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High-Entropy Alloys: A Critical Review

High-entropy alloys (HEAs) are alloys with five or more principal elements. Due to the distinct design concept, these alloys often exhibit unusual properties. Thus, there has been significant interest in these materials, leading to an emerging yet exciting new field. This paper briefly reviews some critical aspects of HEAs, including core effects, phases and crystal structures, mechanical properties, high-temperature properties, structural stabilities, and corrosion behaviors. Current challenges and important future directions are also pointed out.


Introduction
Most conventional alloys are based on one principal element. Different kinds of alloying elements are added to the principal element to improve its properties, forming an alloy family based on the principal element. For example, steel is based on Fe, and aluminum alloys are based on Al. However, the number of elements in the periodic table is limited, thus the alloy families we can develop are also limited. If we think outside the conventional box and design alloys not from one or two 'base' elements, but from multiple elements altogether, what will we get? This new concept, first proposed in 1995, [1] has been named a high-entropy alloy (HEA). [2] HEAs are defined as alloys with five or more principal elements. Each principal element should have a concentration between 5 and 35 at.%. [2,3] Besides principal elements, HEAs can contain minor elements, each below 5 at.%. These alloys are named 'HEAs' because their liquid or random solid solution states have significantly higher mixing entropies than those in conventional alloys. Thus, the effect of entropy is much more pronounced in HEAs. Existing physical metallurgy knowledge and binary/ternary phase diagrams suggest that such multielement alloys may develop several dozen kinds of phases and intermetallic compounds, resulting in complex and brittle microstructures that are difficult to analyze and engineer, and probably have very limited practical value.
behaviors. It also points out current challenges and future directions. The physical (magnetic, electrical, and thermal) properties of HEAs have been reviewed elsewhere [4] and are not discussed in this paper.

Core Effects [3,5]
The multiprincipal-element character of HEAs leads to some important effects that are much less pronounced in conventional alloys. These can be considered as four 'core effects'. This section briefly introduces and discusses the core effects.

High-Entropy Effect.
The high-entropy effect states that the higher mixing entropy (mainly configurational) in HEAs lowers the free energy of solid solution phases and facilitates their formation, particularly at higher temperatures. Due to this enhanced mutual solubility among constituent elements, the number of phases present in HEAs can be evidently reduced. According to G = H − TS (where G is the Gibbs free energy, H is enthalpy, T is temperature, and S is entropy) entropy can stabilize a phase with higher entropy, provided that temperature is sufficiently high. For example, the melting phenomenon of a pure metal is due to the higher entropy of the liquid state relative to the solid state (according to Richard's rule, [6] at melting point the entropy difference between the two states roughly equals the gas constant R). Similarly, among the various types of phases in an alloy, those with higher entropy may also be stabilized at high temperature. In conventional alloys, solid solution phases (including terminal and intermediate solid solutions) have higher mixing entropy than intermetallic compounds do. For HEAs, the difference in entropy between solid solutions and compounds is particularly large owing to the multiprincipal-element design. For example, the configurational entropy of an equimolar quinary random solid solution is 1.61R. This means that the entropy difference between an equimolar quinary solution and a completely ordered phase (whose entropy is negligible) is about 60% larger than the entropy difference of the melting case mentioned above (i.e. 60% larger than the entropy difference between liquid and solid states of pure metals). Thus, it is highly probable that solid solution phases become the stable phase in HEAs at high temperature. It is expected that with the increase of temperature, the overall degree of order in HEAs decreases. Thus, even those alloys that contain ordered phases in their cast state can transform to random solid solutions at high temperature. However, if the formation enthalpy of an intermetallic compound is high enough to overcome the effect of entropy, that intermetallic compound will still be stable at high temperature.
There is much evidence for the high-entropy effect. [7][8][9][10][11][12][13] Here we demonstrate this effect with Figure 1, which shows the XRD patterns of a series of binary to septenary alloys. It is seen that the phases in quinary, senary, and septenary alloys remain rather simple: there are only two major phases, and these phases have simple structures such as BCC and FCC (note that the septenary alloy actually contain minor intermetallic phases, although not seen clearly in the XRD pattern). Such simple structures contradict conventional expectation: formation of various kinds of binary/ternary compounds.
Two myths regarding the high-entropy effect are discussed here. The first myth is that the high-entropy effect guarantees the formation of a simple solid solution phase (order/disordered BCC/FCC) at high temperatures. This is not true. Actually, in the very first paper of HEA, [2] the possibility of compound formation was already mentioned. In the same year, the existence of (Cr, Fe)-rich borides in the AlB x CoCrCuFeNi alloys was observed. [14] As mentioned previously, it is the competition between entropy and enthalpy that determines the phase formation. The large negative formation enthalpy of the borides thus justifies their stability at high temperatures. Similar phenomena have also been observed in other high-temperature annealing experiments. [15][16][17] The second myth is that only simple solid solution phases (BCC/FCC) are the desired phases for practical applications of HEAs. This is also not true. It is true that these multiprincipal-element simple solid solution phases are unique to HEAs, and they can have outstanding properties. However, as shown later, many HEAs containing intermediate phases also have great potential. Thus, the exploration of HEAs should not be limited to simple solid solution phases.

Sluggish Diffusion Effect.
It was proposed that the diffusion and phase transformation kinetics in HEAs are slower than those in their conventional counterparts. [18] This can be understood from two Figure 2. Schematic diagram of the potential energy change during the migration of a Ni atom. The mean difference (MD) in potential energy after each migration for pure metals is zero, whereas that for HEA is the largest. [19] aspects. Firstly, in HEA the neighboring atoms of each lattice site are somewhat different. Thus, the neighbors before and after an atom jump into a vacancy are different. The difference in local atomic configuration leads to different bonding and therefore different local energies for each site. When an atom jumps into a low-energy site, it becomes 'trapped' and the chance to jump out of that site will be lower. In contrast, if the site is a high-energy site, then the atom has a higher chance to hop back to its original site. Either of these scenarios slows down the diffusion process. Note that in conventional alloys with a low solute concentration, the local atomic configuration before and after jumping into a vacancy is, most of the time, identical. Tsai et al. used a seven-bond model to calculate the effect of local energy fluctuation on diffusion. [19] They showed that for Ni atoms diffusing in Co-Cr-Fe-Mn-Ni alloys (which has a single-phase, FCC structure), the mean potential energy difference between lattice sites is 60.3 meV, which is 50% higher than that in Fe-Cr-Ni alloys (see Figure 2). This energy difference leads to a significantly longer occupation time at low-energy sites (1.73 times longer than that at high-energy sites). Indeed, diffusion couple experiments show that the activation energy of diffusion in the Co-Cr-Fe-Mn-Ni alloys is higher than those in other FCC ternary alloys and pure elements (Table 1). [19] The diffusivity of Ni at the respective melting point of each metal is also calculated ( Table 1). The slowest diffusion takes place again in the Co-Cr-Fe-Mn-Ni alloys.
Secondly, the diffusion rate of each element in a HEA is different. Some elements are less active (for example, elements with high melting points) than others so these elements have lower success rates for jumping into vacancies when in competition with other elements. However, phase transformations typically require the coordinated diffusion of many kinds of elements. For example, the nucleation and growth of a new phase require the redistribution of all elements to reach the desired composition. Grain growth also requires the cooperation of all elements so that grain boundaries can successfully migrate. In these scenarios, the slow-moving elements become the rate-limiting factor that impedes the transformation.
The slow kinetics in HEAs allows readily attainable supersaturated state and nano-sized precipitates, even in the cast state. [2,20,21] It also contributes to the excellent performance of HEA coatings as diffusion barriers. [22,23] Moreover, it allows better high-temperature strength and structural stability, which are discussed in Section 5. For the same reason, HEAs are also expected to have outstanding creep resistance.

Severe-Lattice-Distortion Effect. The lattice in
HEAs is composed of many kinds of elements, each with different size. These size differences inevitably lead to distortion of the lattice. Larger atoms push away their neighbors and small ones have extra space around. The strain energy associated with lattice distortion raises the overall free energy of the HEA lattice. It also affects the properties of HEAs. For example, lattice distortion impedes dislocation movement and leads to pronounced

Cocktail Effect.
The properties of HEAs are certainly related to the properties of its composing elements.
For example, addition of light elements decreases the density of the alloy. However, besides the properties of the individual composing elements, the interaction among the component elements should also be considered. For example, Al is a soft and low-melting-point element. Addition of Al, however, can actually harden HEAs. Figure 3 plots the hardness of the Al x CoCrCuFeNi alloy as a function of Al content. It is clearly seen that the alloy hardens significantly with the addition of Al. This is partly because of the formation of a hard BCC phase, and partly due to the stronger cohesive bonding between Al and other elements, and its larger atomic size. Thus, the macroscopic properties of HEA not only come from the averaged properties of the component elements, but also include the effects from the excess quantities produced by inter-elemental reactions and lattice distortion. In the literature, phases in HEAs are usually classified differently. They are often classified as: random solid solution (e.g. FCC, BCC), ordered solid solution (e.g. B2 and L1 2 ), and intermetallic phases (e.g. Laves phases). This classification may lead to some confusion because by definition, intermetallic phases can also be classified as ordered solid solutions-they have composition ranges and are typically ordered. In view of this, we suggest classifying the phases according to their structure (simple/complex) and ordering (ordered/disordered). A phase is said to be simple if its structure is identical to or derived from FCC, BCC or HCP structures. Namely, cI2-W, cF4-Cu, and hP2-Mg structures and their ordered versions (superlattices) such as cP2-CsCl (B2) and cP4-AuCu3 (L1 2 ) structures are simple. [24] If a phase is not simple (e.g. Laves phases), it is said to be complex. Therefore, the above three types now becomes: simple disordered phase (SDP), simple ordered phase (SOP), and complex ordered phase (COP).

Simple Solid Solution or Intermetallics?. One fundamental and important question regarding HEAs is:
what kind of phase and crystal structure will form when we mix so many different elements together? To the surprise of most people, simple structures (SDPs and SOPs) are the most frequently seen in as-cast HEAs (see Figure 1). These simple phases originate from the high-entropy effect mentioned previously. Besides simple phases, different kinds of COPs, such as σ , μ, Laves, etc., are also observed in HEAs. Because simple-solution phases with more than five elements are unique to HEAs, researchers have worked intensely to understand the conditions for their formation.
From the classic Hume-Rothery rules, factors that affect the formation of binary solid solutions include atomic size difference, electron concentration, and difference in electronegativity. [25] Besides these factors, enthalpy and entropy of mixing are the most important phase formation parameters for HEAs. Zhang et al. [26] and Guo et al. [27] studied the effect of these parameters on the phase formation of HEAs and obtained similar conclusions: the formation of simple or complex phases depends mainly on the enthalpy of mixing ( H mix ), entropy of mixing ( S mix ), and atomic size differences (δ). In order to form sole simple phases (i.e. FCC, BCC, and their mixtures, including both ordered/disordered cases), the following conditions have to be met simultaneously: −22 ≤ H mix ≤ 7 kJ/mol, δ ≤ 8.5, and 11 ≤ S mix ≤ 19.5 J/(K mol). [27] This is shown in Figure 4, where the boundary of simple phases is shown. These conditions are quite logical: H mix cannot be too large in value because large positive H mix leads to phase separation and large negative H mix typically leads to Figure 4. Superimposed effect of H mix , δ, and S mix on phase stability in equiatomic multicomponent alloys and BMGs. Notes: Blue symbols indicate the relationship between H mix and δ, while red ones represents that between S mix and δ. Symbol • represents equiatomic amorphous phase-forming alloys; • represents nonequiatomic amorphous phase-forming alloys; represents simple phases and represents intermetallic phases. The region delineated by the dash-dotted lines indicates the requirements for simple phases to form. [27] intermetallic phases. δ has to be small enough since large δ leads to excess strain energy and destabilizes simple structures. S mix has to be large enough because it is the main stabilizing factor for simple phases. If the target of discussion is limited to SDPs only, the conditions are more strict: −15 ≤ H mix ≤ 5 kJ/mol, δ ≤ 4.3, and 12 ≤ S mix ≤ 17.5 J/(K mol). [26] The other criterion to judge whether or not sole simple phases will form in an HEA is the parameter ε = |T S mix / H mix |. [28,29] Based on the concept of entropy-enthalpy competition (see Section 2.1), ε represents the effect of entropy relative to that of enthalpy. Thus, larger ε suggests a higher possibility of forming SDP, which agrees well with the analysis. [28,29] A drawback of the two criterions shown above is that even if an alloy is designed following these criterions, it can still contain intermetallic phases. This can be seen in Figure 4. The 'solid solution' region delineated by the dash-dotted lines still contains triangles, indicating the existence of intermetallic phases. A new thermodynamic parameter has been proposed to solve this problem, and the results will be published soon. [30] Some points of notice should be mentioned here: Firstly, the above analyses are based on the as-cast state. This is natural because most reported HEAs are in their as-cast state. Secondly, the phases observed macroscopically in as-cast HEAs can contain atomic-scale decompositions/inhomogeneities. This is evidenced by detailed transmission electron microscopy (TEM) and atomic probe analysis. [21,31,32] Thirdly, in HEAs similar disordered and ordered versions of the same base structure often co-exist. [8,[33][34][35] For example, coexistence of BCC and ordered BCC (B2) phases was frequently reported. [8,21,34] Coexistence of FCC and ordered FCC (L1 2 ) phases was also observed. [35] When these phases have nearly identical lattice parameter, their XRD peaks overlap (except the superlattice peaks). Therefore, when the amount of the SOP is small, XRD may reflect only the peaks belonging to SDP, even if SOP exists. [34,35] This indicates that TEM analysis is needed to really confirm the existence of SOP. [34,35]

Crystal Structure of Simple Solid Solution
Phases. If an HEA does crystallize into simple phases, what will the structure of the phase be? Virtually all simple solid solution phase observed in HEAs have either BCC or FCC structures. [8,10,36,37] The most critical factor that decides whether an alloy crystallizes into BCC or FCC structure appears to be its VEC (valence electron concentration, the VEC of an alloy is calculated from the Figure 5. Relationship between VEC and the FCC, BCC phase stability for various HEA systems. Notes: Fully closed symbols for sole FCC phases; fully open symbols for sole BCC phase; top-half closed symbols for mixed FCC and BCC phases. [38] weighted average VEC of the constituent components). Guo et al. summarized the relationship between VEC and structure of many HEAs ( Figure 5). [38] They found that when the VEC of the alloy is larger than 8, the FCC structure is stabilized. When the VEC of the alloy is smaller than 6.87, the BCC structure is stabilized. Coexistence of FCC and BCC phase is observed at VEC values between 6.87 and 8. [38] However, the mechanism behind the effect of VEC on phase formation has not been fully understood so far.

High-Entropy Bulk Metallic Glass. Although
HEAs and bulk metallic glasses (BMGs) are both multicomponent and their compositions look similar at first glance, these two classes of materials are based on completely different concepts. For HEAs, more than five principal elements are required. In contrast, most BMGs are based on one or two principal elements. More importantly, BMGs are characterized by their amorphous structure, but there is no structural requirement for HEAs. The concept of HEAs, however, triggered an idea: would there be good glass formers in HEAs? If so, how do we design high-entropy BMGs (HEB-MGs) with good glass-forming ability (GFA) and better properties?
Inoue's empirical rules [39] state that (1) more than three composing element, (2) large atomic size differences, and (3) large negative mixing enthalpies favor the formation of BMGs. The first rule is in line with the design concept of HEAs. Indeed, high configurational entropy in HEAs was found to be beneficial for glass formation. This was evidenced by lower critical cooling rates in BMGs with higher S mix . [27] Another earlier study also reported similar findings. [40] The second and third rules for BMG formation are exactly opposite to the requirements to form simple phases in HEAs (see Section 3.1). This suggests that one should choose HEAs that are predicted to form intermetallic phases rather than simple phases as candidates for HEBMGs. Indeed, very recently Guo et al. analyzed the H mix and δ of some HEAs and BMGs and found that simple-solution-type HEAs and BMGs fall in opposite corners of the δ-H mix plot (shown in Figure 6 [41]). They further pointed out that as long as the competition from intermetallic phases can be ruled out by composition adjustment, HEAs that locate in the amorphous phase region of the δ-H mix plot indeed solidify into metallic glasses. [41] This strategy has been successfully demonstrated in a series of Zr-containing alloys prepared via melt spinning. [41] Not many HEBMGs have been fabricated till now. Most of these HEBMGs have diameters smaller than 3 mm. [42][43][44][45][46] Takeuchi et al. successfully prepared Pd 20 Pt 20 Cu 20 Ni 20 P 20 HEBMG with a maximum diameter of 10 mm. [47] But they suggested that conventional GFA assessment parameters such as T rg (reduced glass Figure 6. The δ-H mix plot delineating the phase selection in HEAs. Note: The dash-dotted regions highlight the individual region to form simple solid solutions, intermetallic phases and amorphous phases. [41] transition temperature normalized with liquidus temperature) cannot evaluate the GFA of HEBMG effectively. Other factors, e.g. Gibbs free energy assessments, need to be taken into account. [47] Apparently more study is needed to understand these issues.
The distinct composition design in HEAs brings with their amorphous structure other special properties. For example, an extremely high crystallization temperature (probably the highest among reported BMG) of over 800 • C and good performance as diffusion barrier between Cu and Si was observed in 20-nm-thick NbSiTaTiZr alloy film. [23] Additionally, Zn 20 Ca 20 Sr 20 Yb 20 (Li 0.55 Mg 0.45 ) 20 shows room temperature homogeneous flow and a plasticity of 25%, [44] which is in sharp contrast to the typical shear banding and brittle behavior of most BMGs. Very recently CaMgZn-SrYb, a high-entropy modification of Ca 65 Mg 15 Zn 20 BMG, has shown improved properties in orthopedic applications. [48] CaMgZnSrYb not only has better mechanical properties, but also promotes osteogenesis and new bone formation after two weeks of implantation. These examples indicate that high-entropy BMGs and amorphous films do have great potential and deserve further exploration-which is good news for both communities.

Phases at Elevated Temperatures.
As mentioned in Section 3.1, our understanding of the phases in HEAs focuses mainly on their as-cast state. However, the cast state is typically not in thermodynamic equilibrium. For alloys that locate in the BCC or FCC phase region suggested by Guo et al. (see Figure 5 [38]), phase transformation may be less significant. For example, no phase formation was found in the Co-Cr-Fe-Mn-Ni FCC alloys during their high-temperature annealing. [14,17] For alloys that locate in the BCC+FCC phase region, Figure 7. Relationship between the VEC and the presence of σ phase after aging for a number of HEAs. Green and red icons indicate the absence and presence of σ phase after aging, respectively. [52] however, phase transformation will take place at higher temperatures. In the Al-Co-Cr-Cu-Fe-Ni system, annealing at ∼800 • C or below increases the fraction of the BCC phase. [49] Annealing at temperatures higher than 800 • C increases the fraction of the FCC phase. Because the FCC phase is ductile and the BCC phase is relatively brittle, annealing can be used to tune the mechanical properties.
Some alloys have phases that are stable at intermediate temperatures only. These alloys may contain only simple phases in their as-cast state due to the higher cooling rate of Cu mold casting. Upon annealing, however, the intermediate-temperature phases will appear. For example, η phase [17,50] and σ phase [12,34,51] are known to form in some HEAs. One should pay great attention to the formation of these intermediate-temperature phases, because their formation can lead to significant changes in mechanical properties. Tsai et al. studied the stability of the σ phase. [52] They showed that the formation of the σ phase is predictable: it is directly related to the VEC of the alloy. As shown in Figure 7, [52] there is a σ -phase-forming VEC range for HEAs based on Al, Co, Cr, Cu, Fe, Mn, Ni, Ti, and V. Alloys whose VEC fall in this range develop σ phase upon annealing at 700 • C.

Computer-Aided Phase Prediction.
The multiprincipal-element nature of HEA translates to an immense amount of possible compositions. It takes huge amounts of time and cost to study all the compositions experimentally. Hence, it is critically important to implement highly efficient, low-cost methods to assist experimental study. For example, phase diagram is one of the most important tools to understand an alloy system. However, developing the phase diagram of just one senary HEA requires tremendous efforts-there are six composition axes! If the phase diagram could be reliably predicted with computational techniques, then experiments are only needed to verify some critical points on the predicted phase diagram, and the work needed is significantly reduced.
The main problem in applying existing phaseprediction techniques (e.g. CALculation of PHAse Diagrams, CALPHAD) on HEAs is probably the lack of a suitable database for multicomponent systems. [53] Thus, Zhang et al. developed a thermodynamic database for the Al-Co-Cr-Fe-Ni alloy system by extrapolating binary and ternary systems to wider composition ranges. [53] This database does not consider quaternary or quinary interaction parameters. Using their database, they obtained phase diagrams of the Al-Co-Cr-Fe-Ni system that agree with available experimental results.
In some cases, it seems that phase diagrams with reasonable accuracy can still be obtained even without the development of new databases. An Ni-based superalloy database was used to predict the equilibrium phases and phase fractions of Al 0. 5 [12,31] Except for some minor discrepancies, [12] experiments and predictions agree with each other. These examples demonstrate the effectiveness of computational methods in predicting the phase of HEAs. Although the accuracy of computational phase predictions for other HEA systems (other than Al-Co-Cr-Cu-Fe-Ni) is still not clear (the only result appears to be not as pleasing [54]), it is greatly expected that these techniques will become important tools for the future design and development of HEAs.

Phase Evolutions in Representative Alloy Systems.
The phase evolutions in some important alloy systems are discussed, along with their mechanical properties, in Sections 4.2 and 4.3. This is because the mechanical properties of HEAs are directly related to the phases contained in the alloys (see Section 4.1). Therefore, the best way to understand the mechanical properties is to discuss phase evolution and mechanical properties together.

Basic Concepts and Deformation Behaviors.
Because of the wide composition range and the enormous number of alloy systems in HEAs, the mechanical properties of HEA can vary significantly. In terms of hardness/strength, the most critical factors are: (1) Hardness/strength of each composing phase in the alloy; (2) Relative volume ratio of each composing phase; (3) Morphology/distribution of the composing phases. The first factor is largely determined by the crystal structure and bonding of each phase. In our experience, the phases can be roughly classified into four categories each having a different hardness range. This is provided in Table 2. It is important to note that the hardness ranges listed in Table 2 are merely typical ranges and there could be exceptions. Valence compounds are based on very strong covalent bonding and are essentially ceramics. Thus, their hardness is the highest. Non-simple intermetallic phases often lack easily accessible slip systems and dislocation activities are thus largely hindered. BCC and BCC-derivative structures (see the second paragraph of Section 3) are harder than their FCC counterparts because BCC-based structures have stronger directional bonding and lack a truly close-packed slip plane. [55,56] The general rule to estimate the hardness/strength of an HEA is straight forward: the harder the phase (and the higher the fraction of the hard phase), the harder the alloy. When two HEAs have phases with similar hardness and relative fraction, the distribution of the phases can also play an important role. This is exemplified in later sections. The ductility of HEA is also related to the phase in the alloy. As can be expected, harder phases usually have lower ductility.
The deformation microstructure and mechanism in most HEAs is unclear. However, such information was revealed in a single-phase, FCC CoCrFeMnNi alloy via detailed TEM study. [57] It was found that at small strains (<2.4%), the deformation is governed by planar slip of 1/2 110 type dislocations on {111} planes. At a strain around 20%, the deformation behavior depends on the deformation temperature. At 77 K, higher strain leads to the prevalence of deformation twinning. At 293 K, however, higher strain leads to the formation of dislocation cell structures. The texture evolution (during cold rolling) and recrystallization behavior in the same alloy have also been studied. [58] The texture after 90% cold rolling was predominantly brass-type, which indicates low stacking fault energy (SFE). Bhattacharjee et al. suggested that the lower SFE in HEAs is due to the higher free energy of the 'perfect crystal' in HEAs. This is explained in further detail in the last paragraph of Section 5. The recrystallization texture in the CoCrFeMnNi alloy retains components from previous cold rolling, which is similar to low-SFE TWIP steels and 316 stainless steel. However, the relative proportions of various texture components between HEA and other low-SFE alloys are different. The main reasons are the slower recrystallization rate in regions with brass texture, and the absence of preferential nucleation of brass-textured grains from shear bands. [58] The low SFE in the CoCrFeMnNi alloy as suggested by Bhattacharjee et al. [58] is indeed confirmed by Zaddach et al. [59] They combined experiments with first-principle calculations and concluded that the SFE of CoCrFeMnNi is between 18.3 and 27.3 mJ/m 2 . This value is close to that of conventional low-SFE alloys such as AISI 304L and brass.
The cold deformation and annealing behaviors of another FCC-based HEA, the Al 0.5 CoCrCuFeNi alloy, was also investigated. [60] Deformation twinning is observed again in this alloy. Prevalence of twinning appears to result from the blockage of local slip by the Widmanstätten precipitates. The above results suggest that in FCC HEAs, dislocation slip and twinning are still the main deformation mechanisms.

The Al-Co-Cr-Cu-Fe-Ni Alloy System.
Of all the HEA systems reported, Al-Co-Cr-Cu-Fe-Ni [7,8,21,35,36,49,[60][61][62][63][64][65][66][67][68] and its Cu-free version Al-Co-Cr-Fe-Ni [69][70][71][72][73][74] have been studied most comprehensively. Most of the reported HEA systems emerge from them. The main difference between the two is the presence of Curich interdendrite in the former. This is because Cu has positive mixing enthalpies with most of the other elements, and is thus repelled to the interdendrite region. Additionally, remnant Cu in the dendrite region also clusters and forms various kinds of precipitates. [21,35,60] The main phases in cast Al-Co-Cr-Cu-Fe-Ni alloy include FCC, BCC/B2, and the Cu-rich phase (also has an FCC structure). The relative volume ratio of these phases depends on the composition. The relative volume of BCC and FCC phases is related to the VEC of the alloy. Higher VEC typically leads to more of the FCC phase, and vice versa. Al has the strongest effect in this regard, [36] and addition of Al leads to the transition of the FCC phase to the BCC phase. Additionally, when there is sufficient Al, BCC tend to further decompose to a (Al, Ni)-rich B2 phase and a (Cr, Fe)-rich BCC phase that have almost identical lattice parameters. [8,21,71,74] The Cu-rich phase appears clearly when the concentration of Cu is higher than ∼10 at.%. Its fraction increases with the content of Cu. In general, hardness of the FCC phase is between HV 100-200. [7,8,[70][71][72] Alloys having a sole FCC phase have ductilities between 20-60%, and usually exhibit significant work hardening. [62,63,69] Hardness of the BCC/B2 phase is typically between HV 500-600. Because addition of Al triggers the formation of hard BCC/B2 phase, and hardness of the alloy increases with Al content [7,8,[70][71][72] (see Figure 3). Alloys having sole BCC/B2 phases probably have ductilities less than 5%. [62] In terms of their plasticity, values of 30% or higher have been reported. [74] Superplasticity was observed in AlCoCrCuFeNi at 800-1000 • C, and the elongation was between 405 and 800%. [65] Effect of annealing on the mechanical properties of the Al-Co-Cr-Cu-Fe-Ni system depends on the phase transformation involved (see Section 3.4). When the selected annealing temperature favors the formation of BCC phase, the alloy becomes harder and more brittle. When the temperature selected increases the fraction of FCC phase, the alloys softens and becomes more ductile. It should be carefully noted that in this alloy system, a long annealing time may lead to the formation of σ phase, [12] which hardens the alloy but remarkably reduces its ductility/plasticity. [12,34,51]

Derivatives of the Al-Co-Cr-Cu-Fe-Ni Alloy
System. This section discusses the mechanical properties of alloys derived from the Al-Co-Cr-Cu-Fe-Ni system. A system is said to be 'derived' from Al-Co-Cr-Cu-Fe-Ni when the number of different principal element between the two systems is less than two. For example, Al-Co-Cr-Cu-Ti is derived from Al-Co-Cr-Cu-Fe-Ni (replace Fe with Ti and remove Ni).
Addition of Ti to Al-Co-Cr-Cu-Fe-Ni usually leads to formation of intermetallic phases such as the Laves phase, [15,[75][76][77][78][79] σ phase, [17,75,76] Heusler phase, [17] η-Ni 3 Ti, [17,50] and R phase. [76] This is because Ti Figure 8. Compressive true stress-strain curves of AlCoCrCuFeNiTi x alloys with x being 0, 0.5, 1, and 1.5. [77] has large negative enthalpy of mixing with rest of the composition. These phases strengthen the alloy. For example, AlCoCrFeNiTi 0.5 has a high yield strength of 2.26 GPa and a plasticity of 23% (Figure 8). [77] The hardness of Al 0.5 CoCrCuFeNiTi x alloy can be significantly increased to HV 600 or higher when x is larger than 1. [75] The hardness of Ti-containing alloys may further increase upon annealing owing to the formation of more intermetallic phases. [17,50] Addition of Mo to the system generally leads to the formation of (Cr, Mo, Co, Fe)-rich σ phase. [80][81][82][83][84][85][86][87] σ phase is a very hard and brittle intermetallic phase. As mentioned previously, it significantly enhances the hardness of the alloy but reduces its ductility/plasticity. For example, addition of 0.5 mole fraction of Mo to AlCoCrFeNi raises its strength from 1,051 to 2,757 MPa. [86] Addition of Mn leads to two representative systems, the CoCrFeMnNi and Al x CrFe 1.5 MnNi 0.5 alloys. The CoCrFeMnNi alloy is a classic sole-FCC HEA and does not experience phase transformation when annealed. Therefore, it is a good model system to study the behavior of multiprincipal-element solid solutions and has been used to study the grain growth, [88] diffusion, [19] dislocation behavior, [57] and texture evolution [58] of FCC HEAs. This alloy is soft (yield strength ∼170 MPa), very ductile (elongation ∼60%), and shows high strain hardening capability at room temperature. At 77 K, its strength nearly doubles and its elongation increases to ∼80%. [57] The Al 0.3 CrFe 1.5 MnNi 0.5 alloy is a representative age-hardening HEA. When aged between 600 and 900 • C, its hardness boosts from about 300 HV to almost 900 HV. [34,51,89] TEM observation shows that the hardening phenomenon is not from precipitation hardening, but from the transformation of the BCC matrix to the σ phase during aging. [34] The hardening phenomenon in this alloy can actually be controlled to occur only near the surface, forming a surface hardening layer that significantly improves wear resistance (see Section 4.5).
Addition of non-metallic elements such as Si and B leads to the formation of valence compounds such as silicides and borides. These compounds also lead to evident hardening. For example, the hardness of AlCoCr-FeNiSi [90] and AlBCoCrCuFeNi [14] are HV 738 and HV 740, respectively. Compound formation also causes loss of toughness. Addition of Si to AlCoCrFeNi by the equimolar ratio reduces its plastic strain from over 25% to ∼1%. [90] It should be noted that although valence compounds are typically harder than intermetallic phases, silicide/boride-containing alloys are not necessarily harder than alloys containing intermetallic phase. In the two alloys mentioned above (AlCoCrFeNiSi and AlBCoCrCuFeNi), the volume fractions of silicides and borides are not high, and the distribution of these compounds is quite localized. Hard phases with such morphology (isolated coarse laths) are less effective in strengthening, particularly when they are embedded in an apparently softer matrix. Besides the above, many other elements have been added to the Al-Co-Cr-Cu-Fe-Ni system (e.g. Zr [91], Nd [91], Nb [92], V [93,94], Y [95], Sn [96,97], Zn [98], and C [99,100]).

Refractory HEA systems.
Some researchers developed HEAs that are based on totally different elements. The most notable system is the refractory HEA system developed by Senkov et al. [10,37,54,[101][102][103][104][105][106] These alloys are composed mainly of refractory elements such as Cr, Hf, Mo, Nb, Ta, V, W, and Zr. The major phase in these alloys is typically BCC. Many of them even have a sole BCC structure. These BCC solid solutions have high strength ranging from 900 to 1,350 MPa. [37,101,105,106] Interestingly, such strength is several times higher than that of the weighted average strength of the composing elements. For example, the rule-of-mixture hardness of MoNbTaVW and HfNbTa-TiZr are 1,596 and 1,165 MPa, respectively. However, the hardness measured for these alloys are more than three times higher: 5,250 and 3,826 MPa, respectively. Such high hardness is a result of the solid solution strengthening effect which originated from the difference in atomic size and modulus among the composing elements. [37] Room temperature plasticity of refractory HEA is closely related to composition. The plastic fracture strain is less than 5% for many alloys. However, some alloys can be compressed to 50% without fracture. [37,103,106] It was suggested that the high plasticity in some alloys is owing to the prevalence of deformation twinning. [37,103] It is worth mentioning that refractory HEAs can show superb mechanical properties at elevated temperatures, which is described in Section 5.

Wear and Fatigue Properties.
Wear properties of HEAs, under both abrasive [7,14,51,75,94,107] and adhesive [50,66,80] conditions, have been tested in a number of systems. The adhesive wear properties of tested HEAs are summarized in Figure 9. HEAs composed solely of SDPs typically do not show better wear properties than conventional alloys with similar hardness (e.g. compare Al 0.5 CoCrCuFeNi and 316SS in Figure 9). Wear resistance is clearly enhanced in the presence of some B2 or COPs (e.g. Al 2 CoCrCuFeNi and AlCoCrFe 1.5 Mo 0.5 Ni). If COPs with high hardness become the main phase, wear resistance is often outstanding. Many HEA systems of this type have comparable or even better performances than conventional wearresistant alloys such as SKD61 tool steel (AISI H13) and SUJ2 bearing steel (AISI 52100). [50,51,80,107] For example, the wear resistance of the AlCoCrFeMo 0.5 Ni alloy is comparable to that of SUJ2 (Figure 9). In particular, the Al 0.2 Co 1.5 CrFeNi 1.5 Ti alloy [50] has significantly Figure 9. Relationship between hardness and adhesive wear resistances of various HEAs and conventional alloys. [50,66,80] better wear resistance than SUJ2 and SKH51 (AISI M2, see Figure 9). Its hardness is similar to that of SUJ2, but its wear resistance is 3.6 times that of the latter. The wear resistance of the Al 0.2 Co 1.5 CrFeNi 1.5 Ti alloy is twice that of SKH51, but the hardness of the latter is a lot higher. The outstanding performance is owing to the higher hardness of HEA at high temperatures. Additionally, because the wear mechanism is mild oxidational wear, the better oxidation resistance in the Al 0.2 Co 1.5 CrFeNi 1.5 Ti alloy is also beneficial to its wear resistance. [50] The Al 0.3 CrFe 1.5 MnNi 0.5 alloy (see also Section 4.3) has a unique surface hardening phenomenon that can increase its wear resistance by ∼50%, but only reduce its bending strain by 2%. This is achieved by limiting the formation of hard σ phase at the surface region. Because the heterogeneous nucleation energy and strain energy near the free surface are lower than those in the interior of the alloy, one can carefully select the aging temperature so that the kinetics of the transformation is sufficient high only at the near-surface region. The process is simple (atmospheric aging), applicable to objects with complex shapes, and offers a markedly better combination of mechanical properties than SUJ2 (30% higher wear resistance and 300% larger maximum bending strain).
Preliminary fatigue results are also encouraging. Al 0.5 CoCrCuFeNi alloy can show long fatigue lives at relatively high stresses, high endurance limits (540-945 MPa) and endurance limit to ultimate tensile strength ratio (0.402-0.703). [67] These values are comparable to or even better than many conventional alloys, such as various steels, Ti alloys, and Zr-based BMGs. [67] Although some degree of scattering in fatigue life was observed, it is probably due to defects introduced during sample preparation. If defects can be carefully controlled, the fatigue resistance characteristics in Al 0.5 CoCrCuFeNi are very promising. Figure 10. Hardness of various HEAs and conventional alloys as functions of temperature. Some of the data is collected from literature. [50,85]

High Temperatures Properties and Structural
Stability Many HEAs have exceptional hightemperature strength/hardness, which originates from their excellent resistance to thermal softening. Figure 10 shows the hot hardness of three HEAs and three conventional alloys as functions of temperature. [50,85] The AlCoCr 2 FeMo 0.5 Ni and Al 0.2 Co 1.5 CrFeNi 1.5 Ti alloys have very high hardness at room temperature (more than two times higher than Inconel 718 superalloy). Their hardness decreases slowly with the increase of temperature. SKH51 and SUJ2 also have very high hardness at room temperature (over HV 700). However, in contrast to the gradual softening behavior for the two HEAs, SKH51 and SUJ2 start to soften drastically at 200 and 600 • C, respectively. This leads to a significant difference in hardness for the four alloys at 800 • C-the two HEAs are still more than two times harder than the Inconel 718 superalloy, while the two conventional alloys become apparently softer than Inconel 718 (Figure 10). At 1000 • C, AlCoCr 2 FeMo 0.5 Ni is already more than three times harder than Inconel 718. Also note that even the soft, sole-FCC Al 0.5 CoCrCuFeNi alloy is quite resistant to thermal softening.
Hsu et al. [85] calculated the high-temperature softening coefficient of the AlCoCr x FeMo 0.5 Ni (x = 0-2.0) Table 3. Softening coefficient of two AlCoCr x FeMo 0.5 Ni alloys and two conventional high-temperature alloys above transition temperature (T T ). [85] Softening coefficient Alloy above T T alloys from the hot hardness vs. temperature curves and found that above the transition temperature (onset temperature for rapid softening), the softening coefficient of these HEAs are evidently lower than that of conventional high-temperature alloys such as Inconel 718 and T800 (Table 3). [85] Softening coefficient of some alloys are even less than half that of Inconel 718. The resistance to thermal softening is mainly a result of the sluggish diffusion in HEAs. Above the transition temperature, deformation proceeds via diffusion-related mechanisms. Thus, the slower diffusion rate in HEAs can effectively reduce the degree of thermal softening. Because the transition temperature and diffusion rate of metallic materials are directly related to the melting point of the material, it can be expected that refractory HEAs should have outstanding properties in this regard. Indeed, W-containing refractory HEAs have excellent resistance to thermal softening. This is shown in Figure 11. [101] The yield strengths The two alloys have extremely high yield strength of over 400 MPa at 1600 • C. Owing to their good strength at high temperatures, some refractory alloys actually have higher specific yield strength than superalloys such as Inconel 718 and Haynes 230 at higher than 1000 • C. [106] Structural stability is an important issue for hightemperature applications. HEAs generally have good structural stability at high temperatures. This is owing to (1) slower kinetics due to the sluggish diffusion effect (Section 2.2); (2) reduced driving force to eliminate defects (e.g. grain boundaries, interfaces, etc.). The driving force to eliminate defects originates from the free energy difference between defect-containing and defectfree crystals. However, the defect-free lattice in HEAs is still highly strained and full of atomic-scale zigzag, owing to the atomic size differences between the elements. The strain energy of the distorted lattice raises the overall free energy, leading to a reduced energy difference between defect-containing crystal/amorphous structure and the defect-free crystal. [23,108] Thus, the driving force to eliminate defects is reduced. There is much evidence for the good structural stability of HEAs. For example, the activation energy of grain growth for the CoCrFeMnNi alloy (which has a sole FCC structure) is 321.7 kJ/mole −1 . [88] This value is twice that of AISI 304LN, indicating a significantly slower kinetics for microstructural coarsening. [88] Growth of precipitates in HEAs is also slow. Precipitates in HEAs treated at moderate temperature (e.g. 700-800 • C) often still remain small-diameters on the scale of tens or hundreds of nanometers have been reported. [34,82] Another example is the NbSiTaTiZr amorphous alloy. The crystallization temperature of NbSiTaTiZr is higher than 800 • C. [23] This temperature of stability is probably the highest among reported amorphous metals. The excellent structural stability comes not just from the above two reasons. NbSiTaTiZr was designed based on a free energy consideration. [23] Its composition was selected so that the non-crystalline random solid solution state has a low free energy. This lowers its driving force to crystallize into silicides. Additionally, the kinetics in this alloy is further slowed due to the high atomic packing density and high melting points of composing elements.

Corrosion Properties
The corrosion resistances of some HEAs have been tested in both NaCl and H 2 SO 4 solutions. [20,73,[109][110][111][112][113][114][115][116][117][118][119][120][121] In both solutions, some HEAs show better corrosion properties than 304 SS and even 304L SS and have good resistance to pitting. [109,[115][116][117] Important factors include alloy composition and microstructure, particularly the amount and distribution of corrosion-resistant elements (such as Cr), and the presence of galvanic corrosion. Most of these tested alloys are based on Co, Cr, Fe and Ni. Therefore, the corrosion behavior of the Co-Cr-Fe-Ni alloys is discussed first, followed by the effect of elemental addition to the alloy system.
Corrosion in NaCl Solution: In NaCl solution, the CoCrFeNi single-phase FCC alloy has markedly better corrosion resistance than 304L SS. [111] This is probably owing to its high Cr and Ni content. Addition of Cu to the CoCrFeNi alloy leads to the formation of Cu-rich interdendrite phase, which suffers from galvanic corrosion and severely degrades the corrosion resistance. [111,118] To make things worse, the passive film on the Curich interdendrite regions does not offer good protection, which further narrows down the passivation region. [111] The corrosion property can be improved by reducing the amount of Cu-rich phase via high-temperature annealing. [118] Addition of 0.5 mole fraction of Al to the CoCrFeNi alloy leads to the formation of (Al, Ni)rich BCC phase and causes galvanic corrosion in NaCl solution. [73] Al is also detrimental to the corrosion resistance of the Al x CrFe 1.5 MnNi 0.5 alloys. Addition of Al significantly reduces the pitting potential and increases the area of localized/pitting corrosion of these alloys in NaCl solution. [114] Although it has higher pitting potential, the passive region is less than that of 304SS. Addition of Mo to the Co 1.5 CrFeNi 1.5 Ti 0.5 Mo x alloy is beneficial because it evidently raises the pitting potential. [115,119] For example, Co 1.5 CrFeNi 1.5 Ti 0.5 Mo 0.1 alloy has a wide passivation region of 1.43 V in 1 M NaCl and does not suffer from any pitting. [116,119] Corrosion in H 2 SO 4 Solution: In H 2 SO 4 solution, CoCrFeNi also has better corrosion resistance than 304 SS. [117] Addition of Cu to the CoCrFeNi alloy deteriorates its corrosion properties. Similar to the case in NaCl solution, this is due to the formation of the Cu-rich phase. The width of the passivation region is also evidently narrowed. [109,113] Addition of Al is also harmful. [114,117] It leads to the formation of a FCC+BCC or BCC+B2 two-phase structure, in which the Ni,Al-rich BCC/B2 phase is preferentially corroded. [74,117] Additionally, the existence of Al turns the passive film porous and less protective. [117] Replacing Co with Mn in the Co-Cr-Fe-Ni alloy degrades its corrosion resistance in H 2 SO 4 solution, and renders its resistance inferior to 304 SS. [114] Addition of B to the Al 0.5 CoCrCuFeNi alloy leads to the formation of Cr-rich borides and apparently lowers the Cr content in the matrix. [113] This renders the matrix much less corrosion resistant. Addition of Ti to the Co-Cr-Fe-Ni alloy does not affect its outstanding corrosion behavior. [119] Addition of Mo to the Co-Cr-Fe-Ni-Ti alloy is detrimental to the corrosion resistance in H 2 SO 4 solution. This is because of the formation of the (Cr, Mo)-rich σ phase, which reduces the concentration of Cr in other phases. [119] The AlCu 0.5 CoCrFeNiSi alloy has better general corrosion resistance than 304 SS, but its passivation region is smaller. [109] 7. Outstanding Issues and Future Directions HEA is an immense field with a countless number of new alloy systems. Our understanding of these new materials is still very limited and preliminary. As mentioned in Section 4, our knowledge mainly focuses on the Al-Co-Cr-Cu-Fe-Ni system, its derivatives, and the refractory HEAs-just a very tiny fraction of the whole HEA world. More elements and combinations need to be explored to further understand the potentials of HEAs. For example, light HEAs are of great interest. [122][123][124] Non-metallic elements such as C, B, and N can also be incorporated to achieve even higher strength. [14,99,100,125,126] Additionally, HEAs can also be prepared by other process and/or in other forms. For example, HEAs can be processed by a wrought process including homogenization, hot/cold working, and annealing [62,64] to eliminate casting defects and improve microstructure. HEA powders can be prepared by ball milling. [98,[127][128][129][130][131] The powders can subsequently be sintered to form bulk HEAs with ultrafine structure. [98,127,129] Single-crystal of the Al 0.3 CoCrFeNi HEA has also been prepared successfully by the Bridgman solidification method. [132] Cemented carbides and cermets could be sintered by a proper HEA binder to reduce cost and improve resistance to softening. HEA coatings/claddings can also be prepared by sputtering or other cladding techniques. [22,23,73,[133][134][135][136][137][138][139][140] In terms of the phase and crystal structure of HEAs, most of our knowledge is based on the as-cast state (after vacuum arc melting). Very little is known about the equilibrium phases and phase diagrams, but these are key to the design and development of HEAs. Significantly more combined experimental and computational works are needed in this regard. It is worth mentioning that although equimolar compositions are usually the easiest access to a new alloy system, they probably do not possess the best combination of properties. Therefore, there is probably far more treasure in those non-equimolar alloys with carefully designed composition and tailored microstructures. A good approach is to start from equimolar alloys and then extract the desired non-equimolar composition from the equimolar alloy.
The mechanical properties of HEAs available are mostly about hardness and compressive properties. Because reasonable ductility is crucial to structural applications, more studies on tensile behaviors are definitely needed. Additionally, except for a few examples, [57][58][59]141] not much is known about the deformation of HEAs, e.g. dislocation behavior, deformation twinning, deformation microstructures, texture evolution, stacking fault energy, etc. This information is the foundation for a thorough understanding of mechanical properties. Microstructural characterization in most works is only conducted by SEM. This is not sufficient for unraveling the structure-property relationship in HEAs because all detailed microstructure observations reveal the existence of nano-sized or even atomic-scale precipitations and/or decompositions. [21,31,32,34,35,98,142] More tests aimed at prolonged service and/or high-temperature applications are also required. These include but are not limited to: fatigue behavior, [67] oxidation resistance, [50,51,104] structural stability, and creep behavior.
HEAs can also be used in other applications. For example, HEAs can be used as hydrogen storage materials, [143,144] radiation resistant materials, [145] diffusion barriers for electronics, [22,23] precision resistors, [4,146] electromagnetic shielding materials, soft magnetic materials, [147] thermoelectric materials, functional coatings, and anti-bacterial materials. Because of the distinct design strategy and unique properties, HEAs and other high-entropy materials (e.g. highentropy ceramics or even polymers) will probably find applications in fields significantly wider than those listed above. This exciting new virgin field awaits exploration by more scientists.