Microstructures and Deformation Mechanisms of FCC-Phase High-Entropy Alloys

*Kaisheng Ming, Shijian Zheng and Jian Wang*

## **Abstract**

Strength and ductility are the most fundamental mechanical properties of structural materials. Most metallurgical mechanisms for enhancing strength often sacrifice ductility, referred to as the strength–ductility trade-off. Over the past few decades, a new family of alloys—high-entropy alloys (HEAs) with multi-principal elements, has appeared great potential to overcome the strength–ductility trade-off. Among various HEAs systems, CrFeCoNi-based HEAs with a face-centered cubic (fcc) structure exhibit a great combination of strength, ductility, and toughness via tailoring microstructures. This chapter summarizes recent works on realizing strength–ductility combinations of fcc CrFeCoNi-based HEAs by incorporating multiple strengthening mechanisms, including solid solution strengthening, dislocation strengthening, grain boundary strengthening, and precipitation strengthening, through compositional and microstructural engineering. The abundant plastic deformation mechanisms of fcc HEAs, including slips associated with Shockley partial dislocation and full dislocations, nanotwinning, martensitic phase transformation, deformation-induced amorphization, and dynamically reversible shear transformation, are reviewed. The design strategies of advanced HEAs are also discussed in this chapter, which provides a helpful guideline to explore the enormous number of HEA compositions and their microstructures to realize exceptional strength–ductility combinations.

**Keywords:** high-entropy alloys, strength, ductility, twinning, phase transformation, amorphization, reversible shear transformation

## **1. Introduction**

Developing structural materials with both high-strength and ductility could mitigate the ecological and economical concerns for decreasing weight and improving energy efficiency. Unfortunately, these properties are generally mutually exclusive, i.e., increasing strength will inevitably lead to ductility loss, an effect referred to as the strength–ductility trade-off [1–4]. Various strengthening mechanisms, including solid solution strengthening, dislocation strengthening, grain boundary strengthening, precipitation strengthening, twinning, and phase transformation-induced hardening, have been widely utilized to produce high-strength and high-ductility alloys that are based on one principal element [5]. Recently, a new idea—high-entropy alloys (HEAs), has appeared that shows great potential to overcome the strength–ductility trade-off, and thus break through the mechanical property limits [6–13]. HEAs were founded independently by two different research groups in 2004 [6, 7], which emphasize the unexplored regions in the center of multi-element phase diagrams, in which all elements are concentrated and there is no base element. HEAs are generally defined as an alloy that is composed of four or more elements in an equiatomic or near-equiatomic composition [7]. Each HEA can be considered as a new alloy base because its properties can be further optimized by minor elemental additions, similar to alloying in convention alloys. HEAs provide near-infinite new alloy bases for designing structural materials with excellent performance. Depending on the composition and microstructure, HEAs display attractive mechanical properties, and possible combinations of some properties, including high strength/hardness, outstanding wear resistance, excellent fatigue resistance, exceptional high-temperature strength, good structural stability, good corrosion and oxidation resistance, and high radiation tolerance [8, 10, 12, 14]. Particularly, single-phase face-centered cubic (fcc) HEAs and medium-entropy alloys (MEAs) based on the transition metal elements Cr, Mn, Fe, Co, or Ni generally display some of the best mechanical properties reported to date [9–12, 15–25]. For example, equiatomic CrMnFeCoNi HEA and CoCrNi MEA possess exceptional combinations of tensile strength and ductility (tensile strength of ~1 GPa as well as ductility exceeding 60%) at 77 K, and ultra-high fracture toughness at both room temperature and 77 K (KJ1C > 200 MPa√m), making them one class of the toughest metallic materials reported so far [17, 23, 26]. Such exceptional mechanical properties are attributed to continuous steady strain-hardening, resulting from extensive dislocation activities and deformation-induced nanotwinning [18, 21, 22, 27]. These fcc-structured HEAs can be used as ideal alloy bases to design high-strength and high-ductility structural materials through further compositional and microstructural engineering.

Intensive studies have been invested in overcoming the strength–ductility tradeoff of the single-phase fcc CrFeCoNi-based HEAs through tailoring the chemical composition and microstructure [9–12, 15–24]. It has been reported that metastable high-entropy dual-phase alloys can overcome the strength–ductility trade-off by interface hardening and transformation-induced hardening, realized by reducing the stacking fault energy (SFE) via tailoring chemical composition [15, 19, 24, 28–33]. The tensile strength and ductility are simultaneously enhanced due to heterogeneous microstructures, such as gradient nanotwins, gradient nano-grains, or recrystallized and non-recrystallized grains arranged in hierarchical structures with characteristic dimensions spanning from submicron scale to micro-scale, that are obtained by coldrolling and annealing [9, 16, 27, 34, 35]. Note that single-phase fcc CrFeCoNi-based HEAs often have a very low SFE, which promotes deformation-induced nanotwinning and martensitic phase transformation [36–42]. The simultaneous increase in tensile strength and ductility is related to the enhanced strain-hardening capability enabled by nanotwinning and/or phase transformation. However, they generally possess very low yield strength since the nanotwinning and phase transformation cannot be activated at the early stages of plastic deformation [12].

Multiple strengthening mechanisms (such as solid solution strengthening, dislocation strengthening, grain boundary strengthening, and precipitation strengthening, et al.) show great potential to enhance the yield strength and tensile strength of fcc CrFeCoNi-based HEAs while maintaining their excellent ductility and strain-hardening capability [43–62]. In addition, some new strengthening mechanisms are also proposed to improve the mechanical properties, such as magnetic hardening [60]. This chapter summarizes recent works on realizing strength–ductility combinations of fcc

#### *Microstructures and Deformation Mechanisms of FCC-Phase High-Entropy Alloys DOI: http://dx.doi.org/10.5772/intechopen.104822*

CrFeCoNi-based HEAs by incorporating multiple strengthening mechanisms, achieved through compositional and microstructural engineering. The abundant plastic deformation modes of fcc HEAs, including slips associated with Shockley partial dislocation and full dislocations, twinning, martensitic phase transformation, deformationinduced amorphization, and dynamically reversible shear transformation, are also summarized. We also demonstrate some HEA design strategies to provide guidelines for exploring the enormous number of HEA compositions and their microstructures to realize exceptional strength–ductility combinations.

## **2. Strengthening via tailoring composition and microstructure**

## **2.1 Ductile nanoprecipitates**

Single-phase fcc HEAs generally exhibit excellent ductility and strain-hardening capability but low yield strength at room temperature. Extensive studies demonstrate that nanoprecipitation strengthening is one of the most effective approaches to strengthen HEAs without apparent ductility loss. For example, Ming et al. [61] introduced highly dispersed nano-sized, coherent precipitates in the grain interior of coarse-grained fcc CrFeCoNi-based HEAs (grain size ~ 1 mm) via alloying a small addition of Al and Ti elements, which realizes exceptional combinations of strength and ductility (**Figure 1**). Yang et al. [47] produce a CrFeCoNi-based HEA with a superb yield strength of above 1.0 GPa while maintaining 50% ductility in tension at room temperature via introducing high-density ductile multicomponent intermetallic nanoparticles with coherent phase boundaries, enabled by alloying a small addition of Ti and Al elements. The size and distribution of nanoprecipitates can be tailored by the content of alloying elements (Ti and Al), aging temperature, and aging time [61]. As shown in **Figure 1a**–**g**, the strength and ductility can be optimized corresponding to the aging time-dependent size and spacing of nanoprecipitates in an Al0.2Co1.5C rFeNi1.5Ti0.3 HEA. When aged at 800°C for 1–5 h, numerous uniformly dispersed, nano-sized, spherical L12 precipitates are formed in coarse grains (grain size ~ 1 mm), which results in a significant increase in both yield strength and ultimate tensile strength without apparent sacrificing of ductility (**Figure 1a, b, f**). With increasing aging time, the average diameter of precipitates increases from 6 nm at 1 h of aging to 51 nm at 100 h of aging (**Figure 1e**). **Figure 1f** presents that aging time does not show a significant effect on the yield strength and ultimate tensile strength, but increasing aging time results in an apparent decrease in ductility. As shown in **Figure 1g**, the strain-hardening rate obviously increases with increasing aging time, in particular, for the aging time equal and longer than 5 h. The samples aged for 1–5 h exhibit superior combinations of strength and ductility (yield strength of ~760 MPa, ultimate tensile strength of ~1160 MPa, elongation of ~40%).

The nature of dislocation interaction with the nanoprecipitates in Al0.2Co1.5CrFe Ni1.5Ti0.3 HEA samples is revealed by using transmission electron microscopy (TEM) analysis, as shown in **Figure 2**. The TEM bright-field image of the sample aging at 800°C for 1 h (**Figure 2a**) shows high-density dislocations tangled with stacking faults after tensile deformation to fracture. As shown in **Figure 2b**, original spherical precipitates have changed to an irregular shape after tensile deformation, which indicates the dislocations cutting through precipitates. In contrast to dislocation cutting through nano-sized precipitates, dislocations bypass relatively large precipitates by looping them in the sample aging at 800°C for 50 h, as demonstrated in **Figure 2c**. **Figure 2d**

#### **Figure 1.**

*TEM images of the Al0.2Co1.5CrFeNi1.5Ti0.3 HEA after aging at 800°C for: (a) 1 h (b) 5 h, (c) 50 h, and (d) 100 h, with the corresponding selected area electron diffraction patterns inset. (e) Influence of aging time on the diameter of the precipitate (D), separation distance (L) between the centers of neighboring precipitates, and edge-to-edge inter-precipitate distance (l). (f) Engineering stress–strain curves of the solution-annealed and aged samples. (g) Variation of strain-hardening rates with plastic strain [61].*

#### **Figure 2.**

*(a) TEM image of the 1 h aged Al0.2Co1.5CrFeNi1.5Ti0.3 HEA at a tensile strain of 36%, (b) HRTEM image of a precipitate with an irregular shape. (c) TEM image of the 50 h aged sample at a tensile strain of 3%. (d) Variation of precipitate shearing stress (σsh) and Orowan dislocation looping stress (σOr) with aging time [61].*

#### *Microstructures and Deformation Mechanisms of FCC-Phase High-Entropy Alloys DOI: http://dx.doi.org/10.5772/intechopen.104822*

schematically illustrates that dislocation cutting through precipitates is the primary strengthening mechanism for the 1 h and 5 h aged HEA with small precipitates (diameter < 10 nm), while the dominant strengthening mechanism is dislocation bypassing precipitates by looping them in HEAs aged for >25 h with large precipitates (diameter > 30 nm). Based on the tensile testing and TEM characterizations, it can be concluded that nanoscale coherent precipitation strengthening is a very effective approach for strengthening coarse-grained fcc-structured HEAs without loss of ductility. The exceptional combinations of strength and ductility can be achieved via introducing uniformly distributed, nano-sized, coherent precipitates in the coarse-grained HEAs, where gliding dislocations cut through nano-sized coherent precipitates. In contrast, with increasing the size of precipitates, the ductility is reduced due to the strong barrier of large precipitates to dislocation motion where gliding dislocations bypass a precipitate by looping it, resulting in dislocations pileups at precipitate-matrix interfaces which generate high stress/strain concentration and micro-crack initiation.

It should be noted that the sluggish diffusion effect in HEAs enables easily tailoring the size of precipitates by adjusting aging processing and optimizing compositions. Recently, the coarsening kinetics and thermal stability of nanoscale precipitates in HEAs at elevated temperatures have also been widely studied to explore their applications at high temperatures [63–66]. In a word, precipitation strengthening is one of the most promising approaches for enhancing the mechanical properties of fcc HEAs to meet the requirements for engineering applications both at room temperatures and high temperatures.

#### **2.2 Brittle intermetallic compounds**

The original concept employed in the HEAs design is to suppress the formation of brittle intermetallic compounds which can lead to poor ductility [6, 7]. Unfortunately, most HEAs reported to date contain various brittle intermetallic compounds [63–66]. Therefore, it is required to manipulate the intermetallic compounds to reconcile the strength and ductility of HEAs. It has been shown that the brittle but hard intermetallic compound μ phase can be effectively used as a strengthening unit in CrFeCoNiMo HEAs while relieving its harmful effect on ductility by manipulating its dimension and distribution via tailoring Mo contents [62, 67–69]. Moreover, by further coupling solid solution hardening and nanotwinning induced hardening, a superb yield strength-tensile strength–ductility combination is realized [62]. As shown in **Figure 3**, the dimension and distribution of the μ phase can be tuned through thermal-mechanical processing and annealing. The μ phases grow mainly at boundaries (grain boundaries, triple junctions, and annealing twin boundaries), and secondarily in the interior of grains. The corresponding engineering stress– strain curves in **Figure 4a, b** demonstrate that dual-phase (fcc matrix + nanoscale μ phase) Cr15Fe20Co35Ni20Mo10 (Mo10) HEA displays high strength (yield strength of 0.8–1.3 GPa and ultimate tensile strength of 1.1–1.4 GPa) but moderate ductility (elongation to fracture is approximately 13–28%). The single-phase fcc Cr12.5Fe20Co4 2.5Ni20Mo5 (Mo5) HEAs have superb ductility (elongation to fracture of 45–75%) but moderate strength (yield strength of 0.3–0.8 GPa and ultimate tensile strength of 0.7–1.0 GPa). **Figure 4c** demonstrates that dual-phase Mo10 HEAs annealed at 850– 1000°C exhibit a higher strain-hardening rate than single-phase Mo5 HEAs, and Mo10 HEA annealed at 1150°C as the true strain is less than half the maximum elongation. The high strain-hardening rate in Mo10 HEAs annealed at 850–1000°C is ascribed to the formation of distributed μ phase precipitates while the low strain-hardening

#### **Figure 3.**

*TEM images of cold-rolled Mo10 HEAs after annealing at 800°C for (a1) 5 min, (a2) 1 h; 900°C for (b1) 5 min, (b2) 1 h; 1000°C for (c1) 5 min, (c2) 1 h [62].*

rate in Mo10 HEAs annealed at 1150°C is due to the decomposition and disappearance of clustered μ phase precipitates. **Figure 4d** compares the mechanical properties of the dual-phase Mo10 HEA and single-phase Mo5 HEA with various other fcc HEAs [70–75]. An exceptional combination of strength and ductility of Mo10 and Mo5 HEAs is attributed to the synergetic effect of solid solution strengthening, precipitation hardening, and twinning-induced hardening. The solution strengthening effect of the Mo addition in single-phase Mo10 alloys is evidenced by the higher strain-hardening rate of single-phase Mo10 alloy after true strain exceeds 13%

(**Figure 4c**). **Figure 5a, b** demonstrates dislocations are bowing out or piled up at the μ phase interfaces. It is noted that μ precipitates are plastically non-shearable by gliding dislocations, which accounts for the high strength of Mo10 alloy according to the precipitation strengthening mechanism and the high strain-hardening rate at the early stage of plastic deformation. As demonstrated in **Figure 5c**, numerous deformation nano-twins are observed in single-phase Mo10 alloy at a strain of 25%, which provides a steady source of strain-hardening by blocking the motion of dislocations.

Other brittle intermetallic compounds, such as the sigma phase, Laves phases, have been also frequently observed in HEAs, which generally deteriorate the ductility *Microstructures and Deformation Mechanisms of FCC-Phase High-Entropy Alloys DOI: http://dx.doi.org/10.5772/intechopen.104822*

#### **Figure 4.**

*Engineering stress–strain curves for the (a) homogenized and recrystallized Mo10 alloys and (b) Mo5 alloys. (c) the strain-hardening rates of Mo10 and Mo5 alloys as a function of plastic strain. (d) Comparing the yield strength and elongation of the Mo10 and Mo5 alloys with various HEAs [62].*

#### **Figure 5.**

*TEM bright-field images of annealed Mo10 alloy at tension strains of 15–25%, (a) bow-out of gliding dislocations, (b) pile-up of dislocations at interfaces, and (c) high-density deformation nano-twins [62].*

[76–80]. Similarly, if the size, distribution, and morphology of intermetallic compounds can be tailored via controlling the chemical composition and thermomechanical treatment, an excellent combination of strength and ductility can be achieved in multi-component HEAs.

#### **2.3 Hierarchically heterogeneous microstructures**

Hierarchically heterogeneous microstructures at scales from a few nanometers to hundreds of micrometers can be easily produced in fcc HEAs by conventional cold work and annealing treatment, which could generate superior mechanical properties to those with simple microstructures due to the hetero-deformation induced strengthening [81]. As shown in **Figure 6**, single-phase fcc Cr20Fe6Co34Ni34Mo6 HEAs with hierarchical microstructures that comprise high-density annealing

nanotwins in recrystallized fine grains (grain size ~ 1 μm), and dislocation walls and stacking faults in non-fully recrystallized fine grains are produced by cold-rolling and annealing [82]. The formation of numerous annealing nanotwins and stacking faults is obviously attributed to the very low SFE of the Cr20Fe6Co34Ni34Mo6 alloy. The addition of Mo in the Cr-Fe-Co-Ni system is found to be very effective in retarding the recrystallization and grain growth, promoting the formation of recrystallized fine grains. Such hierarchical microstructures can be generated in various single-phase fcc HEAs with low SFE by cold-rolling and annealing.

Comparing the mechanical properties of the samples with hierarchical microstructures and fully recrystallized coarse-grained microstructures, the former exhibit exceptional combinations of yield strength-ultimate tensile strength–ductility. As shown in **Figure 7a**, the single-phase HEAs with hierarchical microstructure (annealed

#### **Figure 6.**

*TEM images of the cold-rolled Cr20Fe6Co34Ni34Mo6 HEAs after annealing at (a) 675°C for 1 h; (b) 700°C for 0.5 h; (c) 700°C for 1 h; (d) 800°C for 1 h. (e) Higher magnification micrograph of the interface between fully recrystallized grain and non-fully recrystallized grain, showing high density of dislocations, (f) selected area electron diffraction pattern patterns corresponding to non-fully recrystallized grain in (e) [82].*

### *Microstructures and Deformation Mechanisms of FCC-Phase High-Entropy Alloys DOI: http://dx.doi.org/10.5772/intechopen.104822*

at 675–800°C for 3 min, 0.5 h, 1 h, 5 h, or 10 h) exhibit a very high strength (YS = 0.95– 1.1 GPa and UTS = 1.2–1.3 GPa, **Figure 7a1**) as well as a high ductility (EL is around 30–40%), which is much superior to that of the HEAs (annealed at 800–1150°C for 1 h) with simple fully recrystallized microstructures (**Figure 7a2**). Such high yield strengths exceeding 1.0 GPa at enhanced ductility (30%) in single-phase HEAs are comparable to the nanoprecipitation strengthened HEAs (**Figure 7b1, b2**) [17, 23, 70, 71, 83–89]. This indicates that enhancing yield strength while retaining good ductility of single-phase fcc HEAs can be achieved by developing hierarchical microstructures. The grain boundaries, annealing nanotwins, and dislocation walls play important roles in enhancing the strength, ductility, and strain-hardening capability of the annealed sample with hierarchical microstructure. **Figure 7c** plots the yield strength (σy) as a function of the grain size (d), which follows the well-known Hall–Petch relationship, σy = σ0 + k·d−1/2. The σ0 is 246 MPa. The Hall–Petch coefficient k = 743 MPa·μm1/2 is higher than that of most fcc metals (600 MPa·μm1/2) [90], which indicates that the grain-boundary strengthening mechanism is very effective for improving yield strength. Microstructure characterizations reveal that both twin boundaries and dislocation walls act as strong barriers for dislocation motion, strengthening the alloy, as shown in **Figure 8**. **Figure 8a** shows a large number of stacking faults (marked by red arrows) around twin boundaries. The enlarged HRTEM image in **Figure 8b** shows high-density stacking faults in the twin and matrix, indicating slip transmission associated with dislocations crossing through the coherent twin boundaries (CTBs). Consequently, atomically flat CTBs develop stepped or serrated CTBs that contain interface defects. These interfacial defects along stepped or serrated CTBs then provide sources for nucleating and blocking dislocations. **Figure 8c, d** shows that dislocation

#### **Figure 7.**

*(a1, a2) engineering stress–strain curves of the cold-rolled and annealed Cr20Fe6Co34Ni34Mo6 HEAs at various temperatures of 675–1150°C. (b1, b2) yield strength and ultimate tensile strength versus uniform elongation of annealed alloys compared with various HEAs. (c) the variation of yield strength with average grain size [82].*

**Figure 8.**

*(a, b) TEM micrographs of the recrystallized grains in annealed Cr20Fe6Co34Ni34Mo6 HEAs (700°C for 1 h) at a strain of* ∼*30%. (c, d) TEM micrographs of non-fully recrystallized grains at strains of 15–30% [82].*

walls act as sources for dislocations and deformation twins at large deformation stages, contributing to strain-hardening by continually introducing new interfaces and blocking the motion of dislocations. The enhanced strength enables the production of dislocations at serrated/stepped CTBs, preventing the early onset of necking instability. Therefore, the strength is enhanced without apparent ductility loss.

Other hierarchically heterogeneous microstructures, such as gradient grain structure, gradients in twin and dislocation densities, hierarchical nano-twins, could be introduced into fcc HEAs via surface mechanical grinding treatment, which contributes to the strain hardening capability and thus exceptional strength–ductility combinations, owing to the hetero-deformation induced strengthening. These results demonstrate that engineering the hierarchical microstructure should be an efficient strategy for enhancing the strength and ductility of single-phase fcc HEAs with low and medium SFE.

## **2.4 Shear transformation bands**

Depending on chemical composition and deformation temperature, fcc HEAs with relatively low SFE generate various types of nanoscale shear transformation bands during plastic deformation, including stacking fault bands, nanotwin bands, phase transformation bands, deformation bands with ultrahigh density dislocation, and amorphous bands [10]. The shear transformation banding plays an important role in enhancing strain-hardening capability and thus contributes to enhanced strength and ductility simultaneously. The typical examples are twinning-induced plasticity (TWIP) and transformation-induced plasticity (TRIP) effects in metastable HEAs, which could overcome the long-standing strength–ductility dilemma

## *Microstructures and Deformation Mechanisms of FCC-Phase High-Entropy Alloys DOI: http://dx.doi.org/10.5772/intechopen.104822*

[18, 21, 22, 27]. For example, the strength, ductility, and toughness of single-phase fcc CrMnFeCoNi HEAs can be improved simultaneously with decreasing temperature from room temperature to 77 K, owing to the transition of deformation mechanisms from dislocation slip to nanotwinning [17, 20, 21, 23]. As shown in **Figure 9**, the formation of stacking faults, nanotwinning, and martensitic phase transformation could be generated sequentially in a non-equiatomic CrMnFeCoNi HEA with increasing tensile strains at low temperatures, leading to the formation of various nanoscale bands [91]. Such shear transformation banding, by continually introducing new interfaces and decreasing the mean free path of dislocations during tensile testing ("dynamic Hall–Petch"), produces a high degree of work-hardening and a significant increase in the ultimate tensile strength [92–94].

Recently, amorphous bands (one type of shear transformation bands) have been observed in a non-equiatomic CrMnFeCoNi HEA under tensile deformation at 93 K [91]. These amorphous bands are generated by deformation-induced solid-state amorphization, as summarized in **Figure 10**. The nanoscale amorphous bands, stacking fault bands, nanotwin bands, phase transformation bands can be generated simultaneously in HEAs, which contributes to enhanced strength, ductility, and strain-hardening rate synergistically. In particular, the nanoscale amorphous bands exhibit much higher thermal stability than nanotwins and phase transformation bands. Amorphous bands ensure the enhanced strength and good ductility of hightemperature tempered samples. Amorphous bands can plastically co-deform with the matrix. The interfaces between the amorphous band and fcc matrix provide not only strong barriers for dislocation motion, strengthening materials, but also natural sinks of dislocations, disrupting stress concentrations and delaying decohesion and fracture initiation. These results demonstrate that engineering amorphous bands could be an efficient strategy in remaining enhanced mechanical properties of fcc HEAs at high temperatures.

In addition, Ming et al. design nano-laminated dual-phase structures in a nonequiatomic CrMnFeCoNi HEA via dynamically reversible shear transformations associated with reversible martensitic phase transformation and nanotwinning. The detailed mechanism for dynamically reversible shear transformations can be seen in the following sections. The nano-laminated dual-phase structures could evade the

#### **Figure 9.**

*TEM images of the single-phase fcc CrMnFeCoNi HEA after tension to different strains at 93 K, show the sequential formation of stacking faults, nanotwinning, and martensitic phase transformation [91].*

#### **Figure 10.**

*(a, b) TEM images of the non-equiatomic CrMnFeCoNi HEA after tension to a strain of 36% at 93 K, (c, d) the corresponding high-resolution TEM images showing the amorphous bands. (e–g) TEM images show the interactions of stacking faults, nano-twins, and amorphous bands [91].*

strength–ductility dilemma due to the synergistic operations of the TRIP effect, TWIP effect, and associated with "dynamic Hall–Petch" effect. In a word, the abundant deformation mechanisms of fcc HEAs, as described below, enable us to design various types of microstructures to achieve an exceptional combination of strength and ductility.

## **3. Deformation mechanisms**

Similar to conventional fcc metals and alloys, fcc HEAs plastically deform through three deformation mechanisms: dislocation slip, twinning, and phase transformation. Recently, some unique deformation pathways of deformation-induced amorphization and dynamically reversible shear transformations which rarely occur in conventional materials are found in fcc HEAs.

## **3.1 Deformation-induced amorphization**

Various mechanical processes can partially or fully amorphized crystalline materials, including high-pressure treatments, severe plastic deformation, or mechanical alloying [95–103]. From the thermodynamic driving force perspective, amorphization starts from the massive displacement of atoms into metastable positions and occurs when the free energy of the crystalline phase is higher than that of the amorphous phase [104–106]. From the kinetic hindrance perspective, the formation of a metastable amorphous phase requires the kinetic hindrances which block the formation of a more stable equilibrium crystalline phase [104]. It should be noted that deformation-induced soild-state amorphization under stresses can take place in the elastic regime at an extremely high strain rate, or in the plastic stage at severe strain [95]. At an extremely high strain rate, there is no enough time to activate plastic deformation modes such as dislocation slips or even the faster twinning mode. Thus, the large elastic strains promote the crystalline phase mechanically unstable due to the loss of shear rigidity, which leads to amorphization [107–109].

Interestingly, Ming et al. [91] for the first time observed deformation-induced amorphization in a non-equiatomic Cr26Mn20Fe20Co20Ni14 HEA at cryogenic temperature. Such deformation-induced amorphization was later observed by several groups [110–113]. The deformation-induced amorphization generates extensive nanoscale amorphous bands, as shown in **Figure 10a**–**d**. The formation of extensive nanoscale amorphous bands is attributed to the significant dislocation accumulation in a constrained region inside shear bands, which raises the free energy of the original fcc phase to a point higher than that of the amorphous phase, then the energy difference drives amorphization. Subsequently, the deformation-induced amorphization was also observed in equiatomic CrMnFeCoNi HEA under severe plastic deformation through swaging followed by either quasi-static compression or dynamic deformation in shear [110]. In addition, the deformationinduced localized amorphization can also occur at the tip of cracks, which enhances the toughness significantly by blunting cracks and impeding the expansion of cracks [111]. Subsequently, the nanoscale origin of the mechanism of amorphization is revealed by using molecular dynamics simulation [112, 113]. The amorphization originates from the formation of multi-dislocation junctions due to the low SFE, which results in high lattice resistance to dislocation glide and facilitates nucleation of amorphous nuclei. The deformation mechanisms in the amorphous/crystalline dual-phase regions include highdensity Shockley partial dislocations, multi-dislocation junctions, and nanotwinning in the crystalline region (TEM observations in **Figure 10e**–**g**), as well as radiation-shaped shear bands and amorphous bridges in the amorphous region.

Based on the mechanism of deformation-induced solid-state amorphization in HEAs, nanoscale amorphization bands can be introduced into HEAs to optimize mechanical properties. It is demonstrated that the yield strength of the HEAs could be enhanced without apparent loss of ductility by introducing high-density nanoscale amorphous bands. TEM characterizations (**Figure 10e**-**g**) reveal that introducing nanoscale amorphous bands can achieve three key benefits: (i) amorphous-crystalline interface (ACI) hardening, i.e., ACIs not only act as high-capacity sources for dislocations nucleation but also barriers for dislocation motion; (ii) large stress concentrations at the ACIs can be disrupted and relieved by the amorphous bands since ACIs act as natural sinks of dislocations, averting dislocation pileups at ACIs, which delay decohesion and fracture initiation at the ACIs; (iii) high thermal stability of amorphous bands enables to increase strain-hardening capability through the tempering at relatively higher temperatures without sacrificing the high yield strength. Wang

et al. [111] observed the formation of the amorphous area ahead of the crack tip after amorphization in equiatomic CrMnFeCoNi HEA by using in situ straining TEM experiments and found that the amorphous bridges in the crack wake provided effective toughening of the HEA. Ji et al. [112] provided atomistic insights, via molecular dynamics simulations, into the origin of the solid-state amorphization ahead of a crack tip and report the deformation mechanisms contributing to cryogenic damagetolerance. It is believed that a much better combination of yield strength, ductility, and toughness can be achieved through optimizing the amorphization. These results demonstrate that engineering amorphous bands could be an efficient strategy in remaining enhanced mechanical properties of fcc-structured HEAs.

#### **3.2 Dynamically reversible shear transformations**

Traditional shear transformation banding, such as deformation-induced {111} twinning and martensitic phase transformation (fcc-γ → hcp-ε), has been widely observed in fcc HEAs with low SFE. Recently, Ming et al. [114] found two dynamically reversible shear transformation mechanisms in a CrMnFeCoNi HEA under uniaxial tension at 4.2 K, featured by γ → ε → {1 01 1} twin → γ/γtw and γ → ε → γ/γtw. When deformed at cryogenic temperature, the lower SFE promotes γ → ε shear transformation, forming hcp grains which gradually deplete the original fcc grains. Meanwhile, high-density {0001} stacking faults and {1 01 1} nanotwinning are activated to accommodate plastic deformation, as shown in **Figure 11a, b**. More intriguingly, reverse hcp → fcc shear transformations are stimulated within {1 01 1} twin and surrounding hcp matrix by deformation-induced local dissipative heating (**Figure 11c**). When the {1 01 1} twins transform into fcc structure, Shockley partial dislocations are activated on {111} planes, leading to formation of {111} SFs and {111} nanotwins in the newly formed fcc domain. **Figure 11c** shows two fcc domains with {111} twin orientation inside an {1 01 1} twin. This shear transformation mechanism is described by γ → ε → {1 01 1} twin → γ/γtw.

In addition, for high-density basal stacking faults in the hcp-ε phase, the faulted regions can transform back into fcc structure through correcting SFs via nucleation and glide of Shockley partial dislocations, which is energetically favorable since the process will reduce the density of stacking faults. It is noted that fcc laminates frequently have two orientations, forming a {111} twin orientation relationship (**Figure 11d, e**). This shear transformation mechanism is described by γ → ε → γ/γtw. The reversible fcc ↔ hcp shear transformations and both {1 01 1} and {111} nanotwinning lead to dynamic nano-laminated dual-phase structures, which advance the monotonic "dynamic Hall–Petch" effect in enhancing strength, strain-hardening ability, and ductility by dynamically tailoring the type and width of shear transformation bands.

#### **3.3 Nano-segregation of multi-principal elements**

The unique multiple principal elements endow HEAs with adjustable microstructures and corresponding excellent mechanical properties. Extensive efforts invest in tailoring the chemical compositions of fcc-structured HEAs to optimize the strength and ductility simultaneously. When designing the compositions of HEAs, one should realize two important facts. The first one is that the stronger HEAs are not necessarily the ones with the most elements. The nature of the constituent elements is also important, with the Cr-containing alloys, in general, being the strongest in a family of *Microstructures and Deformation Mechanisms of FCC-Phase High-Entropy Alloys DOI: http://dx.doi.org/10.5772/intechopen.104822*

#### **Figure 11.**

*(a–c) TEM images a CrMnFeCoNi HEA after uniaxial tension to fracture at 4.2 K, showing the reversible shear transformation mechanism featured by γ → ε → {1* 01 *1} twin → γ/γtw. (d, e) the reversible shear transformation mechanism of γ → ε → γ/γtw [114].*

equiatomic binary, ternary and quaternary alloys based on the elements Fe, Ni, Co, Cr, and Mn [71]. Secondly, when designing the compositions of HEAs, nano-segregation of principal elements at grain boundaries (GBs) should be cause for concern since it could deteriorate the ductility significantly. As shown in **Figure 12**, nanoclustering Cr, Ni, and Mn separately at GBs, as detected by atom probe tomography, reduce GB cohesion, and promotes crack initiation along GBs, leading to ductility loss in the CrMnFeCoNi HEA [115]. The GB segregation engineering strategy is then proposed to avoid ductility loss by shifting the fast segregation of principal elements from GBs into preexisting Cr-rich secondary phases. Such GB decohesion by nanoclustering multi-principal elements is a common phenomenon in HEAs. Linlin Li et al. also found that Ni and Mn cosegregate to some regions of the GBs along with the depletion of Fe, Co, and Cr, while Cr is enriched in other regions of the GBs where Ni and Mn are depleted [116–118].

Generally, only nano-segregation of several elements occurs in fully annealed HEAs with coarse and clean grains after short-time annealing treatment, which can lead to GB decohesion and thus ductility loss. However, with very long-time annealing, such as annealing at intermediate temperatures for tens to hundreds of days, intermetallic phases are precipitated at GBs in coarse-grained HEAs [119, 120]. When the grain size of HEAs is decreased to the nanoscale, annealing at intermediate temperatures for mere minutes can lead to precipitation of nanoscale intermetallic phases at GBs, such as Cr-rich phase, NiMn phase, and FeCo phase in equiatomic CrMnFeCoNi HEA [121]. The formation of intermetallic phases at GBs will reduce the ductility significantly. Therefore, the nature of the constituent elements and their nano-segregation behavior should be considered when we design high-strength and high-ductility HEAs.

#### **Figure 12.**

*Atom probe tomography of the GB in a CrMnFeCoNi HEA after uniaxial tension to fracture at 700°C: (a) atom map of the tip showing Cr (24 at %) and Ni (23 at %) isocomposition surfaces viewed with the GB edge-on. (b) 1D compositional profiles along the cyan arrow E and the orange arrow F are indicated in (a). (c) Fracture lateral surface of the tensile sample after tensile tests at 700 °C, showing many cracks along the GBs [115].*

## **4. Conclusion**

This chapter summarizes recent works on realizing strength–ductility combinations in fcc CrFeCoNi-based HEAs via composition and microstructure engineering: (I) Nanoprecipitation strengthening associated with tailoring the size, distribution, and morphology of second phases via alloying a small addition of Ti and Al elements; and (II) The synergistic operations of multiple strengthening mechanisms, such as solid solution strengthening, dislocation strengthening, grain boundary strengthening, precipitation strengthening, TWIP/TRIP effect, and amorphization-induced strengthening. The abundant deformation mechanisms, including slips associated with Shockley partial dislocation and full dislocations, nanotwinning, martensitic phase transformation, deformation-induced amorphization, and dynamically reversible shear transformation, are also summarized. Among them, the recently reported deformation-induced amorphization and dynamically reversible shear transformation are highlighted in terms of their nanoscale origins and strengthening effects for overcoming the strength– ductility trade-off. Finally, this chapter points out that the nature of the constituent

*Microstructures and Deformation Mechanisms of FCC-Phase High-Entropy Alloys DOI: http://dx.doi.org/10.5772/intechopen.104822*

elements and their nano-segregation behavior should be considered when we design high-strength and high-ductility HEAs via engineering the compositions.

## **Acknowledgements**

The authors acknowledge the financial support from the National Natural Science Foundation of China (No. 52002109), the Natural Science Foundation of the Hebei province (No. E2020202088), and the Overseas Scientists Sponsorship Program by Hebei Province (C20210331).

## **Conflict of interest**

The authors declare no competing financial interests.

## **Author details**

Kaisheng Ming1 , Shijian Zheng1 and Jian Wang2 \*

1 State Key Laboratory of Reliability and Intelligence of Electrical Equipment, School of Materials Science and Engineering, Hebei University of Technology, Tianjin, China

2 Mechanical and Materials Engineering, University of Nebraska-Lincoln, Lincoln, NE, USA

\*Address all correspondence to: jianwang@unl.edu

© 2022 The Author(s). Licensee IntechOpen. This chapter is distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

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## **Chapter 2**
