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Precipitation behavior of selective laser melted FeCoCrNiC0.05 high entropy alloy
Intermetallics 106 (2019) 20–25
Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Precipitation behavior of selective laser melted FeCoCrNiC0.05 high entropy alloy
T
Rui Zhoua, Yong Liua,∗, Bin Liua, Jia Lib, Qihong Fangb,∗∗ a b
State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, PR China State Key Laboratory of Advanced Design and Manufacturing for Vehicle Body, Hunan University, Changsha, 410082, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: High entropy alloy Selective laser melting Mechanical properties Precipitation kinetics Hardening mechanism
Interstitial elements are effective to strengthen high entropy alloys (HEAs). In this work, FeCoCrNiC0.05 was prepared by selective laser melting (SLM) followed by annealing. The effects of annealing on the microstructures and mechanical properties of the SLMed FeCoCrNiC0.05 were investigated. Results show that, nano-scale Cr23C6type carbides can precipitate under annealing conditions, leading to higher yielding strength. The SLMed FeCoCrNiC0.05 annealed at 1073 K for 0.5 h has a yielding strength of 787 MPa and an elongation of 10.3%. Precipitation kinetics in SLMed FeCoCrNiC0.05 has been established according to Avrami formula. The high strength can be attributed to solid solution hardening, precipitation hardening and cell-like structures.
1. Introduction High entropy alloys (HEAs) are a type of alloys which have multiple constituents with equimolar or near-equimolar compositions [1,2]. A large number of alloying elements in a wide range of composition can form HEAs. HEAs usually exhibit a combination of good mechanical properties, excellent corrosion and oxidation resistance [3–5]. FeCoCrNi is a widely studied HEA with a simple fcc structure. The alloy has very good ductility, however for high-performance applications, the alloy should be strengthened [6]. Precipitation hardening has shown to be an effective way. A significant increase of strength was observed in FeCoCrNi HEA with the addition of Mo, because Mo can induce the precipitation of fine hard σ and μ phases in fcc matrix [7]. The simultaneous additions of Ti and Al also lead to a pronounced strengthening in FeCoNiCr HEA matrix, for precipitating of nano-sized coherent L12-Ni3(Ti, Al) reinforcing phase [8]. R. Banerjee have also reported Al-rich L12 precipitations in Al0.3CuFeCrNi2 HEA by appropriate thermo-mechanical treatment, leading to uniform distribution of L12(γ’)/Cu-rich precipitates and high mechanical properties [9]. FeCoCrNiMnC and AlFeCoNiC are also strengthened by the precipitation of carbides [10,11]. In order to achieve reasonable precipitation strengthening, a clear understanding of precipitation behavior is necessary. Selective laser melting (SLM) is a rapid manufacturing technique, which allows the components being manufactured directly from alloy
∗
powder [12]. This simplified process shortens the production period, and provides an efficient way to produce complex-shaped components. During the SLM process, HEA powders are melted by laser, and followed by rapid solidification. The fast cooling rate of solidification leads to fine microstructures and uniform chemical compositions [13]. Meanwhile, the fast cooling rate will also induce supersaturated solid solution in HEAs, which are intrinsically diffusion-difficult. Thus, SLMed HEA provides an appropriate model for studying the precipitation behavior. It has been reported that a pure ε phase can precipitate in SLMed CoCrMo alloys after annealing, and increased the hardness [14]. Moreover, a high volume fraction of γ″ phase in SLMed Inconel 718 alloy also enhanced the hardness [15]. However, the comprehensive study on the precipitation behavior of SLMed HEAs is still rare. In this work, FeCoCrNiC0.05 HEA was fabricated by SLM. After heat treatments, nanoparticles strengthened FeCoCrNiC0.05 was obtained. The precipitation behavior and strengthening mechanisms were studied. 2. Experimental 2.1. Material preparation and heat treatment The SLMed samples were fabricated from pre-alloyed powder, and the mean particle size was about 45 μm. The morphology of the prealloyed powder was spherical, and the chemical composition was
Corresponding author. Corresponding author. E-mail addresses: [email protected] (Y. Liu), [email protected] (Q. Fang).
∗∗
https://doi.org/10.1016/j.intermet.2018.12.001 Received 20 July 2018; Received in revised form 20 October 2018; Accepted 2 December 2018 0966-9795/ © 2018 Published by Elsevier Ltd.
Intermetallics 106 (2019) 20–25
R. Zhou et al.
Table 1 Chemical compositions of FeCoCrNiC0.05 powder. Element (at. %)
Fe
Co
Cr
Ni
C
O
FeCoCrNiC0.05
27.1
24.0
22.5
25.2
1.21
0.015
shown in Table 1. An SLM machine (FS271M, Farsoon, China) with an output laser power of 400 W and a scanning speed of 800 mm/s has been employed to fabricate standard-sized tensile specimens. All the samples were prepared in nitrogen atmosphere to minimize oxidation. When a layer had been melted, the laser rotated 67° and started to scan the next layer. For releasing residual stress, the samples were then annealed at 673 K for 3 h and quenched in water. The samples were annealed at 1073 K for 0.5–14 h followed by cooling in air. 2.2. Microstructural characterization Fig. 2. Hardness of the SLMed FeCoCrNiC0.05 HEA annealed at 1073 K for different time and the corresponding distributions of C and Cr elements in EPMA images.
For the phase constitutions, SLMed FeCoCrNiC0.05 annealed at 1073 K for various time were examined using an X-ray diffractometer (XRD) with a Cu Kα radiation. Microstructures were characterized by an FEI Tecnai G2 F20 transmission electron microscopy (TEM) and an FEI Quanta 650 scanning electron microscope (SEM) equipped with an electron backscattered diffraction (EBSD) analyzer. TEM samples were prepared by spark erosion, and then electropolished using a solution of 5% (vol. %) perchloric acid and 95% alcohol with a voltage of 30 V at a temperature of −35 °C. Compositions were analyzed using Electron Probe Micro Analysis (EPMA, JXA-8530F, Japan).
found within the detecting limitation of the XRD. The influence of annealing time conditions on the hardness of FeCoCrNiC0.05 HEA is shown in Fig. 2. The hardness increases significantly, and reaches the highest value of 340 HV at 0.5 h. As the annealing time prolongs from 0.5 to 8 h, the hardness decreases from 337 to 305 HV. Fig. 2 also shows the C and Cr distributions in FeCoCrNiC0.05 annealed for various time. At 0.5 h, a slight segregation of C occurs near the grain boundary, but Cr still distributes uniformly. That may be caused by the small atomic size and fast diffusion of C. After 0.5 h, Cr starts to segregate and large precipitates form in the matrix. At 0.5 h, the significant increment of hardness may be contributed by the segregation of C and Cr on grain boundaries. After 0.5 h, the decrease of the content of carbon dissolved in matrix causes a sustained hardness drop. Moreover, the recovery by annihilation of dislocations may be another reason for the decrease of hardness. Fig. 3 presents the hardness and loading-displacement curves for different phases. Obviously, the carbide phase possesses a much higher hardness (1157 HV) than FCC matrix (440 HV). The newly-formed carbide is generally considered to improve the strength [16]. Fig. 4 shows the microstructures of SLMed FeCoCrNiC0.05 HEA annealed at 1073 K for different time. The grains are equiaxed and randomly orientated. With the prolongation of the annealing time, the sizes of grains are also not changed obviously (about 40–50 μm). Because, compared with HEAs fabricated by conventional methods, SLMed FeCoCrNiC0.05 has a higher density of defects [13]. These
2.3. Mechanical tests Standardized dog-bone-shaped tensile specimens with a thickness of 2.5 mm and a gauge length of 10 mm were prepared by SLM. Tensile tests were conducted using an Instron 3369 machine equipped with laser extensometer at an engineering strain rate of 10−3/s and room temperature. Hardness and microhardness were conducted using a Vickers hardness tester and a nanoindenter (UNHTL + MCT, Switzerland) under a load of 300 g for 15 s and 10 mN for 15 s. At least 5 measurements for each sample were used to determine the average value. 3. Results Fig. 1 shows the XRD patterns for the SLMed FeCoCrNiC0.05 and the alloys annealed at 1073 K for various time. The results of SLMed alloy have a mono-phase of FCC, and no carbide is indexed. After being heattreated, the alloy remains the same single phase, and no other phase is
Fig. 3. Load-displacement curves and hardness of SLMed FeCoCrNiC0.05 annealed at 1073 K for 8 h
Fig. 1. XRD patterns of the FeCoCrNiC0.05 annealed at 1073 K 21
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Fig. 4. EBSD analyses of as-SLMed FeCoCrNiC0.05 alloys annealed at 1073 K for (a) 0 h, (b) 0.5 h, (c) 2 h and (d) 8 h.
defects provide nucleation sites for carbides. Many fine carbides form and pin the grain boundaries. As a result, the grain growth is inhibited. Substructures in the annealed FeCoCrNiC0.05 HEA at 1073 K for 0 h, 0.5 h and 8 h are presented in Fig. 5. Before annealing treatment, lots of cellular structures (around 0.5–1 μm) can be found. After 0.5 h, nanoscale precipitates (80 nm) can be seen in the annealed FeCoCrNiC0.05. Cellular structures disappear and some dislocations are observed around the precipitates (Fig. 5b). These precipitates are square shaped, and the boundary is unclear. At 8 h, the size of precipitates increases to about 140 nm (Fig. 5c). From the analysis of TEM, it can be found that carbides nucleate mainly on boundaries of grains and cell-structures. Because nucleation energy in the sites of defects is relatively low and the precipitates nucleated more easily at these sites. The precipitates present regular hexagon shape, and the boundary is clear, so a high degree of crystallization of precipitates is obtained. Selected area diffraction (SAD) pattern in Fig. 5d shows that the precipitate is Cr23C6type carbide, with a fcc structure. Fig. 6a shows engineering stress-strain curves of the SLMed and annealed FeCoCrNiC0.05 (1073 K for 0.5 h and 8 h). The results reveal that the annealing treatment can lead to an increase both in yield strength and ultimate tensile strength, but a decrease in elongation. Compared with the SLMed FeCoCrNiC0.05 (YS = 708 MPa, UTS = 872 MPa) and FeCoCrNiC0.05 annealed at 1073 K for 8 h (YS = 720 MPa, UTS = 905 MPa), the FeCoCrNiC0.05 annealed at 1073 K for 0.5 h exhibits the highest strength (YS = 787 MPa, UTS = 950 MPa). Fig. 6b shows the strain hardening rates varying with true strain for the SLMed and annealed FeCoCrNiC0.05 HEA. Obviously, the annealing treatments increase the strain-hardening rate. It can be noted that after being annealed at 1073 K for 8 h, SLMed FeCoCrNiC0.05 still has high yield strength. SLMed FeCoCrNiC0.05 maintains good thermal stability of mechanical properties.
4. Discussion 4.1. Precipitation kinetics Avrami formula is effective to describe the precipitation kinetics. The precipitation kinetics in alloys can be described as follows [17].
f = 1 − exp(−kt n )
(1)
where f is the experimental volume fraction of the precipitates corresponding to annealing time, k is a parameter. It depends on the temperature, formation energy and interface energy of precipitation, but is independent from heat treatment time. t is the heat treatment time. Moreover, n works as a time index and it depends on the type of precipitations. According to Yong's work, if the nucleation occurred at dislocations and the rate decreased rapidly to zero, n would work as a value of 1. If the nucleation occurred in FCC matrix and the rate attenuated sharply to zero, n would work as a value of 1.5 [17]. Therefore, n should be chosen as the value of 1 in this work. In practice, for FeCoCrNiC0.05, the content of dissolved carbon and the volume fraction of precipitate both influence the yield strength and ultimate strength. fr defined as relative volume fraction for precipitation is introduced to evaluate the degree of Cr23C6-type carbides, and (1- fr) can be used to evaluate the content of dissolved carbon. fr can be described by the equation as follows,
fr = fexp / fmax
(2)
where fexp is the experimental volume fraction of the precipitates corresponding to annealing time, f max is the maximum volume fraction of precipitates and it is calculated based on the content of solute carbon. The relative volume fraction of precipitates varying with the elongation of annealing time is calculated and shown in Fig. 7. In this work, f r as the relative precipitation volume fraction is calculated on the analysis of microstructures. It can be described by a modified Avarmi formula, 22
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Fig. 5. Bright field TEM images of the as-annealed FeCoCrNiC0.05 HEA at 1073 K for (a) 0 h, (b) 0.5 h, (c) 8 h; and selected area diffraction pattern (d).
fr = M(1 − exp(−kt n ))
4.2. Strengthening mechanisms
(3)
where M is the maximum relative volume fraction of precipitations at definite temperature, t is the heat treatment time and n is the time index. At a certain temperature, k and n are seen as constants. In Fig. 7, the modified model fits quite well with the experimental data, and the fitting degrees (R2) are evaluated to be over 0.99. It means that model can be successfully used on predicting the precipitation behavior.
Strengthening mechanisms in SLMed FeCoCrNiC0.05 annealed at 1073 K can be summarized into four different aspects: solid-solution hardening, precipitation hardening, dislocation hardening and grainboundary hardening. Consequently, the yield strength can be calculated by microstructure-related equation [8]:
σy = σA + σss + σppt + σdis + σgb
(4)
where σA is the yield strength of the FeCoCrNi fabricated by casting [18], σss , σppt , σdis and σgb are the strength increments from solid-
Fig. 6. Typical engineering stress-strain curves (a) and strain-hardening rate curves (b) of the SLMed and annealed FeCoCrNiC0.05 (1073 K for 0.5 h and 8 h). 23
Intermetallics 106 (2019) 20–25
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Moreover, the refinement of grains leads to a high density of grain boundaries, and can impede the dislocation slip. In the SLMed FeCoCrNiC0.05, dislocation-walls play the same role with the grains boundaries. The relationship between the size of cell-forming structures and the yield strength can also be described by the Hall-Petch equation [21]: 1
σy = σ0 + k y / d 2
(6)
where σy represents yield strength, σ0 represents again lattice friction stress, k y represents Hall-Petch slope, and d represents average grain or cell-forming structures diameter. The size of cell-forming structures is far smaller than the grain size. Therefore, according to Eq. (3), yield strength increase (σc ) caused by the formation of cell-forming structure can be expressed as: −1 2
σdis + σgb = σc = k y ⎛dc ⎝ ⎜
Fig. 7. Relative fraction volume of precipitates in the SLMed FeCoCrNiC0.05 HEA annealed at 1073 K for different time and the precipitation kinetics curves fitted with formula.
−1
−1
+ dg12 − dg 22 ⎞ ⎠ ⎟
(7)
where dc represents the size of cell-forming structures of the SLMed FeCoCrNiC0.05 (about 0.5 μm) and dg1 and dg2 represent the size of grain structures of the FeCoCrNi prepared by SLM (about 50 μm), and cast one (about 289.7 μm) [18]. In this work, the value of k y is used as 226 MPa μm1/2 [6]. And, the yield strength improvement caused by cell-forming structure and grain boundary hardening is calculated to be 338 MPa. This result shows that the cell-forming hardening is also an important contribution for the improvement of yield strength. 5. Conclusions (1) Nano-scaled Cr23C6-type carbide can precipitate in the SLMed FeCoCrNiC0.05 being annealed at 1073 K. The precipitates are mainly distributed on boundaries of grains or cell structures. (2) The precipitation behavior can be successfully predicted by the modified Avarmi formula. (3) SLMed FeCoCrNiC0.05 annealed at 1073 K for 0.5 h has the highest yield strength of 787 MPa and a reasonable elongation. The strengthening effect can be quantitatively calculated by considering solid solution hardening, precipitation hardening and cell structure hardening.
Fig. 8. The strength contributions from different hardening mechanisms.
Acknowledgement
solution, precipitation, dislocation and grain-boundary, respectively. In HEA, it has been observed that the yield strength increases linearly with the increase of C content at room temperature [19]. In this work, the increment in yield strength (σy ) per atomic percent of carbon is 155 MPa. Fine Cr23C6-type carbides will produce hardening effect by Orowan mechanism, and the value of strengthening σppt can be calculated by the equation [20]:
σppt = (0.538 ∗ Gbf 0.5 /D)ln(D /2b)
The authors would like to appreciate the support from the National Natural Science Foundation of China under grant No. 51671217. References [1] K.Y. Tsai, M.H. Tsai, J.W. Yeh, Sluggish diffusion in Co–Cr–Fe–Mn–Ni high-entropy alloys, Acta Mater. 61 (13) (2013) 4887–4897. [2] B. Cantor, I.T.H. Chang, P. Knight, A.J.B. Vincent, Microstructural development in equiatomic multicomponent alloys, Mater. Sci. Eng., A 375 (Supplement C) (2004) 213–218. [3] L. Zhang, Y. Zhou, X. Jin, X. Du, B. Li, Precipitation-hardened high entropy alloys with excellent tensile properties, Mater. Sci. Eng., A (2018) 186–191. [4] T. Fujieda, H. Shiratori, K. Kuwabara, M. Hirota, T. Kato, K. Yamanaka, Y. Koizumi, A. Chiba, S. Watanabe, CoCrFeNiTi-based high-entropy alloy with superior tensile strength and corrosion resistance achieved by a combination of additive manufacturing using selective electron beam melting and solution treatment, Mater. Lett. 189 (Supplement C) (2017) 148–151. [5] T.M. Butler, M.L. Weaver, Oxidation behavior of arc melted AlCoCrFeNi multicomponent high-entropy alloys, J. Alloy. Comp. 674 (2016) 229–244. [6] B. Liu, J. Wang, Y. Liu, Q. Fang, Y. Wu, S. Chen, C.T. Liu, Microstructure and mechanical properties of equimolar FeCoCrNi high entropy alloy prepared via powder extrusion, Intermetallics 75 (Supplement C) (2016) 25–30. [7] W.H. Liu, Z.P. Lu, J.Y. He, J.H. Luan, Z.J. Wang, B. Liu, Y. Liu, M.W. Chen, C.T. Liu, Ductile CoCrFeNiMox high entropy alloys strengthened by hard intermetallic phases, Acta Mater. 116 (2016) 332–342. [8] 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, A precipitation-hardened high-entropy alloy with outstanding tensile properties, Acta Mater. 102 (2016) 187–196. [9] B. Gwalani, D. Choudhuri, V. Soni, Y. Ren, M. Styles, J.Y. Hwang, S.J. Nam, H. Ryu, S.H. Hong, R. Banerjee, Cu assisted stabilization and nucleation of L12 precipitates
(5)
where G is the shear modulus, and that can be calculated from G= E/2(1 + v) , which assumes the HEA is elastically isotropic. The value of Poisson's ratio (v) is 0.3 and the Young's modulus (E) can be measured by nanoindentation. b is the burgers vectors of the HEA and f is the volume fraction of precipitates. D is the spatial diameter of precipitates. According to Fig. 6, the SLMed FeCoCrNiC0.05 annealed for 0.5 h has the highest value of yield strength. When the annealing time elongated to 8 h, the yield strength decreased. The reason can be clarified in Fig. 8, which shows the calculated strength contribution from solidsolution hardening and precipitation hardening. At early annealing stage, a sharp increase occurs in the precipitation hardening, but the solid solution hardening decreases slightly. When increasing annealing time to 8 h, the precipitation hardening increases slightly, but the solid solution hardening decreases significantly. As a result, the yield strength decreases. 24
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[16] S. Nam, M.J. Kim, J.Y. Hwang, H. Choi, Strengthening of Al0.15CoCrCuFeNiTix–C (x = 0, 1, 2) high-entropy alloys by grain refinement and using nanoscale carbides via powder metallurgical route, J. Alloy. Comp. 762 (2018) 29–37. [17] Z. Peng, L. Li, S. Chen, X. Huo, J. Gao, Isothermal precipitation kinetics of carbides in undercooled austenite and ferrite of a titanium microalloyed steel, Mater. Des. 108 (2016) 289–297. [18] A. Gali, E.P. George, Tensile properties of high- and medium-entropy alloys, Intermetallics 39 (2013) 74–78. [19] Z. Wang, I. Baker, Z. Cai, S. Chen, J.D. Poplawsky, W. Guo, The effect of interstitial carbon on the mechanical properties and dislocation substructure evolution in Fe40.4Ni11.3Mn34.8Al7.5Cr6 high entropy alloys, Acta Mater. 120 (2016) 228–239. [20] Z. Wang, I. Baker, W. Guo, J.D. Poplawsky, The effect of carbon on the microstructures, mechanical properties, and deformation mechanisms of thermo-mechanically treated Fe40.4Ni11.3Mn34.8Al7.5Cr6 high entropy alloys, Acta Mater. 126 (2017) 346–360. [21] F. Yin, G.J. Cheng, R. Xu, K. Zhao, Q. Li, J. Jian, S. Hu, S. Sun, L. An, Q. Han, Ultrastrong nanocrystalline stainless steel and its Hall-Petch relationship in the nanoscale, Scripta Mater. 155 (2018) 26–31.
in Al0.3CuFeCrNi2 fcc-based high entropy alloy, Acta Mater. 129 (2017) 170–182. [10] N.D. Stepanov, N.Y. Yurchenko, M.A. Tikhonovsky, G.A. Salishchev, Effect of carbon content and annealing on structure and hardness of the CoCrFeNiMn-based high entropy alloys, J. Alloy. Comp. 687 (2016) 59–71. [11] J.T. Fan, L.J. Zhang, P.F. Yu, M.D. Zhang, D.J. Liu, Z. Zhou, P. Cui, M.Z. Ma, Q. Jing, G. Li, R.P. Liu, Improved the microstructure and mechanical properties of AlFeCoNi high-entropy alloy by carbon addition, Mater. Sci. Eng., A 728 (2018) 30–39. [12] R. Li, P. Niu, T. Yuan, P. Cao, C. Chen, K. Zhou, Selective laser melting of an equiatomic CoCrFeMnNi high-entropy alloy: processability, non-equilibrium microstructure and mechanical property, J. Alloy. Comp. 746 (2018) 125–134. [13] R. Zhou, Y. Liu, C. Zhou, S. Li, W. Wu, M. Song, B. Liu, X. Liang, P.K. Liaw, Microstructures and mechanical properties of C-containing FeCoCrNi high-entropy alloy fabricated by selective laser melting, Intermetallics 94 (2018) 165–171. [14] M. Zhang, Y. Yang, C. Song, Y. Bai, Z. Xiao, An investigation into the aging behavior of CoCrMo alloys fabricated by selective laser melting, J. Alloy. Comp. 750 (2018) 878–886. [15] X. Li, J.J. Shi, C.H. Wang, G.H. Cao, A.M. Russell, Z.J. Zhou, C.P. Li, G.F. Chen, Effect of heat treatment on microstructure evolution of Inconel 718 alloy fabricated by selective laser melting, J. Alloy. Comp. 764 (2018) 639–649.
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Contents lists available at ScienceDirect
Intermetallics journal homepage: www.elsevier.com/locate/intermet
Precipitation behavior of selective laser melted FeCoCrNiC0.05 high entropy alloy
T
Rui Zhoua, Yong Liua,∗, Bin Liua, Jia Lib, Qihong Fangb,∗∗ a b
State Key Laboratory of Powder Metallurgy, Central South University, Changsha, 410083, PR China State Key Laboratory of Advanced Design and Manufacturing for Vehicle Body, Hunan University, Changsha, 410082, PR China
A R T I C LE I N FO
A B S T R A C T
Keywords: High entropy alloy Selective laser melting Mechanical properties Precipitation kinetics Hardening mechanism
Interstitial elements are effective to strengthen high entropy alloys (HEAs). In this work, FeCoCrNiC0.05 was prepared by selective laser melting (SLM) followed by annealing. The effects of annealing on the microstructures and mechanical properties of the SLMed FeCoCrNiC0.05 were investigated. Results show that, nano-scale Cr23C6type carbides can precipitate under annealing conditions, leading to higher yielding strength. The SLMed FeCoCrNiC0.05 annealed at 1073 K for 0.5 h has a yielding strength of 787 MPa and an elongation of 10.3%. Precipitation kinetics in SLMed FeCoCrNiC0.05 has been established according to Avrami formula. The high strength can be attributed to solid solution hardening, precipitation hardening and cell-like structures.
1. Introduction High entropy alloys (HEAs) are a type of alloys which have multiple constituents with equimolar or near-equimolar compositions [1,2]. A large number of alloying elements in a wide range of composition can form HEAs. HEAs usually exhibit a combination of good mechanical properties, excellent corrosion and oxidation resistance [3–5]. FeCoCrNi is a widely studied HEA with a simple fcc structure. The alloy has very good ductility, however for high-performance applications, the alloy should be strengthened [6]. Precipitation hardening has shown to be an effective way. A significant increase of strength was observed in FeCoCrNi HEA with the addition of Mo, because Mo can induce the precipitation of fine hard σ and μ phases in fcc matrix [7]. The simultaneous additions of Ti and Al also lead to a pronounced strengthening in FeCoNiCr HEA matrix, for precipitating of nano-sized coherent L12-Ni3(Ti, Al) reinforcing phase [8]. R. Banerjee have also reported Al-rich L12 precipitations in Al0.3CuFeCrNi2 HEA by appropriate thermo-mechanical treatment, leading to uniform distribution of L12(γ’)/Cu-rich precipitates and high mechanical properties [9]. FeCoCrNiMnC and AlFeCoNiC are also strengthened by the precipitation of carbides [10,11]. In order to achieve reasonable precipitation strengthening, a clear understanding of precipitation behavior is necessary. Selective laser melting (SLM) is a rapid manufacturing technique, which allows the components being manufactured directly from alloy
∗
powder [12]. This simplified process shortens the production period, and provides an efficient way to produce complex-shaped components. During the SLM process, HEA powders are melted by laser, and followed by rapid solidification. The fast cooling rate of solidification leads to fine microstructures and uniform chemical compositions [13]. Meanwhile, the fast cooling rate will also induce supersaturated solid solution in HEAs, which are intrinsically diffusion-difficult. Thus, SLMed HEA provides an appropriate model for studying the precipitation behavior. It has been reported that a pure ε phase can precipitate in SLMed CoCrMo alloys after annealing, and increased the hardness [14]. Moreover, a high volume fraction of γ″ phase in SLMed Inconel 718 alloy also enhanced the hardness [15]. However, the comprehensive study on the precipitation behavior of SLMed HEAs is still rare. In this work, FeCoCrNiC0.05 HEA was fabricated by SLM. After heat treatments, nanoparticles strengthened FeCoCrNiC0.05 was obtained. The precipitation behavior and strengthening mechanisms were studied. 2. Experimental 2.1. Material preparation and heat treatment The SLMed samples were fabricated from pre-alloyed powder, and the mean particle size was about 45 μm. The morphology of the prealloyed powder was spherical, and the chemical composition was
Corresponding author. Corresponding author. E-mail addresses: [email protected] (Y. Liu), [email protected] (Q. Fang).
∗∗
https://doi.org/10.1016/j.intermet.2018.12.001 Received 20 July 2018; Received in revised form 20 October 2018; Accepted 2 December 2018 0966-9795/ © 2018 Published by Elsevier Ltd.
Intermetallics 106 (2019) 20–25
R. Zhou et al.
Table 1 Chemical compositions of FeCoCrNiC0.05 powder. Element (at. %)
Fe
Co
Cr
Ni
C
O
FeCoCrNiC0.05
27.1
24.0
22.5
25.2
1.21
0.015
shown in Table 1. An SLM machine (FS271M, Farsoon, China) with an output laser power of 400 W and a scanning speed of 800 mm/s has been employed to fabricate standard-sized tensile specimens. All the samples were prepared in nitrogen atmosphere to minimize oxidation. When a layer had been melted, the laser rotated 67° and started to scan the next layer. For releasing residual stress, the samples were then annealed at 673 K for 3 h and quenched in water. The samples were annealed at 1073 K for 0.5–14 h followed by cooling in air. 2.2. Microstructural characterization Fig. 2. Hardness of the SLMed FeCoCrNiC0.05 HEA annealed at 1073 K for different time and the corresponding distributions of C and Cr elements in EPMA images.
For the phase constitutions, SLMed FeCoCrNiC0.05 annealed at 1073 K for various time were examined using an X-ray diffractometer (XRD) with a Cu Kα radiation. Microstructures were characterized by an FEI Tecnai G2 F20 transmission electron microscopy (TEM) and an FEI Quanta 650 scanning electron microscope (SEM) equipped with an electron backscattered diffraction (EBSD) analyzer. TEM samples were prepared by spark erosion, and then electropolished using a solution of 5% (vol. %) perchloric acid and 95% alcohol with a voltage of 30 V at a temperature of −35 °C. Compositions were analyzed using Electron Probe Micro Analysis (EPMA, JXA-8530F, Japan).
found within the detecting limitation of the XRD. The influence of annealing time conditions on the hardness of FeCoCrNiC0.05 HEA is shown in Fig. 2. The hardness increases significantly, and reaches the highest value of 340 HV at 0.5 h. As the annealing time prolongs from 0.5 to 8 h, the hardness decreases from 337 to 305 HV. Fig. 2 also shows the C and Cr distributions in FeCoCrNiC0.05 annealed for various time. At 0.5 h, a slight segregation of C occurs near the grain boundary, but Cr still distributes uniformly. That may be caused by the small atomic size and fast diffusion of C. After 0.5 h, Cr starts to segregate and large precipitates form in the matrix. At 0.5 h, the significant increment of hardness may be contributed by the segregation of C and Cr on grain boundaries. After 0.5 h, the decrease of the content of carbon dissolved in matrix causes a sustained hardness drop. Moreover, the recovery by annihilation of dislocations may be another reason for the decrease of hardness. Fig. 3 presents the hardness and loading-displacement curves for different phases. Obviously, the carbide phase possesses a much higher hardness (1157 HV) than FCC matrix (440 HV). The newly-formed carbide is generally considered to improve the strength [16]. Fig. 4 shows the microstructures of SLMed FeCoCrNiC0.05 HEA annealed at 1073 K for different time. The grains are equiaxed and randomly orientated. With the prolongation of the annealing time, the sizes of grains are also not changed obviously (about 40–50 μm). Because, compared with HEAs fabricated by conventional methods, SLMed FeCoCrNiC0.05 has a higher density of defects [13]. These
2.3. Mechanical tests Standardized dog-bone-shaped tensile specimens with a thickness of 2.5 mm and a gauge length of 10 mm were prepared by SLM. Tensile tests were conducted using an Instron 3369 machine equipped with laser extensometer at an engineering strain rate of 10−3/s and room temperature. Hardness and microhardness were conducted using a Vickers hardness tester and a nanoindenter (UNHTL + MCT, Switzerland) under a load of 300 g for 15 s and 10 mN for 15 s. At least 5 measurements for each sample were used to determine the average value. 3. Results Fig. 1 shows the XRD patterns for the SLMed FeCoCrNiC0.05 and the alloys annealed at 1073 K for various time. The results of SLMed alloy have a mono-phase of FCC, and no carbide is indexed. After being heattreated, the alloy remains the same single phase, and no other phase is
Fig. 3. Load-displacement curves and hardness of SLMed FeCoCrNiC0.05 annealed at 1073 K for 8 h
Fig. 1. XRD patterns of the FeCoCrNiC0.05 annealed at 1073 K 21
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Fig. 4. EBSD analyses of as-SLMed FeCoCrNiC0.05 alloys annealed at 1073 K for (a) 0 h, (b) 0.5 h, (c) 2 h and (d) 8 h.
defects provide nucleation sites for carbides. Many fine carbides form and pin the grain boundaries. As a result, the grain growth is inhibited. Substructures in the annealed FeCoCrNiC0.05 HEA at 1073 K for 0 h, 0.5 h and 8 h are presented in Fig. 5. Before annealing treatment, lots of cellular structures (around 0.5–1 μm) can be found. After 0.5 h, nanoscale precipitates (80 nm) can be seen in the annealed FeCoCrNiC0.05. Cellular structures disappear and some dislocations are observed around the precipitates (Fig. 5b). These precipitates are square shaped, and the boundary is unclear. At 8 h, the size of precipitates increases to about 140 nm (Fig. 5c). From the analysis of TEM, it can be found that carbides nucleate mainly on boundaries of grains and cell-structures. Because nucleation energy in the sites of defects is relatively low and the precipitates nucleated more easily at these sites. The precipitates present regular hexagon shape, and the boundary is clear, so a high degree of crystallization of precipitates is obtained. Selected area diffraction (SAD) pattern in Fig. 5d shows that the precipitate is Cr23C6type carbide, with a fcc structure. Fig. 6a shows engineering stress-strain curves of the SLMed and annealed FeCoCrNiC0.05 (1073 K for 0.5 h and 8 h). The results reveal that the annealing treatment can lead to an increase both in yield strength and ultimate tensile strength, but a decrease in elongation. Compared with the SLMed FeCoCrNiC0.05 (YS = 708 MPa, UTS = 872 MPa) and FeCoCrNiC0.05 annealed at 1073 K for 8 h (YS = 720 MPa, UTS = 905 MPa), the FeCoCrNiC0.05 annealed at 1073 K for 0.5 h exhibits the highest strength (YS = 787 MPa, UTS = 950 MPa). Fig. 6b shows the strain hardening rates varying with true strain for the SLMed and annealed FeCoCrNiC0.05 HEA. Obviously, the annealing treatments increase the strain-hardening rate. It can be noted that after being annealed at 1073 K for 8 h, SLMed FeCoCrNiC0.05 still has high yield strength. SLMed FeCoCrNiC0.05 maintains good thermal stability of mechanical properties.
4. Discussion 4.1. Precipitation kinetics Avrami formula is effective to describe the precipitation kinetics. The precipitation kinetics in alloys can be described as follows [17].
f = 1 − exp(−kt n )
(1)
where f is the experimental volume fraction of the precipitates corresponding to annealing time, k is a parameter. It depends on the temperature, formation energy and interface energy of precipitation, but is independent from heat treatment time. t is the heat treatment time. Moreover, n works as a time index and it depends on the type of precipitations. According to Yong's work, if the nucleation occurred at dislocations and the rate decreased rapidly to zero, n would work as a value of 1. If the nucleation occurred in FCC matrix and the rate attenuated sharply to zero, n would work as a value of 1.5 [17]. Therefore, n should be chosen as the value of 1 in this work. In practice, for FeCoCrNiC0.05, the content of dissolved carbon and the volume fraction of precipitate both influence the yield strength and ultimate strength. fr defined as relative volume fraction for precipitation is introduced to evaluate the degree of Cr23C6-type carbides, and (1- fr) can be used to evaluate the content of dissolved carbon. fr can be described by the equation as follows,
fr = fexp / fmax
(2)
where fexp is the experimental volume fraction of the precipitates corresponding to annealing time, f max is the maximum volume fraction of precipitates and it is calculated based on the content of solute carbon. The relative volume fraction of precipitates varying with the elongation of annealing time is calculated and shown in Fig. 7. In this work, f r as the relative precipitation volume fraction is calculated on the analysis of microstructures. It can be described by a modified Avarmi formula, 22
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Fig. 5. Bright field TEM images of the as-annealed FeCoCrNiC0.05 HEA at 1073 K for (a) 0 h, (b) 0.5 h, (c) 8 h; and selected area diffraction pattern (d).
fr = M(1 − exp(−kt n ))
4.2. Strengthening mechanisms
(3)
where M is the maximum relative volume fraction of precipitations at definite temperature, t is the heat treatment time and n is the time index. At a certain temperature, k and n are seen as constants. In Fig. 7, the modified model fits quite well with the experimental data, and the fitting degrees (R2) are evaluated to be over 0.99. It means that model can be successfully used on predicting the precipitation behavior.
Strengthening mechanisms in SLMed FeCoCrNiC0.05 annealed at 1073 K can be summarized into four different aspects: solid-solution hardening, precipitation hardening, dislocation hardening and grainboundary hardening. Consequently, the yield strength can be calculated by microstructure-related equation [8]:
σy = σA + σss + σppt + σdis + σgb
(4)
where σA is the yield strength of the FeCoCrNi fabricated by casting [18], σss , σppt , σdis and σgb are the strength increments from solid-
Fig. 6. Typical engineering stress-strain curves (a) and strain-hardening rate curves (b) of the SLMed and annealed FeCoCrNiC0.05 (1073 K for 0.5 h and 8 h). 23
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Moreover, the refinement of grains leads to a high density of grain boundaries, and can impede the dislocation slip. In the SLMed FeCoCrNiC0.05, dislocation-walls play the same role with the grains boundaries. The relationship between the size of cell-forming structures and the yield strength can also be described by the Hall-Petch equation [21]: 1
σy = σ0 + k y / d 2
(6)
where σy represents yield strength, σ0 represents again lattice friction stress, k y represents Hall-Petch slope, and d represents average grain or cell-forming structures diameter. The size of cell-forming structures is far smaller than the grain size. Therefore, according to Eq. (3), yield strength increase (σc ) caused by the formation of cell-forming structure can be expressed as: −1 2
σdis + σgb = σc = k y ⎛dc ⎝ ⎜
Fig. 7. Relative fraction volume of precipitates in the SLMed FeCoCrNiC0.05 HEA annealed at 1073 K for different time and the precipitation kinetics curves fitted with formula.
−1
−1
+ dg12 − dg 22 ⎞ ⎠ ⎟
(7)
where dc represents the size of cell-forming structures of the SLMed FeCoCrNiC0.05 (about 0.5 μm) and dg1 and dg2 represent the size of grain structures of the FeCoCrNi prepared by SLM (about 50 μm), and cast one (about 289.7 μm) [18]. In this work, the value of k y is used as 226 MPa μm1/2 [6]. And, the yield strength improvement caused by cell-forming structure and grain boundary hardening is calculated to be 338 MPa. This result shows that the cell-forming hardening is also an important contribution for the improvement of yield strength. 5. Conclusions (1) Nano-scaled Cr23C6-type carbide can precipitate in the SLMed FeCoCrNiC0.05 being annealed at 1073 K. The precipitates are mainly distributed on boundaries of grains or cell structures. (2) The precipitation behavior can be successfully predicted by the modified Avarmi formula. (3) SLMed FeCoCrNiC0.05 annealed at 1073 K for 0.5 h has the highest yield strength of 787 MPa and a reasonable elongation. The strengthening effect can be quantitatively calculated by considering solid solution hardening, precipitation hardening and cell structure hardening.
Fig. 8. The strength contributions from different hardening mechanisms.
Acknowledgement
solution, precipitation, dislocation and grain-boundary, respectively. In HEA, it has been observed that the yield strength increases linearly with the increase of C content at room temperature [19]. In this work, the increment in yield strength (σy ) per atomic percent of carbon is 155 MPa. Fine Cr23C6-type carbides will produce hardening effect by Orowan mechanism, and the value of strengthening σppt can be calculated by the equation [20]:
σppt = (0.538 ∗ Gbf 0.5 /D)ln(D /2b)
The authors would like to appreciate the support from the National Natural Science Foundation of China under grant No. 51671217. References [1] K.Y. Tsai, M.H. Tsai, J.W. Yeh, Sluggish diffusion in Co–Cr–Fe–Mn–Ni high-entropy alloys, Acta Mater. 61 (13) (2013) 4887–4897. [2] B. Cantor, I.T.H. Chang, P. Knight, A.J.B. Vincent, Microstructural development in equiatomic multicomponent alloys, Mater. Sci. Eng., A 375 (Supplement C) (2004) 213–218. [3] L. Zhang, Y. Zhou, X. Jin, X. Du, B. Li, Precipitation-hardened high entropy alloys with excellent tensile properties, Mater. Sci. Eng., A (2018) 186–191. [4] T. Fujieda, H. Shiratori, K. Kuwabara, M. Hirota, T. Kato, K. Yamanaka, Y. Koizumi, A. Chiba, S. Watanabe, CoCrFeNiTi-based high-entropy alloy with superior tensile strength and corrosion resistance achieved by a combination of additive manufacturing using selective electron beam melting and solution treatment, Mater. Lett. 189 (Supplement C) (2017) 148–151. [5] T.M. Butler, M.L. Weaver, Oxidation behavior of arc melted AlCoCrFeNi multicomponent high-entropy alloys, J. Alloy. Comp. 674 (2016) 229–244. [6] B. Liu, J. Wang, Y. Liu, Q. Fang, Y. Wu, S. Chen, C.T. Liu, Microstructure and mechanical properties of equimolar FeCoCrNi high entropy alloy prepared via powder extrusion, Intermetallics 75 (Supplement C) (2016) 25–30. [7] W.H. Liu, Z.P. Lu, J.Y. He, J.H. Luan, Z.J. Wang, B. Liu, Y. Liu, M.W. Chen, C.T. Liu, Ductile CoCrFeNiMox high entropy alloys strengthened by hard intermetallic phases, Acta Mater. 116 (2016) 332–342. [8] 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, A precipitation-hardened high-entropy alloy with outstanding tensile properties, Acta Mater. 102 (2016) 187–196. [9] B. Gwalani, D. Choudhuri, V. Soni, Y. Ren, M. Styles, J.Y. Hwang, S.J. Nam, H. Ryu, S.H. Hong, R. Banerjee, Cu assisted stabilization and nucleation of L12 precipitates
(5)
where G is the shear modulus, and that can be calculated from G= E/2(1 + v) , which assumes the HEA is elastically isotropic. The value of Poisson's ratio (v) is 0.3 and the Young's modulus (E) can be measured by nanoindentation. b is the burgers vectors of the HEA and f is the volume fraction of precipitates. D is the spatial diameter of precipitates. According to Fig. 6, the SLMed FeCoCrNiC0.05 annealed for 0.5 h has the highest value of yield strength. When the annealing time elongated to 8 h, the yield strength decreased. The reason can be clarified in Fig. 8, which shows the calculated strength contribution from solidsolution hardening and precipitation hardening. At early annealing stage, a sharp increase occurs in the precipitation hardening, but the solid solution hardening decreases slightly. When increasing annealing time to 8 h, the precipitation hardening increases slightly, but the solid solution hardening decreases significantly. As a result, the yield strength decreases. 24
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