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Scripta Materialia 244 (2024) 116031
Available online 13 February 2024
1359-6462/© 2024 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Uniting superior mechanical properties with oxidation resistance in a
refractory high-entropy alloy via Cr and Al alloying
Dingcong Cui
a
, Xin Liu
a
, Zhongsheng Yang
a
, Bojing Guo
a
, Zhijun Wang
a
, Junjie Li
a
,
Jincheng Wang
a
,
*
, Feng He
a
,
b
,
c
,
*
a
State Key Laboratory of Solidication Processing, Northwestern Polytechnical University, Xi’an 710072, China
b
Research & Development Institute of Northwestern Polytechnical University in Shenzhen, Shenzhen 518063, China
c
Collaborative Innovation Center of Northwestern Polytechnical University, Shanghai 201100, China
ARTICLE INFO
Keywords:
Refractory high-entropy alloys
Spinodal decomposition
Mechanical properties
Oxidation resistant
ABSTRACT
Refractory high-entropy alloys (RHEAs) have attracted considerable interest due to their elevated melting points
and remarkable softening resistance. Nevertheless, the ambient-temperature brittleness and inadequate high-
temperature oxidation resistance commonly limit the application of the body-centered-cubic (BCC) RHEAs. In
this study, we achieved a Ti
41
V
27
Hf
11.5
Nb
11.5
Cr
3
Al
6
RHEA with a desirable yield strength of ~1178 MPa and
tensile ductility of ~19.5 %. Exploring the underlying mechanisms, we demonstrated that Cr and Al alloying
induced a nanoscale spinodal structure and generated a signicant lattice mist, resulting in a notable
strengthening effect and pinning behavior. Meanwhile, dislocation congurations involving loops and cross slips
were stimulated by pinning, serving a reliable strain-hardening capability to large strains. Signicantly, Cr and Al
alloying improved oxidation resistance and prevented severe spallation at high temperatures by forming pro-
tective oxide layers. These results provide opportunities to design novel RHEAs.
Refractory high-entropy alloys (RHEAs), a subgroup of extensively
alloyed systems, have garnered increasing attention due to their ca-
pacity to retain signicant yield strength under extremely elevated
temperatures [1–3]. Yet several serious obstacles exist in the RHEAs
with body-centered-cubic (BCC) structures [4,5]. First, the Achilles’ heel
of most BCC RHEAs is their negligible tensile ductility at ambient tem-
perature. For instance, NbMoTaW and VNbMoTaW only exhibit fracture
strain of about 2 % under ambient compression tests [6]. Second, many
discovered that RHEAs suffer from poor oxidation resistance, like pure
refractory metals, owing to substantial scale spallation or even complete
oxidation at elevated temperatures [7,8]. Adding oxidation-resistant
elements (e.g., aluminum (Al), chromium (Cr), and silicon (Si)) to
RHEAs is a promising path to achieve high-temperature oxidation
resistance [9,10]. However, Al, Cr, and Si elements exhibit chemical
activity with refractory metallic elements and commonly promote the
formation of brittle phase, degrading the room temperature ductility
[11,12]. Consequently, the RHEA rush has led to a few alloys exhibiting
application-worthy mechanical performances.
Recently, Wei et al. [13] designed Ti
38
V
15
Nb
23
Hf
24
RHEAs with
as-cast tensile ductility exceeding 20 % and yield strength reaching 800
MPa. However, the catch is that this system’s yield strength and
oxidation resistance are insufcient. In the case of the former, the spi-
nodal structure induced by a high absolute value of mixing enthalpy
(ΔH
m
) can greatly increase the strength of RHEAs [14,15]. This
enhancement can be achieved by alloying with certain metallic elements
like Al [16]. Regarding the latter case, Ouyang et al. [17] reported that
Ti
38
V
15
Nb
23
Hf
24
RHEA cannot form a dense passivation oxide layer to
prevent the rapid diffusion of oxygen. The protective oxide layers, such
as Cr
2
O
3
and Al
2
O
3
, are the prerequisite to achieving an outstanding
oxygen-shielded capability [18].
In this study, we took advantage of the intrinsic characteristics of Cr
and Al elements to develop a novel RHEA with superior mechanical
properties and improve its oxidation resistance. Initially, we adopted an
updated Ti
41
V
27
Hf
16
Nb
16
RHEA as the ductile base alloy. This choice
mitigates the likelihood of brittle phase formation by decreasing the Hf
content [19]. Supplementary Table. S1 illustrates the ΔH
m
between the
atom pairs. The values of ΔH
m
are pretty large for the atom pairs
involving the refractory element (Ti, V, Hf, or Nb) and Cr/Al element
[20]. Hence, alloying with appropriate Cr and Al in this alloy can
probably introduce spinodal structure and improve oxidation resistance.
* Corresponding authors.
E-mail addresses: jchwang@nwpu.edu.cn (J. Wang), fenghe1991@nwpu.edu.cn (F. He).
Contents lists available at ScienceDirect
Scripta Materialia
journal homepage: www.journals.elsevier.com/scripta-materialia
https://doi.org/10.1016/j.scriptamat.2024.116031
Received 11 December 2023; Received in revised form 23 January 2024; Accepted 7 February 2024
Scripta Materialia 244 (2024) 116031
2
Supplementary Table. S2–3 provides empirical phase prediction pa-
rameters for HEAs. The brittle behavior tends to happen when Cr con-
tent is over 6 at.% owing to a large valence electron concentration (VEC
>4.5) [21]. According to the previous research, the ductility would be
reduced in the AlHfNbTiZr RHEA when the Al content surpasses 7 at.%
[22]. Overall, our investigation focused on the mechanical properties
and oxidation resistance of the Ti
41
V
27
Hf
16
Nb
16
RHEA, with the addi-
tion of Cr and Al (≤6 at.%).
Ti
41
V
27
Hf
16
Nb
16
(at.%, denoted as base RHEA) and Ti
41
V
27
Hf
16-(x +
y)/2
Nb
16-(x +y)/2
Cr
x
Al
y
(where x =3, 6; y =6) were prepared in a water-
cooled copper crucible by vacuum arc melting, and their ingots were
remelted ve times to ensure chemical homogeneity. The purity of every
alloying element exceeded 99.9 wt.%. Crystalline structures were
identied using X-ray diffractometry (XRD) (X’Pert PRO) with Cu-K
α
radiation at a scanning rate of 2◦/min. The samples intended for scan-
ning electron microscopy (SEM), energy dispersive spectrometer (EDS),
and electron backscattered diffraction (EBSD) characterization were
ground to 4000-grit SiC paper and polished with SiO
2
suspension, then
characterized by TESCAN MIRA-III equipped with EBSD/EDS detector.
Aztec-HKL software was employed for grain size analysis. The samples
for transmission electron microscope (TEM) observation were mechan-
ically ground to ~40
μ
m thickness, followed by ion milled, and nally
identied by a double Cs corrector TEM (Themis Z).
Ambient-temperature uniaxial tensile testing was conducted on a
TSMT mechanical testing platform at a strain rate of 1 ×10
−3
s
−1
. The
tensile specimens with a gauge geometry of
12.5 mm ×3.0 mm ×2.0 mm were sectioned using electrical dis-
charging. At least three specimens were tested for each composition to
conrm the reproducibility. Before being tested for digital image cor-
relation (DIC) analyses, the speckle patterns were coated on the surface
of the tensile sample, and the resulting deformed characteristics were
captured by a high-speed camera (frame rate, 1 Hz). Zeiss Inspect
Correlate software was used to calculate the local strain proles during
the DIC experiment. The oxidation resistance of the alloys was assessed
at 800 and 1000 ◦C, with exposure times of 2, 4, 6, and 10 h in an
ambient air environment.
We rst examined the ambient-temperature tensile properties of
these RHEAs. As presented in Fig. 1, the base RHEA possesses a yield
strength (
σ
y
) of 953 ±23 MPa with a fracture elongation (
ε
f
) of 25.2 ±
0.2 %. Alloying with 3 at.% Cr enhances the tensile properties of the base
RHEA, while 6 at.% Cr embrittles the alloy (see the engineering stress-
strain curves of RHEAs alloying with Cr in Supplementary Fig. S1).
We proceeded to alloy with 6 at.% Al in the Ti
41
V
27
Hf
14.5
Nb
14.5
Cr
3
RHEA. The
σ
y
of the obtained Ti
41
V
27
Hf
11.5
Nb
11.5
Cr
3
Al
6
RHEA (denoted
as CrAl) exhibited a signicant increase of almost 24 %, reaching a value
of 1178 ±11 MPa. Notably, this increase in
σ
y
was achieved without
compromising its tensile ductility (
ε
f
=19.5 ±1.4 %). SEM images of the
fractured surface show dimples (see inset in Fig. 1(a)), indicating a
ductile manner. After conducting a comparative analysis of the present
RHEAs with other RHEAs [5,15,23-29], it becomes apparent that CrAl
demonstrates an unprecedented combination of yield strength and ten-
sile ductility in the as-cast state (refer to Fig. 1(b)).
The insight into the relationships between microstructure and me-
chanical response will provide a fundamental basis for the perfor-
mances. The EBSD-IPF maps and XRD curves (Supplementary Fig. S2)
indicate that the base and CrAl RHEAs display a fully equiaxed grain
morphology with BCC-phase constitution. Fig. 2(a-c) illustrates the
high-angle annular dark eld-scanning transmission electron micro-
scope (HAADF-STEM) micrographs of CrAl under the [111] zone axis.
The plate-like nanoscale phases (β*) with dimensions of 20~300 nm in
length come into being, exhibiting varying contrast and featuring a
diffusional transition zone instead of a clearly dened interface. The
variation in image contrast can be attributed to the Z-contrast of the
regions rich in light or heavy atoms (Fig. 2(d-f)) [30]. Therefore, the
brighter contrast of β* is due to the concentration of the Hf element,
which has a more signicant atomic number than the others. Laube et al.
[31] found a sharp interface between the decomposed phases in another
Al-containing RHEA, indicating a phase nucleation and growth mecha-
nism. The blurred boundary dominated by ‘uphill’ diffusion in CrAl
should be a case of spinodal decomposition.
The associated fast Fourier transform (FFT) patterns in Fig. 2(c)
conrm that both indicated regions correspond to BCC phases, revealing
the characteristic of spinodal decomposition in CrAl, a phenomenon not
seen in the base RHEA (Supplementary Fig. S3). Fig. S3 shows no evident
chemical heterogeneity or second phase in the base RHEA. Hence,
introducing Cr and Al atoms facilitates the development of spinodal
structure. Additionally, FFT patterns have not yet exhibited extra discs,
indicating the lack or insignicance of long-range/short-range ordering
(L/SRO). According to XRD results in Supplementary Fig. S2(c-d), the
(200) diffraction peak of CrAl shows a sideband due to the decomposed
phases and the periodic difference of their lattice parameters. The lattice
parameters of the matrix phase (β) and β* can be calculated to be 3.273
Å and 3.258 Å, respectively, exhibiting a mist of 0.44 %.
As depicted in Fig. 3(a-b), the shear strain (e
xx
) of the spinodal
structure was mapped using the geometric phase analysis (GPA) method
[32]. The denition of the strain is with respect to the undistorted lattice
[33]. The β* phase resulted in an inhomogeneous strain distribution,
with the highest strain at the coherent interface with the β phase. On the
one hand, spinodal structure enhanced the chemical heterogeneity, thus
increasing the lattice distortion (from XRD analysis and GPA mapping)
and inuencing the strengthening mechanisms. On the other hand, the
spinodal structure created a high density of diffusive boundaries, which
Fig. 1. The tensile strength-ductility combination achieved in the Ti
41
V
27
Hf
16
Nb
16
and Ti
41
V
27
Hf
11.5
Nb
11.5
Cr
3
Al
6
RHEAs at room temperature. (a) Engineering
stress-strain curves with SEM images of the fracture surface of CrAl in the inset. (b) Tensile yield strength versus fracture elongation, compared with the previously
reported BCC RHEAs [5,15,23–29].
D. Cui et al.
Scripta Materialia 244 (2024) 116031
3
Fig. 2. Microstructure of the CrAl RHEA. (a) HAADF-STEM image showing nanoscale spinodal structure. (b) Spinodal phases (β*) distribute in the matrix (β). (c) β*
and β phases show the same BCC crystal structures (see the selected area electron diffraction patterns). (d) Atomic-resolution HAADF-STEM image showing coherency
at the β and β* interface. (e) Zoom-in image of the blue square region in (d), showing varying intensity from atomic column to column. (f) Intensity color map for the
atomic columns corresponding to (d).
Fig. 3. Lattice strain mapping and impact mechanisms of spinodal structure in the CrAl RHEA. (a) Magnied HAADF-STEM image of the spinodal structure. (b) The
corresponding strain distribution before tensile deformation calculated by geometric phase analysis (GPA). (c) Estimated yield strength of the base and CrAl RHEAs
caused by the solid solution strengthening (
σ
ss), grain boundary strengthening (
σ
gb), and spinodal-decomposition strengthening (
σ
sds). The corresponding experi-
mental yield tensile strengths (
σ
yts) are marked by red dots. The inset indicates that
σ
sds provides the enhanced yield tensile of CrAl. (d) Dislocation pinning in the 2 %
deformed CrAl RHEA. (e) Pining points (marked as the white arrows) at the interface between β and β* phases (marked as the red dotted lines). (f) Lattice strain
distribution around the pinning points in (e). (For interpretation of the references to color in this gure legend, the reader is referred to the web version of
this article.).
D. Cui et al.
Scripta Materialia 244 (2024) 116031
4
acted as effective barriers to dislocation motion [34].
Based on the acknowledged solid-solution strengthening models
proposed by Maresca’s model [35] and Kato’s model [36] for
spinodal-modulated BCC alloys, we calculated the theoretical
σ
y
of both
RHEAs considering solid solution strengthening (
σ
ss) contributed by the
alloying elements, Hall-Petch strengthening (
σ
gb) resulting from the
grain renement, and spinodal-decomposition strengthening (
σ
sds).
Fig. 3(c) illustrates the enhancement from these mechanisms (detailed in
Notes. 1). Alloying with Cr and Al decreases
σ
ss by ~132 MPa and
σ
gb by
~7 MPa. The increased
σ
yts in CrAl is attributed to the effective
σ
sds
(~340 MPa). Due to the existence of micro-segregations (Fig. S2), we
should note potential discrepancies in
σ
sds. According to research on
spinodal-modulated Ti
41
V
27
Hf
15
Nb
15
O
2
RHEA, it has been reported that
the large lattice mist of 0.59 % generated by β and β* phases can
provide lattice mist strengthening (~293 MPa) for
σ
sds [15]. CrAl
achieves a comparable lattice mist of 0.44 %, the dominating factor for
enhancing yield strength.
The spinodal strengthening fundamentally linked to the β/β* inter-
face. TEM observations of deformed CrAl at 2 % strain (Fig. 3(d-e))
revealed the interaction between dislocations and spinodal structure.
The dispersed high-strain regions in Fig. 3(f) correspond to the spinodal
interfaces, showing large elastic strains due to the atomic mismatch. As
such, an unusually rugged landscape is introduced, enhancing the
resistance to dislocation motion [37]. The conventional wisdom holds
that the energy would be relaxed through crack propagation if the
dislocation activity remains inactive [38,39]. As a result, pinning typi-
cally increases the strength by promoting dislocation multiplication at
the expense of ductility [40]. The promising ductility of CrAl calls for a
deeper assessment of the underlying deformation mechanisms.
Fig. 4(a) plots the normalized strain-hardening rate (SHR) versus the
true strain curves of CrAl. After experiencing a decline in stage I (0–5%),
the SHR continuously increases throughout stage II (approximately 5–10
%), temporarily stabilizing in stage III and exhibiting a reduction until
the fracture occurs. As shown in Fig. S4, this well-developed SHR is
responsible for high tensile ductility, which is rarely observed in other
RHEAs [25,26,41]. TEM investigations were performed to understand
the origin of SHR from a plastic-ow perspective. The macro digital
image correlation (DIC) analyses prove that deformed CrAl exhibits a
non-uniform strain distribution during tension (Fig. 4(b)). To guarantee
precision, TEM samples were prepared by subjecting them to incre-
mental tensile strains in accordance with the strain prole.
Now, let us return to the deformation mechanisms of CrAl. In stage I,
planar slips prevail, accompanied by abundant wavy dislocations (Fig. 4
(c)). There are two conceivable reasons for the coexistence of planar and
wavy dislocations. First, the spinodal interfaces act as obstructive agents
and cause dislocation pinning. The heavily curved dislocations may
represent the local bow-out at discrete pinning points [42]. Second,
when dislocations with screw type in the planar slips reach high-stress
levels, cross slips are facilitated owing to the strain effect [43]. The
cross-slip easily produced a wavy dislocation morphology [44]. In stage
II (Fig. 4(d)), cross-slip can contribute to obstacle bypass, enabling more
efcient utilization of undeformed regions and promoting deformation
delocalization [45]. Meanwhile, a signicant quantity of dislocation
loops were present at 7 % strain, as evidence of dislocation multiplica-
tion from pinning [14]. The pinning and de-pinning process of disloca-
tion reportedly produced loops and glissile dislocations [46]. Cross-slip
and pinning alone were insufcient for dislocation multiplication. The
formed dislocation junctions acted as new sources for generating mobile
dislocations, effectively giving rise to a rising SHR. Similar dislocation
congurations that can increase the dislocation density were also found
Fig. 4. Mechanical behavior and dislocation evolution of the CrAl RHEA during plastic deformation. (a) True stress-strain curve and plot of strain-hardening rate for
the CrAl RHEA. (b) The local strain distribution for a typical tensile sample. The samples at different strains for TEM characterization are extracted from the cor-
responding positions (marked as red lines). (c-f). Evolution of dislocation patterns. (c) Planar slips at 2 % strain. (d) Cross slips (marked as white arrows) and
dislocation loops (marked as red arrows) at 7 % strain. (e) Double slips and formed dislocation tangles (marked as a red arrow) at 10 % strain. (f) Dipolar walls are
full of dislocation tangles at 15 % strain. (For interpretation of the references to color in this gure legend, the reader is referred to the web version of this article.).
D. Cui et al.
Scripta Materialia 244 (2024) 116031
5
in the TiZrHfNb RHEAs [38]. Fig. 4(e) highlights the proliferation of
second-generation slip bands. Double slips provide opportunities for
dislocation intersection and sow dislocation tangles, thereby compen-
sating for the reduction in SHR caused by cross-slips (stage III) [47]. In
stage IV, the dislocation pattern behaves like dipolar walls full of
dislocation tangles (Fig. 4(f)) [48]. As such, the combined effects of
dislocation multiplication, delocalization, and patterning jointly lead to
the favorable SHR and ductility. Moreover, we note that alloying with
trace interstitial elements (oxygen or nitrogen) can deliver a
spinodal-decomposition strengthening, like in (TiZrNbTa)
99.1
N
0.9
[49]
and (TiVHfNb)
98
O
2
[15], while leading to signicant deterioration of
the ductility. These results revealed that the advantages of Cr and Al
alloying are reected in maintaining ductility even beyond gigapascal
stresses.
The investigation includes an assessment of the oxidation resistance
of the RHEAs (exposed to air at 800 and 1000 ◦C). The mass change
curves in Fig. 5(a) indicate that CrAl exhibits consistently lower mass
gain than base RHEA as the exposure temperature and time increase.
Correspondingly, the samples exposed at 1000 ◦C demonstrated that
CrAl could maintain structural integrity for up to 6 h. Conversely, the
base alloy experienced severe spallation under identical conditions
(Fig. 5(b)). RHEA oxide layers at 1000 ◦C-2 h were selected and inves-
tigated. As depicted in Fig. 5(c), XRD patterns demonstrate the presence
of mixed oxides consisting of TiO
2
, Nb
2
O
5
, HfO
2
, and TiNb
2
O
7
in the
base alloy. In contrast, there exist extensive diffraction peaks of Cr
2
O
3
and Al
2
O
3
in CrAl. As a result, the thickness of the oxidation layer is
much thinner for CrAl (~383
μ
m) compared with the base alloy (~800
μ
m), as shown in Fig. 5(d-e). Refractory elements tend to generate
porous or volatile oxide scales, making oxygen diffuse into the alloy
[18]. As such, Ti, V, Hf, Nb, and O are evenly distributed within the
oxide scales of both alloys. Cr
2
O
3
and Al
2
O
3
scales can enable the in-
hibition of oxygen dissolution and metallic ion external diffusion [8].
That is, the formed Cr
2
O
3
and Al
2
O
3
retard the rapid growth of the outer
scale and postpone spallation in CrAl.
Fig. 5. Oxidation resistance of the base and CrAl RHEAs. (a) The mass gain versus exposure time at different temperatures (800 ℃ and 1000 ℃). (b) The mor-
phologies of two RHEAs after 2 h, 4 h, and 6 h exposure at 1000 ℃. (c) XRD patterns of the RHEAs after 1000
◦C-2 h oxidation. The cross-sectioned microstructures of
1000 ◦C-2 h oxidized (d) base and (e) CrAl RHEAs. The element maps were measured by EDS.
D. Cui et al.
Scripta Materialia 244 (2024) 116031
6
RHEAs containing Hf element frequently show suboptimal resistance
to oxidation. Sheikh et al. [7] reported Ti
1.5
ZrTa
0.5
Hf
0.5
Nb
0.5
suffers
pesting failure within the temperature range of 600–1000 ◦C, where the
bulk alloy disintegrates into powders. In contrast, catastrophic oxidation
has not been seen in CrAl. However, the mass gain gradually increases
with time (Fig. 5(b)). The formed oxide scales in CrAl cannot completely
prevent the diffusion of oxygen and will continuously grow into the
matrix until the alloy is fully oxidized. Oxygen-shielded ability is the
necessity for further improving this alloy system’s oxidation resistance.
In summary, we investigated how mechanical properties and
oxidation resistance can be successfully improved for a class of RHEAs.
Via composition modulation, we put alloying with appropriate Cr and Al
to generate a nanoscale spinodal structure and enhance yield strength to
1178 MPa. Meanwhile, dislocation congurations were induced by the
pinning impact of the spinodal structure. Dislocation intersection and
delocalization guaranteed an excellent working hardening ability and
thus obtained the elongation of ~19.5 %. The improved oxidation
resistance of CrAl RHEA at 1000 ℃ can be attributed to the formation of
protective Cr
2
O
3
and Al
2
O
3
layers, which mitigated severe spallation.
These ndings accelerate the development of spinodal-modulated
RHEAs and provide a new candidate for structural RHEAs.
Declaration of competing interest
The authors declare that they have no known competing nancial
interests or personal relationships that could have appeared to inuence
the work reported in this paper.
Acknowledgement
This work was supported by the National Natural Science Foundation
of China (Granted No. 52001266), the Fundamental Research Funds for
the Central Universities (No. G2022KY05109), the Guangdong Basic and
Applied Basic Research Foundation (No. 2023A1515012703), the
Shanghai “Phosphor” Science Foundation, China (No. 23YF1450900),
the Research Fund of the State Key Laboratory of Solidication Pro-
cessing (NPU), China (Grant No. 2023-QZ-02), the Young Elite Scientists
Sponsorship Program by CAST, (No. 2023QNRC001), and the Innova-
tion Foundation for Doctor Dissertation of Northwestern Polytechnical
University (No. CX2023044). The Analytical & Testing Center of
Northwestern Polytechnical University is acknowledged for providing
characterization facilities.
Supplementary materials
Supplementary material associated with this article can be found, in
the online version, at doi:10.1016/j.scriptamat.2024.116031.
References
[1] O.N. Senkov, J.M. Scott, S.V. Senkova, F. Meisenkothen, D.B. Miracle, C.
F. Woodward, Microstructure and elevated temperature properties of a refractory
TaNbHfZrTi alloy, J. Mater. Sci. 47 (9) (2012) 4062–4074.
[2] C. Lee, G. Kim, Y. Chou, B.L. Musico, M.C. Gao, K. An, G. Song, Y.C. Chou,
V. Keppens, W. Chen, P.K. Liaw, Temperature dependence of elastic and plastic
deformation behavior of a refractory high-entropy alloy, Sci. Adv. 6 (37) (2020)
12.
[3] R. Feng, B.J. Feng, M.C. Gao, C. Zhang, J.C. Neuefeind, J.D. Poplawsky, Y. Ren,
K. An, M. Widom, P.K. Liaw, Superior high-temperature strength in a
supersaturated refractory high-entropy alloy, Adv. Mater. 33 (48) (2021) 9.
[4] L. Wang, J. Ding, S. Chen, K. Jin, Q. Zhang, J. Cui, B. Wang, B. Chen, T. Li, Y. Ren,
S. Zheng, K. Ming, W. Lu, J. Hou, G. Sha, J. Liang, L. Wang, Y. Xue, E. Ma, Tailoring
planar slip to achieve pure metal-like ductility in body-centred-cubic multi-
principal element alloys, Nat. Mater. 22 (2023) 950–957.
[5] Z. Wang, S. Chen, S. Yang, Q. Luo, Y. Jin, W. Xie, L. Zhang, Q. Li, Light-weight
refractory high-entropy alloys: a comprehensive review, J. Mater. Sci. Technol. 151
(2023) 41–65.
[6] O.N. Senkov, G.B. Wilks, J.M. Scott, D.B. Miracle, Mechanical properties of
Nb25Mo25Ta25W25 and V20Nb20Mo20Ta20W20 refractory high entropy alloys,
Intermetallics 19 (5) (2011) 698–706.
[7] S. Sheikh, M.K. Bijaksana, A. Motallebzadeh, S. Shafeie, A. Lozinko, L. Gan, T.
K. Tsao, U. Klement, D. Canadinc, H. Murakami, S. Guo, Accelerated oxidation in
ductile refractory high-entropy alloys, Intermetallics 97 (2018) 58–66.
[8] F. Müller, B. Gorr, H.-J. Christ, J. Müller, B. Butz, H. Chen, A. Kauffmann,
M. Heilmaier, On the oxidation mechanism of refractory high entropy alloys,
Corrosion Sci. 159 (2019) 108161.
[9] X. Pei, Y. Du, H. Wang, T. Li, M. Hu, H. Wang, W. Liu, Effects of Al/Si on the
oxidation behavior of a TiZrV0.5Nb0.5 refractory high entropy alloy at 1000 ℃,
Corrosion Sci. 224 (2023) 111527.
[10] Z. Zhou, X. Peng, W. Lü, S. Yang, H. Li, H. Guo, J. Wang, Ultra-high temperature
oxidation resistant refractory high entropy alloys fabricated by laser melting
deposition: Al concentration regulation and oxidation mechanism, Corrosion Sci.
221 (2023) 111537.
[11] Z.Q. Xu, Z.L. Ma, Y. Tan, X.W. Cheng, Designing TiVNbTaSi refractory high-
entropy alloys with ambient tensile ductility, Scr. Mater. 206 (2022) 6.
[12] X. Huang, J. Miao, A.A. Luo, Order-disorder transition and its mechanical effects in
lightweight AlCrTiV high entropy alloys, Scr. Mater. 210 (2022) 114462.
[13] S.L. Wei, S.J. Kim, J.Y. Kang, Y. Zhang, Y.J. Zhang, T. Furuhara, E.S. Park, C.
C. Tasan, Natural-mixing guided design of refractory high-entropy alloys with as-
cast tensile ductility, Nat. Mater. 19 (11) (2020) 1175–1181.
[14] Y. Chen, Y. Fang, R. Wang, Y. Tang, S. Bai, Q. Yu, Achieving high strength and
ductility in high-entropy alloys via spinodal decomposition-induced compositional
heterogeneity, J. Mater. Sci. Technol. 141 (2023) 149–154.
[15] D. Cui, Y. Zhang, L. Liu, Y. Li, L. Wang, Z. Wang, J. Li, J. Wang, F. He, Oxygen-
assisted spinodal structure achieves 1.5GPa yield strength in a ductile refractory
high-entropy alloy, J. Mater. Sci. Technol. 157 (2023) 11–20.
[16] S. Zeng, Y. Zhou, H. Li, H. Zhang, H. Zhang, Z. Zhu, Microstructure and mechanical
properties of lightweight Ti3Zr1.5NbVAlx (x =0, 0.25, 0.5 and 0.75) refractory
complex concentrated alloys, J. Mater. Sci. Technol. 130 (2022) 64–74.
[17] D. Ouyang, Z.-J. Chen, H.-B. Yu, K.C. Chan, L. Liu, Oxidation behavior of the
Ti38V15Nb23Hf24 refractory high-entropy alloy at elevated temperatures,
Corrosion Sci. 198 (2022) 110153.
[18] B. Gorr, S. Schellert, F. Muller, H.J. Christ, A. Kauffmann, M. Heilmaier, Current
status of research on the oxidation behavior of refractory high entropy alloys, Adv.
Eng. Mater. 23 (5) (2021) 14.
[19] S. Sheikh, S. Shafeie, Q. Hu, J. Ahlstr¨
om, C. Persson, J. Veselý, J. Zýka, U. Klement,
S. Guo, Alloy design for intrinsically ductile refractory high-entropy alloys, J. Appl.
Phys. 120 (16) (2016) 164902.
[20] A. Takeuchi, A. Inoue, Classication of bulk metallic glasses by atomic size
difference, heat of mixing and period of constituent elements and its application to
characterization of the main alloying element, Mater. Trans. 46 (12) (2005)
2817–2829.
[21] J.W. Yeh, Alloy design strategies and future trends in high-entropy alloys, Jom 65
(12) (2013) 1759–1771.
[22] Y.D. Wu, J.J. Si, D.Y. Lin, T. Wang, W.Y. Wang, Y.D. Wang, Z.K. Liu, X.D. Hui,
Phase stability and mechanical properties of AlHfNbTiZr high-entropy alloys,
Mater. Sci. Eng. A. 724 (2018) 249–259.
[23] S. Zhang, X. Ding, X. Gao, R. Chen, W. Cao, Y. Su, J. Guo, Dual enhancement in
strength and ductility of Ti-V-Zr medium entropy alloy by fracture mode
transformation via a heterogeneous structure, Int. J. Plast. 160 (2023) 103505.
[24] W. Huang, J. Hou, X. Wang, J. Qiao, Y. Wu, Excellent room-temperature tensile
ductility in as-cast Ti37V15Nb22Hf23W3 refractory high entropy alloys,
Intermetallics 151 (2022) 107735.
[25] D. Cui, Z. Yang, B. Guo, L. Liu, Z. Wang, J. Li, J. Wang, F. He, Microstructures and
mechanical properties of a precipitation hardened refractory multi-principal
element alloy, Intermetallics 151 (2022) 107727.
[26] Z.H. Han, L.K. Meng, J. Yang, G. Liu, J.G. Yang, R. Wei, G.J. Zhang, Novel BCC
VNbTa refractory multi-element alloys with superior tensile properties, Mater. Sci.
Eng. A. 825 (2021) 4.
[27] O.N. Senkov, D.B. Miracle, S.I. Rao, Correlations to improve room temperature
ductility of refractory complex concentrated alloys, Mater. Sci. Eng. A. 820 (2021)
141512.
[28] S.B. Wang, D.L. Wu, H. She, M.X. Wu, D. Shu, A.P. Dong, H.C. Lai, B.D. Sun, Design
of high-ductile medium entropy alloys for dental implants, Mater. Sci. Eng. C-
Mater. Biol. Appl. 113 (2020) 10.
[29] T.D. Huang, S.Y. Wu, H. Jiang, Y.P. Lu, T.M. Wang, T.J. Li, Effect of Ti content on
microstructure and properties of TixZrVNb refractory high-entropy alloys, Int. J.
Miner. Metall. Mater. 27 (10) (2020) 1318–1325.
[30] Z.B. An, S.C. Mao, T. Yang, C.T. Liu, B. Zhang, E. Ma, H. Zhou, Z. Zhang, L.
H. Wang, X.D. Han, Spinodal-modulated solid solution delivers a strong and ductile
refractory high-entropy alloy, Mater. Horizons 8 (3) (2021) 948–955.
[31] S. Laube, A. Kauffmann, S. Schellert, S. Seils, A.S. Tirunilai, C. Greiner, Y.
M. Eggeler, B. Gorr, H.-J. Christ, M. Heilmaier, Formation and thermal stability of
two-phase microstructures in Al-containing refractory compositionally complex
alloys, Sci. Technol. Adv. Mater. 23 (1) (2022) 692–706.
[32] S.Y. Wang, F.C. Lang, Y.M. Xing, Geometric phase analysis for characterization of
3D morphology of carbon ber reinforced composites, Compos. Sci. Technol. 242
(2023) 110215.
[33] M.J. Hÿtch, E. Snoeck, R. Kilaas, Quantitative measurement of displacement and
strain elds from HREM micrographs, Ultramicroscopy 74 (3) (1998) 131–146.
[34] F. Findik, Improvements in spinodal alloys from past to present, Mater. Des. 42
(2012) 131–146.
[35] B.L. Yin, F. Maresca, W.A. Curtin, Vanadium is an optimal element for
strengthening in both fcc and bcc high-entropy alloys, Acta. Mater. 188 (2020)
486–491.
D. Cui et al.
Scripta Materialia 244 (2024) 116031
7
[36] M. Kato, Hardening by spinodally modulated structure in b.c.c. alloys, Acta. Metall.
29 (1) (1981) 79–87.
[37] T. Xin, Y. Zhao, R. Mahjoub, J. Jiang, A. Yadav, K. Nomoto, R. Niu, S. Tang, F. Ji,
Z. Quadir, D. Miskovic, J. Daniels, W. Xu, X. Liao, L.-Q. Chen, K. Hagihara, X. Li,
S. Ringer, M. Ferry, Ultrahigh specic strength in a magnesium alloy strengthened
by spinodal decomposition, Sci. Adv. 7 (23) (2021) eabf3039.
[38] Y.Q. Bu, Y. Wu, Z.F. Lei, X.Y. Yuan, H.H. Wu, X.B. Feng, J.B. Liu, J. Ding, Y. Lu, H.
T. Wang, Z.P. Lu, W. Yang, Local chemical uctuation mediated ductility in body-
centered-cubic high-entropy alloys, Mater. Today. 46 (2021) 28–34.
[39] Q. Wu, F. He, J. Li, H.S. Kim, Z. Wang, J. Wang, Phase-selective recrystallization
makes eutectic high-entropy alloys ultra-ductile, Nat. Commun. 13 (1) (2022)
4697.
[40] H. Yang, K. Li, Y. Bu, J. Wu, Y. Fang, L. Meng, J. Liu, H. Wang, Nanoprecipitates
induced dislocation pinning and multiplication strategy for designing high
strength, plasticity and conductivity Cu alloys, Scr. Mater. 195 (2021) 113741.
[41] I.A. Su, K.K. Tseng, J.W. Yeh, B. El-Sayed, C.H. Liu, S.H. Wang, Strengthening
mechanisms and microstructural evolution of ductile refractory medium-entropy
alloy Hf20Nb10Ti35Zr35, Scr. Mater. 206 (2022) 6.
[42] Y. Lu, Y.H. Zhang, E. Ma, W.Z. Han, Relative mobility of screw versus edge
dislocations controls the ductile-to-brittle transition in metals, Proc. Natl. Acad.
Sci. U. S. A. 118 (37) (2021) 6.
[43] Q.Q. Ding, Y. Zhang, X. Chen, X.Q. Fu, D.K. Chen, S.J. Chen, L. Gu, F. Wei, H.B. Bei,
Y.F. Gao, M.R. Wen, J.X. Li, Z. Zhang, T. Zhu, R.O. Ritchie, Q. Yu, Tuning element
distribution, structure and properties by composition in high-entropy alloys,
Nature 574 (7777) (2019) 223–227.
[44] S. Zhao, R. Zhang, Y. Chong, X. Li, A. Abu-Odeh, E. Rothchild, D.C. Chrzan,
M. Asta, J.W. Morris, A.M. Minor, Defect reconguration in a Ti–Al alloy via
electroplasticity, Nat. Mater. 20 (4) (2021) 468–472.
[45] T. Li, S. Wang, W. Fan, Y. Lu, T. Wang, T. Li, P.K. Liaw, CALPHAD-aided design for
superior thermal stability and mechanical behavior in a TiZrHfNb refractory high-
entropy alloy, Acta. Mater. 246 (2023) 118728.
[46] M. Stricker, D. Weygand, Dislocation multiplication mechanisms – Glissile
junctions and their role on the plastic deformation at the microscale, Acta. Mater.
99 (2015) 130–139.
[47] F. He, S. Wei, J.L. Cann, Z. Wang, J. Wang, C.C. Tasan, Composition-dependent slip
planarity in mechanically-stable face centered cubic complex concentrated alloys
and its mechanical effects, Acta. Mater. 220 (2021) 117314.
[48] Z.F. Lei, X.J. Liu, Y. Wu, H. Wang, S.H. Jiang, S.D. Wang, X.D. Hui, Y.D. Wu,
B. Gault, P. Kontis, D. Raabe, L. Gu, Q.H. Zhang, H.W. Chen, H.T. Wang, J.B. Liu,
K. An, Q.S. Zeng, T.G. Nieh, Z.P. Lu, Enhanced strength and ductility in a high-
entropy alloy via ordered oxygen complexes, Nature 563 (7732) (2018) 546–550.
[49] R. Wang, Y. Tang, Z. Lei, Y. Ai, Z. Tong, S. Li, Y. Ye, S. Bai, Achieving high strength
and ductility in nitrogen-doped refractory high-entropy alloys, Mater. Des. 213
(2022) 110356.
D. Cui et al.
