Materials Science for Energy Technologies 7 (2024) 35–60
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Materials Science for Energy Technologies
CHINESE ROOTS
GLOBAL IMPACT
journal homepage: www.keaipublishing.com/en/journals/materials-science-for-energy-technologies
An overview of microstructure, mechanical properties and processing of
high entropy alloys and its future perspectives in aeroengine applications
Tushar Sonar a,⇑, Mikhail Ivanov a, Evgeny Trofimov b, Aleksandr Tingaev a, Ilsiya Suleymanova a
a
Department of Welding Engineering, Institution of Engineering and Technology, South Ural State University (National Research University), Chelyabinsk 454080, Russia
Department of Materials Science, Physical and Chemical Properties of Materials, Institution of Engineering and Technology, South Ural State University (National Research
University), Chelyabinsk 454080, Russia
b
A R T I C L E
I N F O
Keywords:
High entropy alloys
Microstructure
Mechanical properties
Processing
Heat treatment
Aeroengine
A B S T R A C T
Modern engineering applications continually strive to develop greater performance mechanical components
with good microstructural stability, improved mechanical properties, corrosion resistance and decreased cost
of repairing and maintenance. This necessitates the broad use of advanced high performance materials like
high entropy alloys (HEAs). These alloys are created by combining five or more alloying elements in equal
or substantial amount. About 5 to 35 at. % of the alloying element is present in HEAs. It is characterized primarily by greater entropy, slow diffusion, severe lattice distortion, and cocktail effects. Due to its advanced
microstructural stability throughout a larger temperature span and for longer length of time, it demonstrates
improved mechanical characteristics at ambient temperature, cryogenic temperature, and elevated temperature. The diversity of elemental contents and significantly higher mixing entropy of HEAs make them mechanically superior to classic metals and alloys. It also shows better strength to weight ratio. Hence, it qualifies as a
possible structural and functional material for aeroengine applications. In this work, the studies on the HEAs
are briefly reviewed. A basic explanation of the four core effects of HEAs is given. Discussion is held on
microstructure and mechanical properties of HEAs. The processing routes for manufacturing of HEAs (arc melting, bridgman solidification, mechanical alloying and vapour deposition) are presented briefly. The influence
of heat treatment on mechanical behavior and microstructure of HEAs is presented. The simulation approach of
CALPHAD modeling for designing of HEAs is discussed briefly. The future scope for research and development
of HEAs in aeroengine applications is briefed.
1. Introduction
Modern engineering applications continually strive to develop
greater performance mechanical components with good microstructural stability, improved mechanical properties, corrosion resistance
and decreased cost of repairing and maintenance. This necessitates
the broad use of advanced high‐performance materials like high
entropy alloys (HEAs). These alloys are created by combining five or
more elements for alloying in equal or substantial amount. About 5
to 35 at. % of the alloying element is present in HEAs. It is characterized primarily by greater entropy, slow diffusion, severe lattice distortion, and cocktail effects [1–5]. The diversity of elemental contents
and significantly higher mixing entropy of HEAs make them mechanically superior to classic metals and alloys. Due to its advanced
microstructural stability throughout a larger temperature and for
longer length of time, it demonstrates improved mechanical capabilities at ambient temperature, cryogenic temperature, and elevated temperature. It shows outstanding resistance to erosion and corrosion,
great strength and hardness, better strength to weight ratio and good
ductility [6,7]. Due to its excellent balance of exceptional fractured
toughness and yielding strength, as demonstrated in Fig. 1, HEAs have
a greater threshold to breakage. Clearly, HEAs have a higher fracture
toughness than the entirety of metals and alloys, but their yield
strengths are comparable to structural ceramics and certain bulk
metallic glasses [8]. Moreover, HEAs have been demonstrated to be
oxidation‐ resistant at high temperatures, expanding their potential
technological applications as potential substitutes for Nickel alloys in
turbine systems [9]. Hence, HEAs have grown to be the most practical
choices for fulfilling the requirements for difficult loading condition,
particularly in the fields of aerospace and power generation sectors
⇑ Corresponding author.
E-mail addresses: tushar.sonar77@gmail.com, sonart@susu.ru (T. Sonar), ivanovma@susu.ru (M. Ivanov), trofimovea@susu.ru (E. Trofimov), tingaevak@susu.ru (A. Tingaev),
suleimanovaii@susu.ru (I. Suleymanova).
https://doi.org/10.1016/j.mset.2023.07.004
Received 4 May 2023; Revised 5 July 2023; Accepted 5 July 2023
Available online 16 July 2023
2589-2991/© 2024 The Authors. Publishing services by Elsevier B.V. on behalf of KeAi Communications Co. Ltd.
This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).
T. Sonar et al.
Materials Science for Energy Technologies 7 (2024) 35–60
[10]. For structural, aeronautical, nuclear, and many other engineering fields, it is therefore a viable engineering material. It discovers utility for discharge nozzles, combustion chambers, and jet engine
turbines.
During 2004, Cantor [11] and Yeh's team of investigators [12–16]
published the very first study on HEAs. HEAs are dependent on the
novel configurational entropy maximization alloy design concepts,
which offer great versatility for developing diverse materials
[17–22]. In contrast with present‐day physical‐metallurgy systems,
the advanced mixing entropy in HEAs favors the creation of solid‐
solution phases such as body‐centered cubic (BCC), face‐centered
cubic (FCC), hexagonal closed packing (HCP), and multiphases rather
than intermetallic compounds (IMCs) and perhaps other intricate
phases [23,24]. Melting by arc, Bridgman method of solidification,
atomization, cladding by laser, manufacturing by additive technology,
alloying by mechanical systems, powder metallurgical route, and
incorporating vapor state elements using deposition of vapor phase
and atomic layers, and sputtering are some of the production processes
that may be utilized to create HEAS [25–30]. Subsequently, investigations have taken place on a variety of HEAs, notably FeNiMnCoCr
[31], AlCoCrFeNi2.1 [32], CoCrFeNi Al0.6 [33], CrCoFeNiNb [34],
CoCrFeNiCu [35], AlCrTiVZr [23], and CrTiNbZrAlx [36], among
others. In this work, the studies on the HEAs (typically CoCrFeNiMn
and AlxCoCrFeNi HEAs) are briefly reviewed. A basic explanation of
the four core effects of HEAs is given. Discussion is held on microstructure and mechanical properties of HEAs. The processing routes for
manufacturing of HEAs (arc melting, bridgman solidification, mechanical alloying and vapour deposition) are presented briefly. The influence of heat treatment on mechanical behavior and microstructure
of HEAs is presented. The simulation approach of CALPHAD modeling
for designing of HEAs is discussed briefly. The future scope for
research and development of HEAs in aeroengine applications is
briefed.
2. Definition of HEAs
Since there are different interpretations of HEAs dependent on
composition and entropy, it is unclear whether multi‐elemental alloys
could be categorized as HEAs.
2.1. Composition based definition
According to compositional definitions, HEAs are alloys with an
atomic percent range of 5 to 35% with at least five major elements
[37,38]. HEAs don't necessarily be equimolar or substantially equimolar in order to equalize various material properties like toughness, oxidation, ductility, strength, creep etc[39,40]. Even small elements are
conceivable in them. The atomic proportion of any minor element present in HEAs is <5% [41,42]. This composition‐based definition alone
specifies the constituent concentrations; it places no restrictions on the
entropy's magnitude. Therefore, this idea does not need the existence
of a single‐phase SS.
2.2. Entropy based definition
The degree of configurational entropy of mixing of such alloys
determines the concept of HEAs in accordance with entropy. Hence,
it distinguishes between high EAs (SSS,ideal greater than 1.61R) with
more than 05 primary elements of alloying, medium EAs
(0.69R < SSS,ideal < 1.61R) with two to 04 primary elements of alloying, and low EAs (SSS,ideal < 0.69R) with classic alloys. Here, R is the
gas constant, and SSS,ideal is the overall configurational molar entropy
in an idealized solid solution [43]. The words multi‐component alloys
(MCAs), multi‐principal element alloys (MPEAs), concentrated solid
solution alloys (CSSs) [44], complex concentrated alloys (CCAs)
[45], and metal buffets (MBs) [46] are often used to designate to these
alloys.
Fig. 1. Comparison of fracture toughness and yielding strength of engineering materials with HEAs [8].
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solid solution, the greater entropy impact outweighed other factors
like atomic radii and packing density [51]. Nonetheless, it ought to
be outlined that the AlCoCrCuFeNi alloy system illustrates that HEAsmay still respond brittlely [52,53]. Fig. 3 illustrates the strength‐
ductility trade‐off [38]; high‐strength alloys frequently have poor ductility, and inversely. Nonetheless, certain HEAs do permit circumventing the strength‐ductility trade‐off due to its diverse principal element
constitution. They encompass extraordinary mechanical properties
such excellent fracture toughness at extremely low temperatures
[8,57], elevated temp mechanical performance [54–56], or specific
stiffness [49,52], among others. Some alloy theories give superparamagnetism and other interesting characteristics [60].
3. Four-core effects of HEAs
Owing to the fact that HEAs are multi‐principal element systems,
they exhibit four main effects: greater entropy, extreme distortion of
lattice, slow diffusion and cocktail effects. It sets HEAs apart from classic alloys. Fig. 2 [3] offers a illustration of key effects in HEAs.
3.1. Greater entropy
The Greater entropy effect constitutes the initial notable HEA
effect. The evolution of solution phases may be advanced because of
this, and the microstructure can become much relatively simple than
initially thought. Due to solution hardening, this aspect may increase
the strength and ductility of solution phases. The term “entropy” refers
mostly to configurational entropy linked to Gibb's free energy. The
“greater entropy effect” prevents the emergence of brittle intermetallic
with extremely ordered structures. Greater configurational entropy is
correlated with it [47]. The phase would stay stable because, in accordance with the Greater entropy theory, the entropy would be much
significant and free energy would subsequently decrease. With the
simultaneous enhancement in the solid solubility of alloying constituents, the amount of phases in the HEAs may also be shown to
be diminishing. It may be understood using Gibbs free energy formulation, G = H‐TS, that expresses Gibbs free energy and the system's
temperature, enthalpy, and entropy in sequence. If the temperature
is elevated enough, entropy would prevail, and a phase would become
prominent at greater entropy [48,49]. A metal has a greater entropy
when it's melted compared to when it remains solid, and Richard's rule
asserts that the distinction between both states is relatively similar to
the gas constant R. The establishment of simple solid solutions is stimulated by the greater entropy effect in HEAs, which causes a substantial reduction in the total number of phases compared to the maximum
number (i.e., n + 1, where n is the number of constituents) predicted
by phase rule of Gibbs'. Therefore, the microstructure is simpler than
initially thought, increasing the possibility that it may display superior
properties [50]. Earlier, it was argued that the ability to stabilize a
3.2. Extreme lattice distortion
It was once believed that the different elements making up lattice
of crystal with different atomic radii were the cause of the lattice
deformation effect because they might induce local atom deflections
related to their locations in diluted alloys. In contrast to ordinary
alloys, it was thought that this would result in enhanced solid solution
hardening. One may determine the level of deformation in each material by looking at the comparative hard‐sphere models of the lattice for
metals, common dilute alloys, and HEAs. A pure metal has the identical type of atoms filling the lattice places, as diagrammatically illustrated in Fig. 4 [61].
Consequently, the lattice is not distorted because of the atomic
locations. In typical dilute alloys, the introduction of a little amount
of a secondary atomic species causes a slight lattice deformation.
Extreme deformation occurs in HEAs because multiple atoms of differing sizes are compelled to cohabit in lattice at random. The massive
lattice deformation in HEAs has supposedly been detected using neutron and X‐ray diffractions [27,38,43,58]. On macroscopic level, the
lattice distortions are assumed to be responsible for substantial
strengthening and lower thermal and electric conductivities because
of elevated electron and phonon dispersion [59,60]. Moreover, the
extreme lattice distortions are ascribed for the superior properties of
Fig. 2. Schematics of 4 core effects in HEAs [3].
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Materials Science for Energy Technologies 7 (2024) 35–60
Fig. 3. Comparing mechanical properties of selected HEAs to alloys of Al and Mg, industry-related steel grades, and metastable MEAs at room temperature [61].
Fig. 4. Lattice distortion is depicted diagrammatically for BCC pure metals, classic dilute alloys, and HEAs. The letters A through E, generally, stand for various
element species [38].
decreased chance of escaping when it approaches a location with
lower energy levels as a consequence. On the flip hand, a jumping site
with far more energy has a higher chance of returning to its initial
location. Each constituent part diffuses into the other characteristic
at a different pace. Some elements have great levels of activity whereas
others exhibit significantly lower levels, dependent on their characteristics [66]. A wider span of atomic configurations could encircle vacancies, that also decelerates elevated‐temperature diffusion and
diffusion‐controlled characteristic like oxidation [55], creep [54],
phase transitions [68], and particulate growth in HEAs compared with
classic alloys [69]. This is because of the consequences of sluggish diffusion [20,67]. This is owing to negative effects of sluggish diffusion
rate [20,67]. The kinetics of diffusion are delayed by such oscillations,
which raises the energy barrier for diffusion. Yeh [15] compared diffusion coefficients of elemental contents in metals, stainless‐steels, and
HEAs to determine how vacancies developed and how the constitution
of HEAs was partitioned. He found that metals, stainless‐steels, and
HEAs have the greatest rate of diffusion in each of the 3 alloy systems:
metals > stainless‐steels > HEAs. Several different sorts of atoms
ought to migrate collaboratively in to accomplish the compositional
partitioning across phases throughout phase shifts in HEAs. The pres-
HEAs, specifically the BCC‐structured HEAs [62,63]. The extreme
lattice‐distortion effect is also coupled to the retarded kinetics and tensile fragility of HEAs [16,64]. The investigators also highlighted the
comparatively inferior strength of 01‐phase FCC‐structured HEAs,
which unquestionably cannot be explained by the extreme lattice distortions idea [65]. The lattice distortions of HEAs have to be quantified
by theoretical study.
3.3. Sluggish diffusion
The diffusion in HEA was disclosed to be sluggish than in classic
alloys. Diffusion is becoming a crucial factor when elevated‐
temperature strength or elevated‐temperature microstructural stability
are being considered. The HEAs have a sluggish rate of diffusion,
which would be indicative of a delayed kinetics or phase transition.
The sluggish diffusion in the context of HEAs varies markedly from
that of typical alloy system because of the varied configurations of
the nearby atoms [47]. The atom introduces an additional neighbour
to the sites of lattice by leaping into a vacancy. The local variations
in arrangement of atoms give rise to different bonds and local energy
to every position in the atoms. An atom remains locked and has
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defects, which are significant microstructural features. Key features
of microscopic or macroscopic imperfections include pores, compositional partitioning, fractures, and residual stresses. Grain boundaries,
vacancies, and dislocations are examples of atomic‐level imperfections. All of these must be noted in order to understand mechanical
characteristics. HEAs have been demonstrated to exhibit high compressive strength and hardness at both room temperature and elevated
temperatures owing to their microstructures, [15,66,75–77]. Greater
tensile integration has been reported for HEAs, comprising both
advanced strength and enough ductility [75,78,79]. In contrast to
FCC‐structured HEAs, which exhibit poor strength and high plasticity,
BCC‐structured HEAs reportedly have great strength and low plasticity. As a consequence, the structural categories play a crucial role in
determining how robust or challenging HEAs are.
ence of vacancies for diffusion by substitution is restricted in HEAs,
just like it is in standard alloy systems, because every vacancy in crystalline HEAs is linked to a + enthalpy of development and advanced
mixing entropy. The struggle between these 2 elements results in a certain steady state of vacancies with lowest free energy of mixing at a
defined temperature [70]. Many elements atom actually encircle and
compete for vacancy in matrix that constitutes the entire solution
throughout diffusion. An atom or vacancy would proceed more slowly
and have a higher activation energy across a fluctuating diffusion
channel [71].
3.4. Cocktail effect
The cocktail effect for metals was first mentioned by Ranganathan
[72], and it has subsequently been confirmed in terms of mechanical
and physical properties. His paper asserts that an alloy's properties
are not only governed by the characteristics of its component elements. Contrarily, the alloy exhibits composite characteristics as a consequence of the relationships between each of its constituents. The
composite characteristics are not subject to the unanticipated mixing
rules. The characteristics are extensive, ranging from atomic scale to
micro‐scale multiphase composite phenomenon [47]. This phenomenon is present in typical alloys as well, although it's more pronounced in HEAs since at minimum 5 essential elements have been
incorporated to advance the properties of materials. Dependent on
their constitution, HEAs can have a single phase, two phases, three
phases, or more. As a consequence, the characteristics of each phase
as well as impacts of morphologies, the distribution of grain sizes,
grain and phase boundaries, and each grain's integrity affect the overall characteristics of the investigated material. Yet, each phase is a
solid mixture comprised of several primary components and may be
compared to an atomic‐scale composite. In conjunction to the basic
characteristics of the components as specified by the mixture rule, it
also exhibits considerable deformation and interrelations between all
of its solutes. The interaction and lattice distortions would add extra
quantities to mixing rule's predicted variables. Scales ranging from
micro‐scaled multi‐phase composite influence to atomic‐scaled
multiprincipal‐element composite effects are covered by term “cocktail
effect” [50]. Consequently, it is essential for alloying designer to comprehend the associated factors involved prior to actually selecting an
acceptable mixture and processes depending on cocktail effect. The
cocktail effect was illustrated by Wu et al. [73] in their finding’s presentation from 2006. The experiment proves that the aluminum inclusion helps AlxCoCrCuFeNi HEA transition from FCC to BCC. Al
increases the HEA's hardness while simultaneously being soft and having lowered melting point. Senkov et al. [54,74] evaluated NbMoTaWV and VNbMoTaW, refractory HEAs with melting temperatures
somewhere around 2600 °C, in accordance with the rule of combined
effects. While the alloy has a superior softening resistance relative to
superalloys, its yield strength is significant, approaching 400 MPa at
1600 °C. The refractory alloy has a great deal of promise in applications requiring elevated temperatures. In addition to widely recognized four primary effects, typical alloys and HEAs have a number of
significant differences.
4.1. Tensile properties
The HEAs disclosed good yielding strength and ductility at ambient
temperature, according to Fig. 5 [80], that compiles all tensile data
[53,75,81–89] for the HEAs. It follows the same trend of greater
strength‐lower ductility and vice‐versa as in classic alloys. Compared
to Ti‐6Al‐4 V and Inconel 713, AlCoCrCuFeNi and Al0.5CoCrCuFeNi
have greater yielding strength. The CoCrFeMnNi alloy is more ductile
when comparing to prominent alloys like 304 stainless steel, Inconel
713, Ti‐6Al‐4 V, and 5083 Al. The Al0.5CoCrCuFeNi HEA disclosed
greater yield strength and ductility compared to 304 stainless steel.
The impact of temperature on the mechanical performance of Al0.5CoCrCuFeNi [75] and CoCrFeMnNi [90] systems is analyzed. When
temperature goes up, the yield strength decreases (Fig. 6) [80]. The
Al0.5CoCrCuFeNi and AlCoCrCuFeNi HEAs showed significant reduction in yield strength at increased levels of temperature compared to
CoCrFeMnNi HEA. The CoCrFeMnNi and CoCrFeNi alloys show substantial temperature‐dependent strength loss from 77 to 1,273 K and
a slight strain‐rate reliance at 10−3 and 10−1 per second, accordingly
(Fig. 7) [85].
The route of synthesis has substantial impact on ductility of HEAs
in additional to test conditions (temperatures and strain rates). The
AlCoCrCuFeNi alloy was much more ductile and harder following
hot working compared to the as‐cast [75,82]. This material's hardening and toughening capabilities are principally owing to its extremely
small grain size (1.5 µm). Between 1073 and 1273 K, the AlCoCrCuFeNi alloy (hot‐worked) exhibited a very remarkable superplastic
behavior. The ductility advanced by 400% and reached 860% at
1273 K [79] (Fig. 8). Regarding HEAs with room temperature mechanical characteristics, the yield strength can extend from 300 MPa for
FCC alloys, such as CoCrFeNiCuTix system, to around 3000 MPa for
BCC alloys, like AlCoFeCrNiTix system [91,92]. HEAs may additionally
incorporate modest concentrations of alloying elements to modify
their strength, ductility, hardness, etc., much like other classic alloys.
The body‐centered‐cubic (BCC) solid solution phase and the Laves
phase of (CoCr)Nb type were found to exist in the generated
AlCoCrFeNbxNi HEAs by Ma and Zhang [93] in their assessment of
Nb alloying effect (Fig. 9a). The micrographs of the alloy series extend
between hypoeutectic and hypereutectic, and the Vickers hardness and
compressive yield strength increase almost in line with rising Nb content (Fig. 9b).
Zhou et al. [91] examined the effect of Ti alloying on AlCoCrFeNiTix using the equiatomic ratio and greater mixing entropy. The bulk of
the alloy system, which exhibits remarkable compressive mechanical
properties at room temperature, is composed of the BCC solid solution.
In comparison to other high strength alloys, particularly bulk metallic
glasses (BMGs), the AlCoFeCrNiTi0.5 alloy especially had greater yield
strength, fracture strength, and plastic strain of 2.26 GPa, 3.14 GPa,
and 23.3%, correspondingly (Fig. 10). The HEA approach was implemented to create new refractory alloys that contain several key alloying elements at almost equiatomic percentages using innovative
4. Mechanical properties of HEAs
Mechanical characteristics are significantly influenced by both
microstructure and composition. Dislocation behaviors are regulated
by the elastic characteristics and atom‐to‐atom interactions defined
by constitution. The constitution also determines the volume percentages of phases, and intrinsic attributes of these phases affect other
parameters. Even with specific composition and phase content, characteristics can be drastically altered by changing the size, shape, and distribution of phases. Mechanical properties are greatly modified by
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Fig. 5. Most of the investigated HEAs' tensile yield stress vs ductility at room temperature compared to regularly used alloys including 304 stainless steel, Inconel
713, 5083 aluminum and Ti-6Al-4 V alloy [80].
Fig. 6. Temperature-dependent yield stress for three types of HEAs [80].
298 and 77 K are intergranular and transgranular, correspondingly,
the fracture strains only minimally vary. This shows that the DBTT
of the AlCoCrFeNi BCC HEA is lower below −196 °C. (Fig. 12).
metallic materials with higher melting temperatures, such as refractory molybdenum (Mo) and niobium (Nb) alloys [53,76,94,95]. The
compressive characteristics of AlCoFeCrNi‐based single phase BCC
HEA was studied by Qiao et al. [96]. The AlCoCrFeNi HEA does not
readily transition from becoming ductile to brittle even if the temperature is lowered to 77 K. When comparing to the compressing properties at various temperatures displayed in Fig. 11, the yield and fracture
strengths of AlCoFeCrNi HEA was advanced by 29.7 and 19.9% as the
temperatures decline from 298 to 77 K. While the fracture patterns at
4.2. Hardness
HEAs exhibited hardness ranging from 140 to 900 HV which is
dependent on the alloy systems and preparation methods
[15,41,53,54,75–78,81,83,94,97–110]. The hardness values of the
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Fig. 7. Effect of temperature on yield strength of CrMnFeCoNi and CrFeCoNi HEAs at engineering strain rates of a) 10−3 per s and b) 10−1 per s.
Fig. 8. Tensile curves of a) as-cast and b) hot forged AlCuCrFeNiCo HEAs deformed at different temperature levels and 10−3 s−1 initial strain rate [79].
Fig. 9. A) change in the initial phase composition caused by the addition of nb elements causes the formation of the ordered laves phase in addition to the solid
solution phase; b) compressive stress-strain curves for samples of AlCoCrFeNiNbx rods with a diameter of 5 mm for × = 0, 0.1, 0.25, and 0.5 [93].
MoNbTaVW, Al0.4Hf, MoNbTaW, and HfNbTaTiZrTi have corresponding hardness of 591 HV, 535 HV, 500 HV, 454 HV, and 390
HV. Defining the alloy system, modifying the constituent ratio inside
the alloy, and selecting an alloy processing method are essential steps
in determining the hardness of HEAs [80].
20 HEAs that have been studied the most are contrasted with those of
classic alloys in Fig. 12 [80]. The level of hardness varies greatly
between every alloy system. Hardness of AlCoCrCuFeNi ranges from
154 to 658 HV. It significantly influenced by the specific elemental
composition, manufacturing processes, and subsequent thermal operations. AlCrFeMnNi and AlCrFeMoNi are two alloy systems that have
tougher properties than classic alloys. Yet the as‐cast HEAs which
often include FCC phase (i.e., CoCrFeNiMn CoCrCuFeNi and CoCrFeNi) have low hardness values at room temperature [95], although
those that include a great deal of Al and Ti had greater hardness
because stronger second phases develop in those alloys. Typical hardness for refractory HEAs, especially those who have BCC phase, were
considerably large; for instance, the alloys AlMo0.5NbTa0.5TiZr,
4.3. Fatigue properties
Several HEA applications, like components of aviation engines,
often endure cyclic stress. In terms of seeking for applicability in the
aerospace industry or other fields, the fatigue behavior and lifetime
projection are two among the most crucial elements that must be
researched and analyzed but are scarcely documented [17]. Hemphill
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Materials Science for Energy Technologies 7 (2024) 35–60
respect to their UTS, the fatigue ratios are being used, as shown in
Fig. 14b. This explains relationship between UTS and the fatigue
endurance limit. It's possible that HEAs' excellent tensile strength contributes to their improved fatigue strengths. It implies that when UTS
increases, endurance limit would increase in step with it and grow linearly as well, typically equivalent to 0.5 for the majority of materials
[111]. According to this approach, HEAs even surpass it, as seen in
Fig. 14b, where the limit is 0.703. The excellent fatigue durability
and prospects for an extended fatigue life of HEAs, even at loads that
are approaching to ultimate stress, are strongly supported by these
data. Due to the dearth of publications on the fatigue behavior of
HEAs, current study should concentrate on the data sets that show
an exceptionally long fatigue life.
If such crucial information on fatigue life can be found and a prediction model for fatigue specimens can be established, the future of
HEAs in diverse applications for sections in fatigue contexts is bright
[17]. Fig. 15 [111] shows the fractographs of fatigue specimen at
900 MPa stress range following 555, 235 cycles. Typically, imperfections at the sample's corner or surface, in which the stress is greatest,
are the source of cracks. According to Fig. 15a and b, the initiation
sites have been observed at microcracks. It is possible to see the characteristic fatigue zones, such as the initiation, extension, and rupture
areas. Red arrows serve as markers to indicate certain areas. It is possible to find distinctive markings inside the fatigue propagating zone
that reveal the pattern of fracture development, which is frequently
aligned to the striation in Fig. 15c. The dimple‐like pattern of fracture
surface in Fig. 15d demonstrates ductile fracture.
Fig. 10. AlCoCrFeNiTix HEA with 5 mm diameter: compressive stress-strain
curves [91].
et al. [111] compared the results of their study on the fatigue behavior
of the Al0.5CoCrCuFeNi HEA with that of many classic alloys, such as
steels, titanium alloys, and advanced BMGs. The fatigue characteristics
of the Al0.5CoCrCuFeNi HEA are compared to those of various classic
alloys and BMGs in Fig. 13, which displays a typical stress range vs.
the number of cycles to failure (S‐N) curve. The lower limit of the fatigue ratios of HEAs is well compared to that of steels, titanium, and
nickel alloys, in addition to various Zr‐based BMGs and zirconium
alloys. Moreover, certain materials, such as ultra‐high strength steels
and wrought aluminum alloys, exhibit lower fatigue ratios regardless
of their greater tensile strengths because of their brittle character.
The strong group of HEAs often surpasses these materials by displaying
a greater fatigue ratio over materials with equal tensile strengths
because of the reduced fault density. Given that HEAs' upper fatigue
limit is significantly higher in comparison to BMGs and other classic
alloys, it is possible that such materials may be outperformed in load
– bearing applications by HEAs with improved fabrication and processing [17].
The fatigue‐endurance limits of the Al0.5CoCrCuFeNi HEA are
depicted in Fig. 14a as a measure of UTS. One element that contributes
to the exceptional fatigue strength of HEAs is their greater tensile
strength. It is clear that when the UTS rises, the fatigue‐endurance
limit—which for the bulk of materials is approximately equivalent to
0.5—increases linearly [111]. HEAs have an upper limit of 0.703,
which is comparable to and even above this pattern. To more properly
assess the fatigue performance of HEAs compared to other materials in
4.4. Wear behaviour
Chuang et al. [41] reported that Co1.5CrFeNi1.5Ti and Al0.2Co1.5CrFeNi1.5Ti alloys exhibit higher wear resistance than standard wear‐
resistant steels of equivalent hardness, (Fig. 16). Because when molar
ratio of Ti grows, a discernible increase in the volume percentage of
precipitates occurs. According to Chuang [41], the inclusion of Al
leads to the development of the needle‐like phase with Widmanstätten
features in the Al‐rich interdendritic (ID) portions and a decrease in
the amount of ƞ coarse phase. This phenomenon originates from the
difficulty of transferring Al effectively. The hardness of the alloys
Co1.5CrFeNi1.5Ti and Al0.2Co1.5 CrFeNi1.5Ti are HV 654 and HV 717,
correspondingly, owing to the strengthening effect of the hard phase.
It is proposed that the remarkable anti‐oxidation property and thermal
softening durability of these HEAs are responsible for their superior
wear resistance. Alloying has the potential to affect the wear behavior
of HEAs. Wu et al. [73] examined the adhesive wear behavior of AlxCoCrCuFeNi HEAs as a proportion of Al content. He found that the
smoothness of worn surface and generation of tiny debris with a high
oxygen content are a consequence of higher Al concentration which
significantly increases wear resistance (Fig. 17).
Fig. 11. The AlCoCrFeNi HEA's compressive true stress-strain curves at (a) 298 and (b) 77 K. As relative to the similar mechanical characteristics at room
temperature, the yield strengths and fracture strengths at cryogenic temperatures enhanced noticeably [96].
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Fig. 12. Comparing 20 most studied HEAs' hardness to that of prominent alloys made of Al, Cu, Co, Fe, Cr, Ni, V and Ti [80].
Fig. 13. Graphs showing comparison of the Al0.5CoCrCuFeNi HEA's fatigue ratios and fatigue endurance limits as variables of UTS of structural materials and
BMGs [111].
Fig. 14. A) fatigue endurance limits and b) fatigue ratios of al0.5CoCrCuFeNi HEA as a function of ultimate tensile strength of other engineering materials and
BMGs [111].
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Fig. 15. A) sem image of fatigue tested fractured surfaces after 5,55,235 cycles at 900 MPa; b) microcracks noticed at surface prior fatigue testing; c) fatigue
striations in the zone of crack-propagation; d) dimples in the region of ultimate fracture [111].
Fig. 16. A) tests on samples of al00ti05 (co1.5CrFeNi1.5Ti0.5), Al02Ti05 (Al0.2Co1.5CrFeNi1.5Ti0.5), Al00Ti10 (Co1.5CrFeNi1.5Ti), Al02Ti10 (Al0.2Co1.5CrFeNi1.5Ti1.5Ti), SUJ2 (AISI 52100), and SKH51 (AISI) were conducted. The data is shown as weight gain following a test per unit area; b) Plots of the hot hardness of
Al00Ti10, Al02Ti10, SUJ2 and SKH51 specimens from ambient temperature to 900 °C.
there are increased chances of element segregation in multicomponent
alloys. Grain growth can result in planar, cellular, or dendritic morphologies depending on the growth circumstances. According to
XRD data, the CoCrFeMnNi alloy generates a uniform FCC solid solution, and EDX reveals that the solid solution's CoCrFeMnNi composition is basically homogeneous [8,11]. The BSE image shows that
annealing process produced completely recrystallized, equiaxed grains
that were 6 µm in size. 2 µm diameter uniformly distributed Cr‐ and
Mn‐rich particles were detected in solid‐solution (Fig. 18).
The tendency of HEA to form simple solid‐solution was supported
by a previous investigation on AlxCoCrCuFeNi alloys [15,77] with a
range of Al concentrations (molar ratios from × = 0 to 3.0). The
FCC and BCC structures were differentiated using classic XRD
(Fig. 19). It was discovered that the FCC structure changes into a
BCC structure when the Al concentration was advanced. As the Al content rises, the microstructures of the as‐cast AlxCoCrCuFeNi alloys
5. Microstructure of HEAs
The FCC structure occurs the most commonly among all single‐
phase HEAs recorded, closely followed by the BCC structure, whereas
reports of the HCP structure are far less common [51]. The compounds
of intermetallics that have been observed in multi‐phase HEAs include
B2 (ClCs, AlNi, cP2), C14 (Fe2Ti, hexagonal Laves phase, 2hP12,
MgZn), A5 (β‐Sn, tI4), C16 (tI12, Al2Cu), A9 (graphite, hP4), A12
(α‐Mn, cI58,), C15 (cubic Laves phase, Cu2Mg, cF24), D8b (σ‐CrFe,
tP30), D022 (tI8, Al3Ti), D02 (Li2MgSn, BiF3, cF16), D024 (hP16,
Ni3Ti), DO11 (oP16, Ni3Si), D2b (AlFe3Zr, Mn12Th, tI26), D85
(FeMo, Fe7W6, Co‐Mo, hR13), L12 (AuCu3, cP4), L21 (AlCu2Mn,
Heusler, cF16) and L10 (AuCu, tP2). The effects of these intermetallic
compounds on the properties of HEAs might be either favorable or
unfavourable [10]. Due to the disparity in melting temperatures, densities, and other physical characteristics of the constituent elements,
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HEAs is significantly influenced by the rates of solidification and
cooling.
6. Processing of HEAs
6.1. Fabrication routes
The approach that has been most often documented in literature is
the arc‐melting process. The concentrated solid solution structures and
distinctive mechanical properties of the HEAs with multiprincipal
alloying elements are readily apparent. As a result of the frequent presence of a dendritic microstructure in the as‐cast condition, annealing
at higher temperatures is necessary to produce equiaxed grains and
get rid of chemical inhomogeneity. The single‐crystalline HEAs with
excellent tensile plasticity are effectively synthesized using the Bridgman solidification method. The laser cladding HEA coatings exhibit
highly fine microstructures and improved properties, such as high‐
temperature softening resistance, corrosion resistance, and oxidation
resistance, as a result of the substantially greater cooling rates utilized.
High‐entropy alloys may be made using a variety of techniques,
including mechanical alloying of elemental powders. The as‐milled
grain structures of powder are generally nanocrystalline. Depending
on the specific alloy, annealing the as‐milled powders may or may
not cause changes in the as‐milled crystal structure. When using
mechanical alloying to examine the difference between amorphous
and crystalline solid solutions, care must be taken because extended
milling may result in the introduction of enough free energy from
the defects at the grain boundaries to transform a crystalline structure
into an amorphous phase. A number of materials may be effectively
prepared into thin films using physical vapor deposition. Based on
its simple accessibility, magnetron sputtering deposition has been successfully used to create multicomponent HEA thin films with possibly
advantageous features. While molecular beam epitaxy shows potential
for the production of single‐crystalline HEA films, atomic layer depo-
Fig. 17. Vickers hardness and wear coefficients of AlxCoCrCuFeNi alloys with
varied Al contents [73].
become increasingly complicated. Although the partitioning of Cu in
the small volume section of the interdendrites, the EDS data showed
that the dendrites were composed of many primary components, particularly at low Al concentrations (x 0.8). When the Al concentration
increases to × = 1/41, the dendrite develops a spinodal structure
(a periodic architecture brought on by elemental percentage fluctuations) (Fig. 20). The alloy phases for minimum aluminium% are FCC
(a Cu‐enriched FCC phase in interdendritic area and a multi‐
principal element FCC phase in dendrites), while for increased Al content, the alloy phases were mixed‐phase FCC/BCC/B2. This would be
determined by the relationship between both the SEM microstructural
evolution and XRD tests. The microstructural development of as‐cast
Fig. 18. Microstructure of annealed CoCrFeMnNi alloy: backscattered electron (BSE) micrograph, EDS and XRD results [11].
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Fig. 19. AlxCoCrCuFeNi alloy with varied Al percentages, according to XRD analysis (x values). The designs are indexed using the FCC (open triangles) and BCC
(solid diamonds) phases. The ordered peaks of the B2 phase are also displayed [77].
sition is more easily able to achieve film growth with regulated atomic
layers than sputtering deposition [112]. The processing routes for
HEAs are shown in Fig. 21.
The most popular method for creating HEAs [17] is called arc melting, which involves completely combining all of the component elements in a liquid state before solidifying them in a copper crucible.
To guarantee the chemical homogeneity of the alloys, several recurrent melting and solidification processes are frequently used. In the
copper crucible with a bowl‐like form, the solidified ingots have a
button‐like shape. Another option is to use the pre‐machined hole in
the copper crucible's bottom to allow the melt to flow into a copper
mold. This process makes it possible to produce cylindrical ingots with
a faster cooling rate. Additionally, it is simple to prepare the samples
into specimens for tensile/compressive testing. A number of HEAs
with distinctive characteristics have been discovered using this casting
technique. Zhou et al. [91] used this method to create AlCoCrFeNiTix
(x in molar ratio) HEAs, which displayed body‐centered cubic (BCC)
solid‐solution structures with exceptional compressive attributes such
ultrahigh fracture strength and strong work‐hardening capacity. The
issue with this method is that rapid solidification makes it difficult
to control the solidification process. As a result, alloy samples have
varying microstructure characteristics from the surface to the center,
such as an uneven distribution of as‐cast dendrites in terms of size
and morphology, ranging from fine grains to columnar dendrites,
and then to coarse equiaxed dendrites (Fig. 22). The mechanical properties of HEAs may also be negatively impacted by a number of
unavoidable as‐cast defects, including as elemental segregation, suppression of equilibrium phases, residual stresses, cracks, and porosities. There should be actions done to lessen or get rid of these
defects in the HEAs.
Singh et al. [113] analyzed the decomposition of AlCoCrCuFeNi
HEA developed using splat quenching and casting and observed that
the greater rate of cooling promoted the evolution of single phase
structure in HEA. The splat‐quenced alloy disclosed the evolution of
signle BCC phase. However, the cast alloyed showed the evolution of
one BCC and two FCC phases (Fig. 23a). The TEM micrographs showed
that the splat quenched alloy exhibited the domain like struacture with
imperfectly ordered BCC phase that were having the size of few
nanometers (Fig. 23b). The cast alloy exhibted the evolution of many
BCC and FCC phases within the dendritic regions (Fig. 23c). The interdendritic regions were observed to be enriched with Cu. HEAs made
from elements like Al‐Ni, Cu‐Ni, and Fe‐Cr that have substantial
enthalpy of mixing variations are prone to segregation during
solidification.
The Bridgman solidification casting (BSC), as opposed to conventional casting, may be utilized to control the microstructure and optimize the properties of HEAs. In particular, for rod‐shaped samples
obtained by BSC, the direction of thermal conduction and extraction
is intensely along the longitudinal direction, assuring the direction
of microstructural growth. Inductive or resistive heating is used to
melt the samples in a crucible after which the molten alloys are progressively drawn down to the liquid metal. Water is used to further
cool the outside of the liquid metal (Fig. 24) [114]. Two crucial processing parameters, namely, the temperature gradient and growth rate,
can be precisely controlled by adjusting the power of heating and withdrawal velocity, assuring the desired microstructure. Zhang et al.
[115] studied the transition of morpology of microstrucrual features
in BSCed AlCoCrFeNi HEA and observed that the BSC results in a
change from flowery dendrites to equiaxed grains, with the average
grain size being between 100 and 150 µm. This modification was
due to the higher ratio of temperature gradient (G) and growth velocity (V) in Bridgman solidification. The ratio of G/V is often low for
copper‐mould casting because G is unpredictable and V is typically
extremely high. The higher values of G/V tends to reduce the constitutional undercooling of the HEA, which makes it feasible to prevent the
development of dendrites. Fig. 25.
The heat source for laser cladding is a focused laser beam, which
keeps the substrate's heat‐affected zone relatively shallow by concen46
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Fig. 20. The SEM micrographs show as-cast AlxCoCrCuFeNi alloys with varying amounts of Al (x values): a) 0, b) 0.3, c) 0.5, d) 0.8, and e) 1, where
DR = dendrite, ID = interdendrites, and SD = spinodal decomposition. [77].
Fig. 21. Processing routes for HEAs.
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Fig. 22. Optical microstructure of arc melted and coper mould casted AlCoCrFeNi HEA cylindrical samples: a) and b) central equiaxed dendrites; c) typical as-cast
structure with fine grains viewed in the inset; d) transitional columnar dendrites [91].
Fig. 23. A) xrd pattern and b) and c) tem micrographs of splat-quenched and as-cast equiatomic alcocrcufeni hea[113].
trating on a very tiny region. This characteristic reduces the possibility
of cracking, voids, and deformation and produces a metallurgical bond
with fine microstructure and improved bond strength compared to
thermal spray. The effects of laser fast solidification on the microstructure and phase structure in HEA coatings were meticulously examined
by Zhang et al. [115–117]. Because of the different atomic radii in the
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synthesizing a range of equiatomic and non‐equiatomic HEAs beginning from blended elemental or pre‐alloyed powders. In three steps,
MA takes place. The alloy components are first mixed and processed
to a fine powder in a ball mill. The particles are subsequently compressed and sintered using a hot‐isostatic pressing (HIP) procedure.
Any residual internal tensions from any possible cold compaction are
reduced by a final heat‐treatment cycle. Alloys appropriate for high‐
heat turbine blades and other aerospace components have been successfully manufactured using this MA method. The working mechanism of MA is shown in Fig. 26.
AlCoCrCuFeNi that is equiatomic was milled by Zhang et al. [120].
The as‐milled powders had the BCC (major phase) and FCC (minor
phase) structures and were solid solutions (Fig. 27). The grain size
was incredibly small, measuring only 7 nm at the nanoscale. At
600 °C, annealing produced BCC and FCC phases. At 1000 °C, another
FCC phase emerged after annealing. This was analogous to the structures seen in AlCoCrCuFeNi that had been cast and arc‐melted.
AlCoCrCuFe and CoCrCuFeNi equiatomic alloys were developed by
Praveen et al. [121] using mechanical alloying and spark plasma sintering (SPS) at 900 °C to compact the particles. AlCoCrCuFe exhibited
predominantly BCC structure with a little FCC peak by X‐ray diffraction after mechanical alloying for 15 h. Nevertheless, CoCrCuFeNi
mostly displayed the FCC phase with very few BCC traces. In AlCoCrCuFe, the ordered BCC (B2) phase predominated, with minor quantities of the Cu‐rich FCC phase and σ phase available (Fig. 28a and b);
in CoCrCuFeNi, the FCC phases predominated, with minor amounts
of sigma phase available. After MA, simple FCC and BCC phases developed, but after SPS, Cu‐rich FCC and σ phase phases developed alongside FCC and BCC phases (Fig. 28c and d). AlCoCrCuFe and
CoCrCuFeNi were found to have hardness values of 770 HV and 400
HV, respectively. Phase development in high entropy alloys after sintering suggests that configurational entropy is insufficient to inhibit
the development of Cu‐rich FCC and phases, and that enthalpy of mixing likely plays a significant role in this process.
Physical vapor deposition (PVD), which involves the condensation
of a vaporized form of the desired film material on various workpiece
Fig. 24. Schematic representation of bridgman solidification [114].
HEA composition, which increases the solid–liquid interface energy
and makes it more difficult for atoms to diffuse over long distances
in the crystal lattice, the growth of intermetallic compounds will be
slowed down if the solidification rate is high enough, according to
their calculations of the nucleation incubation times for various competing phases [118].
A high‐energy ball mill is used in the mechanical alloying (MA)
process, which involves repeatedly cold welding, fracture, and rewelding powder particles. MA has now been demonstrated to be capable of
Fig. 25. Microstructure of Al2CoCrFeNiSi HEA coatings: a) CoCrCuFeNi [116]; b–d) SEM, EBSD, and angle distribution of grain boundaries [118].
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Fig. 26. Planetary ball milling and different stages in MA [119].
diffusion was discovered to be prevented by the Cu/AlMoNbSiTaTiVZr/Si sandwich structure up to 700 °C for 30 min. Accordingly, it
is possible for the current HEA to function as an efficient diffusion barrier for copper metallization. It is believed that even improved performance is possible with advancements in HEA composition. The effects
of substrate bias, deposition temperature, and post‐deposition annealing on the structure and characteristics of the (AlCrMoSiTi)N were
studied by Chang et al. [123].
6.2. Heat treatment of HEAs
Heat treatment can be employed to enhance the mechanical and
microstructural properties of HEAs without changing its shape or elemental constitution. Moravcik et al. [124] studied the effect of heat
treatment on microstructure and mechanical properties of
AlCoCrFeNiTi0.5 HEA. The HEA was developed through the fabrication
route of mechanical alloying. The HEA was then subjected to spark
plasma sintering (SPS) at the sintering temperature of 1100 °C, pressure of 60 MPa and holding time of 8 min. The SPSed HEA were then
heat treated at the temperatures of 1100 and 1250 °C for 2 h
respectively.
The HEA showed the evolution of Cr‐dependent supersaturated
solid solution with BCC structure. The HEA after SPS exhibited the
decomposition of solid solution into nano‐grained microstructure that
consists of solid solution of FCC and ordered BCC phases, σ phases, and
in‐situ evolved nanoparticles of TiC (Fig. 30a – c). After the HT at
1100 °C and 1250 °C, the phase composition of SPS HEAs remained
essentially the same. But after the HT, there was no longer any evidence of the σ phase (Fig. 1d and e), which was most likely caused
by its conversion into the FCC phase during the HT. Due to its extreme
brittleness, the σ phase is typically thought to be harmful to the characteristics of alloys; hence, its removal could improve alloy performance. After HT at 1100 °C, the HEA showed the evolution of FCC
grains and B2 phases with finer TiC nano‐particles dispersed in the
matrix. The TiC particles were found to be becoming coarser at HT
of 1250 °C (Fig. 30d – f). The HEA after SPS showed the hardness of
762 HV. The hardness was also lowered as a result of the HT. The
HEA showed the hardenss of 603 HV and 533 HV when HTed at
Fig. 27. XRD of solid solution of AlCoCrCuFeNi HEA made by ball milling at
60 h under different annealing temperatures [120].
surfaces, offers a range of vacuum deposition techniques used to create
thin films. It comprises sputtering, evaporative deposition, pulsed laser
deposition, cathodic arc deposition, and electron beam physical vapor
deposition. The process of fabricating HEA films via a magnetron sputtering deposition is the most common of them. Sputtered atoms that
are near the atomic fractions of the deposited HEA are expelled off
the sputtering target based on the ion or atom bombardment from
the sputtering gas (Fig. 29). The sputtered atoms are randomly distributed on the substrate, but the parameters such as the type of material source, power, base pressure, composition of atmosphere, bias
voltage of substrate, and temperature of workpiece control the nucleation, growth, and consequently microstructural evolution of the HEA
films. This method has so far been widely used in the creation of a
number of HEA films. AlMoNbSiTaTiVZr HEA layer was successfully
synthesized by Tsai et al. [122], and studied the layer's characteristics
as a diffusion barrier between silicon and copper. The Cu and Si inter‐
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Fig. 28. Micrographs of SPS pellets: a) & b) AlCoCrCuFe; c) & d) CoCrCuFeNi [121].
treatment, however distinct phase transitions took place, particularly
in the inter‐dendritic areas. The HEAs when heat treated to
650–975 °C exhibited the BCC to σ phase transformation in interdendritc regions which increases it hardness (Fig. 32a–c). The HEA on
heat treating to 1100 °C showed the σ phase to BCC structure transformation in interdenrtic region which softens the HEA (Fig. 32e and f).
The HEA was able to re‐enter the gap of miscibility and decompose to
BCC matrix with B2 precipitation in dendrite core and inter‐dendritic
areas after heat treatment at 1200 °C due to the partial dissolution
of phases and homogenization.
Wang et al. [127] developed non‐equiatomic Al0.6CoFeNiCr0.4 HEA
using arc melting. The HEAs were heat treated at 550, 650, 750 and
850 °C for 24 h. The as‐cast HEA showed dual phase structure of
FCC and BCC. With increase in temperature of ageing, the HEA’s crsystaline phase transforms from FCC + disordered BCC to FCC + disordered BCC + ordered B2 phase. Finally it transforms to
FCC + ordered B2 phase. The B2 spherical phase was found when
aging was done at 650 °C. The HEA when aged at 850 °C showed
the dissapearance of disordered BCC phase (Fig. 33a–f). The hardening
was observed when aged at 650 °C. It refers to the disordered BCC
phase growth and the precipitation of B2 spherical phases. The HEA
aged at 650 °C and 750 °C showed superior tensile properties
(Fig. 33g). It is attributed to the optimum decomposition of spinoidal
FCC (Fe‐Cr) phases and precipitation of B2 spherical phases. Further
increase in ageing temperature to 850 °C showed reduction in tensile
properties of HEAs. It is correalted to the excessive decoposition of spinoidal FCC phases reduces the mechanical properties of HEAs. The
reduction in corrosion resistance of HEAs was observed in 3.5 wt%
of NaCl solution with increase in volume fraction of B2 phases. This
infers that the crysraline phases of HEAs that contains greater % of
Al are susceptible to corrosion (Fig. 33h).
Shabani et al. [128] studied the effect of HT temperature on
microstructure and tensile properties of FeCrCuMnNi HEA. The HEA
was developed using vacuum induction casting process. It was sub-
Fig. 29. Schematic representation of sputtering deposition [112].
1100 and 1250 °C. It is attributed to the coarsening of grains in the
micrstructure of HEAs in HT condition compared to SPS. According
to Niu et al. [125], after being heated at 650 °C for 8 h, Al0.5CoCrFeNi
HEAs precipitated nanoscale B2 phase (Fig. 31a and b), that improved
the UTS and YS of HEAs. As‐cast HEA exhibited UTS, YS, and EL of
714 MPa, 355 MPa, and 41.6%, accordingly. The UTS, YS, and EL of
HEAs that had a 650 °C/8‐hour heat treatment was 1220 MPa,
834 MPa, and 25%, correspondingly (Fig. 31c).
Munitz et al. (1 2 6) studied the effect of heat treatment on
microstructural evolution and mechanical properties of AlCoCrFeNi
HEA. Arc melting was used to create the AlCoCrFeNi HEA. The HEA
showed dendritic microstructure. The interdendritic regions were
enriched with Co, Cr, Fe and Al whereas the dendritic core was
enriched with Ni. The dendrite core of as‐cased HEA was soft and
showed nano‐sized precipitates. The interdendritc regions were comparatively hard and showed larger nano‐size precipitates than the dendrite core. The dendritic morphology was unaffected by heat
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Fig. 30. MAed and SPSed AlCoCrFeNiTi0.5 HEA: a) XRD spectrum of HEA; b) and c) Bright field TEM image and SAED pattern of HEA; SEM micrograph of HEA
after d) SPS, e) SPS + HT at 1100 °C, f) SPS + HT at 1250 °C [124].
mous. Therefore, it is crucial to use highly effective, affordable techniques to support experimental research. CALPHAD modeling can be
considered as the simplest approach for alloy design because it
involves the global reduction of the system's Gibbs free energy as a
function of temperature and composition, However, it has issues with
incompletely reliable databases that cover all edge binaries and ternaries for HEA systems. Otto et al. [129] investigated the stability of the
phases in six different alloys, including CoCrFeMnNi, CoCrCuFeMn,
CoCrMnNiV, CoFeMnMoNi, CoFeMnNiV, and CrFeMnNiTi, and discovered that only CoCrFeMnNi HEA developed a single‐phase FCC
solid solution after annealing at 1000 °C or 800 °C. With the exception
of CoCrCuFeMn and CoFeMnNiV, when TTNI8 database was utilized,
all of the edge binaries and some ternaries of these alloys are covered
by the TCNI7 database, the CALPHAD calculations qualitatively replicate the experimental finding. Only in CoCrFeMnNi was a single FCC
phase predicted by both databases. In CoFeMnMoNi, an FCC matrix
phase and a minor phase were accurately predicted, while in CrFeMnNiTi, a complex microstructure with four phases was anticipated. The
FCC phase of CoFeMnNiV was anticipated to be stable across a constrained temperature range, whereas the minor phase develops at
lower temperatures (Fig. 35). Since Otto et al. [129] detected it in
CoCrCuFeMn and CoCrMnNiV, it appears that both databases significantly underestimate the thermal stability of the phase in the systems.
This suggests that both databases still need to be improved, particularly the descriptions of all ternary systems.
A thermodynamic dataset for the Al‐Co‐Cr‐Fe‐Ni alloy system was
generated by Zhang et al. [130] by extending binary and ternary systems to larger composition ranges. Quaternary and quinary interaction
parameters are not taken into account by this database. They developed Al‐Co‐Cr‐Fe‐Ni phase diagrams utilizing their database that are
jected to differebt HT temperatures of 500–1000 °C for 4 h to analyze
the effect of HT temperature. The HEA exhibited debdritic microstructure. It was stable upto 800 °C. The dendritic microstructure of BCC
phase was dissolved in the matrix of HEA at 1000 °C and the speherodized FCC2 phases were observed (Fig. 34a – d). The HEA showed
reduction in strength with increase in HT temperature (Fig. 34e). It is
attributed to the BCC phase dissolution. Also the HEA exhibited softening due to the grain growth and stress relieving at increased HT temperatures. The HEA showed drastic drastic reduction in tensile
strength and elongation due to the evolution of harder σ (Cr5Fe6Mn8)
phases at 800 °C. The SEM tensile fractographs of the HEA showed
ductile failure mode in as‐cast and HT of 600 °C. The brittle failure
mode was observed for the HEA heat treated at 800 °C. However,
the HEA showed full elongation after HT at 1000 °C (Fig. 34f – i).
He et al. [42] showed that the precipitation of L12‐Ni3(Ti, Al) hardening phase in the matrix of FCC enabled the YS of (CoCrFeNi)94Al4Ti2
HEA to significantly rise from 503 MPa to 1005 MPa after being
heated at 650 °C for 4 h immediately following water quenching. Thus,
the mechanical characteristics of HEAs can be further improved using
heat treatment. The heat treatment showed greater effect on
microstructural characteristics and mechanical properties of HEAs.
The research papers documented on heat treatment of HEAs are limited. The variables of HT need to be optimized to attain desired
microstructure and enhance the mechanical properties of HEAs.
7. Simulation of HEAs
The fact that HEA has several primary components leading to a
huge variety of potential compositions. The expense and effort
required to investigate every composition experimentally are enor52
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Fig. 31. A) sem micrograph of heat treated al0.5cocrfeni hea at 650 °C for 0.5 h, 1 h, 2 h, 4 h and 8 h; b) TEM micrograph of Al0.5CoCrFeNi alloy heat-treated at
650 °C/8h; c) tensile curves of as casted and heat-treated HEA [125].
ditions, a particular material is chosen, such as various grades of stainless steels, dispersion‐strengthened steels, nickel‐based superalloys,
etc. Due to its increased microstructural stability throughout a larger
temperature range and for a longer length of time, the HEAs demonstrates superior functionality at room temperature, cryogenic temperature, and elevated temperature. This makes it suitable as a potential
engineering material for high temperature and corrosive sections of
power plants. The HEAs must be able to be joined with nickel‐based
superalloys and stainless steel. For the practicality of employing HEAs
in real‐world engineering applications, a thorough, systematic examination of the weldability of HEAs with conventional engineering alloys
is required.
Massive structural and mechanical parts in power generation sectors are being redesigned to mitigate complex and expensive casting
and machining operations and successive transportation from vendor
to industrial sectors owing to of the growing competitive pressure in
material processing and manufacturing. Manufacturing substructures
and joining them together using cutting‐edge welding techniques is
one such development in the design and production of structural sections. Moreover, weld repair of engineering parts and assemblies is
gaining importance as a way to prolong service life and reduce replacement costs [37,133]. Such materials must be capable to be joined satisfactorily to be able to meet the essential criteria of design and
consistent with the results of the existing experiments. In some circumstances, it appears that phase diagrams can still be obtained with some
degree of precision without the creation of new databases. The equilibrium phases and phase fractions of the alloys Al0.5CoCrCuFeNi,
Al23Co15Cr23Cu8Fe15Ni16, and Al8Co17Cr17Cu8Fe17Ni33 were
predicted using a Ni‐based superalloy database [52,131]. Experiments
and predictions are in agreement [52] with a few small exceptions.
These illustrations show how computational techniques may precisely
predict the phase of HEAs. It is anticipated that these techniques will
become crucial tools for the design and development of HEAs in the
future, even though the accuracy of computational phase predictions
for other HEA systems (other than Al‐Co‐Cr‐Cu‐Fe‐Ni) is still unclear
[132].
8. Future scope of HEAs in aeroengine sector
HEAs provides a promising approach to be employed as both structural and functional materials in aeroengines. The HEAs can be joined
with conventional alloy sytems to fulfill the requirements of diverse
operating environments. The differing operational temperatures, pressures, and stress levels experienced by aeroengines force the use of
diverse metallic systems. For each component and set of operating con53
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Fig. 32. BSE images of interdendritic region of as-cast AlCoCrFeNi HEA at different temperatures of HT for 3 h: a) as-cast; b) 850; c) 975; d) 1100; e) and f) 1200 °
C [126].
mance of HEAs based on their process structure and property
relationships.
Modern aeroengines employ materials that can withstand high
rotational speeds, creep, fatigue fractures, and high operating temperatures. As a result, the manufactured components must be small and
extremely strong at high temperatures, as well as resistant to fatigue,
chemical degradation, wear, and oxidation. HEAs are characteristics
of cutting‐edge high‐performance materials. These alloys are differentiated from conventional alloys because of their intricate compositions
of different alloying elements. This makes it more stable at elevated
temperatures owing to their greater entropy. This feature provides
potential opportunities as a structural material in aeroengines by permitting suitable alloying elements to advance the mechanical properties of the HEAs depending on four core effects. The HEAs are
mostly prepared using arc melting. However, as it is difficult to heat
the elements all at once, hypoeutectic HEA, which separates from
the other elements, is prone to form. Defects are also known to be
introduced during the casting process. But a powder‐based laser additive manufacturing technique known as Laser Engineering Net Shaping
(LENSTM) and Selective Laser Melting (SLM) offers adaptability, geometric accuracy, and the capacity to construct 3‐dimensional dense
structures layer by layer while avoiding production mistakes [134].
manufacture. The strength performance and microstructural integrity
of structural and mechanical parts to be used in critical applications
of aeroengines have been improved by latest innovations in welding
technologies and high‐performance engineering materials like HEAs.
Similar and dissimilar HEAs can be welded to improve their industrial
usage in aeroengine applications, particularly for the manufacture of
complex and large‐scale components [10].
The tensile properties, microhardness, and cryogenic tensile properties of HEAs are provided in the literature review. Research on the
notched sections of HEAs, such as the notch tensile performances
and notched Charpy impact toughness, as well as high temperature
characteristics like hot tensile, creep, and stress rupture properties has
not yet been documented. Research on high temperature characteristics of HEA including as hot tensile strength, creep and stress rupture
life, and hot corrosion resistance, is crucial for its use in the aeroengines and must be undertaken owing to. The mechanical behavior and
microstructural growth of HEAs were greatly affected by the heat
treatment. There have been very few studies reported on heat treatment of HEAs. To achieve the desired microstructure and improve
the mechanical properties of HEAs, the heat treatment must also be
optimised. There is a great deal of demand for research into the mechanisms underlying the strengthening and enhanced strength perfor54
Materials Science for Energy Technologies 7 (2024) 35–60
T. Sonar et al.
Fig. 33. Heat treatment of Al0.6CoFeNiCr0.4 HEA: SEM micrographas of HEAs in a) as-cast, b) 550 °C, c) 650 °C, d) 750 °C, e) 850 °C; f) XRD spectrum; g) tensile
curves; h) Tafel plot [127].
Fig. 34. HT of FeCrCuMnNi HEA: SEM microstructure of HEA for a) as-cast and b) 600 °C, c) 800 °C and d) 1000 °C HT conditions; e) tensile curves; SEM tensile
fractographs of HEAs for f) as-cast, g) 600 °C, h) 800 °C and i) 1000 °C HT conditions [128].
55
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Materials Science for Energy Technologies 7 (2024) 35–60
Fig. 35. Calculated equilibrium phase mole fraction vs temperature for a) CoCrFeMnNi, b) CoCrCuFeMn, c) CoCrMnNiV, d) CoFeMnMoNi, e) CoFeMnNiV, and f)
CrFeMnNiTi [129].
Fig. 36. The sections of high temperature in aeroengine [9].
56
Materials Science for Energy Technologies 7 (2024) 35–60
T. Sonar et al.
6. The HEAs must be capable of being welded and additively manufactured in similar and dissimilar configuration for its feasibility
in practical engineering applications. This demands extensive
research work on similar and dissimilar welding of HEAs.
7. CALPHAD modeling can be considered as the simplest approach for
alloy design because it involves the global reduction of the system's
Gibbs free energy as a function of temperature and composition,
However, it has issues with incompletely reliable databases that
cover all edge binaries and ternaries for HEA systems.
8. HEAs provides a promising approach to be employed as both structural and functional materials in aeroengines. The HEAs can be
joined with conventional alloy sytems to fulfill the requirements
of diverse operating environments.
HEAs were discovered as a result of material breakthroughs and
technical developments in the method of synthesis [135]. With a
greater ratio of strength‐to‐weight, excellent resistance to oxidation
and fatigue, thermal resistance, greater strength at elevated temperature, wear and resistance to creep and stress rupture, HEAs are great
materials for aeroengine applications such as turbines blades, vanes,
stator and rotors, combustion chambers, exhaust nozzles, and gas turbine cases as shown in Fig. 36. The four primary components of aeroengines are the compressor, combustor, turbine blade, and injector.
The development of innovative materials for aeroengines that meet
standards for thrust, weight, safety, fuel efficiency, life cycle costs,
and environmental considerations has lately been challenged by competition from the aerospace sector. The development and implementation of structural materials that would provide higher performance and
be more cost‐effective to manufacture as well as repair than present
components is required by contemporary current improvements and
progression in the aeroengine sector. The ideal alloy that could withstand extreme temperatures and still be lightweight will be chosen
based on the working environment. The material distribution of an
aeroengine consists of steels, titanium alloys, nickel superalloys, aluminum alloys, and more recently, HEAs [136].
Success in the creation and realization of engineering materials in
the appropriate shapes and conditions is directly correlated with success in aeronautical science and technology. The lack of engineering
materials with the appropriate characteristics frequently prevents
advancements and enhancements to the current aeronautical systems.
It is a major challenge for materials scientists to create materials that
can operate effectively in harsh environments, and this necessitates a
fundamental knowledge of how materials react to extremely high temperatures, high stresses, high strain rates, and corrosive environments,
such as those found in rocket engines. Contrary to many other research
disciplines, materials processing technologies are expensive and time‐
consuming while also being infrastructure‐sensitive. As a result, careful preparation is necessary before beginning any activity using new
material systems [137]. System designers must specify the HEAs
needed up front so that development work may start well in advance
and move forward to meet deadlines.
CRediT authorship contribution statement
Tushar Sonar: Conceptualization, Methodology, Resources, Investigation, Data Curation, Formal Analysis, Writing – Original draft,
Writing – review & editing. Mikhail Ivanov: Conceptualization,
Methodology, Funding acquisition, Resources, Investigation, Formal
Analysis, Writing – review & editing. Evgeny Trofimov: Conceptualization, Methodology, Funding acquisition, Resources, Investigation,
Formal Analysis, Writing – review & editing. Aleksandr Tingaev: Conceptualization, Methodology, Methodology, Resources, Investigation,
Formal Analysis, Writing – review & editing. Ilsiya Suleymanova:
Conceptualization, Methodology, Resources, Investigation, Formal
Analysis, Writing – review & editing.
Declaration of Competing Interest
The authors declare that they have no known competing financial
interests or personal relationships that could have appeared to influence the work reported in this paper.
Acknowledgements
This work is supported by the Ministry of Science and Higher Education of the Russian Federation (Grant No. FENU‐2023‐0013).
9. Conclusions
References
1. HEAs are alloys with an atomic percentage range of between 5 and
35% or at least five major alloying elements. Four main effects are
induced by the multi‐principal element characteristics of HEAs:
greater entropy, extreme lattice distortion, sluggish diffusion, and
cocktail effects. It sets HEAs apart from classic alloys.
2. HEAs integrate tensile characteristics better than other materials,
improving both strength and ductility. In contrast to FCC‐
structured HEAs, which have low strength and high plasticity,
BCC‐structured HEAs have great strength and low plasticity. The
strength and hardness of HEAs are therefore primarily determined
by their structural characteristics.
3. HEAs demonstrated excellent mechanical performance at both
ambient, cryogenic and elevated temperatures, as well as strong
corrosion resistance and endurance limit, in addition to their superior mechanical characteristics.
4. The diversity of elemental contents and significantly higher mixing
entropy of HEAs make them mechanically superior to classic metals
and alloys. It also shows better strength to weight ratio. Hence, it qualifies as a possible structural material for engineering applications.
5. The microstructural development and mechanical characteristics of
HEAs were greatly affected by the heat treatment. To accomplish
the suitable microstructural features and enhance the mechanical
integrity of HEAs, the heat treatment must be optimized. The influence of duration, quenching medium, and heating and cooling rates
on the microstructure and mechanical characteristics of HEA are also very crucial.
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