**Corrosion, Erosion and Wear Behavior of Complex Concentrated Alloys: A Review**

#### **Aditya Ayyagari 1, Vahid Hasannaeimi 1, Harpreet Singh Grewal 2, Harpreet Arora <sup>2</sup> and Sundeep Mukherjee 1,\***


Received: 18 June 2018; Accepted: 24 July 2018; Published: 3 August 2018

**Abstract:** There has been tremendous interest in recent years in a new class of multi-component metallic alloys that are referred to as high entropy alloys, or more generally, as complex concentrated alloys. These multi-principal element alloys represent a new paradigm in structural material design, where numerous desirable attributes are achieved simultaneously from multiple elements in equimolar (or near equimolar) proportions. While there are several review articles on alloy development, microstructure, mechanical behavior, and other bulk properties of these alloys, then there is a pressing need for an overview that is focused on their surface properties and surface degradation mechanisms. In this paper, we present a comprehensive view on corrosion, erosion and wear behavior of complex concentrated alloys. The effect of alloying elements, microstructure, and processing methods on the surface degradation behavior are analyzed and discussed in detail. We identify critical knowledge gaps in individual reports and highlight the underlying mechanisms and synergy between the different degradation routes.

**Keywords:** corrosion; surface degradation; wear; high entropy alloys; complex concentrated alloys; potentiodynamic polarization; erosion-corrosion; slurry-erosion; oxidation wear; highly wear resistant coatings

#### **1. Introduction**

Development of materials having superior surface degradation resistance has been a major thrust area of research in modern metallurgy. Loss of material in the form of corrosion, erosion, and wear results in economic impact in the range of billions of dollars worldwide by some estimates. Several technologies have not realized their full potential due to the lack of materials that can withstand surface degradation in critical applications. Specific examples include core walls, diverters, and reactor vessels in nuclear reactors that can withstand hot corrosion, contact with molten metals, high-pressure water, and exposure to super critical temperatures. Similarly, there is high demand for developing materials with improved wear resistance for enhancing the energy efficiency of turbines, windmill rotors, and automobiles. These call for structural components that can withstand high torque and resist metallurgical changes that are caused by frictional heat and high pressure. Problems, such as white matter in bearings, spalling, and deterioration in mechanical properties under operating conditions remain as a major challenge. In addition, synergistic combination of different surface degradation mechanisms leads to accelerated material loss. Examples include simultaneous wear and corrosion seen in food processing and chemical handling industries. Erosion is another

significant source of material loss, where particulate materials, such as sand and debris entrapped in a moving/impinging liquid, degrade the surface integrity of materials.

Traditionally, development of materials that simultaneously meet multiple application requirements has been done by adding minor proportions of alloying elements to the base material and tailoring the heat treatment. Examples include aluminum alloys, where tempering treatments are used to obtain a balance in mechanical properties and corrosion resistance. In that regard, multi-principal element alloys represent a new paradigm in structural material design, where numerous desirable attributes are achieved simultaneously from multiple elements in equimolar (or near equimolar) proportions [1–7]. These alloys are typically referred to as high entropy alloys (HEAs) or more generally as complex concentrated alloys (CCAs). High configurational entropy leads to single-phase solid solutions in a certain subset of these multi-component systems. It was initially believed that the core effects, such as high configurational entropy [8], lattice distortion [9], and sluggish diffusion [10] may have resulted in a gamut of attractive properties including high strength-ductility combination [6,10,11], resistance to oxidation, corrosion and wear properties [12,13]. However, recent reports suggest that these may not be the only structure-property determining parameters, thus leaving a large scope for understanding the physical metallurgy of complex concentrated alloys [14–16]. Another advantage of the complex concentrated approach is the vast number of alloy systems that can be developed from a small palette of elements by focusing on the central region of the multi-component phase space, rather than the edges [16].

With exponentially growing interest in complex concentrated (or high entropy) alloys, there are several reports in literature on the surface degradation behavior of these multi-component systems. The corrosion behavior of high entropy alloys has been discussed in a recent review [17]. However, a clear understanding of the underlying mechanisms and synergy between the different surface degradation routes is lacking. Here, we provide a comprehensive overview of corrosion, erosion, and wear behavior of complex concentrated alloys to elucidate the similarities and in certain cases the unique differences in response to different environments. The effect of alloying elements, microstructure, and processing methods on the different surface degradation routes are analyzed and discussed in detail.

#### **2. Evaluation of Surface Degradation Mechanisms**

In this section, the methods used in literature for evaluation of corrosion, wear and erosion behavior of complex concentrated (or high entropy) alloys are summarized along with the pertinent metrics for quantifying the extent of damage.

#### *2.1. Corrosion Characterization*

Corrosion behavior of complex concentrated alloys has been evaluated by immersion (or mass loss/gain) test, open circuit potential measurement with time, potentiodynamic polarization, and anodic polarization. Immersion test is the simplest, where the change in mass of the sample is measured by assessing the damage that is caused by the environment in which it is immersed (ASTM G31). The corrosion rate is calculated as:

$$\text{Corrosion rate } = \frac{(K \times W)}{(A \times T \times D)} \tag{1}$$

where, *K* is a constant, *T* is time of exposure in hours, *A* is area in cm2, *W* is mass loss/gain in g, and *D* is the density in gm/cm3.

Accelerated assessment of corrosion performance can be made using electrochemical corrosion tests. When no external current or potential is applied to a metal immersed in an electrolyte, the system eventually reaches equilibrium and the net current measured is zero. The potential developed on the surface of the electrode when the metal is immersed into the electrolyte is called the open circuit potential (OCP). For potentiodynamic polarization, three types of reference electrodes are typically used, namely saturated calomel electrode (SCE), Ag/AgCl electrode, and standard hydrogen electrode (SHE). In a three-electrode set up, one of the aforementioned electrodes is connected as reference electrode, the sample as working electrode, and platinum or graphite as counter electrode. In the potentiodynamic polarization test, the sample is subjected to a potential sweep typically from −250 mV with respect to OCP to at least +250 mV at a scan rate of 0.16 mV/s. Scanning beyond 250 mV above OCP may cause further anodic reactions, such as breakdown of the protective surface oxides and pitting. Potentiodynamic polarization tests are extensively used to identify critical corrosion parameters such as pitting potential, passivation range, corrosion current, and re-passivation potentials in an accelerated way. Corrosion rate is calculated as:

$$\text{Corrosion Rate} = \frac{K \times i\_{\text{corr}} \times EW}{\text{Density}} \tag{2}$$

where, *<sup>K</sup>* is 3.27 × <sup>10</sup>−<sup>3</sup> mm g/(μA·cm·year), *<sup>i</sup>*corr is the corrosion current density, and *EW* is the equivalent weight. Equivalent weight is calculated from the expression:

$$\text{Equivalent weight} = \left\{ \sum \frac{f\_i \times n\_i}{w\_i} \right\}^{-1} \tag{3}$$

where, *fi* is the mass fraction, *wi* is the atomic weight, and *ni* is the valence of the *i*th element in the alloy [18].

#### *2.2. Wear Testing*

The wear behavior of complex concentrated alloys has been evaluated using three techniques, namely sliding reciprocating wear test, pin on disc test, and modified pin on disc test (pin-on-belt test). The fundamental working principle is the same in all three tests—a normal load is applied on to a sample while it is in contact with a reference material. Depending on the test type, either the reference material or the sample are moved to cause a relative motion between the surfaces. In pin-on-disc and pin-on-belt tests, the sample is made into the form of a stationary cylinder called "pin", which is brought in contact with a rotating disc made of hardened steel. A test load is applied normal to the pin, producing wear at the interface of the two materials as shown in Figure 1a. The rotating steel disc may be replaced with a moving belt, typically coated with Alumina or Silica abrading media, as shown in Figure 1b. In the sliding reciprocating test, the sample is made in the form of a flat plate and loaded under a hard counterface, such as WC or Si3N4 ball indenter or steel pin, as shown in Figure 1c. The stage slides at a set frequency and stroke length. The wear volume loss is quantified while using weight loss, contact profilometry, or interferometry.

Quantification of loss during wear test is done from the volume of worn material removed (*V*w) and relating it to total sliding distance (*L*) and load (*F*). Wear volume loss for most engineering materials increases with decreasing hardness, as given by Archard's relation:

$$V\_{\rm w} = K \frac{L \times F}{H} \tag{4}$$

where, *K* is the dimensionless wear coefficient and *H* is the hardness. Certain high entropy alloys followed the Archard's wear relation in sliding wear test. The engineering unit of wear resistance is measured as wear volume loss per unit distance of sliding and expressed in the dimensions of [*L*]/[*L*] 3.

**Figure 1.** Illustration of (**a**) pin-on-disc test setup; (**b**) pin-on-belt setup; and (**c**) sliding reciprocating wear stage.

#### *2.3. Erosion and Erosion Corrosion Characterization*

Erosion is a form of material degradation characterized by the progressive loss of material from a solid surface due to mechanical interaction between the surface and a fluid or impinging liquid containing solid particles. There are very limited number of reports on erosion behavior of complex concentrated (or high entropy) alloys. Erosion that is caused by the impact of solid particles entrained in gaseous medium is termed as solid particle erosion. On the other hand, if a liquid is used as a carrier medium, the process is termed as slurry erosion [19,20]. The impact of the entrained abrasive particles results in micro-cutting and severe plastic deformation of the target surface depending on the operating parameters (Table 1). Other forms of erosion, which result from the interaction between a solid surface and fluid alone, are cavitation erosion and liquid droplet erosion [21–23]. In the case of cavitation erosion, degradation takes place due to implosion of cavities/bubbles in the liquids. Implosion of such cavities results in the formation of high velocity micro jets or shockwaves affecting the solid surfaces. Contact pressures at the point of impact can reach several hundred Giga-Pascals, which is sufficient for the deformation and removal of material. Several parameters influence the erosion processes, which may be classified into flow related, erodent related, and materials related parameters (Table 1) [24–28]. Chemical/electrochemical interactions (corrosion) are also possible along with erosion depending on the working environment. The synergy between erosion and corrosion can further aggravate the material degradation. The synergistic effect in tribo-corrosion process due to interaction between erosion and corrosion is given as [25]:

$$S = \mathcal{W} - (E + \mathbb{C})\tag{5}$$

where, *S* is the synergy, *W* is the material removal rate by combined erosion and corrosion process, *E* is the material removal rate by pure erosion process, and *C* is the material removal rate by pure corrosion process. The synergy, *S*, is further composed of: (1) erosion induced corrosion (Δ*CE*) and (2) corrosion induced erosion *(*Δ*EC*). The factors that are responsible for erosion-induced corrosion are increase in surface area due to roughening effect, increased strain hardening and dislocation density, mechanical damage of the passive layer, and increased local temperatures. The factors contributing towards corrosion-induced erosion may be dislodgement of hard particles due to corrosion of the matrix, weakened grain boundaries, inter-granular pitting, and accelerated cracking due to crevice corrosion.


**Table 1.** Process parameters that affect the erosion process [24].

#### **3. Corrosion Behavior of Complex Concentrated Alloys**

Majority of the complex concentrated alloys studied so far for their corrosion behavior are based on the CoCrFeNi equimolar system. The observed corrosion behavior in these alloys may be broadly classified based on their composition and the resulting surface passivation layers, microstructural heterogeneity, phase segregation and associated galvanic corrosion, and finally, the test environment. CoCrFeNi-Cu*<sup>x</sup>* (where "*x*" indicates varying proportions) was one of the earliest developed alloys, where the effect of increasing copper content on the microstructure and corrosion properties was reported [29]. Immersion and potentiodynamic polarization tests were conducted in 3.5 wt% NaCl solution. The as-cast CoCrFeNi-Cu*<sup>x</sup>* alloys showed face centered cubic (FCC) phase mixture having distinct dendritic (copper lean) and inter-dendritic (copper rich) phases. In this alloy, the bright inter-dendritic regions were Cu rich as shown in Figure 2a. The X-ray diffraction (XRD) results show a single set of FCC peaks for several of these compositions, although the microstructure shows segregation between dendrites. This may be due to the very close *d*-spacing of the two phases that could not be resolved in XRD. The corrosion behavior of these alloys was comparable to SS304 stainless steel, with the copper free composition showing highest resistance to pitting (Figure 2b). The *x* = 0.5 alloy showed higher corrosion current density and pitting density, as shown in Figure 2c. This was attributed to higher galvanic action prompted from the Cu segregated in the inter-dendritic regions. Galvanic coupling results in initiation and propagation of localized corrosion pits causing rapid dissolution of the more anodic phase (in this case the bright Cu rich phase).

**Figure 2.** (**a**) As-cast microstructure of CoCrFeNi-Cu alloy showing dendritic microstructure. The copper rich interdendritic regions appear white, while the dendrites are darker; and, (**b**) Potentiodynamic polarization plots of the alloys in 3.5% NaCl. The Cu free CoCrFeNi alloy showed highest resistance compared to other two complex concentrated alloys (CCAs) and SS304L; (**c**) Microstructure after corrosion tests showing that Cu rich inter-dendritic regions corroded faster as compared to the dendritic regions lean in Cu. This may be due to the galvanic effect arising from the difference in composition [29] (reprinted with permission from Elsevier).

The effect of Al addition to CoCrFeNi has also been systematically studied and the resulting microstructure and corrosion behavior has been reported [30]. The Al0.1CoCrFeNi alloy shows a single-phase FCC structure with good microstructural stability. The XRD patterns for the alloy in recrystallized and as-cast state are shown in Figure 3a,b, respectively. The corresponding SEM microstructures are shown in Figure 3c,d.

**Figure 3.** X-ray diffraction curves for Al0.1CoCrFeNi alloy in (**a**) recrystallized state heat-treated at 900 ◦C for 20 h, as compared to its (**b**) as-cast state. Corresponding back scattered electron scanning electron microscopy (SEM) microstructures of the alloys in (**c**) recrystallized and (**d**) as-cast states [31] (reprinted with permission from Elsevier).

The corrosion behavior of Al0.1CoCrFeNi has been reported in as-cast and recrystallized states in 3.5 wt% NaCl solution [31,32]. The corrosion behavior of the alloy was superior to SS304 steel as seen in Figure 4a. The corrosion potential, corrosion current density, and pitting resistance (referred to as *E*BD in [32]) were comparable between as-cast and recrystallized states for Al0.1CoCrFeNi alloy. Minor variations between reported values in literature may be explained based on the microstructural differences between as-cast and recrystallized samples. Surface finish also plays an important role in determining the corrosion behavior. Potentiodynamic polarization for Al0.1CoCrFeNi alloy was compared with another single phase HEA, CoCrFeMnNi, as shown in Figure 4b. Both alloys showed wide passivation region and transient pitting, which may be an indication of local corrosion and the re-passivation on the surface.

**Figure 4.** Potentiodynamic polarization curves of (**a**) as-cast Al0.1CoCrFeNi CCA versus SS304 steel in 3.5 wt% NaCl solution [32] (**b**) rolled and recrystallized Al0.1CoCrFeNi CCA as compared to CoCrFeMnNi CCA [31] (reprinted with permission from Elsevier).

*Metals* **2018**, *8*, 603

The pitting resistance (Δ*E*) or breakdown resistance (*E*BD) of Al0.1CoCrFeNi alloy measured as the difference of pitting potential (*E*PIT) and corrosion potential (*E*CORR) in both conditions was ~1 V. Transient pitting was observed in both conditions around 0.6 V (highlighted with circle on the polarization curve), which may be an indicator of localized surface instability. More events of transient pitting were seen for the as-cast condition as compared to the recrystallized sample, which may be explained from the microstructural heterogeneity. Both as-cast and wrought alloys showed a high passivation resistance, 193 kΩ [31] and 115.5 kΩ [32], respectively, when tested for their EIS response. The superior corrosion resistance of Al0.1CoCrFeNi alloy has been explained based on the relatively high content of Cr and Ni that form a strong passivating surface layer. Pitting resistance is typically quantified based on the wt% of passivating elements that are present in the alloy. Particularly, Cr, Mo, and Ni enhance pitting resistance of most engineering alloys. Since CCAs have passivating elements as high as 20%, are reported to have excellent pitting and corrosion resistance, provided that there are no extraneous corrosion promoters, such as galvanic phases or physical surface aberrations. Al0.1CoCrFeNi high entropy alloy showed unique corrosion microstructures, as shown in Figure 5. Corrosion was initially observed to occur in the form of tiny pits, as shown in Figure 5a. Unique hierarchical features developed as a result of extensive grain boundary corrosion as well as micro/nano porosity formation within the grains, as shown in Figure 5c,d.

**Figure 5.** Pitting morphologies in Al0.1CoCrFeNi alloy after polarization test in 3.5% NaCl [32]. (**a**) large pitting on the sample when tested to current density of 10 mA/cm2; (**b**) low magnification image showing grain boundary corrosion; (**c**) high magnification image showing grain boundary corrosion and micro-porosity formation; (**d**) high magnification image showing small micro-porous structures on the surface after corrosion [31] (reprinted with permission from Elsevier).

The effect of increasing Al content on the corrosion behavior of Al*x*CoCrFeNi alloy system was investigated in sulfuric acid [33] as well as in NaCl solution [34]. Besides the effect of alloying elements, the effect of experimental variable, i.e., scanning rate and temperature of the alloys in the corrosive media was also studied. This is an important metric to be systematically studied since scan rate can significantly alter the values of corrosion rate measured [33,34]. Increasing Al content induced microstructural changes to Figure 7. Pitting morphology hat significantly affected the corrosion

behavior. The single phase FCC alloys for *x* = 0, *x* = 0.25, and *x* = 0.3 were more corrosion resistant when compared to the alloys containing higher fraction of Al. This was primarily attributed to phase separation induced by increasing Al content. Figure 6a shows secondary passive region for alloys with *x* = 0.5 and *x* = 1.0, which was attributed to the selective corrosion of dual phase face centered cubic (FCC)–body centered cubic (BCC) alloy (for *x* = 0.5) and BCC-ordered structure (for *x* = 1.0). In comparison, Figure 6b shows the potentiodynamic polarization charts for alloys with *x* = 0.5 and *x* = 0.7, displaying multiple transient pitting sites and continuous corrosion, both of which indicate the corrosion of secondary phases and partial passivation behavior of the matrix.

**Figure 6.** Potentiodynamic polarization curves of Al*x*CoCrFeNi alloys in (**a**) 0.5 M H2SO4 solution at Al content 0, 0.25, 0.5, 1.0; as compared with SS304, test performed at 25 ◦C; (**b**) 3.5 wt% NaCl solution for Al = 0.30, 0.50, 0.70 alloys. Both tests show gradual deterioration in corrosion in corrosion behavior with increasing Al content. Tests in NaCl solution resulted in extensive unstable pit formation on the sample, seen as short current spikes on the anodic branch [33,34] (reprinted with permission from Elsevier).

The single phase FCC alloys for *x* = 0 to 0.3 in Al*x*CoCrFeNi system showed significantly lower pitting density and pit depth when compared to the two phase alloys resulting from higher Al content. The *x* = 0.5 alloy showed pits on the inter-phase boundary between the FCC and BCC phases, indicating the formation of a galvanic couple. For *x* = 0.7 and *x* = 1.0, the BCC phase was observed to significantly/completely dissolved in H2SO4 solution and NaCl solution. Increasing Al content promotes the formation of BCC phase in Al*x*CoCrFeNi alloys, which undergoes selective dissolution. Increasing Al content likely results in the formation of porous Al oxide on the surface at the expense of more compact and passivating Cr oxide. A clear pattern of evolution of pitting morphology is seen as the Al content in gradually increased. Al = {0.1–0.3}: uniform pitting → Al = {0.5}: interphase galvanic corrosion → Al = 0.70–1.0 complete dissolution of BCC phase. Therefore, increasing the Al content beyond a threshold value resulted in higher pitting susceptibility, as seen in Figure 7.

**Figure 7.** Pitting morphology on Al*x*CoCrFeNi alloys after polarization experiments. The microstructure of the alloys with (**a**) Al = 0; (**b**) Al = 0.25; and (**c**) Al = 0.3 was reported to be of single phase, and consequent absence of galvanic corrosion sites. The microstructures with Al = 0.5–1.0 showed two-phase microstructure. This promoted accelerated corrosion at the interface between the two-phases (**d**) pits on face centered cubic-face centered cubic (FCC-FCC) interphase boundary; (**e**) BCC-BCC interphase boundary and (**f**) corrosion on ordered BCC phases [34] (reprinted with permission from Elsevier).

Heat treatment of Al0.5CoCrFeNi alloy resulted in phase separation and formation of BCC + FCC phases from a single-phase parent FCC cast alloy [35]. The overall corrosion resistance of the alloy was lower when compared to SS304 steel. Corrosion morphology on the single-phase FCC phase alloy comprised mostly of hemispherical pits nucleating randomly on the surface. This is an indicator of no preferred pit initiation site, while the hemispherical morphology indicates an equal propensity for pit to propagate into the material (Figure 8a). No dendritic coring or secondary pitting was seen. However, in contrast, the alloy with FCC + BCC phases showed preferred pitting along the interphase boundary. This is a clear indication of galvanic coupling between the two phases, governed by composition difference and the partitioning of elements between the two phases. Figure 8b shows the random pitting morphology preferentially occurring along the grain boundaries of the two phases.

**Figure 8.** Pitting morphology for Al0.5CoCrFeNi alloy in (**a**) as-cast condition and (**b**) after heat treated at 800 ◦C. Heat treatment resulted in phase separation forming BCC phase. Corrosion was observed to preferentially nucleate along the interphase boundaries [35] (reprinted with permission from Elsevier).

In addition to intrinsic chemistry and crystal structure, melt-solidification history was found to affect corrosion resistance of CCAs [36]. Understanding the effect of re-melting on the microstructure and consequent corrosion properties can help in casting alloys with superior chemical homogeneity and properties. A multicomponent AlCoCrFeNiTi alloy was prepared by induction melting in Ar atmosphere. The alloy was subsequently re-melted several times in order to homogenize the distribution of elements. The re-melting may have eliminated macro-segregation arising from

incomplete melting of elemental metal chunks used in alloy making. However, this may not have completely eliminated micro-segregation in the form of coring. This was evident in the form of a dendritic microstructure with micro-segregation in the inter-dendritic areas. The dendrites were rich in Al, Co, Ni, and Ti, while the inter-dendritic regions contained a higher fraction of Fe and Cr. The microstructure consisted of BCC phases along with complex intermetallics such as AlFe3. Despite the complex microstructure, the corrosion performance of the alloy was better than SS410 alloy. The addition of Ti improved the corrosion rate of the alloy (0.0216 mm/year) by almost a factor of four compared to the Ti-free alloy (0.08 mm/year). This improvement may have resulted from the complex surface oxides that promote strong passivation. Re-melting the alloy homogenized the microstructure by removing macro-segregation, which contributed to improved corrosion resistance.

With increasing interest in the additive manufacture of these complex alloys, the first step is to be able to process the alloys using power-technology. The corrosion behavior of AlCoCrFeNi alloy system was studied as a function of Cu addition via the powder metallurgy route. The corrosion properties were evaluated in 1 mol/L NaCl [37]. The alloy produced using powder metallurgy route showed a microstructure that is similar to the conventional casting route. Microstructure of the AlCoCrFeNi–Cu alloy after sintering is shown in Figure 9a. The microstructure was complex and showed a two phase-mixture of FCC + BCC phases. The potentiodynamic polarization curve for this alloy is shown in Figure 9b. The corrosion potential was −0.012 V and the corrosion current density was 3.23 nA/cm2. The alloy did not show any pitting up to potentials as high as 1.5 V versus saturated calomel electrode (SCE). This observation is insightful since powder-technology route was observed to possess improved corrosion resistance compared to conventional melt route. This may have resulted from the fact that powder particles individually possess oxide on the surface that are retained when the compacted is sintered. The larger surface oxide present on the bulk of the material may have imparted a nobler corrosion resistance as compared to its fused counter parts. The highly symmetric pattern seen in Figure 9a may have its origins in the parent (oxide covered) powder particles that explains the improved corrosion resistance that is seen in Figure 9b.

**Figure 9.** (**a**) Microstructure of AlCoCrFeNi-Cu alloy obtained by powder-metallurgy route; (**b**) Potentiodynamic polarization curve in 1 mol/L NaCl showing corrosion potential close to 0 V with respect to saturated calomel electrode [37] (reprinted with permission from Elsevier).

Increasing B content led to the precipitation of boride containing phases as seen in Figure 10a–d [38]. The corrosion behavior of the alloys was tested in1NH2SO4 [39]. The corrosion current density of the alloy increased from 787 μAmps/cm<sup>2</sup> to 2848 μAmps/cm<sup>2</sup> with the increase in boron content and boride phase fraction. The increasing boride phase fraction promoted the formation of "stringy precipitates" rich in Cr, Fe, and Co borides. This difference in composition led to the formation of local micro galvanic couples, making them susceptible to corrosion, as seen in Figure 10e–h. The phases rich in strongly passivating elements (Cr, Co) showed high pitting resistance while the matrix and inter-dendritic regions preferentially corroded. The morphology of the secondary phase changed with progressively increasing B content. Increasing B promoted stronger phase separation and formation of anodic regions that corroded more aggressively. The precipitation of hard boride phases may improve other surface properties, such as hardness and wear resistance (as discussed in subsequent sections), but certainly deteriorated the corrosion resistance due to galvanic corrosion. A balance of mechanical degradation resistance and galvanic corrosion resistance must be achieved by properly tailoring the composition to suite the application requirements.

**Figure 10.** Microstructure of the as-cast alloys with increasing Boron content: (**a**) boron free alloy; (**b**) Boron = 0.2; (**c**) Boron = 0.6; and (**d**) Boron = 1.0. Corroded features on respective microstructures. Preferred corrosion of (**e**) inter-dendritic phase (**f**) of stringer features (**g**) and (**h**) borides [38,39] (reprinted with permission Journal of the Electrochemical Society and Springer).

The corrosion behavior resulting from addition of Cr and Ti to the base composition of AlCoCuFeNi has been reported and corresponding microstructures are shown in Figure 11a–d [40]. The bright phase in the images is Cu-rich FCC, whereas the BCC phases are rich in Al and Ni and they show a darker contrast. Adding Cr resulted in the formation of dendrites, whereas BCC formed into lamellar Widmanstatten type structures. Ti caused Al-, Co-, Ni-, and Ti-rich BCC phases (A2/B2), whereas adding both Cr and Ti refined the grain structure and led to copper segregation. The corrosion behavior of the alloys studied in 0.5 mol/L H2SO4 solution showed very low corrosion current densities—in the range of 5–8 μA/cm2 for the Cr containing alloy. In contrast, the Ti and Ti-Cr containing alloys showed much higher corrosion activity. Cr and Ti typically improve corrosion resistance. The anomalous finding in this study may be due to the heterogeneous microstructure of the alloy that resulted in the formation of micro-anode and micro-cathode regions that accelerated corrosion. The corroded microstructures are shown in Figure 11e–g, indicating that the Cu-rich FCC phase was highly susceptible to corrosion. Corrosion of Cu-rich phases may be due to the higher galvanic character of the alloy, thus making it susceptible to dissolution.

In contrast to adding passivating elements, such as Cr and Ti, in the aforementioned study, the effect of changing corroding species, such as Cu and Al content on mechanical and corrosion behavior was studied for Al*x*CoCrCu0.5FeNi system [41]. The corrosion resistance of this alloy was studied in 0.5 M H2SO4 and 0.5 M NaCl solution. The *x* = 0.5 alloy showed single-phase FCC structure, while alloys with *x* = 1.0 and *x* = 1.5 showed a mixture of FCC and BCC phases. The microstructure of the two-phase alloys was a BCC-rich dark matrix and light inter-dendritic FCC phase, as shown in Figure 12a. Potentiodynamic polarization curves showed that the alloys with two-phase microstructure had lower corrosion resistance due to the formation of micro-galvanic couples. NaCl environment caused pitting, while the alloy showed passivation behavior in H2SO4 solution. Solutionizing heat treatment improved the corrosion resistance of the *x* = 1.5 alloy as the FCC phase dissolved leaving

behind a largely BCC alloy. The potentiograms in Figure 12b,c show the alloys' behavior in 0.5 M NaCl and 0.5 M H2SO4 solution.

**Figure 11.** As-cast (**a**–**d**) and corroded (**e**–**g**) microstructures of AlCoCuFeNi-Cr/Ti alloys. Inter-dendritic phases containing Cu were observed to corrode and dissolve rapidly. In the case of alloys containing both Cu and Ti, the lighter FCC phase dissolved leaving behind rounded dendrite features [40] (reprinted with permission from Elsevier).

**Figure 12.** (**a**) As-cast microstructure of Al1.5CoCrCu0.5FeNi CCA alloy, potentiodynamic polarization plots of Al*x*CoCrCu0.5FeNi (*x* = 0.5, 1.0, and 1.5) in (**b**) 0.5 M/L NaCl, and (**c**) 0.5 M/L H2SO4 [41]. The potentiograms show better corrosion resistance of the alloys as compared to SS321 alloy (reprinted with permission from Elsevier).

The pitting corrosion resistance of CoCrFeNiTiMo*<sup>x</sup>* was evaluated as a function of Mo content, varying from *x* = 0 to *x* = 0.8 [42,43]. The corrosion behavior was tested in acidic, basic, and saline solution. The as-cast microstructures of CoCrFeNiTiMo*<sup>x</sup>* are shown in Figure 13a–d. Increasing Mo content altered the microstructure to result in phase-partitioning—a dark dendritic phase and a bright interdendritic (ID) phase. Mo was observed be uniformly distributed between the two phases at 0.1 at%, however, increasing the Mo content partitioned in to the interdendritic (ID) regions. This may be due to the strong single-phase forming tendency of Co-Cr-Fe-Ni system, as established by various studies. While the average composition of Co, Cr, Fe, and Ti varied by a mere 2–4% between the two regions, Ni content variation between the dendritic and interdendritic phase was as high 50%, as measured by EDS. A broad observation is that the interdendritic region is rich in Mo and lean in Ni. This information in conjunction with individual binary phase diagrams of Mo and Co, Cr, Fe, Ti, (elements in the ID region), and enthalpy of mixing values suggests that the σ phase that is formed in the alloy might not be a simple Mo-Cr phase, akin to SS316L. The changing phase composition and partitioning of elements between dendritic and ID regions may have resulted in galvanic coupling and the consequent increased corrosion of Mo containing alloys as compared to Mo free/lean alloys.

**Figure 13.** Corrosion and pitting of CoCrFeNiTiMo*x* (*x* = 0, 0.1, 0.5, and 0.8) alloys (**a**–**d**) As-cast, (**e**–**h**) after corrosion in H2SO4, (**i**–**l**) after corrosion in NaCl [42] (reprinted with permission from Elsevier).

The corrosion behavior by immersion test was studied for two CCAs, namely, CrCu0.5FeMnNi and Cr0.5CuFeMnNi [44]. The as-cast microstructure of the two alloys are shown in Figure 14a,d. Both of the alloys showed dendritic microstructure with FCC or FCC + BCC solid solution phases. The corrosion behavior of the alloys was characterized by an immersion test and potentiodynamic polarization in 1 M H2SO4. The microstructures after immersion test are shown in Figure 14b,e, while that after accelerated corrosion are shown in Figure 14c,f. Both tests showed preferred corrosion of the inter-dendritic phase due to galvanic coupling from the partitioning of alloying elements. Superior corrosion resistance of the dendritic phase was explained by the passive layers of NiO, Ni(OH)2, NiSO4, and Cr2O3. The corrosion resistance of both the alloys was superior to stainless steel. Between the two alloys, the one with lower Cu content showed lesser elemental segregation and higher corrosion resistance.

Complex concentrated (or high entropy) alloys have been synthesized in the form of coatings by several processing routes, including melt cladding and deposition, chemical vapor deposition (CVD), physical vapor deposition (PVD), electro spark processing [45], direct current magnetron sputtering [46], and laser cladding techniques [47]. Surface clad CCAs showed a desirable microstructure because of rapid solidification, good metallurgical bonding to substrate, and lesser compositional segregation [48,49]. In the form of coating, CoCrFeMnNi CCA showed spontaneous passivation in NaCl solution. Although this CCA coating showed *i*corr value similar to 304SS, the passivation potential window for 304SS being wider. The CCA showed better corrosion resistance than 304SS in H2SO4, with a stable passive film formation. The initiation of corrosion for CoCrFeMnNi coating started with the depletion of chromium between the dendrites, and the subsequent weakening of the microstructure.

**Figure 14.** (**a**,**d**) Typical microstructures of Cr0.5CuFeMnNi and CrCu0.5FeMnNi system, respectively; (**b**,**e**) after immersion test in1MH2SO4 and (**c**,**f**) after polarization test in1MH2SO4 [44] (reprinted with permission from Wiley).

Addition of Ti up to a certain percentage to AlCoCrFeNi CCA coating resulted in better corrosion and cavitation erosion performance in NaCl [28]. Further increase in Ti content resulted in the formation of Ti2-Ni and NiAl intermetallic compounds and decreased the passivation resistance. The alloy with the highest Ti content showed the worst corrosion resistance. Cavitation erosion behavior is primarily dictated by mechanical strength of a material in a non-corrosive medium. Therefore, the coating with the highest Ti content showed improved cavitation erosion resistance in distilled water, because the intermetallic compounds acted as a deformation barrier on the surface. In contrast, the same alloy (highest Ti content) showed the worst cavitation resistance in NaCl solution due to the synergistic effect of cavitation and corrosion in NaCl medium. Interestingly, AlCoCrFeNi coating without Ti showed remarkable cavitation erosion resistance, better than 304 stainless steel in NaCl with a lower *i*corr value [50]. Ti addition up to a certain percentage to Al2CrCoCuFeNiTi*<sup>x</sup>* coating fabricated by laser cladding resulted in good corrosion performance in HNO3 [51]. Ti promoted the formation of a BCC phase in this CCA coating and affected both the corrosion and wear properties. Due to rapid cooling rates that were achieved during laser cladding, lesser segregation and uniformly refined grains (down to nanoscale) were reported [52]. The homogenous microstructure resulted in better corrosion resistance. Addition of Ti to AlCoCuFeNi CCA produced by arc melting resulted in a two-phase heterogeneous microstructure and micro-anode/cathode regions in the electrolyte [53]. The corrosion resistance was found to decrease for this alloy at both 298 K and 366 K in H2SO4. Laser processed AlCoCuFeNi CCA coating showed similar corrosion current densities to the coatings containing Ti [40,54]. The improvement in corrosion performance was attributed to the reduced dilution rate and formation of a compact CCA phase by controlling the laser parameters. Niobium also showed a similar effect as Ti in CCA coatings. The addition of Nb prevented less noble elements from dissolving in corrosive environments [54]. Corrosion performance of CoCrCuFeNi CCA coating increased considerably with the addition of Nb due to modification of microstructure and the formation of a very stable passive film [55]. Addition of Nb reduced the Cu segregation in interdendrite regions, resulted in the formation of a finer FCC phase and very stable surface oxide films.

Al addition up to a certain percentage showed improved corrosion current density for Al*x*CoCuFeNi CCA coating made by laser cladding [56]. A monotonic increase in corrosion resistance was reported for Al*x*CoCrFeNiTi [57]. A similar trend was observed for addition of Ni to Al2CoCrCuFeNi*x*Ti laser cladded CCA coating [58]. Increasing Ni content in this CCA coating, with *x* values up to 2, resulted in better corrosion performance in NaOH and NaCl solutions. In summary, the presence of corrosion resistant elements, such as Cr, Ti, Al, and Ni in limited quantity improved corrosion resistance in CCA coatings. However, beyond a certain mole fraction, microstructural segregation and lattice distortion led to a worsening of corrosion resistance.

The presence of Co in CCA coatings typically resulted in improved corrosion performance. In some CCA coatings, Co formed a passive film of CoO, which after exposure to corrosive medium, formed Co(OH)2 [59]. Co(OH)2 acted as a protective passivation layer and prevented corrosive species, such as Cl<sup>−</sup> and O2<sup>−</sup>, from diffusing into the coating. AlCoCrCuFe-*X*0.5 CCA coating in which *X* was Si, Mo, and Ti, showed no passivation in NaCl [60]. In addition, extensive pitting was observed for the alloy containing Mo and Ti on the Cr-depleted Fe2Mo and Fe2Ti phases, as shown in Figure 15. However, the dendritic regions enriched with Cr remained passivated after polarization tests.

**Figure 15.** (**a**) Potentiodynamic polarization plots of AlCoCrCuFe-*X*0.5 CCA coatings, SEM microgarphs of surface morphologies after polarization tests for (**b**) *X* = Cu, (**c**) *X* = Si05, (**d**) *X* = Mo05, and (**e**) *X* = Ti05 [60] (reprinted with permission from Elsevier).

The electrochemical behavior of CCA coatings has been reported to be different from their bulk counterparts with identical chemical composition [18,61,62]. Due to rapid cooling rates, CCA coatings possess more homogenous microstructure with lesser elemental segregation. In contrast, bulk as-cast HEAs typically consist of dendrites and inter-dendritic regions with different chemical compositions, resulting in micro-galvanic cells that accelerate the corrosion process. This effect was clearly demonstrated for AlCoCrFeNi CCA coating fabricated through electro-spark method. Relatively uniform corrosion was seen for the coating (Figure 16a), while a non-uniform attack was seen for the as-cast alloy (Figure 16b). The inhomogeneous corrosion of the cast CCA was attributed to the micro-galvanic coupling between the matrix precipitates and the matrix itself. However, the CCA coating was free from intercellular segregation and precipitates.

**Figure 16.** (**a**) Uniform corroded surface of AlCoCrFeNi CCA coating processed by electrospark after polarization test in NaCl, (**b**) Non-uniform corroded surface of a cast AlCoCrFeNi CCA in the same solution [62] (reprinted with permission from Springer).

Direct current magnetron sputtering has also been used for the fabrication of CCA coating with a uniform and homogenous microstructure consisting of very fine grains and low levels of segregation. Coatings fabricated via this method typically showed amorphous microstructure at initial stages of deposition, which crystallized with the increase in deposition time. AlCoCrCuFeMn CCA coating fabricated by magnetron sputtering with a thickness of 1–2 μm showed better corrosion resistance than 201 stainless steel in NaCl, NaOH, and H2SO4, with a wide passive region due to fine grains and limited segregation in the microstructure.

Overall, the corrosion resistance of several CCAs are reported to be comparable or better than stainless steels. This may be attributed to the larger fraction of constituent elements, such as Co, Cr, and Ni in the alloy that improve the pitting resistance and improve passivation. Addition of copper was found to induce phase separation and formation of galvanic couples. Vast majority of CCAs reported so far have corrosion potential between −200 to −400 mV and corrosion current density less than 2 μAmps/cm2 as summarized in Figure 17. The corrosion current density is lower than stainless steels although corrosion potentials are comparable. Another metric for evaluating the corrosion behavior is the pitting resistance (Δ*E*), measured as the difference between corrosion potential and pitting potential. The pitting resistance of several CCAs are compared with stainless steels in Figure 18. Some CCAs show two times higher pitting resistance when compared to stainless steel. Table 2 is a summary of reported CCAs, their microstructure, corrosion environment, and type of polarization along with the major finding in each case.

**Figure 17.** Corrosion current density versus corrosion potential for CCAs/high entropy alloy (HEAs) in 3.5 wt% NaCl solution [18,28,29,31,32,34,51,63–65].

The overall corrosion behavior of CCAs was observed to be dependent on three major factors. First, the composition of the alloys—this in turn affects the nature of the passivation layers, and the relative galvanic characteristic of the constituent phases; second, the environment in which corrosion is being evaluated; and third, the processing parameters. Most of the CCAs investigated showed better corrosion resistance as compared to stainless steels. This is primarily due to the high content of elements that form a passivating oxide layer. For example, SS316 has ~18% Cr, ~12% Ni, and ~2% Mo. In contrast, most of the CCAs that are made of equimolar proportions have at least 20% Cr, 20% Co, and 20% Ni, all of which provide strong passivating effect that translates into better corrosion resistance. Further, the HEA subset showed high resistance to uniform corrosion since these alloys form a single phase structure devoid of galvanic coupling. The general observation of lowering of corrosion resistance with multi-phase CCAs is in line with the galvanic series of alloys. Cu was observed to be particularly detrimental in several of these alloys since it is not only anodic with respect to the passivating elements, but also precipitated in the form of secondary phases that acted as the preferred corrosion sites. No particular relationship between the crystal structure (FCC or BCC) and corrosion resistance was observed. However, phase mixtures had lower corrosion resistance when compared to isomorphous systems. Intermetallic phases, such as borides, aluminides, and Nickel Titanates acted as corrosion initiation sites. The matrix region around the intermetallics dissolve due to anodic character that cause the particles to dislodge. Rupture of passive layer promotes rapid dissolution of the underlying alloy and associated material degradation. Similar effects were observed in Al, B, Mo, and Ti; however, the extent of deterioration varied significantly. Phase morphology was also found to play an important role. Secondary phases with needle and plate-like features dissolved more rapidly when compared to uniformly distributed equiaxed phases, likely because of unfavorable anode to cathode ratio at the tips.

The test environment and corroding species determined the electrochemical kinetics. In general, Cl− containing solutions caused more corrosion damage as compared to acidic or alkaline solutions. There are limited reports on in vitro and in vivo corrosion studies for bio-medical applications. Processing parameters affect the microstructure, which in turn affects the corrosion behavior of CCAs. Powder processing route was observed to produce more corrosion resistant alloys due to homogeneous elemental distribution, whereas lower corrosion performance was seen in alloys that were produced via melt-casting routes due to coring and segregation. There are significant knowledge gaps on the response of CCAs to welding and joining treatments and associated weld-induced sensitization.





#### **4. Erosion and Erosion Corrosion of CCAs**

There are very limited number of studies on the erosion behavior of CCAs/HEAs. The two alloy systems that have been studied the most are Al*x*CoCrFeNi and Al*x*CoCrCuFeNi. Slurry-erosion behavior of Al3CrCoFeNi laser cladded CCA was compared with conventionally used 17-7 precipitation hardened (PH) stainless steel. The Al3CrCoFeNi CCA coating showed excellent erosion resistance when compared to 17-7 PH stainless steel with seven times higher resistance at 15◦ impingement angle [71]. Maximum erosion rate for both CCA and 17-7 PH steel were observed at 45◦ impingement angle. Thereafter, the erosion rates were more or less constant with a further increase in impingement angle. Higher erosion resistance of the Al3CrCoFeNi CCA was due to high hardness (~750 HV) of the BCC phase and severe lattice strains. Significant lattice distortion was attributed to the high mole fraction of large sized Al atom. The effect of heat treatment on erosion behavior of Al3CrCoFeNi CCA coating was also investigated. Increased erosion resistance (~15% compared to untreated HEA) was seen with increase in annealing temperature with maximum corresponding to the 950 ◦C heat-treated sample. Authors attributed the enhanced erosion resistance of the CCA coating treated at 950 ◦C to increased hardness (765 HV), resulting from Cr3Ni2 precipitation and reduced roughness.

In Al*x*CoCrFeNi alloy system, decrease in Al content results in transition from pure BCC to BCC + FCC and finally to pure FCC structure [72]. Al0.1CoCrFeNi alloy shows a single-phase solid solution of face centered cubic (FCC) structure with good thermal stability. Slurry erosion behavior of Al0.1CrCoFeNi CCA was evaluated at different impingement angles (30◦ to 90◦) and a constant impact velocity (20 m/s). Despite the low hardness of 150 HV, the cast Al0.1CrCoFeNi alloy displayed erosion resistance that is comparable to or better than mild steel (of hardness 205 HV) at acute angles (Figure 19a). At normal impingement, Al0.1CrCoFeNi showed much better erosion resistance when compared to mild steel, which was attributed to the significant work hardening ability of the alloy. Continuous impact of abrasive particles during the erosion test results in significant work hardening of the HEA due to its low stacking fault energy and nano-twin formation [73]. The stacking fault energy (SFE) for Al0.1CrCoFeNi high entropy alloy is reported to be about 30 mJ/m2 [74]. However, Al0.1CoCrFeNi alloy showed lower erosion resistance when compared to stainless steel SS316L due to higher hardness and strength of the later. Correlation with different mechanical properties showed that the ultimate strength and ultimate resilience significantly affected the erosion behavior in these multi-component metallic systems.

The slurry erosion-corrosion behavior of AlCrCoCuFeNi CCA after annealing at different temperatures (600 ◦C and 1000 ◦C) was studied [27]. Both untreated and heat-treated AlCrCoCuFeNi alloy showed high erosion resistance compared to SS304 stainless steel. Untreated AlCrCoCuFeNi alloy also showed higher corrosion resistance compared to SS304 stainless steel. However, sample annealed at 600 ◦C showed significantly reduced corrosion resistance, which was attributed to precipitation of intermediate phases. Further increase in annealing temperature improved corrosion resistance from the resulting microstructural homogeneity. In contrast to corrosion studies, the combined erosion-corrosion test showed distinctly different behavior for the sample annealed at 600 ◦C, exhibiting the lowest mass loss. The improvement in erosion-corrosion resistance was predominantly due to increased hardness (~500 HV) due to the formation of ordered B2 or disordered A2 structures from the annealing process. Addition of higher Al fraction in the AlCrCoCuFeNi system was observed to increase hardness due to the formation of BCC/B2 structure and improved erosion-corrosion resistance. However, lowering the Al content improved the corrosion behavior [34]. The Al0.1CrCoFeNi HEA showed high slurry erosion-corrosion resistance (~1.8 times higher) as compared to SS316L stainless steel with significant negative synergy, as shown in Figure 19b. The negative synergy for Al0.1CrCoFeNi CCA indicates the positive contribution of corrosion in lowering the mass loss during the erosion-corrosion test. The Al0.1CrCoFeNi CCA also showed high pitting and protection potentials as compared to SS316L steel, indicating the formation of stable passive layer. Stability of the passive layer was partly attributed to the high mixing entropy resulting in high activation energy for diffusion.

Ti addition was found to enhance the corrosion and cavitation erosion resistance of complex concentrated alloys [28]. Cavitation erosion-corrosion behavior of laser cladded AlCrCoFeNiTi*<sup>x</sup>* CCA (*x* = 0.5 to 2) was evaluated. Maximum cavitation erosion resistance was observed for AlCrCoFeNiTi2 HEA. However, the trend reversed completely for the cavitation erosion-corrosion test, with AlCrCoFeNiTi2 CCA showing the least resistance. This trend reversal was explained by the formation of Ti2Ni and NiAl intermetallic compounds. These intermetallics significantly enhanced the hardness and reduced erosion rates due to increased resistance to plastic deformation. At the same time, formation of intermetallic compounds degraded the corrosion resistance due to the formation of localized galvanic cells and unstable passive layer.

The cavitation erosion-corrosion performance of AlCrCoFeNi laser cladded CCA was compared with 304 stainless steel [50]. The AlCrCoFeNi coating showed 7.6 times better cavitation erosion-corrosion resistance compared to 304 stainless steel. The better performance of the CCA was attributed to combined effect of high hardness and corrosion resistance. The high hardness resulted from the BCC solid solution and corrosion resistance from the homogeneous microstructure without any intermetallic phases. When compared to *Epit* of 96 mV observed for SS304 steel, AlCrCoFeNi laser cladded CCA showed significantly high pitting potential of 257 mV indicating higher passive layer stability of the later.

**Figure 19.** (**a**) Slurry erosion rate [26]; (**b**) slurry erosion-corrosion rate and [75]; (**c**) cumulative volume loss under cavitation erosion and erosion-corrosion of Al0.1CrCoFeNi high entropy alloy compared to stainless steel SS316L [76].

Al0.1CrCoFeNi alloy showed remarkable resistance to cavitation erosion and erosion-corrosion compared to 316L stainless steel as shown in Figure 19c [76]. Additionally, the alloy showed a much longer incubation period of 6.5 h as compared to 2.5 h for 316L stainless steel. This was attributed to comparatively greater degree of work hardening and superior corrosion resistance of Al0.1CrCoFeNi alloy. The strain hardening exponent for Al0.1CrCoFeNi (*n* = 0.77) was more the two times that of the 316L stainless steel (*n* = 0.3). Higher strain hardening increased the incubation period and lowered the erosion rates by effectively increasing the flow stresses. Superior erosion-corrosion resistance of the alloy may also be explained by its higher pitting potential (*Epit* = 490 mV) and protection potential (*Epp* = 184 mV) when compared to SS316L (*Epit* = 359.8 mV; *Epp* = −10.58 mV).

Scanning electron microscopy (SEM) images of Al0.1CrCoFeNi alloy and 316L stainless steel after slurry erosion and cavitation tests under identical conditions are shown in Figure 20 [26,76]. The SS316L steel and Al0.1CrCoFeNi CCA both showed ductile mode of erosion in slurry and cavitation erosion. In the case of slurry erosion, micro-cutting and ploughing were the prominent material removal mechanisms observed at oblique angles. For normal impingement, material removal was mainly through formation and removal of platelets (platelet mechanism). When compared to micro-cutting, the material removal for Al0.1CrCoFeNi alloy at an oblique angle was mainly through ploughing mechanism due to higher ductility as compared to SS316L steel. In addition, few micro-indents were also observed for samples that were tested at 90◦. Micro-indentation resulted in severe plastic deformation and the removal of the strained material once accumulated strain reached a critical value [77]. From the cavitation erosion-corrosion test, the formation of craters and pits were observed

as the primary damage mechanism (Figure 21). The size of the craters were significantly larger for SS316L steel, while they were virtually absent for the CCA as seen in Figure 21. High ductility and strain hardening for Al0.1CrCoFeNi alloy played an important role in limiting crack formation/propagation and material loss.

**Figure 20.** Scanning electron microscope images showing the damage mechanism due to slurry erosion of (**a**,**b**) stainless steel 316L and (**c**,**d**) Al0.1CrCoFeNi high entropy alloy at different impingement angles. SEM images of cavitation eroded (**e**) SS316L steel and (**f**) Al0.1CrCoFeNi CCA samples tested for 20 h [26,76] (reprinted with permission from Elsevier and Wiley).

**Figure 21.** Macrographs showing the (**a**) SS316L steel and (**b**) Al0.1CrCoFeNi high entropy alloy (HEA) samples after cavitation erosion-corrosion testing for 20 h.

Complex concentrated alloy coatings have also been studied for their erosion behavior [28,45,78–81]. Al*x*CrCoFeNi (*x* = 0.1 to 3) CCA coatings were developed using microwave processing on SS316L steel substrate, with microstructure consisting of intermetallic phases, as shown in Figure 22 [82]. The matrix in these coatings was composed of either FCC or BCC phases depending on the Al fraction. The average micro-hardness showed a direct correlation with Al fraction in the coatings. The average hardness for Al0.1CrCoFeNi coating was 438 HV, which was significantly higher when compared to the bulk counterpart, which is mainly due to the difference in microstructure [26]. Maximum hardness of 624 HV was obtained for Al3CrCoFeNi CCA coating. Slurry erosion studies of these coating showed a significantly higher erosion resistance at an oblique impingement angle (30◦) when compared to SS316L steel. The maximum erosion resistance was observed for the equimolar composition. However, for normal impingement all of the coatings showed an increase in erosion rate indicating brittle behavior. The increase in erosion rate at normal impingement was attributed to the brittle intermetallic σ and B2 phases. Lowest erosion rate observed for the equimolar AlCrCoFeNi

CCA coating at both acute angles and normal impingement was attributed to the combination of higher hardness resulting from the secondary phases and fracture toughness. SEM image of the slurry eroded CCA coating that was tested at 90◦shows the presence of large craters for the non-equimolar compositions as shown in Figure 22. The cracks resulted in disintegration of the secondary phase, which is most prominent for Al3CrCoFeNi. In contrast, the equimolar composition showed lesser tendency for brittle fracture.

**Figure 22.** Microwave processed coatings of Al*x*CrCoFeNi high entropy alloys (*x*= 0.1 to 3). (**a**–**c**) cellular microstructure of the synthesized coatings with dendrite region (DR) and inter dendrite region (ID); and, (**d**–**f**) microstructures after slurry erosion test at normal impingement angle (90◦) for Al0.1CrCoFeNi, AlCrCoFeNi and Al3CrCoFeNi high entropy alloys [82].

A comparative analysis of cavitation erosion and erosion-corrosion behavior of CCAs with respect to conventional structural materials [28,50,83–87] is shown in Figure 23 in terms of the mean depth erosion rates (MDER). For both of the test conditions, CCAs show much lower MDER compared to conventional alloys, such as stainless steels. Therefore, the use of CCAs in applications demanding high erosion and erosion-corrosion resistance can effectively improve the durability and service life of the susceptible components.

The existing literature provides a fair understanding on the erosion and erosion-corrosion behavior of Al*x*CrCoFeNi-*X* high-entropy alloy systems. There is a large scope for investigation of erosion-corrosion behavior of alloys for several other compositions with high hardness and corrosion resistance. The alloys in the Al*x*CrCoFeNi-*X* system tend to have higher erosion-corrosion resistance as compared to stainless steels, which may be due to the higher content of passivating elements. The studies reported so far have investigated the room-temperature properties of the alloys. However, critical knowledge gaps exist for high-temperature erosion-corrosion behavior. For example, boiler tubes are susceptible to extreme conditions and high temperature erosion-corrosion. In addition, the synergistic effects in cavitation erosion and slurry erosion due to the presence of corrosive media need to be investigated for better understanding of degradation in marine environments. Biofouling behavior of HEAs is another area for future research with high impact.

**Figure 23.** Comparison of high entropy alloys with conventional structural materials for erosion and erosion-corrosion resistance. "Conventional materials" in the figure refers to a broad range of materials developed for erosion/corrosion applications such as SS304 [28,50], SS304L, mild steel, Bainitic Steel [86] and Copper Alloys [87]. "Coatings" refer to coatings on AA6061 [88] for erosion mitigation.

#### **5. Wear Behavior of CCAs**

All of the complex concentrated (high entropy) alloy systems that have been studied so far for their wear behavior are summarized in Figure 24. They are broadly classified based on the alloy chemistry and processing. In addition to wear behavior, the hardening response, phase stability, and hot hardness was reported for some alloys. The copper containing alloys were typically single phase or mixture of two simple phases. The copper free alloys were based on AlCoCrFeNi system and modified with Ti or Mo. The copper free alloys showed higher wear resistance when compared to the copper containing alloys. In addition, wear behavior of alloys that are composed of purely refractory elements have also been reported. Surface modification was done for some conventional steels and Ti alloys by LASER cladding, tungsten inert gas (TIG), and sputtering to enhance the wear properties.

**Figure 24.** Summary of complex concentrated alloys studied for their tribological behavior.

The wear behavior of Al*x*CoCrCuFeNi [89] alloy system has been systematically investigated by varying the Al content from *x* = 0.5 to *x* = 2. In the Al*x*CoCrCuFeNi cast alloys, there was Cu segregation in the inter-dendritic region, while other elements enriched the dendrites. At *x* = 0.5, the dendritic and inter-dendritic phases were both FCC, which changed with increasing Al content. Increasing Al content stabilized BCC phases and resulted in hardness increase by several times The Archard's wear relation was found to hold true in the case of Al*x*CoCrCuFeNi alloy system with improving wear resistance from the increased hardness.

The softer composition in Al*x*CoCrCuFeNi system with *x* = 0.5 showed ductile deformation, grooving, and disc-like wear debris, as shown in Figure 25, all of which are in line with ductile character of the FCC-rich alloy. On the other hand, the BCC-rich alloys showed smoother surface deformation, finer wear debris that were enriched with oxygen, indicating oxidative wear. The steady state friction value was the lowest for Al0.2CoCrCuFeNi, which is in line with the surface oxidation and high hardness that protected the alloy from tribo-degradation. Based on variation in hardness, ductile deformation, and steady state friction values in Al*x*CoCrCuFeNi alloy system, the Al0.5CoCrCuFeNi composition was chosen for further modification. Addition of V to Al0.5CoCrCuFeNi alloy resulted in a phase mixture consisting of FCC, BCC, and sigma phase [90]. Fixing aluminum content at 0.5 and gradually increasing V content increased the BCC phase fraction. The change in microstructure with increasing V mole fraction is shown in Figure 26. Increasing the vanadium concentration increased BCC phase fraction and hardness of the alloy, but this did not translate into improved wear resistance. There was marginal increase in wear resistance even though the hardness increased from 200 HV to a peak hardness of ~650 HV. The optimum composition range for enhanced mechanical and tribological properties was in the range of 1.0 to 1.2 mole fraction of Vanadium. Therefore, multiple competing factors affected the wear behavior, including complex microstructure, elemental segregation, BCC phase fraction, and morphology.

**Figure 25.** Deformation and wear mechanisms for Al0.5CoCrCuFeNi alloys: (**a**) Ductile deformation showing long grooves with minimal lateral cracks (**b**) wear particles showing large flakes indicating delamination (**c**) smooth surface characteristic of bcc and high hardness alloys, (**d**) smaller wear particles showing significantly higher oxidation seen from lighter shades from oxide charging [89] (reprinted with permission from Elsevier).

**Figure 26.** Microstructure of the Al0.5CoCrCuFeNi alloy at (**a**) 0.0, (**b**) 1.0 and (**c**) 2.0 mole fraction V. The initial microstructure had the lowest hardness value, which was observed to increase with increasing vanadium content. All three microstructures showed segregation of copper into inter-dendritic regions [90] (reprinted with permission from Springer).

Boron is known to be a BCC phase stabilizer in ferrous alloys. In contrast, this effect was not seen with boron addition to Al0.5CoCrCuFeNi. The alloy retained its FCC structure when B was less than 10%, while small quantities of ordered FCC phases evolved when the content was increased to ~15% [38]. In contrast to Vanadium addition, changing B content led to a significant strengthening effect. Addition of boron to the AlCoCrCuFeNi system was observed to increase the hardness and wear resistance of the alloy to 736 HV and 1.76 m/mm3, respectively. The wear resistance of the alloy was superior when compared to wear resistant SUJ bearing steels. Hardness increased from ~300 HV to 750 HV along with nearly doubling of wear resistance when boron content was increased from 0% to 15%. The changes in hardness and wear resistance with the addition of Vanadium, Boron, and Aluminum are summarized in Figure 27.

**Figure 27.** Summary of mechanical, tribological and metallurgical properties of Al0.5CoCrCuFeNi-*X* alloy system. Improved wear resistance of the alloys with increasing proportions of (**a**) Vanadium (**b**) Boron and (**c**) Aluminum; (**d**) relative comparison and summary of wear resistance of Al0.5CoCrCuFeNi-*X* alloys where *X* is V, B and Al*y*CoCrCuFeNi where *y* is 0.3, 1.0 and 1.5 [38,89,90].

The CoCrFeNi alloy system has been shown to be very versatile in terms of compositional and microstructural modifications. Addition of Ti results in the formation of Ni3Ti intermetallic compounds that impart exceptional high temperature strength. Therefore, high entropy forming CoCrFeNi composition may be significantly strengthened by these intermetallic phases. Extensive studies have been done to understand the competing effects of Al and Ti in the Al*x*Co1.5CrFeNi1.5Ti*<sup>y</sup>* system by creating a series of alloys [91]. Al-free Co1.5CrFeNi1.5Ti0.5 and Co1.5CrFeNi1.5Ti1.0 alloys were developed for isolating the effect of Ti, while Al0.2Co1.5CrFeNi1.5Ti0.5 and Al0.2Co1.5CrFeNi1.5Ti1.0 were developed to identify the effect of Al.

Addition of Ti1.0 improved the hardness by about 100 HV when compared to the Ti0.5 bearing alloys. Addition of Al caused slight drop in hardness due to suppression of (Ni,Co)3Ti formation. The overall hardness of Al0.2Co1.5CrFeNi1.5Ti1.0 improved by over 200 HV as compared to Al free counterparts. Among the four alloy systems Al0.2Co1.5CrFeNi1.5Ti1.0 showed highest wear resistance, being as high as 5500 m/mm3, which is nearly double the values on the other alloys. The tribological properties of the alloys can be compared to commercial wear resistant steels, such as SUJ2 and SKH51. As far as the mechanisms are concerned, the softer alloys displayed characteristic features of wear on ductile materials, such as delamination, grooving, and plastic deformation, as shown in Figure 28. The degradation mechanism for the high Ti content alloys was predominantly oxidative wear in the tribo-system, shallow wear tracks, and marks on the surface. These oxide layers typically protect the underlying alloy from meta-metal contact, thereby reducing adhesive wear. Although the tests were conducted at room temperature, the contact temperature during wear may rise to temperatures where softening and oxidation may be of concern. Therefore, high temperature hardness is important in these applications. Unlike other alloys derived from the CoCrFeNi base system, the AlCoCrFeNiTi alloy forms high temperature deformation resistant intermetallic particles. These particles effectively improve the alloy performance and contribute to improved tribological properties.

**Figure 28.** (**a**–**d**) Wear track morphology and wear particle morphology showing clear (**e**,**f**) delamination wear and (**g**,**h**) partial oxidation wear on the Al*x*Co1.5CrFeNi1.5Ti*<sup>y</sup>* alloys [91] (reprinted with permission from Elsevier).

Al and Ti promoted the formation of ordered phases and had a significant effect on the hardness and wear resistance of alloys. Extensive microstructure characterization was done and wear behavior was studied for the Al0.25CoCrFeNiTi0.75 alloy [92]. The Al0.25CoCrFeNiTi0.75 alloy was seen to have a Cr-Fe-Co rich phase, a Chi (χ) phase with BCC structure, a Ni2AlTi based L21 ordered phase, and FCC minor phase. A scanning electron microscopy image of the alloy along with the EDS elemental maps is shown in Figure 29. The EDS maps showed that the lighter contrast phase is chi (χ) phase, the darker contrast phase as the FCC based ordered phase, and the greyish phase as disordered FCC. This is distinct from the Al0.2Co1.5CrFeNi1.5Ti1.0 alloy that showed blocky complex η-(Ni,Co)3Ti phase with needle like morphology. The hardness of the alloy was also observed to be around

570 HV for the (L21) phase, while, lighter contrast matrix regions showed a hardness of 1090 HV. Comparatively, these values are lower than the Al0.2Co1.5CrFeNi1.5Ti1.0 that showed η phase with Widmanstätten structures having a hardness of 1200 HV, and an overall hardness of 717 HV. The wear behavior of the alloy when tested in sliding reciprocating mode showed long grooves running parallel to the wear track. The wear tracks were shallow, with little to no surface oxidation. The wear behavior of the alloy followed Archard's relation. Between the two phases, the softer dark phase composed of Al-Ti-Ni was observed to wear off preferentially. The size of the wear tracks increased with increasing test load. These features are shown in Figure 30. The Al0.2Co1.5CrFeNi1.5Ti1.0 was reported to have improved wear resistance than commercial high hardness wear resistant steels, such as SUJ2 and SKH51 (~65 HRC), whereas the wear performance of Al0.25CoCrFeNiTi0.75 CCA was more comparable to SS 440C (~55 HRC).

**Figure 29.** Scanning electron microscope images of Al0.25CoCrFeNiTi0.75 alloy in as-cast condition. (**a**) Low magnification scanning electron image of the alloy; (**b**) high magnification image showing the three distinct contrasts - lighter contrast phase is chi (χ) phase, the darker contrast phase as the FCC based ordered phase, and the greyish phase as disordered FCC; EDS elemental maps showing distribution of (**c**) Co (**d**) Ti (**e**) Ni (**f**) Al (**g**) Fe and (**h**) Cr [92] (reprinted with permission from Elsevier).

**Figure 30.** Wear behavior of Al0.2Co1.5CrFeNi1.5Ti1.0 alloy. The alloy showed increased wear volume loss with increased load, corresponding well with archard's law. High magnification SEM images, and EDS maps shows oxides on the surface, indicating oxidative wear operating on the sample surface [92] (reprinted with permission from Elsevier).

The wear behavior of high entropy alloys in marine conditions was evaluated for CoCrFeNiMn and Al0.1CoCrFeNi alloys. Here, isolating the effects of wear and corrosion that are acting simultaneously on the sample are important. Such problems can be approached by individually assessing the wear in dry condition, corrosion using electrochemial or immersion tests, and comparing the results to marine wear tests. A weighted summation of material loss from each of the wear tests would reveal the predominant mechanism of material loss between the competing mechanisms. In the case of CoCrFeNiMn and Al0.1CoCrFeNi alloys, the synergy between wear and corrosion was observed to be negative—implying that corrosion did not aggravate material loss in the alloys. Material loss during marine wear was lower than the summation of dry wear loss and material loss from corrosion [31]. The wear tracks imaged using white light interferometry (Figure 31a,b for dry and Figure 30g,h for wet conditions) show higher wear volume loss on the CoCrFEMnNi alloy in dry and marine condition. High magnification images confirm micro-grooving on CoCrFeNiMn, while spalling and fatigue wear on the Al0.1CoCrFeNi alloy. The Al0.1CoCrFeNi alloy showed lower wear loss in both dry and wet test conditions (Figure 31c,d for dry and Figure 31i,j for marine condition). Parallel grooves in the wear track indicate two body or three body wear to be operative in the alloys.

**Figure 31.** Interferometry images of dry wear track for (**a**) CoCrFeMnNi and (**b**) Al0.1CoCrFeNi. Scanning Electron Microscope images of wear track on (**c**) CoCrFeMnNi and (**d**) Al0.1CoCrFeNi; Higher magnification images of wear tracks for (**e**) CoCrFeMnNi showing coarse microabrasion/microcutting and (**f**) Al0.1CoCrFeNi showing finer microabrasion, deformation and delamination of oxide layer; The corresponding interferometry images of wear tracks generated during marine wear for (**g**) CoCrFeMnNi and (**h**) Al0.1CoCrFeNi. Scanning Electron Microscope images of wear track due to corrosive wear on (**i**) CoCrFeMnNi and (**j**) Al0.1CoCrFeNi; Higher magnification images of wear tracks for (**k**) CoCrFeMnNi showing fine microabrasion/microcutting and shallow deformation and (**l**) Al0.1CoCrFeNi showing corrosive wear in terms of break-down of surface passive layers [31] (reprinted with permission from Elsevier).

Wear resistance of the alloys that were surveyed in this study have been plotted against the reported hardness values in Figure 32. The crystal structure of the alloys have also been marked. It can be seen that irrespective of chemistry, FCC alloys typically are softer and have lower wear resistance, followed by two-phase alloys, and highest wear resistance was seen for BCC alloys. This trend changes in case of high wear resistance materials, as shown in Figure 32b. The two-phase alloys showed significantly higher hardness and wear resistance than the BCC alloys. However, there was no significant correlation for the BCC and two-phase alloys. For example, the Al0.5CoCrCuFeNiV1.6 alloy has a hardness of 600 HV with a wear resistance of 1.1 m/mm3, whereas the AlCoCrFe1.5MoNi has a wear resistance of over 1000 m/mm3 with similar hardness values, although both of the alloys have BCC crystal structure. This implies that a materials response to sliding wear is governed by factors more than just hardness and crystal structure. These critical knowledge gaps need to be addressed in future studies. The effect of lattice distortion from complex compositions on the hardness, wear resistance and friction evolution need further investigations.

**Figure 32.** Hardness-wear resistance relationship classified with respect to their magnitude of wear resistance (**a**) less than 2.5 m/mm3; (**b**) between 500 and 6000 m/mm3. The wear performance of the alloys do not show a particular dependency on the crystal structure.

In contrast to corrosion behavior, there are limited reports on wear behavior of CCA coatings. Surface cladding was typically employed for wear performance enhancement. AlCoCrNiW and AlCoCrNiSi CCA cladded layers through gas tungsten arc welding (GTAW) showed enhanced wear resistance than AISI 1050 medium carbon steel [93]. The superior wear performance of these coatings was attributed to the strong mechanical interlocking between the dense dendrites and the matrix. During the wear test, dense dendrites can strengthen the structure and prevent plastic flow. The wear performance of AlCoCrNiW layer exceeded that of AlCoCrNiSi due to stronger mechanical interlocking and a more complex microstructure. Mechanical and wear behavior of Al0.5CoCrFe2MoNiSi CCA coating fabricated by the GTAW method were reported as a function of silicon addition [94]. Superior wear resistance of this cladding layer was attributed to the formation of strong bonds between Si and the other elements in dendritic region and nanoscale precipitation in the inter-dendritic region.

Tungsten inert gas (TIG) was also used to produce CoCrFeMnNbNi CCA coating and showed much lower wear loss when compared with AISI 304 steel due to presence of a FCC Nb-rich Laves phase with nanoscale lamellar spacing [95]. The Laves phase resisted damage during sliding and protected the coating surface from plastic deformation. CuNiSiTiZr CCA coating fabricated by vacuum arc melting showed almost 2.5 times higher hardness than TC11 (typically used in aerospace industry), and superior wear resistance due to various effects, such as solid solution strengthening, precipitation strengthening, and nanocomposite strengthening [96]. A similar phenomenon was observed when AlCrSiTiV CCA coating was deposited on Ti-6Al-4V substrate via the laser cladding technique. CCA coating showed improved wear rate as compared with the Ti-6Al-4V substrate as well as higher hardness values (Figure 33).

**Figure 33.** (**a**) SEM micrograph of AlCrSiTiV CCA coating on Ti-6Al-4V substrate and (**b**) specific wear rate of two materials after dry sliding wear test under various frequencies [97] (reprinted with permission from Elsevier).

The high wear resistance of AlCrSiTiV CCA coating was attributed to the formation of hard intermetallic phase in a relatively ductile BCC matrix. The softer matrix limited brittle crack propagation during abrasive and adhesive wear. Plasma-spray has also been used to synthesize several CCA coatings, including AlCo0.6Cr*y*Fe0.2Ni*x*SiTi0.2, AlCoCrCuFeNi, AlCoCrFeNi, CoCrFeMnNi, and AlCoCrFeNiTi [98]. The effect of temperature on the wear behavior of AlCoCrFeNiTi CCA coating has also been reported. Adhesive wear with minor abrasion was the main mechanism for AlCoCrFeNiTi coating wear at 25 and 500 ◦C [98]. More severe adhesive wear was observed at higher temperatures due to the decrease in hardness of the coating. At temperatures higher than 500 ◦C, both the morphology of wear track and mechanism changed due to oxidation processes and the softening of the coating. The main wear mechanism at 700 ◦C was tribo-oxidation wear with wide grooves on the wear surface. The presence of micro-cracks and pores in the coating facilitated the diffusion of oxygen at higher temperatures and it caused more accelerated tribo-oxidation wear. AlCoCrFeNiTi CCA coating showed a volume wear loss of about one-ninth of 316 stainless steel at 700 ◦C [12].

Thermally sprayed AlCoCrFeMo0.5NiSiTi and AlCrFeMo0.5NiSiTi CCA coatings, consisting of a BCC dendrite and a FCC inter-dendritic microstructure, exhibited very good wear resistance [12,99]. Annealed CCA coatings showed even better wear resistance with minimized weight loss. The hardness value for these coatings was lower than bearing SUJ2 and hot-die tool steel SKD61. However, their wear resistance after annealing at 800 ◦C was much better than these wrought steels. Laser alloying has also been used for synthesizing thicker HEA coatings with strong metallurgical bonding to the substrate [48,97]. AlCoCrCuFe CCA coating fabricated by laser surface alloying showed much better specific wear rate and lower coefficient of friction (COF) than that of Q235 steel substrate, as well as three times higher hardness. It was shown that the addition of certain elements, such as Boron to AlB*x*CoCrFeNi coating fabricated by laser cladding, changes the wear mechanism from adhesive to abrasive wear, as the hardness increased. A summary of wear resistance and hardness is shown in Table 3.


**Table 3.** Sumamry of wear behavior of alloys reported in literature.

#### **6. Conclusions**

The number of complex concentrated (or high entropy) alloy systems being reported in recent years has exploded because of their tunable microstructures and desirable properties. Improving the surface degradation characteristics of these alloys will make them very attractive in wide ranging commercial applications. Some of the main corrosion, erosion, and wear characteristics in these emerging materials are summarized below:


#### **7. Future Opportunities and Outlook**

Complex concentrated alloys present a plethora of opportunities for the development of next generation materials. The scope is not just limited to bulk materials and melt-deposition coatings but also in the form of thin films and powder-metallurgy products. However, critical knowledge gaps in surface degradation mechanisms need to be assessed prior to determining the true application worthiness of these alloys. The effect of processing on surface degradation mechanisms are not well understood for welding and joining, severe-plastic deformation, and hot working. Another area with very limited number of studies includes extreme environments, such as molten/fused salts, heavy ion/neutron irradiation, and high temperatures.

Understanding the nature and chemistry of the surface passivation layer is critically important for corrosion, erosion, and wear applications. This has not been done in a comprehensive way. Surface characterization using X-ray photoelectron spectroscopy (XPS) and ultra-violet photoemission spectroscopy (UPS) could reveal valuable information about the chemistry and electronic structure of the surface passivation layer and help in fundamental understanding of the underlying mechanisms. This knowledge may be utilized to develop specific surface treatments to produce strongly passivating and non-porous oxides that can offer exceptional surface degradation resistance. Slight changes in composition (micro-alloying) or processing conditions have shown large variations in properties, compounding the complexity in analyzing these multi-component systems. In complex precipitation hardened CCAs, composition fluctuations at multiple length-scales (atomic, nano, micro) may lead to "local" effects that nucleate the breakdown of passivation layers. These effects may be captured by phase-specific corrosion and wear tests at the microstructural length-scales, including scanning electrochemical microscopy (SECM) and nano-scratch/wear tests.

There are very limited reports on the lubricity and friction behavior of complex concentrated alloys. This may be of interest in tribology for developing super-lubricity complex composition coatings. Phase-specific friction studies will provide significant insights into surface degradation from multi-body wear in multi-phase CCAs [101]. Ni free complex alloys might be attractive for biomedical applications. Evaluating in-vitro wear behavior and quantifying cytotoxicity of the wear products will significantly help in developing new biomaterials. Recently reported refractory CCAs may be attractive for highly stressed bearing applications, where high temperature wear behavior, evolution of oxide layers, kinetics of spalling, and pesting are of significant fundamental interest.

Advanced additive manufacturing, LASER melt-deposition, and combinatorial development using powder bed and powder feed techniques for CCAs hold tremendous potential towards meeting long-standing challenges like highly corrosion resistant surface materials and thermal barrier coatings. CCA claddings via melt-deposition is yet to be explored. Electrodeposition of CCA/HEA coatings via co-deposition and auto-catalytic reactions could dramatically enhance functional applications of these alloys. CCA thin films could be potentially transformative as diffusion barriers in integrated circuit (IC) manufacturing because of extremely sluggish diffusion and lattice distortion. The future opportunities and outlook for complex concentrated alloys in different areas are summarized in Figure 34.

**Figure 34.** Future opportunities and outlook for complex concentrated alloys.

**Author Contributions:** A.A. and S.M. conceived and designed the layout of the review paper; A.A. analyzed corrosion and wear sections; V.H. analyzed surface and coatings techniques; H.S.G. and H.A. analyzed erosion, corrosion and erosion behaviors. A.A. and S.M. wrote the manuscript.

**Funding:** This research received no external funding.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Article* **Corrosion Behavior of Selectively Laser Melted CoCrFeMnNi High Entropy Alloy**

**Jie Ren 1,**†**, Chaitanya Mahajan 2,**†**, Liang Liu 1, David Follette 3, Wen Chen 1,\* and Sundeep Mukherjee 2,\***


Received: 21 August 2019; Accepted: 19 September 2019; Published: 23 September 2019

**Abstract:** CoCrFeMnNi high entropy alloys (HEAs) were additively manufactured (AM) by laser powder bed fusion and their corrosion resistance in 3.5 wt% NaCl solution was studied by potentiodynamic polarization and electrochemical impedance spectroscopy tests. A systematic study of AM CoCrFeMnNi HEAs' porosity under a wide range of laser processing parameters was conducted and a processing map was constructed to identify the optimal laser processing window for CoCrFeMnNi HEAs. The near fully dense AM CoCrFeMnNi HEAs exhibit a unique non-equilibrium microstructure consisting of tortuous grain boundaries, sub-grain cellular structures, columnar dendrites, associated with some processing defects such as micro-pores. Compared with conventional as-cast counterpart, the AM CoCrFeMnNi HEAs showed higher pitting resistance (Δ*E*) and greater polarization resistance (*R*p). The superior corrosion resistance of AM CoCrFeMnNi HEAs may be attributed to the homogeneous elemental distribution and lower density of micro-pores. Our study widens the toolbox to manufacture HEAs with exceptional corrosion resistance by additive manufacturing.

**Keywords:** CoCrFeMnNi high entropy alloys; additive manufacturing; corrosion behavior; non-equilibrium microstructure; micro-pores

#### **1. Introduction**

In recent years, high entropy alloys (HEAs) have received remarkable attention from both academia and industry due to a portfolio of unusual properties including high specific strength [1–6], high fracture resistance [7,8], and excellent corrosion and oxidation resistance [9,10]. The superior corrosion behavior has been attributed to the locally disordered chemical environment obtained from random arrangement of multi-principal elements in solid solution [9,11]. Because of the formation of protective passive films on the surface, the corrosion behavior of Cr-, Ni-, and Mo-based HEAs have been widely investigated [12–16]. Despite the technical potential, processing of HEAs is generally a challenge. Most HEAs suffer from inferior flowability or formability in the liquid state to be shaped into useful components by conventional manufacturing routes such as casting [17]. Recently, the rapid development of additive manufacturing (AM) has enabled precise and versatile manufacturing of geometrically complex components necessary for practical applications [18]. Here, we apply a laser powder bed fusion based AM technique, namely selective laser meting (SLM), to produce CoCrFeMnNi HEAs. Unlike the conventional manufacturing technique, SLM prints materials and components

directly from a computer-aided design file and offers unique advantages of design freedom for three-dimensional complex geometry. The highly localized melting, strong temperature gradient, and high cooling rate (~10<sup>6</sup> K/s) [19] during laser melting often give rise to unique ultra-fine and non-equilibrium microstructures that deliver superior mechanical properties such as high strength and high ductility that are not accessible via conventional methods [20,21]. For example, AM CoCrFeNiMn HEA demonstrated a yield strength of 510 ± 10 MPa with a uniform elongation of 32.4%, in contrast to 205 ± 5 MPa and 50.2% for the as-cast counterpart [21].

Corrosion resistance is often an additional core performance factor for structural metals. Fundamental understanding of the microstructure-electrochemical property is therefore critical to the AM technology. The corrosion resistance of AM 316L stainless steels [22–26] and Ti-6Al-4V alloys [27,28] have so far been widely studied. Contradictory results have been reported for the corrosion resistance of AM 316L stainless steels [22–26]. For example, Sun et al. [22] and Zi ˛etala et al. [23] showed that the corrosion behavior of AM 316L stainless steels were similar to that of conventionally manufactured counterparts and the AM samples were more susceptible to pitting corrosion. Trelewicz et al. [24] and Geenen et al. [25] reported reduced corrosion resistance of AM 316L stainless steels. In contrary, Kazemipour et al. [26] reported superior pitting resistance and reduced metastable pitting rate in SLM-fabricated samples compared to wrought alloy. While for Ti-6Al-4V alloy, the constituent phase is the dominant factor for corrosion resistance, followed by grain size and morphology. Therefore, the SLM alloy with more acicular α and less β–Ti phase exhibited worse corrosion resistance than Grade 5 alloy [27,28]. However, for CoCrFeMnNi, one of the most notable HEAs till date, corrosion studies have been mainly focused on as-cast bulk samples [29,30] and laser cladded coatings [31], for which elemental segregation makes Cr-depleted inter-dendrites especially vulnerable to pitting corrosion. Therefore, there is limited understanding of the effect of highly non-equilibrium microstructure obtained in AM on the corrosion resistance of this alloy.

In this study, the corrosion behavior of AM CoCrFeMnNi HEA was compared with the as-cast counterpart in 3.5 wt% NaCl solution by potentiodynamic polarization and electrochemical impedance spectroscopy measurements. The microstructures and surface morphologies of both alloys were analyzed to explain the difference in corrosion resistance. Our study reveals that the AM CoCrFeMnNi HEA with unique non-equilibrium microstructure, homogeneous elemental distribution, and smaller defect (micro-pore) density exhibit superior corrosion resistance to the as-cast counterpart.

#### **2. Materials and Methods**

The SLM as-printed samples were manufactured by M290 (EOS GmbH, Munich, Germany), which is equipped with a Yb-fiber laser with a maximum power of 400 W and a focus diameter of 100 μm. Gas atomized equiatomic CoCrFeMnNi HEA powders with the particle size ranging from 15 to 53 μm (D10 = 22.8 μm, D50 = 36.8 μm, D90 = 58.4 μm) were used. As shown in Figure 1, 36 cubes of 8 × 8 <sup>×</sup> 6 mm<sup>3</sup> were printed with different processing parameters by bi-directional and chessboard scan strategies at a constant layer thickness of 0.04 mm and a hatching distance of 0.08 mm with varying laser powers (250–370 W) and scan speeds (500–2500 m/s). To reduce the anisotropy of mechanical properties, laser scan direction was rotated alternately 90◦ for successive layers. The as-cast counterpart with the same composition was prepared by arc melting constituent elements with 99.99% purity under argon atmosphere. The ingots were re-melted at least five times to assure chemical homogeneity and subsequently sucked into a water-cooled copper mold.

The relative density of the as-printed samples was measured by gas pycnometer (AccuPyc II 1340, Micromeritics, Norcross, GA, USA). Optical microscope (OM, BX53M, Olympus, Tokyo, Japan) and scanning electron microscopy (SEM, Magellan 400 XHR, FEI, Hillsboro, OR, USA) were used to examine the microstructures of the samples. The metallurgical samples were polished with SiC abrasive papers with different grits of 400, 800, and 1200 respectively, and finally with 1 μm diamond suspension. To reveal the melt pool boundaries and sub-microstructures, the samples were etched in a mixture of 50% aqua regia and 50% ethanol (vol.%).

**Figure 1.** (**a**) As-printed samples with different processing parameters; (**b**) Schematics of scan strategies in selective laser melting process.

Prior to the corrosion experiments, samples were cleaned with deionized water, acetone, and then dried in air. The specimen surface was masked with a tape exposing 0.283 cm<sup>2</sup> area to 3.5 wt.% NaCl electrolyte solution. Electrochemical studies were performed with a potentiostat (Reference 3000, Gamry, Warminster, PA, USA) at room temperature. A graphite rod with an area of 4.9 cm<sup>2</sup> was used as the counter electrode and saturated calomel electrode (SCE) was used as the reference electrode. Electrochemical impedance spectroscopy (EIS) was performed after the open circuit potential (OCP) stabilized after ~6000 s. EIS was started with an AC voltage amplitude of 10 mV and frequency was swept from 10 mHz to 100 kHz. The impedance data was interpreted using Gamry software and equivalent circuits followed by cyclic polarization tests which were performed starting from −0.25 mV versus OCP to the upper threshold limit of 10 mA·cm−<sup>2</sup> for the reverse scan. The scan rate for both forward and reverse scan was 0.25 mV·s<sup>−</sup>1. The corrosion rate (*CR*) was calculated by [30,32]:

$$CR = \frac{I\_{\text{corr}} \times K \times EW}{\rho} \tag{1}$$

where, *I*corr is the corrosion current density, *K* is a constant for use in Faraday's penetration rate equation which equals 3.27 <sup>×</sup> 10−<sup>3</sup> mm·g/(μA·cm·year), <sup>ρ</sup> is the density, and *EW* is the equivalent weight which can be regarded as the mass of alloy in grams that will be oxidized by the passage of one Faraday (96489 ± 2 C) of electric charge:

$$EW = \frac{1}{\sum \frac{n\_i f\_i}{W\_i}} \,\tag{2}$$

where *fi* is the mass fraction of the *i*th element in the alloy, *Wi* is the atomic weight of the *i*th element in the alloy, and *ni* is the valence of the *i*th element of the alloy.

#### **3. Results and Discussion**

#### *3.1. Porosity and Relative Density*

With the bulk density of 8.05 g/cm3, the relationship of relative density of AM CoCrFeMnNi samples versus volume energy density (VED) is shown in Figure 2. When VED is in the range of 62.5–115.6 J/mm3, the relative densities of AM CoCrFeMnNi samples were all above 99.5%. When the VED was further increased, the relative density decreased dramatically due to keyhole effect [33] and partial evaporation of some constituent elements [34], particularly Mn which has higher vapor pressure and lower melting point compared to the other elements [35]. Chemical evaporation and over-penetration of laser beam in the high VED regime (>115.6 J/mm3) can cause entrapped gas and formation of micro-pores inside the melt pool. In contrast, at low VED (<62.5 J/mm3), inadequate penetration of the melt pool into the previously deposited layers may leave voids in the sample [36]. In this study, the samples with relative density of 99.5% (VED of 77.08 J/mm3, laser power of 370 W, scan speed of 1500 mm/s, and bi-directional scan strategy) were selected for microstructural characterization and electrochemical corrosion tests.

**Figure 2.** Variation of relative density of additively manufactured (AM) CoCrFeMnNi samples with different volume energy densities.

#### *3.2. Microstructural Characterization*

As shown in Figure 3, the gas-atomized CoCrFeMnNi powders are composed of irregular needle-like dendrites after etching. Some internal micro-pores were also observed, which is one root of the micro-pores (1–18 μm) in the eventual printed samples (Figure 4a,b). After etching, more pores were observed especially at the melt pool boundary (Figure 4c,d), which indicates the microstructure near the melt pool boundary has weaker chemical resistance. The purple arrows in Figure 4c display the epitaxial growth of the elongated grain with tortuous grain boundaries. Melt pools with the width and depth of about 165 μm and 124 μm, respectively, were generated during the rapid solidification.

**Figure 3.** (**a**) Optical microscopic (OM) cross-section image of gas-atomized CoCrFeMnNi powders with pores inside; (**b**) SEM cross-section image of a single gas-atomized CoCrFeMnNi powder with needle-like dendrites (after etching) and internal pores. Micro-pores are highlighted by red circles.

**Figure 4.** Non-equilibrium microstructure of AM CoCrFeMnNi as-cast sample: (**a,b**) OM images showing the porosity profile of the side and top surfaces of the AM sample; (**c**) OM image showing the microstructures of the side surface of the AM sample. The laser melting layers and elongated grains are depicted by white lines and purple arrows, respectively; (**d**) high magnification OM image showing the heterogeneous microstructures of the side surface of the AM sample. The melt pool boundaries, micro-pores, columnar and equiaxed cellular sub-structures are represented by orange dash lines, blue circles, red and white arrows, respectively; (**e**) Secondary electrons (SE) mode SEM image showing the melt pool boundaries as well as columnar and equiaxed cellular sub-structures; (**f**) high magnification SEM image of equiaxed cellular structures.

As displayed in Figure 4d,e, the AM sample is primarily composed of columnar dendrites and equiaxed cellular sub-structures. The columnar sub-grain structures in Figure 4d display epitaxial growth across the melt pool centerlines and the boundary. The average size of the equiaxed cellular structure is in the range of 0.69–0.91 μm. Compared to the as-cast sample, the refinement of the microstructure was attributed to the high cooling rate (~10<sup>6</sup> K/s) [19] and the re-melting of the previous deposited layer. Based on solidification theory [37], the metallurgical morphology and cooling rate are dependent on the temperature gradient (G) and solid/liquid interface growth rate (R). As shown in Figure 4d,e, the cellular sub-grain structures mainly formed near the melt pool boundary where the temperature gradient (G) is higher and the solid/liquid interface growth rate (R) is relative lower. While at the melt pool center where the G/R ratio is lower than the boundary, the columnar sub-structures were observed.

#### *3.3. Electrochemical Corrosion Behavior*

The open circuit potential (OCP) for the two alloys are shown in Figure 5a. For both the alloys, the potential increased to positive (more noble) values indicating the formation of passive film in 3.5 wt.% NaCl solution [38]. Both the alloys showed multiple small spikes in open circuit potential, suggesting breakdown and re-passivation prior to the stabilization of the OCP after ~6000 s. The OCP value was very similar for the alloys prepared by the two different routes, indicating overall identical chemical composition. The passivation range and the corrosion current density were extracted from cyclic polarization tests (Figure 5b). The corrosion potential *(E*corr) for AM and as-cast HEAs were −189 mV and −179 mV, respectively. The values of corrosion current density (*I*corr) were similar for both alloys. In both samples, several spikes in current density were observed indicating metastable pitting followed by re-passivation. Note that the AM HEA showed a wider passive range (Δ*E*resistance = *E*pit − *E*corr) of ~386 mV compared to the as-cast HEA (~200 mV). Figure 5b shows that both samples did not re-passivate, which suggests that no protective passive film was formed at the large active pits.

**Figure 5.** (**a**) Open-circuit potential (*E*OCP) with respect to saturated calomel electrode (*E*SCE) as a function of time; (**b**) cyclic polarization plots at a scan rate of 0.25mV/s; (**c**) Nyquist plot; and (**d**) the equivalent circuit used for fitting is shown, where R.E. is the reference electrode, *R*s is the solution resistance, *R*p is the polarization resistance, *C* is the constant phase element and W.E. is the working electrode.

The Nyquist plots depicted in Figure 5c was obtained from electrochemical impedance spectroscopy (EIS) in steady-state condition in which the capacitive arc trend corresponds to the double layer and

passive film formation on the surface. The radius of the arc is directly proportional to the polarization resistance of the material. A modified Randles circuit (Figure 5d) was fit to experimental data to calculate the electrochemical parameters which includes solution resistance (*R*s), a double layer capacitor (*C*) to represent the double layer charge capacitance absorbed onto the sample surface, and polarization resistance (*R*p). All the parameters are summarized in Table 1. The simulation showed that the polarization resistance of AM HEA alloy was higher than the as-cast alloy, while the solution resistance and double layer capacitance values were approximately similar.


**Table 1.** Electrochemical and equivalent circuit parameters for AM and as-cast CoCrFeMnNi alloys in 3.5 wt% NaCl solution.

#### *3.4. Morphology of Corroded Surfaces*

SEM micrographs of the alloys before and after corrosion tests are shown in Figure 6. The AM alloy exhibited tortuous grain boundaries and micro-pores (Figure 6a). The as-cast alloy showed large grains of the order of 50–100 μm in diameter and high density of micro-pores randomly distributed in the microstructure (Figure 6c). Figure 6b,d show corrosion attack by pitting mechanism in both HEAs which may be correlated with the differences in porosity in microstructure and the elemental distribution between the AM and the as-cast alloy. The AM alloy demonstrated less and smaller current fluctuations in the metastable pitting region due to the lower density of large pores which acted as active anodic sites and prevented nucleation and growth of new pits. As the potential reaches the breaking point of the passive film, the stable pitting of the surface likely initiates at the existing large pores. The as-cast alloy on the other hand showed extensive metastable pitting due to the presence of high density of micro-pores. In another study, CoCrFeMnNi was 72% cold rolled and annealed at 900 ◦C for 20 h resulting in a much more homogenized microstructure with equiaxed grains. This microstructure showed similar corrosion current density but much larger pitting resistance (Δ*E* ~420 mV) [39] compared to the AM CoCrFeMnNi in the present study.

**Figure 6.** Back scattered electrons(BSE) mode SEM images of CoCrFeMnNi surface before and after cyclic polarization test in 3.5 wt.% NaCl solution: (**a**) AM HEA (top-surface) before corrosion; (**b**) AM HEA (top-surface) after corrosion; (**c**) As-cast HEA before corrosion; (**d**) As-cast HEA after corrosion. The inset in (**b**) and (**d**) are high magnification images.

Elemental mapping was conducted by energy-dispersive spectroscopy (EDS) for both AM and as-cast HEAs. Table 2 shows the EDS quantitative analysis of the both alloys, where impurities such as C, N, and O were less than 0.1 at.%. The AM HEA showed uniform distribution of the elements throughout the surface (Figure 7) which may be due to better homogenization from repeated melting or thermal annealing of previous layers during the multi-layer deposition process. The higher pitting resistance of the AM HEA can be attributed to the lower density of large pores and the uniform distribution of Cr, which leads to the formation of a more stable passive film over the surface. In contrast, the as-cast HEA exhibited non-uniform elemental distribution with significant chemical micro-segregation (Figure 8). There were two distinct segregation regions, one with higher Cr, Co, and Fe and the other with higher Mn and Ni. Similar segregation behavior was reported for as-cast CoCrFeMn0.5Ni [40], which is typical for multi-principal element alloys that often have a wide freezing range [41]. The elemental map for the as-cast HEA revealed Cr depleted regions acting as the active sites for corrosion initiation and leading to the formation of non-uniform and less stable passive film. The EDS mapping also confirmed the presence of less Cr content at the micro-pores, which is potentially the main reason for the severe corrosion attack by pitting mechanism on as-cast HEA.

**Figure 7.** Energy-dispersive spectroscopy (EDS) map of additive manufactured CoCrFeMnNi HEA surface.

**Figure 8.** EDS map of as-cast CoCrFeMnNi HEA surface.


**Table 2.** EDS quantitative analysis of AM and as-cast CoCrFeMnNi alloy.

#### **4. Conclusions**

In this study, CoCrFeMnNi HEA was additively manufactured by SLM and its corrosion resistance in 3.5 wt% NaCl solution was evaluated by potentiodynamic polarization and electrochemical impedance spectroscopy measurements and compared with the as-cast counterpart. The main conclusions are as follows:


**Author Contributions:** Conceptualization, W.C. and S.M.; Methodology, W.C. and S.M.; Formal analysis, J.R. and C.M.; Investigation, J.R., C.M., L.L. and D.F.; Writing—original draft preparation, J.R. and C.M.; Writing—review and editing, L.L., W.C. and S.M.; Supervision, L.L., W.C. and S.M.; Project administration, W.C. and S.M.

**Funding:** This research received no external funding.

**Acknowledgments:** W.C. acknowledges the support by the NASA Marshall Space Flight Center Cooperative Agreement Notice (80MSFC19M0030) and UMass Amherst faculty startup.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


© 2019 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

### *Perspective* **High Entropy Alloys: Ready to Set Sail?**

#### **Indranil Basu 1,2,\* and Je**ff **Th. M. De Hosson 1,\***


Received: 24 December 2019; Accepted: 27 January 2020; Published: 29 January 2020

**Abstract:** Over the past decade, high entropy alloys (HEAs) have transcended the frontiers of material development in terms of their unprecedented structural and functional properties compared to their counterpart conventional alloys. The possibility to explore a vast compositional space further renders this area of research extremely promising in the near future for discovering society-changing materials. The introduction of HEAs has also brought forth a paradigm shift in the existing knowledge about material design and development. It is in this regard that a fundamental understanding of the metal physics of these alloys is critical in propelling mechanism-based HEA design. The current paper highlights some of the critical viewpoints that need greater attention in the future with respect to designing mechanically and functionally advanced materials. In particular, the interplay of large compositional gradients and defect topologies in these alloys and their corresponding impact on overall mechanical response are highlighted. From the point of view of functional response, such chemistry vis-à-vis topology correlations are extended to novel class of nano-porous HEAs that beat thermal coarsening effects despite a high surface to volume ratio owing to retarded diffusion kinetics. Recommendations on material design with regards to their potential use in diverse applications such as energy storage, actuators, and as piezoelectrics are additionally considered.

**Keywords:** serrated flow; thermal coarsening; actuators; phase transformation; nanoporous metals and alloys

#### **1. Introduction**

The classical rationale behind complex concentrated alloys (CCAs), more popularly referred to as high entropy alloys, comprises the addition of four or five elements in equiatomic and near-equiatomic proportions that eventually generate a single-phase solid solution [1,2]. The theoretical feasibility of such a counterintuitive alloy design stems from the concept of entropic stabilization, wherein a large number of elements would raise the configurational entropy and simultaneously overcome the enthalpies associated with the formation of intermetallic compounds. Owing to the inherent complexity of entropic stabilization, such multicomponent alloys have been observed to display peculiar characteristics that were summarized by Yeh and co-workers [1] as the four core effects: (i) The maximization of configurational entropy, owing to which the term high entropy alloys (HEAs) was coined; (ii) lattice strain due to large variation in atomic sizes of constituent elements; (iii) sluggish diffusion kinetics due to frustrated crystal structures; and iv) unusual properties displayed due to diverse interatomic interactions, also referred to as the "cocktail" effect [3].

However, the universality of occurrence of the abovementioned core effects in HEAs is debatable [4–9]. For instance, the very theory describing entropic stabilization has found little experimental validation, indicated by the fact that majority of HEAs that have fabricated till now either exist as multiphase alloys or decompose to more than one phase at thermodynamic equilibrium [6–8,10–12]. This is primarily because of the fact that the postulation by Yeh et al. on the maximization of configurational entropy remains valid at melting temperatures [13]. On the other hand, most experimental alloys are characterized at room temperatures, which renders significant microstructural and phase reordering driven by local stress/strain distributions, local compositional fluctuations, and interatomic interactions among constituent elements under diverse thermomechanical processing schemes [13,14]. Nevertheless, due to the widespread popularity of the postulated core effects within the community, they are still being actively associated with HEAs.

While the search for stable single-phase HEAs via combinatorial approaches continues [9,15], a large part of the research in the field delves into exploiting multiphase HEAs as means to design materials with significantly enhanced structural and functional properties [10,13]. Of special note in this regard are the less stringent alloying routes that provide access to a much larger compositional space in comparison with conventional alloys. This facilitates a mechanism driven compositional tuning of alloys, wherein the high solid solution content of individual elements can be tailored in order to generate unique features, such as spinodal structures [16], nano-scale coherent precipitates [17], and the modification of stacking fault energies [18,19], that give rise to interesting metal physics at intrinsic length scales. The advantages of this are clearly visible from the wide-spread future applications proposed for HEAs with unprecedented combinations of structural and functional properties [13,20–23].

In light of the aforementioned aspects of HEAs, it has become imperative to understand the independent and interdependent effects of the spatial distribution of linear and planar defects (i.e., defect topology) and compositional fluctuations existing in these alloys on the local mechanical and functional responses.

#### **2. Crystallographic Defects versus Compositional Variation: A Tale of Two E**ff**ects**

A critical aspect that influences mechanical responses in materials is the dependence of defect topologies and compositional gradients upon length scale dynamics. When considering crystal defects, the particular game players with respect to local plasticity are line defects and surface/interface defects, as well as their mutual interactions. Adding to this, the effect of local compositional fluctuations that could be considerable in HEAs further adds another layer of complexity in terms of local stress evolution and corresponding strain accommodation. Owing to such compositional diversity, metal physics in HEAs therefore offers an interesting playground wherein a superposition of multiple strengthening mechanisms that can interact and augment each other's contribution takes place. It must be understood that even though the observed metal physics in HEAs are derived from those that are seen for conventional alloys and superalloys, the compositional flexibility and simultaneous activation of different strengthening modes in the former potentially places them as more suitable candidates for mechanism-based alloy design.

The influence of local chemistry on meso-scale microstructures serves as a key design aspect in HEAs in regard to tailoring alloys with significantly improved structural responses. Owing to the absence of a primary solvent, local elemental partitioning in these alloys often results in the formation of phase interfaces. Depending upon whether such interfaces are crystallographically similar or dissimilar, the corresponding impact on the plasticity and strengthening behavior is distinct. For instance, one of the commonly studied systems in this regard is the AlxCoCrFeNi, wherein increasing the Al content drives a crystallographic transition from an Face-centered cubic (FCC) solid solution to a Body-centered cubic (BCC) phase [1]. Moreover, the BCC phases that exist in this alloy are further known to spinodally decompose into ordered B2 (enriched with Al and Ni) and disordered BCC structures [16,23–25]. It must be highlighted that the underlying strength contribution and metal physics are entirely different when considering such different interface types, and these become critical to appraise when engineering HEAs with strengthening across different length scales.

#### *2.1. Defect Generation and Strengthening Behavior across Crystallographically Similar Interfaces*

In a recent study [16], the nano-indentation response of BCC grains in Al0.7CoCrFeNi alloys indicated series of random displacement bursts. The observed pop-ins showed varying amplitudes

and seemed to be more obvious at lower indentation loads and smaller penetration depths (Figure 1c). Serrated flow characteristics are the fingerprint of jerky dislocation kinetics, as these arise from intermittent intervals of obstructed dislocation motion. In the present case, it was shown that the spinodally decomposed ordered B2 and disordered A2 phases generated interfaces that gave rise to simultaneous spinodal hardening and order-hardening effects (Figure 1a,b). These effects typically manifest at deformation length scales that are comparable to the mean size of the A2 phase, that is ~100–200 nm (Figure 1d,e). Considering the fact that spinodal strengthening in BCC crystal structures can be significantly larger than FCC spinodal alloys due to the sizable contribution of both elastic coherency strains and hardening from the modulus differential in the former, it has been shown that the strengthening potential in spinodal BCC HEAs can be as high as 0.5 GPa. In another study [17], it was shown that the addition of Ti and Al to single phase FCC CoCrFeNi HEAs leads to precipitation hardening effects due to presence of ordered FCC precipitates in a random FCC matrix, and these contribute to a strengthening increment between 0.3 and 0.4 GPa, which is significantly larger than counterpart contributions from strain hardening, grain boundary hardening, and solid solution strengthening. Lately, the concept of utilizing such spinodally-induced strengthening and order hardening effects in HEAs has given rise to a new generation of modulated, nano-phase structured, BCC-refractory HEAs that mimic super alloy type microstructures [26,27].

**Figure 1.** (**a**) Indent strain profile and (**b**) corresponding microstructure showing the dislocation motion across an A2/B2 interface. Dislocations experience strengthening when moving from the A2 to the ordered B2 phase, owing to spinodal and order hardening effects. The corresponding jerky dislocation motion is highlighted in nano-indentation load-displacement curves in (**c**). (**d**) shows the strain distribution with bright A2 phases being more plastic than the elastically stiffer dark B2 phase. (**e**) Corresponding misorientation profile and kernel average misorientation gradient on moving away from the interphase into the B2 phase. Figure 1d,e was adapted from [16] with permission from Elsevier, 2018.

#### *2.2. Strengthening Mechanisms across Crystallographically Dissimilar Phase Boundaries*

Unlike grain boundaries, the mechanics of strengthening across heterophase interfaces involve diverse contributing factors. Classical grain boundaries govern strain transmission, primarily on the basis of grain boundary geometry and the alignment of active slip systems in pile-up and emission grains [28–31]. On the other hand, interphase boundaries involve interface-dependent strengthening mechanisms [32] that add to the overall extent of dislocation pile-up and internal stress configurations.

In theory, heterophase interphases derive strength from three main contributing mechanisms, apart from the geometric slip transmission criterion.

#### 2.2.1. Image forces or Koehler forces:

The mismatch in shear moduli between neighboring phases gives rise to a Koehler force barrier between the dislocation and the interface [33,34]. The underlying effect that is responsible for this is the variation of strain energy per unit length of dislocation with changing modulus, such that a dislocation that moves from a stiffer grain to a softer grain would experience an attractive force, and the opposite scenario would result in repulsion between the incoming dislocation and the interface (Figure 2c). Mathematically, Koehler forces (τ*Koehler*) at an interphase boundary between phase *A* and phase *B* can be expressed as

$$\tau\_{Kochler} = \frac{G\_A(G\_B - G\_A)}{4\pi(G\_B + G\_A)} \cdot \frac{b}{h} \tag{1}$$

where *GA* and *GB* are the shear moduli values of incident and emission grains, respectively; *b* is the magnitude of the Burgers vector of active slip system in the incident grain; and *h* is the normal distance between the dislocation and the interface. The exerted force is hypothetically similar to the stress field that is exerted by a negative image dislocation that is positioned at the other side of the interface, hence the term of image forces. The presence of Koehler forces significantly impacts the dislocation pile-up characteristics at the phase boundary that subsequently influence the strengthening that is imparted from the interfaces, as shown in Figure 2a,b. For instance, it has been shown in a BCC/FCC dual-phase AlxCoCrFeNi alloy that the elastic modulus difference in FCC and ordered BCC phases (EBCC = 275 GPa vs. EFCC = 252 GPa) results in an attractive image force on incoming BCC dislocations and repulsive image forces on incoming FCC dislocations [35].

#### 2.2.2. Misfit Stresses

Crystallographically dissimilar phase boundaries also result in interfacial stresses that arise from lattice parameter mismatch (Δ*a*) between adjacent phases (Figure 2d). The size misfit is compensated by a grid of van der Merwe dislocations that give rise to coherency strain hardening effects at the interface [36]. Coherency stresses typically dampen as a function of 1/λ upon moving away from the interface, and they are mathematically given as,

$$
\tau\_{misfit} = 0.5G^\* \sqrt{\frac{2b(\delta - \varepsilon)}{\lambda}} \tag{2}
$$

where δ = <sup>Δ</sup>*<sup>a</sup> <sup>a</sup>* ; *a* is the mean lattice parameter (*aphase*\_*<sup>A</sup>* + *aphase*\_*B*)/2, ε = 0.76δ is the residual elastic strain that was determined to agree for most heterophase interface types, *G*∗ is the average shear modulus for the two phases, and λ is the grain dimension over which misfit stresses are determined, i.e., the distance between a dislocation and the interface. Coherency stresses exert a Peach–Koehler force on incoming glide dislocations, which can be either attractive or repulsive depending on the sense of applied stress with respect to the dislocation slip system. Apart from affecting the dislocation glide stresses, the coherency stresses aid in strengthening by additionally influencing the non-glide stress components of the dislocation stress field, whereby they can locally modify the dislocation core energy that directly influences the ease of a dislocation in overcoming an obstacle.

#### 2.2.3. Chemical Mismatch Effect

Another aspect that contributes to interfacial strengthening is a mismatch in chemical energy or gamma surfaces, as this mismatch directly determines the stacking fault energies in adjacent phases [37]. When the leading partial in a stacking fault moves across an interface, the dislocation

configuration experiences an abrupt change in stacking fault energy. This manifests as an effective stress that is exerted upon the leading dislocation in the pile-up (Figure 2e). The resultant stacking fault strengthening stresses is described as

$$
\tau\_{\text{chemical}} = \frac{\Delta \mathcal{V}}{b} \tag{3}
$$

where Δγ is the stacking fault energy differential between neighboring phases. Overall, the change in energy of dislocation as it moves across the phase boundary involves an elastic energy contribution (that is a combination of τ*Koehler* and τ*mis fit*) and a chemical contribution in the form of τ*chemical*.

**Figure 2.** (**a**) BCC–FCC interface in an AlxCoCrFeNi high entropy alloy, with indent profile. (**b**) Variation of hardness and residual stress as a function of distance from phase boundary. The role of image forces is highlighted by local maximum in stress values on the BCC side (attractive image forces) and a local minimum on the FCC side (repulsive image forces). (**c**–**e**) illustrate different interfacial-dependent strengthening mechanisms. Figure 2b was adapted from [2] with permission from Elsevier, 2018.

Owing to the strong compositional fluctuations and propensity of single-phase decomposition in HEAs, the contribution of interface-dependent strengthening could be significant in terms of augmenting overall material strength at both local and global scales. Recent studies have now made effective use of such hardening mechanisms to tailor microstructural designs that give rise to substantially stronger HEAs in comparison with their single-phase counterparts or with respect to conventional alloys. It has been shown that the BCC–FCC interfaces in HEAs could give rise to strengths of the order of 4 GPa that are nearly four times the measured values of conventional

BCC–FCC interfaces [35]. The underlying contribution for the augmented strengthening in HEAs has primarily been attributed to the enhanced interfacial-dependent strengthening that is caused by the strong compositional gradients in these alloys. Investigations on HEAs that comprised of BCC/FCC multilayers has indicated yield strengths of the order of 3.3. GPa, with more than two-thirds of this strength coming from interfacial strengthening effects and the remaining coming from solid solution strengthening [38]. In another study [39], it was shown that BCC/FCC interfacial strengthening mechanisms could be further enhanced by tuning the multilayer thickness in HEAs, whereby strengths of the order of ~13 GPa can be reached.

#### **3. An Outlook to HEAs: Structural Properties**

Compositionally, HEAs can be described as a concoction of multiple elements, an arrangement that often results in a frustrated crystal structure. Moreover, chemical gradients further trigger local rearrangements and the shuffling of elements, thus influencing the stability of the existing phases.

In short, HEAs are considerably more prone to phase transformation under applied temperature or stress, which could be a potent mechanism to trigger interesting plasticity mechanisms as well as to accommodate larger strains. For instance, in a seminal work by Li et al. on non-equiatomic compositions [10] based on the FCC single phase cantor alloy, it was shown that under plastic deformation, the dynamic transformation from an FCC to an HCP crystal structure is achieved that simultaneously enhances strength and ductility. In a more recent study [16], dynamic indentation-induced phase transition from BCC to FCC was observed in BCC Al0.7CoCrFeNi HEAs (Figure 3a,b). The underlying reason behind the transformation was attributed to the spinodal decomposition of the BCC phase in Al, Ni rich-ordered B2 phases and random A2 phases (Figure 3a). Under applied stress, the A2 phases that are locally depleted in Al content could displacively transform and revert back to the more stable and ductile FCC phase (Figure 3c).

The results once again provide an opportunity to exploit the compositional fluctuations in tandem with thermomechanical treatment to dynamically trigger strength and ductility enhancing mechanisms. Displacive phase transformation effects or TRIP effects in HEAs could be exciting focal points in novel advances of HEAs in structural properties and applications.

Another mechanistic design criterion that employs compositional fluctuations is through the intrinsic modification of stacking fault energies. It was shown by Ritchie and co-workers [40,41] that by tuning local chemical ordering in HEAs, considerable variation in the intrinsic and extrinsic stacking fault energy values can be realized. Such design pathways are critical in terms of triggering new deformation mechanisms such as deformation twinning, whereby additional strain accommodation mechanisms are dynamically activated. Twinning not only contributes to plasticity but can also promote dynamic Hall–Petch-driven strengthening behavior, owing to grain fragmentation caused by twin boundary formation. For instance, it was shown by Deng et al. that for non-equiatomic Fe40Mn40Co10Cr10 HEA deformation, twinning is triggered as an additional mechanism for higher strains, thereby contributing to the overall strength–ductility increment [18].

Finally, the possibility to exploit simultaneous TRIP and TWIP effects along with multiscale strengthening effects that encompass interfacial strengthening and solid solution hardening effects can be envisioned in upcoming HEAs. In this regard, it has been shown that for the non-equiatomic FeMnCoCr alloy, dilute additions of C (~0.6 at. %) already trigger simultaneous twinning and phase transformation along with an interstitial hardening response [19]. Similar alloying strategies have been implemented on non-equiatomic BCC HfNbTaTiZr, wherein simultaneous plasticity-induced displacive transformation from the BCC phase to the HCP phase was observed, along with twinning in the HCP phase [42].

**Figure 3.** (**a**) EDS maps showing selective partitioning of Al and Cr into B2 and A2 phases, respectively. (**b**) Nano-indentation-induced phase transformation of A2 BCC phase (shown in red) to FCC phase (shown in green). Grain boundaries shown in white in the left image, and the non-indexed areas that are shown in white in the right image correspond to the experimentally made indents; local compositional fluctuations of Al and Cr indicated in (**c**), highlighting the instability of A2 BCC phases in regions of depleted Al, whereby such dynamic phase transformation to FCC at room temperature deformation is facilitated. Experimental data for the figures were derived from [16]. Figure 3b was adapted from [16] with permission from Elsevier, 2018.

The abovementioned impact of compositional fluctuations and strengthening modes related to phase formation also needs to be incorporated into current solid solution strengthening models in HEAs. There are now sufficient studies that have revealed that the strengthening of dislocation motion in HEAs is strongly dependent upon its susceptibility to display either short or long range ordering effects rather than simple lattice friction-induced hardening responses [43,44]. This was validated by a recent study by Robert Maaß and collaborators, wherein the peak dislocation velocities in FCC Al0.3CoCrFeNi and pure Au did not show much difference, indicating that dislocation motion was not significantly sluggish in single phase solid solution HEAs (Rizzardi et al. [45]). Moreover, the contributions of interfacial-dependent strengthening and solute strengthening modes need to be appraised, as these could be critical in driving application-based future multiphase HEA alloy design. In this regard, greater efforts are needed in understanding the influence of alloying chemistry on engineering interphase boundaries in HEAs rather than focusing upon solid-solution strengthening as the primary strength contributor in these alloys. Indeed, the outcomes look promising and may open a new paradigm of structurally advanced HEAs, as was recently shown in study [46] where a compositionally graded AlxCoCrFeNi bar was additively manufactured with increasing Al contents from x = 0.3 to x = 0.7 along the longitudinal direction, such that one end of the material was a single phase FCC and the other end formed a dual phase B2–FCC microstructure. From the point of view of mechanical response, the dual phase microstructure clearly highlighted the positive role of interfaces with significantly larger strengthening potentials compared to the single-phase FCC solid solutions.

#### **4. An Outlook to HEAs: Functional Properties**

Thus far, the emphasis in our feature paper has lied upon mechanical performance, plasticity and damage control. Attractive and rather unexplored frontiers of HEAs concern applications of functional properties, e.g., magnetic and electrical including thermoelectricity (Seebeck effect) [40], in the field of microelectronics and bio-medicine [23,47].

Several ideas are currently under investigation, and, in particular, we like to refer to possibilities of HEAs as cellular/porous materials that can be used in 'energy materials' (H-storage) but also explored as radiation resistant sensors and actuators. Actuation refers to mechanical displacement due to an electric signal. The opposite is also possible when an electrical current is generated by mechanical deformation, which creates what is known as piezoelectric materials. In general, these piezoelectrics need high voltages, with an order of 100 V, and, at present, methods are being developed in medicine and biology to manufacture high-precision actuators that work at lower voltages on cell manipulation [48]. Structural stability over a large range of temperatures is essential.

In particular, highly porous metallic systems can mimic the properties of muscles upon an outside stimulus, and they have been coined 'artificial muscles' in analogy to human skeletal muscles, which are ideal actuators with a high energy efficiency, fast strain-rate response, and high durability. The common use of existing materials as actuators like piezoceramics and electroactive polymers are limited by several factors, including low energy efficiency, low strain amplitudes, fatigue limits, and the high actuation voltages needed. In our recent work, we have shown that nanoporous organometallic materials can operate as actuators, thereby offering a unique combination of relatively large strain amplitudes, high stiffness and strength, and importantly (see above) low operating voltages—say, at a few volts. However, a serious concern in this field of applications is (thermal) stability and the effects of coarsening.

Because of their high porosity and surface areas, the stability of cellular/porous systems is a major issue. In fact, depending on temperature (low versus high) and environment, the stability might be questionable, and, as a consequence, the functional properties might be not very stable and may deteriorate over time. Because of the suppression of the diffusional processes of defects at the surface of a (nano) ligament, we believe that HEAs might provide a very interesting and effective remedy to address these essential problems for applications of unique properties of functional HEA materials.

#### *4.1. Porous*/*Cellular Systems*

To illustrate the problem in a bit more detail, we discuss here the characteristics of cellular and porous media that possess a lower density and a higher surface area-to-volume ratio. The terminology is in macrofoams a bit different from that of nanofoams. In the former, besides pores, the material is made up of struts, and, in the latter, we call the struts ligaments. The topology of nodes and struts/ligaments can be anisotropic as well as isotropic. Besides these structural differences, nanoporous foams have been applied in nanofiltration systems, drug delivery platforms, catalysis, sensing and actuation [49–54]. In contrast, macro foams have been explored in macroscopic applications of the transport, automotive and aerospace industries.

A popular way of making metal nanofoams is based on dealloying through leaching. Preferably, the base material is a solid solution of a noble and a less noble element. Unfortunately, many alloys form intermetallics and many metals do not easily form solid solutions, which limits the dealloying methodology. Recently, nanoporous HEAs were produced through a rather novel method [55] by using liquid metal dealloying (LMD), a technique to fabricate non-noble porous materials by suppressing oxidation in a metallic melt [56–62]. It turned out that the structure of nanoporous TiVNbMoTa HEAs [55] can be described as nanoscale ligaments of a solid-solution phase, the stability of which is due to suppressed surface diffusion.

#### *4.2. Suppression of Coarsening*

It should be realized that the actuation mechanism based on nanoporous metals with high surface-to-volume ratios is different to piezoceramics. For details, one can look to [63]. The physical principles in the case of metals are based on the lower coordination of the surface atoms. Therefore, to gain 'density,' the atoms move inwards, and a positive displacement is necessary to bring them back to the equilibrium interatomic distances of the situation in the bulk. Therefore, a positive, i.e., tensile stress state at the surface, is generated that is compensated for by a compressive stress state inside the

strut/ligament. Clearly, a positive charge injection can equilibrate the existing excess negative surface charge. As a consequence, a positive charge lowers the positive tensile stress and, importantly, relaxes the negative compressive stress in the ligament, i.e., the negative compressive stress becomes a bit more 'positive', i.e., a positive charge will generate a positive displacement that can be detected by optical means by using a small laser. It is noteworthy that these rather small displacements as a result of stress relaxations are not detectable in macrofoams, but nanofoams are completely different because a much large surface area to volume ratio exists. To give a rough estimate: If the size of the ligaments is of the order of 5–10 nm, a substantial fraction of the total number of atoms, say 20%–10%, is on the surface, meaning that a 100 nm thick ligament already reflects the bulk situation, and the actuating properties are hardly detectable. It is important to realize that the electronic charge distribution at a nanoporous metal interface can effectively be controlled during cycling voltammetry experiments. Only small electrical voltages of the order of 1 V are needed to bring positive or negative charge carriers (ions) from the electrolyte to the nanoporous metal [64,65]. Of course, working with liquid electrolytes can be cumbersome, and [65] we recently demonstrated that metallic muscles can operate in a dry environment, even at high strain rates, i.e., much higher than allowed for in electrochemical artificial muscles [65].

Nonetheless, electrochemical processes may lead to severe coarsening (undesired growth) of the ligaments [66–68], which is a major concern because actuation is hampered or completely lost, as is shown in Figure 4, where strain amplitudes are plotted as a function of the average ligament size. As expected, there has been a strong size effect and the strain amplitude recorded on nanoporous systems with different ligament sizes, and these may decrease upon increasing the ligament size (Figure 4). Even as a sensor (i.e., without external applied electric fields), the response is not stable due to coarsening effects (Figure 5), and, at higher temperatures, these effects are amplified tremendously.

**Figure 4.** Ligament size-dependence of the charge-induced strain in nanoporous metals (Au). The strain amplitude that was recorded on five specimens with different ligament sizes decreased with increasing ligament size. This shows that ligament growth during electrochemical actuation is undesirable [67,68].

**Figure 5.** (**a**) Typical scanning electron micrograph of a nanoporous material (Au) that was synthesized by the dealloying process. The diameter of the ligaments was about 20 nm. (**b**) Display of the changes in relative humidity versus time for alternations of humid and dry air (blue curve refers to right ordinate) and corresponding strain versus time (red curve refers to left ordinate). (**c**) Responses of relative humidity versus time for long alternations of dry and humid air (blue curve refers to right ordinate) and corresponding strain versus time (red curve refers to left ordinate [67]).

The solution to these instabilities is basically to reduce surface diffusion, and, for that reason, HEAs may offer a suitable solution. Recently, Soo-Hyun Joo and collaborators found an exceptional stability against coarsening of an MoNbTaTiV nanoporous HEA at elevated temperatures [55]. The ligament size and distribution of the HEA versus dealloying time at various temperatures are displayed in Figure 6.

**Figure 6.** Ligament size in a TiVNbMoTa nanoporous high entropy alloy (HEA) versus dealloying time at various temperatures. The dealloying time, *t*, and ligament size, *d*, are correlated through a power function *d<sup>n</sup>* = *ktD*, where *n* is the coarsening component, *k* is a constant, and *D* is the surface diffusivity. By plotting the *ln*[*d*(*t*)] vs. ln *t* curve, the coarsening exponent, *n*, can be obtained. Error bars denote the distribution of ligament sizes. (Figure reprinted from [55] with permission from Wiley, 2019.)

To make the coarsening behavior a bit more quantitative: Dealloying time, *t*, and ligament size, *d*, are written as a power function: *dn* = *ktD*, where *n* is the coarsening component, *k* is a constant, and *D* is the surface diffusivity. By plotting the *ln*[*d*(*t*)] vs. *ln t* curve, the coarsening exponent, *n*, can be measured. Figure 6 displays the measured specific surface areas of 55.7 (<10 nm), 38.8 (14 nm), and 3.6 (155 nm) m2/g, depending on the ligament size as obtained in [55] based on an analytical model we designed in the past [68]. The adsorption/desorption isotherm curve (inset in Figure 7) corresponds to a ligament size of about 10 nm. The hysteresis loop in the isotherm is associated with capillary condensation in nanopores (<50 nm).

**Figure 7.** Specific surface areas measured with the Brunauer–Emmett–Teller (BET) method as a function of the ligament size in a TiVNbMoTa nanoporous HEA. The lines (dashed and dashed-dot) correspond to the predictions based on the analytic model *S* = *C*/ρ*d*, as designed and tested in [68]. The inset shows the nitrogen adsorption/desorption curves with an average ligament size of ~10 nm. The solid bulk density <sup>ρ</sup> was assumed to be 7.7 g·cm−<sup>3</sup> from the average atomic weight and the radius of the constituent elements (Figure reprinted from [55] with permission from Wiley, 2019.)

As far as the physical explanations of these phenomena are concerned, a first and easy conclusion would be: The suppression of coarsening is due to what has been called sluggish diffusion in HEA. However, this might be a too simple and naïve reason. There is no literature on the details of diffusional processes in nano-porous HEA, but many papers have been published on bulk HEA [69], e.g., CoCrFeMnNi [70]. Dezso Beke [71] analyzed the diffusion coefficients of elements in CoCrFeMnNi HEA. Interestingly, sluggish diffusion could be explained as not being based on high activation energies but on correlation effects. Indeed [72], it has been reported that the diffusion of Ni in both HEAs follow an Arrhenius behavior. In fact, the tracer diffusion in HEAs does not become sluggish at absolute temperature; it only becomes so if it is considered at a homologous temperature. As a sequence, it can be concluded that diffusion in HEAs is not sluggish, as such, but other factors such as frequency factors can explain a slower diffusion rate in HEAs. From a scientific viewpoint, more in-depth analyses regarding the diffusional processes are necessary, not only for bulk HEA but also for surface diffusional processes in nanoporous HEAs.

A stimulating and exciting field of research would be to devise specific rules for the design of highly porous HEAs that can be used over a large range of temperatures in the field of sensors and actuators, in 'energy materials' (H-storage), radiation resistant materials, and many other scenarios. It is important to note that in these cellular and highly porous HEAs materials, it is not just the specific materials properties of strength and surface diffusion matter—the local topology and connectivity of struts and nano-ligaments also does [73].

As a consequence, topology provides an additional design parameter, in addition to the extrinsic and intrinsic material size-dependent properties [73,74]. This approach fits the term 'architected material,' which was, to the best of our knowledge, proposed for the first time by Mike Ashby and Yves Bréchet and which bridges the structural engineering of topology and good practice in architecture [75]. More recently, a special issue of excellent contributions to the Materials Research Society (MRS) Bulletin, edited by Julia R. Greer and Vikram S. Deshpande [76], was published on the design, fabrication and mechanical performance of three-dimensional architected materials and structures. It would be exciting to explore HEAs among these lines of 'architected materials' for the optimization of the 'integral' of both structural and functional properties together for novel applications.

**Funding:** This research received no external funding.

**Acknowledgments:** The work was supported by the Applied Physics-Materials Science group of the Zernike Institute for Advanced Materials of the University of Groningen, the Netherlands. Discussions and contributions by Eric Detsi (now at UPENN, Philadelphia, USA) and Václav Ocelík (Groningen, the Netherlands) are gratefully acknowledged.

**Conflicts of Interest:** The authors declare no conflict of interest.

#### **References**


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