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Cost-Effective Fe-Rich High-Entropy Alloys: A Brief Review

Written By

Yu Yin, Andrej Atrens, Han Huang and Ming-Xing Zhang

Reviewed: 27 April 2022 Published: 08 June 2022

DOI: 10.5772/intechopen.105081

High Entropy Materials IntechOpen
High Entropy Materials Microstructures and Properties Edited by Yong Zhang

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High Entropy Materials - Microstructures and Properties

Edited by Yong Zhang

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Abstract

High-entropy alloys (HEAs) have attracted increased attention due to their extraordinary properties. However, the multicomponent characteristic of equiatomic HEAs inevitably leads to high material costs, which thus limits their widespread industrial applications. Although HEAs are claimed to be suitable for applications in extreme environment due to their comprehensive properties, the actual properties of HEAs dramatically vary with compositions and processes. Therefore, the development of cost-effective HEAs with comprehensive properties is indispensable for industrial uses. Till now, although comprehensive review papers on HEAs are available, few works focused on the cost-effectiveness of HEAs, particularly Fe-rich HEAs recently developed. This review thus aims to fill this gap by reviewing the current research progress in Fe-rich HEAs with a focus on the composition design, microstructure, and properties, including mechanical properties, and resistances for oxidation, wear, and corrosion. The challenges for applying cost-effective Fe-rich HEAs into industries are also arising as future research topics.

Keywords

  • cost-effective
  • high-entropy alloys
  • multicomponent alloys
  • composition complex alloys
  • alloy design
  • microstructure
  • properties
  • application

1. Introduction

Since the concept of high-entropy alloys (HEAs) was first proposed by Yeh and Cantor in 2004 [1, 2, 3], this type of new metallic material has drawn increasingly more attention due to their unique characteristics (e.g. high configuration entropy, sluggish atomic diffusion, and large lattice distortion) and resultant extraordinary properties (e.g. high strength, high low-temperature fracture toughness, good corrosion resistance and high-temperature properties) [4, 5, 6]. Yeh and co-workers firstly defined HEAs as alloys composed of five or more principal elements in equiatomic ratios [1], with single-phase solid-solution (SS) microstructure in either face-centered cubic (FCC) or body-centered cubic (BCC) or hexagonal close-packed (HCP) lattice [4, 5, 6, 7, 8, 9], due to the “high-entropy effect” [10, 11]. At the early stage of research on HEAs, the majority of the works focused on the exploration of single-phase HEAs and their corresponding microstructure and mechanical properties. In recent years, the coverage of HEAs has been extended to include the multiple phase structure with less than five non-equiatomic principal elements [5, 11]. As a result, numerous non-equiatomic alloys were designed and studied based on the equiatomic single-phase HEA system in pursuit of superior mechanical properties [12, 13]. Different terms have been used for these non-equiatomic variants, including multicomponent alloys, multi-principal element alloys and composition complex alloys (CCAs).

It is noteworthy that the current composition design strategy for HEAs usually results in alloys with a high content of expensive metals (e.g. Co, Ta, Hf, V, Nb, W, and Mo). For instance, the above-mentioned HEAs usually contain a Co element, which is common in most of the HEAs [5]. Obviously, using expensive elements increases the overall cost of HEAs. On the other hand, high-entropy alloys usually exhibit comprehensive properties due to their combined effects of multi-principal elements, which makes them suitable for applications in extreme environments. However, majority of the current researches focuses on the mechanical properties only while other properties are ignored. Thus, the development of cost-effective HEAs with comprehensive properties is critical to promote industrial applications of HEAs.

In the past decade, some cost-effective Fe-rich HEAs have been designed and developed, such as the FeNiCrMo alloy that has comprehensive properties [14] and the metastable FeMnCoCr that exhibits superior mechanical properties [12]. Fe-rich HEAs are also named as Fe-rich CCAs [15], Fe-rich medium-entropy alloy (MEA) [16], high-entropy steel [17] and compositionally complex steels [18]. As the entropy value of this type of alloys is not at the high-entropy level, they are generally classified as medium-entropy alloys according to the classification proposed by Yeh [7]. Nevertheless, the entropy of the Fe-rich alloys is indeed higher than traditional steels due to their multicomponent characteristic. To avoid the ambiguity, Fe-rich HEAs are used in this review.

Several reviews on HEAs [4, 5, 6, 19, 20] focused on the microstructure and mechanical properties of different types of HEAs, but there is a lack of reviews with a focus on cost-effective HEAs, particularly on Fe-rich HEAs, that is badly needed to guide industrial applications. The present work would fill the gap through overviewing the composition design, manufacturing, processing, microstructure, and properties of the cost-effective Fe-rich HEAs. Based on the review, key challenges for promoting the industry applications of the cost-effective HEAs are proposed.

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2. Alloy design strategy

In the past decades, various methods have been developed for composition design of traditional alloys, including physical model methods, computational methods and machine learning methods [21, 22]. These design strategies have also been proved effective to assist the development of HEAs, which has been reviewed several times. More details can be found in Refs. [14, 22, 23]. Using these design strategies, many cost-effective HEAs have been developed. For instance, the Co-free HEAs with relatively lower cost include the equiatomic FeMnCrNi, AlCrFeNi [24], FeCrNiMnCu [25], NiMnFeCu [26], AlFeMnSi [27] and FeCrNiTiAl [28] alloys. But these alloy systems are equiatomic or near-equiatomic, there is less potency for further exploitation of the new alloys with better performance. In recent years, several non-equiatomic HEAs with one enriched element have been developed, which includes the Fe-rich HEAs [14], Ni-rich HEAs [29, 30], Cr-rich HEAs [31, 32], Co-rich HEAs [33, 34, 35], Ti-rich HEAs [36] and Al-rich HEAs [37, 38]. Among these non-equiatomic HEAs, Fe-rich HEAs show high potential for industrial applications due to their superior properties and relatively low cost. To the best of the authors’ knowledge, there is no definition for Fe-rich HEAs. Generally, Fe content of these Fe-rich HEAs is in the range of 35–50%, while the content of expensive elements (e.g. Co, Mo, V) is less than 10 at.%. As aforementioned, Fe-rich HEAs are also named Fe-rich CCAs [15], Fe-rich medium-entropy alloy (MEA) [16], high-entropy steel [17] and compositionally complex steels [18]. Up to now, a few Fe-rich HEAs have been developed based on the equiatomic 3d transition metal HEAs, including the FeNiCrMo, FeMnCoCr, FeMnNiCr, FeCoNiCr, FeNiMnAlCr, FeMnNiAlCr and FeMnCrSiNi, through increasing the Fe content in the equiatomic HEAs system. The compositions of these Fe-rich HEAs are listed in Table 1.

Alloy systemComposition (at. %)Phase constituentStatusRef.
FeNiCrMoFe50−xNi25Cr25MoxFCC + IMsAs-cast[14]
Fe50−2xNiCrxMoxFCC + IMsAs-cast[14]
Fe35Ni35Cr25Mo5FCCSolid solution treatment[39]
FeMnCoCrFe80−xMnxCo10Cr10FCC + HCPHot rolling + homogenization[12]
Fe49.5Mn30Co10Cr10C0.5FCC + HCPHot rolling + homogenization[40]
Fe49.5Mn30Co10Cr10C0.5FCCHot rolling + homogenization[41]
FeMnNiCrFe50Mn27Ni10Cr13FCCAs-cast/cold rolling/rolling + annealing[42]
FeMnNiCoCrFe40Mn27Ni26Co5Cr2FCCAs-cast/homogenization[43]
Fe34Mn20Co20Cr20Ni6FCC + HCPHot rolling+homogenization[13]
FeMnNiAlCr(Fe40.4Mn34.8Ni11.3Al7.5Cr6)100-xCxFCCAs-cast[44]
(Fe40.4Mn34.8Ni11.3Al7.5Cr6)100-xBxFCCAs-cast[44]
FeMnNiAlCrFe36Mn21Cr18Ni15Al10BCC + B2As-cast[45]
FeNiCrMoTiCFe40.5Ni22.5Cr22.5Mo4.5Ti5C5FCC + IMs + CarbidesAs-cast/As-aged[46]

Table 1.

The currently reported Fe-rich HEAs.

Previous research work has confirmed that the increasing Fe content may result in different effects on the microstructure and properties in different HEA systems. For instance, Yu and co-workers proposed a new strategy to develop Fe-rich HEAs with eutectic structures [47]. Based on the current pseudo-binary design strategy of eutectic HEAs, the eutectic composition can be experimentally identified by adjusting the atomic content ratio of the intermetallic forming elements to the face-centered cubic (FCC) forming elements in an equiatomic HEA system [47]. As shown in Figure 1, a series of cost-effective Fe-rich HEAs with various compositions (Fe50−xNi25Cr25Mox) and corresponding structures (i.e. the hyper-eutectic, eutectic, hypo-eutectic and near single FCC structure) were produced through decreasing the ratio of the intermetallic forming elements (Mo) to the FCC forming elements (Fe) from 1 to 3/7 [47].

Figure 1.

(a–d) Microstructure of as-cast Fe50−xNi25Cr25Mox alloys: (a) x = 20, hyper-eutectic; (b) x = 15, fully-eutectic; (c) x = 10, hypo-eutectic; (d) x = 5, divorced-eutectic; (e) XRD analyses of Fe50−xNi25Cr25Mox HEAs [47].

The strategies proposed by Yu and co-workers [47] is suitable for HEAs with intermetallic-dominated phase. For HEAs with a single solid solution phase, the increase in the ratio of Fe to other elements may not result in the formation of a eutectic structure. Instead, it may change the phase constituents of the solid solution phase matrix. For instance, by increasing the Fe content in the Fe80−xMnxCo10Cr10 [12] and Fe40−xMn20Co20Cr20Nix [13] HEAs, the structure of the alloys shifted from single FCC phase to dual-phase (FCC and HCP). By increasing the content of Fe or decreasing the content of other elements in an equiatomic HEA, the thermal stability and mechanical stability can be decreased due to the reduced “high-entropy effect”, which may result in precipitation during heat treatment [48] or TRIP effect during deformation [13], which will be further discussed in Sections 3 and 4.

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3. Microstructure

3.1 Microstructure of the as-cast and homogenization treated alloys

The currently reported Fe-rich HEAs can be classified into different groups in terms of their phase constituent, namely single or near single FCC phase alloys, FCC/HCP dual-phase alloys and BCC/B2 dual-phase alloys. The FCC single phase and FCC/HCP dual-phase Fe-rich HEAs have been comprehensively studied while very few researches focus on the BCC/B2 dual-phase Fe-rich HEAs.

As shown in Table 1, the Fe-rich HEAs with single FCC phase include the Fe50Mn27Ni10Cr13, Fe40Mn27Ni26Co5Cr2 and Fe40.4Mn34.8Ni11.3Al7.5Cr6. The Fe50−xNi25Cr25Mox HEAs have a near single FCC structure matrix with an intermetallic phase. Yao and co-workers developed a non-equiatomic Fe40Mn27Ni26Co5Cr2 alloy based on the equiatomic FeMnNiCoCr Cantor alloy and their microstructure is shown in Figure 2a–f [42]. The EBSD and EDS results of the recrystallized alloy in Figure 2 reveals that the non-equiatomic variant of Cantor alloy has a homogenized solid solution matrix despite its comparably low configurational entropy [42], indicating that equimolarity may not be a compulsory requirement to achieve single-phase solid solution in multicomponent systems [42]. A similar conclusion can be drawn from Yu and co-workers’ work. As shown in Figure 1, with the increasing of the ratio of Fe to Mo in the FeNiCrMo HEA, the equiatomic FeNiCrMo was transformed into a non-equiatomic alloy, resulting in a lower configuration entropy value of the alloy system. However, the equiatomic FeNiCrMo is composed of an intermetallic phase while the non-equiatomic Fe45Ni25Cr25Mo5 alloy with the lowest entropy value shows the near FCC phase with low fraction of intermetallic (Figure 2g–i). Obviously, this cannot be explained using the “high-entropy effect” proposed by Yeh and co-workers. This can be related to the decreased atom size difference due to the increasing of Fe content, which is usually ignored by the current “high-entropy effect”. As a consequence, much more compositions of non-equiatomic HEAs can be designed to gain single solid solution phase by decreasing the atom size difference even the maximum entropy value may not be achieved.

Figure 2.

(a–f) Microstructure of the recrystallized Fe40Mn27Ni26Co5Cr2: a. OM image, b. SE image, c and d. EBSD phase maps and e and f. EDS maps [42]; (g–i) Microstructure of the as-cast Fe45Ni25Cr25Mo5: g. SEM image, h. EDS mapping and i. XRD spectrum [15]; (j–m) Microstructure of the homogenized and water-quenched Co20Cr20Fe34Mn20Ni6 HEA: j. EBSD phase map; k and l. ECCI image; m. EDS maps [13]; (n and o) Microstructure of the Fe36Mn21Cr18Ni15Al10 alloy in the as-cast condition: a. BSE image, c. bright-field TEM image [45].

The current dual-phase Fe-rich HEAs were developed based on the equiatomic FeMnCoCr and FeMnCoCrNi. Increasing the ratio of Fe to Mn in the FeMnCoCr HEA and the ratio of Fe to Ni in the FeMnCoCrNi HEA decreased the phase stability of single solid solution phase. Hence, both the non-equiatomic Fe80−xMnxCo10Cr10 (x < 30 at.%) and Fe40−xMn20Co20Cr20Nix (x < 6 at.%) systems involved partial martensitic transformation from FCC to the HCP phase during quenching (Figures 2j–m and 3). As shown in Figure 2j–m, the homogenized Co20Cr20Fe34Mn20Ni6 alloy after water quenching shows an FCC and HCP dual-phase structure with homogeneous composition. The SEM-BSE image reveals a high density of stacking faults in the FCC phase, indicating a very low stacking fault energy of the new alloy. The effect of the metastable phase on the mechanical properties will be further discussed in Section 4.1.

Figure 3.

XRD patterns and EBSD phase maps of Fe80−xMnxCo10Cr10 (x = 45 at.%, 40 at.%, 35 at.% and 30 at.%) HEAs [12]; XRD patterns of homogenized Co20Cr20Fe40−xMn20Nix (x = 20 at.%, 6 at.% and 0) HEAs reveal the variations of phase configurations with changing the x value [13].

In addition, in contrast to the FeMnCrNiAl variant (Fe40.4Mn34.8Ni11.3Al7.5Cr6) with a single FCC phase, a Fe-rich HEA (Fe36Mn21Cr18Ni15Al10) with BCC/B2 dual-phase structure has been developed. As shown in Figure 2n–o, the as-cast alloy has a dual-phase structure consisting of a BCC matrix and homogeneously distributed cuboidal B2 ordered particles. The BCC/B2 dual-phase structure has been widely reported in the refractory HEAs [49, 50], which usually exhibit superior high-temperature properties. However, nearly all the current refractory HEAs (e.g. MoNbTaW, MoNbTaVW and HfNbTaTiZr) contain expensive elements, such as the Ta, Hf, V, Nb, Mo, W. In contrast, the newly developed BCC/B2 dual-phase HEA (Fe36Mn21Cr18Ni15Al10) is Fe-enriched and only contains inexpensive elements, which sheds light on the development of cost-effective refractory HEAs.

3.2 Microstructure after annealing

Although HEAs with single-phase were initially considered more stable due to the high-entropy effect, such as the CoCrFeMnNi HEA, they are recognized as a supersaturated solid solution at intermediate temperature in recent years because precipitation of the second phase or decomposition in the single solid solution phase occurs in the alloys. Decomposition was observed in both the FCC and BCC HEAs, such as the CoCrFeMnNi [51] with FCC structure and HfNbTaTiZr HEA [52, 53] with BCC structure, after long-period annealing at intermediate temperatures. It is believed that the annealing at intermediate temperatures for a long-time results in a weakened “high-entropy effect” due to the temperature-dependent contribution of mixing entropy to total Gibbs free energy and thus decreases the phase stability of the solid solution phase.

For Fe-rich HEAs, the phase stability of their solid solution can be further weakened due to the non-equiatomic and reduced entropy value, including thermal stability and mechanical stability. As aforementioned, the decreased thermal stability of the solid solution phase results in phase transformation during cooling, such as the partial martensitic transformation from FCC to the HCP phase when quenching as shown in Figures 2j and 3. The decreased mechanical stability will introduce phase transformation during deformation. Both the reduced thermal stability and mechanical stability contribute to the improvement of mechanical properties, which will be further discussed in Section 4.1. In addition, the decreased thermal stability may result in the precipitation of the second phase in the metastable solid solution phase, which can be used to manipulate the microstructure and thus their mechanical properties. For instance, Yu and co-workers investigated the precipitation behavior of the cost-effective Fe45Ni25Cr25Mo5 HEA with a face-cantered cubic (FCC) matrix [48]. As shown in Figure 4, with the increasing aging time at 900°C, the volume fraction of needle-shaped precipitates increased significantly in the FCC matrix of Fe45Ni25Cr25Mo5 HEA. This resulted in a high age-hardening effect, which raised the hardness from 192 HV5 to nearly 300 HV5. However, the peak-aged sample exhibited room-temperature brittleness due to the precipitation of a large needle-shaped intermetallic phase. To decrease the brittleness, thermomechanical processing can be applied to decrease the size of precipitates, which will be discussed in Section 3.3.

Figure 4.

(a and b) Novel heterogeneous lamella (HL) structure in the Fe35Ni35C25Mo5 HEA: STEM bright-field image of the thin-foil sample. The yellow arrows indicate the σ precipitates (a) [39]; Schematic illustration of the HL microstructure with nanoprecipitates and twins (b). (c) Schematic plot of producing hierarchical grain structure in the Fe49.5Mn30Co10Cr10C0.5 alloy [41].

Precipitation or phase decomposition is also reported in the dual-phase Fe-rich HEAs, for instance, the Fe36Mn21Cr18Ni15Al10 HEA with BBC and B2 phases. Interestingly, annealing at a high temperature of 1200°C for 24 h barely changed the structure (Figures 2n–o and 5h), but annealing at a relatively lower temperature of 1000°C resulted in precipitation of the FCC phase along the grain boundaries and within the matrix (Figure 5i–j). Due to the presence of the soft and ductile FCC phase at temperatures lower than 1000°C, the alloy is softened and thus may not be suitable for applications under 1000°C. However, the nanometer scaled cuboid B2 phase remains unchanged at 1200°C for 24 h, which is abnormal. Generally, high-temperature treatment will result in the coarsening of the second phase. This suggests a high stability of the nanometer scaled dual-phase structure, and therefore maintaining superior mechanical properties at temperatures close to 1200°C. Future work is needed to verify the high stability and evaluate the high-temperature (>1000°C) properties.

Figure 5.

(a–f) SEM images of the Fe45Ni25Cr25Mo5 samples aged at 900°C for different hours; (g) EDS map of the peaked aged Fe45Ni25Cr25Mo5 HEA [48]; (h–j) Microstructure of the Fe36Mn21Cr18Ni15Al10 alloy after annealing at 1200 (h) and 1000°C (i, j) for 24 h: h and i. BSE images, j. EBSD phase map (FCC phase in red and BCC/B2 phase in green) [45].

3.3 Microstructure after the thermomechanical process

In contrast to the simplicity of microstructure in the as-cast and as-aged Fe-rich HEAs, microstructure after thermomechanical processing can be more diverse, which provides much more opportunities for microstructure and mechanical properties manipulation. Similar to traditional alloys or equiatomic HEAs, thermomechanical processing is performed to control the grain size, phase constituents or phase fraction in Fe-rich HEAs (Fe80-xMnxCo10Cr10 and Fe49.5Mn30Co10Cr10C0.5) [12, 13, 40, 54], which corresponds to the grain-refinement strengthening, TWIP or TRIP effect. Except for the above traditional microstructure control methods, heterogeneous or hierarchical microstructure design (e.g. heterogeneous lamella structures, gradient structures, laminate structures, and harmonic structures) [19, 55, 56] has been proved effective in the property improvement of HEAs. This has been an active research topic.

Unique heterogeneous lamella (HL) structure was introduced in a cost-effective FCC HEA (Fe35Ni35Cr25Mo5) through a single-step heat treatment (800°C for 1 h) after cold rolling. As shown in Figure 4a and b, the HL structure consists of alternative layers of coarse-grained FCC matrix, and ultra-fine grains or subgrains layer with nanoprecipitates and annealing twins (ATs). According to Yu and co-workers, the preferential precipitation of σ phase at the shear bands with a high density of lattice defects (e.g. high-density dislocation walls and nano deformation twins) restrict the growth of recrystallized grains, resulting in the partial recrystallization and thus the formation of HL structure [39]. In addition, Su and co-workers [41] demonstrated a hierarchical microstructure design strategy to improve the mechanical properties of an Fe-rich HEA (Fe49.5Mn30Co10Cr10C0.5) by a thermomechanical processing. As shown in Figure 4c, three distinguished regions with different levels of dislocation density formed in the as-homogenized FCC matrix after cold rolling, including the deformation-induced HCP phase with the lowest dislocation density, residual FCC region with medium dislocation density, and severe shear bands region with the highest dislocation density. After annealing at a temperature above 400°C, the HCP phase is reversed to FCC phase while twins which co-existed with martensite lamellae got thickened in the parent grains. Due to the different dislocation densities and thus different recrystallization kinetics, trimodal grain structures were finally produced and characterized by small recrystallized grains associated with shear bands, medium-sized grains recrystallized from parent grains and unrecrystallized large grains. Such a grain size hierarchy promotes the variation in phase stability and results in a joint activation of transformation-induced plasticity (TRIP) and twinning-induced plasticity (TWIP) effects upon loading and thus a good strength-ductility synergy [41].

The above findings indicate that the non-equiatomic alloys (e.g. Fe-rich HEAs) not only enable more alloys with different compositions to be designed, but also broaden the window for microstructure and properties tuning due to the reduced thermal and mechanical stability of the solid solution phase.

3.4 Microstructure during deformation

As aforementioned, the decreased thermal stability and mechanical stability of Fe-rich HEAs can be applied to manipulate the microstructure via different deformation mechanisms (e.g. dislocation slip, twinning and the formation of stacking faults) and strengthening effects (e.g. TRIP effect, precipitation strengthening or integrated strengthening effect), and thus to achieve better mechanical properties. For instance, Zhiming and co-workers [12] reported that the improvement in the mechanical property of the metastable DP HEA (Fe50Mn30Co10Cr10) is related to the thermally induced DP structure and mechanically induced HCP phase [12]. As shown in Figure 6a and b, a large number of stacking faults is observed in the metastable FCC phase in the Fe50Mn30Co10Cr10 HEA, which acted as phase-formation nuclei (FCC → HCP) during deformation. When strain is lower than 30%, the stress-induced transformation from the FCC to HCP phase is the dominant deformation mechanism. With increasing strain to over 30%, the density of stacking faults and nano-twins increased in both the initial and mechanically induced HCP phase, which contributes significantly to strain hardening. Further increase of the strain to over 45% resulted in the formation of a high density of dislocations. Therefore, the thermally and mechanically induced HCP phase plays an important role in plastic accommodation and hardening at later stages of deformation via multiple deformation mechanisms, including dislocation slip, twinning, and the formation of stacking faults [12]. In addition, Yu and co-workers [39] introduced a unique heterogeneous lamella (HL) grain structure with a high density of nanoprecipitates, annealing twins and low angle boundaries in the Fe35Ni35Cr25Mo5 HEA (Figure 4a and b), resulting in a superior tensile property, with yield strength over 1.0 GPa and total elongation of ~13%. It is reported that the superior tensile properties were resulted from the hetero-deformation induced (HDI) strengthening, precipitation strengthening and the coexistence of multiple deformation mechanisms [39]. As shown in Figure 6c–f, the as-deformed Fe35Ni35Cr25Mo5 HEA shows a complex hierarchical microstructure including stacking faults (SFs), dislocation, annealing/deformation twins (DTs), low-angle grain boundaries (LAGBs), nanoprecipitates and HL interfaces. During deformation, the hierarchical microstructure (a high density of LAGBs, HL interfaces and precipitates) in the HL Fe35Ni35Cr25Mo5 alloy interact with the subsequently activated dislocations, SFs and DTs and thus significantly promote the strain hardening ability and prevent early necking.

Figure 6.

(a and b) Deformation micro-mechanisms in the TRIP-DP-HEA with increasing tensile deformation at room temperature [12]; (c–f) Microstructural evolution upon tensile deformation in the HL Fe35Ni35Cr25Mo5 HEA [39].

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4. Properties

4.1 Mechanical properties

The room-temperature tensile properties (yield strength and fracture strain) of the current cost-effective Fe-rich HEAs/CCAs and a few commercial steels (for comparison purpose) are summarized in Figure 7. Obviously, the tensile properties of the current Fe-rich HEAs significantly vary, depending on their composition and microstructure. The widely reported strength-ductility trade-off can be identified in both the Fe-rich HEAs and commercial steels. Most Fe-rich HEAs or steels with high yield strength correspond to low elongation and vice versa. For instance, the ductile Fe50Mn30Co10Cr10 alloy with high elongation of over 70% has a yield strength of less than 500 MPa [12], while the Fe35Ni35Cr25Mo5 HEAs [39] and Fe49.5Mn30Co10Cr10C0.5 [41] after thermomechanical process exhibit a very high strength over 1 GPa but with ductility around 10%. Some HEAs exhibit balanced strength and ductility, such as the Fe49.5Mn30Co10Cr10C0.5 HEA [41]. After cold rolling and annealing at 650°C for 3 min, a good combination of yield strength (824 MPa), UTS (1.05 GPa) and ductility (33%) are attained due to the hierarchical grain structure (Figure 4c) and the sequential activation of transformation-induced plasticity (TRIP) and twinning-induced plasticity (TWIP) effects during deformation [41]. Some HEAs also exhibit quite low tensile performance, for instance, the BCC/B2 dual-phase Fe36Mn21Cr18Ni15Al10 HEA with severe brittleness, the elongation of which is less than 3%. Furthermore, most FCC Fe-rich HEAs exhibit very low strength (<250 MPa) at as-cast condition, which is close to that of the 316 L stainless steel.

Figure 7.

Tensile properties of most currently reported Fe-rich HEAs [12, 13, 14, 39, 40, 41, 42, 43, 44, 45, 46] compared with different types of steels (Data of steels is from Ref. [57, 58]).

Compared with commercial steels, the mechanical properties of the currently reported Fe-rich HEAs overlap most of the commercial steels, but are still lower than some high strength steels, particularly the maraging steels (number 2 in Figure 6) and D&P steels (number 1 in Figure 6) with ultra-high strength [57, 59, 60]. For instance, the low-cost D&P steel (FeMn9.95C0.44Al1.87V0.67) possesses a superior tensile property, with the yield strength and elongation of nearly 2 GPa and 22.0%, respectively [60]. Thus, it is essential to improve the mechanical properties of Fe-rich HEAs to further enhance their application potential.

4.2 Wear resistance

Except for the tensile or compressive properties, data of other mechanical properties, such as wear resistance, impact toughness, fatigue properties, or creep resistance, of Fe-rich HEAs is very limited. Based on the cost-effective age-hardenable Fe45Ni25Cr25Mo5 HEA, Yu and co-workers developed an intermetallics and carbides reinforced Fe40.5Ni22.5Cr22.5Mo4.5Ti5C5 HEA with superior wear resistance [46]. As shown in Figure 8a–d, the as-cast FeNiCrMoTiC alloy consists of an FCC solid solution matrix with randomly-distributed carbides and FCC/intermetallics eutectic structures. Aging at 800°C for 96 h effectively increases the hardness (Figure 8e) and wear resistance (Figure 8f) of the alloy due to the precipitation strengthening of intermetallics. It is noteworthy that, although the hardness of high-chromium cast iron (HCCI) is much higher than that of the peak-aged FeNiCrMoTiC alloy, the wear resistance of the latter is superior (Figure 8f). This is attributed to their different wear behavior during dry friction (Figure 8g). In contrast to the severe delamination in the HCCI, the peak-aged FeNiCrMoTiC alloy shows moderate abrasive wear and minor delamination under sliding friction due to its combined effect of ductile FCC matrix and fine reinforced particles, including the in-situ formed carbides/eutectic structures and precipitates formed during aging treatment [46]. On one hand, the fine reinforced particles with lower cracking susceptibility effectively strengthen the soft FCC matrix and thus reduce both the material loss by abrasive wear and severe brittle delamination. Moreover, the spalled fine particles (e.g. carbides or intermetallics) were found welded back into the FCC phase during friction and thus further strengthening the matrix and decreasing the abrasive wear. On the other hand, the propagation of micro-cracks from the brittle particles is inhibited by the ductile FCC matrix, which also suppresses the severe brittle delamination [46].

Figure 8.

SEM-BSE images (a–b) and the corresponding EDS map (c) and line-scan results (d) of the as-cast Fe40.5Ni22.5Cr22.5Mo4.5Ti5C5 HEC; (e) Hardness aging curves of the HEC at 800°C for up to 240 h, with inserted BSE images of the as-cast sample and peak-aged (800°C, 96 h) HEC; (f) Wear resistance of the HECs and destabilized HCCI are plotted against their hardness values; (g) Comparison of wear mechanism of the HCCI and as-aged HEC [46].

4.3 Oxidation and corrosion resistance

As shown in Table 1, all the currently reported Fe-rich HEAs contain Cr, most of which is over 10 at.%. Generally, alloys with high content of Cr show high oxidation and corrosion resistance. It is reasonable to assume that most currently reported Fe-rich HEAs possess superior oxidation and corrosion resistance. However, corrosion behavior of Fe-rich HEAs is rarely reported. Yu and co-workers evaluated the high-temperature oxidation resistance of the Fe-rich Fe45Ni25Cr25Mo5 HEA in comparison to two commercial alloys (i.e. 316 L stainless steel and Inconel 625 superalloy) [15]. In contrast to the catastrophic oxidation behavior of 316 L stainless steel and severe oxide spallation on Inconel 625 superalloy, the Fe-rich HEA showed outstanding oxidation resistance at 1200°C, including low oxidation rate and high spallation resistance. As shown in Figure 9a–l, catastrophic oxidation and spallation occurred on the 316 L stainless steel at 1200°C mainly due to the high oxidation rate and the intergranular cracking in the coarse-grained FeO scale. For Inconel 625 superalloy, the high densities of pores, cracks and Nb oxides in the Cr2O3 scale promoted the oxidation process. Meanwhile, the selective formation of Nb2O5 along the Cr2O3 scale/substrate interface resulted in severe oxide spallation and thus reduction of the oxidation resistance [15]. In contrast, an exclusive and compact chromia scale with better mechanical properties (e.g. high hardness and elastic modulus) formed on the Fe-rich HEA, protecting the matrix from further oxidation and high resistance against interface cracking during the oxidation process. As a result, the Fe-rich HEA showed the lowest oxidation rate at 1200°C compared with that of 316 L stainless steel and Inconel superalloy (Figure 9m) [15].

Figure 9.

EBSD analysis of the FeO oxide scale formed on the 316 L SS (a–d), on the Inconel 625 superalloy (e–h) and the Fe45Ni25Cr25Mo5 CCA (i-l) after oxidation at 1200°C in the air for 48 h: (a, e, and i) band contrast map; (b, f, and j) IPF map; (c, g, and k) phase map; (d, h, and l) KAM map; The isothermal oxidation kinetics of the Fe-rich CCA as compared with the 316 L SS and the Inconel 625 alloy at 1200°C in the air for 48 h [15].

Yu and co-workers also qualitatively evaluated the corrosion resistance of the Fe45Ni25Cr25Mo5 and Fe40.5Ni22.5Cr22.5Mo4.5Ti5C5 alloys using the immersion experiment [46, 48], both of which have superior corrosion resistance in an acid environment. As shown in Figure 10a, after immersion in the 5% HCl solution at room temperature for 10 days, obvious surface corrosion occurred on the 316 L stainless steel, but no visible corrosion was found on the Fe45Ni25Cr25Mo5 HEA [48]. Similarly, catastrophic corrosion was identified in the stainless steel while only minor corrosion pits are visible on the Fe40.5Ni22.5Cr22.5Mo4.5Ti5C5 HEA after immersion in 20% HCl solution for 150 days (Figure 10b) [46]. In-depth future work about the corrosion mechanisms of the Fe-rich HEAs is necessary.

Figure 10.

(a) Surface morphologies of the Fe45Ni25Cr25Mo5 HEA and the 316 L SS after immersion in the 5%HCl solution at room temperature for 10 days. (b) Surface morphologies of the aged Fe40.5Ni22.5Cr22.5Mo4.5Ti5C5 HEA and 316 L SS before and after immersion in the 20% HCl solution at room temperature for 150 days [46, 48].

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5. Conclusions and outlook

Fe-rich HEAs exhibit a high potential for industrial applications owing to their superior properties and relatively low cost. Upon reviewing the currently reported Fe-rich HEAs the following conclusions are achieved.

  1. The Fe-rich HEAs reported most recently were developed based on the 3d transition-metal HEAs, including single or near single FCC, FCC/HCP dual-phase and BCC/B2 dual-phase alloys. The FCC and FCC/HCP dual-phase Fe-rich HEAs have been widely studied, but few researches focused on the BCC/B2 dual-phase Fe-rich HEAs.

  2. The compositions of various Fe-rich HEAs can be designed by simply increasing the Fe content in the equiatomic HEAs system, which may result in different effects on the microstructure and properties of different HEA systems. By increasing the content of Fe or decreasing the content of other elements in equiatomic HEAs, the thermal stability and mechanical stability are reduced because of lowering the “high-entropy effect”. The decreased thermal stability of the solid solution phase facilitates precipitation during heat treatment or phase transformation (e.g., FCC to HCP transformation) during cooling. The decreased mechanical stability enables stress-induced phase transformation during deformation. Like conventional alloys, these phenomena can be used to tailor the microstructure, introduce different deformation mechanisms (e.g., dislocation slip, twinning and the formation of stacking faults) and strengthening effects (e.g. TRIP effect, precipitation strengthening or integrated strengthening effect), and thus achieve improved mechanical properties.

  3. Like the processing of traditional alloys or equiatomic HEAs, a thermomechanical process can be used to tailor the grain size, phase constituents or fraction of the Fe-rich HEAs (e.g., Fe80-xMnxCo10Cr10 and Fe49.5Mn30Co10Cr10C0.5), in order to improve the mechanical properties of the alloys.

  4. The tensile properties of the currently reported Fe-rich HEAs vary within a wide range, depending on their composition and microstructure. The widely reported strength-ductility trade-off can be identified in both the Fe-rich HEAs and commercial steels. The mechanical properties of the currently reported Fe-rich HEAs are better than most of the commercial steels, but are still lower than some high strength steels, such as the maraging steels and D&P steels with ultra-high strength.

  5. Except for the tensile or compressive properties, data of other mechanical properties of Fe-rich HEAs is limited, such as wear resistance, impact toughness, fatigue properties or creep resistance. In addition, although most Fe-rich HEAs are supposed to possess superior oxidation and corrosion resistance, very few research results are reported.

Some Fe-rich HEAs exhibit superior properties including mechanical properties, and oxidation and corrosion resistance. However, the greatest challenge is how to commercialize such alloys for industrial application. Systematic and comprehensive research is needed, which should focus more on the aspects of composition design, microstructure control, and properties evaluation and improvement. Here, combining the opinions proposed by other experts in the field of HEAs [23, 61] the authors propose a few topics that are of particular significance for the application of Fe-rich HEAs.

  1. Except for the current 3d transition metal Fe-rich HEAs with FCC or FCC/HCP dual-phase structure, the Fe-rich HEAs with BCC structure should also be developed. Such studies are expected to shed some light on the development of low-cost refractory alloys. This can be achieved by introducing refractory elements (i.e., Ti, V, Zr, Nb, Hf, and W) or other elements (e.g., Al, Si) with a low valence electron concentration (VEC) value. Due to the high strength of the BCC phase at elevated temperatures, the cost-effective Fe-rich HEAs with BCC structure can be a new type of promising refractory HEAs. For instance, the Fe-rich HEAs (Fe36Mn21Cr18Ni15Al10) exhibit a very stable fine BCC/B2 DP structure at 1200°C, which can be a cost-effective alloy for high-temperature application [45].

  2. The mechanical properties of the Fe-rich HEAs developed need further improvement. Different strategies should be explored to optimize their composition, microstructure, and process. For instance, other than C, the addition of other interstitial elements e.g., N, B, O, in Fe-rich HEAs may be used for property improvement. In addition, thermomechanical processing can be an effective method for microstructure control and properties manipulation.

  3. The properties of Fe-rich HEAs need to be evaluated comprehensively to explore their potential applications in extreme environments. Except for the most widely reported room-temperature tensile properties, other mechanical properties like low/high-temperature tensile properties, impact toughness, fatigue resistance, creep resistance and wear resistance should be characterized. Further exploration of other properties like oxidation and corrosion resistance should also be conducted.

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Acknowledgments

This work was financially supported by the ARC Discovery Project (No. DP200101408).

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Conflict of interest

The authors declare no conflict of interest.

References

  1. 1. Yeh JW, Chen SK, Lin SJ, Gan JY, Chin TS, Shun TT, et al. Nanostructured high-entropy alloys with multiple principal elements: Novel alloy design concepts and outcomes. Advanced Engineering Materials. 2004;6(5):299-303
  2. 2. Cantor B, Chang ITH, Knight P, Vincent AJB. Microstructural development in equiatomic multicomponent alloys. Materials Science and Engineering A. 2004;375:213-218
  3. 3. Huang PK, Yeh JW, Shun TT, Chen SK. Multi-principal-element alloys with improved oxidation and wear resistance for thermal spray coating. Advanced Engineering Materials. 2004;6(12):74-78
  4. 4. George EP, Raabe D, Ritchie RO. High-entropy alloys. Nature Reviews Materials. 2019;4(8):515-534
  5. 5. Miracle DB, Senkov ON. A critical review of high entropy alloys and related concepts. Acta Materialia. 2017;122:448-511
  6. 6. Zhang Y, Zuo TT, Tang Z, Gao MC, Dahmen KA, Liaw PK, et al. Microstructures and properties of high-entropy alloys. Progress in Materials Science. 2014;61:1-93
  7. 7. Ye YF, Wang Q , Lu J, Liu CT, Yang Y. High-entropy alloy: Challenges and prospects. Materials Today. 2016;19(6):349-362
  8. 8. Chen J, Zhou X, Wang W, Liu B, Lv Y, Yang W, et al. A review on fundamental of high entropy alloys with promising high-temperature properties. Journal of Alloys and Compounds. 2018;760:15-30
  9. 9. George EP, Curtin WA, Tasan CC. High entropy alloys: A focused review of mechanical properties and deformation mechanisms. Acta Materialia. 2020;188:435-474
  10. 10. Tsai MH, Yeh JW. High-entropy alloys: A critical review. Materials Research Letters. 2014;2(3):107-123
  11. 11. Yeh J-W. Recent progress in high-entropy alloys. Annales de Chimie Science des Matériaux. 2006;31(6):633-648
  12. 12. Li Z, Pradeep KG, Deng Y, Raabe D, Tasan CC. Metastable high-entropy dual-phase alloys overcome the strength-ductility trade-off. Nature. 2016;534(7606):227-230
  13. 13. Li ZM, Kormann F, Grabowski B, Neugebauer J, Raabe D. Ab initio assisted design of quinary dual-phase high-entropy alloys with transformation-induced plasticity. Acta Materialia. 2017;136:262-270
  14. 14. Yin Y. Development of cost-effective high-entropy alloys with superior mechanical properties. University of Queensland. 2021:1-224
  15. 15. Yin Y, Tan Q , Zhao Y, Sun Q , Shi Z, Bermingham M, et al. A cost-effective Fe-rich compositionally complicated alloy with superior high-temperature oxidation resistance. Corrosion Science. 2021;180:109190
  16. 16. Bae JW, Seol JB, Moon J, Sohn SS, Jang MJ, Um HY, et al. Exceptional phase-transformation strengthening of ferrous medium-entropy alloys at cryogenic temperatures. Acta Materialia. 2018;161:388-399
  17. 17. Raabe D, Tasan CC, Springer H, Bausch M. From high-entropy alloys to high-entropy steels. Steel Research International. 2015;86(10):1127-1138
  18. 18. Wang Z, Lu W, Zhao H, Liebscher CH, He J, Ponge D, et al. Ultrastrong lightweight compositionally complex steels via dual-nanoprecipitation. Science Advances. 2020;6(46):eaba9543
  19. 19. Sathiyamoorthi P, Kim HS. High-entropy alloys with heterogeneous microstructure: Processing and mechanical properties. Progress in Materials Science. 2020;123:100709
  20. 20. Li Z, Zhao S, Ritchie RO, Meyers MA. Mechanical properties of high-entropy alloys with emphasis on face-centered cubic alloys. Progress in Materials Science. 2019;102:296-345
  21. 21. Li J-H, Tsai M-H. Theories for predicting simple solid solution high-entropy alloys: Classification, accuracy, and important factors impacting accuracy. Scripta Materialia. 2020;188:80-87
  22. 22. Murty BS, Yeh J-W, Ranganathan S, Bhattacharjee P. High-ntropy Alloys. Amsterdam: Elsevier; 2019
  23. 23. Körmann F, Li Z, Raabe D, Sluiter MHF. Iron-rich high entropy alloys. In: Rana R, editor. High-Performance Ferrous Alloys. Cham: Springer International Publishing; 2021. pp. 389-421
  24. 24. Wu H, Xie J, Yang H-Y, Shu D-L, Hou G-C, Li J-G, et al. Comparative study of mechanical and corrosion behaviors of cost-effective AlCrFeNi high entropy alloys. Journal of Materials Engineering and Performance. 2022;396:1-11
  25. 25. Rao ZY, Wang X, Zhu J, Chen XH, Wang L, Si JJ, et al. Affordable FeCrNiMnCu high entropy alloys with excellent comprehensive tensile properties. Intermetallics. 2016;77:23-33
  26. 26. Kumar A, Mucalo M, Bolzoni L, Li Y, Kong F, Yang F. Fabrication, microstructure, mechanical, and electrochemical properties of NiMnFeCu high entropy alloy from elemental powders. Metals. 2022;12(1):167
  27. 27. O’Brien SP, Christudasjustus J, Esteves L, Vijayan S, Jinschek JR, Birbilis N, et al. A low-cost, low-density, and corrosion resistant AlFeMnSi compositionally complex alloy. npj Materials Degradation. 2021;5(1):12
  28. 28. Liu Y, Qu N, Zhao X, Chen J, Zhu J, Lai Z. Stability of FeCrNiTiAl high-entropy alloy at high temperature. Heat Treatment and Surface Engineering. 2021;3(1):29-36
  29. 29. Jin X, Bi J, Zhang L, Zhou Y, Du X, Liang Y, et al. A new CrFeNi2Al eutectic high entropy alloy system with excellent mechanical properties. Journal of Alloys and Compounds. 2019;770:655-661
  30. 30. Jin X, Zhou Y, Zhang L, Du X, Li B. A novel Fe 20 Co 20 Ni 41 Al 19 eutectic high entropy alloy with excellent tensile properties. Materials Letters. 2018;216:144-146
  31. 31. Fang W, Yu H, Chang R, Zhang X, Ji P, Liu B, et al. Microstructure and mechanical properties of Cr-rich Co-Cr-Fe-Ni high entropy alloys designed by valence electron concentration. Materials Chemistry and Physics. 2019;238:121897
  32. 32. Cho K, Fujioka Y, Nagase T, Yasuda HY. Grain refinement of non-equiatomic Cr-rich CoCrFeMnNi high-entropy alloys through combination of cold rolling and precipitation of sigma phase. Materials Science and Engineering A. 2018;735:191-200
  33. 33. Wei D, Li X, Heng W, Koizumi Y, He F, Choi W-M, et al. Novel Co-rich high entropy alloys with superior tensile properties. Materials Research Letters. 2019;7(2):82-88
  34. 34. Coury FG, Santana D, Guo Y, Copley J, Otani L, Fonseca S, et al. Design and in-situ characterization of a strong and ductile co-rich multicomponent alloy with transformation induced plasticity. Scripta Materialia. 2019;173:70-74
  35. 35. Hsu C-Y, Wang W-R, Tang W-Y, Chen S-K, Yeh J-W. Microstructure and mechanical properties of new AlCoxCrFeMo0.5Ni high-entropy alloys. Advanced Engineering Materials. 2010;12(1-2):44-49
  36. 36. Hori T, Nagase T, Todai M, Matsugaki A, Nakano T. Development of non-equiatomic Ti-Nb-Ta-Zr-Mo high-entropy alloys for metallic biomaterials. Scripta Materialia. 2019;172:83-87
  37. 37. Li C, Zhao M, Li JC, Jiang Q. B2 structure of high-entropy alloys with addition of Al. Journal of Applied Physics. 2008;104(11):113504
  38. 38. Yang T, Xia S, Liu S, Wang C, Liu S, Zhang Y, et al. Effects of AL addition on microstructure and mechanical properties of AlxCoCrFeNi High-entropy alloy. Materials Science and Engineering: A. 2015;648:15-22
  39. 39. Yin Y, Tan Q , Sun Q , Ren W, Zhang J, Liu S, et al. Heterogeneous lamella design to tune the mechanical behaviour of a new cost-effective compositionally complicated alloy. Journal of Materials Science & Technology. 2022;96:113-125
  40. 40. Li Z, Tasan CC, Springer H, Gault B, Raabe D. Interstitial atoms enable joint twinning and transformation induced plasticity in strong and ductile high-entropy alloys. Scientific Reports. 2017;7:40704
  41. 41. Su J, Raabe D, Li Z. Hierarchical microstructure design to tune the mechanical behavior of an interstitial TRIP-TWIP high-entropy alloy. Acta Materialia. 2019;163:40-54
  42. 42. Nutor RK, Azeemullah M, Cao QP, Wang XD, Zhang DX, Jiang JZ. Microstructure and properties of a Co-free Fe50Mn27Ni10Cr13 high entropy alloy. Journal of Alloys and Compounds. 2021;851:156842
  43. 43. Yao MJ, Pradeep KG, Tasan CC, Raabe D. A novel, single phase, non-equiatomic FeMnNiCoCr high-entropy alloy with exceptional phase stability and tensile ductility. Scripta Materialia. 2014;72-73:5-8
  44. 44. Wang Z, Baker I. Interstitial strengthening of a f.c.c. FeNiMnAlCr high entropy alloy. Materials Letters. 2016;180:153-156
  45. 45. Shaysultanov DG, Salishchev GA, Ivanisenko YV, Zherebtsov SV, Tikhonovsky MA, Stepanov ND. Novel Fe36Mn21Cr18Ni15Al10 high entropy alloy with bcc/B2 dual-phase structure. Journal of Alloys and Compounds. 2017;705:756-763
  46. 46. Yin Y, Zhao Y, En Koey K, Tan Q , Zhang M-X, Huang H. In-situ synthesized age-hardenable high-entropy composites with superior wear resistance. Composites Part B: Engineering. 2022;235:109795
  47. 47. Yin Y, Zhang JQ , Tan QY, Zhuang W, Mo N, Bermingham M, et al. Novel cost-effective Fe-based high entropy alloys with balanced strength and ductility. Materials & Design. 2019;162:24-33
  48. 48. Yin Y, Chen Z, Mo N, Kent D, Candella AR, Koey KE, et al. High-temperature age-hardening of a novel cost-effective Fe45Ni25Cr25Mo5 high entropy alloy. Materials Science and Engineering: A. 2020;788:139580
  49. 49. Senkov ON, Miracle DB, Chaput KJ, Couzinie J-P. Development and exploration of refractory high entropy alloys—A review. Journal of Materials Research. 2018;33(19):3092-3128
  50. 50. Soni V, Gwalani B, Alam T, Dasari S, Zheng Y, Senkov ON, et al. Phase inversion in a two-phase, BCC+B2, refractory high entropy alloy. Acta Materialia. 2019;185:89-97
  51. 51. Otto F, Dlouhý A, Pradeep KG, Kuběnová M, Raabe D, Eggeler G, et al. Decomposition of the single-phase high-entropy alloy CrMnFeCoNi after prolonged anneals at intermediate temperatures. Acta Materialia. 2016;112:40-52
  52. 52. Chen SY, Tong Y, Tseng KK, Yeh JW, Poplawsky JD, Wen JG, et al. Phase transformations of HfNbTaTiZr high-entropy alloy at intermediate temperatures. Scripta Materialia. 2019;158:50-56
  53. 53. Stepanov ND, Yurchenko NY, Zherebtsov SV, Tikhonovsky MA, Salishchev GA. Aging behavior of the HfNbTaTiZr high entropy alloy. Materials Letters. 2018;211:87-90
  54. 54. Li ZM, Tasan CC, Pradeep KG, Raabe D. A TRIP-assisted dual-phase high-entropy alloy: Grain size and phase fraction effects on deformation behavior. Acta Materialia. 2017;131:323-335
  55. 55. Zhu Y, Ameyama K, Anderson PM, Beyerlein IJ, Gao H, Kim HS, et al. Heterostructured materials: Superior properties from hetero-zone interaction. Materials Research Letters. 2021;9(1):1-31
  56. 56. Hasan MN, Liu YF, An XH, Gu J, Song M, Cao Y, et al. Simultaneously enhancing strength and ductility of a high-entropy alloy via gradient hierarchical microstructures. International Journal of Plasticity. 2019;123:178-195
  57. 57. He BB, Hu B, Yen HW, Cheng GJ, Wang ZK, Luo HW, et al. High dislocation density-induced large ductility in deformed and partitioned steels. Science. 2017;357(6355):1029-1032
  58. 58. Huang H, Wu Y, He J, Wang H, Liu X, An K, et al. Phase-transformation ductilization of brittle high-entropy alloys via metastability engineering. Advanced Materials. 2017;29(30):1701678
  59. 59. Xu SS, Li JP, Cui Y, Zhang Y, Sun LX, Li J, et al. Mechanical properties and deformation mechanisms of a novel austenite-martensite dual phase steel. International Journal of Plasticity. 2020;128:102677
  60. 60. Liu L, Yu Q , Wang Z, Ell J, Huang MX, Ritchie RO. Making ultrastrong steel tough by grain-boundary delamination. Science. 2020;368(6497):1347-1352
  61. 61. Li Z, Raabe D. Strong and ductile non-equiatomic high-entropy alloys: Design, processing, microstructure, and mechanical properties. JOM. 2017;69(11):2099-2106

Written By

Yu Yin, Andrej Atrens, Han Huang and Ming-Xing Zhang

Reviewed: 27 April 2022 Published: 08 June 2022

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

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