- Email: [email protected]
Mechanical properties and deformation behavior of dual-phase Al0.6CoCrFeNi high-entropy alloys with heterogeneous structure at room and cryogenic temperatures
Journal of Alloys and Compounds 816 (2020) 152663
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Mechanical properties and deformation behavior of dual-phase Al0.6CoCrFeNi high-entropy alloys with heterogeneous structure at room and cryogenic temperatures Q. Li a, b, T.W. Zhang a, b, J.W. Qiao c, S.G. Ma a, b, *, D. Zhao a, b, P. Lu a, b, Z.H. Wang a, b, ** a b c
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 Institute of High-entropy Alloys, College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan, 030024, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 August 2019 Received in revised form 8 October 2019 Accepted 12 October 2019 Available online 16 October 2019
Dual-phase Al0.6CoCrFeNi high-entropy alloys (HEAs) with heterogeneous structure are obtained by cold rolling and annealing. Tensile experiments show that this alloy exhibits excellent combinations of yield strength, tensile strength and tensile elongation at room (298 K) and cryogenic (77 K) temperatures. Through the thermodynamic formula, the stacking fault energy of face-centered cube (fcc) phases in Al0.6CoCrFeNi alloys are calculated to be 49.33 mJ/m2 and 28.69 mJ/m2 at 298 K and 77 K, respectively. Combining with EBSD (electron backscattering diffraction) and TEM (transmission electron microscopy), it is concluded that the deformation of dual-phase Al0.6CoCrFeNi alloys with heterogeneous structure at 298 K is mainly dominated by the dislocation slip in fcc plus bcc (body-centered cube) phases, and the deformation mode at 77 K can be divided into two stages: (1) when the true strain is between 0 and 3.8%, the plasticity is mainly caused by dislocation slip of fcc and bcc phases; (2) when the true strain is between 3.8% and samples fracture, the plasticity is mainly induced by deformation twins in fcc phase and dislocation slip in bcc phase. © 2019 Published by Elsevier B.V.
Keywords: High-entropy alloys Heterogeneous structure Mechanical properties Cryogenic temperature Deformation twins
1. Introduction It is well known that conventional metallic alloys usually consist of one or two principal elements. In order to overcome this limitation, a new kind of alloys named multi-principal component alloys or high-entropy alloys (HEAs) has been proposed recently. They generally consist of five or more principal elements, and the concentration of each element ranges from 5 to 35 at. % [1].Due to their high mixing entropy of HEAs, it is easy to form solid solutions of face-centered cube (fcc), body-centered cube (bcc) or hexagonal close-paced (hcp) structures under as-cast conditions. In addition, HEAs have exhibited many potential properties, such as high hardness, high strength, thermal stability and excellent wear resistance [2]. Because of its excellent physical and chemical properties, HEAs have the potential to be used as structural
* 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 address: [email protected] (Z.H. Wang). https://doi.org/10.1016/j.jallcom.2019.152663 0925-8388/© 2019 Published by Elsevier B.V.
materials in military, aerospace, automotive and other industries. Generally, fcc-single phase HEAs exhibit a good ductility but a low strength, while bcc-single phase HEAs have good strength but poor plasticity. Nowadays, many scholars are based on single-phase alloys, in order to obtain HEAs with good plasticity and high strength, many measures are implemented. Such as, precipitationhardening [3], twinning-induced plasticity effect (TWIP) [4], and transformation-induced plasticity effect (TRIP) [5]. Although alloys with better comprehensive properties can be obtained by these strategies, it is important that the matrix materials have excellent comprehensive performance and the fcc plus bcc dual-phase HEAs conforms to this characteristic. For example, the yield strength, tensile strength and tensile elongation of the heat-treated duplephase Al0.5CoCrFeNi HEAs are 834 MPa, 1220 MPa and 25%, respectively. This excellent property is very rare in fcc or bcc singlephase alloys. In addition, researchers usually study alloys with fully recrystallized microstructures, that is, alloys with homogeneous structure. However, alloys with homogeneous structure generally have a good ductility and a low yield strength. Recently, some scholars have found that heterogeneous microstructure with recrystallization and non-recrystallization regions can be obtained
2
Q. Li et al. / Journal of Alloys and Compounds 816 (2020) 152663
by cold rolling, followed by partial recrystallization annealing. Nonrecrystallized grains can cause strain hardening, thereby can contribute to a high strength, while recrystallized grains can increase the ductility in materials with heterogeneous structure [6]. Therefore, alloys with heterogeneous structure tend to have high strength without sacrificing the ductility. At present, there are many studies on the low temperature mechanical properties of fcc or bcc single phase HEAs [7,8]. In their view, for cubic structure, a decrease in temperature leads to a decrease in the stacking fault energy (SFE) of the materials, thus activating more deformation twins to improve the plasticity of the alloys [9]. This theory is called the “twinning-induced plasticity” effect (TWIP) [4]. In the present study, the dual-phase Al0.6CoCrFeNi HEAs with heterogeneous structure was obtained by cold rolling 50%, and followed by an annealing at 1000 C for 1 h. The newly-developed alloys with a heterogeneous structure exhibit high yield and tensile strength, together with good uniform elongation and fracture elongation at room and low temperatures. Microstructure evolution, tensile properties, and deformation mechanisms of dual-phase Al0.6CoCrFeNi HEAs with heterogeneous structure at room and low temperatures are shown in the results and discussion. 2. Experimental Al0.6CoCrFeNi (the subscript is atomic ratio) HEAs was prepared by arc melting five high-purity metal elements (Al, Co, Cr, Fe, Ni elements purity larger than 99.9%) in a Ti-gettered high-purity argon atmosphere. Each alloy ingot was melted at least five times to ensure the homogeneity of chemical composition, and then the molten alloy was sucked into a copper mold cavity with a size of 50 mm (length) 10 mm (width) 2 mm (depth). A total of two alloy ingots were prepared. The obtained alloy ingots were cold rolled 50% along the length direction and final thickness of the samples were 1 mm. Then the cold rolled alloy sheets were annealed at 1000 C for 1 h, subsequently by water quenching. 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 gauge length and width of the tensile specimens were 10 mm and 4 mm, respectively. Quasi-static uniaxial tensile tests at 298 K and 77 K were conducted using an Instron 5969 machine at a strain rate of 5 103/s. In addition, for the accuracy of tensile test, at least three samples were tested at each temperature. The grips and tensile specimens need to be immersed in liquid nitrogen for 30 min before the tensile tests at 77 K for that the tensile specimens were held at the target temperature. The chemical compositions and structures of the phases were identified by energy-dispersive X-ray spectrometry (EDS) and X-ray diffraction (XRD, diffractometer using Cu-Ka radiation). Microstructure evolution before and after deformation were characterised by backscattered-electron (BSE), electron backscattering diffraction (EBSD) and transmission electron microscopy (TEM, JEM-2010F). 3. Results and discussion 3.1. Initial microstructures analyses The XRD patterns of Al0.6CoCrFeNi HEAs before and after cold rolling and annealing show the main presence of simple fcc plus bcc solid solutions, (see Fig. 1(a)). The lattice constants of the fcc and bcc phases in the as-cast sample were calculated by XRD to be 3.6038Å and 2.8725 Å, respectively, and the lattice constants of the fcc and bcc phases of the sample after cold rolling and annealing were 3.6006 Å and 2.8824 Å, respectively. The bright-field (BF) TEM
Fig. 1. XRD patterns of Al0.6CoCrFeNi HEAs before and after cold rolling and annealing (a); BF TEM image with the corresponding SAED patterns of Al0.6CoCrFeNi HEA after cold rolling and annealing (b).
images and selected-area-electron diffraction (SAED) patterns of Al0.6CoCrFeNi HEAs after cold rolling and annealing, as shown in Fig. 1(b), further manifest that Al0.6CoCrFeNi HEAs remains fcc plus bcc structure after cold rolling and annealing. In addition, EDS point analysis are performed on the two grains indicated by the white arrows in Fig. 1(b), and the results are shown in Table 1. It is found that the fcc phase is mainly enriched in Co, Cr and Fe, while the bcc phase is mainly enriched in Ni and Al. The BSE image with attached EDS analyses present microstructures after cold rolling and annealing, shown in Fig. 2. It is worth mentioning that, after the cold rolling of the alloy sheet, the internal stress of the alloy is not uniform, resulting in unequal storage energy, that is, the driving force for subsequent recrystallization formed inside the alloy is not equal. So, the HEAs possesses a complex heterogeneous structure with two types of microstructures (recrystallization microstructures and non-recrystallization microstructures), when annealed at the partial recrystallization temperature of 1000 C. In addition, EDS analysis shows that in the BSE image, the white matrix is fcc phase while the black regions are bcc phases. It is apparent that in the recrystallized region, the fcc and the bcc phases are separated, and in the non-recrystallized region, the two phases are entangled with each other. To further explore the alloys with heterogeneous structures, Fig. 3 presents the EBSD maps of the annealed samples.
Table 1 EDS chemical composition analysis in Fig. 1(b) (at. %). Phase
Al
Co
Cr
Fe
Ni
FCC BCC
5.23 32.41
24.21 16.55
29.64 7.50
25.09 10.53
15.82 33.00
Q. Li et al. / Journal of Alloys and Compounds 816 (2020) 152663
3
whereas those of the non-recrystallized regions are of 0.5 e4.5 . This difference indicates that the dislocation density in the nonrecrystallized regions is higher than that in the recrystallized regions. Therefore, the alloys also exhibit heterogeneity in dislocation density. According to the EBSD IPF and grain-size maps, recrystallized grains can be divided into medium-sized recrystallized grains (a2, 0.6~1 mm) far from the non-recrystallized regions and smaller grains (a3, ~0.6 mm) nucleated in the vicinity of the nonrecrystallized regions. The un-recrystallized grains retain the bulk grains (a1, 1~5 mm) formed by the distortion of the original coarse grains along the rolling direction after deformation. Overall, when annealed at the 1000 C, Al0.6CoCrFeNi HEAs possesses a heterogeneous microstructure with three types of grains (a1, a2, a3) at multiple length scales, and dislocation-density heterogeneity (see Fig. 3).
3.2. Mechanical properties
Fig. 2. BSE image of Al0.6CoCrFeNi HEA after cold-rolling and annealing; (b) BSE image of the red rectangle in (a); (ceg) the corresponding EDS elemental maps. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
The heterogeneous structure in alloys and dislocation structure formed by high-density dislocations in the non-recrystallized regions are clearly shown in the IPF map. Moreover, local misorientations across the non-recrystallized grains and recrystallized grains are observed, as evidenced in the EBSD kernel average misorientation (KAM) map shown in Fig. 3(b). According to the KAM map, two types of misorientations are clearly distinguished. The KAM values of the recrystallized regions are less than 0.5 ,
Fig. 4(a) shows representative engineering stress-strain curves of the uniaxial tensile test of Al0.6CoCrFeNi HEAs at 77 K and 298 K. The average yielding strength is determined to be about 786 MPa and 964 MPa, in addition, the ultimate tensile strength is determined to be about 1101 MPa and 1422 MPa for the alloys tested at 298 K and 77 K respectively. Compared with the elongation at room temperature, the elongation at 77 K were reduced slightly, but still reached 22.3%. Overally, Al0.6CoCrFeNi HEAs with heterogeneous structure have high strength and good plasticity at both 77 K and 298 K. This is mainly attributed to the high-density dislocations in the non-recrystallized regions and the ultrafine grains in the recrystallized region. Ultrafine grains in recrystallized regions have the potential of plastic deformation to provide plasticity for the alloys, while high dislocation density in the non-recrystallized regions can improve work-hardening ability of the materials and hence the strength of the alloys. According to Fig. 4(a), the corresponding true stress-strain curves are obtained, as shown in Fig. 4(b), and then the work hardening rate-true strain curves are also obtained according to the true stress-strain curves, which are shown in Fig. 4(c). The newly developed alloys show a very high work hardening behavior, especially at 77 K, the work hardening rate of the alloy still reaches 3.4 GPa before the sample is necked. Therefore, compared to 298 K, the alloys have a higher work hardening behavior at 77 K, so that the sample can reach a higher stress value before necking, thereby increasing the strength of the alloys. In addition, we compare tensile properties of Al0.6CoCrFeNi HEAs with those of reported fcc/L12 þ bcc/B2 dual-phase HEAs and some advanced traditional duplex stainless steels in Fig. 5 [10e18].
Fig. 3. EBSD results of Al0.6CoCrFeNi HEA after cold-rolling and annealing: (a) IPF map overlaid with IQ map; (b) KAM map; (c) the distributions of grain size.
4
Q. Li et al. / Journal of Alloys and Compounds 816 (2020) 152663
Fig. 4. (a) Engineering tensile strain-stress, (b) true tensile stress-strain and (c) work hardening rate versus true train curves of Al0.6CoCrFeNi HEAs in tension mode at 289 K and 77 K.
It is clear that in terms of product of strength and ductility and yield strength, Al0.6CoCrFeNi HEAs exhibit a better combination over existing dual-phase HEAs and traditional duplex stainless steels. 3.3. Deformation mechanism In order to characterize the underlying deformation mechanisms of Al0.6CoCrFeNi HEAs with heterogeneous structure under uniaxial tension at 298 K and 77 K. It is necessary to reveal the microstructural evolution during plastic deformation. Fig. 6(a and b) shows BF TEM images for the samples deformed at 298 K and 77 K. Other literatures have reported that the SFE of Al0.6CoCrFeNi HEAs is 147.4e162.0 mJ/m2 [19]. In addition, the SFE not only has a strong influence on the deformation mechanism, but it also controls the slip mode of the dislocations. A high SFE usually promotes cross-slip of dislocations [20]. So, a large amount of crossdislocation slip can be observed in samples that are deformed at room temperature. From the engineering stress-strain curve (Fig. 4(a)) at room temperature, it can be seen that the sample has a relatively high work hardening ability, while cross-slip of dislocations has little effect on the work hardening ability [20]. In addition, new slip planes will be required to accommodate the excess strain when initial slip planes are blocked, thus yielding the formation of high-dense dislocation walls (DWs) and small-sized dislocation cells (DCs) [21]. The high-dense DWs and small-sized DCs increase the resistance to dislocation movement, and hence enhance the work hardening ability of the material. This is to say; dislocation slip dominated the deformation of the sample at room
temperature. Some high-density dislocation structures are also observed in Fig. 6(b), but the difference is that a large number of deformation twins are activated at 77 K, and the average twin size is 10 nm. To further explore the microstructure evolution of Al0.6CoCrFeNi HEAs at 77 K, EBSD patterns after tensile deformation at 77 K are shown in Fig. 7. Obviously, a large number of deformed twins are generated in the tensile deformation at 77 K (highlighted by yellow lines in Fig. 7(a)), in addition, there are some deformation twins mainly concentrated in the fcc phase. According to Narita and Takamura [22], the critical resolved shear stress required for twinning is shown as:
sT ¼ 2gSFE
mbp
(1)
where sT is the critical stress for twinning formation, gSFE is the stacking fault energy, m is the average Schmid factor of about 0.33 and bp is the Burgers vector of the partial dislocations and bp ¼ 0.147 nm [6]. In addition, according to the thermodynamic analysis, without considering the interaction between the fcc and bcc phases in the Al0.6CoCrFeNi HEAs, the ideal gSFE of fcc phases in the alloys is then expressed by Ref. [23].
gSFE ¼ 2rDGg/εþ2s
(2)
where r is the molar surface density along {111} planes, DGg/ε is the molar Gibbs energy of the austenite to ε-martensite phase transformation gfcc /εMs and s is the interfacial energy per unit hcp area of the phase boundary. The molar surface density r is
Fig. 5. Yield strength versus product of strength and elongation of various traditional duplex stainless steels and some FCC/L12 þ BCC/B2 dual-phase HEAs including Al0.6CoCrFeNi HEAs.
Q. Li et al. / Journal of Alloys and Compounds 816 (2020) 152663
5
Fig. 6. TEM images of Al0.6CoCrFeNi HEA after tensile deformation at 298 K (a) and 77 K (b, c, d).
Fig. 7. EBSD results of Al0.6CoCrFeNi HEA after tensile deformation at 77 K: (a) Phase map (twins are highlighted by yellow lines), (b) misorientation angle. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
g/ε
calculated by Allain et al. as follows [24]:
4
1 3 a2 N
r ¼ pffiffiffi
(3)
where a is the lattice parameter of the alloy and N is Avogadro’s number. Considering a regular solution model, DGg/ε can be expressed for the present Al0.6CoCrFeNi alloys as follow [25e27]:
DGg/ε ¼
X 1X g/ε g/ε Xi DGi þ Xi Xj Uij 2 i ij
(4)
ði; j ¼ Al; Co; Cr; Fe; Ni; i s jÞ g/ε
where Xi and DG i represent the molar fraction and the differg/ε ence of free energy between fcc and hcp of pure metals. Uij is an interaction energy parameter for components i and j. Because of the content of Al in the fcc phases of the alloys are low, the interaction energy parameter for the Al and other elements can be ignored. Accordingly values of the parameters in the above formula are
g/ε
g/ε
g/ε
shown in Table 2. In addition, DGCo , UCoCr , UNiCo , UFeCo have not been reported in other literatures, so in this paper, according to equations (1) and (4) and the stacking fault energy of similar alloys that have been reported in other literatures, the following results
Table 2 Numerical values and functions used for the calculations. Parameters
Functions used (units)
g/ε DGFe g/ε DGCr g/ε DGNi g/ε DGAl g/ε UFeCr g/ε UFeNi g/ε UFeAl g/ε UCrNi
2243.38 þ 4.309T (J mol1)
[36]
1370e0.163T (J mol1)
[37]
1046 þ 1.255T (J mol1)
[36]
4190 (J mol1)
[35]
a
3.6006 Å 0.008 (J m2)
This work [36]
s
Reference
2800þ5T (J mol1)
[38]
2095 (J mol1)
[39]
2095 (J mol1)
[38]
3328 (J mol1)
[38]
6
Q. Li et al. / Journal of Alloys and Compounds 816 (2020) 152663
are obtained:
DGgCo/ε ¼ 3384:8 þ 0:35T Jmol1
/ε UgCoCr ¼ 2158:24 Jmol1
/ε UgNiCo ¼ 6777:24 Jmol1
/ε UgCoFe ¼ 6562:64 Jmol1
(5)
(6)
4. Conclusion
(7)
(8)
Substituting the obtained data into formulas (2), (3) and (4), the relationship between the stacking fault energy of the alloys and temperature is finally determined as follows:
gSFE ¼ 21:5 þ 0:0943T
(9)
Then, 77 K and 298 K are substituted into formula (9), and the stacking fault energy of the alloys at two temperatures is determined as follows:
.
gSFE ¼ 28:69 mJ m2 77K .
gSFE ¼ 49:33 mJ m2 298K
This is due to the “dynamic HallePetch” effect, which promotes the work hardening ability of Al0.6CoCrFeNi HEAs by activating the twin system at 77 K. Moreover, according to considerer’s criterion [35], the twinning-induced work hardening ability increases the flow stress of the necking of the sample to a higher stress value. This results in a higher strength of the alloy at 77 K compared to 298 K.
(10)
In conclusion, dual-phase Al0.6CoCrFeNi HEAs with heterogeneous structure are successfully prepared by arc melting, cold rolling, and partial recrystallization annealing. Tensile tests reveal that the present alloys exhibit excellent combination of strength and ductility for both 298 K and 77 K loading. Specifically, at 77 K, the yield strength and tensile strength reach 964 MPa and 1422 MPa, respectively, while the plasticity has a slight decrease. At 298 K, dislocation slip dominates the overall plastic deformation, while at 77 K deformation twins prevail (in addition to dislocations) when the true train is over 3.8% based on the critical stress formula. According to considerer’s criterion, the twinning-induced work hardening renders the onset of necking instability to a higher stress value, which results in the higher tensile strength at low temperature. Declaration of competing interest
(11)
Substituting formulas (10) and (11) into formula (1), the critical stress for twinning formation of the fcc region at 298 K and 77 K for the Al0.6CoCrFeNi HEAs are 2034 MPa and 1183 MPa respectively. It can be seen from Fig. 4(b) that the maximum tensile true stress of the sample is 1427 MPa at 298 K, which is much smaller than the critical stress for twinning formation. At 77 K, the maximum tensile true stress of the sample is 1740 MPa, which has exceeded the critical stress of twinning formation. This is to say, in the tensile deformation at 77 K, when the true strain of the sample reaches 3.8%, the formation of twins has begun. So, the generation of twins is not seen in Fig. 6(a), while a large number of twins can be seen in Fig. 6(b). Therefore, in the later stage of deformation at 77 K, the plastic deformation in the fcc phase are mainly induced by deformation twins. In addition, Table 1 can be known that the bcc phase are rich in a large amount of Al. In other literatures, it has been reported that Al has a high SFE of about 86 mJ/m2 [28,29]. Other literatures had also mentioned that the average stacking fault energy of Al0.6CoCrFeNi HEA is about 150 mJ/m2 [30], and we have confirmed that fcc phase have lower SFE, therefore, it can be inferred that the SEF of bcc phase is much larger than that of fcc phase. This is why no deformation twins are observed in the bcc phase. Although no deformation twins were observed in the bcc phase during the later deformation process, but dislocation slip must be generated to coordinate the plasticity increase caused by the deformation twins in the fcc phase. Therefore, the plastic deformation of Al0.6CoCrFeNi HEAs at 77 K can be divided into two stages: (1) when the true strain is between 0 and 3.8%, the plasticity is mainly caused by dislocation slip of fcc and bcc phases; (2) when the true strain is between 3.8% and sample fracture, the plasticity is mainly induced by deformation twins in fcc phase and dislocation slip in bcc phase. In addition, it has been reported in other literatures that the formation of deformation twins can induce the increase of work hardening [31,32]. According to “dynamic HallePetch” effect, the formation of deformation twin boundaries can also led to grain refinement, which reduced the dislocation mean free path and increased the strength of alloys [33,34]. In Fig. 4(c), the work hardening ability at 77 K is significantly higher than that at 298 K.
The manuscript submitted was the result of my independent research. Except as noted in the text, this article does not contain research results that have been published or written by other individuals or groups. The legal responsibility of this statement is borne by me. 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, Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes, Adv. Eng. Mater. 5 (2004) 299e304. [2] D.B. Miracle, O.N. Senkov, A critical review of high entropy alloys and related concepts, Acta Mater. 122 (2017) 448e511. [3] Y.J. He, H. Wang, H.L. Huang, A precipitation-hardened high-entropy alloy with outstanding tensile properties, Acta Mater. 102 (2016) 187e196. [4] Y. Deng, C.C. Tasan, K.G. Pradeep, H. Springer, A. Kostka, D. Raabe, Design of a twinning-induced plasticity high entropy alloy, Acta Mater. 94 (2015) 124e133. [5] Z.M. Li, K.G. Pradeep, Y. Deng, Metastable high-entropy dual-phase alloys overcome the strength-ductility trade-off, Nature 534 (2016) 227e230. [6] S.W. Wu, G. Wang, Q. Wang, Y.D. Jia, J. Yi, Q.J. Zhai, Enhancement of strengthductility trade-off in a high-entropy alloy through a heterogeneous structure, Acta Mater. 165 (2019) 444e458. [7] J.W. Qiao, S.G. Ma, E.W. Huang, C.P. Chuang, P.K. Liaw, Y. Zhang, Microstructural characteristics and mechanical behaviors of AlCoCrFeNi high-entropy alloys at ambient and cryogenic temperatures, Mater. Sci. Forum 688 (2011) 419e425. [8] B. Gludovatz, A. Hohenwarter, D. Catoor, E.H. Chang, E.P. George, R.O. Ritchie, A fracture-resistant high-entropy alloy for cryogenic applications, Science 345 (2014) 1153e1158.
Q. Li et al. / Journal of Alloys and Compounds 816 (2020) 152663 [9] T. Ericsson, The temperature and concentration dependence of the stacking fault energy in the Co-Ni system, Acta Metall. 14 (1966) 853e865. [10] K. Chandra, R. Singhal, V. Kain, V.S. Raja, Low temperature embrittlement of duplex stainless steel: correlation between mechanical and electrochemical behavior, Mater. Sci. Eng. A 527 (2010) 3904e3912. , Behavior of duplex stainless-steel casting [11] S. Jayet-Gendrot, P. Gilles, C. Migne defects under mechanical loadings, Nucl. Eng. Des. 197 (2000) 141e153. [12] E. Medina, J.M. Medina, A. Cobo, D.M. Bastidas, Evaluation of mechanical and structural behavior of austenitic and duplex stainless-steel reinforcements, Constr. Build. Mater. 78 (2015) 1e7. [13] M.K. Mishra, G. Gunasekaran, A.G. Rao, B.P. Kashyap, N. Prabhu, Effect of multipass friction stir processing on mechanical and corrosion behavior of 2507 super duplex stainless steel, Constr. Build. Mater. 26 (2017) 849e860. [14] M. Okayasu, D. Ishida, Effect of microstructural characteristics on mechanical properties of austenitic, ferritic, and g-a duplex stainless steels, Metall. Mater. Trans. A 50 (2019) 1380e1388. [15] C.S. Marchi, J. Zelinski, X. Tang, G.H. Schiroky, Mechanical properties of super duplex stainless steel 2507 after gas phase thermal precharging with hydrogen, Metall. Mater. Trans. A 38 (2007) 2763e2775. [16] L. Chen, F.P. Yuan, P. Jiang, X.L. Wu, Mechanical properties and nanostructures in a duplex stainless steel subjected to equal channel angular pressing, Mater. Sci. Eng. A 551 (2012) 154e159. [17] S. Niu, H. Kou, T. Guo, Y. Zhang, J. Wang, J. Li, Strengthening of nanoprecipitations in an annealed Al0.5CoCrFeNi high entropy alloy, Mater. Sci. Eng. A 671 (2016) 82e86. [18] Y.J. Liang, L.J. Wang, Y.R. Wen, B.Y. Cheng, Q.L. Wu, T.Q. Cao, Qian Xiao, Y.Y. Xue, G. Sha, Y.D. Wang, Y. Ren, X.Y. Li, L. Wang, F.C. Wang, H.N. Cai, Highcontent ductile coherent nanoprecipitates achieve ultrastrong high-entropy alloys, Nat. Commun. 9 (2018) 1e8. [19] T.S. Cao, J.L. Shang, J. Zhao, C.Q. Cheng, R. Wang, H. Wang, The influence of Al elements on the structure and the creep behavior of AlxCoCrFeNi high entropy alloys, Mater. Lett. 164 (2016) 344e347. [20] E. Welsch, D. Ponge, S.M. Hafez Haghighat, S. Sandlobes, P. Choi, M. Herbig, S. Zaefferer, D. Raabe, Strain hardening by dynamic slip band refinement in a high-Mn lightweight steel, Acta Mater. 116 (2016) 188e199. [21] F. Otto, A. Dlouhy, Ch Somsen, H. Bei, G. Eggeler, E.P. George, The influences of temperature and microstructure on the tensile properties of a CoCrFeMnNi high-entropy alloy, Acta Mater. 61 (2013) 5743e5755. [22] T.S. Byun, On the stress dependence of partial dislocation separation and deformation microstructure in austenitic stainless steels, Acta Mater. 51 (2003) 3063e3071. [23] G.B. Olson, M. Cohen, A General mechanism of martensitic nucleation: part III. kinetics of martensitic nucleation, Metall. Mater. Trans. A 7 (1976) 1915e1923.
7
€nninen, Thermodynamic [24] S. Curtze, V.-T. Kuokkala, A. Oikari, J. Talonen, H. Ha modeling of the stacking fault energy of austenitic steels, Acta Mater. 59 (2011) 1068e1076. [25] W.S. Yang, C.M. Wan, The influence of aluminium content to the stacking fault energy in Fe-Mn-Al-C alloy system, J. Mater. Sci. 25 (1990) 1821e1823. [26] Y.Q. Wang, B. Liu, K. Yan, M.S. Wang, S. Kabra, Y.L. Chiu, D. Dye, P.D. Lee, Y. Liu, B. Cai, Probing deformation mechanisms of a FeCoCrNi high-entropy alloy at 293 and 77K using in situ neutron diffraction, Acta Mater. 154 (2018) 79e89. [27] Y.H. Zhang, Y. Zhuang, A. Hu, J.J. Kai, C.T. Liu, The origin of negative stacking fault energies and nano-twin formation in face-centered cubic high entropy alloys, Scr. Mater. 130 (2017) 96e99. [28] B. Hammer, K.W. Jacobsen, V. Milman, M.C. Payne, Stacking fault energies in aluminium, J. Phys. Condens. Matter 4 (1992) 10453e10460. [29] L.L. Ma, L. Wang, Z.H. Nie, F.C. Wang, Y.F. Xue, J.L. Zhou, T.Q. Cao, Y.D. Wang, Y. Ren, Reversible deformation-induced martensitic transformation in Al0.6CoCrFeNi high-entropy alloy investigated by in situ synchrotron-based high-energy X-ray diffraction, Acta Mater. 128 (2017) 12e21. [30] T.S. Cao, J.L. Shang, J. Zhao, C.Q. Cheng, R. Wang, H. Wang, The influence of Al elements on the structure and the creep behavior of AlxCoCrFeNi high entropy alloys, Mater. Lett. 164 (2016) 344e347. [31] H. Beladi, I.B. Timokhina, Y. Estrin, J. Kim, B.C. De Cooman, S.K. Kim, Orientation dependence of twinning and strain hardening behaviour of a high manganese twinning induced plasticity steel with polycrystalline structure, Acta Mater. 59 (2011) 7787e7799. [32] I. Gutierrez-Urrutia, D. Raabe, Dislocation and twin substructure evolution during strain hardening of an Fee22 wt.% Mne0.6 wt.% C TWIP steel observed by electron channeling contrast imaging, Acta Mater. 59 (2011) 6449e6462. [33] O. Bouaziz, N. Guelton, Modelling of TWIP effect on work-hardening, Mater. Sci. Eng. A 319 (2001) 246e249. , [34] I. Karaman, H. Sehitoglu, A.J. Beaudoin, Y.I. Chumlyakov, H.J. Maier, C.N. Tome Modeling the deformation behavior of Hadfield steel single and polycrystals due to twinning and slip, Acta Mater. 48 (2000) 2031e2047. [35] A. Gali, E.P. George, Tensile properties of high- and medium-entropy alloys, Intermetallics 39 (2013) 74e78. [36] K. Ishida, Direct estimation of stacking fault energy by thermodynamic analysis, Phys. Status Solidi A 36 (1979) 717e728. [37] G. Inden, The role of magnetism in the calculation of phase diagrams, Phys. BþC 103 (1981) 82e100. [38] Y.H. Zhang, Y. Zhuang, A. Hu, J.J. Kai, C.T. Liu, The origin of negative stacking fault energies and nano-twin formation in face-centered cubic high entropy alloys, Scr. Mater. 130 (2017) 96e99. [39] I.A. Yakubtsov, A. Ariapour, D.D. Perovic, Effect of nitrogen on stacking fault energy of fcc iron-based alloys, Acta Mater. 47 (1999) 1271e1279.
Contents lists available at ScienceDirect
Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Mechanical properties and deformation behavior of dual-phase Al0.6CoCrFeNi high-entropy alloys with heterogeneous structure at room and cryogenic temperatures Q. Li a, b, T.W. Zhang a, b, J.W. Qiao c, S.G. Ma a, b, *, D. Zhao a, b, P. Lu a, b, Z.H. Wang a, b, ** a b c
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 Institute of High-entropy Alloys, College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan, 030024, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 17 August 2019 Received in revised form 8 October 2019 Accepted 12 October 2019 Available online 16 October 2019
Dual-phase Al0.6CoCrFeNi high-entropy alloys (HEAs) with heterogeneous structure are obtained by cold rolling and annealing. Tensile experiments show that this alloy exhibits excellent combinations of yield strength, tensile strength and tensile elongation at room (298 K) and cryogenic (77 K) temperatures. Through the thermodynamic formula, the stacking fault energy of face-centered cube (fcc) phases in Al0.6CoCrFeNi alloys are calculated to be 49.33 mJ/m2 and 28.69 mJ/m2 at 298 K and 77 K, respectively. Combining with EBSD (electron backscattering diffraction) and TEM (transmission electron microscopy), it is concluded that the deformation of dual-phase Al0.6CoCrFeNi alloys with heterogeneous structure at 298 K is mainly dominated by the dislocation slip in fcc plus bcc (body-centered cube) phases, and the deformation mode at 77 K can be divided into two stages: (1) when the true strain is between 0 and 3.8%, the plasticity is mainly caused by dislocation slip of fcc and bcc phases; (2) when the true strain is between 3.8% and samples fracture, the plasticity is mainly induced by deformation twins in fcc phase and dislocation slip in bcc phase. © 2019 Published by Elsevier B.V.
Keywords: High-entropy alloys Heterogeneous structure Mechanical properties Cryogenic temperature Deformation twins
1. Introduction It is well known that conventional metallic alloys usually consist of one or two principal elements. In order to overcome this limitation, a new kind of alloys named multi-principal component alloys or high-entropy alloys (HEAs) has been proposed recently. They generally consist of five or more principal elements, and the concentration of each element ranges from 5 to 35 at. % [1].Due to their high mixing entropy of HEAs, it is easy to form solid solutions of face-centered cube (fcc), body-centered cube (bcc) or hexagonal close-paced (hcp) structures under as-cast conditions. In addition, HEAs have exhibited many potential properties, such as high hardness, high strength, thermal stability and excellent wear resistance [2]. Because of its excellent physical and chemical properties, HEAs have the potential to be used as structural
* 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 address: [email protected] (Z.H. Wang). https://doi.org/10.1016/j.jallcom.2019.152663 0925-8388/© 2019 Published by Elsevier B.V.
materials in military, aerospace, automotive and other industries. Generally, fcc-single phase HEAs exhibit a good ductility but a low strength, while bcc-single phase HEAs have good strength but poor plasticity. Nowadays, many scholars are based on single-phase alloys, in order to obtain HEAs with good plasticity and high strength, many measures are implemented. Such as, precipitationhardening [3], twinning-induced plasticity effect (TWIP) [4], and transformation-induced plasticity effect (TRIP) [5]. Although alloys with better comprehensive properties can be obtained by these strategies, it is important that the matrix materials have excellent comprehensive performance and the fcc plus bcc dual-phase HEAs conforms to this characteristic. For example, the yield strength, tensile strength and tensile elongation of the heat-treated duplephase Al0.5CoCrFeNi HEAs are 834 MPa, 1220 MPa and 25%, respectively. This excellent property is very rare in fcc or bcc singlephase alloys. In addition, researchers usually study alloys with fully recrystallized microstructures, that is, alloys with homogeneous structure. However, alloys with homogeneous structure generally have a good ductility and a low yield strength. Recently, some scholars have found that heterogeneous microstructure with recrystallization and non-recrystallization regions can be obtained
2
Q. Li et al. / Journal of Alloys and Compounds 816 (2020) 152663
by cold rolling, followed by partial recrystallization annealing. Nonrecrystallized grains can cause strain hardening, thereby can contribute to a high strength, while recrystallized grains can increase the ductility in materials with heterogeneous structure [6]. Therefore, alloys with heterogeneous structure tend to have high strength without sacrificing the ductility. At present, there are many studies on the low temperature mechanical properties of fcc or bcc single phase HEAs [7,8]. In their view, for cubic structure, a decrease in temperature leads to a decrease in the stacking fault energy (SFE) of the materials, thus activating more deformation twins to improve the plasticity of the alloys [9]. This theory is called the “twinning-induced plasticity” effect (TWIP) [4]. In the present study, the dual-phase Al0.6CoCrFeNi HEAs with heterogeneous structure was obtained by cold rolling 50%, and followed by an annealing at 1000 C for 1 h. The newly-developed alloys with a heterogeneous structure exhibit high yield and tensile strength, together with good uniform elongation and fracture elongation at room and low temperatures. Microstructure evolution, tensile properties, and deformation mechanisms of dual-phase Al0.6CoCrFeNi HEAs with heterogeneous structure at room and low temperatures are shown in the results and discussion. 2. Experimental Al0.6CoCrFeNi (the subscript is atomic ratio) HEAs was prepared by arc melting five high-purity metal elements (Al, Co, Cr, Fe, Ni elements purity larger than 99.9%) in a Ti-gettered high-purity argon atmosphere. Each alloy ingot was melted at least five times to ensure the homogeneity of chemical composition, and then the molten alloy was sucked into a copper mold cavity with a size of 50 mm (length) 10 mm (width) 2 mm (depth). A total of two alloy ingots were prepared. The obtained alloy ingots were cold rolled 50% along the length direction and final thickness of the samples were 1 mm. Then the cold rolled alloy sheets were annealed at 1000 C for 1 h, subsequently by water quenching. 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 gauge length and width of the tensile specimens were 10 mm and 4 mm, respectively. Quasi-static uniaxial tensile tests at 298 K and 77 K were conducted using an Instron 5969 machine at a strain rate of 5 103/s. In addition, for the accuracy of tensile test, at least three samples were tested at each temperature. The grips and tensile specimens need to be immersed in liquid nitrogen for 30 min before the tensile tests at 77 K for that the tensile specimens were held at the target temperature. The chemical compositions and structures of the phases were identified by energy-dispersive X-ray spectrometry (EDS) and X-ray diffraction (XRD, diffractometer using Cu-Ka radiation). Microstructure evolution before and after deformation were characterised by backscattered-electron (BSE), electron backscattering diffraction (EBSD) and transmission electron microscopy (TEM, JEM-2010F). 3. Results and discussion 3.1. Initial microstructures analyses The XRD patterns of Al0.6CoCrFeNi HEAs before and after cold rolling and annealing show the main presence of simple fcc plus bcc solid solutions, (see Fig. 1(a)). The lattice constants of the fcc and bcc phases in the as-cast sample were calculated by XRD to be 3.6038Å and 2.8725 Å, respectively, and the lattice constants of the fcc and bcc phases of the sample after cold rolling and annealing were 3.6006 Å and 2.8824 Å, respectively. The bright-field (BF) TEM
Fig. 1. XRD patterns of Al0.6CoCrFeNi HEAs before and after cold rolling and annealing (a); BF TEM image with the corresponding SAED patterns of Al0.6CoCrFeNi HEA after cold rolling and annealing (b).
images and selected-area-electron diffraction (SAED) patterns of Al0.6CoCrFeNi HEAs after cold rolling and annealing, as shown in Fig. 1(b), further manifest that Al0.6CoCrFeNi HEAs remains fcc plus bcc structure after cold rolling and annealing. In addition, EDS point analysis are performed on the two grains indicated by the white arrows in Fig. 1(b), and the results are shown in Table 1. It is found that the fcc phase is mainly enriched in Co, Cr and Fe, while the bcc phase is mainly enriched in Ni and Al. The BSE image with attached EDS analyses present microstructures after cold rolling and annealing, shown in Fig. 2. It is worth mentioning that, after the cold rolling of the alloy sheet, the internal stress of the alloy is not uniform, resulting in unequal storage energy, that is, the driving force for subsequent recrystallization formed inside the alloy is not equal. So, the HEAs possesses a complex heterogeneous structure with two types of microstructures (recrystallization microstructures and non-recrystallization microstructures), when annealed at the partial recrystallization temperature of 1000 C. In addition, EDS analysis shows that in the BSE image, the white matrix is fcc phase while the black regions are bcc phases. It is apparent that in the recrystallized region, the fcc and the bcc phases are separated, and in the non-recrystallized region, the two phases are entangled with each other. To further explore the alloys with heterogeneous structures, Fig. 3 presents the EBSD maps of the annealed samples.
Table 1 EDS chemical composition analysis in Fig. 1(b) (at. %). Phase
Al
Co
Cr
Fe
Ni
FCC BCC
5.23 32.41
24.21 16.55
29.64 7.50
25.09 10.53
15.82 33.00
Q. Li et al. / Journal of Alloys and Compounds 816 (2020) 152663
3
whereas those of the non-recrystallized regions are of 0.5 e4.5 . This difference indicates that the dislocation density in the nonrecrystallized regions is higher than that in the recrystallized regions. Therefore, the alloys also exhibit heterogeneity in dislocation density. According to the EBSD IPF and grain-size maps, recrystallized grains can be divided into medium-sized recrystallized grains (a2, 0.6~1 mm) far from the non-recrystallized regions and smaller grains (a3, ~0.6 mm) nucleated in the vicinity of the nonrecrystallized regions. The un-recrystallized grains retain the bulk grains (a1, 1~5 mm) formed by the distortion of the original coarse grains along the rolling direction after deformation. Overall, when annealed at the 1000 C, Al0.6CoCrFeNi HEAs possesses a heterogeneous microstructure with three types of grains (a1, a2, a3) at multiple length scales, and dislocation-density heterogeneity (see Fig. 3).
3.2. Mechanical properties
Fig. 2. BSE image of Al0.6CoCrFeNi HEA after cold-rolling and annealing; (b) BSE image of the red rectangle in (a); (ceg) the corresponding EDS elemental maps. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
The heterogeneous structure in alloys and dislocation structure formed by high-density dislocations in the non-recrystallized regions are clearly shown in the IPF map. Moreover, local misorientations across the non-recrystallized grains and recrystallized grains are observed, as evidenced in the EBSD kernel average misorientation (KAM) map shown in Fig. 3(b). According to the KAM map, two types of misorientations are clearly distinguished. The KAM values of the recrystallized regions are less than 0.5 ,
Fig. 4(a) shows representative engineering stress-strain curves of the uniaxial tensile test of Al0.6CoCrFeNi HEAs at 77 K and 298 K. The average yielding strength is determined to be about 786 MPa and 964 MPa, in addition, the ultimate tensile strength is determined to be about 1101 MPa and 1422 MPa for the alloys tested at 298 K and 77 K respectively. Compared with the elongation at room temperature, the elongation at 77 K were reduced slightly, but still reached 22.3%. Overally, Al0.6CoCrFeNi HEAs with heterogeneous structure have high strength and good plasticity at both 77 K and 298 K. This is mainly attributed to the high-density dislocations in the non-recrystallized regions and the ultrafine grains in the recrystallized region. Ultrafine grains in recrystallized regions have the potential of plastic deformation to provide plasticity for the alloys, while high dislocation density in the non-recrystallized regions can improve work-hardening ability of the materials and hence the strength of the alloys. According to Fig. 4(a), the corresponding true stress-strain curves are obtained, as shown in Fig. 4(b), and then the work hardening rate-true strain curves are also obtained according to the true stress-strain curves, which are shown in Fig. 4(c). The newly developed alloys show a very high work hardening behavior, especially at 77 K, the work hardening rate of the alloy still reaches 3.4 GPa before the sample is necked. Therefore, compared to 298 K, the alloys have a higher work hardening behavior at 77 K, so that the sample can reach a higher stress value before necking, thereby increasing the strength of the alloys. In addition, we compare tensile properties of Al0.6CoCrFeNi HEAs with those of reported fcc/L12 þ bcc/B2 dual-phase HEAs and some advanced traditional duplex stainless steels in Fig. 5 [10e18].
Fig. 3. EBSD results of Al0.6CoCrFeNi HEA after cold-rolling and annealing: (a) IPF map overlaid with IQ map; (b) KAM map; (c) the distributions of grain size.
4
Q. Li et al. / Journal of Alloys and Compounds 816 (2020) 152663
Fig. 4. (a) Engineering tensile strain-stress, (b) true tensile stress-strain and (c) work hardening rate versus true train curves of Al0.6CoCrFeNi HEAs in tension mode at 289 K and 77 K.
It is clear that in terms of product of strength and ductility and yield strength, Al0.6CoCrFeNi HEAs exhibit a better combination over existing dual-phase HEAs and traditional duplex stainless steels. 3.3. Deformation mechanism In order to characterize the underlying deformation mechanisms of Al0.6CoCrFeNi HEAs with heterogeneous structure under uniaxial tension at 298 K and 77 K. It is necessary to reveal the microstructural evolution during plastic deformation. Fig. 6(a and b) shows BF TEM images for the samples deformed at 298 K and 77 K. Other literatures have reported that the SFE of Al0.6CoCrFeNi HEAs is 147.4e162.0 mJ/m2 [19]. In addition, the SFE not only has a strong influence on the deformation mechanism, but it also controls the slip mode of the dislocations. A high SFE usually promotes cross-slip of dislocations [20]. So, a large amount of crossdislocation slip can be observed in samples that are deformed at room temperature. From the engineering stress-strain curve (Fig. 4(a)) at room temperature, it can be seen that the sample has a relatively high work hardening ability, while cross-slip of dislocations has little effect on the work hardening ability [20]. In addition, new slip planes will be required to accommodate the excess strain when initial slip planes are blocked, thus yielding the formation of high-dense dislocation walls (DWs) and small-sized dislocation cells (DCs) [21]. The high-dense DWs and small-sized DCs increase the resistance to dislocation movement, and hence enhance the work hardening ability of the material. This is to say; dislocation slip dominated the deformation of the sample at room
temperature. Some high-density dislocation structures are also observed in Fig. 6(b), but the difference is that a large number of deformation twins are activated at 77 K, and the average twin size is 10 nm. To further explore the microstructure evolution of Al0.6CoCrFeNi HEAs at 77 K, EBSD patterns after tensile deformation at 77 K are shown in Fig. 7. Obviously, a large number of deformed twins are generated in the tensile deformation at 77 K (highlighted by yellow lines in Fig. 7(a)), in addition, there are some deformation twins mainly concentrated in the fcc phase. According to Narita and Takamura [22], the critical resolved shear stress required for twinning is shown as:
sT ¼ 2gSFE
mbp
(1)
where sT is the critical stress for twinning formation, gSFE is the stacking fault energy, m is the average Schmid factor of about 0.33 and bp is the Burgers vector of the partial dislocations and bp ¼ 0.147 nm [6]. In addition, according to the thermodynamic analysis, without considering the interaction between the fcc and bcc phases in the Al0.6CoCrFeNi HEAs, the ideal gSFE of fcc phases in the alloys is then expressed by Ref. [23].
gSFE ¼ 2rDGg/εþ2s
(2)
where r is the molar surface density along {111} planes, DGg/ε is the molar Gibbs energy of the austenite to ε-martensite phase transformation gfcc /εMs and s is the interfacial energy per unit hcp area of the phase boundary. The molar surface density r is
Fig. 5. Yield strength versus product of strength and elongation of various traditional duplex stainless steels and some FCC/L12 þ BCC/B2 dual-phase HEAs including Al0.6CoCrFeNi HEAs.
Q. Li et al. / Journal of Alloys and Compounds 816 (2020) 152663
5
Fig. 6. TEM images of Al0.6CoCrFeNi HEA after tensile deformation at 298 K (a) and 77 K (b, c, d).
Fig. 7. EBSD results of Al0.6CoCrFeNi HEA after tensile deformation at 77 K: (a) Phase map (twins are highlighted by yellow lines), (b) misorientation angle. (For interpretation of the references to colour in this figure legend, the reader is referred to the Web version of this article.)
g/ε
calculated by Allain et al. as follows [24]:
4
1 3 a2 N
r ¼ pffiffiffi
(3)
where a is the lattice parameter of the alloy and N is Avogadro’s number. Considering a regular solution model, DGg/ε can be expressed for the present Al0.6CoCrFeNi alloys as follow [25e27]:
DGg/ε ¼
X 1X g/ε g/ε Xi DGi þ Xi Xj Uij 2 i ij
(4)
ði; j ¼ Al; Co; Cr; Fe; Ni; i s jÞ g/ε
where Xi and DG i represent the molar fraction and the differg/ε ence of free energy between fcc and hcp of pure metals. Uij is an interaction energy parameter for components i and j. Because of the content of Al in the fcc phases of the alloys are low, the interaction energy parameter for the Al and other elements can be ignored. Accordingly values of the parameters in the above formula are
g/ε
g/ε
g/ε
shown in Table 2. In addition, DGCo , UCoCr , UNiCo , UFeCo have not been reported in other literatures, so in this paper, according to equations (1) and (4) and the stacking fault energy of similar alloys that have been reported in other literatures, the following results
Table 2 Numerical values and functions used for the calculations. Parameters
Functions used (units)
g/ε DGFe g/ε DGCr g/ε DGNi g/ε DGAl g/ε UFeCr g/ε UFeNi g/ε UFeAl g/ε UCrNi
2243.38 þ 4.309T (J mol1)
[36]
1370e0.163T (J mol1)
[37]
1046 þ 1.255T (J mol1)
[36]
4190 (J mol1)
[35]
a
3.6006 Å 0.008 (J m2)
This work [36]
s
Reference
2800þ5T (J mol1)
[38]
2095 (J mol1)
[39]
2095 (J mol1)
[38]
3328 (J mol1)
[38]
6
Q. Li et al. / Journal of Alloys and Compounds 816 (2020) 152663
are obtained:
DGgCo/ε ¼ 3384:8 þ 0:35T Jmol1
/ε UgCoCr ¼ 2158:24 Jmol1
/ε UgNiCo ¼ 6777:24 Jmol1
/ε UgCoFe ¼ 6562:64 Jmol1
(5)
(6)
4. Conclusion
(7)
(8)
Substituting the obtained data into formulas (2), (3) and (4), the relationship between the stacking fault energy of the alloys and temperature is finally determined as follows:
gSFE ¼ 21:5 þ 0:0943T
(9)
Then, 77 K and 298 K are substituted into formula (9), and the stacking fault energy of the alloys at two temperatures is determined as follows:
.
gSFE ¼ 28:69 mJ m2 77K .
gSFE ¼ 49:33 mJ m2 298K
This is due to the “dynamic HallePetch” effect, which promotes the work hardening ability of Al0.6CoCrFeNi HEAs by activating the twin system at 77 K. Moreover, according to considerer’s criterion [35], the twinning-induced work hardening ability increases the flow stress of the necking of the sample to a higher stress value. This results in a higher strength of the alloy at 77 K compared to 298 K.
(10)
In conclusion, dual-phase Al0.6CoCrFeNi HEAs with heterogeneous structure are successfully prepared by arc melting, cold rolling, and partial recrystallization annealing. Tensile tests reveal that the present alloys exhibit excellent combination of strength and ductility for both 298 K and 77 K loading. Specifically, at 77 K, the yield strength and tensile strength reach 964 MPa and 1422 MPa, respectively, while the plasticity has a slight decrease. At 298 K, dislocation slip dominates the overall plastic deformation, while at 77 K deformation twins prevail (in addition to dislocations) when the true train is over 3.8% based on the critical stress formula. According to considerer’s criterion, the twinning-induced work hardening renders the onset of necking instability to a higher stress value, which results in the higher tensile strength at low temperature. Declaration of competing interest
(11)
Substituting formulas (10) and (11) into formula (1), the critical stress for twinning formation of the fcc region at 298 K and 77 K for the Al0.6CoCrFeNi HEAs are 2034 MPa and 1183 MPa respectively. It can be seen from Fig. 4(b) that the maximum tensile true stress of the sample is 1427 MPa at 298 K, which is much smaller than the critical stress for twinning formation. At 77 K, the maximum tensile true stress of the sample is 1740 MPa, which has exceeded the critical stress of twinning formation. This is to say, in the tensile deformation at 77 K, when the true strain of the sample reaches 3.8%, the formation of twins has begun. So, the generation of twins is not seen in Fig. 6(a), while a large number of twins can be seen in Fig. 6(b). Therefore, in the later stage of deformation at 77 K, the plastic deformation in the fcc phase are mainly induced by deformation twins. In addition, Table 1 can be known that the bcc phase are rich in a large amount of Al. In other literatures, it has been reported that Al has a high SFE of about 86 mJ/m2 [28,29]. Other literatures had also mentioned that the average stacking fault energy of Al0.6CoCrFeNi HEA is about 150 mJ/m2 [30], and we have confirmed that fcc phase have lower SFE, therefore, it can be inferred that the SEF of bcc phase is much larger than that of fcc phase. This is why no deformation twins are observed in the bcc phase. Although no deformation twins were observed in the bcc phase during the later deformation process, but dislocation slip must be generated to coordinate the plasticity increase caused by the deformation twins in the fcc phase. Therefore, the plastic deformation of Al0.6CoCrFeNi HEAs at 77 K can be divided into two stages: (1) when the true strain is between 0 and 3.8%, the plasticity is mainly caused by dislocation slip of fcc and bcc phases; (2) when the true strain is between 3.8% and sample fracture, the plasticity is mainly induced by deformation twins in fcc phase and dislocation slip in bcc phase. In addition, it has been reported in other literatures that the formation of deformation twins can induce the increase of work hardening [31,32]. According to “dynamic HallePetch” effect, the formation of deformation twin boundaries can also led to grain refinement, which reduced the dislocation mean free path and increased the strength of alloys [33,34]. In Fig. 4(c), the work hardening ability at 77 K is significantly higher than that at 298 K.
The manuscript submitted was the result of my independent research. Except as noted in the text, this article does not contain research results that have been published or written by other individuals or groups. The legal responsibility of this statement is borne by me. 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, Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes, Adv. Eng. Mater. 5 (2004) 299e304. [2] D.B. Miracle, O.N. Senkov, A critical review of high entropy alloys and related concepts, Acta Mater. 122 (2017) 448e511. [3] Y.J. He, H. Wang, H.L. Huang, A precipitation-hardened high-entropy alloy with outstanding tensile properties, Acta Mater. 102 (2016) 187e196. [4] Y. Deng, C.C. Tasan, K.G. Pradeep, H. Springer, A. Kostka, D. Raabe, Design of a twinning-induced plasticity high entropy alloy, Acta Mater. 94 (2015) 124e133. [5] Z.M. Li, K.G. Pradeep, Y. Deng, Metastable high-entropy dual-phase alloys overcome the strength-ductility trade-off, Nature 534 (2016) 227e230. [6] S.W. Wu, G. Wang, Q. Wang, Y.D. Jia, J. Yi, Q.J. Zhai, Enhancement of strengthductility trade-off in a high-entropy alloy through a heterogeneous structure, Acta Mater. 165 (2019) 444e458. [7] J.W. Qiao, S.G. Ma, E.W. Huang, C.P. Chuang, P.K. Liaw, Y. Zhang, Microstructural characteristics and mechanical behaviors of AlCoCrFeNi high-entropy alloys at ambient and cryogenic temperatures, Mater. Sci. Forum 688 (2011) 419e425. [8] B. Gludovatz, A. Hohenwarter, D. Catoor, E.H. Chang, E.P. George, R.O. Ritchie, A fracture-resistant high-entropy alloy for cryogenic applications, Science 345 (2014) 1153e1158.
Q. Li et al. / Journal of Alloys and Compounds 816 (2020) 152663 [9] T. Ericsson, The temperature and concentration dependence of the stacking fault energy in the Co-Ni system, Acta Metall. 14 (1966) 853e865. [10] K. Chandra, R. Singhal, V. Kain, V.S. Raja, Low temperature embrittlement of duplex stainless steel: correlation between mechanical and electrochemical behavior, Mater. Sci. Eng. A 527 (2010) 3904e3912. , Behavior of duplex stainless-steel casting [11] S. Jayet-Gendrot, P. Gilles, C. Migne defects under mechanical loadings, Nucl. Eng. Des. 197 (2000) 141e153. [12] E. Medina, J.M. Medina, A. Cobo, D.M. Bastidas, Evaluation of mechanical and structural behavior of austenitic and duplex stainless-steel reinforcements, Constr. Build. Mater. 78 (2015) 1e7. [13] M.K. Mishra, G. Gunasekaran, A.G. Rao, B.P. Kashyap, N. Prabhu, Effect of multipass friction stir processing on mechanical and corrosion behavior of 2507 super duplex stainless steel, Constr. Build. Mater. 26 (2017) 849e860. [14] M. Okayasu, D. Ishida, Effect of microstructural characteristics on mechanical properties of austenitic, ferritic, and g-a duplex stainless steels, Metall. Mater. Trans. A 50 (2019) 1380e1388. [15] C.S. Marchi, J. Zelinski, X. Tang, G.H. Schiroky, Mechanical properties of super duplex stainless steel 2507 after gas phase thermal precharging with hydrogen, Metall. Mater. Trans. A 38 (2007) 2763e2775. [16] L. Chen, F.P. Yuan, P. Jiang, X.L. Wu, Mechanical properties and nanostructures in a duplex stainless steel subjected to equal channel angular pressing, Mater. Sci. Eng. A 551 (2012) 154e159. [17] S. Niu, H. Kou, T. Guo, Y. Zhang, J. Wang, J. Li, Strengthening of nanoprecipitations in an annealed Al0.5CoCrFeNi high entropy alloy, Mater. Sci. Eng. A 671 (2016) 82e86. [18] Y.J. Liang, L.J. Wang, Y.R. Wen, B.Y. Cheng, Q.L. Wu, T.Q. Cao, Qian Xiao, Y.Y. Xue, G. Sha, Y.D. Wang, Y. Ren, X.Y. Li, L. Wang, F.C. Wang, H.N. Cai, Highcontent ductile coherent nanoprecipitates achieve ultrastrong high-entropy alloys, Nat. Commun. 9 (2018) 1e8. [19] T.S. Cao, J.L. Shang, J. Zhao, C.Q. Cheng, R. Wang, H. Wang, The influence of Al elements on the structure and the creep behavior of AlxCoCrFeNi high entropy alloys, Mater. Lett. 164 (2016) 344e347. [20] E. Welsch, D. Ponge, S.M. Hafez Haghighat, S. Sandlobes, P. Choi, M. Herbig, S. Zaefferer, D. Raabe, Strain hardening by dynamic slip band refinement in a high-Mn lightweight steel, Acta Mater. 116 (2016) 188e199. [21] F. Otto, A. Dlouhy, Ch Somsen, H. Bei, G. Eggeler, E.P. George, The influences of temperature and microstructure on the tensile properties of a CoCrFeMnNi high-entropy alloy, Acta Mater. 61 (2013) 5743e5755. [22] T.S. Byun, On the stress dependence of partial dislocation separation and deformation microstructure in austenitic stainless steels, Acta Mater. 51 (2003) 3063e3071. [23] G.B. Olson, M. Cohen, A General mechanism of martensitic nucleation: part III. kinetics of martensitic nucleation, Metall. Mater. Trans. A 7 (1976) 1915e1923.
7
€nninen, Thermodynamic [24] S. Curtze, V.-T. Kuokkala, A. Oikari, J. Talonen, H. Ha modeling of the stacking fault energy of austenitic steels, Acta Mater. 59 (2011) 1068e1076. [25] W.S. Yang, C.M. Wan, The influence of aluminium content to the stacking fault energy in Fe-Mn-Al-C alloy system, J. Mater. Sci. 25 (1990) 1821e1823. [26] Y.Q. Wang, B. Liu, K. Yan, M.S. Wang, S. Kabra, Y.L. Chiu, D. Dye, P.D. Lee, Y. Liu, B. Cai, Probing deformation mechanisms of a FeCoCrNi high-entropy alloy at 293 and 77K using in situ neutron diffraction, Acta Mater. 154 (2018) 79e89. [27] Y.H. Zhang, Y. Zhuang, A. Hu, J.J. Kai, C.T. Liu, The origin of negative stacking fault energies and nano-twin formation in face-centered cubic high entropy alloys, Scr. Mater. 130 (2017) 96e99. [28] B. Hammer, K.W. Jacobsen, V. Milman, M.C. Payne, Stacking fault energies in aluminium, J. Phys. Condens. Matter 4 (1992) 10453e10460. [29] L.L. Ma, L. Wang, Z.H. Nie, F.C. Wang, Y.F. Xue, J.L. Zhou, T.Q. Cao, Y.D. Wang, Y. Ren, Reversible deformation-induced martensitic transformation in Al0.6CoCrFeNi high-entropy alloy investigated by in situ synchrotron-based high-energy X-ray diffraction, Acta Mater. 128 (2017) 12e21. [30] T.S. Cao, J.L. Shang, J. Zhao, C.Q. Cheng, R. Wang, H. Wang, The influence of Al elements on the structure and the creep behavior of AlxCoCrFeNi high entropy alloys, Mater. Lett. 164 (2016) 344e347. [31] H. Beladi, I.B. Timokhina, Y. Estrin, J. Kim, B.C. De Cooman, S.K. Kim, Orientation dependence of twinning and strain hardening behaviour of a high manganese twinning induced plasticity steel with polycrystalline structure, Acta Mater. 59 (2011) 7787e7799. [32] I. Gutierrez-Urrutia, D. Raabe, Dislocation and twin substructure evolution during strain hardening of an Fee22 wt.% Mne0.6 wt.% C TWIP steel observed by electron channeling contrast imaging, Acta Mater. 59 (2011) 6449e6462. [33] O. Bouaziz, N. Guelton, Modelling of TWIP effect on work-hardening, Mater. Sci. Eng. A 319 (2001) 246e249. , [34] I. Karaman, H. Sehitoglu, A.J. Beaudoin, Y.I. Chumlyakov, H.J. Maier, C.N. Tome Modeling the deformation behavior of Hadfield steel single and polycrystals due to twinning and slip, Acta Mater. 48 (2000) 2031e2047. [35] A. Gali, E.P. George, Tensile properties of high- and medium-entropy alloys, Intermetallics 39 (2013) 74e78. [36] K. Ishida, Direct estimation of stacking fault energy by thermodynamic analysis, Phys. Status Solidi A 36 (1979) 717e728. [37] G. Inden, The role of magnetism in the calculation of phase diagrams, Phys. BþC 103 (1981) 82e100. [38] Y.H. Zhang, Y. Zhuang, A. Hu, J.J. Kai, C.T. Liu, The origin of negative stacking fault energies and nano-twin formation in face-centered cubic high entropy alloys, Scr. Mater. 130 (2017) 96e99. [39] I.A. Yakubtsov, A. Ariapour, D.D. Perovic, Effect of nitrogen on stacking fault energy of fcc iron-based alloys, Acta Mater. 47 (1999) 1271e1279.