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Superior tensile properties of Al0.3CoCrFeNi high entropy alloys with B2 precipitated phases at room and cryogenic temperatures
Materials Science & Engineering A 767 (2019) 138424
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Superior tensile properties of Al0.3CoCrFeNi high entropy alloys with B2 precipitated phases at room and cryogenic temperatures
T
Q. Lia,b, T.W. Zhanga,b, J.W. Qiaoc, S.G. Maa,b,∗, D. Zhaoa,b, P. Lua,b, B. Xud, Z.H. Wanga,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 d School of Mechanics, Civil Engineering and Architecture, Northwestern Polytechnical University, Xian, 710129, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: High-entropy alloys Mechanical properties Cryogenic temperature Deformation twins Work hardening
Al0.3CoCrFeNi high-entropy alloys (HEAs) with B2 ordered phase mainly precipitated at grain boundaries are obtained by cold rolling, annealing and aging treatment. Tensile experiments show that both CRSA (annealed at 1100 °C for 1 h) and CRSA-600-24 (annealed at 1100 °C for 1 h and subsequently aging treatment at 600 °C for 24 h) samples have excellent combinations of yield strength, tensile strength and tensile elongation at room temperature, and more importantly, simultaneous enhancements in strength and ductility at liquid nitrogen (77 K) are attained for the latter condition. Analyses reveal that after aging treatment, the precipitation of B2 phase with ultrafine-grained or even nanoscale sizes, as well as moderate grain refinement, yields the significant increase in tensile strength and fascinating tensile plasticity. By transmission electron microscopy (TEM) characterization, it is found that the deformation mode dominated by dislocation glide at room temperature is transformed to that combined with dislocation glide plus nanoscale twinning at 77 K. Moreover, the twinninginduced work hardening renders the onset of necking instability to a higher strain/stress value, which therefore improves the strength and ductility simultaneously.
1. Introduction High-entropy alloys (HEAs) with multi-principal elements are generally composed of five or more elements with concentrations in the range of 5–35 at. % [1]. The HEAs tend to form simple solid solution phases, such as face-centered cubic (fcc), body-centered cubic (bcc) or hexagonal close-paced (hcp) structures, due to high mixing entropy reducing the Gibbs free energy and thus broadening the solid-solution limits especially at high temperatures. Moreover, extensive interests have been received in the HEAs field owing to their exceptional properties, such as high strength and good ductility, excellent resistances to oxidation, corrosion, fatigue and radiation, outstanding fracture toughness [2,3]. These properties qualify HEAs as potential candidates for the future engineering materials, to meet the demanding requirements for selected extreme applications, primarily in the automotive, defensive, marine and aerospace industries. It is known that brittlement is generally problematic when metals or
alloys are subjected to extreme conditions such as cryogenic temperatures or impact loadings. For instance, the low-temperature strength is obviously higher than that at room temperature, whereas the elongation shows an opposite trend in some austenitic high-Mn steels [4]. Amazingly, the recently-reported HEAs especially having low stackingfault energy (SFE) [5], effectively overcome the strength-ductility trade-off due to the nanotwinning formation during plastic deformation. This phenomenon is known as the “twinning-induced plasticity” effect (TWIP) [6]. Recently, Li et al. found that the incidence of deformation twins depended on the SFE and their frequency increased with the decrease in SFE [7]. Moreover, for cubic crystals, it is found that the SFE also decreases with decreasing the temperature [8]. Therefore, nanotwins may become more prevalent as cubic-structured and low-SFE HEAs are subjected to low-temperature loading. As a result, simultaneous improvement of strength and ductility could be realized due to the prominent TWIP effect. In fact, it has been reported that the fracture strength and tensile elongation of fcc-structured
∗ 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.msea.2019.138424 Received 26 June 2019; Received in revised form 14 September 2019; Accepted 14 September 2019 Available online 16 September 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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Fig. 1. SEM images of Al0.3CoCrFeNi HEAs: (a) CRSA sample, (b) CRSA-600-24 sample, and the insets are partially enlarged views of the (a, b) diagrams; (c) statistical results of the size distribution of B2 phase.
tests were performed with an Instron 5969 machine at a strain rate of 1 × 10−3/s and two temperatures of 298 K and 77 K. For the 77 K experiments, both the grips and tensile specimens were completely immersed in liquid nitrogen for about 20 min before the tensile tests, and the specimens were always submerging in the liquid nitrogen environment during tensile tests to ensure that the tensile specimens are held at the target temperature. Tensile tests at all temperatures were repeated at least three times. Microstructure evolution and phase constitution of the CRSA and CRSA-600-24 specimens before and after deformation were identified by scanning electron microscope (SEM, JEOM JSM-7100F) and transmission electron microscopy (TEM, JEM2010F). The chemical composition of the phase is analyzed by an attached energy dispersive spectrometer (EDS).
Al0.1CoCrFeNi HEAs respectively reached 635 MPa and 58.8% at 298 K, while at 77 K, the fracture strength increased to 1042 MPa, and what's more amazing was that the plasticity had not decreased, but increased to 81.6% [9]. In addition, He et al. [10] reported that the precipitation of secondary phase significantly improved the tensile strength in (CoCrFeNi)94Ti2Al4 HEA. In particular, the precipitation-hardened/toughed Al0.3CoCrFeNi HEAs have been widely reported, in which multiscalesized precipitates yield a good balance of strength and ductility [7]. Therefore, based on the above analyses, the fcc-structured Al0.3CoCrFeNi HEA with the precipitation of B2 ordered phases are obtained by cold rolling and subsequent aging, accompanied by room temperature and liquid nitrogen tensile tests. Finally, microstructure evolution, tensile properties, and deformation mechanisms at room and low temperatures are presented in results and discussion.
3. Results 2. Materials and methods
3.1. Initial microstructure analyses
Ingots with a nominal composition of Al0.3CoCrFeNi (the subscript is atomic ratio) were synthesized via arc-melting a mixture of highpurity elements (> 99.9 wt %) in a Ti-gettered argon atmosphere. The ingots were flipped and re-melted at least five times to ensure the homogeneity of chemical composition, and then suck into a watercooled copper mould to obtain plates with a dimension of 50 mm (length) × 10 mm (width) × 2 mm (depth). The actual chemical composition of the alloy plates is determined by EDS analysis to be 6.07Al17.52Co-18.02Cr-17.54Fe-17.16Ni (wt. %). The plates were cold rolled into a final thickness of 0.8 mm (~60% thickness reduction), annealed at 1100 °C for 1 h, and then water quenched (referred to as CRSA), and moreover, are further subjected to aging treatment at 600 °C for 24 h followed by water quenching (referred to as CRSA-600-24). Flat dog-bone-shaped tensile specimens were machined from alloy sheets by electric discharge machining with their longitudinal axes along with the rolling direction. The gage length and width of the tensile specimens are 10 mm and 4 mm, respectively. Uniaxial tensile
The initial microstructures of CRSA and CRSA-600-24 specimens are presented in Fig. 1. It is shown that both samples exhibit polycrystalline microstructures with irregular grains and visible annealed twins. Specifically, as shown in Fig. 1(c), the grain structure appears to be more uniform in CRSA-600-24 than CRSA, and the average grain size is about 20 μm. It is known that along with the recrystallization process during annealing (1100 °C for 1 h), dislocation recovery significantly occurs and strain energy substantially releases, which result in the minor driving force for the grain-boundary migration. Therefore, the speed of grain growth is very slow when samples are subjected to low-temperature aging treatment (600 °C for 24 h). In addition, recrystallization may also occur in the portion that has not been recrystallized during annealing, thus leading to the microstructure being relatively uniform. Interestingly, seen from Fig. 1(b, d), very fine or even nanoscale precipitates in these two samples are precipitated near or at the grain boundaries (as grain boundaries can be regarded as the diffusion path of atoms that effectively overcome the slow diffusion effect of HEAs 2
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Fig. 2. (a) A TEM BF image from the CRSA sample, (b) an SAED pattern obtained from the matrix of (a), (c) an SAED pattern obtained from the precipitation of (a), (d) a TEM BF image from the CRSA-600-24 sample, (e) an SAED pattern obtained from the matrix of (d), and (f) an SAED pattern from the precipitation of (d). (the inserted green tables in Figs. 2(b, d) are the EDS results of the precipitation). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
24 h, with lower annealing temperature and prolonged annealing time, more Al and Ni atoms tend to be precipitated due to the enhanced phase instability, thereby forming B2 phase of the lath-like particles. In addition, recrystallized grain boundaries or large angle grain boundaries can accelerate the diffusion ability of atoms [12], which can effectively overcome the slow diffusion ability of HEAs [13]. Thus, the B2 phase tends to be precipitated at the grain boundaries.
[10,11]), and the amount of precipitates for the latter condition (~7% in volume fraction) is much higher than the former condition (~2%) due to the additional aging effect. Moreover, Fig. 1(e) depicts the size distribution of the precipitates for both samples. One can see that a wide size distribution ranging from ~50 to 450 nm is attained, in which over 50% of precipitates have the size of 50–100 nm. Next, TEM images of CRSA and CRSA-600-24 samples, as shown in Fig. 2, further manifest the feature of the precipitation as well as the fcc matrix for both samples, in terms of bright-field (BF) images with attached EDS analyses (Fig. 2(a, d)) and selected-area-electron diffraction (SAED) patterns (Fig. 2(b-c, e-f)). It is noted that the NiAl-type B2 (ordered bcc) precipitates, enriched in Al and Ni elements as compared to the fcc matrix, are mainly in forms of irregular or lath-like particles. This is because the present solid-solution alloy is in a saturated state, and Al and Ni elements have no time to nucleate and precipitate due to high-temperature melting and rapid solidification, thus forming supersaturated solid solutions with simple fcc structure. When the alloys are annealed at 1100 °C for 1 h, minor B2 phase is precipitated due to the very negative enthalpy and larger size difference of Al and Ni atoms as compared to other constitutive elements. After aging at 600 °C for
3.2. Mechanical properties Fig. 3(a) shows engineering tensile stress-strain curves of the CRSA and CRSA-600-24 samples at 298 K. It is found that for the aged alloy, the yield tensile strength and ultimate tensile strength increased from 330 MPa to 460 MPa and 746 MPa–913 MPa, respectively, as compared with the annealed counterpart. Although the plasticity decreased slightly, it still reached 62.5%. Moreover, Fig. 3(d) shows engineering tensile stress-strain curves of the CRSA-600-24 samples at 298 K and 77 K. For comparison, from 298 K to 77 K, the yield strength increased from 460 MPa to 730 MPa, nearly increased by 60%. Meanwhile, the ultimate tensile strength increased from 913 MPa to 1355 MPa, nearly 3
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Fig. 3. Engineering strain-stress curves (a) and true stress-strain curves (b), as well as work hardening rate-true strain curves (c) of CRSA and CRSA-600-24 samples in tension mode at 298 K; (d) engineering strain-stress curves of CRSA-600-24 samples in tension mode at 298 K and 77 K.
Fig. 4. The maps of tensile strength versus ductility (a) and yield strength versus product strength and elongation (b) of various advanced steels or alloys including CRSA-600-24 HEAs.
austenitic stainless steels [40,41], titanium alloys [42] and AlxCoCrFeNi HEAs [43–48]. It is seen that in comparison with other materials, the present alloys tend to be located at the top right corner of the maps and exhibit an excellent combination of strength and ductility particularly upon 77 K loading. In fact, the product of elongation and strength for CRSA-600-24 at 298 K and 77 K reach 57 GPa% and 89 GPa%, respectively. Moreover, the fracture-surface morphology from CRSA-600-24
increased by 50%. It is even amazing that the plasticity of the samples does not decrease but increases from 62.5% to 66.2%. Specifically, Fig. 4 compares the elongation versus strength and the strength-elongation product versus yield strength of CRSA-600-24 and other alloys, e.g., magnesium alloys [14–17], aluminum alloys [18–22], interstitial-free steels [23–27], transformation induced plasticity steels [28–32], dual-phase steels [33–37], ferritic stainless steels [38,39], 4
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Fig. 5. The dimpled fracture surface of CRSA-600-24 samples in tension mode at 298 K (a) and 77 K (b).
Fig. 6. True stress-true strain curves (a) and work hardening rate versus true strain curves (b) as well as work hardening rate versus true stress curves (c) of the CRSA600-24 samples in tension mode at 298 K and 77 K.
There may be two reasons for this: (1) moderate grain refinement and relatively uniform microstructure, as shown in Fig. 1(b), could facilitate localized stress release and strain accommodation between grains, which is more or less helpful for the strength increment and plasticity retention [49]. (2) On the other hand, when the dislocation moves on the slip surface, it will be hindered and bent by the precipitates if it encounters hard and large precipitates. With the increase of the applied stress, the dislocation curvature of the hindered part increased. When the dislocation lines meet on both sides of the precipitated phase,
samples tested at 298 K and 77 K is shown in Fig. 5(a and b). Clearly, all samples exhibit ductile dimples, which is consistent with the good plasticity as shown in Fig. 3(d).
4. Discussion Seen from Fig. 3(a), after aging at 600 °C for 24 h, the yield tensile strength and ultimate tensile strength are significantly improved, but the elongation remains basically unchanged at room temperature. 5
Materials Science & Engineering A 767 (2019) 138424
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Fig. 7. TEM images of CRSA-600-24 samples after tensile deformation at 298 K (a, b, c) and 77 K (d, e, f).
arrows, respectively, indicating a transition from planar to wavy slip. This is because when the initial slip surfaces are blocked, new slip surface must be activated to accommodate the strain mismatch, so that dislocations are substantially entangled at multiple slip surfaces [53]. In addition, a small amount of deformation twins is detected in Fig. 7(c). These deformation twins may be caused by special grain orientations, which is common in TWIP steels that are easy to activation of twins [54,55]. This feature indicates that dislocation slip dominates the tensile deformation of CRSA-600-24 samples at room temperature. Similarly, from Figs. 7(d) and 7(e), it is also observed that dislocation lines and HDDWs appear in the tensile samples upon 77 K loading. However, different from 298 K loading, a large number of nanotwins are detected at 77 K and the width of nanotwins is about 10–20 nm, as shown in Fig. 7(f). According to Narita and Takamura [56], the critical resolved shear stress required for twinning is shown as:
positive and negative dislocations around the particles cancel each other and formed a dislocation ring around the particles. The remaining dislocation lines are continuing to move forward. This mechanism was known as “Orowan mechanism” [50]. Considering the size range and stiff structure of the B2 phase (see Figs. 1 and 2) in current alloys, the movement mechanism of dislocations should be mainly the bypassed mechanism (Orowan mechanism). Moreover, the shear stress required for dislocations to bypass the precipitation phase is shown as [51]:
τ = Gb/ λ
(1)
where G is the shear modulus, b is the burgers vector, λ is the spacing between two precipitates. As mentioned above, the volume fraction of the B2 phase is distinctly increased after aging treatment, which therefore yields the lower spacing between precipitates, and thus increased shear stress for dislocation movements. In fact, seen from Figs. 3(b) and 3(c), significant work-hardening ability, in terms of steady-state work-hardening-rate (WHR) values (over 1 GPa), dominates the overall plastic deformation, particularly for the aged alloy. Surprisingly, upon 77 K loading, the CRSA-600-24 HEA exhibits much stronger work-hardening ability than room-temperature loading, as shown in Fig. 6(a), which is true stress-strain curves obtained from Fig. 3(d). The corresponding WHR curves are presented in Fig. 6(b), in which the ultimate WHR value ahead of fracture reaches 2.5 GPa. Fig. 6(c) further displays the relationship between WHR and true stress, where the gray area denotes the necking zone of the specimen during plastic deformation, i.e. dσ/dε ≤ σ. It is noted that when stretched at room temperature, the strain value of necking onset is 0.43, and at 77 K, it reaches 0.467, indicating that the enhanced WHR at cryogenic temperature can delay necking onset to a higher strain value (Considerer's criterion) [52]. Moreover, the increased work hardening also pushes the flow stress to a higher value. At 298 K, the true stress reaches 1402 MPa and necking occurs, while at 77 K, the true stress delays until 2162 MPa. This feature again confirms that simultaneous enhancement between strength and ductility can be realized in the current alloys when subjected to 77 K loading. In order to further explore the underlying deformation mechanism of the CRSA-600-24 HEA after tension, TEM images from deformed samples at 298 K and 77 K are presented in Fig. 7. Fig. 7(a) shows a large number of parallel slip lines, which is typical of planar slip. Meanwhile, dislocation cells (DCs) and high-density dislocation walls (HDDWs) are also observed in Fig. 7(b), highlighted by white and red
τT = γSF /(kbS )
(2)
where τT is the critical resolved shear stress (CRSS), γSF is the SFE, k is a constant related to shear modulus, bS is the Burgers vector of the Shockley partial dislocation. It was reported that SFE decreased with decreasing temperature in the cubic materials and increased in the hexagonal materials [8]. Accordingly, it can be concluded that for fccstructured CRSA-600-24, twinning tends to be easily activated possibly due to the decreasing SFE at cryogenic temperature. Therefore, for 77 K loading, dislocation slip plus deformation twins together contribute to the plastic deformation process. It has been reported that the formation of twinning could reduce the dislocation mean free path, and thus produce the dynamic Hall–Petch effect [53,57,58]. Therefore, the activation of twin systems at low temperature plays an important role in the strength and ductility improvement. Additionally, the precipitation of B2 phase in CRSA-600-24 also yields multiple interactions between precipitates, dislocations, twins and grain boundaries. These features ultimately lead to the significant strengthening effect as described in Fig. 3(d). 5. Conclusion Al0.3CoCrFeNi HEAs with precipitated phases were prepared by arc melting, cold rolling, subsequent annealing, and aging treatments. Tensile properties of the alloys with two temperatures are investigated, and microstructures before and after deformation are analyzed by SEM 6
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and TEM. Based on this investigation, the following conclusions are drawn:
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(1) After 60% cold rolling and annealing at 1100 °C for 1 h, minor irregular B2 phases are precipitated in the annealed Al0.3CoCrFeNi HEAs. Subsequent aging at 600 °C for 24 h leads to more B2 phases in forms of lath-like particles in the aged counterparts. (2) Compared with the CRSA sample, the combination of fine grain strengthening and precipitation strengthening results in an increase in tensile strength and yield strength of the CRSA-600-24 sample, and a slight decrease in plasticity. Specifically, the yield strength, ultimate tensile strength and plasticity of CRSA-600-24 samples at 77 K increase from 460 MPa to 730 MPa, 913 MPa–1355 MPa, 62.5%–66.5%, respectively. (3) For the CRSA-600-24 HEA, dislocation slip dominates the tensile plastic deformation at 298 K, while at 77 K, the deformation mode is composed of slip plus nanotwins. Moreover, at 77 K, nanotwins induce an increase in work hardening, and the increased work hardening also delays the plastic instability to a higher stress/strain value. Acknowledgements The authors thank the National Natural Science Foundation of China (Nos. 51501123 and 11602158), the Youth Natural Science Foundation of Shanxi (No. 201601D021026), and 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 support. References [1] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.H. Tsau, S.Y. Chang, Adv. Eng. Mater. 6 (2004) 299–303. [2] D.B. Miracle, O.N. Senkov, Acta Mater. 122 (2017) 448–511. [3] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Prog. Mater. Sci. 61 (2014) 1–93. [4] A.S. Hamada, L.P. Karjalainen, M.C. Somani, R.M. Ramadan, Mater. Sci. 550 (2007) 217–222. [5] H. Zhang, Y.Z. He, Y. Pan, S. Guo, J. Alloy. Comp. 600 (2014) 210–214. [6] Y. Deng, C.C. Tasan, K.G. Pradeep, Acta Mater. 94 (2015) 124–133. [7] Z.Z. Li, S.T. Zhao, R.O. Ritchie, M.A. Meyers, Prog. Mater. Sci. 102 (2019) 296–345. [8] T. Ericsson, Acta Metall. 14 (1966) 853–865. [9] D.Y. Li, Y. Zhang, Intermetallics 70 (2016) 24–28. [10] J.Y. He, H. Wang, H.L. Huang, X.D. Xu, M.W. Chen, Y. Wu, X.J. Liu, T.G. Nieh, K. An, Z.P. Lu, Acta Mater. 102 (2016) 187–196. [11] Z. Li, S. Zhao, H. Diao, P.K. Liaw, M.A. Meyers, Sci. Rep. 7 (2017) 42742. [12] K. Chen, Y.T. Yang, G.J. Shao, K.J. Liu, Steel Res. Int. 82 (2011) 1325–1331. [13] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Prog. Mater. Sci. 61 (2014) 1–93. [14] J.X. Chen, L.L. Tan, X.M. Yu, I.P. Etim, M. I brahim, K. Yang, J. Mech. Behav. Biomed. 87 (2018) 68–79. [15] X.P. Luo, D.Q. Fang, Q.S. Li, Y.S. Chai, J. Rare Earths 34 (2016) 1134–1138.
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Superior tensile properties of Al0.3CoCrFeNi high entropy alloys with B2 precipitated phases at room and cryogenic temperatures
T
Q. Lia,b, T.W. Zhanga,b, J.W. Qiaoc, S.G. Maa,b,∗, D. Zhaoa,b, P. Lua,b, B. Xud, Z.H. Wanga,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 d School of Mechanics, Civil Engineering and Architecture, Northwestern Polytechnical University, Xian, 710129, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: High-entropy alloys Mechanical properties Cryogenic temperature Deformation twins Work hardening
Al0.3CoCrFeNi high-entropy alloys (HEAs) with B2 ordered phase mainly precipitated at grain boundaries are obtained by cold rolling, annealing and aging treatment. Tensile experiments show that both CRSA (annealed at 1100 °C for 1 h) and CRSA-600-24 (annealed at 1100 °C for 1 h and subsequently aging treatment at 600 °C for 24 h) samples have excellent combinations of yield strength, tensile strength and tensile elongation at room temperature, and more importantly, simultaneous enhancements in strength and ductility at liquid nitrogen (77 K) are attained for the latter condition. Analyses reveal that after aging treatment, the precipitation of B2 phase with ultrafine-grained or even nanoscale sizes, as well as moderate grain refinement, yields the significant increase in tensile strength and fascinating tensile plasticity. By transmission electron microscopy (TEM) characterization, it is found that the deformation mode dominated by dislocation glide at room temperature is transformed to that combined with dislocation glide plus nanoscale twinning at 77 K. Moreover, the twinninginduced work hardening renders the onset of necking instability to a higher strain/stress value, which therefore improves the strength and ductility simultaneously.
1. Introduction High-entropy alloys (HEAs) with multi-principal elements are generally composed of five or more elements with concentrations in the range of 5–35 at. % [1]. The HEAs tend to form simple solid solution phases, such as face-centered cubic (fcc), body-centered cubic (bcc) or hexagonal close-paced (hcp) structures, due to high mixing entropy reducing the Gibbs free energy and thus broadening the solid-solution limits especially at high temperatures. Moreover, extensive interests have been received in the HEAs field owing to their exceptional properties, such as high strength and good ductility, excellent resistances to oxidation, corrosion, fatigue and radiation, outstanding fracture toughness [2,3]. These properties qualify HEAs as potential candidates for the future engineering materials, to meet the demanding requirements for selected extreme applications, primarily in the automotive, defensive, marine and aerospace industries. It is known that brittlement is generally problematic when metals or
alloys are subjected to extreme conditions such as cryogenic temperatures or impact loadings. For instance, the low-temperature strength is obviously higher than that at room temperature, whereas the elongation shows an opposite trend in some austenitic high-Mn steels [4]. Amazingly, the recently-reported HEAs especially having low stackingfault energy (SFE) [5], effectively overcome the strength-ductility trade-off due to the nanotwinning formation during plastic deformation. This phenomenon is known as the “twinning-induced plasticity” effect (TWIP) [6]. Recently, Li et al. found that the incidence of deformation twins depended on the SFE and their frequency increased with the decrease in SFE [7]. Moreover, for cubic crystals, it is found that the SFE also decreases with decreasing the temperature [8]. Therefore, nanotwins may become more prevalent as cubic-structured and low-SFE HEAs are subjected to low-temperature loading. As a result, simultaneous improvement of strength and ductility could be realized due to the prominent TWIP effect. In fact, it has been reported that the fracture strength and tensile elongation of fcc-structured
∗ 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.msea.2019.138424 Received 26 June 2019; Received in revised form 14 September 2019; Accepted 14 September 2019 Available online 16 September 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.
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Fig. 1. SEM images of Al0.3CoCrFeNi HEAs: (a) CRSA sample, (b) CRSA-600-24 sample, and the insets are partially enlarged views of the (a, b) diagrams; (c) statistical results of the size distribution of B2 phase.
tests were performed with an Instron 5969 machine at a strain rate of 1 × 10−3/s and two temperatures of 298 K and 77 K. For the 77 K experiments, both the grips and tensile specimens were completely immersed in liquid nitrogen for about 20 min before the tensile tests, and the specimens were always submerging in the liquid nitrogen environment during tensile tests to ensure that the tensile specimens are held at the target temperature. Tensile tests at all temperatures were repeated at least three times. Microstructure evolution and phase constitution of the CRSA and CRSA-600-24 specimens before and after deformation were identified by scanning electron microscope (SEM, JEOM JSM-7100F) and transmission electron microscopy (TEM, JEM2010F). The chemical composition of the phase is analyzed by an attached energy dispersive spectrometer (EDS).
Al0.1CoCrFeNi HEAs respectively reached 635 MPa and 58.8% at 298 K, while at 77 K, the fracture strength increased to 1042 MPa, and what's more amazing was that the plasticity had not decreased, but increased to 81.6% [9]. In addition, He et al. [10] reported that the precipitation of secondary phase significantly improved the tensile strength in (CoCrFeNi)94Ti2Al4 HEA. In particular, the precipitation-hardened/toughed Al0.3CoCrFeNi HEAs have been widely reported, in which multiscalesized precipitates yield a good balance of strength and ductility [7]. Therefore, based on the above analyses, the fcc-structured Al0.3CoCrFeNi HEA with the precipitation of B2 ordered phases are obtained by cold rolling and subsequent aging, accompanied by room temperature and liquid nitrogen tensile tests. Finally, microstructure evolution, tensile properties, and deformation mechanisms at room and low temperatures are presented in results and discussion.
3. Results 2. Materials and methods
3.1. Initial microstructure analyses
Ingots with a nominal composition of Al0.3CoCrFeNi (the subscript is atomic ratio) were synthesized via arc-melting a mixture of highpurity elements (> 99.9 wt %) in a Ti-gettered argon atmosphere. The ingots were flipped and re-melted at least five times to ensure the homogeneity of chemical composition, and then suck into a watercooled copper mould to obtain plates with a dimension of 50 mm (length) × 10 mm (width) × 2 mm (depth). The actual chemical composition of the alloy plates is determined by EDS analysis to be 6.07Al17.52Co-18.02Cr-17.54Fe-17.16Ni (wt. %). The plates were cold rolled into a final thickness of 0.8 mm (~60% thickness reduction), annealed at 1100 °C for 1 h, and then water quenched (referred to as CRSA), and moreover, are further subjected to aging treatment at 600 °C for 24 h followed by water quenching (referred to as CRSA-600-24). Flat dog-bone-shaped tensile specimens were machined from alloy sheets by electric discharge machining with their longitudinal axes along with the rolling direction. The gage length and width of the tensile specimens are 10 mm and 4 mm, respectively. Uniaxial tensile
The initial microstructures of CRSA and CRSA-600-24 specimens are presented in Fig. 1. It is shown that both samples exhibit polycrystalline microstructures with irregular grains and visible annealed twins. Specifically, as shown in Fig. 1(c), the grain structure appears to be more uniform in CRSA-600-24 than CRSA, and the average grain size is about 20 μm. It is known that along with the recrystallization process during annealing (1100 °C for 1 h), dislocation recovery significantly occurs and strain energy substantially releases, which result in the minor driving force for the grain-boundary migration. Therefore, the speed of grain growth is very slow when samples are subjected to low-temperature aging treatment (600 °C for 24 h). In addition, recrystallization may also occur in the portion that has not been recrystallized during annealing, thus leading to the microstructure being relatively uniform. Interestingly, seen from Fig. 1(b, d), very fine or even nanoscale precipitates in these two samples are precipitated near or at the grain boundaries (as grain boundaries can be regarded as the diffusion path of atoms that effectively overcome the slow diffusion effect of HEAs 2
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Fig. 2. (a) A TEM BF image from the CRSA sample, (b) an SAED pattern obtained from the matrix of (a), (c) an SAED pattern obtained from the precipitation of (a), (d) a TEM BF image from the CRSA-600-24 sample, (e) an SAED pattern obtained from the matrix of (d), and (f) an SAED pattern from the precipitation of (d). (the inserted green tables in Figs. 2(b, d) are the EDS results of the precipitation). (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
24 h, with lower annealing temperature and prolonged annealing time, more Al and Ni atoms tend to be precipitated due to the enhanced phase instability, thereby forming B2 phase of the lath-like particles. In addition, recrystallized grain boundaries or large angle grain boundaries can accelerate the diffusion ability of atoms [12], which can effectively overcome the slow diffusion ability of HEAs [13]. Thus, the B2 phase tends to be precipitated at the grain boundaries.
[10,11]), and the amount of precipitates for the latter condition (~7% in volume fraction) is much higher than the former condition (~2%) due to the additional aging effect. Moreover, Fig. 1(e) depicts the size distribution of the precipitates for both samples. One can see that a wide size distribution ranging from ~50 to 450 nm is attained, in which over 50% of precipitates have the size of 50–100 nm. Next, TEM images of CRSA and CRSA-600-24 samples, as shown in Fig. 2, further manifest the feature of the precipitation as well as the fcc matrix for both samples, in terms of bright-field (BF) images with attached EDS analyses (Fig. 2(a, d)) and selected-area-electron diffraction (SAED) patterns (Fig. 2(b-c, e-f)). It is noted that the NiAl-type B2 (ordered bcc) precipitates, enriched in Al and Ni elements as compared to the fcc matrix, are mainly in forms of irregular or lath-like particles. This is because the present solid-solution alloy is in a saturated state, and Al and Ni elements have no time to nucleate and precipitate due to high-temperature melting and rapid solidification, thus forming supersaturated solid solutions with simple fcc structure. When the alloys are annealed at 1100 °C for 1 h, minor B2 phase is precipitated due to the very negative enthalpy and larger size difference of Al and Ni atoms as compared to other constitutive elements. After aging at 600 °C for
3.2. Mechanical properties Fig. 3(a) shows engineering tensile stress-strain curves of the CRSA and CRSA-600-24 samples at 298 K. It is found that for the aged alloy, the yield tensile strength and ultimate tensile strength increased from 330 MPa to 460 MPa and 746 MPa–913 MPa, respectively, as compared with the annealed counterpart. Although the plasticity decreased slightly, it still reached 62.5%. Moreover, Fig. 3(d) shows engineering tensile stress-strain curves of the CRSA-600-24 samples at 298 K and 77 K. For comparison, from 298 K to 77 K, the yield strength increased from 460 MPa to 730 MPa, nearly increased by 60%. Meanwhile, the ultimate tensile strength increased from 913 MPa to 1355 MPa, nearly 3
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Fig. 3. Engineering strain-stress curves (a) and true stress-strain curves (b), as well as work hardening rate-true strain curves (c) of CRSA and CRSA-600-24 samples in tension mode at 298 K; (d) engineering strain-stress curves of CRSA-600-24 samples in tension mode at 298 K and 77 K.
Fig. 4. The maps of tensile strength versus ductility (a) and yield strength versus product strength and elongation (b) of various advanced steels or alloys including CRSA-600-24 HEAs.
austenitic stainless steels [40,41], titanium alloys [42] and AlxCoCrFeNi HEAs [43–48]. It is seen that in comparison with other materials, the present alloys tend to be located at the top right corner of the maps and exhibit an excellent combination of strength and ductility particularly upon 77 K loading. In fact, the product of elongation and strength for CRSA-600-24 at 298 K and 77 K reach 57 GPa% and 89 GPa%, respectively. Moreover, the fracture-surface morphology from CRSA-600-24
increased by 50%. It is even amazing that the plasticity of the samples does not decrease but increases from 62.5% to 66.2%. Specifically, Fig. 4 compares the elongation versus strength and the strength-elongation product versus yield strength of CRSA-600-24 and other alloys, e.g., magnesium alloys [14–17], aluminum alloys [18–22], interstitial-free steels [23–27], transformation induced plasticity steels [28–32], dual-phase steels [33–37], ferritic stainless steels [38,39], 4
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Fig. 5. The dimpled fracture surface of CRSA-600-24 samples in tension mode at 298 K (a) and 77 K (b).
Fig. 6. True stress-true strain curves (a) and work hardening rate versus true strain curves (b) as well as work hardening rate versus true stress curves (c) of the CRSA600-24 samples in tension mode at 298 K and 77 K.
There may be two reasons for this: (1) moderate grain refinement and relatively uniform microstructure, as shown in Fig. 1(b), could facilitate localized stress release and strain accommodation between grains, which is more or less helpful for the strength increment and plasticity retention [49]. (2) On the other hand, when the dislocation moves on the slip surface, it will be hindered and bent by the precipitates if it encounters hard and large precipitates. With the increase of the applied stress, the dislocation curvature of the hindered part increased. When the dislocation lines meet on both sides of the precipitated phase,
samples tested at 298 K and 77 K is shown in Fig. 5(a and b). Clearly, all samples exhibit ductile dimples, which is consistent with the good plasticity as shown in Fig. 3(d).
4. Discussion Seen from Fig. 3(a), after aging at 600 °C for 24 h, the yield tensile strength and ultimate tensile strength are significantly improved, but the elongation remains basically unchanged at room temperature. 5
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Fig. 7. TEM images of CRSA-600-24 samples after tensile deformation at 298 K (a, b, c) and 77 K (d, e, f).
arrows, respectively, indicating a transition from planar to wavy slip. This is because when the initial slip surfaces are blocked, new slip surface must be activated to accommodate the strain mismatch, so that dislocations are substantially entangled at multiple slip surfaces [53]. In addition, a small amount of deformation twins is detected in Fig. 7(c). These deformation twins may be caused by special grain orientations, which is common in TWIP steels that are easy to activation of twins [54,55]. This feature indicates that dislocation slip dominates the tensile deformation of CRSA-600-24 samples at room temperature. Similarly, from Figs. 7(d) and 7(e), it is also observed that dislocation lines and HDDWs appear in the tensile samples upon 77 K loading. However, different from 298 K loading, a large number of nanotwins are detected at 77 K and the width of nanotwins is about 10–20 nm, as shown in Fig. 7(f). According to Narita and Takamura [56], the critical resolved shear stress required for twinning is shown as:
positive and negative dislocations around the particles cancel each other and formed a dislocation ring around the particles. The remaining dislocation lines are continuing to move forward. This mechanism was known as “Orowan mechanism” [50]. Considering the size range and stiff structure of the B2 phase (see Figs. 1 and 2) in current alloys, the movement mechanism of dislocations should be mainly the bypassed mechanism (Orowan mechanism). Moreover, the shear stress required for dislocations to bypass the precipitation phase is shown as [51]:
τ = Gb/ λ
(1)
where G is the shear modulus, b is the burgers vector, λ is the spacing between two precipitates. As mentioned above, the volume fraction of the B2 phase is distinctly increased after aging treatment, which therefore yields the lower spacing between precipitates, and thus increased shear stress for dislocation movements. In fact, seen from Figs. 3(b) and 3(c), significant work-hardening ability, in terms of steady-state work-hardening-rate (WHR) values (over 1 GPa), dominates the overall plastic deformation, particularly for the aged alloy. Surprisingly, upon 77 K loading, the CRSA-600-24 HEA exhibits much stronger work-hardening ability than room-temperature loading, as shown in Fig. 6(a), which is true stress-strain curves obtained from Fig. 3(d). The corresponding WHR curves are presented in Fig. 6(b), in which the ultimate WHR value ahead of fracture reaches 2.5 GPa. Fig. 6(c) further displays the relationship between WHR and true stress, where the gray area denotes the necking zone of the specimen during plastic deformation, i.e. dσ/dε ≤ σ. It is noted that when stretched at room temperature, the strain value of necking onset is 0.43, and at 77 K, it reaches 0.467, indicating that the enhanced WHR at cryogenic temperature can delay necking onset to a higher strain value (Considerer's criterion) [52]. Moreover, the increased work hardening also pushes the flow stress to a higher value. At 298 K, the true stress reaches 1402 MPa and necking occurs, while at 77 K, the true stress delays until 2162 MPa. This feature again confirms that simultaneous enhancement between strength and ductility can be realized in the current alloys when subjected to 77 K loading. In order to further explore the underlying deformation mechanism of the CRSA-600-24 HEA after tension, TEM images from deformed samples at 298 K and 77 K are presented in Fig. 7. Fig. 7(a) shows a large number of parallel slip lines, which is typical of planar slip. Meanwhile, dislocation cells (DCs) and high-density dislocation walls (HDDWs) are also observed in Fig. 7(b), highlighted by white and red
τT = γSF /(kbS )
(2)
where τT is the critical resolved shear stress (CRSS), γSF is the SFE, k is a constant related to shear modulus, bS is the Burgers vector of the Shockley partial dislocation. It was reported that SFE decreased with decreasing temperature in the cubic materials and increased in the hexagonal materials [8]. Accordingly, it can be concluded that for fccstructured CRSA-600-24, twinning tends to be easily activated possibly due to the decreasing SFE at cryogenic temperature. Therefore, for 77 K loading, dislocation slip plus deformation twins together contribute to the plastic deformation process. It has been reported that the formation of twinning could reduce the dislocation mean free path, and thus produce the dynamic Hall–Petch effect [53,57,58]. Therefore, the activation of twin systems at low temperature plays an important role in the strength and ductility improvement. Additionally, the precipitation of B2 phase in CRSA-600-24 also yields multiple interactions between precipitates, dislocations, twins and grain boundaries. These features ultimately lead to the significant strengthening effect as described in Fig. 3(d). 5. Conclusion Al0.3CoCrFeNi HEAs with precipitated phases were prepared by arc melting, cold rolling, subsequent annealing, and aging treatments. Tensile properties of the alloys with two temperatures are investigated, and microstructures before and after deformation are analyzed by SEM 6
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and TEM. Based on this investigation, the following conclusions are drawn:
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(1) After 60% cold rolling and annealing at 1100 °C for 1 h, minor irregular B2 phases are precipitated in the annealed Al0.3CoCrFeNi HEAs. Subsequent aging at 600 °C for 24 h leads to more B2 phases in forms of lath-like particles in the aged counterparts. (2) Compared with the CRSA sample, the combination of fine grain strengthening and precipitation strengthening results in an increase in tensile strength and yield strength of the CRSA-600-24 sample, and a slight decrease in plasticity. Specifically, the yield strength, ultimate tensile strength and plasticity of CRSA-600-24 samples at 77 K increase from 460 MPa to 730 MPa, 913 MPa–1355 MPa, 62.5%–66.5%, respectively. (3) For the CRSA-600-24 HEA, dislocation slip dominates the tensile plastic deformation at 298 K, while at 77 K, the deformation mode is composed of slip plus nanotwins. Moreover, at 77 K, nanotwins induce an increase in work hardening, and the increased work hardening also delays the plastic instability to a higher stress/strain value. Acknowledgements The authors thank the National Natural Science Foundation of China (Nos. 51501123 and 11602158), the Youth Natural Science Foundation of Shanxi (No. 201601D021026), and 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 support. References [1] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.H. Tsau, S.Y. Chang, Adv. Eng. Mater. 6 (2004) 299–303. [2] D.B. Miracle, O.N. Senkov, Acta Mater. 122 (2017) 448–511. [3] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Prog. Mater. Sci. 61 (2014) 1–93. [4] A.S. Hamada, L.P. Karjalainen, M.C. Somani, R.M. Ramadan, Mater. Sci. 550 (2007) 217–222. [5] H. Zhang, Y.Z. He, Y. Pan, S. Guo, J. Alloy. Comp. 600 (2014) 210–214. [6] Y. Deng, C.C. Tasan, K.G. Pradeep, Acta Mater. 94 (2015) 124–133. [7] Z.Z. Li, S.T. Zhao, R.O. Ritchie, M.A. Meyers, Prog. Mater. Sci. 102 (2019) 296–345. [8] T. Ericsson, Acta Metall. 14 (1966) 853–865. [9] D.Y. Li, Y. Zhang, Intermetallics 70 (2016) 24–28. [10] J.Y. He, H. Wang, H.L. Huang, X.D. Xu, M.W. Chen, Y. Wu, X.J. Liu, T.G. Nieh, K. An, Z.P. Lu, Acta Mater. 102 (2016) 187–196. [11] Z. Li, S. Zhao, H. Diao, P.K. Liaw, M.A. Meyers, Sci. Rep. 7 (2017) 42742. [12] K. Chen, Y.T. Yang, G.J. Shao, K.J. Liu, Steel Res. Int. 82 (2011) 1325–1331. [13] Y. Zhang, T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Prog. Mater. Sci. 61 (2014) 1–93. [14] J.X. Chen, L.L. Tan, X.M. Yu, I.P. Etim, M. I brahim, K. Yang, J. Mech. Behav. Biomed. 87 (2018) 68–79. [15] X.P. Luo, D.Q. Fang, Q.S. Li, Y.S. Chai, J. Rare Earths 34 (2016) 1134–1138.
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