered cubic high-entropy alloy
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Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Revealing the Hall-Petch relationship of Al0.1CoCrFeNi high-entropy alloy and its deformation mechanisms J. Yang a, b, J.W. Qiao c, S.G. Ma a, b, *, G.Y. Wu a, b, D. Zhao a, b, Z.H. Wang a, b, ** a
Institute of Applied Mechanics, College of Mechanical and Vehicle Engineering, Taiyuan University of Technology, Taiyuan 030024, China Shanxi Key Laboratory of Material Strength and Structural Impact, Taiyuan University of Technology, Taiyuan 030024, China c Institute of High-entropy Alloys, College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan 030024, China b
a r t i c l e i n f o
a b s t r a c t
Article history: Received 22 February 2019 Received in revised form 28 April 2019 Accepted 30 April 2019 Available online 4 May 2019
Al0.1CoCrFeNi high-entropy alloys with fully recrystallized single-phase structures and various grain sizes ranging from 4.7 to 59.5 mm are fabricated through cold rolling and subsequent annealing. Tensile tests reveal that the yielding strength has a linear relationship with the inverse square root of grain size, i.e. a Hall-Petch type relationship, shown as sYTS ¼ 83 þ 464d1=2 . Moreover, similar deformation textures, i.e. strong <111> fiber, moderate <100> fiber, and relative stable Rotated cube and Brass components between fine-grained (FG) and coarse-grained (CG) alloys, in combination with TEM analyses, suggest analogous deformation mechanisms for both alloys and more active twinning behavior of the CG alloy. © 2019 Elsevier B.V. All rights reserved.
Keywords: High-entropy alloy Grain size Hall-Petch relationship Deformation mechanism
1. Introduction High-entropy alloys (HEAs) have drawn great interests in the materials field due to amazing properties, such as high strength and good ductility, potential low- and/or high-temperature properties, good wear and fatigue resistances, and excellent fracture toughness [1e10]. Distinct from traditional alloys that are based on one or two major elements, the HEAs usually contain five or more principal elements in equimolar or near-equimolar ratio. The unique alloy design concept often yields the formation of random solid solutions upon solidification, rather than complex phases or intermetallic compounds, probably due to the unique high mixing entropy effect [1e3]. Therefore, the HEAs exhibit a good combination of great academic values and potential application prospects. In fact, a large number of investigations on phase formation, microstructure evolution, mechanical property and deformation mechanisms of the HEAs have been reported [4e11]. Among various HEAs, the alloys with a face-centered cubic (FCC)
* Corresponding author. Institute of Applied Mechanics, College of Mechanical and Vehicle Engineering, Taiyuan University of Technology, Taiyuan 030024, China. ** Corresponding author. Institute of Applied Mechanics, College of Mechanical and Vehicle Engineering, Taiyuan University of Technology, Taiyuan 030024, China. E-mail addresses: [email protected] (S.G. Ma), [email protected] (Z.H. Wang). https://doi.org/10.1016/j.jallcom.2019.04.333 0925-8388/© 2019 Elsevier B.V. All rights reserved.
solid-solution structure, especially for the low stacking-faultenergy (SFE) systems, have been the current research hotspots [11e19]. Plenty studies have shown that the SFE could significantly affect the deformation modes of materials when subjected to external loading. For example, plastic deformation mainly occurs by dislocation slip in the high SFE alloy, however, deformation twins (DTs) tend to be more dominated when the SFE is low [18e20]. What's more, the formation of DTs often causes a good balance of strength and ductility in FCC alloys with low SFE [20e22]. However, the underlying deformation mechanisms originated from DTs have not been well documented especially for the low SFE HEAs, in which abundant DTs and excellent toughness have been reported in various conditions [23e25]. In this paper, tensile behavior and deformation mechanisms of the Al0.1CoCrFeNi HEAs with various grain sizes are studied, in terms of twin thickness, twin spacing, and twin-dislocation interactions. Analyses about mechanical response, Hall-Petch relationship, strain hardening rate and twinning behavior are presented in Results and discussion. 2. Experimental procedures Button ingots with a nominal composition of Al0.1CoCrFeNi (in atomic ratio) are prepared by arc-melting mixtures of the constituent elements (~99.99 wt percent in purity) under a high-purity
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argon atmosphere. Each master ingot is remelted at least five times to ensure its chemical homogeneity, and then cast into a watercooled copper mold with a dimension of 2 10 100 mm3. Then, the 2-mm-thick plates are further cold rolled into 0.5-mm-thick sheets that are subsequently annealed at varying temperatures ranging from 1073 K to 1373 K and followed by water quenching. Tensile specimens with a gauge dimension of 10 2 0.5 mm3, in accordance with the rolling (RD), transverse (TD), and normal directions (ND), respectively, are cut from the cold-rolled and annealed sheets by electrical discharge machine. Tensile tests are performed at a nominal strain rate of 1 103 s1 with an Instron5969 materials testing machine. Phase structures are characterized by X-ray diffraction (XRD) using a PHILIPS APD-10D diffractometer with Cu-Ka radiation. Electron backscattered diffraction (EBSD), with a JEOL JSM-7100F scanning electron microscope (SEM) operated at 20 kV, is applied to microstructural characterizations of the alloys before and after tension. Transmission electron microscope (TEM, JEM-2010F, operated at 200 kV) observation is used to further examine the deformation substructures.
3. Results and discussion Fig. 1(a) shows the XRD patterns for Al0.1CoCrFeNi HEAs at varying annealed statuses. Reflection patterns reveal that only single-phase FCC solid solutions are obtained for the annealed alloys. Fig. 1(b)-1(c) display EBSD maps of fully recrystallized microstructures annealed at 1073 K (1 h) and 1373 K (1 h), respectively, where equiaxed grain structures embedded with many annealing twins are clearly exhibited. Grain sizes are measured through mean-linear-intercept method in which annealing twins are regarded as subgrains. Accordingly, the average grain sizes are obtained as 4.7 mm for 1073 K (1 h), 7.5 mm for 1073 K (3 h), 15.2 mm for 1173 K (1 h), 32.5 mm for 1273 K (1 h), and 59.5 mm
for 1373 K (1 h), respectively. That is to say, with an increase in annealing temperature (or time), an obvious gradient effect on grain size ranging from fine-grained (FG) state to coarse-grained (CG) state occurs in the annealed alloys. Fig. 2(a) shows the tensile true stress - true train curves of Al0.1CoCrFeNi HEAs with various grain sizes. It is noted that with decreasing the grain size, the yielding tensile strength (YTS) and ultimate tensile strength (UTS) increase obviously while the uniform tensile elongation (UTE) decreases modestly. In particular, the 1073 K (1 h)-annealed alloy (~4.7 mm in grain size) realizes a good balance of strength and ductility with a product of ~30 GPa%, revealing pronounced strain hardening and high energy absorption. Fig. 2(b) further exhibits a Hall-Petch type relationship between YTS and grain size that can be expressed as:
sYTS ¼ sYTS;0 þ kYTS d1=2
(1)
where sYTS is the YTS, sYTS;0 is the initial friction stress, kYTS is the strengthening coefficient, and d is the average grain size. Accordingly, the received sYTS;0 and kYTS , extracted from the fitting curve in Fig. 2(b), are 83 MPa and 464 MPa mm1/2, respectively. The kYTS value is very close to the reported 494 MPa mm1/2 for the CoCrFeNiMn HEA also with a single FCC structure [26]. Moreover, strainhardening response extracted from Fig. 2(a), in terms of strainhardening rates (SHR, i.e. ds/dε), as a measure of the ability to strain hardening [27,28], are plotted in Fig. 2(c). Clearly, three representative zones are presented in the SHR diagrams for various grain sizes. The SHR of Zone 1 exhibits a dramatic decrease from over 3 GPa to about 1.5 GPa due to the elastic to plastic transition process. Interestingly, relatively-steady SHR values (i.e. ~1.5e1.7 GPa) can be seen in Zone 2, particularly for the large grain structure, indicating a relatively-stable competitive relationship between strain hardening and strain softening, probably due to the generation of DTs during plastic deformation. Next, Zone 3 again demonstrates a gradually-decreasing trend in SHR that
Fig. 1. (a) XRD patterns for Al0.1CoCrFeNi HEAs at varying annealed statuses; EBSD maps for Al0.1CoCrFeNi HEAs annealed at (b) 1073 K (1 h) and (c) 1373 K (1 h).
Fig. 2. (a) Tensile true stress - true train curves of Al0.1CoCrFeNi HEAs with various grain sizes, (b) yielding tensile strength versus grain size, and (c) strain-hardening rate versus true strain.
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Fig. 3. EBSD maps of (a) FG (~4.7 mm) and (e) CG (~59.5 mm) HEAs after tensile deformation (the tensile direction//RD), Pole figures (PFs) of (b) FG and (f) CG alloys on the crystal plane of (110) and (111), Inverse pole figures (IPFs) of (c) FG and (g) CG alloys, and orientation distribution function (ODF) sections at F2 ¼ 0 , 45 , and 65 of (d) FG and (h) CG alloys.
corresponds to a dislocation-mediated plastic deformation as reported in twinning-induced plasticity (TWIP) steels [29,30]. It is worth noting that the ultimate SHR value ahead of fracture is up to 1 GPa, revealing excellent strain-hardening capacity for the current alloys. To analyze the texture evolution after tensile deformation,
representative EBSD maps, pole figures (PFs), inverse pole figures (IPFs), and orientation distribution function (ODF) figures obtained from the FG (~4.7 mm) and CG (~59.5 mm) alloys, are presented in Fig. 3. The tensile deformation direction is parallel to RD. It is seen that the grains with modest elongation along tensile direction occur in both alloys, as shown in Fig. 3(a, e). Furthermore, Fig. 3(b) -
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Fig. 4. Volume fractions of the main texture components developed during uniaxial tension.
(d) and (f) - (h) reveal similar deformation textures between FG and CG alloys, and quantitative analyses for the main texture components are presented in Fig. 4. It is noted that in addition to extensive random textures, strong <111> fiber orientations, i.e. Copper ({112}
<111>) and A ({110}<111>) components, occupy the largest volume fraction. The Cube ({001}<100>) and Goss ({110}<100>) components of the <100> fiber orientations, as well as the S component ({123}<634>), exhibit moderate volume percentage. In addition, a relative stable changing trend in Rotated cube ({001}<110>) and Brass ({110}<112>) components is found for both alloys. This feature is similar to the texture evolution of Fee24Mne3Ale2Sie1Nie0.06C steel during uniaxial tension [31]. Nevertheless, as shown in Fig. 4, different from rolling textures [32], the FG and CG alloys exhibit distinct randomization of the texture during tensile testing (the former randomization level is slightly higher than the latter), possibly due to the multiple grain-boundary resistance to dislocation movement upon tensile loading. Moreover, representative bright-field (BF), dark-field (DF), and high resolution TEM (HRTEM) images, as displayed in Fig. 5, further reveal the underlying deformation mechanisms after tensile testing for both FG (~4.7 mm) and CG (~59.5 mm) alloys. As shown in Fig. 5(a, d), high density of dislocation entanglements, in combination with the generation of DT bundles (evidenced by the insets, corresponding selected-area diffraction patterns), indicating that a mixed deformation mode of dislocation slip plus mechanical twinning occurs in the present alloys [20]. In addition, the DF (Fig. 5(b, e)) and HRTEM (Fig. 5(c, f)) images again confirm that nanoscale DTs with several to tens of nanometers are produced upon tensile loading. Furthermore, Fig. 5 (g, h) respectively depict
Fig. 5. (a) Bright-field (BF) image with corresponding selected-area diffraction pattern (SADP), (b) dark-field (DF) image, and (c) high-resolution TEM (HRTEM) image of FG (~4.7 mm) HEA; (d) BF image with corresponding SADP, (e) DF image, and (f) HRTEM image of CG (~59.5 mm) HEA; statistics of twin thickness (g) and twin spacing (h); (i) schematic illustrations of the contribution of twin boundaries (TBs) and grain boundaries (GBs) to the tensile stress of FG and CG HEAs.
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the twin thickness and twin spacing of FG and CG alloys. It is suggested that the latter generally has wider twin thickness and smaller twin spacing relative to the former, indicating more active twin behaviors for the CG alloy [24]. It is widely accepted that a low SFE could facilitate the generation of DTs [20], and for AlxCoCrFeNi HEAs, the SFE parameter is closely related with the Al concentration. For instance, the reported SFE for Al0.1CoCrFeNi is about 6e21 mJ/m2, which is significantly smaller than those of Al0.3CoCrFeNi (~51 mJ/m2) and Al0.6CoCrFeNi (~150 mJ/m2 with various temperatures) alloys [33e35]. This is because Al has an intrinsic SFE of ~86 mJ/m2 and the addition of Al may restrict the formation of partial dislocations due to local chemical ordering and phase separation, thus yielding a high SFE value [33,36]. Therefore, for the latter two alloys, DTs are not easy to be produced upon static loading at room temperature, conversely often observed at extreme conditions such as high-speed impact loading and/or cryogenic temperatures [8,37,38]. In other words, for Al0.1CoCrFeNi, a good combination of strength and ductility, induced by the TWIP effect, can be realized by simple roomtemperature static loading. DTs, acting as barriers to dislocation motion analogous to grain boundaries (GBs), referred to as a dynamic Hall-Petch effect, have been reported in FCC metals and alloys [39e42]. It is believed that twin boundaries (TBs) could reduce the mobility of dislocations and result in greater mechanical forces required to initiate further dislocation movements, thereby rendering a material stronger. Thus, twins may be treated as a compensation for grain boundaries to impede dislocation motion and strengthen the material. Therefore, for the current alloys, the CG sample should have greater compensatory ability than the FG counterpart due to the more active twin behavior. This elucidation is consistent with the strain-hardening response from Zone 2 (a near-platform region) in Fig. 2(c), in which DTs effectively compensate the annihilation of opposite dislocations and build a relative balance between dislocation storage and dislocation recovery [42]. Meanwhile, this balanced work hardening also postpones the onset of necking instability, which contributes to the enhanced ductility for CG alloy [43]. In addition, by carefully comparing the three SHR zones shown in Fig. 2(c), it is easy to be noted that the SHR difference gradually decreases from Zone 1 to Zone 3 with increasing the grain size, indicating that DTs tend to be initiated at large strains, and vice versa, dislocation slip easily appear at small strains for the CG alloy than the FG counterpart. This phenomenon is also observed in the low SFE Cu-Al alloys [20]. Nevertheless, over the entire deformation stage, the FG alloys generally have larger stress and SHR values during stretching (shown in Fig. 2), revealing that GBs still play an important role in blocking dislocation motions and rendering dislocation storage relative to TBs, as illustrated in Fig. 5 (i). 4. Conclusion In conclusion, fully recrystallized single-phase structured Al0.1CoCrFeNi HEAs with different grain sizes are investigated. A good balance of tensile strength and tensile ductility (~30 GPa%) is realized in the FG samples, and the ultimate SHR value ahead of fracture is up to 1 GPa, revealing excellent strain-hardening capacity for the present alloys. Interestingly, a platform region (i.e. Zone 2) occurs in the SHR curves, particularly for the CG alloys, which is in agreement with the more active twinning behavior (wider twin thickness and smaller twin spacing). This balanced work hardening postpones the onset of necking instability and contributes to the enhanced ductility for CG alloys. Combining similar deformation textures, high density of dislocation entanglements and nanoscale DTs, deformation mechanisms of the current alloys are roughly analogous. Nevertheless, the FG alloys with
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more GBs have larger stress values during stretching, revealing that GBs still play an important role in blocking dislocation motions and rendering dislocation storage. Acknowledgements The authors thank the National Natural Science Foundation of China (Nos. 51501123 and 11602158), the Youth Natural Science Foundation of Shanxi (No. 201601D021026), Scientific and Technological Innovation Programs of Higher Education Institutions in Shanxi (No. 2015127), the Top Young Academic Leaders of Shanxi, the “1331 project” fund and Key Innovation Teams of Shanxi Province, the Youth Academic Backbone Cultivation Project from Taiyuan University of Technology, the Sanjin Young Scholars Project of Shanxi Province, and the opening project from the National Key Laboratory for Remanufacturing (No. 61420050204) for financial supports. References [1] J.W. Yeh, S.K. Chen, S.J. Lin, et al., Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes, Adv. Eng. Mater. 6 (2004) 299e303. [2] B. Cantor, I.T.H. Chang, P. Knight, et al., Microstructural development in equiatomic multicomponent alloys, Mater. Sci. Eng. A 375 (2004) 213e218. [3] F. Otto, Y. Yang, H. Bei, et al., Relative effects of enthalpy and entropy on the phase stability of equiatomic high-entropy alloys, Acta Mater. 61 (2013) 2628e2638. [4] T. Yang, Y.L. Zhao, Y. Tong, et al., Multicomponent intermetallic nanoparticles and superb mechanical behaviors of complex alloys, Science 362 (2018) 933e937. [5] O.N. Senkov, G.B. Wilks, D.B. Miracle, et al., Refractory high-entropy alloys, Intermetallics 18 (2010) 1758e1765. [6] C.Y. Hsu, J.W. Yeh, S.K. Chen, et al., Wear resistance and high-temperature compression strength of Fcc CuCoNiCrAl0.5Fe alloy with boron addition, Metall. Mater. Trans. A 35 (2004) 1465e1469. [7] B. Gludovatz, A. Hohenwarter, D. Catoor, et al., A fracture-resistant highentropy alloy for cryogenic applications, Science 345 (2014) 1153e1158. [8] L. Wang, J.W. Qiao, S.G. Ma, et al., Mechanical response and deformation behavior of Al0.6CoCrFeNi high-entropy alloys upon dynamic loading, Mater. Sci. Eng. A 727 (2018) 208e213. [9] D.B. Miracle, O.N. Senkov, A critical review of high entropy alloys (HEAs) and related concepts, Acta Mater. 122 (2017) 448e511. [10] Y. Zhang, T.T. Zuo, Z. Tang, et al., Microstructures and properties of highentropy alloys, Prog. Mater. Sci. 61 (2014) 1e93. [11] J.Y. He, C. Zhu, D.Q. Zhou, et al., Steady state flow of the FeCoNiCrMn high entropy alloy at elevated temperatures, Intermetallics 55 (2014) 9e14. [12] F. Otto, A. Dlouhý, C. Somsen, et al., The influences of temperature and microstructure on the tensile properties of a CoCrFeMnNi high-entropy alloy, Acta Mater. 61 (2013) 5743e5755. [13] N. Stepanov, M. Tikhonovsky, N. Yurchenko, et al., Effect of cryo-deformation on structure and properties of CoCrFeNiMn high-entropy alloy, Intermetallics 59 (2015) 8e17. €lker, et al., Mechanical properties, micro[14] B. Schuh, F. Mendez-Martin, B. Vo structure and thermal stability of a nanocrystalline CoCrFeMnNi high-entropy alloy after severe plastic deformation, Acta Mater. 96 (2015) 258e268. [15] F. Otto, A. Dlouhy, C. Somsen, et al., The influences of temperature and microstructure on the tensile properties of a CoCrFeMnNi high-entropy alloy, Acta Mater. 61 (2013) 5743e5755. [16] G. Chen, L.T. Li, J.W. Qiao, et al., Gradient hierarchical grain structures of Al0.1CoCrFeNi high-entropy alloys through dynamic torsion, Mater. Lett. 238 (2019) 163e166. [17] M. Komarasamy, N. Kumar, Z. Tang, et al., Effect of microstructure on the deformation mechanism of friction stir-processed Al0.1CoCrFeNi high entropy alloy, Mater. Res. Lett. 3 (2015) 30e34. [18] A.J. Zaddach, C. Niu, C.C. Koch, et al., Mechanical properties and stacking fault energies of NiFeCrCoMn high-entropy alloy, JOM 65 (2013) 1780e1789. [19] S. Huang, W. Li, S. Lu, et al., Temperature dependent stacking fault energy of FeCrCoNiMn high entropy alloy, Scripta Mater. 108 (2015) 44e47. [20] Y.Z. Tian, L.J. Zhao, N. Park, et al., Revealing the deformation mechanisms of Cu-Al alloys with high strength and good ductility, Acta Mater. 110 (2016) 61e72. [21] Y. Deng, C.C. Tasan, K.G. Pradeep, et al., Design of a twinning-induced plasticity high entropy alloy, Acta Mater. 94 (2015) 124e133. [22] Z. Wu, C.M. Parish, H. Bei, Nano-twin mediated plasticity in carbon-containing FeNiCoCrMn high entropy alloys, J. Alloys Compd. 647 (2015) 815e822. [23] S.J. Sun, Y.Z. Tian, H.R. Lin, et al., Transition of twinning behavior in CoCrFeMnNi high entropy alloy with grain refinement, Mater. Sci. Eng. A 712 (2018) 603e607.
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