Nanoparticles-strengthened high-entropy alloys for cryogenic applications showing an exceptional strength-ductility synergy

Abstract

We designed a novel nanoparticles-strengthened high-entropy alloy, Ni30Co30Fe13Cr15Al6Ti6, which exhibits excellent strength-ductility combinations at both ambient and cryogenic temperatures. Especially at 77 K, an exceptional strength-ductility synergy can be observed, showing an ultrahigh tensile strength of 1.7 GPa and a large ductility of 51%, accompanied by a distinctive three-stage strain-hardening response. The precipitation-hardening behaviors and deformation micro-mechanisms were carefully investigated by the atom probe tomography and transmission electron microscope. The dynamic formation of nano-spaced stacking faults contributed to improved strain-hardening capacities, resulting in simultaneous enhancements of plastic deformation stability and tensile ductility at such high-strength levels.

Full text

Nanoparticles-strengthened high-entropy alloys for cryogenic
applications showing an exceptional strength-ductility synergy
T. Yang
a,b
,Y.L.Zhao
c,d
, J.H. Luan
a,b
,B.Han
c,d
,J.Wei
c,d
, J.J. Kai
c,d
,C.T.Liu
a,b,c,
⁎
a
Department of Materials Science and Engineering, City University of Hong Kong, Hong Kong, China
b
Center for Advanced Structural Materials, City University of Hong Kong, Hong Kong, China
c
Department of Mechanical Engineering, City University of Hong Kong, Hong Kong, China
d
Center for Advanced Nuclear Safety and Sustainable Development, City University of Hong Kong, Hong Kong, China
abstractarticle info
Article history:
Received 16 November 2018
Accepted 19 January 2019
Available online xxxx
We designed a novel nanoparticles-strengthened high-entropy alloy, Ni
30
Co
30
Fe
13
Cr
15
Al
6
Ti
6
, which exhibits ex-
cellentstrength-ductility combinations at both ambientand cryogenic temperatures.Especially at 77K, an excep-
tional strength-ductility synergy can be observed, showing an ultrahigh tensile strength of 1.7GPa and a large
ductility of 51%, accompanied by a distinctive three-stage strain-hardening response. The precipitation-
hardening behaviors and deformation micro-mechanisms were carefully investigated by the atom probetomog-
raphy and transmission electron microscope. The dynamic formation of nano-spaced stacking faults contributed
to improvedstrain-hardening capacities, resulting in simultaneous enhancements of plasticdeformation stability
and tensile ductility at such high-strength levels.
© 2019 Published by Elsevier Ltd on behalf of Acta Materialia Inc.
Keywords:
High-entropy alloy
Nanoparticle strengthening
Mechanical properties
Stacking-fault-induced plasticity
Metallic materials with superior mechanical properties at both am-
bient and cryogenic temperatures are highly demanded for modern en-
gineering applications, especially in extreme service conditions like
aerospace, marine shipbuilding, and natural gas industries. Recently,
multi-component high-entropy alloys (HEAs) have attracted intensive
interests in the materials field [1–3]. Enlarged compositional choices
in HEA systems offer us a great space to tailor phase structures and de-
formation mechanisms of alloys for achieving desired strength-ductility
combinations [4–6]. Among them, the single-phase HEAs, especially
those with the face-centered-cubic (fcc) structure and low stacking
fault energies (SFEs), have been actively studied, exhibiting unique mi-
crostructural and mechanical responses at cryogenic temperatures. Me-
chanical twinning and transformation-induced martensite have been
demonstrated as the dominated deformation modes, which act to en-
hance the strain-hardening capacity of alloys and delay the onset of
local necking to higher plastic strains [7–10]. A striking example is the
equiatomic CoCrFeNiMn HEA, which displays a remarkable tensile duc-
tility (~70%) and fracture toughness (~200 MPa·m
1/2
)at77K[7]. The
observed mechanical performance can be comparable to the best cryo-
genic alloys like austenitic stainless steels. However, these single-
phased HEAs generally exhibit insufficient yield strengths at room tem-
perature, which significantly reduce their engineering practicality and
applicability [7,8].
As compared to single-fcc HEAs, precipitation-hardening HEAs
strengthened by coherent L1
2
-type nanoparticles show a great potential
to achieve enhanced mechanical performance at room temperature, es-
pecially the higher yield strengths with a considerable tensile ductility
[11,12]. Also, given the fcc-type crystalline structure of both the matrix
and precipitates, excellent cryogenic performance could be expected in
these L1
2
-strengthed HEAs. However, present studies of these alloys
mostly focus on their mechanical properties at room and elevated tem-
peratures, while the cryogenic behaviors are rarely investigated
[13–15]. The strain-hardening behaviors and underlying deformation
micro-mechanisms at cryogenic temperatures are also not well under-
stood. In addition, although some advances have been achieved, the un-
desired precipitation of brittle intermetallic phases (such as the Heuser
phase [16] and eta phase [17]) can be frequently observed in many
existing L1
2
-strengthed HEAs, which inevitably damage their mechani-
cal properties to a certain extent.
In this study, we designed a novel precipitation-strengthened HEA,
which exhibits a pure “fcc + L1
2
”dual-phase structure without involv-
ing any brittle phases. We first carefully evaluated the Ni-Co-Fe-Cr-Al-Ti
(at. %) system based on the CALPHAD (CALculation of PHAse Diagrams)
technique. Figs. 1(a) and 1(b) present the equilibrium phase diagrams
calculated by the Thermo-Cal software with a TTNI8 database, and the
rationale for selecting this non-equiatomic Ni
30
Co
30
Fe
13
Cr
15
Al
6
Ti
6
(at.
%) HEA can be primarily described as follows: (1) 6 at. % of Al and Ti
were added to catalyze the formation of high-density L1
2
nanoparticles
for precipitation hardening, without the formation of BCC-type inter-
metallic phases; (2) the concentrations of Ni and Co were controlled
Scripta Materialia 164 (2019) 30–35
⁎Corresponding author at: Department of Materials Science and Engineering, City
University of Hong Kong, Hong Kong, China.
E-mail address: chainliu@cityu.edu.hk (C.T. Liu).
https://doi.org/10.1016/j.scriptamat.2019.01.034
1359-6462/© 2019 Published by Elsevier Ltd on behalf of Acta Materialia Inc.
Contents lists available at ScienceDirect
Scripta Materialia
journal homepage: www.elsevier.com/locate/scriptamat

to be 30 at. % to stabilize the fcc matrix. (3) Cr was alloyed to decrease
the SFEs of the fcc matrix due to its strong preference to partitioning
into the matrix [8,18,19]. The Cr concentration was elaborately tailored
to be about 15 at. %, since the excessive addition of which will signifi-
cantly promote the formation of brittle σphase, accompanied by a
drastic reduction of the fcc matrix (Fig. 1(b)). Subsequently, the micro-
structure evolutions, mechanical properties and associated deformation
micro-mechanisms at both 293 K and 77 K were systematically
investigated.
To start with, alloy ingots with a nominal composition of
Ni
30
Co
30
Fe
13
Cr
15
Al
6
Ti
6
(at. %) were produced by arc-melting in a Ti-
gettered argon atmosphere. The as-cast ingots were first solution-
treated at1150 °Cfor 2 h and then cold rolled along the longitudinal di-
rection with a total thickness reduction of ~65%, followed by a recrystal-
lization process at 1150 °C for 2 min. The resulting sheets were
subsequently annealed at 800 °C for 4 h to introduce precipitation. All
heat treatments were finished by oil quenching. Phase identification
was conducted on a Rigaku X-ray diffractometer (XRD) equipped with
a monochromator. Microstructure characterizations were carried out
using the scanning electron microscope (SEM, Quanta FEG450) and
transmission electron microscope (TEM, JEOL 2100F). The TEM speci-
mens were firstly mechanically ground to a thickness of ~50 μm,
followed by a twin-jet electro-polishing. Electron backscattered diffrac-
tion (EBSD) characterizations were analyzed by the TSL OIM data-
collection software. The elemental distributions were investigated
using the energy-dispersive X-ray spectroscopy (EDX) and atom
probe tomography (APT). Needle-shaped specimens required for APT
were prepared by lift-outs and then annular milled in a FEI Scios focused
ion beam/scanning electron microscope (FIB/SEM). The APT character-
izations were performed in a local electrode atom probe (CAMEACA
LEAP 5000 XR). The APT specimens were analyzed at 70 K in the voltage
mode, with a pulse repetition rate of 200 KHz, a pulse fraction of 20%,
and an evaporation detection rate of 0.2% atom per pulse. AnIntegrated
Visualization and Analysis Software (IVAS) protocol was employed
to reconstruct the 3D atomic maps. Flat dog-bone-shaped specimens
with a gauge length of 12.5 mm, a width of 3.2 mm and a thickness
of ∼1.5 mm were fabricated by electro-discharge machining for
tensile tests. The uniaxial tensile tests were performed on a Material
Testing System (MTS, Alliance RT30) tension machine with a strain
rate of 1 × 10
−3
s
−1
.Toconfirm the data reproducibility, three tensile
Fig. 1. (a)Equilibrium phasediagram of the Ni
30
Co
30
Al
6
Ti
6
Fe
28-x
Cr
x
(at. %) alloysystem, highlighting the alloying effectof Cr element. (b) The phasestructure evolutionsas a function of the
Cr content at various temperatures. (c) EBSD mapping indicating the polycrystalline structure containing micron-sized equiaxed grains. (d) X-ray diffraction pattern of the aged HEA
showing the targeted “FCC + L1
2
”dual-phase microstructures. (e) TEM image showing the morphology and distribution of L1
2
nanoparticles. (f) HR-TEM image demonstrating the
highly coherent interfaces between the L1
2
nanoparticles and FCC matrix.
31T. Yang et al. / Scripta Materialia 164 (2019) 30–35

specimens for each condition were tested. Fracture surfaces were exam-
ined by scanning electron microscopy. The deformation mechanisms
were investigated by TEM at the gauge length section of the deformed
specimens with different strains.
Typical microstructures of the Ni
30
Co
30
Fe
13
Cr
15
Al
6
Ti
6
(at. %) HEA are
presented in Fig. 1(c)-(f). The EBSD image reveals a homogenous and
random distribution of grains with an average size of 43 ± 9 μm. Fig. 1
(d) shows the corresponding XRD pattern, exhibiting a pristine “fcc
+L1
2
”dual-phase microstructure without any detections of other in-
termetallic phases. Fig. 1(e) shows the representative dark-field TEM
(DF-TEM) image of the L1
2
-type nanoparticles and the corresponding
selected-area-diffraction (SAD) pattern. High-density near-spherical
L1
2
nanoparticles with an average diameter of 29 ± 5 nm were uni-
formly distributed in thefcc matrix. Subsequently, the nanoparticle/ma-
trix interfacial feature was identified by the high-resolution TEM (HR-
TEM), as shown in Fig. 1(f). It is evident that the L12 nanoparticles are
highly coherent with the fcc matrix. The lattice mismatch is small,
which was measured to be about 0.16% based on the XRD analyses.
Elemental distributions of the present HEA were investigated by the
combined analyses of TEM-EDX and 3D-APT. The TEM-EDX mapping
presented in Fig. 2(a) clearly reveals that the L1
2
-type nanoparticles
are enriched with Ni, Ti and Al while depleted with Cr, Co and Fe
atoms. To give a more quantitative compositional analysis, nano-sized
needle tips were carefully lifted out from this alloy and then further
characterized by the APT. Fig. 2(b) shows the typical atom maps of Fe,
Co, Cr and Ni, Aland Ti elements, exhibiting a similar elemental distribu-
tion to that obtained by the TEM-EDX. The chemical composition of
each individual phase was quantitatively analyzed by a mean of proxim-
ity histogram calculated from a 10 at. % Ti iso-concentration surface, and
the detailed results were summarized in Table 1. Evidently, the Ni, Ti
and Al atoms strongly partition into the L1
2
nanoparticles, whereas
the Co atom (~20 at. %) is marginally depleted and the Cr, Fe atoms
(~6 at. % in total) are largely depleted. An opposite characteristic can
be observed in the fcc matrix. Given the stoichiometry of L1
2
phase
with the A
3
B structure, these nanoparticles can be identified as the
multi-component (Ni,Co,Fe,Cr)
3
(Ti,Al) phase. The Co and Fe atoms gen-
erally substitute for the A site in the L1
2
nanoparticles, which can stabi-
lize the L1
2
crystallize structure by decreasing its valence electron
concentration [20]. Meanwhile, they can also help to increase the intrin-
sic ductility of the L1
2
-type nanoparticle by decrease the ordering en-
ergy [21], leading to an enhanced damage tolerance of the resultant
precipitation-hardening HEA. Furthermore, based on the lever rule,
the volume fraction of L1
2
nanoparticles can be determined to be
~37.3%, which is significantly higher than of the previously reported
equiatomic FeCoNiCr- and CoCrNi-based alloys [16,22]. The predicted
results based on the thermodynamic calculations were also presented
here for a direct comparison, which shows a small deviation between
the APT results, demonstrating the high reliability of CALPHAD tech-
nique for the alloy design of L1
2
-strengthened HEAs.
Fig. 3(a) shows the excellent mechanical properties of our present
nanoparticles-strengthened HEA at 293 K and 77 K. A high yield
strength (YS) of ~925 MPa and ultimate tensile strength (UTS) of
~1.31 GPa together with a large tensile elongation of ~43% can be ob-
tained at 293 K. More strikingly, at 77 K, a surprising strength-ductility
synergy can be observed, exhibiting a higher UTS, up to ~1.7 GPa, and
a combined increase in the ductility to ~51%. Numerous dimples can
be observed in the fracture surface with almost no macroscopic necking
(inset in Fig. 3(a)), confirming the ductile nature and excellent plastic
stability of this alloy at 77 K. In comparison with other existing cryo-
genic alloys [8,10,23–30], as shown in Fig. 3(c) and 3(d), it is evident
Fig. 2. (a) STEM-EDX mapping of the present HEA, showing the elemental distribution between the L1
2
nanoparticles and FCC matrix. (b) Atom maps of Fe, Co, Ni, Cr, Al andTi elements.
(c) Proximity histogram constructed across the interfaces between the matrix and nanoparticles, giving a quantitative compositional analysis.
32 T. Yang et al. / Scripta Materialia 164 (2019) 30–35

that our present HEA shows superior strength-ductility combinations at
both 77 K and 293 K, making it a more attractive candidate for many
critical engineering applications. The high yield strength can be mainly
ascribed to the pronounced ordering hardening induced be the high-
density L1
2
nanoparticles, which can be estimated by Δσ
order
=0.81M
(γ
APB
/2b)(3πf/8)
1/2
[11,22,31], in which Mis the Taylor factor (3.06 for
the fcc polycrystalline matrix), γ
APB
is the antiphase boundary energy
of nanoparticles (estimated to be about 200 mJ·m
2
[32]), bis the mag-
nitude of the Burgers vector of the matrix (~0.254 nm) and fis the vol-
ume fraction of the precipitate (0.373). Based on the above calculations,
the contribution of ordering hardening to the yield strength can be de-
termined to be ~646 MPa, which accounts for a high proportion (~70%)
of its total yield strength. Moreover, the strength contribution from the
lattice friction of the fcc matrix can be roughly estimated to be
200–300 MPa [8]. Assuming additivity of these two parts, the total
yield strength amounts to 846–946 MPa, which generally agrees with
the above experimental result. We attributed the excellent ductility
at such high yield strengths to their exceptional strain-hardening
capabilities, manifested by distinctive three-stage strain-hardening be-
haviors (Fig. 3(b)). Clearly, the abnormal increase of strain-hardening
rate in Stage B makes a crucial contribution to sustain the plastic stabil-
ity for a large ductility.
To unravel the deformation micro-mechanism underlying such im-
pressive mechanical properties, detailed TEM characterizations were
conducted. Fig. 4(a) shows the deformation microstructure of the sam-
ple stretched to ~15% true strain (Stage B) at 293 K. The deformation
feature at this stage is characterized by parallel sets of stacking faults
(SFs) along the (111) slip planes with an average spacing size of
~100.7 nm. As the deformation proceeds to ~30% true strain (Stage C),
as shown in Fig. 4(b), SFs multiplied and intersected significantly, pro-
ducing the nano-spaced SF networks. Upon deformed at 77 K, a larger
number of SFs were triggered and proliferated rapidly during theplastic
deformation process. An earlier onset of the SF networks can be acti-
vated within the Stage B (see Fig. 4(c)). With further increasing the
strain, these SF architectures were increasingly refined with a greatlyre-
duced spacing size. Specifically, as presented in Fig. 4(d), the mean
Table 1
Chemical compositions and phase structures of the present HEA based on the 3D-APT analyses and CALPHAD predictions.
a
Phases Compositions (at. %) V
exp
V
cal
Fe Co Ni Cr Al Ti
FCC matrix 3D-APT 19.69 ± 0.14 34.25 ± 0.17 17.86 ± 0.13 23.14 ± 0.15 3.63 ± 0.07 1.43 ± 0.04 37.3% 40%
CALPHAD 19.74 34.98 17.90 23.82 2.69 0.87
L1
2
nanoparticle 3D-APT 3.36 ± 0.09 19.78 ± 0.21 50.79 ± 0.26 2.53 ± 0.08 10.17 ± 0.16 13.37 ± 0.18
CALPHAD 3.02 22.63 48.16 1.70 10.89 13.58
a
V
exp
and V
cal
correspond to the mole fraction of the L1
2
nanoparticles determined by experimental observation and thermodynamic calculation, respectively.
Fig. 3. (a)Engineering tensilecurves of our present HEA at ambient and cryogenic temperatures. Insets show the typical fracturesurfaces of the present HEAdeformed at 77 K, exhibiting
ductile dimpled microstructures with almost no macroscopic necking. (b) Strain-hardening responses of our present HEA. Comparison of the mechanical properties of our present HEA
with other selected cryogenic alloys [8,10,23–30] obtained at (c) 77 K and (d) 293 K, respectively. Alloys with strong texture and heterogeneous grain structures are not included here
due to the serious anisotropy of properties.
33T. Yang et al. / Scripta Materialia 164 (2019) 30–35

spacing between adjacent SFs was decreased to about 14.2 nm. HR-TEM
image further gives a close-up view of the nano-spaced SFs. Despite the
abundant SFs, mechanical twins, which have been frequently identified
in many single-fcc medium- and high-entropy alloys (MEAs and HEAs),
are scarcely observed, as shown in Fig. 4(f). It has been well accepted
that the single-fcc solid-solution alloys with low stacking faults energy
(SFE) (about 10–40 mJ·m
−2
) show a high propensity of deformation-
induced twins [9,33–35]. According to the APT compositional analysis,
large amounts of Cr and Co elements were partitioned into the fcc ma-
trix, the SFE of which can be qualitatively evaluated to be about
16 mJ·m
−2
based on the work by Xie et al. [19]. It is thus expected
that a similar twinning behavior would be generated. However, as
described above, high-density SFs instead of deformation twins were
detected, even under a large strain level at 77 K. It indicates that the
twinning formation was significantly inhibited in our present
precipitation-hardened alloy. One major reason for this phenomenon
could be ascribed to a high SFE of the L1
2
-type nanoparticles. Based
on previous studies of Breidi et al., the SFE of the Ni
3
(Al,Ti)-type nano-
particles of our present alloy can be estimated to be about
200–250 mJ·m
−2
[36], which is substantially higher than that of the
fcc matrix.As a result, the incorporation of thesenanoparticles will dra-
matically increase the critical stress for twinning nucleation. Further-
more, it has been also reported that smaller grain size will impede the
twinningformation. For our present alloy, with the high-content precip-
itation of nanoparticles, the width of fcc matrix channel was narrowed
down to about 14.5 nm, leading to a greatly increased stress level to ini-
tiate the transformation from SFs to deformation twins [22,37].
Based on above analyses, it is thus concluded that the nano-spaced
SFs, instead the deformation twins, play the dominated role in sustain-
ing the strain hardening of our present HEA. Previous studies revealed
that the strain-hardening efficiency of SFs is highly dependent on their
spacing sizes [38,39]. The dynamic formation and refinement of these
SF networks can prevent dislocations moving by decreasing the mean
free path of dislocations. On the other hand, these nano-spaced SF net-
works would also contribute to an increased work-hardening response
by the so-called dynamic Hall-Petch effect, which is similar to the
twinning-induced plasticity, namely the TWIP effect [33,40]. Further-
more, the mutual interactions between the co-planar SFs would lead
to the extensive formation of sessile dislocations, generally known as
the immobile Lomer-Cottrell locks (L-C locks) [41,42], which can also
act as strong barriers to other mobile dislocations and promote the dis-
location accumulation, leading to a steady and continuous plastic defor-
mation to large strains at the cryogenic temperature.
In summary, we successfully developed a high-performing
nanoparticle-strengthened Ni
30
Co
30
Fe
13
Cr
15
Al
6
Ti
6
(at. %) HEA by using
a computational-aided CALPHAD method. Based on the combined ex-
perimental analyses, the elemental distributions, microstructures and
the associated mechanical behaviors at both room and cryogenic tem-
peratures have been carefully studied. A pristine “fcc + L1
2
”dual-
phase nano-composite structure was clearly identified in this HEA
Fig. 4. TEM micrographs showing the microstructural evolutions of the present HEA deformed to: (a) ~15% true strain at 293 K, (b) ~30% true strain at 293 K, (c) ~20%truestrainat77K,
and (d) ~40% true strain at 77 K, respectively. (e) Representative HR-TEM images of the deformation-induced stacking faults. (f) TEM image and corresponding SAD pattern of the
mechanical nano-twins that occasionally observed in a local region.
34 T. Yang et al. / Scripta Materialia 164 (2019) 30–35

without other brittle phases.As compared to theroom-temperature de-
formation, an exceptional strength-ductility synergy can be obtained at
77 K, showing an ultrahigh tensile strength of 1.7 GPa combined with a
large elongation of 51%. The three-stage strain-hardening behaviors
were carefully analyzed by TEM. Distinct from the twinning and
transformation-induced plasticity in previous HEAs and MEAs, the
nano-spaced stacking faults are demonstrated as a new effective tough-
ening mechanism for achieving large plastic strains in high-strength
precipitation-hardening materials. More importantly, our present find-
ings are expected to open up a new area for the future development of
nanoparticles-strengthened HEAs for advanced engineering applica-
tions, especially for the extreme conditions with low temperatures.
Acknowledgments
This work was supported by the Hong Kong Research Grant Council,
University Grants Committee (RGC) for financial support (CityU
11209314 and CityU 11205515). The APT research was supported by
the Collaborative Research Fund (CityU 11205515) from the RGC of
Hong Kong.
References
[1] J.W. Yeh, S.K. Chen, S.J. Lin, J.Y. Gan, T.S. Chin, T.T. Shun, C.H. Tsau, S.Y. Chang, Adv.
Eng. Mater. 6 (2004) 299–303.
[2] B. Cantor, Entropy 16 (2014) 4749–4768.
[3] Y. Zhang,T.T. Zuo, Z. Tang, M.C. Gao, K.A. Dahmen, P.K. Liaw, Z.P. Lu, Prog. Mater. Sci.
61 (2014) 1–93.
[4] I. Baker,H. Wu, X. Wu, M.K. Miller, P.R. Munroe, Mater. Charact. 62 (2011) 952–958.
[5] Z. Li, K.G. Pradeep, Y. Deng, D. Raabe, C.C. Tasan, Nature 534 (2016) 227–230.
[6] O.N. Senkov, J.D. Miller, D.B. Miracle, C. Woodward, Nat. Commun. 6 (2015) 6529.
[7] B. Gludovatz, A. Hohenwarter, D. Catoor, E.H. Chang, E.P. George, R.O. Ritchie, Sci-
ence 345 (2014) 1153–1158.
[8] Z. Wu, H. Bei, G.M. Pharr, E.P. George, Acta Mater. 81 (2014) 428–441.
[9] S. Huang, H. Huang, W. Li, D. Kim, S. Lu, X. Li, E. Holmstrom,S.K. Kwon, L. Vitos, Nat.
Commun. 9 (2018) 2381.
[10] J.W. Bae, J.B. Seol, J. Moon, S.S. Sohn, M.J. Jang, H.Y. Um, B.-J. Lee, H.S. Kim, Acta
Mater. 161 (2018) 388–399.
[11] J.Y. He, H. Wang, H.L. Huang, X.D.Xu, M.W. Chen, Y. Wu, X.J. Liu, T.G. Nieh, K. An, Z.P.
Lu, Acta Mater. 102 (2016) 187–196.
[12] T. Yang, Y. Zhao, W. Liu, J. Kai, C. Liu, J. Mater. Res. 33 (2018) 2983–2997.
[13] Y.-J. Chang, A.-C. Yeh, J. Alloys Compd. 653 (2015) 379–385.
[14] J.Y. He, H. Wang, Y. Wu, X.J. Liu, T.G. Nieh, Z.P. Lu, Mater. Sci. Eng. A 686 (2017)
34–40.
[15] T.K. Tsao, Y.J. Chang, K.C. Chang, J.W. Yeh, M.S. Chiou, S.R. Jian, C.M. Kuo, W.R. Wang,
H. Murakami, Int. J. Metall. Mater. Eng. 1 (2015).
[16] J.Y. He, H. Wang, Y. Wu, X.J. Liu,H.H. Mao, T.G. Nieh, Z.P. Lu, Intermetallics79 (2016)
41–52.
[17] Y.L. Zhao, T. Yang, J.H. Zhu, D. Chen, Y. Yang, A. Hu, C.T. Liu, J.J. Kai, Scr. Mater. 148
(2018) 51–55.
[18] T.-K. Tsao, A.-C. Yeh, H. Murakami, Metall. Mater. Trans. A 48 (2017) 2435–2442.
[19] X. Xie, G. Chen, P. McHugh, J. Tien, Scr. Metall. 16 (1982) 483–488.
[20] T. Yang, Y. Zhao, W. Liu, J. Zhu, J. Kai, C. Liu, Math. Res. Lett. 6 (2018) 600–606.
[21] A. Chiba, S. Hanada, S. Watanabe, Mater. Sci. Eng. A 152 (1992) 108–113.
[22] Y.L. Zhao, T. Yang, Y. Tong, J. Wang, J.H. Luan, Z.B. Jiao, D. Chen, Y. Yang, A. Hu, C.T.
Liu, J.J. Kai, Acta Mater. 138 (2017) 72–82.
[23] T. Bhattacharjee, R. Zheng, Y. Chong, S. Sheikh, S. Guo, I.T. Clark, T. Okawa, I.S. Wani,
P.P. Bhattacharjee, A. Shibata, N. Tsuji, Mater. Chem. Phys. 210 (2018) 207–212.
[24] P. Czarkowski, A.T. Krawczynska, R. Slesinski, T. Brynk, J. Budniak, M. Lewandowska,
K.J. Kurzydlowski, Fusion Eng. Des. 86 (2011) 2517–2521.
[25] D. Li, C.Li, T. Feng, Y. Zhang,G. Sha, J.J. Lewandowski, P.K.Liaw, Y. Zhang, ActaMater.
123 (2017) 285–294.
[26] D. Li, Y. Zhang, Intermetallics 70 (2016) 24–28.
[27] L.F. Li 614 (2002) 165–169.
[28] C. Zheng, W. Yu, Mater. Sci. Eng. A 710 (2018) 359–365.
[29] S.S.Sohn, S. Hong, J. Lee, B.-C. Suh, S.-K. Kim, B.-J.Lee, N.J. Kim, S. Lee,Acta Mater. 100
(2015) 39–52.
[30] H.-S. Shin, H.-M. Lee, M.-S. Kim, Int. J. Impact Eng. 24 (2000) 571–581.
[31] T. Gladman, Mater. Sci. Technol. 15 (1999) 30–36.
[32] M. Vittori, A. Mignone, Mater. Sci. Eng. 74 (1985) 29–37.
[33] I. Gutierrez-Urrutia, D. Raabe, Acta Mater. 59 (2011) 6449–6462.
[34] S. Zhao, G.M. Stocks, Y. Zhang, Acta Mater. 134 (2017) 334–345.
[35] J. Kim, B.C. De Cooman, Metall. Mater. Trans. A 42 (2011) 932–936.
[36] A. Breidi, J. Allen, A. Mottura, Acta Mater. 145 (2018) 97–108.
[37] L. Murr, A. Ayala, C. Niou, Mater. Sci. Eng. A 216 (1996) 69–79.
[38] Y.Z. Tian, L.J. Zhao, S. Chen, A. Shibata, Z.F. Zhang, N. Tsuji, Sci. Rep. 5 (2015) 16707.
[39] W. Jian, G. Cheng, W. Xu, H. Yuan, M. Tsai, Q. Wang, C. Koch, Y. Zhu, S. Mathaudhu,
Math. Res. Lett. 1 (2013) 61–66.
[40] Z. Wu, C.M. Parish, H. Bei, J. Alloys Compd. 647 (2015) 815–822.
[41] X.L. Wu, Y.T. Zhu, Y.G. Wei, Q. Wei, Phys. Rev. Lett. 103 (2009) 205504.
[42] X. Xu, P. Liu, Z. Tang, A. Hirata, S. Song, T. Nieh, P. Liaw, C. Liu, M. Chen, Acta Mater.
144 (2018) 107–115.
35T. Yang et al. / Scripta Materialia 164 (2019) 30–35