Python/耐烧蚀合金智能体/wiki/raw/sources/.cache/高低温条件下高熵合金性能与机理的最新进展.pdf.txt

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Academic Editor: Hideyuki
Murakami
Received: 9 December 2024
Revised: 11 January 2025
Accepted: 13 January 2025
Published: 15 January 2025
Citation: Xi, R.; Li, Y. Recent
Advances in the Performance and
Mechanisms of High-Entropy Alloys
Under Low- and High-Temperature
Conditions. Coatings 2025, 15, 92.
https://doi.org/10.3390/
coatings15010092
Copyright: © 2025 by the authors.
Licensee MDPI, Basel, Switzerland.
This article is an open access article
distributed under the terms and
conditions of the Creative Commons
Attribution (CC BY) license
(https://creativecommons.org/
licenses/by/4.0/).
Review
Recent Advances in the Performance and
Mechanisms of High-Entropy Alloys Under Low- and
High-Temperature Conditions
Rui Xi 1 and Yanzhou Li 2,*
1 School of Mechanical Engineering, North China University of Water Resources and Electric Power,
Zhengzhou 450045, China; xirui@ncwu.edu.cn
2 School of Mechanical and Vehicle Engineering, West Anhui University, Lu’an 237010, China
* Correspondence: liyanzhou9336@163.com; Tel.: +86-150-4421-1299
Abstract: High-entropy alloys, since their development, have demonstrated great potential
for applications in extreme temperatures. This article reviews recent progress in their
mechanical performance, microstructural evolution, and deformation mechanisms at low
and high temperatures. Under low-temperature conditions, the focus is on alloys with
face-centered cubic, body-centered cubic, and multi-phase structures. Special attention
is given to their strength, toughness, strain-hardening capacity, and plastic-toughening
mechanisms in cold environments. The key roles of lattice distortion, nanoscale twin formation, and deformation-induced martensitic transformation in enhancing low-temperature
performance are highlighted. Dynamic mechanical behavior, microstructural evolution,
and deformation characteristics at various strain rates under cold conditions are also summarized. Research progress on transition metal-based and refractory high-entropy alloys is
reviewed for high-temperature environments, emphasizing their thermal stability, oxidation resistance, and frictional properties. The discussion reveals the importance of precipitation strengthening and multi-phase microstructure design in improving high-temperature
strength and elasticity. Advanced fabrication methods, including additive manufacturing and high-pressure torsion, are examined to optimize microstructures and improve
service performance. Finally, this review suggests that future research should focus on
understanding low-temperature toughening mechanisms and enhancing high-temperature
creep resistance. Further work on cost-effective alloy design, dynamic mechanical behavior
exploration, and innovative fabrication methods will be essential. These efforts will help
meet engineering demands in extreme environments.
Keywords: high-entropy alloys; low-temperature performance; high-temperature performance; strengthening mechanisms
1. Introduction
As science and technology advance swiftly, the application settings in critical areas
like aerospace, energy, transportation, and defense are growing more intricate. Taking
the aerospace industry as an example, with the continuous increase in engine operating
temperatures, they have approached or even exceeded the melting points of traditional hightemperature alloys, creating an urgent need for structural materials with higher thermal
resistance [1–8]. Applications like deep space, deep sea exploration, and superconductivity
demand materials capable of maintaining strength and toughness under low and ultra-low
temperature conditions to endure extreme environments [9–17]. Simultaneously, in new
Coatings 2025, 15, 92 https://doi.org/10.3390/coatings15010092


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Coatings 2025, 15, 92 2 of 32
application areas such as space debris mitigation and fusion reactors, service conditions
involving high loads and variable temperatures require enhanced mechanical properties
of materials over extensive temperature spans [18–25]. These application requirements
necessitate materials that exhibit excellent physical properties under extreme temperature
conditions and maintain good mechanical performance and long-term stability.
Beyond bulk materials, coatings have also been the focus of numerous studies for their
role in enhancing material performance under such extreme conditions. Advanced coatings,
including ceramic and metallic coatings, show potential to improve oxidation resistance,
thermal insulation, and wear resistance of substrate materials at high temperatures [26–28].
For instance, thermal barrier coatings are widely used in aero-engines to protect superalloys
from thermal degradation and extend their service life [29–31]. Similarly, low-temperature
applications require coatings that can mitigate embrittlement, reduce friction, and enhance
surface durability, ensuring long-term stability in cryogenic environments [32,33].
However, the impact of temperature changes on materials is highly complex. At
low temperatures, the plasticity and toughness of alloys typically decrease significantly,
dislocation movement becomes more difficult, and alloys are more prone to brittle failure [34–45]. At high temperatures, alloys may experience issues such as grain coarsening,
phase transitions, creep, and oxidation, leading to decreased material strength, reduced
flexibility, and an increased risk of fatigue and fracture [46–54].
Within this framework, the idea of multi-principal-element high-entropy alloys
(HEAs), proposed and developed over the past two decades, has provided new ideas
for material design in extreme application environments [55–67]. Researchers have found
that a multi-element mixture strategy can create complex and diverse microstructures, synergistically activating multiple strengthening mechanisms to enhance the material’s overall
performance [68–80]. HEAs have shown extensive application possibilities in harsh environments, including extremely low and high temperatures, broad temperature fluctuations,
high-velocity impacts, radiation-induced damage, and cyclic loading [81–85].
This paper will review the latest research advancements and theoretical progress on
HEAs fabricated using various methods under low and ultra-high temperature conditions.
The common microstructure types of HEAs and their deformation mechanisms under
extreme loading environments will be discussed, along with important unresolved research
issues, and the paper will conclude with an outlook on future research directions.
2. Introduction and Structural Characteristics of HEAs
2.1. Introduction to HEAs
In 2004, Cantor et al. and Yeh et al. introduced HEAs as a novel category of metallic
materials that question conventional alloy design principles [86,87]. Conventional alloys
usually contain one or two main components, supplemented by other elements to enhance their characteristics. Conversely, HEAs incorporate four or more constituents in
approximately equal atomic proportions, resulting in a straightforward lattice structure
with extensive chemical randomness. Typically, these alloys develop as a single-phase
solid solution. Including numerous primary elements in HEAs inhibits the development
of intermetallic compounds (IC). By suppressing IC formation, HEAs address the key
limitations found in conventional alloy systems, offering enhanced design flexibility. The
discovery of HEAs has disrupted the conventional understanding of phase behaviors in
metals and significantly expanded the scope for alloy composition design.
Figure 1 illustrates the development history of HEAs [88–90]. Initial studies primarily
examined the architecture and fundamental characteristics of HEAs. With the expansion of
industrial requirements, additional functional attributes—including resistance to radiation
and corrosion, hydrogen storage capabilities, and catalytic efficiency—started to garner


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interest [91–102]. Further research discovered that HEAs exhibit excellent mechanical
properties under low and high-temperature conditions. This has led to significant advances in traditional applications and broadened their potential in extreme temperature
environments such as low and high temperatures, driving their widespread prospects
in engineering.
Coatings 2025, 15, 92 3 of 34 
radiation and corrosion, hydrogen storage capabilities, and catalytic efficiency—started to 
garner interest [91–102]. Further research discovered that HEAs exhibit excellent mechanical properties under low and high-temperature conditions. This has led to significant advances in traditional applications and broadened their potential in extreme temperature 
environments such as low and high temperatures, driving their widespread prospects in 
engineering. 
Figure 1. Overview of the historical evolution and milestones in HEA research.
2.2. Chemical Disorder Structure of HEAs 
HEAs comprise multiple primary elements, and their phase structure governs the 
material properties. Entropy is a major factor affecting the structure and a measure of the 
system’s stability. The thermodynamic mixing entropy refers to the configurational entropy of the system. Elevated mixing entropy decreases the phase count in HEAs relative 
to that anticipated by Gibbs’ phase rule, thereby improving the interaction among the primary constituents [90]. As stated by the Gibbs free energy equation (ΔGmixed = ΔHmixed 
− TΔSmixed, here, ΔGmixed represents the alloy system’s free energy, ΔHmixed denotes 
the mixing enthalpy, ΔSmixed signifies the mixing entropy, and T stands for the temperature), higher mixing entropy helps lower the system’s free energy, maintaining the solid 
solution framework and inhibiting the creation of IC. As a result, HEAs preserve a consistent solid solution architecture at elevated temperatures and display chemical randomness traits, as shown in Figure 2 [90].
Figure 2. Schematic diagram of lattice distortion in HEAs [90]. 
The development of solid solution phases is additionally affected by elements like 
atomic size disparity (δ), electronegativity difference, and mixing enthalpy among the 
Figure 1. Overview of the historical evolution and milestones in HEA research.
2.2. Chemical Disorder Structure of HEAs
HEAs comprise multiple primary elements, and their phase structure governs the
material properties. Entropy is a major factor affecting the structure and a measure of the
system’s stability. The thermodynamic mixing entropy refers to the configurational entropy
of the system. Elevated mixing entropy decreases the phase count in HEAs relative to that
anticipated by Gibbs’ phase rule, thereby improving the interaction among the primary
constituents [90]. As stated by the Gibbs free energy equation (∆Gmixed = ∆Hmixed −
T∆Smixed, here, ∆Gmixed represents the alloy system’s free energy, ∆Hmixed denotes the
mixing enthalpy, ∆Smixed signifies the mixing entropy, and T stands for the temperature),
higher mixing entropy helps lower the system’s free energy, maintaining the solid solution
framework and inhibiting the creation of IC. As a result, HEAs preserve a consistent solid
solution architecture at elevated temperatures and display chemical randomness traits, as
shown in Figure 2 [90].
Coatings 2025, 15, 92 3 of 34 
radiation and corrosion, hydrogen storage capabilities, and catalytic efficiency—started to 
garner interest [91–102]. Further research discovered that HEAs exhibit excellent mechanical properties under low and high-temperature conditions. This has led to significant advances in traditional applications and broadened their potential in extreme temperature 
environments such as low and high temperatures, driving their widespread prospects in 
engineering. 
Figure 1. Overview of the historical evolution and milestones in HEA research.
2.2. Chemical Disorder Structure of HEAs 
HEAs comprise multiple primary elements, and their phase structure governs the 
material properties. Entropy is a major factor affecting the structure and a measure of the 
system’s stability. The thermodynamic mixing entropy refers to the configurational entropy of the system. Elevated mixing entropy decreases the phase count in HEAs relative 
to that anticipated by Gibbs’ phase rule, thereby improving the interaction among the primary constituents [90]. As stated by the Gibbs free energy equation (ΔGmixed = ΔHmixed 
− TΔSmixed, here, ΔGmixed represents the alloy system’s free energy, ΔHmixed denotes 
the mixing enthalpy, ΔSmixed signifies the mixing entropy, and T stands for the temperature), higher mixing entropy helps lower the system’s free energy, maintaining the solid 
solution framework and inhibiting the creation of IC. As a result, HEAs preserve a consistent solid solution architecture at elevated temperatures and display chemical randomness traits, as shown in Figure 2 [90].
Figure 2. Schematic diagram of lattice distortion in HEAs [90]. 
The development of solid solution phases is additionally affected by elements like 
atomic size disparity (δ), electronegativity difference, and mixing enthalpy among the 
Figure 2. Schematic diagram of lattice distortion in HEAs [90].
The development of solid solution phases is additionally affected by elements like
atomic size disparity (δ), electronegativity difference, and mixing enthalpy among the
major elements. When δ ≤ 6.6%, reduced lattice distortions promote the establishment of
solid solutions [103–107]. Valence electron concentration (VEC) serves as a predictor for de-


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termining the crystal structure of HEAs. For VEC values exceeding eight, an FCC structure
is probable [108–110]. The majority of HEAs generally develop crystalline architectures,
where the FCC structure offers excellent plasticity yet lower strength. Conversely, the BCC
structure provides enhanced strength but increased brittleness [111–113]. Due to the slower
atomic diffusion kinetics in HEAs, they may also form amorphous or nanocrystalline
phases [114–116].
Furthermore, Luan et al. [117] highlighted that adding more elements directly amplifies the number of potential IC within the alloy matrix. The impact of mixing enthalpy on
low-temperature environments becomes more significant, and the formation of IC reduces
the system’s free energy. As a result, most HEAs tend to form multi-phase structures. The
combined influence of various factors makes the phase formation behavior of HEAs highly
complex, offering extensive opportunities for controlling their microstructural characteristics. This complexity also provides a broad space for designing and optimizing HEAs with
tailored properties.
Recent research has confirmed that alloys with non-equal atomic ratios can maintain a
stable solid solution structure and that variations in specific elements can notably impact
their properties [118,119]. This “cocktail effect” enables HEAs to retain phase stability
across a broader range of compositions, which expands their potential for applications in
challenging and extreme environments [120,121].
Additionally, studies have revealed that HEAs often display elemental segregation, as
seen in ordered oxygen complexes in TiZrNbHf alloys and other ordered phases in similar
materials [122–124]. These nanoscale-ordered structures challenge the traditional concept
of chemical disorder in HEAs, offering a new understanding of solid material deformation
mechanisms. This creates new opportunities for developing stronger and tougher alloys,
with important scientific implications.
3. Low-Temperature Properties of HEAs
In their 2011 study on the low-temperature performance of HEAs, Qiao et al. [125]
were the pioneers in examining the behavior of the AlCoCrFeNi HEA, which possesses a
single-phase BCC structure. They found that this alloy exhibited exceptional compressive
performance at low temperatures, with a yield strength reaching 1.88 GPa, and displayed
serrated flow behavior. In 2013, a research team led by George E.P. [126] from Oak Ridge
National Laboratory investigated the low-temperature properties of CoCrFeNiMn and
CoCrFeNi alloys (later known as Cantor alloys), revealing that both strength and ductility
markedly enhanced with decreasing temperatures. Specifically, the tensile strength (TS) of
the CoCrFeNi HEA surpassed 1 GPa at 77 K, with an elongation greater than 60%. As research on HEAs has advanced, the scope has expanded beyond the low-temperature tensile
and compressive performance of equimolar HEAs. For example, studies on non-equimolar
HEAs at low temperatures, such as AlCoCrFeNi2.1 [127,128], CrMnFeCoNi2Cu [129], and
Al0.5CoCrCuFeNi [130,131], have been increasingly performed. Furthermore, research has
also covered various aspects of HEAs under low-temperature conditions, including elastic
properties, corrosion resistance, and electromagnetic behaviors. This chapter will provide a
classified review of the research developments based on the lattice structures of HEAs.
3.1. FCC-Phase HEAs
3.1.1. CoCrFeMnNi HEA
Recent research into the plastic deformation of HEAs has shown that, in comparison
to traditional materials, certain FCC-phase HEAs exhibit a unique characteristic of “increasing strength and toughness at lower temperatures”. CoCrFeMnNi, a widely studied
single-phase FCC HEA, was one of the earliest alloys investigated for its low-temperature


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mechanical properties. As mentioned earlier [126], it was found that the tensile strength(TS)
and flexibility of CoCrFeMnNi improved significantly at 77 K. Building upon this, in
study [132], the impact of grain size on the low-temperature mechanical characteristics
and deformation mechanisms of HEAs was further investigated. FCC HEAs undergo the
formation of nanotwins during the later stages of low-temperature deformation, where
the abundance of low-energy interfaces generated by twinning refines the grains. These
low-energy interfaces effectively hinder dislocation motion, increasing the material’s workhardening rate and delaying necking behavior. The simultaneous effects of dislocation slip
and twinning mechanisms significantly boosted the strength and ductility of Cantor alloys
at 77 K.
Building upon earlier findings, later research has uncovered how nanotwinning contributes to the improvement in low-temperature mechanical properties in HEAs. The
Ritchie research team from the University of California, Berkeley, reported in 2014 that [133]
CrMnFeCoNi alloys exhibited excellent fracture toughness at liquid nitrogen temperatures,
confirming that the formation of numerous nanotwins was the microstructural origin of
their superior low-temperature mechanical performance. The appearance of nanotwins
refined the grains, and the abundance of low-energy interfaces further improved the
mechanical properties at low temperatures. The research also found that the lower stacking fault energy in HEAs leads to twinning deformation characteristics, where dislocation motion is hindered at low temperatures, and the synergistic effect of twinning and
dislocations enables HEAs to exhibit the “stronger and tougher at lower temperatures”
behavior [133,134]. In 2014, Gludovatz et al. [133] evaluated the fracture toughness of
single-phase FCC HEAs at both ambient and low temperatures. Their findings revealed
that the alloy exhibited exceptional toughness at low temperatures, achieving fracture
toughness values above 200 MPa·m1/2 during crack initiation and exceeding 300 MPa·m1/2
during crack propagation at 77 K.
A study on the CoCrFeMnNi HEA examined its mechanical properties under multidirectional loading across temperatures ranging from 113 K to 1273 K, emphasizing
its performance under directional stress. Kim et al. [135] found that although the alloy
had a strong <111> or <001> crystallographic texture, the elongation to fracture (Ef) was
affected by the loading direction. Concurrently, both yield strength (TS) and toughness
exhibited minimal dependence on the direction of loading. Figure 3 depicts the correlations
among tensile YS, TS, and elongation to fracture (Ef) as they relate to temperature and
loading direction. YS and TS were highest at 113 K, indicating excellent strength at low
temperatures, with the loading direction having minimal impact on these properties. At
773 K, the as-cast alloy showed the highest Ef, offering better flexibility than the wrought
alloy at 673 K. However, above 973 K, Ef decreased sharply, signaling reduced ductility at
elevated temperatures.
V-notch Charpy impact test results indicated that, between −150 ◦C and 150 ◦C, the ascast HEA absorbed about half the impact energy compared to the wrought alloy. Despite the
reduced absorbed energy, the as-cast alloy did not demonstrate considerable temperature
dependence or low-temperature embrittlement. Along different loading directions, the
WL direction showed the highest absorbed impact energy, while the WS and H directions
were relatively lower, indicating a degree of anisotropy. However, overall, the anisotropy
observed in both tensile and impact tests was insignificant. Moreover, under reduced
temperatures, the as-cast alloy’s impact toughness exceeded that of numerous conventional
alloys, displaying strong embrittlement resistance.


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Figure 3. Variations in (a) mechanical strength and (b) fracture strain across temperatures ranging 
from 113 K to 1273 K and different loading directions [135].
V-notch Charpy impact test results indicated that, between −150 °C and 150 °C, the 
as-cast HEA absorbed about half the impact energy compared to the wrought alloy. Despite the reduced absorbed energy, the as-cast alloy did not demonstrate considerable 
temperature dependence or low-temperature embrittlement. Along different loading directions, the WL direction showed the highest absorbed impact energy, while the WS and 
H directions were relatively lower, indicating a degree of anisotropy. However, overall, 
the anisotropy observed in both tensile and impact tests was insignificant. Moreover, under reduced temperatures, the as-cast alloy’s impact toughness exceeded that of numerous conventional alloys, displaying strong embrittlement resistance.
Analysis of strain hardening behavior indicated that the as-cast HEA underwent 
three discernible stages across 113 K, ambient temperature 298 K, 773 K, and 973 K. In 
Stage I, the hardening rate declined; in Stage II, it rose owing to deformation twin formation; and in Stage III, it dropped again following twin saturation. Notably, at 773 K, 
twin formation was highly active, contributing to the improved flexibility at this temperature. At 700 °C, however, only a limited quantity of twin boundaries was detected, and 
the number of recrystallized grains increased, as the deformation mechanism transitioned 
from planar to wavy slip, a marked reduction in flexibility ensued. This indicates that twin 
activity was greatly suppressed at high temperatures, with recrystallization and dislocation slip becoming the dominant deformation mechanisms.
Further investigations have highlighted twinning’s contribution to improving HEA 
properties at low temperatures. For example, Kireeva et al. [136] examined the deformation behavior of CoCrFeMnNi HEA single crystals. Their findings indicate that at 296 
K, the [137]-oriented crystal primarily deforms via dislocation slip up to 30% strain, after 
which twinning appears. In contrast, the [123]- and [011]-oriented crystals deform predominantly through dislocation slip until failure. At 77 K, twinning becomes the primary 
deformation mechanism. In the [137]-oriented crystal, deformation transitions to twinning after a 5% slip, while in the [111]-oriented crystal, twinning initiates after a 20% slip. 
The strain hardening rate decreased in tandem, as shown in Figure 4. At 77 K and 20% 
strain, [111]-oriented crystals displayed dense dislocations alongside intersecting twins 
with around 15–20 nm thicknesses. By 35% strain, twinning predominantly takes place in 
a single system, and after treatment, twins are not visible on the surface, indicating they 
are very thin and uniformly distributed throughout the crystal. Compared to RT(RT), the 
alloy exhibits significantly improved flexibility and plasticity at 77 K, highlighting the crucial role of twinning in enhancing the toughness and strength of the HEA at low temperatures.
Figure 3. Variations in (a) mechanical strength and (b) fracture strain across temperatures ranging
from 113 K to 1273 K and different loading directions [135].
Analysis of strain hardening behavior indicated that the as-cast HEA underwent three
discernible stages across 113 K, ambient temperature 298 K, 773 K, and 973 K. In Stage I,
the hardening rate declined; in Stage II, it rose owing to deformation twin formation; and
in Stage III, it dropped again following twin saturation. Notably, at 773 K, twin formation
was highly active, contributing to the improved flexibility at this temperature. At 700 ◦C,
however, only a limited quantity of twin boundaries was detected, and the number of
recrystallized grains increased, as the deformation mechanism transitioned from planar
to wavy slip, a marked reduction in flexibility ensued. This indicates that twin activity
was greatly suppressed at high temperatures, with recrystallization and dislocation slip
becoming the dominant deformation mechanisms.
Further investigations have highlighted twinning’s contribution to improving HEA
properties at low temperatures. For example, Kireeva et al. [136] examined the deformation
behavior of CoCrFeMnNi HEA single crystals. Their findings indicate that at 296 K,
the [137]-oriented crystal primarily deforms via dislocation slip up to 30% strain, after which
twinning appears. In contrast, the [123]- and [011]-oriented crystals deform predominantly
through dislocation slip until failure. At 77 K, twinning becomes the primary deformation
mechanism. In the [137]-oriented crystal, deformation transitions to twinning after a 5%
slip, while in the [111]-oriented crystal, twinning initiates after a 20% slip. The strain
hardening rate decreased in tandem, as shown in Figure 4. At 77 K and 20% strain, [111]-
oriented crystals displayed dense dislocations alongside intersecting twins with around
15–20 nm thicknesses. By 35% strain, twinning predominantly takes place in a single
system, and after treatment, twins are not visible on the surface, indicating they are very
thin and uniformly distributed throughout the crystal. Compared to RT(RT), the alloy
exhibits significantly improved flexibility and plasticity at 77 K, highlighting the crucial
role of twinning in enhancing the toughness and strength of the HEA at low temperatures.
Coatings 2025, 15, 92 7 of 34 
Figure 4. Twin interactions in [111]-oriented CoCrFeMnNi HEA crystals subjected to 20% tensile 
deformation at 77 K are shown as: (a) a bright-field image, (b) a dark-field image, and (c) a diffraction pattern [136].
In addition to these findings, Rusakova et al. [138] compared the mechanical properties of initial coarse-grained samples and nanocrystalline variants before and after highpressure torsion (HPT) at low temperatures in CoCrFeNiMn HEAs. The as-cast CoCrFeNiMn HEA was annealed and rolled to achieve an average grain size of four microns. Fol
Figure 4. Twin interactions in [111]-oriented CoCrFeMnNi HEA crystals subjected to 20% tensile
deformation at 77 K are shown as: (a) a bright-field image, (b) a dark-field image, and (c) a diffraction
pattern [136].


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In addition to these findings, Rusakova et al. [138] compared the mechanical properties
of initial coarse-grained samples and nanocrystalline variants before and after high-pressure
torsion (HPT) at low temperatures in CoCrFeNiMn HEAs. The as-cast CoCrFeNiMn HEA
was annealed and rolled to achieve an average grain size of four microns. Following HPT
treatment at RT, the grain size diminished to 40–80 nm, yielding a uniform microstructure,
and only a few grains exceeded 200 nm. Low-temperature HPT treatment resulted in a
more heterogeneous structure, with nanocrystalline grains and larger grains up to 500 nm
in size. Within the examined temperature interval (77 K to 290 K), microhardness increased
as the temperature decreased, indicating that thermally activated processes significantly
influenced the plastic deformation mechanism of the material at low temperatures. For the
cryo-HPT-treated sample, the variation in microhardness along the disk diameter showed
distinct differences. The microhardness rose approximately 35%–40% as it extended from
the sample’s center to a radius of 1.5–2 mm, increasing from 3 GPa up to 4.25 GPa. In the
RT-HPT-treated specimen, microhardness elevated from 3.5 GPa to 4.75 GPa, following
a comparable trend, although the cryo-HPT-induced microhardness increase was less
pronounced. Microhardness surged from an initial 1.4 GPa to 4.25 GPa with cryo-HPT
treatment and to 4.75 GPa following RT-HPT processing. In all studied temperature ranges
from 77 K to 290 K, microhardness increased as the temperature decreased, indicating that
thermally activated processes substantially governed the material’s plastic deformation
mechanisms under low-temperature conditions.
3.1.2. Equimolar CoCrFe-Based HEAs
While CoCrFeMnNi remains the most extensively studied FCC phase HEA for lowtemperature applications, research has expanded to include other equimolar FCC phase
HEAs, exploring their low-temperature behaviors. Tan et al. [139] investigated lattice
distortions, the evolution of magnetic properties, and associated behaviors in CrCoNi, CrFeCoNi, and CrMnFeCoNi HEA across various temperatures. As decreasing temperatures,
the lattice parameters of each alloy significantly contracted. The lattice contraction for
CrCoNi MEA was 1.425%, for CrFeCoNi HEA, 2.278%, and for CrMnFeCoNi HEA, it was
the greatest, at 3.476%. As the temperature dropped from 293 K to 123 K, the local lattice
distortions in each of the three alloys generally diminished. CrFeCoNi HEA showed the
most dramatic changes, especially below 173 K, where changes in the local structure were
observed. In terms of magnetism, the CrFeCoNi HEA shifted from a paramagnetic to a
ferromagnetic state near 34 K and retained its ferromagnetic properties at 25 K and 5 K,
exhibiting soft magnetic behavior. Throughout the entire low-temperature spectrum (77 K,
25 K, and 5 K), both CrCoNi MEA and CrMnFeCoNi HEA stayed paramagnetic, exhibiting
only minor increases in magnetization. At low temperatures, the stress–strain profiles for
all three alloys displayed comparable serrated patterns, suggesting analogous deformation
mechanisms. As the temperature declined from 293 K to 123 K and down to 5 K, the lattice
distortions in all three alloys underwent substantial alterations. Additionally, the CrFeCoNi
HEA experienced a notable magnetic transition below 173 K, shifting from paramagnetic
to ferromagnetic. In contrast, CrCoNi MEA and CrMnFeCoNi HEA did not undergo such
a transition.
The impact of alloying elements such as Cu on the magnetic properties at low temperatures has also been explored. Chaudhary et al. [140] examined how adding copper affects
the low-temperature magnetic properties of FeCoNiCrCu HEA. At RT, the FeCoNiCr alloy
maintains a single-phase FCC. The alloy exhibits paramagnetic behavior with a Curie temperature of 85 K. As the copper content rises from 0 to 0.5, Ms increases from 30.7 emu/g to
32.7 emu/g. However, the XRD patterns reveal the presence of two FCC phases, one of
which is enriched with copper. The copper-rich regions primarily appear as small clusters,


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ranging from 3 to 5 nm in size, and these clusters are distributed within the FCC matrix,
forming a copper-enriched nanostructure.
At low temperatures, the saturation magnetization of FeCoNiCrCux alloys shows a
noticeable enhancement. When the copper content increases from 0 to 0.5, Ms increases
from 30.7 emu/g to 32.7 emu/g. This increase in saturation magnetization indicates that
the addition of copper significantly enhances the low-temperature magnetism of the alloy.
The improved magnetic properties suggest that copper affects the alloy’s microstructure
and strengthens the magnetic exchange interactions within the alloy, thereby enhancing
its magnetic performance. Figure 5 shows Cu’s distribution and structural characteristics
in the FeCoNiCrCu0.5 alloy. Figure 5a displays the individual ion distributions of Fe,
Co, and Cu, illustrating that copper forms nanoscale clusters within the FCC matrix,
each approximately 3 to 5 nm in size. Figure 5b analyzes the composition variation
between the fcc matrix and Cu clusters using a proximity histogram, revealing that the Cu
clusters comprise around 92% Cu and 5% Ni, along with approximately 1% Fe, Co, and
Cr. Figure 5c displays the morphology of Cu clusters of different sizes obtained through
atom probe tomography reconstruction, and further 1-D compositional analysis indicates
that the cluster size influences its composition, with larger Cu clusters containing a higher
Cu content.
pyg gpp
peratures has also been explored. Chaudhary et al. [140] examined how adding copper 
affects the low-temperature magnetic properties of FeCoNiCrCu HEA. At RT, the FeCoNiCr alloy maintains a single-phase FCC. The alloy exhibits paramagnetic behavior with 
a Curie temperature of 85 K. As the copper content rises from 0 to 0.5, Ms increases from 
30.7 emu/g to 32.7 emu/g. However, the XRD patterns reveal the presence of two FCC 
phases, one of which is enriched with copper. The copper-rich regions primarily appear 
as small clusters, ranging from 3 to 5 nm in size, and these clusters are distributed within 
the FCC matrix, forming a copper-enriched nanostructure.
At low temperatures, the saturation magnetization of FeCoNiCrCux alloys shows a 
noticeable enhancement. When the copper content increases from 0 to 0.5, Ms increases 
from 30.7 emu/g to 32.7 emu/g. This increase in saturation magnetization indicates that 
the addition of copper significantly enhances the low-temperature magnetism of the alloy. 
The improved magnetic properties suggest that copper affects the alloy’s microstructure 
and strengthens the magnetic exchange interactions within the alloy, thereby enhancing 
its magnetic performance. Figure 5 shows Cu’s distribution and structural characteristics 
in the FeCoNiCrCu0.5 alloy. Figure 5a displays the individual ion distributions of Fe, Co, 
and Cu, illustrating that copper forms nanoscale clusters within the FCC matrix, each approximately 3 to 5 nm in size. Figure 5b analyzes the composition variation between the 
fcc matrix and Cu clusters using a proximity histogram, revealing that the Cu clusters 
comprise around 92% Cu and 5% Ni, along with approximately 1% Fe, Co, and Cr. Figure 
5c displays the morphology of Cu clusters of different sizes obtained through atom probe 
tomography reconstruction, and further 1-D compositional analysis indicates that the cluster size influences its composition, with larger Cu clusters containing a higher Cu content. 
Figure 5. (a) Elemental distributions of Cr, Fe, Ni, Co, and Cu in the FeNiCoCrCu0.5 HEA; (b) Composition variations between the FCC matrix and Cu-rich areas; (c) atom probe tomography reconstruction illustrating Cu distribution, enlarged cluster images, and a one-dimensional composition 
analysis [140].
3.1.3. Non-Equimolar HEAs 
Building on the understanding gained from equimolar HEAs, researchers have explored non-equimolar HEAs by adjusting alloying element ratios to develop compositions 
with improved low-temperature performance. Tabachnikova et al. [129] explored the mechanical properties and fracture behaviors of CrMnFeCoNi2Cu alloys under low-temperature conditions. The CrMnFeCoNi2Cu alloy has an FCC crystal structure, and in state II, 
Figure 5. (a) Elemental distributions of Cr, Fe, Ni, Co, and Cu in the FeNiCoCrCu0.5 HEA; (b) Composition variations between the FCC matrix and Cu-rich areas; (c) atom probe tomography reconstruction illustrating Cu distribution, enlarged cluster images, and a one-dimensional composition
analysis [140].
3.1.3. Non-Equimolar HEAs
Building on the understanding gained from equimolar HEAs, researchers have explored non-equimolar HEAs by adjusting alloying element ratios to develop compositions with improved low-temperature performance. Tabachnikova et al. [129] explored
the mechanical properties and fracture behaviors of CrMnFeCoNi2Cu alloys under lowtemperature conditions. The CrMnFeCoNi2Cu alloy has an FCC crystal structure, and in
state II, after 60% rolling deformation, the alloy exhibits a bimodal FCC structure. Compression and tensile tests were conducted in the 300 K to 4.2 K temperature range. The
results showed that the alloy exhibits high plasticity under compression, especially at low
temperatures below 15 K, where it displays jump-like plastic flow. Under compression, the
jump depth ∆τ is approximately 44 MPa initially. As the strain progresses from 1% to 2%,
the jump depth escalates to around 115 MPa. During tensile deformation, the alloy forms a
necking phenomenon, and fracture occurs along a plane at approximately 45◦to the tensile
axis. Even at low temperatures of 4.2 K, the alloy maintains ductile fracture characteristics.


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Figure 6 depicts the tensile fracture surface morphology of the CrMnFeCoNi2Cu alloy
across various temperatures, from 300 K to 4.2 K. At 300 K, the alloy displays standard
ductile fracture features, including numerous flat areas and elongated protrusions on the
fracture surface. As the temperature decreases to 200 K and 100 K, the fracture surface
maintains ductile features but with more pronounced plastic deformation traces, and the
size and shape of the protrusions change. Upon further lowering the temperature to
50 K, while ductile features are still present, the brittle component increases, and more
microcracks appear on the fracture surface. High magnification images show that at 100 K
and 4.2 K, local inhomogeneities and microcracks become more noticeable, indicating
that slip localization and microcrack formation are more active at lower temperatures.
Remarkably, even at an ultralow temperature of 4.2 K, the alloy maintains ductile fracture
traits, showcasing evident plastic deformation and protrusions on the fracture surface. This
indicates that the alloy retains substantial plasticity and toughness under low-temperature
conditions. In general, decreasing temperatures lead to a gradual reduction in the alloy’s
ductility and an increase in its brittleness. Nevertheless, the alloy continues to exhibit
considerable plastic deformation capacity even at low temperatures.
peratures below 15 K, where it displays jumplike plastic flow. Under compression, the 
jump depth Δτ is approximately 44 MPa initially. As the strain progresses from 1% to 2%, 
the jump depth escalates to around 115 MPa. During tensile deformation, the alloy forms 
a necking phenomenon, and fracture occurs along a plane at approximately 45° to the 
tensile axis. Even at low temperatures of 4.2 K, the alloy maintains ductile fracture characteristics.
Figure 6 depicts the tensile fracture surface morphology of the CrMnFeCoNi2Cu alloy 
across various temperatures, from 300 K to 4.2 K. At 300 K, the alloy displays standard 
ductile fracture features, including numerous flat areas and elongated protrusions on the 
fracture surface. As the temperature decreases to 200 K and 100 K, the fracture surface 
maintains ductile features but with more pronounced plastic deformation traces, and the 
size and shape of the protrusions change. Upon further lowering the temperature to 50 K, 
while ductile features are still present, the brittle component increases, and more microcracks appear on the fracture surface. High magnification images show that at 100 K 
and 4.2 K, local inhomogeneities and microcracks become more noticeable, indicating that 
slip localization and microcrack formation are more active at lower temperatures. Remarkably, even at an ultralow temperature of 4.2 K, the alloy maintains ductile fracture 
traits, showcasing evident plastic deformation and protrusions on the fracture surface. 
This indicates that the alloy retains substantial plasticity and toughness under low-temperature conditions. In general, decreasing temperatures lead to a gradual reduction in 
the alloy’s ductility and an increase in its brittleness. Nevertheless, the alloy continues to 
exhibit considerable plastic deformation capacity even at low temperatures. 
Figure 6. Fracture surface morphology of CrMnFeCoNi2Cu alloy under tensile deformation at various temperatures. (a) 300 K. (b) Partial enlarged view of (a). (c) 77 K. (d) Partial enlarged view of (c). 
(e) 10 K. (f) Partial enlarged view of (e). (g) 4.2 K. (h) Partial enlarged view of (g) [129].
Investigating the combined effects of low temperatures and high strain rates is essential for comprehending the mechanical boundaries of HEAs. Jiang et al. [141] examined 
Cr26Mn20Fe20Co20Ni14 HEA under low-temperature and high-strain rate environments. As 
the temperature lowers and the strain rate rises, the alloy’s YS and UTS experience substantial increases, whereas its ductility diminishes. At RT with a strain rate of 103 s⁻1, the 
Cr26Mn20Fe20Co20Ni14 HEA showed a yield strength (YS) of 212 MPa, an ultimate TS of 565 
MPa, and an elongation of 0.73. Under LNT and a strain rate of 101 s⁻1, the YS and UTS 
Figure 6. Fracture surface morphology of CrMnFeCoNi2Cu alloy under tensile deformation at
various temperatures. (a) 300 K. (b) Partial enlarged view of (a). (c) 77 K. (d) Partial enlarged view of
(c). (e) 10 K. (f) Partial enlarged view of (e). (g) 4.2 K. (h) Partial enlarged view of (g) [129].
Investigating the combined effects of low temperatures and high strain rates is essential
for comprehending the mechanical boundaries of HEAs. Jiang et al. [141] examined
Cr26Mn20Fe20Co20Ni14 HEA under low-temperature and high-strain rate environments.
As the temperature lowers and the strain rate rises, the alloy’s YS and UTS experience
substantial increases, whereas its ductility diminishes. At RT with a strain rate of 103s
−1
,
the Cr26Mn20Fe20Co20Ni14 HEA showed a yield strength (YS) of 212 MPa, an ultimate
TS of 565 MPa, and an elongation of 0.73. Under LNT and a strain rate of 101s
−1
, the
YS and UTS surged by 210% and 137%, reaching 656 MPa and 1337 MPa. Concurrently,
the elongation decreased by 30%, down to 0.51. In LNT and high strain rate scenarios,
the alloy’s deformation mechanisms altered, with the development of slip bands (SF),
twinning, and HCP phase transformations playing pivotal roles in the deformation process.
In particular, at LNT, forming a nanotwin HCP phase composite structure helped optimize
stress and strain distribution, reducing damage nucleation. However, these deformation
mechanisms enhanced the strength and led to embrittlement, making interfaces (such as
grain and phase boundaries more susceptible to cracking.


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Semerenko et al. [142] investigated how martensitic phase transformation (DIMT)
affects the properties of the Co17.5Cr12.5Fe55Ni10Mo5 alloy during deformation at low temperatures. At RT, this alloy displays an FCC phase structure. Following annealing, the
alloy undergoes complete recrystallization with randomly oriented grains, preserving the
FCC crystal structure. At RT, this alloy displays an FCC. Following annealing, the alloy
undergoes complete recrystallization with randomly oriented grains, preserving the FCC
crystal structure. With an increasing proportion of the BCC phase, the alloy’s dynamic
Young’s modulus rises linearly, suggesting that phase transformation boosts the mechanical
properties of the alloy. At 4.2 K, the deformed alloy shows excellent mechanical performance
with a YS of 1043 MPa and a TS of 1748 MPa while maintaining high plasticity. The shift
from FCC to BCC and HCP phases is regarded as the primary mechanism enhancing the
low-temperature strength and plasticity of the Co17.5Cr12.5Fe55Ni10Mo5 alloy.
DIMT has been demonstrated to be pivotal in improving the mechanical properties
of HEAs under low-temperature conditions. Moon et al. [143] performed tensile tests on
the Co17.5Cr12.5Fe55Ni10Mo5 alloy at temperatures of 77 K, 4.2 K, 2.1 K, and 0.5 K. They
discovered that DIMT is the fundamental mechanism behind the alloy’s high strength and
excellent ductility. In low-temperature deformation, the FCC phase converts into HCP and
BCC phases. Upon reaching a true strain of 10%, the FCC phase initiates transformation
into HCP and BCC phases. As deformation progresses, the BCC proportionally increases
substantially, whereas the HCP remains below 10%. The Mo5 alloy showcases a high YS
of 1075 MPa and a TS of 1651 MPa. It undergoes ductile fracture, characterized by typical
dimpled fracture surfaces at both 77 K and 0.5 K, signifying excellent elasticity. As the
temperature decreases, the alloy exhibits DPF, and adiabatic heating during deformation
promotes dislocation cross-slip and DIMT. Nevertheless, as the temperature continues to
drop, these deformation mechanisms are progressively inhibited. At temperatures below
4.2 K, the alloy’s YS becomes highly temperature dependent. Dislocation motion seems to
be affected by inertial mechanisms, resulting in atypical dislocation behavior and further
complicating deformation processes. DIMT is vital for the ductility of the Mo5 alloy, as
it effectively mitigates strain localization, thereby sustaining excellent flexibility even at
extremely low temperatures.
Scientists have undertaken studies to tailor the low-temperature properties of HEAs
through the incorporation of non-metallic modifying elements. Yang et al. [137] investigated deformation mechanisms of Fe-based HEAs incorporating carbon and silicon across
various temperatures. By incorporating 1% carbon and silicon into the Fe53Mn29Co9Cr9
HEA, they modified the stacking fault energy (SFE) to 33.96 mJ/m2, achieving a singlephase FCC structure. The De-HEAs showcased high strength and excellent ductility at
RT. When tested at 223 K, the alloy’s YS rose to 907 MPa, and its elongation enhanced to
69.6%. The improvement in the alloy’s performance was mainly due to alterations in its
deformation mechanisms under low-temperature conditions. Lowering the temperature
led to a notable rise in the alloy’s strain hardening rate, causing the deformation mechanism
to shift from dislocation slip to twinning. As the strain increased, secondary nanotwins
began to form and activate. Secondary nanoscale twins developed within the alloy’s dislocation framework, promoting grain refinement and enhancing the material’s strength and
plasticity. Refer to Figure 7 for details, TEM analysis indicated that secondary nanotwins
formed within dislocation clusters during low-temperature deformation, increasing the
strain-hardening capacity and contributing to improved hardness and flexibility. At low
temperatures, the De-HEAs demonstrated a higher strain-hardening ability than conventional metals, especially through the formation of twins and the interactions between
dislocations. KAM map analysis revealed that the increase in secondary nanotwins and


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dislocation density significantly enhanced the strain hardening, improving the alloy’s
mechanical performance.
pasiciyeeo igue oeais, Eaaysis iicaeasecoay aois 
formed within dislocation clusters during low-temperature deformation, increasing the 
strain-hardening capacity and contributing to improved hardness and flexibility. At low 
temperatures, the De-HEAs demonstrated a higher strain-hardening ability than conventional metals, especially through the formation of twins and the interactions between dislocations. KAM map analysis revealed that the increase in secondary nanotwins and dislocation density significantly enhanced the strain hardening, improving the alloy’s mechanical performance.
Figure 7. Tensile deformation structures of the alloy at 223 K. (a) Primary deformation twin formation; (b,c) secondary nanoscale twin development; (d) indexed results illustrating twin orientations and strain conditions [137].
Regarding manufacturing techniques, researchers [144] have utilized laser melting 
deposition (LMD) technology to prepare non-equimolar CoCrFeNiMo0.2 HEA and investigate their low-temperature mechanical properties. The findings indicated that varying 
laser power notably influenced the microstructural characteristics of the alloy. As the laser 
power increased, the columnar grains grew and exhibited noticeable anisotropy. At 1400 
W, the orientation and growth of the columnar grains became more pronounced. At 293 
K, enhancing the laser power led to a considerable increase in plasticity, from 37% to 51%, 
along with an improvement in TS. When tested at 77 K, the alloy exhibited a substantial 
increase in TS to 928 MPa, and flexibility reached 60% under a laser power of 1400 W. 
Examination of the fracture surfaces revealed characteristic ductile fracture patterns at 
both 293 K and 77 K. At 77 K, the fracture surfaces exhibited larger dimples and deeper 
tear marks, indicating an enhanced ability to deform plastically at low temperatures. Fracture surface examination revealed (Figure 8) typical ductile fracture features at both 293 
K and 77 K. At 77 K, the fracture surfaces displayed larger and deeper dimples and tear 
marks, indicating enhanced plastic deformation at low temperatures. Additionally, the 
side surface at 77 K showed more slip traces compared to 293 K, suggesting the activation 
of additional slip systems and increased plasticity. The study highlights the influence of 
laser power on the microstructure and performance of the alloy, demonstrating the potential of CoCrFeNiMo0.2 high-entropy alloy as a coating material for low-temperature applications. 
Figure 7. Tensile deformation structures of the alloy at 223 K. (a) Primary deformation twin formation;
(b,c) secondary nanoscale twin development; (d) indexed results illustrating twin orientations and
strain conditions [137].
Regarding manufacturing techniques, researchers [144] have utilized laser melting
deposition (LMD) technology to prepare non-equimolar CoCrFeNiMo0.2 HEA and investigate their low-temperature mechanical properties. The findings indicated that varying
laser power notably influenced the microstructural characteristics of the alloy. As the laser
power increased, the columnar grains grew and exhibited noticeable anisotropy. At 1400 W,
the orientation and growth of the columnar grains became more pronounced. At 293 K,
enhancing the laser power led to a considerable increase in plasticity, from 37% to 51%,
along with an improvement in TS. When tested at 77 K, the alloy exhibited a substantial
increase in TS to 928 MPa, and flexibility reached 60% under a laser power of 1400 W.
Examination of the fracture surfaces revealed characteristic ductile fracture patterns at both
293 K and 77 K. At 77 K, the fracture surfaces exhibited larger dimples and deeper tear
marks, indicating an enhanced ability to deform plastically at low temperatures. Fracture
surface examination revealed (Figure 8) typical ductile fracture features at both 293 K
and 77 K. At 77 K, the fracture surfaces displayed larger and deeper dimples and tear
marks, indicating enhanced plastic deformation at low temperatures. Additionally, the side
surface at 77 K showed more slip traces compared to 293 K, suggesting the activation of
additional slip systems and increased plasticity. The study highlights the influence of laser
power on the microstructure and performance of the alloy, demonstrating the potential of
CoCrFeNiMo0.2 high-entropy alloy as a coating material for low-temperature applications.
3.2. BCC-Phase HEAs
Crystal dislocation movement is more challenging in the BCC than in the FCC. BCC
crystals have only a few slip systems, making them less likely to undergo large-scale
plastic deformation at RT. However, under low-temperature conditions, with appropriate
composition design, BCC single-phase HEAs can enhance their plasticity and toughness
through mechanisms such as twinning or phase transformation, thereby improving their
resistance to brittle fracture. The earliest study [125] of BCC-phase HEAs at cryogenic
temperatures was conducted by Huang et al. in 2011 on AlCoCrFeNi HEA. The study
examined the compressive and tensile behaviors of the BCC-phase alloy at both room and
cryogenic temperatures. It was observed that lowering the temperature from 298 K to 77 K


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Coatings 2025, 15, 92 12 of 32
resulted in a 29.7% increase in YS and a 19.9% rise in fracture strength. Additionally, the
fracture mode transitioned from intergranular at 298 K to transgranular at 77 K.
Coatings 2025, 15, 92 12 of 34 
Figure 8. Fracture surface appearances under various conditions. (a) Entire fracture view at 293 K; 
(b) detailed fracture morphology at 293 K; (c) side view of the sample at 293 K; (d) entire fracture 
view at 77 K; (e) detailed fracture morphology at 77 K; (f) side view of the sample at 77 K [144]. 
3.2. BCC-Phase HEAs 
Crystal dislocation movement is more challenging in the BCC than in the FCC. BCC 
crystals have only a few slip systems, making them less likely to undergo large-scale plastic deformation at RT. However, under low-temperature conditions, with appropriate 
composition design, BCC single-phase HEAs can enhance their plasticity and toughness 
through mechanisms such as twinning or phase transformation, thereby improving their 
resistance to brittle fracture. The earliest study [125] of BCC-phase HEAs at cryogenic temperatures was conducted by Huang et al. in 2011 on AlCoCrFeNi HEA. The study examined the compressive and tensile behaviors of the BCC-phase alloy at both room and cryogenic temperatures. It was observed that lowering the temperature from 298 K to 77 K 
resulted in a 29.7% increase in YS and a 19.9% rise in fracture strength. Additionally, the 
fracture mode transitioned from intergranular at 298 K to transgranular at 77 K.
Leveraging the mechanical characteristics of HEAs in cold environments, researchers 
have explored how low-temperature conditions influence the machining quality and mechanical properties of FeCoCrNiAl0.6 HEA. The research indicates that higher cutting 
speeds enhance the surface quality of specimens at RT, whereas greater cutting depths 
lead to a deterioration in surface finish. In low-temperature conditions at 153 K, increasing 
cutting speed also improves surface quality, which is superior to machining at RT. Increasing the speed from 2000 to 2800 mm/min notably diminishes chip adhesion and pileup edges. Additionally, the area covered by surface pits and defects is reduced, resulting 
in significantly better surface quality. Figure 9 displays the alloy’s surface morphology 
post-machining under various temperature conditions. Figure 9a,b illustrate the alloy’s 
surface after machining at RT, showing prominent built-up edges and pit defects that 
compromise surface integrity. Conversely, Figure 7c,d present the surface morphology 
following machining at low temperatures. Here, the surface quality is notably enhanced, 
with fewer built-up edges and pit defects, leading to a smoother finish.
Figure 8. Fracture surface appearances under various conditions. (a) Entire fracture view at 293 K;
(b) detailed fracture morphology at 293 K; (c) side view of the sample at 293 K; (d) entire fracture
view at 77 K; (e) detailed fracture morphology at 77 K; (f) side view of the sample at 77 K [144].
Leveraging the mechanical characteristics of HEAs in cold environments, researchers
have explored how low-temperature conditions influence the machining quality and mechanical properties of FeCoCrNiAl0.6 HEA. The research indicates that higher cutting
speeds enhance the surface quality of specimens at RT, whereas greater cutting depths
lead to a deterioration in surface finish. In low-temperature conditions at 153 K, increasing
cutting speed also improves surface quality, which is superior to machining at RT. Increasing the speed from 2000 to 2800 mm/min notably diminishes chip adhesion and pile-up
edges. Additionally, the area covered by surface pits and defects is reduced, resulting
in significantly better surface quality. Figure 9 displays the alloy’s surface morphology
post-machining under various temperature conditions. Figure 9a,b illustrate the alloy’s
surface after machining at RT, showing prominent built-up edges and pit defects that
compromise surface integrity. Conversely, Figure 7c,d present the surface morphology
following machining at low temperatures. Here, the surface quality is notably enhanced,
with fewer built-up edges and pit defects, leading to a smoother finish.
Coatings 2025, 15, 92 13 of 34 
Figure 9. Surface morphology of the alloy post-machining. (a) Cutting speed (vc) = 2000 mm/min at 
RT; (b) vc = 2200 mm/min at RT; (c) vc = 2000 mm/min at low temperature; (d) vc = 2200 mm/min at 
low temperature [125].
In terms of fatigue performance, increasing cutting speed enhances fatigue life. At a 
cutting speed of 2000 mm/min, the fatigue life at RT is 142,100 cycles, while at 153 K, it 
increases to 319,071 cycles. Although increasing cutting depth reduces fatigue life, fatigue 
life under low-temperature cutting conditions remains significantly better than at RT. 
Concerning mechanical properties, when the cutting speed is set to 2000 mm/min, both 
Figure 9. Surface morphology of the alloy post-machining. (a) Cutting speed (vc) = 2000 mm/min at
RT; (b) vc = 2200 mm/min at RT; (c) vc = 2000 mm/min at low temperature; (d) vc = 2200 mm/min
at low temperature [125].


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In terms of fatigue performance, increasing cutting speed enhances fatigue life. At a
cutting speed of 2000 mm/min, the fatigue life at RT is 142,100 cycles, while at 153 K, it
increases to 319,071 cycles. Although increasing cutting depth reduces fatigue life, fatigue
life under low-temperature cutting conditions remains significantly better than at RT.
Concerning mechanical properties, when the cutting speed is set to 2000 mm/min, both
YS and TS at low temperatures are enhanced by 6.2% and 10.4%, respectively, relative
to those measured at RT. Therefore, machining FeCoCrNiAl0.6 HEA in low-temperature
environments is feasible and can improve machining quality, mechanical properties, and
fatigue life.
The refractory HEA iZrHfNbTa comprises a single BCC phase. According to Wang
et al.’s [145], at 77 K, the alloy preserves a tensile elongation of 20.8%, demonstrating
increase in YS, reaching 1549 MPa, without any noticeable ductile–brittle transition. At
RT, the alloy’s deformation mainly occurs through dislocation slip. In contrast, at low
temperatures, the deformation mechanisms include the formation of nanoscale twins,
phase transformation induced by deformation, and dislocation slip. Under applied stress,
ω-phase precipitates emerge along the twin boundaries. Mechanical twinning and ω-phase
transformation begin within a temperature window of approximately 227–277 K. HU
et al. [146] observed that when tested at 77 K, the alloy retained a ductile fracture pattern
despite being subjected to high strain rates. Compression testing revealed that elevating
the strain rate from 400 s−1to 2600 s−1raised the flow stress from 1294 MPa to 1760 MPa,
demonstrating a clear strain-rate strengthening effect. At 77 K and a strain rate of 2600 s−1,
the TiZrHfNbTa HEA exhibited a mixed-mode fracture surface, combining features of both
ductile and brittle behavior.
3.3. Multi-Structural HEAs
By controlling the composition of HEAs, multi-structural alloys with excellent properties can be produced. These alloys consist of a softer matrix, typically FCC, with harder
phases like σ, L12, BCC, and sometimes HCP structures. Small grain size and precipitated
hard phases at low temperatures help suppress twinning, making the material more stable.
The strength and toughness of these alloys arise from several mechanisms, including stacking fault strengthening, phase transformation toughening, second-phase strengthening,
and precipitation strengthening. Together, these mechanisms provide high strength and
toughness at low temperatures, enhancing resistance to brittle fracture. Hee et al. [147]
designed an iron-rich VCxMnFeCoNi HEA. As the manganese content rose, structural
stability diminished, leading to the emergence of a brittle intermetallic σ phase. At ambient
conditions, this HEA has a YS of 544 MPa, which climbs to 766 MPa at 77 K. The elongation
at RT is 46.2%, improving to 54.1% at low temperatures. However, as the manganese
content increases, the improvement in elongation at low temperatures becomes smaller.
Introducing the σ phase enhances alloy strength by anchoring grain boundaries and impeding their movement. Still, it also suppresses the formation of twinning, which reduces the
flexibility at low temperatures. The σ phase’s pinning effect on the grain boundaries helps
to inhibit grain sliding and dislocation motion, thereby enhancing the material’s strength.
Such phase transformation mechanisms and the influence of grain boundary pinning
are crucial factors to designing HEAs with enhanced low-temperature properties, as demonstrated in the study by Hee et al. [147]. They designed an iron-rich VCxMnFeCoNi HEA.
As Mn concentration rose, the alloy’s crystalline framework grew less stable, eventually
producing a brittle σ intermetallic phase. Under room-temperature conditions, the alloy’s
YS measures 544 MPa, rising to 766 MPa when cooled to 77 K. The elongation at RT is 46.2%,
improving to 54.1% at low temperatures. However, as the manganese content increases,
the improvement in elongation at low temperatures becomes smaller. By forming the σ


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phase, the alloy gains strength as these precipitates effectively anchor grain boundaries and
hinder their motion. Still, it also suppresses the formation of twinning, which reduces the
flexibility at low temperatures. The σ phase’s pinning effect on the grain boundaries helps
to inhibit grain sliding and dislocation motion, thereby enhancing the material’s strength.
Non-equimolar composition design has also been studied to improve low-temperature
properties in high-entropy alloys. Zhang et al. [148] engineered a Fe28.2Ni7Co11Al2.5Ta0.04B
HEA and employed thermomechanical treatments to achieve a heterogeneous microstructure comprising an FCC matrix and a γ’ phase, where NiAl B2-type particles precipitated
along grain boundaries. Under ambient conditions, the alloy’s TS reaches 1.43 GPa, accompanied by 21% elongation. Cooling to 77 K elevates the TS to around 2.2 GPa, while leaving
the alloy’s ductility largely unaffected. The strain hardening rates at 77 K and 298 K are
over 15 GPa and 4 GPa, respectively.
Figure 10 illustrates the martensitic transformation process in the Fe28.2Ni7Co11Al2.5Ta0.04B
HEA under different tensile strains. At approximately 7% tensile strain (Figure 10a), residual
martensitic regions are mainly observed in the fine-grain areas, as indicated by the blue arrows.
These findings imply that localized stress concentrations in fine-grained regions propel the
transformation, highlighting the heterogeneous structure’s uneven deformation behavior. As
the strain increases to about 12% (Figure 10b), martensite almost completely extends throughout
the fine-grain regions. Also, it spreads into the coarse-grain regions, as indicated by the red
arrows. By adopting a staged martensitic transformation sequence, the alloy averts a premature
grain-boundary failure, ultimately improving its overall ductility. The images in Figure 10c,e
show two different martensitic variants, with the thickness of the martensite in the coarsegrain region being approximately 50 nm and in the fine-grain region approximately 200 nm.
Figure 10d,f show a dense network of nanotwins forming in the martensitic regions, featuring
twin spacings of roughly 6–10 nm and minimal stacking fault presence at twin interfaces.
The twin boundaries serve as potent barriers to dislocation motion, thereby reinforcing the
alloy’s strain-hardening capacity under cryogenic conditions. Moreover, at 77 K the alloy
demonstrates pseudoelasticity: following martensitic transformation under load, it reverts to
its original austenitic state once unloaded, exhibiting superelastic-like recovery characteristics.
Coatings 2025, 15, 92 15 of 34 
Figure 10. The microstructure of Fe28.2Ni7Co11Al2.5Ta0.04B HEA after tensile testing at 77 K. (a) Martensite forms mainly in fine grains at ~7% strain. (b) At ~12% strain, martensite spreads to fine and 
coarse grains. (c) TEM shows thin-plate martensite in coarse grains. (d) High-resolution TEM shows 
nanotwins in martensite. (e) TEM shows thin-plate martensite in fine grains. (f) Nanotwins in finegrain martensite [147].
Drawing upon insights into low-temperature phase transformations, Bhattacharjee 
et al. [127] conducted a comprehensive analysis of the microstructure and mechanical behavior of the AlCoCrFeNi2.1 EHEA under various deformation temperatures. The EHEA 
develops a finely layered microstructure featuring alternating L12 and B2 phases with 
mean thicknesses of about 600 nm and 200 nm. Cooling triggers a phase shift in the L12 
regions, causing them to revert into a disordered FCC configuration. At temperatures 
around 363 K and 423 K, the B2 phase retains its ordered BCC structure. However, at 
Figure 10. The microstructure of Fe28.2Ni7Co11Al2.5Ta0.04B HEA after tensile testing at 77 K.
(a) Martensite forms mainly in fine grains at ~7% strain. (b) At ~12% strain, martensite spreads to fine
and coarse grains. (c) TEM shows thin-plate martensite in coarse grains. (d) High-resolution TEM
shows nanotwins in martensite. (e) TEM shows thin-plate martensite in fine grains. (f) Nanotwins in
fine-grain martensite [147].
Drawing upon insights into low-temperature phase transformations, Bhattacharjee
et al. [127] conducted a comprehensive analysis of the microstructure and mechanical
behavior of the AlCoCrFeNi2.1 EHEA under various deformation temperatures. The EHEA


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Coatings 2025, 15, 92 15 of 32
develops a finely layered microstructure featuring alternating L12 and B2 phases with mean
thicknesses of about 600 nm and 200 nm. Cooling triggers a phase shift in the L12 regions,
causing them to revert into a disordered FCC configuration. At temperatures around 363 K
and 423 K, the B2 phase retains its ordered BCC structure. However, at lower temperatures,
particularly at 77 K, the L12 phase completely disorder and transforms into a disordered
FCC phase. TEM analyses (Figure 11) confirm that the B2/FCC phase boundary remains
notably smooth and uninterrupted, the yellow arrows indicate the lattice bands.
gaiaeie []
Drawing upon insights into low-temperature phase transformations, Bhattacharjee 
et al. [127] conducted a comprehensive analysis of the microstructure and mechanical behavior of the AlCoCrFeNi2.1 EHEA under various deformation temperatures. The EHEA 
develops a finely layered microstructure featuring alternating L12 and B2 phases with 
mean thicknesses of about 600 nm and 200 nm. Cooling triggers a phase shift in the L12 
regions, causing them to revert into a disordered FCC configuration. At temperatures 
around 363 K and 423 K, the B2 phase retains its ordered BCC structure. However, at 
lower temperatures, particularly at 77 K, the L12 phase completely disorder and transforms into a disordered FCC phase. TEM analyses (Figure 11) confirm that the B2/FCC 
phase boundary remains notably smooth and uninterrupted, the yellow arrows indicate 
the lattice bands. 
Figure 11. (a) The AlCoCrFeNi2.1 alloy fractured in tension at 77 K, (b) a SAED pattern from the 
area indicated by a green marker, and (c) an SAED pattern from the region indicated by a red 
marker. At higher magnifications, dislocation arrangements are clearly visible in both the (d) FCC 
and (e) B2 domains [127].
Figure 11. (a) The AlCoCrFeNi2.1 alloy fractured in tension at 77 K, (b) a SAED pattern from the area
indicated by a green marker, and (c) an SAED pattern from the region indicated by a red marker. At
higher magnifications, dislocation arrangements are clearly visible in both the (d) FCC and (e) B2
domains [127].
As deformation temperature drops, yield and ultimate TSs both rise. Meanwhile,
total elongation, holding steady at around 17%–20%, shows little sensitivity to the colder
environment. At low temperatures, the dislocation movement in the alloy is suppressed,
making plastic deformation more difficult, enhancing strain hardening and increasing
strength. Additionally, transforming the L12 phase to the disordered FCC phase at low
temperatures makes the FCC phase more ductile and tougher, contributing to improved
strength and retained ductility.
Although the B2 phase is typically harder to deform, it also experiences strain hardening at low temperatures, likely due to the interaction between the B2 and FCC phases
at their interfaces. A profusion of dislocations accumulates along the B2/FCC interface,
suggesting that their mutual interaction aids the B2 phase in accommodating plastic deformation under cryogenic conditions. Cooperative strain-hardening contributions from
both the disordered FCC and B2 phases confer exceptional low-temperature mechanical
properties to the AlCoCrFeNi2.1 EHEA.
4. High-Temperature Properties of HEAs
High-temperature investigations into HEAs initially targeted the stringent requirements for robust strength and superior structural stability, particularly crucial in aerospace
applications [149]. Researchers aimed to develop high-temperature materials superior
to traditional high-nickel superalloys. Starting in 2010, refractory HEAs [150] became a
key focus of study. Subsequently, researchers [135] employed advanced manufacturing
techniques, such as powder metallurgy and laser melting deposition, to develop HEAs
suitable for high-temperature environments. The scope of research has expanded, includ-


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ing applications in coatings and low-density alloys, to meet various industries’ specific
requirements. In high-temperature environments, the performance of HEAs includes not
only strength and structural stability but also key properties such as thermal stability,
friction performance, and oxidation resistance. Refining these attributes is pivotal for ensuring dependable performance in aerospace, energy systems, and other high-temperature
technologies. This section surveys key investigations into both transition metal-based
and refractory HEAs at elevated temperatures, organizing the findings according to their
elemental placements in the periodic table.
4.1. Transition Metal HEAs
In studying the high-temperature performance of transition metal HEAs, single-phase
HEAs mainly focus on the FCC-phase CoCrFeMnNi alloy. However, due to the relatively
low strength of the FCC phase, research has primarily concentrated on enhancing the
alloy’s high-temperature performance through methods such as precipitate strengthening
and phase transformation strengthening. This section will present the related research
findings, focusing on single-phase and multi-phase HEAs.
4.1.1. Single-Phase HEA
As previously noted, Kim et al. [135] examined the tensile behavior of CoCrFeMnNi
HEAs over an extensive range of temperatures. Later efforts concentrated on uncovering
how the alloy’s intrinsic deformation mechanisms evolve under high-temperature conditions. Ghosh et al. [151] delved into the dynamic recrystallization (DRX) processes in
CoCrFeMnNi HEAs, examining temperatures from 950 ◦C to 1100 ◦C under strain rates of
10−2s
−1 and 10−1
s
−1
. Figure 12 presents the true stress–true strain profiles across various
temperatures (K) and strain rates. Under a 10−2s
−1
strain rate, increasing the temperature
progressively lowers the flow stress. The stress–strain response shows distinct peak stresses
at strains of about 0.18, 0.20, and 0.23.
Coatings 2025, 15, 92 17 of 34 
Figure 12. Stress–strain relationships for the HEA under varying thermal conditions at strain rates 
of (a) 10−1 s⁻1 and (b) 10−2 s⁻1 [135]. 
After reaching the peak value, the flow stress begins to soften and eventually reaches 
a steady state, with steady-state strains of 0.45, 0.5, and 0.6. This indicates that the DRX 
process is more complete at lower strain rates, leading to stress softening and dynamic 
equilibrium. Similarly, at 10−1 s⁻1, the flow stress curve peaks before declining continuously without stabilizing, dropping off until the strain approaches roughly 0.7. This suggests that the DRX process is not fully developed at higher strain rates, and the dislocation 
density has not been effectively reduced. Overall trends indicate that heightened strain 
rates raise flow stress, while elevated temperatures diminish it. Lower strain rates and 
higher K promote the occurrence of dynamic recrystallization, leading to stress softening 
and the emergence of steady-state behavior. It can be observed that at high K, dynamic 
recrystallization causes changes in grain size and induces significant changes in the microtexture. A higher DRX volume fraction and more pronounced microstructural evolution are observed under low strain rate conditions.
Jiang et al. [152] also probed the high-temperature oxidation characteristics of the 
Figure 12. Stress–strain relationships for the HEA under varying thermal conditions at strain rates of
(a) 10−1s
−1 and (b) 10−2
s
−1
[135].
After reaching the peak value, the flow stress begins to soften and eventually reaches
a steady state, with steady-state strains of 0.45, 0.5, and 0.6. This indicates that the DRX
process is more complete at lower strain rates, leading to stress softening and dynamic equilibrium. Similarly, at 10−1
s
−1
, the flow stress curve peaks before declining continuously
without stabilizing, dropping off until the strain approaches roughly 0.7. This suggests that
the DRX process is not fully developed at higher strain rates, and the dislocation density has
not been effectively reduced. Overall trends indicate that heightened strain rates raise flow
stress, while elevated temperatures diminish it. Lower strain rates and higher K promote
the occurrence of dynamic recrystallization, leading to stress softening and the emergence
of steady-state behavior. It can be observed that at high K, dynamic recrystallization causes


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changes in grain size and induces significant changes in the microtexture. A higher DRX
volume fraction and more pronounced microstructural evolution are observed under low
strain rate conditions.
Jiang et al. [152] also probed the high-temperature oxidation characteristics of the
CrMnFeCoNi HEA. Post-oxidation examinations uncovered numerous pores dispersed
within the oxide scale. While these pores correlate strongly with oxide crystallization, they
cannot be attributed to the Kirkendall effect. Moreover, mismatches in thermal expansion
coefficients between the substrate and oxide layers induce residual stresses, ultimately
causing cracks and partial delamination of the oxide scale. The alloy undergoes three
main oxidation stages at high temperatures. At even higher temperatures (1050–1150 ◦C),
thermal stresses and expansion disparities lead to spallation, progressively thinning the
oxide layer. In the medium-temperature stage (950–1050 ◦C), the oxide layer gradually
thickens due to the chemisorption of oxygen atoms. Finally, in the high-temperature
stage (1050–1150 ◦C), the oxide layer peels off due to thermal stress and differences in the
coefficient of thermal expansion, resulting in a thinning of the oxide layer. Cross-sectional
analysis revealed a layered oxide structure: an outermost zone enriched in MnFe2O4
and Mn3O4, an intermediate (Mn, Cr)xO4 region, and an inner Cr2O3 layer. While FCCphase HEAs show some resistance to high-temperature oxidation, there is still significant
potential for improvement. It is evident that the high-temperature performance of singlephase FCC HEAs still requires further improvement to meet the demands of applications
in extreme environments.
4.1.2. Multi-Phase HEA
In contrast to single-phase alloys, multi-phase HEAs gain reinforcement through
minor additions (e.g., Al, Ti, Ta, Nb) that precipitate strengthening second phases. This
approach allows the alloy to achieve higher strength and exhibit excellent thermal stability
and oxidation resistance within specific temperature ranges. For instance, incorporating
Al and Ti fosters the formation of γ
′ precipitates (e.g., Ni3Al), bolstering both the alloy’s
high-temperature resilience and thermal stability. Moreover, certain alloy systems further
improve performance by forming eutectic structures. By adjusting the composition, EHEAs
can introduce fine reinforcing phases into the matrix, effectively hindering dislocation
motion and improving the alloy’s YS, creep resistance, and high-temperature durability.
As previously discussed, single-phase HEAs like CoCrFeMnNi have shown potential
in high-temperature applications. However, they often require further enhancement due
to their relatively low strength in high-temperature environments. In this context, Joseph
et al. [153] investigated the design of Co-Cr-Fe-Ni-Al-Ti HEAs (HEAs), particularly focusing
on alloys with up to 20% Al and Ti content. At 1073 K, these alloys achieve a stable γ + γ
′
phase equilibrium structure, resulting in a high-entropy γ phase matrix. While raising the
(Al + Ti) concentration elevates γ
′ phase volume and strength, it must remain under 18%
to maintain uniform particle dispersion and prevent performance degradation. Keeping
the Al/Ti ratio within 0.8–3 effectively inhibits the formation of brittle intermetallics (e.g.,
NiAl, Ni2AlTi, Ni3Ti), thereby preserving the alloy’s favorable mechanical properties.
Moreover, restricting Fe, Cr, and Co within certain compositional boundaries helps
avert the emergence of detrimental phases such as σ. This alloy exhibits strength retention
of over 873 MPa at 1073 K, with an anti-coarsening performance comparable to some
commercial high-temperature alloys. Furthermore, it does not contain expensive elements
such as Re or Ta, which offers a cost-effectiveness advantage.
Based on this, Joseph et al. [154] developed Ni51Co18Fe5Cr10Al16-XTiX HEAs. Exceeding 60% in γ
′
-phase volume, this alloy maintains a robust γ/γ
′
equilibrium field between
1073 K and 1173 K. Notably, no residual brittle intermetallics (e.g., Ni2AlTi, Ni(Al, Ti), Ni3Ti,


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Coatings 2025, 15, 92 18 of 32
σ) are detected, confirming its strong thermal stability at elevated temperatures. By selecting an Al-to-Ti ratio between 0.8 and 3, the alloy achieves a substantial γ
′
-phase fraction, a
γ
′
solvus above 1457 K, and a broad equilibrium range for the γ/γ
′ phases. Mechanically,
this composition surpasses 800 MPa in TS across the 1073–1173 K interval. Although raising
Ti levels can initially lower the solidus temperature, adjusting the Ti content or the Al/Ti
ratio can boost both the γ
′
fraction and the solidus temperature, granting more flexibility
in tuning the alloy’s high-temperature behavior. The study also revealed that reducing the
aluminum-to-titanium ratio improves the YS. In Ni51Co18Fe5Cr10Al8Ti8, pronounced lattice
mismatch between γ and γ
′ phases, together with fine precipitates, augments precipitation
hardening and elevates its YS. In summary, fine tuning the Al/Ti ratio refines the alloy’s
microstructure and enhances performance by boosting γ
′
-phase content, widening the γ
′
stability domain, and raising YS.
Similarly, Zhang et al. [155] examined γ-reinforced Ni45x(FeCoCr)40(AlTi)15Hfx
(x = 0, 0.2) HEAs, featuring a high-entropy γ matrix and uniformly dispersed γ
′ precipitates in their as-cast state. Their γ
′ precipitates average about 52 nm within dendritic
arms and roughly 90 nm in interdendritic areas. Enriched with γ
′ precipitates, these alloys
maintain superior mechanical attributes at both ambient and elevated temperatures. For
example, Ni44.8(FeCoCr)40(AlTi)15Hf0.2 delivers an impressive 961.6 MPa at 750 ◦C and
retains 696.6 MPa at 1123 K. Remarkably, the Hf0.2 variant achieves strength levels rivaling
traditional Ni-based superalloys—without Mo, W, or carbide additions. The γ
′ phase
fortifies the alloy at high temperatures through coherency-driven reinforcement and the
stability of its ordered lattice.
Jaladurgam et al. [128] examined as-cast AlCoCrFeNi2.1 EHEA at 973 K, investigating
how load sharing shifts between L12 and B2 phases under high-temperature deformation. Their findings reveal a temperature-dependent role reversal: at low temperatures
(77–673 K), L12 is weaker, but at 973 K it emerges as the primary strengthening phase. In
contrast, the B2 phase exhibited low strength and nearly ideal plastic behavior, indicating
a significant change in its deformation mechanism at elevated temperatures. Figure 13
depicts the alloy’s lamellar architecture in the as-cast state alongside its 973 K stress–strain
response, featuring a 216 MPa YS.
Coatings 2025, 15, 92 19 of 34 
significant change in its deformation mechanism at elevated temperatures. Figure 13 depicts the alloy’s lamellar architecture in the as-cast state alongside its 973 K stress–strain 
response, featuring a 216 MPa YS. 
Figure 13. (a) The AlCoCrFeNi2.₁ microstructure, (b) compositional mappings, and (c) its stress–
strain curve at 973 K [135]. 
They also established that cube slip predominates in L12, causing (111)-oriented 
grains—aligned perpendicularly to the tensile direction—to yield prematurely in early 
deformation. As deformation progressed, these grains gradually lost their load-carrying 
capacity, with the load being progressively transferred to grains with (200) and (311) orientations. In contrast, the stress distribution across orientations in the B2 phase showed 
no significant load transfer, differing markedly from its behavior at low temperatures. 
Additionally, the experiments detected possible twinning in the alloy, but its influence 
appeared to be limited, with cube slip remaining the primary deformation mechanism. 
In the field of coatings, Joseph et al. [156] studied the wear resistance and mechanisms 
of AlxCoCrFeNi coatings (x = 0.3, 0.6, 1.0 at %) across a temperature range from room 
temperature to 900 °C. The steady-state friction coefficient decreased with increasing temperature for all HEA coatings. Among them, AlCoCrFeNi demonstrated the highest wear 
Figure 13. (a) The AlCoCrFeNi2.1 microstructure, (b) compositional mappings, and (c) its stress–strain
curve at 973 K [135].
They also established that cube slip predominates in L12, causing [111]-oriented
grains—aligned perpendicularly to the tensile direction—to yield prematurely in early
deformation. As deformation progressed, these grains gradually lost their load-carrying
capacity, with the load being progressively transferred to grains with [200] and [311]
orientations. In contrast, the stress distribution across orientations in the B2 phase showed
no significant load transfer, differing markedly from its behavior at low temperatures.


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Coatings 2025, 15, 92 19 of 32
Additionally, the experiments detected possible twinning in the alloy, but its influence
appeared to be limited, with cube slip remaining the primary deformation mechanism.
In the field of coatings, Joseph et al. [156] studied the wear resistance and mechanisms
of AlxCoCrFeNi coatings (x = 0.3, 0.6, 1.0 at %) across a temperature range from room
temperature to 900 ◦C. The steady-state friction coefficient decreased with increasing
temperature for all HEA coatings. Among them, AlCoCrFeNi demonstrated the highest
wear resistance at all tested temperatures. At 600 ◦C and 800 ◦C, its wear performance
matched that of Inconel 718, and at 900 ◦C, the AlxCoCrFeNi coatings (x = 0.3, 0.6, 1.0)
outperformed Inconel 718. At 900 ◦C, it was observed that coatings with lower aluminum
content developed an outer Cr2O3 oxide layer and an inner Al2O3 layer, whereas higher
aluminum content resulted in a dense Al2O3 oxide layer. Sub-surface analysis beneath the
oxide layer revealed the presence of Cr-rich precipitates in AlCoCrFeNi coatings between
600 ◦C and 900 ◦C, with precipitate size increasing at higher temperatures. Furthermore,
prolonged exposure to temperatures between 600 ◦C and 960 ◦C led to the formation of a
body-centered tetragonal σ phase, which enhanced the mechanical properties of CoCrFeNi.
Ma et al. [157] studied W0.5Ta0.3MoNbVAlTi1−xZrx (x = 0, 0.25, 0.5, 0.75, 1.0) HEA
coatings and found that increasing x had little effect on hardness but greatly improved
oxidation resistance. After 20 h at 800 ◦C, the coating with x = 0.75 showed the best
performance, about 6.2 times better than x = 0. The coatings developed a two-layer oxide
structure. The outer layer mainly contained Fe2O3, FeO, and V2O5, while the inner layer
composition varied with x. Higher x reduced V in the outer layer and decreased TiO2 and
MoO3 in the inner layer. The high PBR values of V2O5 and MoO3 indicated zirconium
minimized unstable oxides. Additionally, the dense Al2O3 film in the inner layer prevented
metal diffusion, enhancing stability.
4.2. Refractory HEAs
Conventional high-temperature alloys typically consist of an unordered FCC matrix,
which can incorporate nanoscale precipitates or an ordered FCC (L12) structure, enabling
good performance at elevated temperatures. In contrast, refractory HEAs feature an
unordered BCC matrix and form similar microstructures, further enhancing their hightemperature properties. Additionally, these alloys can be strengthened using traditional
high-temperature methods, such as heat treatment, which precipitates hard secondary
phases that effectively hinder dislocation movement, thereby increasing strength. These
structures and processing techniques allow refractory HEAs to withstand greater forces
and pressures at high temperatures while maintaining superior mechanical properties. As a
result, refractory HEAs have risen to prominence as advanced high-temperature structural
materials, prompting extensive research into their mechanical robustness, thermal stability,
and oxidation resilience.
4.2.1. Mechanical Properties
In line with the findings on refractory HEAs, studies have explored these alloys’
mechanical performance and deformation behavior at high temperatures. For example,
Wang et al. [158] showed that VxNbMoTa, a single-phase BCC refractory HEA, maintains
structural consistency from its melting point down to 623 K. Higher V concentrations
lead to markedly finer grain dimensions. For instance, boosting V content from 0.25 to
1.0 shrinks the grain size from roughly 830 µm to about 250 µm. At 1273 K, VNbMoTa
attains a YS of 811 MPa, while at room temperature it can sustain more than 25% tensile
fracture strain, exemplifying remarkable strength–ductility synergy.
In parallel work, Couzinié et al. [159] probed deformation mechanisms in Al0.5Nb
Ta0.8Ti1.5V0.2Zr at 873 K. They reported a 1186 MPa YS at 873 K and observed the nucle-


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Coatings 2025, 15, 92 20 of 32
ation and growth of B2 precipitates during mechanical loading. Meanwhile, dislocations
interacted with the precipitated B2 phase during deformation, particularly when a pair of
a/2<111> type dislocations sheared the B2 precipitates during their motion. This shearing interaction strengthened the alloy at high temperatures, enhancing its mechanical
properties and strength.
Liu et al. [160] synthesized NbTaHfTiZrV0.5 and evaluated its mechanical attributes
between ambient conditions and 1273 K. The alloy maintained a BCC structure at different
deformation temperatures, with lattice constants slightly increasing with temperature. The
alloy exhibited a typical dendritic structure, and its microstructure remained essentially unchanged after cold rolling and high-temperature tensile tests, with identical microhardness
both inside and between dendrites. The alloy was able to maintain high YS between 873 K
and 1073 K, with YSs of approximately 1141.55 MPa at 873 K, 1168.89 MPa at 973 K, and
1189.75 MPa at 1073.
Figure 14a presents tensile stress-strain profiles over various test temperatures, each
initially rising to a peak before steadily tapering off as deformation proceeds. Figure 14b
plots how YS, ultimate tensile strength (UTS), and fracture elongation (FE) evolve with
increasing temperature. TAs temperature climbs, YS wanes from over 1200 MPa to under
1000 MPa. Interestingly, between 873 K and 1073 K, a plateau persists above 800 MPa,
indicating a stable strength regime. This behavior correlates closely with changes in
dislocation activity. At elevated temperatures, deformation mainly hinges on dislocation
slip, making dislocation morphology and density key factors in determining mechanical
performance. For instance, at 673 K the exceptionally dense, intricate dislocation networks
underpin the alloy’s high ductility. For instance, at 673 K the exceptionally dense, intricate
dislocation networks underpin the alloy’s high ductility.
Coatings 2025, 15, 92 21 of 34 
Figure 14. (a) Tensile stress–strain curves at various temperatures. (b) YS, UTS, and fracture elongation versus temperature [160].
Further exploring the wear resistance of refractory HEAs, Alvi et al. [161] utilized 
spark plasma sintering to fabricate the CuMoTaWV HEA, whose phase structure consists 
of 80% BCC solid solution and 20% FCC phase. Hardness testing revealed that the vanadium-rich regions have a hardness of approximately 900 Hv, while the high-entropy 
phase regions exhibit a hardness of around 600 Hv. At 298 K, this material exhibited a 
remarkably low wear rate (~5 × 10⁻7 mm3/Nm) and a friction coefficient of 0.5 against steel 
balls (700–880 Hv hardness). However, the COF and wear rate showed significant temperature dependence when subjected to sliding wear tests using Si3N4 balls with a hardness of 1550 Hv at various temperatures. At 298 K, the wear rate was 4.0 × 10⁻3 mm3/Nm 
with a COF of 0.45. Increasing the temperature to 473 K resulted in a substantial rise in 
wear rate to 2.3 × 10⁻2 mm3/Nm and an increase in COF to 0.59, primarily due to adhesive 
wear and severe galling caused by material transfer.
At 673 K, the wear rate declined to ~5 × 10⁻3 mm3/Nm while COF rose to 0.67, a shift 
linked to Cu oxidation generating a lubricious CuO trilayer in the wear path. At 873 K, 
although COF dropped to 0.54, wear rate climbed to about 4.5 × 10⁻2 mm3/Nm due to extensive oxidation of Cu, Ta, W, and the emergence of V2O5 in V-enriched zones. The wear 
rate of the opposing ball also changed with temperature, rising from 2.19 × 10⁻6 mm3/Nm 
at 298 K to 4.94 × 10⁻6 mm3/Nm at 473 K, then dropping to 2.46 × 10⁻6 mm3/Nm at 673 K 
Figure 14. (a) Tensile stress–strain curves at various temperatures. (b) YS, UTS, and fracture elongation versus temperature [160].
Further exploring the wear resistance of refractory HEAs, Alvi et al. [161] utilized
spark plasma sintering to fabricate the CuMoTaWV HEA, whose phase structure consists of
80% BCC solid solution and 20% FCC phase. Hardness testing revealed that the vanadiumrich regions have a hardness of approximately 900 Hv, while the high-entropy phase
regions exhibit a hardness of around 600 Hv. At 298 K, this material exhibited a remarkably
low wear rate (~5 × 10−7 mm3/Nm) and a friction coefficient of 0.5 against steel balls
(700–880 Hv hardness). However, the COF and wear rate showed significant temperature
dependence when subjected to sliding wear tests using Si3N4 balls with a hardness of
1550 Hv at various temperatures. At 298 K, the wear rate was 4.0 × 10−3 mm3/Nm with a
COF of 0.45. Increasing the temperature to 473 K resulted in a substantial rise in wear rate
to 2.3 × 10−2 mm3/Nm and an increase in COF to 0.59, primarily due to adhesive wear
and severe galling caused by material transfer.
At 673 K, the wear rate declined to ~5 × 10−3 mm3/Nm while COF rose to 0.67, a shift
linked to Cu oxidation generating a lubricious CuO trilayer in the wear path. At 873 K,


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Coatings 2025, 15, 92 21 of 32
although COF dropped to 0.54, wear rate climbed to about 4.5 × 10−2 mm3/Nm due to
extensive oxidation of Cu, Ta, W, and the emergence of V2O5 in V-enriched zones. The wear
rate of the opposing ball also changed with temperature, rising from 2.19 × 10−6 mm3/Nm
at 298 K to 4.94 × 10−6 mm3/Nm at 473 K, then dropping to 2.46 × 10−6 mm3/Nm at
673 K and 1.39 × 10−6 mm3/Nm at 873 K. Microstructural analysis indicated that at RT,
a tribolayer rich in Ta and W oxides formed within the wear track. The wear rate of
the opposing ball also changed with temperature, rising from 2.19 × 10−6 mm3/Nm at
298 K to 4.94 × 10−6 mm3/Nm at 473 K, then dropping to 2.46 × 10−6 mm3/Nm at 673 K
and 1.39 × 10−6 mm3/Nm at 873 K. By 873 K, Ta/W oxides and expanded V2O5 regions
lowered COF, but rampant oxidation marginally raised the wear rate again. Overall,
the CuMoTaWV HEA exhibited adaptive wear behavior across the temperature range
of 298 K to 873 K. Its microstructure, and the formation of protective oxide layers at
different temperatures significantly influenced its tribological performance, endowing it
with excellent wear-resistant properties in high-temperature environments.
4.2.2. High-Temperature Thermal Stability
Metal materials can experience thermal degradation at high temperatures, such as
softening, creep, and oxidation, which reduce their strength, hardness, and toughness.
However, an effective alloy design can maintain refractory HEAs’ stability and performance under high temperatures. Subsequent investigations emphasize how compositional
control, microstructural refinement, and tailored processing enhance refractory HEAs’
high-temperature stability.
For example, Jhong et al. [162] investigated a ternary alloy system composed of aluminum (Al), niobium (Nb), and vanadium (V), and constructed the phase diagram of the
system. By performing a one-month post-annealing treatment on the alloy at temperatures
of 1073 K or 1273 K, they determined the equilibrium state of the alloy. They mapped the
distribution regions of various phases (such as BCC solid solution, AlNb2, AlNb3, Al3Nb,
and Al3V) in the phase diagram. For example, Al30V35Nb35 forms nanoscale AlNb2 precipitates that lower its thermal conductivity. In the Al25V25Nb25Cr25 and Al20V20Nb20Cr20Ti20
alloys, the Laves C14 phase also contributes to a reduction in thermal conductivity. He
et al. [163] discovered that CoCrFeNiNbx HEAs retain a stable lamellar architecture at
annealing temperatures under 873 K, underscoring their structural persistence. Even when
heated to 1127 K, the lamellar framework retained commendable mechanical performance.
Elevating the annealing temperature from 873 K to 1127 K softened the alloy, reducing its
hardness from approximately 600 HV to 500 HV. Despite this softening, the material’s compressive strength remained near 2.3 GPa, suggesting minimal alterations to its mechanical
attributes at elevated temperatures.
These studies primarily focus on phase stability and microstructural evolution. At
the same time, other research also investigates the influence of specific structural features,
such as antiphase boundaries and nanoscale precipitates, on high-temperature properties.
Pang et al. [164] examined how the Nb40Ti25Al15V10Ta5Hf3W2 HEA resisted structural and
mechanical degradation at elevated temperatures. Aging at 1023 K for 120 h prompted
the development of nanoscale, Hf-rich segregations along grain boundaries. Throughout
the entire aging interval (923–1123 K), the alloy largely preserved its single-phase BCC
configuration. After aging for 120 h at 923 K, a high density of antiphase boundaries (APBs)
appeared within the alloy. Following aging at 1023 K for 120 h, nanoscale segregations
enriched with hafnium (Hf) formed at the grain boundaries. Moreover, aging at 1123 K for
the same duration induced the appearance of minor tetragonal (Nb, Hf)2Al precipitates
near grain boundaries and Hf-enriched zones. Figure 15 presents transmission electron
microscopy (DF-TEM) images under different temperatures and aging times, specifically


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Coatings 2025, 15, 92 22 of 32
showing the high density of APBs after aging for 120 h at 923 K (Figure 15a), the Hf-rich
nanoscale segregations after aging for 120 h at 1023 K (Figure 15b), and both the Hf-rich
segregations and tetragonal phase precipitates after aging for 120 h at 1123 K (Figure 15c).
Although antiphase boundaries (APBs) were abundant at 923 K, the alloy’s YS showed
no meaningful reduction. Figure 15d–j is the compositional scanning image of Nb, Ti, Al,
V, Ta, Hf, and W, respectively. Even with a high density of APBs at 923 K, the YS of the
alloy did not exhibit significant changes. Across 923–1123 K, the Nb40Ti25Al15V10Ta5Hf3W2
alloy exhibited remarkable structural and mechanical resilience. Its stability stemmed
from a robust high-temperature microstructure and the sparse formation of low-volume
precipitates, ensuring negligible effects on mechanical integrity.
figuration. After aging for 120 h at 923 K, a high density of antiphase boundaries (APBs) 
appeared within the alloy. Following aging at 1023 K for 120 h, nanoscale segregations 
enriched with hafnium (Hf) formed at the grain boundaries. Moreover, aging at 1123 K 
for the same duration induced the appearance of minor tetragonal (Nb, Hf)2Al precipitates 
near grain boundaries and Hf-enriched zones. Figure 15 presents transmission electron 
microscopy (DF-TEM) images under different temperatures and aging times, specifically 
showing the high density of APBs after aging for 120 h at 923 K (Figure 15a), the Hf-rich 
nanoscale segregations after aging for 120 h at 1023 K (Figure 15b), and both the Hf-rich 
segregations and tetragonal phase precipitates after aging for 120 h at 1123 K (Figure 15c). 
Although antiphase boundaries (APBs) were abundant at 923 K, the alloy’s YS showed no 
meaningful reduction. Figure 15d–j is the compositional scanning image of Nb, Ti, Al, V, 
Ta, Hf, and W, respectively. Even with a high density of APBs at 923 K, the YS of the alloy 
did not exhibit significant changes. Across 923–1123 K, the Nb40Ti25Al15V10Ta5Hf3W2 alloy 
exhibited remarkable structural and mechanical resilience. Its stability stemmed from a 
robust high-temperature microstructure and the sparse formation of low-volume precipitates, ensuring negligible effects on mechanical integrity. 
Figure 15. TEM micrographs of the Nb40Ti25Al15V10Ta5Hf3W2 alloy after 120 h of aging at 923 K (a), 
1023 K (b), and 1123 K (c), illustrating APB development and Hf-enriched segregations [164]. 
Beyond experimental efforts, computational research has probed how localized 
chemical ordering influences phase stability and microstructural evolution. For instance, 
Qiu et al. [165] employed simulations to unravel how local compositional ordering affects 
Figure 15. TEM micrographs of the Nb40Ti25Al15V10Ta5Hf3W2 alloy after 120 h of aging at 923 K (a),
1023 K (b), and 1123 K (c), illustrating APB development and Hf-enriched segregations [164].
Beyond experimental efforts, computational research has probed how localized chemical ordering influences phase stability and microstructural evolution. For instance, Qiu
et al. [165] employed simulations to unravel how local compositional ordering affects
both structural robustness and mechanical behavior in AlNbVTiZr alloys. Their findings
revealed that imposing greater atomic order boosted TS from 3.9 GPa (disordered) to
4.1 GPa (ordered), demonstrating that enhanced chemical order bolsters alloy strength.
Ordering influenced thermodynamics as well: between 0 and 1500 K, the ordered arrangement showed lower configurational entropy (9.87 J/(mol·K)) than the disordered variant
(13.38 J/(mol·K)), underscoring the ordered state’s greater thermodynamic stability. Therefore, the ordered structure exhibits better thermodynamic stability at high temperatures
than the disordered structure.
4.2.3. High-Temperature Oxidation Resistance
Refractory HEAs have improved oxidation resistance due to factors like composition,
crystal structure, and surface conditions, which are influenced by alloy composition, heat
treatment, and surface treatment. Elements like Cr, Al, and Ti form dense oxide layers that
prevent oxygen infiltration and enhance oxidation resistance. Incorporating high-meltingpoint components (Hf, Mo, Nb, Ta, W, Zr) together with oxidation-resistant elements (Al,
Cr, Ti, Si) notably improves the oxidation tolerance of refractory HEAs. Gorr et al. [166]
investigated oxidation kinetics in W-Mo-Cr-Ti-Al, Nb-Mo-Cr-Ti-Al, and Ta-Mo-Cr-Ti-Al
refractory HEAs at 1273 K and 1373 K. They found that all three alloys formed uneven, thick,
porous oxide layers on their surfaces. Of these alloys, only Ta-Mo-Cr-Ti-Al showed robust
oxidation resistance at both temperatures, adhering to a parabolic oxidation-rate profile.
This performance arose from the formation of a thin, tightly packed alumina-enriched
oxide film.
For another refractory TiZrHfNbTaV HEA [167], the focus was on its high-temperature
oxidation behavior between 1073 K and 1273 K, particularly the structural characteristics
of the oxide scale and its evolution at different temperatures. Oxidation tests in air were
conducted at 1073 K, 1173 K, and 1273 K for one hour each. After high-temperature
oxidation, the alloy maintained a BCC structure, demonstrating high stability at elevated
temperatures. Under 1073 K and 1173 K conditions, the oxide surface developed a loose,


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whisker-like morphology that facilitated ongoing oxygen penetration. At 1173 K, Ti and
Nb were enriched in the outermost layer, promoting the formation of TiNb2O7, but the
formation of HfO2 and ZrO2 simultaneously accompanied this. These brittle oxides tend
to form microcracks and pores during the oxidation process. At 1273 K, the oxide layer
was divided into an outer layer and an inner layer, with the inner layer exhibiting high
porosity, which further accelerated oxygen permeation into the alloy substrate. Under
these circumstances, (Hf, Zr) O2 crystallized into a dendritic pattern that encouraged pore
formation and undermined the oxide layer’s protective capabilities.
Moreover, Nb and Ti segregated at the oxide surface as oxidation progressed. At the
same time, Zr and Hf were depleted within the oxide layer, resulting in uneven elemental
distribution within the alloy. Such elemental clustering diminished oxide density and
protective efficacy, eroding the HEA’s overall oxidation resistance. These findings also
confirm that refractory HEAs need to optimize their alloy composition and microstructure
in high-temperature oxidative environments to enhance the density and stability of the
oxide layer, thereby improving their overall oxidation resistance.
5. Summary and Outlook
The present review summarizes how HEAs adapt under extreme environments and
discusses how thermal treatments at both ends of the temperature spectrum refine their performance. In low-temperature environments, rational compositional design enables HEAs
to achieve high flexibility and strength, with the plasticity and toughness of single-phase
alloys even improving at low temperatures. Multi-phase HEAs maintain high strength,
flexibility, and toughness at low temperatures. Activating deformation mechanisms such as
twinning, phase transformation, and stacking faults helps to achieve superior mechanical
properties with smaller strains and reduce plasticity loss. Low-temperature treatments and
processing techniques can further enhance performance by controlling temperature and
applying appropriate force and deformation.
HEAs exhibit superior or comparable high-temperature stability in high-temperature
environments compared to traditional nickel- and cobalt-based high-temperature alloys.
In transition metal multi-phase HEAs, γ
′ precipitate-strengthened alloys have a higher
degree of alloying, significantly enhancing high-temperature strength and creep resistance.
Adjustments in lattice distortion and antiphase boundary energies within the γ
′ phase
enhance high-temperature behavior. Because of their lofty melting points and robust
thermal stability, insoluble HEAs retain strong wear and oxidation resistance under hot
operating conditions. The mechanical properties and durability under high-temperature
conditions can be further optimized through high-temperature treatments and advanced
alloy design methods.
Beyond bulk materials, coatings based on HEAs have emerged as a promising approach to enhancing performance under extreme conditions. HEA coatings demonstrate
excellent thermal stability, wear resistance, and oxidation resistance at elevated temperatures, making them highly suitable for applications such as turbine blades, aerospace
components, and nuclear reactors. In cryogenic environments, HEA coatings can mitigate surface embrittlement, reduce friction, and improve wear resistance, ensuring the
long-term stability and durability of structural materials. Future research should focus on
developing advanced fabrication methods for HEA coatings, such as additive manufacturing and thermal spraying, to enhance their performance and adaptability in extreme
environments. Additionally, understanding the interfacial bonding between coatings and
substrates and optimizing the microstructure of HEA coatings will be critical to improving
their mechanical properties and reliability.


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Although HEAs have made progress in compositional tunability and diverse deformation mechanisms, breaking through the performance limits of traditional materials, there is
still room for optimization in their application under extreme conditions. Precisely controlling the microstructure and optimizing deformation mechanisms to maximize performance
under extreme environments remain frontier topics in HEA research. Novel manufacturing
processes such as additive manufacturing and high-pressure torsion have opened new
avenues for developing HEAs while optimizing processing techniques and alloy design to
better adapt to extreme environments. This is an important future research direction. Additionally, as multi-component alloys in the intermediate phase region, the compositional
optimization of HEAs remains challenging. The following points need further attention:
1. The specific impact of multiple deformation mechanisms on alloy performance under
extreme environments.
Numerous studies have introduced traditional material toughening mechanisms into
HEAs, significantly improving their service performance. Nonetheless, pinpointing the
exact roles and relative contributions of various deformation processes in harsh settings
remains essential for engineering stronger, tougher HEAs. Moreover, though precipitate
strengthening improves strength, prolonged aging at elevated temperatures can trigger
abnormal precipitate growth and irregular dispersion, diminishing both strength and
creep resistance at high temperatures. Hence, sustaining higher dislocation densities
alongside effective precipitate strengthening becomes imperative. At lower temperatures,
precipitate-induced reinforcement also boosts strength, yet aging can drastically lower
matrix dislocation densities, ultimately depressing YS. Thus, enhancing the dislocation
density in the matrix is necessary while maintaining the effect of precipitate strengthening.
With the advancement of experimental and computational technologies, a deeper understanding of the interactions between residues and the matrix during deformation, dynamic
recrystallization behavior, and the diffusion and partitioning mechanisms of elements at
grain boundaries and phase interfaces will be achieved. This will aid in designing HEAs
with excellent mechanical properties under extreme environments.
In the context of HEA coatings, various microscopic deformation mechanisms also
interact with each other. Under high-temperature conditions, if precipitation strengthening
occurs within the HEA coating, it is necessary to consider the effects of cyclic thermal
loading or high-temperature aging on the precipitates, as well as the impact of the thermal
expansion mismatch between the coating and substrate on interfacial bonding. At low
temperatures, attention should be paid to the effect of precipitates on dislocation motion,
the stacking fault energy, and microcrack formation within the coating. By optimizing
the microstructure and precipitation distribution within the coating, the durability of the
coating and the overall service performance of the substrate can be synergistically enhanced.
2. Mechanisms of toughening and plasticizing high-strength HEAs at high and low
temperatures.
Although some HEAs achieve YS of up to 2 GPa, most exhibit low-temperature
YSs below 1.5 GPa and do not yet outperform traditional low-temperature engineering
materials. Therefore, further research on HEA’s fundamental structures and deformation
mechanisms is essential to enhance their low-temperature performance. Key areas include
understanding the deformation mechanisms and toughening principles of ultra-highstrength FCC and BCC HEAs at low temperatures, examining the effects of different matrix
structures on toughness and plasticity, analyzing phase and twin interfaces, investigating
dislocation motion within the microstructure, and optimizing alloy compositions to balance
toughness and strength.


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At high temperatures, high-strength HEAs also face challenges with insufficient toughness and plasticity due to creep and phase transformations. Enhancing high-temperature
toughness requires optimized microstructure design, such as introducing high temperatureresistant strengthening phases, designing stable gradient structures, and creating multiphase composites. Nanoscale strengthening phases can impede dislocation movement and
crack propagation, while gradient structures help disperse stress concentrations, enhancing
overall toughness. Multi-phase composites synergistically improve ductility and creep
resistance through interactions between different phases. Understanding these toughening
mechanisms, including high-temperature dislocation dynamics and phase boundary behavior, will provide the foundation for designing HEAs with enhanced high-temperature
toughness and plasticity.
HEA coatings face similar issues in terms of toughening and plasticizing. The coating
may crack or delaminate under low-temperature impact loading. In high-temperature
service, phase transformations or creep can reduce adhesion and durability. Therefore,
strategies such as introducing nanoscale strengthening phases into the coating, optimizing
interfacial bonding, and designing multilayer or gradient coating architectures can improve
the coating’s toughness and plasticity, thereby enhancing its overall performance under
extreme temperatures.
3. Development and application of low-cost iron-based HEAs.
Due to their low cost and abundant resources, iron-based HEAs have broad application
prospects. However, the compositional design of iron-based HEAs is complex, and their
phase transformation behavior is more complicated to control accurately than martensitic
transformations in steel. Additionally, how to reduce the number and content of alloying
elements while maintaining performance remains a key issue to be addressed. Moreover,
Co, Ni, and Cr, which are in the same period as iron, are high-cost elements. During
development, a key research focus is how to reduce these high-cost elements while ensuring
the alloy’s performance under extreme environments. Iron-based HEAs can also be used
to fabricate coatings, thereby reducing overall manufacturing costs and leveraging the
advantages of iron-based systems, which are conducive to large-scale production.
4. Dynamic mechanical behavior and applications of HEAs at high and low temperatures.
The mechanical behavior of HEAs under dynamic loads, especially their responses
under high and low-temperature environments, requires further in-depth study. Dynamic
mechanical behavior involves the material’s elastic and plastic responses, energy absorption,
damage evolution, and fatigue performance. Integrating empirical tests with theoretical
modeling can offer a comprehensive understanding of HEAs’ dynamic mechanical responses across diverse temperatures and strain rates, establishing a scientific basis for their
use in severe service environments. For instance, tailoring HEAs’ dynamic mechanical
traits for aerospace or high-velocity transport applications can improve their structural
durability and operational lifespan.
The dynamic mechanical behavior of HEA coatings at high and low temperatures is
even more complex. On one hand, the interface between the coating and substrate is prone
to delamination or cracking under high-speed impact or abrupt temperature changes. On
the other hand, the strength, ductility, and toughness of the coating itself may undergo
irreversible changes under the combined effects of dynamic loading and thermal shock.
Consequently, it is necessary to integrate advanced experimental methods with multiscale
numerical simulations to conduct in-depth studies of the behavior of HEA coatings under
high strain rates and in extreme temperature environments, thereby guiding the design,
fabrication, and application of these coatings.


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Author Contributions: Conceptualization, Y.L; methodology, Y.L.; validation, R.X.; formal analysis,
R.X.; investigation, Y.L. and R.X.; resources, Y.L.; data curation, Y.L.; writing—original draft preparation, R.X.; writing—review and editing, R.X. All authors have read and agreed to the published
version of the manuscript.
Funding: This study was supported by North China University of Water Resources and Electric
Power Doctoral Research Funding (Grant Nos. 201903010). 2024 Anhui Provincial University Scientific Research Project (Natural Science Category, Key Project, No. 2024AH052003); 2024 Provincial
Department of Education Science and Engineering Teachers’ Internship Program in Enterprises
(No. 2024jsqygz76).
Conflicts of Interest: The authors declare no conflicts of interest.
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