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Precipitation hardening in CoCrFeNi-based high entropy alloys
Materials Chemistry and Physics xxx (2017) 1e10
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Precipitation hardening in CoCrFeNi-based high entropy alloys W.H. Liu, T. Yang, C.T. Liu* Centre for Advanced Structural Materials, Department of Mechanical and Biomechanical Engineering, City University of Hong Kong, Hong Kong, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 3 May 2017 Accepted 10 July 2017 Available online xxx
High entropy alloys (HEAs) with a face-centered cubic (fcc) structure have received an increasing attention in recent years due to their exceptional ductility and fracture toughness. On the other hand, they generally show lower strengths at ambient and elevated temperatures. In this review, we attempt to highlight recent advances in the design of computation-aided HEAs being alloyed with Al, Al/Ti, Nb and Mo additions on the precipitation hardening behavior of this fcc-type CoCrFeNi alloy, and to summarize the corresponding mechanical properties of the newly developed precipitation-strengthened HEAs under different thermomechanical processing. In particular, our emphasis is played on elucidating the correlation between precipitation behaviors and mechanical properties, which may serve as a demonstrative summary for the current progress in the scientific research of precipitation-strengthened HEAs. Finally, future research areas for these precipitation-strengthened HEAs are critically assessed. © 2017 Elsevier B.V. All rights reserved.
Keywords: High entropy alloys Alloying effect Precipitation strengthening Intermetallic phases Mechanical properties
1. Introduction The continuous development of high-performance materials is critical to move toward a sustainable future, among which the design of metallic alloys has been playing a central role over thousands of years. The traditional strategy for alloy development was usually based on a single principal element with various additional elements to enhance particular properties, such as iron or aluminum-based alloys. After extensive efforts on these traditional alloys, their development is approaching to the limit. Recently, a new alloy design concept named as high entropy alloys (HEAs), which abandons the traditional “single-element” base idea and adopts mixed multiple elements in an equimolar or nearequimolar composition for alloying, has been proposed and inspired the exploration of the vast composition space [1,2]. Although HEAs are complex in alloy compositions, some of these alloys can crystallize as single-phase solid solutions with simple crystal structures, such as fcc or body-centered cubic (bcc), which attribute to their distinctive mechanical and physical properties [3e5]. As compared with conventional alloys, the fcc-type HEAs exhibit unique properties, such as the exceptional ductility and fracture toughness at cryogenic temperatures [6e10], which have generated
* Corresponding author. E-mail address: [email protected] (C.T. Liu).
significant interest among the material community. However, recent studies indicate that a HEA matrix alone, especially the single-phase fcc structure, is insufficiently strong for engineering applications at room and elevated temperatures. The extensively studied CoCrFeNi alloy with a simple fcc structure is a case of particular interest [11e13], and an analysis of the true stress (s) and strain (ε) curve of the as-cast CoCrFeNi alloy in Fig. 1 shows that it has a relatively low yield strength (YS) of 160.0 MPa but a quite high ultimate strength (UTS) of 718 MPa (¼ 4.5 YS) with an elongation to failure (EL) of ~50%, indicating an excellent uniform deformation and very high working hardening ability [14]. Such characteristic makes it an excellent base for further hard particle strengthening, the strategy of which has been proved to be one of the most effective approaches to enhance the strength of metallic materials, and the degree of hardening can be controlled by compositional optimization and/or thermomechanical processing. Therefore, significant research efforts have been dedicated to improve the strength of this fcc-type HEA at the present time. There have been significant advances in the precipitationstrengthened HEAs, and their strength is often associated with alloying elements with large differences in atomic radius and large negative enthalpies of mixing with the constituent elements of HEAs. In the current review, we aim at highlighting the recent development in the computation-aided HEA design of the fcc-type CoCrFeNi alloy hardened by precipitation of second-phase particles containing Al [15], Al/Ti [16,17], Nb [18e22] or Mo [14,23,24] additions. Emphasis will be place on the precipitation behavior of the
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Fig. 1. The engineering/true strain stress curves of the as cast CoCrFeNi alloy [14].
CoCrFeNi alloy and its correlation between their microstructures and mechanical properties under different thermal-mechanical processing. Noted that the thermodynamic and physicochemical characteristics of the alloying and constituent elements from Refs. [25e28] are listed in Table 1. 2. Computation-aided alloy design of precipitationstrengthened HEAs In general, the design of new alloys is based on a ‘trial and error’ approach, which requires a systematical production of large numbers of alloys with varied compositions and heat treatments. This approach is usually costly and time-intensive, requiring months to produce and characterize a large number of candidate alloys, especially for multicomponent HEAs. The integrated computational-prediction and experimental-validation approach, which has been already applied in traditional alloy designs, can dramatically accelerate the characterization of the compositioneprocessingestructureeproperty relationship of multicomponent alloys [29]. For the HEA design, the computational thermodynamic software, Thermo-Calc, and the TCHEA1 thermodynamic database for HEAs, developed using a CALPHAD approach, have been applied to understand HEAs' phase relationships. The success of this computation-aided alloy design approach has been demonstrated in several studies. He et al. [17] successfully predicted the phase diagram of the (CoCrFeNi)100-x-yTixAly (at.%) alloy system using TCHEA1, as presented in Fig. 2. Thermodynamic equilibrium calculations agree well with the experimental
Table 1 Thermodynamic and physiochemical properties of the alloying and constituent elements of HEAs [25e28], valence electron concentration (VEC), Pauling electronegativity (PE). Element
Atom size,nm
AB ; (kJ/mol) DHMix
Co Cr Fe Ni Nb Al Ti Mo
0.1251 0.1249 0.1241 0.1246 0.1429 0.1432 0.1462 0.1363
Co-Al Cr-Al Fe-Al Ni-Al Co-Ti Cr-Ti Fe-Ti Ni-Ti
19 10 11 22 28 7 17 35
Co-Nb Cr-Nb Fe-Nb Ni-Nb Co-Mo Cr-Mo Fe-Mo Ni-Mo
25 7 16 30 5 0 2 7
PE
VEC
1.88 1.66 1.83 1.91 1.6 1.61 1.54 2.16
9 6 8 10 5 3 4 6
observations; two types of precipitates are formed: nanosized coherent L12-Ni3(Ti,Al) g0 precipitates and the L21-(Ni,Co)2TiAl Heusler phase. The amount of L12-Ni3(Ti,Al)-type precipitates increases with the Ti and Al contents, but decreases with increased temperatures. For the Heusler phase, the formation temperature gradually expands, whilst the equilibrium mole fraction also increases with the Ti and Al contents. Furthermore, considering that the CoCrFeNi alloy possesses a single fcc disorder solid solution without compositional fluctuation or long-range chemical ordering, He et al. have made a further development (using the simple Thermo-Calc) to show that the CoCrFeNi alloy could serve as an integrated pseudo-element for alloy design [22]. Applying this modified Thermo-Calc program, a series of CoCrFeNiNbx eutectic HEAs were developed [22]. In general, the pseudo binary phase diagram presented in Fig. 3a agrees well with the experimental determined results. The CoCrFeNiNb0.1 alloy consists of a typical divorced eutectic structure, the CoCrFeNiNb0.25 and CoCrFeNiNb0.5 alloys show a hypo-eutectic structure, and the CoCrFeNiNb0.8 alloy has a typical hyper-eutectic structure. Liu et al. [14] used the pseudo binary phase diagram to develop the s and m particles strengthened CoCrFeNiMo0.3 alloy, as presented in Fig. 3b. Noted that the experimental phase relationship of the CoCrFeNiMox alloy system agrees well with the calculated resulted shown in Fig. 3b. Tong [ref as unpublished data] successfully predicted the formation of the intermetallic h phase in the CoCrFeNiTix alloy system using the CoCrFeNiTix pseudo binary phase diagram, as presented in Fig. 3c. All of these suggest that the pseudo binary phase diagrams actually provide a useful guideline for the prediction of precipitation phases in the CoCrFeNi HEA. On the other hand, the first-principles calculation based on the density functional theory (DFT) is able to calculate the ground state electronic structure and total energy of the HEA system. For example, Yu et al. [30] applied the simplified atomic models to investigate the physical mechanism of the nano-scale phase separation and precipitation behavior in the AlxCoCrCuFeNi HEAs [31]. The simulation results suggest the inevitability of the nano-scale phase separation in the AlxCoCrCuFeNi HEAs. It shows the fcc Curich domain surrounded by the L12 (Fe,Co,Ni,Cr)-rich phase coexisted with the disorder fcc solid solution phase, as shown in Fig. 4. Furthermore, the calculations agree well with the observations by state-of-the-art scanning transmission electron microscopy (STEM) [30]. 3. Precipitation-strengthened fcc-type HEAs It is well known that precipitation strengthening is critically dependent on the precipitate type, size, number density and distribution. As a result, understanding the alloying effect on the precipitation behavior in the CoCrFeNi matrix is the key to further control its microstructures and mechanical properties. It is noteworthy that the extensive deformation ability of the CoCrFeNi HEA [14] (see Fig. 1) makes it to be an excellent base for various intermetallic particles strengthening. In this review, the strengthening effect of g0 , Laves, s and m intermediate phases on the fcc-type CoCrFeNi alloy is carefully assessed. 3.1. Nanostructure-strengthened Al modified HEAs The Al content has significant impacts on the microstructural and mechanical properties of both as-cast and as-annealed AlxCoCrFeNi HEA system. Intensive studies show that the addition of Al in AlxCoCrFeNi initially promotes the formation of bcc phases for hardening in the as cast compositions, and the structure further changed from the initial single fcc phase (x < 0.5) to a duplex fcc plus bcc phases (0.5 x < 0.9) and then simple bcc phases (x 0.9)
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Fig. 2. The equilibrium phase mole fraction as a function of temperature calculated using the thermal-calc method; the curves of g0 and Heusler phase for the four TAxy HEAs are separated in (a) and (b) in order to show the alloying effect of Ti and Al. Noted that TAxy stands for TixAly [17].
as the continuous increase of Al content [15]. Furthermore, all the bcc phases in the as-cast alloys have a nano-scale two-phase structure (a mixed microstructure of both disordered bcc and B2 structures) formed by the spinodal decomposition mechanism. With the even further increased Al content, the as cast AlxCoCrFeNi alloys become much stronger, resulting from the formation of bcc type structures [32]. All of these phase changes lead to the hardness values varying from 100 to 740 HV [33,34]; however, the strength enhancement is at the cost of reduced ductility. At high Al contents, the B2 phase dominates and the alloys become brittle. Thus, an optimal Al ratio should be carefully controlled for such Al-modified HEA system in order to achieve optimized mechanical properties. An optimal heat-treatment process of the low Al content alloys is a promising method to improve mechanical behavior. Recent studies by various researchers clearly indicate that the low-Al AlxCoCrFeNi display a significant age-hardening phenomenon. For example, an aging of the Al0.3CrFeCoNi alloy at 550e600 C leads to the precipitation of highly refined ordered L12 Ni3(Ti,Al)-type intermetallic precipitates (~5 nm) as shown in Fig. 5a, and the precipitates are coherent with the fcc-matrix [35]. However, these g0 precipitates are only stable up to temperatures of 550e600 C, and they dissolve at temperatures above ~700 C, and are replaced
by coarser ordered B2 precipitates (~50 nm), as illustrated in Fig. 5b. The increase in YS of the Al0.3CoFeCrNi alloy after the precipitation of the ordered phases is evident [36], as shown in Fig. 5d and Table 2. The improvement in YS due to the second phase precipitation within the fcc matrix was Ds ~126 MPa for the L12 precipitates, and Ds ~56 MPa for the B2 precipitates. A co-precipitation strengthening, which is more attractive than single-phase precipitation of nanoparticles, can be achieved via a combination precipitation of both coherent fine precipitates of L12 within the fcc matrix and semi-coherent or incoherent grain boundary or near grain boundary precipitates of B2 after 620 C aging for 50 h, as shown in Fig. 5c. A significant improvement in YS (~490 MPa) was achieved, while maintaining a high ductility (~45%) in Fig. 5e. Furthermore, Niu et al. [37] shows that an aging at 650 C of the higher Al content Al0.5CoCrFeNi HEA has more significant hardening effect than that of the Al0.3CoCrFeNi, exhibiting a YS of ~834 MPa with an elongation of 25%. The excellent tensile properties were attributed to the nano-sized B2 phase in the interdendritic region and precipitates in the dendritic region. Noted that the YS of the Al0.5CoCrFeNi HEA strongly depends on aging time, varying between 355 and 834 MPa with different aging times. As discussed in the Al0.3CoCrFeNi alloy, L12 is more effective
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nano-scaled precipitates are identified to be the L12 Ni3(Ti,Al)-type g0 phase. An extraordinary balanced tensile properties at room temperature were achieved of a 645 MPa YS with an outstanding 39% elongation, as shown in Fig. 6a. A lower temperature aging at 650 C for 4 h leads to more uniform precipitates throughout the entire matrix, with a diameter ranging from 20 to 100 nm, this resulting in a more promising mechanical properties of YS over 1 GPa and a reliable 17% tensile ductility. Such a remarkable combination of YS and ductility is noteworthy for an fcc-based HEA. In summary, it is interesting to point out that the alloying of Al into transition-metal HEAs, such as CoCrFeNiMn and CoCrCuFeNi, usually changes the primary solidification phase from fcc to fcc þ bcc to bcc-based [38e40]. Such transition is usually attributed to the larger atomic size of Al atoms in contrast with those of Co, Fe, Mn, Cu, Cr and Ni atoms. Consequently, the large size misfit destabilizes the fcc lattice and introduces a bcc structure [32,38e40]. Furthermore, the electronic configuration of Al ([Ne]3s23p1) make it to have the characteristic of both the metal and nonmetal [32,41]. Consequently, Al has a strong tendency to form covalently bonded intermetallic compounds with the partially filled d-shell transitional elements like Ni ([Ar]4s23 d8), Co([Ar]4s23 d7) and Ti([Ar] 4s23 d2) via the sp-d hybridization [32]. These compounds have bcc-ordered structures like B2 [32,38e40] and possibly bcc-ordered L21 [42]. On the other hand, an optimal Al ratio is required to be carefully studied for each Al-modified HEA system to achieve the optimized mechanical properties. Meanwhile, the thermomechanical processing, which could result in different microstructures and mechanical properties than as-cast alloys, should also be optimized to obtain desired mechanical properties. 3.2. Eutectic HEA systems strengthened by Nb additions
Fig. 3. The pseudo binary phase diagram of (a) CoCrFeNiNbx, (b) CoCrFeNiMox and (c) CoCrFeNiTix alloys [14,22].
than B2 on strength enhancement [3635]; however it is only stable up to temperatures of 550e600 C [3535].He et al. [16] expanded the g þ g0 prime (L12) microstructure to a higher temperature of 800 C by alloying a minor Ti into the (FeCoNiCr)94Ti2Al4 (at.%) HEA.An aging at 800 C for 18 h resulted in a mixture of two precipitate morphologies: Region I consists of nano-precipitates less than 40 nm in size, while Region II near grain boundaries consists of particles coarser than 100 nm, as shown in Fig. 6b. All of these
An alloy with lamellar-type eutectic structures containing two phases as an in-situ composite could be promising for the balanced strength and ductility. Thus, the exploitation of the eutectic transformation has been considered in HEAs in order to achieve the balance of adequate fracture strength and ductility. Furthermore, besides the characteristic of sluggish diffusion kinetics and high softening resistance at elevated temperatures, eutectic HEAs with the composite structure are regarded as highly promising candidates for elevated-temperature structural materials, because the eutectic solidification structure has the advantage of nearequilibrium microstructures which resist the structural change at high temperatures. Guo et al. had successfully designed a highly uniform AlCoCrFeNi2.1 eutectic HEA, with an unprecedented combination of cast ability, high tensile ductility and high fracture strength at temperatures at least to 700 C [43]. Liu et al. [18] and He et al. [22] developed a eutectic Nb-modified HEA system with a promising mechanical performance, with the aid of the pseudo binary (CoCrFeNi)-Nb diagram (see Fig. 3a). The additions of Nb into the CoCrFeNi alloy induce a severe lattice distortion, due to the larger atom size of Nb compared to other principle elements, and the formation of a fine laminar eutectic structure composing of a ductile fcc phase and a hard Nb-enriched Laves phase, as shown in Fig. 7a [18]. Note that the Laves phase is a well-known size compound. Meanwhile, with the increased Nb content, the volume fraction of the Laves phase increases accordingly (Fig. 7b), and the microstructures change from an initial single-phase fcc structure (x ¼ 0) to a hypoeutectic microstructure with a primary fcc phase (x ¼ 0.1, 0.25), then to a full eutectic microstructure (x ¼ 0.45) and finally to a hypereutectic microstructure with a primary Laves phase (0.45 x 1.2) [19,22]. Furthermore, with the increase of Nb content, the increased hard/brittle Laves phase leads to a decrease of the plasticity and an increases of the Vickers hardness [19]. Wherein a good balance
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Fig. 4. Atomic models mimicking the nano-scale phase separations of the Al0.45CoCrCuFeNi alloy system [30].
Fig. 5. The microstructures and tensile properties of the Al0.3CoCrFeNi alloy under different aging temperature [36].
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Table 2 Composition and principle strengthening phases of representative precipitation-strengthened HEAs [14,16,38,39]. Composition
Aging condition
Strengthening phase
YS, MPa
UTS, MPa
EL,%
CoCrFeNi Al0.3CoCrFeNi
As-cast 1115 C 700 C/50 h 550 C/150 h 620 C/50 h As cast 650 C/8 h 800 C/18 h As-cast As-cast As-cast 850 C/1 h 900 C/5 h 900 C/5 h þ 700 C/5 h
nil nil B2 L12 B2 þ L12 bcc bcc þ B2 L12 nil nil
~155 ~159 ~215 ~285 ~490 ~355 ~834 ~645 ~198.8 ~254.7 ~305.3 ~815.5 ~646.7 ~683.7
~472 ~410 ~520 ~540 ~825 ~714 ~1220 1094 ~472.4 ~589.6 ~709.7 ~1186.5 ~1042.0 ~1066.6
~59 ~60 ~38 ~50 ~48 ~42 ~25% ~39 ~51 ~55 ~49 ~19 ~33 ~30
Al0.5CoCrFeNi (FeCoNiCr)94Ti2Al4 CoCrFeNiMo0.1 CoCrFeNiMo0.2 CoCrFeNiMo0.3
s sþm sþm sþm
Fig. 6. (a) Room temperature tensile properties of alloys CoCrFeNi (A), (CoCrFeNi)94Ti2Al4 (B), (CoCrFeNi)94Ti2Al4 30%-rolled, subsequent annealing at 1273 K for 2 h and aging at 1073 K for 18 h (P1), and (CoCrFeNi)94Ti2Al4 70%-rolled and subsequent aging at 923 K for 4 h (P2), and microstructures of the alloys (b) P1 and (c) P2 [16].
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Fig. 7. (a) The as-cast microstructure of the CoCrFeNiNb0.309 alloy, (b) the volume fraction of the Nb-enriched laves phase and (c) the tensile strain-stress curves of the as cast CoCrFeNiNbx (x ¼ 0, 0.155, 0.206 and 0.309) [18].
between strength and ductility is achieved in the as cast CoCrFeNiNb0.155 alloy (Fig. 7a) containing a 9.3% Nb-enriched Laves phase, which shows the excellent integrated mechanical properties with the tensile YS and plastic strain as high as 321 MPa and 21.3%, respectively [18]. The balanced tensile properties of the CoCrFeNiNb0.155 alloy result from a combination strengthening of the large Nb atom solid-solution and the precipitation of the Nb-induced Laves phase. However, the continuously increased Nb contents leaded to the formation of the interconnected brittle Laves-phase patches, which resulted in a sharp drop of the ductility in the CoCrFeNiNb0.206 alloy, as shown in Fig. 7c. Because cracks, formed in the ordered Nb-enriched Laves-phase, would propagate immediately along the continuous brittle phase network and result in brittle failure. He et al. [20] further investigated the effect of heat treatment on the microstructural and mechanical properties of the CoCrFeNiNb0.25 HEA, and they found that an aging at 750 C of this HEA promoted a precipitation of a lath-shaped fcc precipitate, which randomly distributed in the initial fcc phase. Fortunately, the fcc precipitate benefited the strength of the CoCrFeNiNb0.25 alloy extensively without sacrificing the compressive ductility. The
compressive YS of the CoCrFeNiNb0.25 alloy aged at 750 C was doubled as compared to that of the as-cast alloy, with the ductility almost unchanged. This is because the interface between the lathshaped fcc phase and the initial fcc phase could effectively inhibit the movement of dislocation, resulting in the improvement of strength without scarifying ductility. The eutectic HEAs are expected to be promising candidates for high temperature materials. He et al. [21] further examined the stability of the eutectic structures of the Nb-modified eutectic HEAs at elevated temperatures, and they reported that the lamellar structures of the eutectic CoCrFeNiNbx (x ¼ 0.5, 0.65 and 0.8) HEAs were stable below 750 C but they coarsened rapidly at temperatures higher than 750 C. The eutectic phases in this alloy system are always stable between 600 and 900 C but starts to spheroize when the temperature goes up to 900 C. It is noteworthy that, different from the conventional alloys, CoCrFeNiNbx eutectic HEAs still maintain good comprehensive mechanical properties even the lamellar structures are spheroized after annealing at 900 C. The good properties can be due to the mixture of the ductile fcc and hard Laves phases. In summary, Nb additions in the CoCrFeNi alloy have a strong
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tendency to induce the formation of eutectic structures, and this is closely related to the limited solubility of Nb in the CoCrFeNi matrix (Fig. 3a). The limited solubility of Nb is due to the fact that its atomic size is much larger than other elements in the CoCrFeNi base (see Table 1). Furthermore, the Fe-Nb, Cr-Nb, Ni-Nb and CoeNb alloys all can gain the eutectic structures around Nb-14% atomic percentage according to the corresponding binary phase diagrams. Consequently, Nb additions into the CoCrFeNi alloy tend to induce the formation of the eutectic structures.
3.3. Duplex particle strengthening of CoCrFeNi by Mo additions It was expected that the discrete, fine and hard topologically close-packed (TCP) phases would play a more effective hardening effect than the Nb-induced eutectic structure in the CoCrFeNi matrix. Liu et al. developed a CoCrFeNiMox alloy system strengthened by the fine and hard s and m phases, with the aid of the pseudo binary (CoCrFeNi)-Mo phase diagram [14]. It was reported that the additions of Mo in the CoCrFeNiMox (x ¼ 0.3 and 0.5) alloys induce the formation of a (Cr,Mo)-rich s phase in the fcc matrix [23]. Also, the (Mo,Cr)-rich m phase forms on the fringes of the s phase in the CoCrFeNiMo0.85 alloy [24]. They reported that the s phase is a (Cr,Mo)-rich phase corresponding to the stoichiometric (Cr,Mo) (Co,Fe,Ni), and the m phase has the stoichiometric (Mo,Cr)7(Co,Fe,Ni)6 [23,24]. This result is in consistent with the calculated pseudo binary (CoCrFeNi)-Mo phase diagram (see Fig. 3b). The solid solution of Mo in the CoCrFeNiMo0.3 alloy leads to a moderate YS improvement from 198.8 MPa to 305.3 MPa while the extremely good ductility was sustained [14]. In the as-cast state, the precipitates (a mixture of the s and fcc phases) formed upon solidification is too large (~2 mm) to provide an effective hardening. A heat treatment between 850 and 900 C leads to the precipitation of a substantial amount of fine s and m intermetallic phases with size of 50e100 nm (the inset of Fig. 8a). In agreement with the expectation, these s and m particles are much harder than that of the matrix, having a hardness at least higher than 8 GPa (see Fig. 8a), and it is interesting to know that these precipitates greatly harden the CoCrFeNiMo0.3 alloy without causing a serious embrittlement. For example, an aging at 1123 K for 1 h of the CoCrFeNiMo0.3 alloy leads to a significant improvement in UTS (1.2 GPa) with an attractive EL of ~19%, as shown in Fig. 8b. The exceptional mechanical behavior of the precipitationhardened CoCrFeNiMo0.3 alloy is associated with several beneficial factors. First of all, these particles are discretely distributed in the fcc matrix, thus the brittle failure associated with interconnected brittle phase was ruled out. Second, those discrete particles, which dissolved out during the annealing process at 850e900 C, are much fine and hard phase which could effectively harden the ductile fcc matrix [44,45]. Most important, the fcc matrix has an unusually large work hardening exponent of 0.75, as shown in Fig. 8c, which promotes an extensive deformation and suppresses the propagation of microcracks associated with these particles. All of these metallurgical factors contribute to a high fracture strength and decent ductility of the CoCrFeNiMo0.3 alloy hardened by both s and m particles. In summary, Mo additions in the CoCrFeNi matrix induced the formation of the discrete intermetallic phases. The precipitate morphology is closely related to the significant solubility of Mo (see Fig. 3b) [14], resulting from the moderately large atomic size as compared to the constitutions. For example, the CoCrFeNiMo0.3 alloy shows a single fcc phase at high temperatures, but on quenching and aging at intermediate temperatures, the intermetallic s and m phases can precipitate out as discrete particles.
Fig. 8. Nano-hardness of the s, m and fcc matrix of the CoCrFeNiMo0.3 alloy 60%-rolled and annealed at 1273 K for 1 d, (b) room-temperature tensile engineering strainestress curves for the as cast Mox (x ¼ 0, 0.2 and 0.3) alloys and the 60%-rolled CoCrFeNiMo0.3 alloy annealed at different conditions, (c) the variation of strain hardening exponent versus true strain of the as-cast Mo0 alloy, the inset of (a) is the microstructure of the CoCrFeNiMo0.3 alloy 60%-rolled and annealed at 1123 K for 1 h [14].
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4. Control of intermetallic compounds Based on the above analysis, alloy additions with Al, Al/Ti, Nb and Mo into the CoCrFeNi alloy induce gamma prime, Laves, sigma, Mu intermediate phases, and their formation is closely related with the solubility limit of the alloying elements. The solubility limit of the alloying elements depends critically on atom-size difference, valence electron and electronegativity with the constituent elements. When an alloying element is added to the CoCrFeNi matrix in such quantities that the limit of its solid solubility is exceeded, a secondary intermediate phase will start to form. This phase can be another solid solution, a chemical compound or a phase with a structure other than the one of the primary solid solution. First of all, the addition of Al to the transition-metal base of the CoCrFeNi alloy causes an inevitable change of its electronic structure and, thus, the lattice structures and properties upon excessive alloying. The element of Al with a valance electron concentration of 3 is effective in decreasing the electron concentration and stabilizing the bcc structure. This agrees with the conclusion of Guo et al. [46], which states that valence electron concentration plays a decisive role in the phase formation of as-cast HEAs. However, the alloying of the transition metals of Nb and Mo, which have a similar electronic configuration with the other constituent elements, would essentially induce the formation of intermediate phases without changing the lattice type. The precipitations and morphologies of these intermediate phases are closely related to the solubility of Nb and Mo in the CoCrFeNi matrix. The atomic size difference is of primary importance in determining the solubility of alloys, as shown in Fig. 3a and b, where the larger atom of Nb (as compared with Mo) would have a limited solubility in the CoCrFeNi matrix, only ~2% at 1600 K. The large lattice mismatch of rNb =r ðCo;Cr;Fe; NiÞ ¼ 1.145 between Nb and the constituent elements leads to a very limited solubility, resulting in the precipitation of the size-factor Laves phase and formation of the eutectic structure. It should be noted that the Nb-enriched Laves phase identified as the (Co,Cr,Fe,Ni)2(Cr,Nb) type with a hexagonal close-packed (hcp) lattice structure in the CoCrFeNiNbx alloy is a AB2 size factor compound. Laves phases are generally formed with a large sizedifference among alloying and constituent atoms, with the ratio of rA =rB approximately in the range of 1.1e1.6 [47]. The alloying of Mo in the CoCrFeNi alloy induced the precipitation of individual Mo-enriched s and m phases, which are essentially “electron compounds” rather than size compounds. This is because this phase formation is closely related to the electron concentration of the outer shell. The s phase forms between so called A and B elements, where A is usually in the group VB or VIB elements and B is usually in the VIIB or VIIIB elements [48]. The m phase starts to appear with increasing average electron concentration. Earlier observation shows that a larger difference in atom size is not favorable for the electron compound formation [47]. The limiting values of rA =rB ranges between 0.93 and 1.15, while the atomic ratio of rMo =r ðCo;Cr;Fe; NiÞ is 1.093, this benefiting the formation of the s and m phases. All of these evidences suggest that the type of precipitation employed to harden the CoCrFeNi base can be precisely controlled by the alloying element and composition. 5. Future work To further expedite the development of precipitationstrengthened HEAs and promote their wide engineering applications, future efforts should be focused on control of the precipitation of various intermetallic compounds, their tensile and creep properties at elevated-temperatures, all of which will lead to a guide to the broad application of HEAs. First of all, it is essentially important to optimize the mechanical
9
properties of the precipitation-strengthened HEAs through the control of the precipitation of various intermetallic compounds, because the strengthening efficiency highly depends on the type of precipitates. Furthermore, more work should be focused on creep resistance and oxidation properties at elevated temperatures, which will lead a guide to the broader application of HEAs. The effects of heat treatments and processing have to be optimized to develop HEAs with desirable microstructures and mechanical properties for high temperature applications. Actually, the structural feature of the (FeCoNiCr)94Ti2Al4 HEA bears a great resemblance to that of g0 -g Ni superalloys [16]. In contrast to the singlephase FeCoNiCrMn HEA, this precipitate-hardened (FeCoNiCr)94Ti2Al4 alloy exhibited a significant improvement in hightemperature mechanical properties [49]. Finally, it is our hope that this brief review would stimulate more interest in the development of precipitation-hardened HEAs in the future for hightemperature structural applications. 6. Conclusion This paper presents a comprehensive review of precipitation hardening in the fcc-type CoCrFeNi base HEA by alloying with additions of Al, Al/Ti, Nb and Mo elements, with emphasis on the correlation of their microstructures with the corresponding mechanical properties under different thermal-mechanical processes. The ultimate goal is to understand the compositioneprocessingemicrostructureeproperty relationship of these newly developed precipitation-strengthened HEAs. Computational techniques based on the Thermo-Calc and DFT first-principles calculations have provided a useful guidance for the design of the precipitates strengthened CoCrFeNi HEAs with additions of Al, Al/ Ti, Nb and Mo elements by introducing gamma prime, Laves, sigma, Mu intermediate phases. Our review indicates that the formation of coherent L12 nanostructured Ni3(Al)-type precipitates has the most promising strengthening effect on achieving a balance between ductility and strength. Among these precipitates, the Mo-induced discrete s and m phases are most effective in strengthening the ductile fcc CoCrFeNi matrix without scarifying too much ductility through proper thermal-mechanical processing. In this case, this fcc matrix promotes an extensive deformation and suppresses the microcrack propagation associated with these hard intermetallic particles. To further speed up the development of new precipitation-strengthened HEAs and promote their wide-range industrial applications, future work should be focused on the study of the control of precipitations of various intermetallic phases, creep resistance and oxidation properties at elevated temperatures. Acknowledge This research was supported by the Hong Kong Government through the General Research Fund (GRF) with the grant number of CityU 11209314. References [1] J.W. Yeh, et al., Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes, Adv. Eng. Mater. 6 (5) (2004) 299e303. [2] B. Cantor, I.T.H. Chang, P. Knight, A.J.B. Vincent, Microstructural development in equiatomic multicomponent alloys, Mater. Sci. Eng. A 213 (2004) 375e377. [3] Y.F. Ye, Q. Wang, J. Lu, C.T. Liu, Y. Yang, High-entropy alloy: challenges and prospects, Mater. Today 19 (2016) 349e362. [4] O.N. Senkov, G.B. Wilks, D.B. Miracle, C.P. Chuang, P.K. Liaw, Refractory highentropy alloys, Intermetallics 18 (2010) 1758e1765. [5] A. Takeuchi, K. Amiya, T. Wada, K. Yubuta, W. Zhang, High-entropy alloys with a hexagonal close-packed structure designed by equi-atomic alloy strategy
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Precipitation hardening in CoCrFeNi-based high entropy alloys W.H. Liu, T. Yang, C.T. Liu* Centre for Advanced Structural Materials, Department of Mechanical and Biomechanical Engineering, City University of Hong Kong, Hong Kong, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 3 May 2017 Accepted 10 July 2017 Available online xxx
High entropy alloys (HEAs) with a face-centered cubic (fcc) structure have received an increasing attention in recent years due to their exceptional ductility and fracture toughness. On the other hand, they generally show lower strengths at ambient and elevated temperatures. In this review, we attempt to highlight recent advances in the design of computation-aided HEAs being alloyed with Al, Al/Ti, Nb and Mo additions on the precipitation hardening behavior of this fcc-type CoCrFeNi alloy, and to summarize the corresponding mechanical properties of the newly developed precipitation-strengthened HEAs under different thermomechanical processing. In particular, our emphasis is played on elucidating the correlation between precipitation behaviors and mechanical properties, which may serve as a demonstrative summary for the current progress in the scientific research of precipitation-strengthened HEAs. Finally, future research areas for these precipitation-strengthened HEAs are critically assessed. © 2017 Elsevier B.V. All rights reserved.
Keywords: High entropy alloys Alloying effect Precipitation strengthening Intermetallic phases Mechanical properties
1. Introduction The continuous development of high-performance materials is critical to move toward a sustainable future, among which the design of metallic alloys has been playing a central role over thousands of years. The traditional strategy for alloy development was usually based on a single principal element with various additional elements to enhance particular properties, such as iron or aluminum-based alloys. After extensive efforts on these traditional alloys, their development is approaching to the limit. Recently, a new alloy design concept named as high entropy alloys (HEAs), which abandons the traditional “single-element” base idea and adopts mixed multiple elements in an equimolar or nearequimolar composition for alloying, has been proposed and inspired the exploration of the vast composition space [1,2]. Although HEAs are complex in alloy compositions, some of these alloys can crystallize as single-phase solid solutions with simple crystal structures, such as fcc or body-centered cubic (bcc), which attribute to their distinctive mechanical and physical properties [3e5]. As compared with conventional alloys, the fcc-type HEAs exhibit unique properties, such as the exceptional ductility and fracture toughness at cryogenic temperatures [6e10], which have generated
* Corresponding author. E-mail address: [email protected] (C.T. Liu).
significant interest among the material community. However, recent studies indicate that a HEA matrix alone, especially the single-phase fcc structure, is insufficiently strong for engineering applications at room and elevated temperatures. The extensively studied CoCrFeNi alloy with a simple fcc structure is a case of particular interest [11e13], and an analysis of the true stress (s) and strain (ε) curve of the as-cast CoCrFeNi alloy in Fig. 1 shows that it has a relatively low yield strength (YS) of 160.0 MPa but a quite high ultimate strength (UTS) of 718 MPa (¼ 4.5 YS) with an elongation to failure (EL) of ~50%, indicating an excellent uniform deformation and very high working hardening ability [14]. Such characteristic makes it an excellent base for further hard particle strengthening, the strategy of which has been proved to be one of the most effective approaches to enhance the strength of metallic materials, and the degree of hardening can be controlled by compositional optimization and/or thermomechanical processing. Therefore, significant research efforts have been dedicated to improve the strength of this fcc-type HEA at the present time. There have been significant advances in the precipitationstrengthened HEAs, and their strength is often associated with alloying elements with large differences in atomic radius and large negative enthalpies of mixing with the constituent elements of HEAs. In the current review, we aim at highlighting the recent development in the computation-aided HEA design of the fcc-type CoCrFeNi alloy hardened by precipitation of second-phase particles containing Al [15], Al/Ti [16,17], Nb [18e22] or Mo [14,23,24] additions. Emphasis will be place on the precipitation behavior of the
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Fig. 1. The engineering/true strain stress curves of the as cast CoCrFeNi alloy [14].
CoCrFeNi alloy and its correlation between their microstructures and mechanical properties under different thermal-mechanical processing. Noted that the thermodynamic and physicochemical characteristics of the alloying and constituent elements from Refs. [25e28] are listed in Table 1. 2. Computation-aided alloy design of precipitationstrengthened HEAs In general, the design of new alloys is based on a ‘trial and error’ approach, which requires a systematical production of large numbers of alloys with varied compositions and heat treatments. This approach is usually costly and time-intensive, requiring months to produce and characterize a large number of candidate alloys, especially for multicomponent HEAs. The integrated computational-prediction and experimental-validation approach, which has been already applied in traditional alloy designs, can dramatically accelerate the characterization of the compositioneprocessingestructureeproperty relationship of multicomponent alloys [29]. For the HEA design, the computational thermodynamic software, Thermo-Calc, and the TCHEA1 thermodynamic database for HEAs, developed using a CALPHAD approach, have been applied to understand HEAs' phase relationships. The success of this computation-aided alloy design approach has been demonstrated in several studies. He et al. [17] successfully predicted the phase diagram of the (CoCrFeNi)100-x-yTixAly (at.%) alloy system using TCHEA1, as presented in Fig. 2. Thermodynamic equilibrium calculations agree well with the experimental
Table 1 Thermodynamic and physiochemical properties of the alloying and constituent elements of HEAs [25e28], valence electron concentration (VEC), Pauling electronegativity (PE). Element
Atom size,nm
AB ; (kJ/mol) DHMix
Co Cr Fe Ni Nb Al Ti Mo
0.1251 0.1249 0.1241 0.1246 0.1429 0.1432 0.1462 0.1363
Co-Al Cr-Al Fe-Al Ni-Al Co-Ti Cr-Ti Fe-Ti Ni-Ti
19 10 11 22 28 7 17 35
Co-Nb Cr-Nb Fe-Nb Ni-Nb Co-Mo Cr-Mo Fe-Mo Ni-Mo
25 7 16 30 5 0 2 7
PE
VEC
1.88 1.66 1.83 1.91 1.6 1.61 1.54 2.16
9 6 8 10 5 3 4 6
observations; two types of precipitates are formed: nanosized coherent L12-Ni3(Ti,Al) g0 precipitates and the L21-(Ni,Co)2TiAl Heusler phase. The amount of L12-Ni3(Ti,Al)-type precipitates increases with the Ti and Al contents, but decreases with increased temperatures. For the Heusler phase, the formation temperature gradually expands, whilst the equilibrium mole fraction also increases with the Ti and Al contents. Furthermore, considering that the CoCrFeNi alloy possesses a single fcc disorder solid solution without compositional fluctuation or long-range chemical ordering, He et al. have made a further development (using the simple Thermo-Calc) to show that the CoCrFeNi alloy could serve as an integrated pseudo-element for alloy design [22]. Applying this modified Thermo-Calc program, a series of CoCrFeNiNbx eutectic HEAs were developed [22]. In general, the pseudo binary phase diagram presented in Fig. 3a agrees well with the experimental determined results. The CoCrFeNiNb0.1 alloy consists of a typical divorced eutectic structure, the CoCrFeNiNb0.25 and CoCrFeNiNb0.5 alloys show a hypo-eutectic structure, and the CoCrFeNiNb0.8 alloy has a typical hyper-eutectic structure. Liu et al. [14] used the pseudo binary phase diagram to develop the s and m particles strengthened CoCrFeNiMo0.3 alloy, as presented in Fig. 3b. Noted that the experimental phase relationship of the CoCrFeNiMox alloy system agrees well with the calculated resulted shown in Fig. 3b. Tong [ref as unpublished data] successfully predicted the formation of the intermetallic h phase in the CoCrFeNiTix alloy system using the CoCrFeNiTix pseudo binary phase diagram, as presented in Fig. 3c. All of these suggest that the pseudo binary phase diagrams actually provide a useful guideline for the prediction of precipitation phases in the CoCrFeNi HEA. On the other hand, the first-principles calculation based on the density functional theory (DFT) is able to calculate the ground state electronic structure and total energy of the HEA system. For example, Yu et al. [30] applied the simplified atomic models to investigate the physical mechanism of the nano-scale phase separation and precipitation behavior in the AlxCoCrCuFeNi HEAs [31]. The simulation results suggest the inevitability of the nano-scale phase separation in the AlxCoCrCuFeNi HEAs. It shows the fcc Curich domain surrounded by the L12 (Fe,Co,Ni,Cr)-rich phase coexisted with the disorder fcc solid solution phase, as shown in Fig. 4. Furthermore, the calculations agree well with the observations by state-of-the-art scanning transmission electron microscopy (STEM) [30]. 3. Precipitation-strengthened fcc-type HEAs It is well known that precipitation strengthening is critically dependent on the precipitate type, size, number density and distribution. As a result, understanding the alloying effect on the precipitation behavior in the CoCrFeNi matrix is the key to further control its microstructures and mechanical properties. It is noteworthy that the extensive deformation ability of the CoCrFeNi HEA [14] (see Fig. 1) makes it to be an excellent base for various intermetallic particles strengthening. In this review, the strengthening effect of g0 , Laves, s and m intermediate phases on the fcc-type CoCrFeNi alloy is carefully assessed. 3.1. Nanostructure-strengthened Al modified HEAs The Al content has significant impacts on the microstructural and mechanical properties of both as-cast and as-annealed AlxCoCrFeNi HEA system. Intensive studies show that the addition of Al in AlxCoCrFeNi initially promotes the formation of bcc phases for hardening in the as cast compositions, and the structure further changed from the initial single fcc phase (x < 0.5) to a duplex fcc plus bcc phases (0.5 x < 0.9) and then simple bcc phases (x 0.9)
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Fig. 2. The equilibrium phase mole fraction as a function of temperature calculated using the thermal-calc method; the curves of g0 and Heusler phase for the four TAxy HEAs are separated in (a) and (b) in order to show the alloying effect of Ti and Al. Noted that TAxy stands for TixAly [17].
as the continuous increase of Al content [15]. Furthermore, all the bcc phases in the as-cast alloys have a nano-scale two-phase structure (a mixed microstructure of both disordered bcc and B2 structures) formed by the spinodal decomposition mechanism. With the even further increased Al content, the as cast AlxCoCrFeNi alloys become much stronger, resulting from the formation of bcc type structures [32]. All of these phase changes lead to the hardness values varying from 100 to 740 HV [33,34]; however, the strength enhancement is at the cost of reduced ductility. At high Al contents, the B2 phase dominates and the alloys become brittle. Thus, an optimal Al ratio should be carefully controlled for such Al-modified HEA system in order to achieve optimized mechanical properties. An optimal heat-treatment process of the low Al content alloys is a promising method to improve mechanical behavior. Recent studies by various researchers clearly indicate that the low-Al AlxCoCrFeNi display a significant age-hardening phenomenon. For example, an aging of the Al0.3CrFeCoNi alloy at 550e600 C leads to the precipitation of highly refined ordered L12 Ni3(Ti,Al)-type intermetallic precipitates (~5 nm) as shown in Fig. 5a, and the precipitates are coherent with the fcc-matrix [35]. However, these g0 precipitates are only stable up to temperatures of 550e600 C, and they dissolve at temperatures above ~700 C, and are replaced
by coarser ordered B2 precipitates (~50 nm), as illustrated in Fig. 5b. The increase in YS of the Al0.3CoFeCrNi alloy after the precipitation of the ordered phases is evident [36], as shown in Fig. 5d and Table 2. The improvement in YS due to the second phase precipitation within the fcc matrix was Ds ~126 MPa for the L12 precipitates, and Ds ~56 MPa for the B2 precipitates. A co-precipitation strengthening, which is more attractive than single-phase precipitation of nanoparticles, can be achieved via a combination precipitation of both coherent fine precipitates of L12 within the fcc matrix and semi-coherent or incoherent grain boundary or near grain boundary precipitates of B2 after 620 C aging for 50 h, as shown in Fig. 5c. A significant improvement in YS (~490 MPa) was achieved, while maintaining a high ductility (~45%) in Fig. 5e. Furthermore, Niu et al. [37] shows that an aging at 650 C of the higher Al content Al0.5CoCrFeNi HEA has more significant hardening effect than that of the Al0.3CoCrFeNi, exhibiting a YS of ~834 MPa with an elongation of 25%. The excellent tensile properties were attributed to the nano-sized B2 phase in the interdendritic region and precipitates in the dendritic region. Noted that the YS of the Al0.5CoCrFeNi HEA strongly depends on aging time, varying between 355 and 834 MPa with different aging times. As discussed in the Al0.3CoCrFeNi alloy, L12 is more effective
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nano-scaled precipitates are identified to be the L12 Ni3(Ti,Al)-type g0 phase. An extraordinary balanced tensile properties at room temperature were achieved of a 645 MPa YS with an outstanding 39% elongation, as shown in Fig. 6a. A lower temperature aging at 650 C for 4 h leads to more uniform precipitates throughout the entire matrix, with a diameter ranging from 20 to 100 nm, this resulting in a more promising mechanical properties of YS over 1 GPa and a reliable 17% tensile ductility. Such a remarkable combination of YS and ductility is noteworthy for an fcc-based HEA. In summary, it is interesting to point out that the alloying of Al into transition-metal HEAs, such as CoCrFeNiMn and CoCrCuFeNi, usually changes the primary solidification phase from fcc to fcc þ bcc to bcc-based [38e40]. Such transition is usually attributed to the larger atomic size of Al atoms in contrast with those of Co, Fe, Mn, Cu, Cr and Ni atoms. Consequently, the large size misfit destabilizes the fcc lattice and introduces a bcc structure [32,38e40]. Furthermore, the electronic configuration of Al ([Ne]3s23p1) make it to have the characteristic of both the metal and nonmetal [32,41]. Consequently, Al has a strong tendency to form covalently bonded intermetallic compounds with the partially filled d-shell transitional elements like Ni ([Ar]4s23 d8), Co([Ar]4s23 d7) and Ti([Ar] 4s23 d2) via the sp-d hybridization [32]. These compounds have bcc-ordered structures like B2 [32,38e40] and possibly bcc-ordered L21 [42]. On the other hand, an optimal Al ratio is required to be carefully studied for each Al-modified HEA system to achieve the optimized mechanical properties. Meanwhile, the thermomechanical processing, which could result in different microstructures and mechanical properties than as-cast alloys, should also be optimized to obtain desired mechanical properties. 3.2. Eutectic HEA systems strengthened by Nb additions
Fig. 3. The pseudo binary phase diagram of (a) CoCrFeNiNbx, (b) CoCrFeNiMox and (c) CoCrFeNiTix alloys [14,22].
than B2 on strength enhancement [3635]; however it is only stable up to temperatures of 550e600 C [3535].He et al. [16] expanded the g þ g0 prime (L12) microstructure to a higher temperature of 800 C by alloying a minor Ti into the (FeCoNiCr)94Ti2Al4 (at.%) HEA.An aging at 800 C for 18 h resulted in a mixture of two precipitate morphologies: Region I consists of nano-precipitates less than 40 nm in size, while Region II near grain boundaries consists of particles coarser than 100 nm, as shown in Fig. 6b. All of these
An alloy with lamellar-type eutectic structures containing two phases as an in-situ composite could be promising for the balanced strength and ductility. Thus, the exploitation of the eutectic transformation has been considered in HEAs in order to achieve the balance of adequate fracture strength and ductility. Furthermore, besides the characteristic of sluggish diffusion kinetics and high softening resistance at elevated temperatures, eutectic HEAs with the composite structure are regarded as highly promising candidates for elevated-temperature structural materials, because the eutectic solidification structure has the advantage of nearequilibrium microstructures which resist the structural change at high temperatures. Guo et al. had successfully designed a highly uniform AlCoCrFeNi2.1 eutectic HEA, with an unprecedented combination of cast ability, high tensile ductility and high fracture strength at temperatures at least to 700 C [43]. Liu et al. [18] and He et al. [22] developed a eutectic Nb-modified HEA system with a promising mechanical performance, with the aid of the pseudo binary (CoCrFeNi)-Nb diagram (see Fig. 3a). The additions of Nb into the CoCrFeNi alloy induce a severe lattice distortion, due to the larger atom size of Nb compared to other principle elements, and the formation of a fine laminar eutectic structure composing of a ductile fcc phase and a hard Nb-enriched Laves phase, as shown in Fig. 7a [18]. Note that the Laves phase is a well-known size compound. Meanwhile, with the increased Nb content, the volume fraction of the Laves phase increases accordingly (Fig. 7b), and the microstructures change from an initial single-phase fcc structure (x ¼ 0) to a hypoeutectic microstructure with a primary fcc phase (x ¼ 0.1, 0.25), then to a full eutectic microstructure (x ¼ 0.45) and finally to a hypereutectic microstructure with a primary Laves phase (0.45 x 1.2) [19,22]. Furthermore, with the increase of Nb content, the increased hard/brittle Laves phase leads to a decrease of the plasticity and an increases of the Vickers hardness [19]. Wherein a good balance
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Fig. 4. Atomic models mimicking the nano-scale phase separations of the Al0.45CoCrCuFeNi alloy system [30].
Fig. 5. The microstructures and tensile properties of the Al0.3CoCrFeNi alloy under different aging temperature [36].
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Table 2 Composition and principle strengthening phases of representative precipitation-strengthened HEAs [14,16,38,39]. Composition
Aging condition
Strengthening phase
YS, MPa
UTS, MPa
EL,%
CoCrFeNi Al0.3CoCrFeNi
As-cast 1115 C 700 C/50 h 550 C/150 h 620 C/50 h As cast 650 C/8 h 800 C/18 h As-cast As-cast As-cast 850 C/1 h 900 C/5 h 900 C/5 h þ 700 C/5 h
nil nil B2 L12 B2 þ L12 bcc bcc þ B2 L12 nil nil
~155 ~159 ~215 ~285 ~490 ~355 ~834 ~645 ~198.8 ~254.7 ~305.3 ~815.5 ~646.7 ~683.7
~472 ~410 ~520 ~540 ~825 ~714 ~1220 1094 ~472.4 ~589.6 ~709.7 ~1186.5 ~1042.0 ~1066.6
~59 ~60 ~38 ~50 ~48 ~42 ~25% ~39 ~51 ~55 ~49 ~19 ~33 ~30
Al0.5CoCrFeNi (FeCoNiCr)94Ti2Al4 CoCrFeNiMo0.1 CoCrFeNiMo0.2 CoCrFeNiMo0.3
s sþm sþm sþm
Fig. 6. (a) Room temperature tensile properties of alloys CoCrFeNi (A), (CoCrFeNi)94Ti2Al4 (B), (CoCrFeNi)94Ti2Al4 30%-rolled, subsequent annealing at 1273 K for 2 h and aging at 1073 K for 18 h (P1), and (CoCrFeNi)94Ti2Al4 70%-rolled and subsequent aging at 923 K for 4 h (P2), and microstructures of the alloys (b) P1 and (c) P2 [16].
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Fig. 7. (a) The as-cast microstructure of the CoCrFeNiNb0.309 alloy, (b) the volume fraction of the Nb-enriched laves phase and (c) the tensile strain-stress curves of the as cast CoCrFeNiNbx (x ¼ 0, 0.155, 0.206 and 0.309) [18].
between strength and ductility is achieved in the as cast CoCrFeNiNb0.155 alloy (Fig. 7a) containing a 9.3% Nb-enriched Laves phase, which shows the excellent integrated mechanical properties with the tensile YS and plastic strain as high as 321 MPa and 21.3%, respectively [18]. The balanced tensile properties of the CoCrFeNiNb0.155 alloy result from a combination strengthening of the large Nb atom solid-solution and the precipitation of the Nb-induced Laves phase. However, the continuously increased Nb contents leaded to the formation of the interconnected brittle Laves-phase patches, which resulted in a sharp drop of the ductility in the CoCrFeNiNb0.206 alloy, as shown in Fig. 7c. Because cracks, formed in the ordered Nb-enriched Laves-phase, would propagate immediately along the continuous brittle phase network and result in brittle failure. He et al. [20] further investigated the effect of heat treatment on the microstructural and mechanical properties of the CoCrFeNiNb0.25 HEA, and they found that an aging at 750 C of this HEA promoted a precipitation of a lath-shaped fcc precipitate, which randomly distributed in the initial fcc phase. Fortunately, the fcc precipitate benefited the strength of the CoCrFeNiNb0.25 alloy extensively without sacrificing the compressive ductility. The
compressive YS of the CoCrFeNiNb0.25 alloy aged at 750 C was doubled as compared to that of the as-cast alloy, with the ductility almost unchanged. This is because the interface between the lathshaped fcc phase and the initial fcc phase could effectively inhibit the movement of dislocation, resulting in the improvement of strength without scarifying ductility. The eutectic HEAs are expected to be promising candidates for high temperature materials. He et al. [21] further examined the stability of the eutectic structures of the Nb-modified eutectic HEAs at elevated temperatures, and they reported that the lamellar structures of the eutectic CoCrFeNiNbx (x ¼ 0.5, 0.65 and 0.8) HEAs were stable below 750 C but they coarsened rapidly at temperatures higher than 750 C. The eutectic phases in this alloy system are always stable between 600 and 900 C but starts to spheroize when the temperature goes up to 900 C. It is noteworthy that, different from the conventional alloys, CoCrFeNiNbx eutectic HEAs still maintain good comprehensive mechanical properties even the lamellar structures are spheroized after annealing at 900 C. The good properties can be due to the mixture of the ductile fcc and hard Laves phases. In summary, Nb additions in the CoCrFeNi alloy have a strong
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tendency to induce the formation of eutectic structures, and this is closely related to the limited solubility of Nb in the CoCrFeNi matrix (Fig. 3a). The limited solubility of Nb is due to the fact that its atomic size is much larger than other elements in the CoCrFeNi base (see Table 1). Furthermore, the Fe-Nb, Cr-Nb, Ni-Nb and CoeNb alloys all can gain the eutectic structures around Nb-14% atomic percentage according to the corresponding binary phase diagrams. Consequently, Nb additions into the CoCrFeNi alloy tend to induce the formation of the eutectic structures.
3.3. Duplex particle strengthening of CoCrFeNi by Mo additions It was expected that the discrete, fine and hard topologically close-packed (TCP) phases would play a more effective hardening effect than the Nb-induced eutectic structure in the CoCrFeNi matrix. Liu et al. developed a CoCrFeNiMox alloy system strengthened by the fine and hard s and m phases, with the aid of the pseudo binary (CoCrFeNi)-Mo phase diagram [14]. It was reported that the additions of Mo in the CoCrFeNiMox (x ¼ 0.3 and 0.5) alloys induce the formation of a (Cr,Mo)-rich s phase in the fcc matrix [23]. Also, the (Mo,Cr)-rich m phase forms on the fringes of the s phase in the CoCrFeNiMo0.85 alloy [24]. They reported that the s phase is a (Cr,Mo)-rich phase corresponding to the stoichiometric (Cr,Mo) (Co,Fe,Ni), and the m phase has the stoichiometric (Mo,Cr)7(Co,Fe,Ni)6 [23,24]. This result is in consistent with the calculated pseudo binary (CoCrFeNi)-Mo phase diagram (see Fig. 3b). The solid solution of Mo in the CoCrFeNiMo0.3 alloy leads to a moderate YS improvement from 198.8 MPa to 305.3 MPa while the extremely good ductility was sustained [14]. In the as-cast state, the precipitates (a mixture of the s and fcc phases) formed upon solidification is too large (~2 mm) to provide an effective hardening. A heat treatment between 850 and 900 C leads to the precipitation of a substantial amount of fine s and m intermetallic phases with size of 50e100 nm (the inset of Fig. 8a). In agreement with the expectation, these s and m particles are much harder than that of the matrix, having a hardness at least higher than 8 GPa (see Fig. 8a), and it is interesting to know that these precipitates greatly harden the CoCrFeNiMo0.3 alloy without causing a serious embrittlement. For example, an aging at 1123 K for 1 h of the CoCrFeNiMo0.3 alloy leads to a significant improvement in UTS (1.2 GPa) with an attractive EL of ~19%, as shown in Fig. 8b. The exceptional mechanical behavior of the precipitationhardened CoCrFeNiMo0.3 alloy is associated with several beneficial factors. First of all, these particles are discretely distributed in the fcc matrix, thus the brittle failure associated with interconnected brittle phase was ruled out. Second, those discrete particles, which dissolved out during the annealing process at 850e900 C, are much fine and hard phase which could effectively harden the ductile fcc matrix [44,45]. Most important, the fcc matrix has an unusually large work hardening exponent of 0.75, as shown in Fig. 8c, which promotes an extensive deformation and suppresses the propagation of microcracks associated with these particles. All of these metallurgical factors contribute to a high fracture strength and decent ductility of the CoCrFeNiMo0.3 alloy hardened by both s and m particles. In summary, Mo additions in the CoCrFeNi matrix induced the formation of the discrete intermetallic phases. The precipitate morphology is closely related to the significant solubility of Mo (see Fig. 3b) [14], resulting from the moderately large atomic size as compared to the constitutions. For example, the CoCrFeNiMo0.3 alloy shows a single fcc phase at high temperatures, but on quenching and aging at intermediate temperatures, the intermetallic s and m phases can precipitate out as discrete particles.
Fig. 8. Nano-hardness of the s, m and fcc matrix of the CoCrFeNiMo0.3 alloy 60%-rolled and annealed at 1273 K for 1 d, (b) room-temperature tensile engineering strainestress curves for the as cast Mox (x ¼ 0, 0.2 and 0.3) alloys and the 60%-rolled CoCrFeNiMo0.3 alloy annealed at different conditions, (c) the variation of strain hardening exponent versus true strain of the as-cast Mo0 alloy, the inset of (a) is the microstructure of the CoCrFeNiMo0.3 alloy 60%-rolled and annealed at 1123 K for 1 h [14].
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4. Control of intermetallic compounds Based on the above analysis, alloy additions with Al, Al/Ti, Nb and Mo into the CoCrFeNi alloy induce gamma prime, Laves, sigma, Mu intermediate phases, and their formation is closely related with the solubility limit of the alloying elements. The solubility limit of the alloying elements depends critically on atom-size difference, valence electron and electronegativity with the constituent elements. When an alloying element is added to the CoCrFeNi matrix in such quantities that the limit of its solid solubility is exceeded, a secondary intermediate phase will start to form. This phase can be another solid solution, a chemical compound or a phase with a structure other than the one of the primary solid solution. First of all, the addition of Al to the transition-metal base of the CoCrFeNi alloy causes an inevitable change of its electronic structure and, thus, the lattice structures and properties upon excessive alloying. The element of Al with a valance electron concentration of 3 is effective in decreasing the electron concentration and stabilizing the bcc structure. This agrees with the conclusion of Guo et al. [46], which states that valence electron concentration plays a decisive role in the phase formation of as-cast HEAs. However, the alloying of the transition metals of Nb and Mo, which have a similar electronic configuration with the other constituent elements, would essentially induce the formation of intermediate phases without changing the lattice type. The precipitations and morphologies of these intermediate phases are closely related to the solubility of Nb and Mo in the CoCrFeNi matrix. The atomic size difference is of primary importance in determining the solubility of alloys, as shown in Fig. 3a and b, where the larger atom of Nb (as compared with Mo) would have a limited solubility in the CoCrFeNi matrix, only ~2% at 1600 K. The large lattice mismatch of rNb =r ðCo;Cr;Fe; NiÞ ¼ 1.145 between Nb and the constituent elements leads to a very limited solubility, resulting in the precipitation of the size-factor Laves phase and formation of the eutectic structure. It should be noted that the Nb-enriched Laves phase identified as the (Co,Cr,Fe,Ni)2(Cr,Nb) type with a hexagonal close-packed (hcp) lattice structure in the CoCrFeNiNbx alloy is a AB2 size factor compound. Laves phases are generally formed with a large sizedifference among alloying and constituent atoms, with the ratio of rA =rB approximately in the range of 1.1e1.6 [47]. The alloying of Mo in the CoCrFeNi alloy induced the precipitation of individual Mo-enriched s and m phases, which are essentially “electron compounds” rather than size compounds. This is because this phase formation is closely related to the electron concentration of the outer shell. The s phase forms between so called A and B elements, where A is usually in the group VB or VIB elements and B is usually in the VIIB or VIIIB elements [48]. The m phase starts to appear with increasing average electron concentration. Earlier observation shows that a larger difference in atom size is not favorable for the electron compound formation [47]. The limiting values of rA =rB ranges between 0.93 and 1.15, while the atomic ratio of rMo =r ðCo;Cr;Fe; NiÞ is 1.093, this benefiting the formation of the s and m phases. All of these evidences suggest that the type of precipitation employed to harden the CoCrFeNi base can be precisely controlled by the alloying element and composition. 5. Future work To further expedite the development of precipitationstrengthened HEAs and promote their wide engineering applications, future efforts should be focused on control of the precipitation of various intermetallic compounds, their tensile and creep properties at elevated-temperatures, all of which will lead to a guide to the broad application of HEAs. First of all, it is essentially important to optimize the mechanical
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properties of the precipitation-strengthened HEAs through the control of the precipitation of various intermetallic compounds, because the strengthening efficiency highly depends on the type of precipitates. Furthermore, more work should be focused on creep resistance and oxidation properties at elevated temperatures, which will lead a guide to the broader application of HEAs. The effects of heat treatments and processing have to be optimized to develop HEAs with desirable microstructures and mechanical properties for high temperature applications. Actually, the structural feature of the (FeCoNiCr)94Ti2Al4 HEA bears a great resemblance to that of g0 -g Ni superalloys [16]. In contrast to the singlephase FeCoNiCrMn HEA, this precipitate-hardened (FeCoNiCr)94Ti2Al4 alloy exhibited a significant improvement in hightemperature mechanical properties [49]. Finally, it is our hope that this brief review would stimulate more interest in the development of precipitation-hardened HEAs in the future for hightemperature structural applications. 6. Conclusion This paper presents a comprehensive review of precipitation hardening in the fcc-type CoCrFeNi base HEA by alloying with additions of Al, Al/Ti, Nb and Mo elements, with emphasis on the correlation of their microstructures with the corresponding mechanical properties under different thermal-mechanical processes. The ultimate goal is to understand the compositioneprocessingemicrostructureeproperty relationship of these newly developed precipitation-strengthened HEAs. Computational techniques based on the Thermo-Calc and DFT first-principles calculations have provided a useful guidance for the design of the precipitates strengthened CoCrFeNi HEAs with additions of Al, Al/ Ti, Nb and Mo elements by introducing gamma prime, Laves, sigma, Mu intermediate phases. Our review indicates that the formation of coherent L12 nanostructured Ni3(Al)-type precipitates has the most promising strengthening effect on achieving a balance between ductility and strength. Among these precipitates, the Mo-induced discrete s and m phases are most effective in strengthening the ductile fcc CoCrFeNi matrix without scarifying too much ductility through proper thermal-mechanical processing. In this case, this fcc matrix promotes an extensive deformation and suppresses the microcrack propagation associated with these hard intermetallic particles. To further speed up the development of new precipitation-strengthened HEAs and promote their wide-range industrial applications, future work should be focused on the study of the control of precipitations of various intermetallic phases, creep resistance and oxidation properties at elevated temperatures. Acknowledge This research was supported by the Hong Kong Government through the General Research Fund (GRF) with the grant number of CityU 11209314. References [1] J.W. Yeh, et al., Nanostructured high-entropy alloys with multiple principal elements: novel alloy design concepts and outcomes, Adv. Eng. Mater. 6 (5) (2004) 299e303. [2] B. Cantor, I.T.H. Chang, P. Knight, A.J.B. Vincent, Microstructural development in equiatomic multicomponent alloys, Mater. Sci. Eng. A 213 (2004) 375e377. [3] Y.F. Ye, Q. Wang, J. Lu, C.T. Liu, Y. Yang, High-entropy alloy: challenges and prospects, Mater. Today 19 (2016) 349e362. [4] O.N. Senkov, G.B. Wilks, D.B. Miracle, C.P. Chuang, P.K. Liaw, Refractory highentropy alloys, Intermetallics 18 (2010) 1758e1765. [5] A. Takeuchi, K. Amiya, T. Wada, K. Yubuta, W. Zhang, High-entropy alloys with a hexagonal close-packed structure designed by equi-atomic alloy strategy
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Please cite this article in press as: W.H. Liu, et al., Precipitation hardening in CoCrFeNi-based high entropy alloys, Materials Chemistry and Physics (2017), http://dx.doi.org/10.1016/j.matchemphys.2017.07.037