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mater.scichina.com link.springer.com Published online 29 July 2020 | https://doi.org/10.1007/s40843-020-1377-2
Size-dependent deformation behavior of dual-phase,
nanostructured CrCoNi medium-entropy alloy
Yujie Chen1,2,3, Xianghai An3*, Zhifeng Zhou4, Paul Munroe5, Sam Zhang1*, Xiaozhou Liao3and
Zonghan Xie2,6
ABSTRACT The mechanical size effect of nanostructured,
dual-phase CrCoNi medium-entropy alloy (MEA) was in-
vestigated by combining in-situ micro-compression testing
with post-mortem electron microscopy analysis. The alloy
possesses a superior yield strength up to ~4 GPa, primarily
due to its hierarchical microstructure including column na-
nograins, preferred orientation, a high density of planar de-
fects and the presence of the hexagonal close packed (HCP)
phase. While the yield strength of the alloy has shown size-
independency, the deformation behaviour was strongly de-
pendent on the sample size. Specifically, with decreasing the
pillar diameters, the dominant deformation mode changed
from highly localized and catastrophic shear banding to ap-
parently homogeneous deformation with appreciable plasti-
city. This transition is believed to be governed by the size-
dependent critical stress required for a shear band traversing
the pillar and mediated by the competition between shear-
induced softening and subsequent hardening mechanisms. In
addition, an unexpected phase transformation from HCP to
face-centered cubic (FCC) was observed in the highly localized
deformation zones, leading to strain softening that con-
tributed to accommodating plasticity. These findings provide
insights into the criticality of sample dimensions in influen-
cing mechanical behaviors of nanostructured metallic mate-
rials used for nanoelectromechanical systems.
Keywords: medium-entropy alloy, size effect, shear banding,
phase transformation, nanostructure
INTRODUCTION
High- and medium-entropy alloys (HEAs and MEAs),
often referred to as multi-component alloys, have
emerged as a new class of metallic materials with excellent
physical and mechanical properties over the last decade
[1–4]. The original design concept of these novel alloys
was proposed to maximize the configurational entropy
for the formation of single-phase solid solution by mixing
multiple alloying elements having near equiatomic con-
centrations [5,6]. It is increasingly realized that the Gibbs
free energy in lieu of this type of entropy plays pre-
dominant roles in the phase formation and its stability,
enabling that the original strict restriction for the devel-
opment of HEA/MEAs is thermodynamically relaxed [7–
9]. This can tremendously increase compositional space
for this kind of novel alloys and catalyze novel mechan-
ism-driven HEAs/MEAs design. For instance, twinning-
induced plasticity (TWIP) and transformation-induced
plasticity (TRIP) effects can be introduced into these al-
loys through the optimization of their stacking fault en-
ergy (SFE) via tuning their chemical compositions and
concentrations, leading to the improved mechanical
properties [10–13]. It is well acknowledged that the
CrCoNi MEA with face-centered cubic (FCC) structure
exhibits an unprecedented combination of strength,
ductility, and toughness, as compared with the prototype
of the CrCoNiFeMn HEA [14]. Extensive investigation
proposed that these extraordinary mechanical properties
might be attributed to its lower SFE, enabling the ex-
tensive twinning activities and the development of hex-
agonal close packed (HCP) nanolaths in CrCoNi via an
FCC to HCP phase transformation [15–17].
More recently, the design principles are dedicated to
1Centre for Advanced Thin Films and Devices, School of Materials and Energy, Southwest University, Chongqing 400715, China
2School of Mechanical Engineering, The University of Adelaide, Adelaide, SA 5005, Australia
3School of Aerospace, Mechanical and Mechatronic Engineering, The University of Sydney, Sydney, NSW 2006, Australia
4Department of Mechanical Engineering, City University of Hong Kong, Kowloon, Hong Kong, China
5School of Materials Science and Engineering, The University of New South Wales, Sydney, NSW 2052, Australia
6School of Engineering, Edith Cowan University, Perth, WA 6027, Australia
*Corresponding authors (emails: xianghai.an@sydney.edu.au (An X); samzhang@swu.edu.cn (Zhang S))
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incorporating other strengthening mechanisms into this
kind of novel alloys, such as microstructure modification
[18–20], precipitation strengthening [21,22], ductile
multicomponent intermetallic nanoparticles [23], and
interstitial solid solution strengthening (by adding boron
[24], carbon [25], nitrogen [26,27] and oxygen [28]), to
push the property boundary of possibility apart from
their inherent substitutional solid solution strengthening.
Similarly, we successfully fabricated a nanostructured
dual-phase (HCP and FCC) CrCoNi MEA using mag-
netron sputtering, which possesses ultrahigh nanohard-
ness of ~10 GPa obtained by nanoindentation that is
much higher than that of conventional FCC structured
counterpart [29,30]. It was revealed that this ultrahigh
hardness originated from its unique hierarchical micro-
structures enabled by the special processes of magnetron
sputtering [29,30]. However, the detailed mechanical
behaviors of the newly fabricated dual-phase CrCoNi
MEA in terms of stress-strain curves, deformability and
underlying deformation mechanism have not been sys-
tematically explored and especially, the ductility achiev-
able under such high hardness remains mysterious, both
of which will be essentially uncovered in the current in-
vestigation.
As well known, the sample size can be a crucial factor in
regulating the mechanical behaviour of structural mate-
rials when the sample size is reduced to the micron scale
and below [31–33]. Notable observations include the in-
versely proportional relationship between strength and
sample sizes, i.e., “smaller is stronger”, which is generally
associated with appreciable reduction in structural defects
and dislocation starvation [32,33] caused by the inability
of sub-micron or nanoscale samples to store dislocations.
In addition, the deformation mode of single crystal pillars
is also size-dependent. For example, the deformation in
Al and Ni single crystal pillars was dominated by
homogeneous plastic flow at the micron scale, while large
strain bursts and dramatic collapse took place in sub-
micron pillars as a result of dislocation avalanche [31,34].
However, when the complex microstructures featuring
grain boundaries (GBs), planar defects and precipitates
are incorporated in the materials, the presence of these
structural defects in small-scale specimens may preclude
known mechanisms responsible for the size-dependent
plasticity of materials [31,32]. For instance, the size effect
on the mechanical behavior of alloys with a high density
of growth planar defects remains largely unexplored. In
light of the high performance of the MEAs and their
massive potential in the application of nano- or micro-
electromechanical systems (NEMS/MEMS) [35], it is
imperative to understand the mechanical response of
MEAs containing complex microstructures as a function
of sample size, which will not only enrich our under-
standing of size-related deformation of novel materials,
but also provide valuable guideline to the design and
development of reliable, high-performance NEMS/MEMS
[35].
In this study, a nanostructured, dual-phase CrCoNi
alloy containing a large number of growth planar defects
was prepared using magnetron sputtering. In-situ com-
pression testing of micro-pillars with diameters ranging
from 2 to 0.23 µm was performed. The size effect on the
mechanical behaviors of the alloy was systematically ex-
plored. Compared with conventional FCC structured
CrCoNi, this newly created alloy showed a superior yield
strength of ~4 GPa. While the yield strength of the alloy
was found to be essentially size-independent, the de-
formation mode was strongly dependent on the sample
size. As the diameter decreases, a transition in the de-
formation mode occurred in the alloy from brittle failure,
caused by unsteady large shear banding, to ‘ductile’ de-
formation (i.e., uniform deformation) accompanied by
minor shear events. Moreover, an unexpected HCP to
FCC phase transformation was identified, indicating the
HCP phase is unstable under applied stress.
EXPERIMENTAL SECTION
The CrCoNi alloy was deposited on AISI M2 steel sub-
strates (hardened to HRC 62, and polished to a final
surface roughness of 0.01 μm) using a direct current (DC)
magnetron sputtering system. A CrCoNi target (Cr:Co:
Ni=1:1:1, at%) with the purity of >99.9% and dimensions
of 345 mm×145 mm×5 mm was used for the alloy de-
position. The substrates were first ultrasonically cleaned
before mounting on a stationary holder. The target-to-
substrate distance was set at 170 mm. The background
chamber was pumped down to a pressure of 4×10−4Pa.
Prior to deposition, the substrates were ion cleaned for
30 min under a bias voltage of −450 V to remove any
contaminants on the substrate surfaces. The argon (Ar)
working gas (99.995% purity) was fed into the chamber at
a constant flow rate of 50 sccm by a MKS mass flow
controller and the chamber pressure was kept at 0.17 Pa
during sputtering. No external heating was used and a
bias voltage of −60 V was maintained during the coating
deposition. The CrCoNi DC was maintained at 4.0 A
(sputtering power of ~1.5 kW), corresponding to a
nominal deposition rate of ~72 nm min−1. The CrCoNi
film was deposited with a nominal thickness of 5 μm
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directly onto the steel substrates.
Micro-pillars for uniaxial micro-compression tests were
fabricated using a dual beam scanning electron micro-
scope/focused ion beam (SEM/FIB) system (FEI Helios
Nanolab 600) [36]. Three steps were involved; namely,
coarse milling, fine milling and a final superfine polish-
ing. The beam current was reduced to 9.7 pA at the final
polishing step to minimize FIB damage. The FIB-milled
pillars had varying diameters (d) ranging from ~230 nm
to ~2 µm, and aspect ratios of the height (h) to d, ranging
from ~2.5 to ~3. The tapering angle of the pillars ranged
from ~2° to ~ 4°. Diameters measured at the top of the
pillars were used to calculate the engineering stress, since
for most of the pillars the deformation was mainly loca-
lized in the upper region [37,38]. The compression tests
were performed in an SEM equipped with a nanoindenter
(PI 85 SEM PicoIndenter, Hysitron Inc.) containing a
5-µm diameter diamond flat punch. The compression
tests were performed in the displacement control mode at
a prescribed strain rate of 10−3s−1. The pillars were de-
formed to a prescribed displacement (or strain) followed
by an incremental unloading.
The overall morphology of the post-compressed pillars
was characterized by a dual beam SEM/FIB at 10 kV.
Thin-foil transmission electron microscopy (TEM) sam-
ples of both the as-deposited films and the pillars after
compression were prepared by using an in-situ lift-out
technique [39] inside the same dual beam SEM/FIB sys-
tem. The TEM lamellae were finally polished with a 5 kV
and 24 pA Ga+ion beam to minimize ion beam damage
to the samples. Bright-field diffraction contrast imaging
and selected area diffraction (SAED) were performed by
using a Philips CM200 TEM operating at 200 kV. High-
angle annular dark-field (HAADF) scanning TEM
(STEM) observations and compositional mapping were
conducted by using an aberration-corrected FEI Titan
Themis TEM operating at 200 kV.
RESULTS
Hierarchical nanostructures of the dual-phase CrCoNi
MEA
As illustrated in Fig. 1, the as-deposited CrCoNi alloy
exhibits a hierarchical nanostructure. The characteristic
microstructural feature of the alloy is the columnar grains
with widths ranging from ~50 to 150 nm (Fig. 1a). At the
atomic scale, typical STEM images of two columnar
‘grains’ (Fig. 1b, c) reveal alternating HCP and FCC
phases, as well as the presence of a high density of planar
defects, including stacking faults (SFs) and twin bound-
aries (TBs). The FCC phase has a consistent orientation
relationship with the HCP phase, specifically,
{111} (0001)
F H
and
110 1120
F H
. The grain tex-
ture was identified with the [0001]Hand
111 F
direction
oriented along the film growth direction. Extensive STEM
investigations further confirmed the presence of HCP and
FCC segments, each having a thickness of ~2 to 8 nm and
alternatively distributed through the columnar grains.
The volume percentages of HCP and FCC phases were
statistically estimated to be ~60% and ~40%, respectively,
based on the STEM images of 15 columnar grains. In
principle, the hierarchical nanostructures of the as-
deposited MEA with dual phases are significantly distinct
from those of the coarse-grained counterpart with a single
FCC phase structure fabricated by conventional me-
tallurgical methods [14,17], endowing it with ultrahigh
hardness and unique deformation behavior [29].
Size-dependent deformation of the dual-phase CrCoNi
MEA
To comprehensively explore the mechanical behavior of
the CrCoNi alloy, in-situ compression tests of pillars were
conducted using samples with diameters ranging from 2
to 0.23 µm. It is apparent that the sample sizes have re-
Figure 1 (a) A bright field TEM image of the as-deposited CrCoNi alloy showing the columnar structure containing a high density of planar defects;
(b, c) STEM images of two separate grains showing the coexistence of HCP and FCC phases, as well as SFs and TBs.
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markable effects on the deformation behavior of the alloy,
as revealed in Fig. 2. When the sample diameters were
larger than 1 µm, pillars failed abruptly through major
shear banding. In striking contrast, pillars with diameters
below 1 µm exhibited appreciable ductility controlled by
steady shear banding. Fig. 2a, b show the undeformed
morphology of a pillar with a diameter of 1.45 µm (pillar
1) and its compression failure mode. The pillar initially
showed elastic deformation, followed by limited strain
hardening with the yield stress reaching ~4.2 GPa before
the stress drop (Fig. 2c). The sudden drop in stress was
caused by shearing of the top surface toward the right-
hand side of the pillar. Finally, the pillar failed catastro-
phically with a strain-to-failure of ~5% through the for-
mation of multiple major shear bands, indicated by white
arrows, and the majority of the pillar body sheared to-
ward the left-hand side. The compression test was per-
formed under displacement-control, therefore extremely
rapid propagation of these shear bands left the punch
behind (moving at the preset displacement rate), resulting
in a stress drop straight to zero.
In comparison, an as-fabricated pillar with a diameter
of 0.88 µm (pillar 2) and its post-compression morphol-
ogy are presented in Fig. 2d, e, respectively. The corre-
sponding engineering stress-strain curve is shown in
Fig. 2f. Owing to the taper geometry of the FIB-fabricated
pillar, the plastic deformation commenced from the pillar
top surface that came into contact with the indenter. The
top part of the pillar experienced a uniform plastic flow
before the shear banding started. The stress decreased
after point 1 due to the formation and development of the
first shear band (SB1). The stress decreased to a trough
(point 2) when the SB1ceased to propagate. Following
that, the stress increased again until a second decrease
Figure 2 SEM images of undeformed and post-compressed pillars with diameters of (a, b) 1.45 µm, (d, e) 0.88 µm, (g, h) 0.23 µm; (c, f, i) engineering
stress-strain curves corresponding to the pillars in (a, d, g), respectively.
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occurred at point 3 caused by the initiation of the second
shear band, as indicated by SB2(Fig. 2e). The indenter
retracted automatically when it reached the prescribed
maximum displacement. Different from the catastrophic
shear failure of pillar 1, the propagation of the shear
bands in Fig. 2e was relatively slow and no large dis-
placement bursts was observed. As shown in the inset in
Fig. 2f, the curve does not exhibit a smooth stress re-
sponse. The intermittent stress drops were presumably
due to successive controlled steady shears in the same
band.
With further reductions in the pillar diameter, uniform
and steady plastic flow became dominant during com-
pression testing. Fig. 2g, h display the morphology of a
pillar with a diameter of 0.23 µm (pillar 3) before and
post deformation. The corresponding engineering stress-
strain curve is shown in Fig. 2i. Plastic deformation was
mainly confined at the top of the slightly tapered pillar,
which changed the pillar shape to a “mushroom-like”
morphology, indicative of considerable plasticity in the
pillar top. The stress response after a strain of 5% presents
“apparent hardening” in the engineering stress-strain
curve, which may result from the increase of the contact
area between the indenter and the pillar tip with con-
tinued compression. The pillar shows ‘ductile’ deforma-
tion with some intermittent stress drops apparent in
Fig. 2i.
The yield strength was measured and plotted in Fig. 3
as a function of pillar diameter. Different from the
“smaller is stronger” phenomenon now well-known for
most crystalline materials [31,40], it turns out that for the
hierarchically nanostructured CrCoNi alloy there is no
apparent dependence of yield strength on the pillar dia-
meter over the size range covered in this work.
Deformation mechanisms of the dual-phase CrCoNi MEA
To obtain a better understanding of the deformation
mechanisms at the atomic level, the microstructure of the
regions both within and around the shear bands of the
post-compressed pillars was analyzed by using STEM
Figure 3 Yield stress versus pillar diameter for nanostructured CrCoNi
alloy pillars.
Figure 4 (a) A STEM image of the post-compressed pillar with a diameter of 1.45 µm failed catastrophically via shear banding. STEM images of the
regions enclosed by (b) red rectangle, and (c) green rectangle in (a) showing the structure of shear bands. Magnified STEM images and the
corresponding FFT (inset) of three selected regions: (d) near the shear band and (e) inside the shear band shown in (b), and (f) inside the shear band
shown in (c).
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imaging. Fig. 4 exhibits the postmortem microstructures
of the pillar presented in Fig. 2b (1.45 µm diameter) that
failed catastrophically during the compression testing.
Three shear bands, as indicated by white arrows, are
detected in the part of the pillar that sheared to the left.
The shear band regions enclosed by red and green rec-
tangles are magnified and shown in Fig. 4b, c, respec-
tively. The grains inside the shear bands still possess a
columnar structure but are sheared to the left, forming an
angle with respect to the less deformed columnar grains.
In the region next to a shear band enclosed by the red
rectangle in Fig. 4b, similar to those in the as-deposited
film, the HCP and FCC structures still coexist as revealed
in the magnified STEM micrograph presented in Fig. 4d
and its corresponding fast Fourier transform (FFT) pat-
tern (inset). However, only the FCC phase was observed
within the shear band, as evidenced in the STEM image in
Fig. 4e taken from the same grain as Fig. 4d, f obtained
from another shear band presented in Fig. 4c, as well as
their corresponding FFT patterns (inset). These findings
signal that plasticity localized within the shear bands and
that the phase transformation from HCP to FCC un-
expectedly occurred during shear banding.
Fig. 5a presents an STEM image of pillar 2 (0.88 µm
diameter) shown in Fig. 2e, indicating the pillar was de-
formed mainly by two shear bands (SB1and SB2marked
by white dash lines). The slight loss of the top-right
corner of the pillar was due to FIB milling. A magnified
image presented in Fig. 5b demonstrates the micro-
structures within the SB2(the bright contrast region) with
a thickness of ~40 nm enclosed by a red rectangle shown
in Fig. 5a. The atomic structure within the shear band
enclosed by a green rectangle in Fig. 5a and the corre-
sponding FFT pattern (inset), as exhibited in Fig. 5c,
validate the existence of a single FCC structure in lieu of
the FCC and HCP phases, which is in line with those in
Fig. 4e, f. In addition to the phase transformation from
HCP structure to FCC structure, other dislocation activ-
ities were also frequently detected within the shear bands.
As indicated by the high-resolution inverse FFT (IFFT)
images (Fig. 5d, e), Lomer-Cottrell (L-C) locks formed in
the shear bands. In general, when two extended disloca-
tions on two intercrossing {111} slip planes meet each
other, the two leading partials form a stair-rod dislocation
by the following reaction:
1
6[121] + 1
6[211] 1
6[110].
The stair-rod dislocation and the two trailing partial
dislocations form a sessile L-C lock dislocation with a
Burgers vector of
1
2[110]
. Extra planes can be seen for
Figure 5 (a) An STEM image of the deformed pillar with a diameter of 0.88 µm in Fig. 2d, showing the two shear bands (SB1and SB2). (b) A
magnified STEM image of part of SB2. (c) An STEM image and its corresponding FFT (inset) exhibiting the FCC structure inside the shear band. (d, e)
IFFT images of an L-C lock dislocation configurations inside the shear band. (f) An STEM image of the boundary between the sheared and un-sheared
region indicating the misorientation and dislocations at the shear band boundary. (g) A shockley partial dislocation at FCC/HCP interfacial region. (h)
An STEM image and its corresponding FFT (inset) demonstrating the highly deformed region beneath the top surface of the pillar that exhibits FCC-
phased nanograins with the twin structure.
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both the
(111)
and the
(111)
planes. It has been reported
that L-C locks play a critical role in the strain hardening
of FCC metals, because they are effective barriers to
mobile dislocations [41]. Therefore, a higher stress is
required during the later stages of shear band propaga-
tion. Since L-C locks formed via the reaction of leading
partials of extended dislocations, their formation are
substantially catalyzed in low SFE materials, such as the
CrCoNi MEA [42].
The structure of the shear band boundary can be found
in Fig. 5f. The regions inside (top) and right below
(bottom) the shear band, exhibit a single FCC phase with
a misorientation angle of 21°, indicating an anticlockwise
crystallographic rotation inside the shear band. It is most
likely that the anticlockwise lattice rotation is caused by
the geometry of the shear deformation. A large shear
strain is introduced to the rotated
{111}F
or
(0001)H
plane inside the shear band, since it is parallel to the shear
direction, facilitating the dislocation activities on the
{111}F
or
(0001)H
plane. Further, the shear band
boundary consists of high density of dislocations, as
marked by the white ⊥signs, which mainly leads to the
misorientation across the shear band. The STEM image of
an FCC/HCP interfacial region (Fig. 5g), enclosed by a
violet rectangle in Fig. 5f, can substantiate that the FCC
phase is formed via the glide of Shockley partial dis-
locations with a Burgers vector of
1
31010
. In addition to
the shear bands, a highly deformed region enclosed by the
blue rectangle beneath the top surface in Fig. 5a was also
carefully characterized, as demonstrated by an STEM
image and corresponding FFT pattern (inset) shown in
Fig. 5h. FCC-structured nanograins with twins can be
observed in this region. The twin thickness in the de-
formed regions is ~10 to 20 nm, which is much thicker
than those in the as-deposited samples, indicative of the
possible occurrence of detwinning. Meticulous examina-
tion of the structure in the highly deformed regions from
four pillars with different diameters found that the HCP
phase rarely presented, verifying the occurrence of a
phase transformation from the HCP to FCC phase in
these areas, which was essentially distinct from the
transformation from FCC to HCP that occurred in the
coarse-grained counterparts [17,43].
DISCUSSION
Size-independent ultra-high strength
The experiments described above show that the nanos-
tructured CrCoNi pillars exhibit a size-independent yield
strength of ~4 GPa, which is a factor of ~10 (i.e., an order
of magnitude) higher than that of bulk FCC CrCoNi
MEA [14,17]. The hardness of the CrCoNi film, measured
by nanoindentation, is ~10 GPa [29], about five times
increase as compared with that of the as-cast CrCoNi
(~1.86 GPa) [44]. The ultra-high yield strength and
hardness are believed to arise from the hierarchical na-
nostructural features of the CrCoNi film, including col-
umn nanograins, a high density of planar defects and
HCP/FCC phase interfaces, strong [0001]Hand
111 F
directional texture, and a large fraction of the HCP phase.
It is well known that the GB strengthening efficiency of
CrCoNi-based HEAs and MEAs is significantly larger
than conventional FCC alloys [45], which reasonably
enables the high strength of nanostructured HEA and
MEA. In addition, within individual column nanograins,
the high density of planar defects including SFs, TBs and
phase boundaries (PBs) plays crucial roles in strength-
ening the materials since they are effective in impeding
the transmission of dislocations on the slip systems in-
clined to these defects [46–49].
Recent investigations revealed that the deformation-
induced HCP phase in bulk HEA and MEA are stronger
than their FCC counterparts [50]. Therefore, the presence
of HCP phases in the nanostructured CrCoNi films un-
ambiguously results in a higher yield strength than its
single FCC structure. Moreover, the growth direction of
the CrCoNi film is dominated by [0001]Hand directions
(i.e., it exhibits a strong texture). The easy glide plane in
the HCP phase is the (0001) basal plane, which is normal
to the loading direction, thus resulting in a lack of re-
solved shear stress for dislocation motion on the basal
plane. Further, the activation of <c> and/or <c+a> dis-
locations on the pyramidal or prismatic slip planes can be
inhibited by the high density of SFs. Similarly, the nu-
cleation of dislocations is essentially suppressed in the
FCC phase with profuse SFs and TBs due to their ex-
tremely small spacing [51,52]. Although dislocations can
also glide along TBs that are orthogonal to the loading
direction as well, there is no resolved shear stress for the
migration of dislocation along TBs in the present study.
The strong [0001]Hand
111 F
texture therefore increases
the yield strength of both the HCP and the FCC phases.
In principle, the ultrahigh strength of nanostructured
CrCoNi MEA is endowed by the integrative effects
stemming from hierarchical, heterogeneous micro-
structures across length scales from atomic level to nan-
ometer, rather than by any specific microstructural
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features.
Although sample size effects on the strength of metallic
materials have been widely acknowledged, there is no
dependence of the yield strength on the pillar diameter
observed over the size scale of the nanostructured
CrCoNi alloy examined in this study. As mentioned
above, the ultra-high yield strength in the CrCoNi pillars
originates from the difficulty of dislocation motion due to
the densely spaced planar defects. Actually, the well-
known “smaller is stronger” phenomenon observed in
nano-sized crystalline materials originates from sophis-
ticated dislocation activities [32,35]. Therefore, apparent
size effects might not be expected here because disloca-
tions are trapped inside the pillars by planar defects and
phase boundaries and thus hardly move and escape from
the free surface to reach a dislocation starvation state.
Although size-induced weakening phenomena were also
found in nanograined metal pillars due to the GB-assisted
deformation [53], the special GB features endowed by the
column grained structures can largely inhibit the GB
migration and GB sliding. Therefore, the yield strength of
current nanostructured CrCoNi MEA is essentially size-
independent, which stems from its unique micro-
structures.
Size-dependent deformability
Unlike dislocation-mediated plastic flow observed in bulk
CrCoNi, pillars made of a CrCoNi film in this study
deform mainly through shear banding. Generally, when
conventional modes of deformation with crystallographic
features are inhibited, shear banding, as a form of local
instability, will be activated to accommodate the plasti-
city, which is frequently observed in high-strength ma-
terials such as bulk metallic glasses and nanocrystalline
materials [54–57]. It is also well known that localized
deformation tends to occur when the ratio between the
strain hardening rate and the applied stress drops below a
critical value [58]. Therefore, the ultrahigh strength of the
nanostructured CrCoNi film and its limited strain hard-
ening at the early stage of deformation will essentially
promote the localized deformation in the form of shear
banding that appears at sites of stress concentration.
Although shear banding is the primary mechanism of
plastic deformation, the remarkable sample size effects on
the deformability of CrCoNi pillars are readily observed
in the current study. With the diameters of the pillars
gradually decrease from 2 to 0.23 µm, the deformation
shows a transition from (1) highly localized and cata-
strophic shear banding to (2) slowly propagated and
steady shear banding accompanied by softening, and
eventually (3) “homogeneous-like” deformation with ap-
preciable plasticity (i.e., a brittle to ‘ductile’ transition).
The size effects on the plasticity of CrCoNi pillars do not
follow the trend of that of single crystal pillars [31,34], but
are similar to that of metallic glass pillars [54–56]. Ac-
tually, the size effects on the deformability of CrCoNi
pillars can be understood by considering the critical stress
required for shear band propagation to fracture [59,60].
While the initiation of shear banding is controlled by the
shear stress, the propagation of the shear band is gov-
erned by the release of stored elastic energy [54]. For
materials with a high density of obstacles, slip is confined
in such a way that the elastic strain energy cannot be
easily released by means of dislocation motion. The total
elastic strain energy Ustored in a pillar with a diameter d
and a height hdue to the applied stress σcan be obtained
as [59]:
U d h= / 8. (1)
2
The consumption or release of the elastic energy, ΔU,
during the shear-banding process can be calculated by
[59]:
U d= 2 / 4, (2)
2
where Γis the energy per unit area of shear band, which is
assumed to be a constant, regardless of the sample size.
Similar to what takes place during crack propagation
(Griffith criterion) [61], the stored elastic energy Uwould
be completely relieved by the catastrophic shear band that
traverses the pillar. Then, the size effect on deformation is
rendered by the dimensional origin (i.e., three-dimen-
sional energy stored in volume versus two-dimensional
energy consumption by surface). Therefore, ΔU=Ugives
an estimate of the critical stress σcrequired for propa-
gating a single shear band to fracture the pillar [59]:
h
E
ad
=2 2 / =2 2 , (3)
c
where Eis Young’s modulus, ais the aspect ratio (height/
diameter). The critical stress required to raise the strain
energy so that it was sufficient to enable the shear band
propagation gradually to increase as the pillar diameter
decreases (Fig. 6). Note that the yield stress is measured to
be size-independent. Regarding the limited strain hard-
ening in pillars with different diameters, it can be rea-
sonably inferred that the flow stress that initiates shear
band also exhibits size-independency. That means, there
exists a critical diameter, dc, where the transition from
brittle to ‘ductile’ occurs (Fig. 6), since below this size,
catastrophic shear banding becomes suppressed due to
the insufficient elastic strain energy to fuel the expansion
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of a single shear band traversing the pillar. By equating
Equations (1) and (2), Γis estimated to be ~71 J m−2.
Using the averaged flow stress of ~4.2 GPa as σc,
Γ=71 J m−2and E=244.9 GPa [62], the sample dcwas
calculated from Equation (3) to be ~0.93 µm (Fig. 6),
agreeing well with the experimental observations.
There are some similarities in the size-dependency of
the deformation modes in the CrCoNi pillars identified in
this study with that of metallic glass pillars [59]. However,
different from metallic glass pillars, the steady shear
banding behavior was observed in the CrCoNi pillars,
which could be contributed to the hardening effect that
slowed down the propagation of the shear bands. As
shown in the inset in Fig. 2f, the curve does not show a
smooth stress response. The intermittent growth behavior
of the single shear band suggests the coexistence of the
softening and subsequent hardening events. In the early
stages of shear banding, the dominant deformation me-
chanisms, such as detwinning, SF elimination, and HCP
to FCC phase transformation, are considered to be strain
softening processes, resulting in the stress drops to the
valley of point 2 in Fig. 2f. Strain softening is not always
detrimental, since it is conducive to accommodate plas-
ticity, and the high stress required to activate these soft-
ening mechanisms ensures the high yield strength. At the
later stages, hardening mechanisms, such as the forma-
tion of L-C locks and dislocation interactions, would kick
in inside the shear band. The hardening mechanism is
working to mitigate the softening effect, which could
arrest the shear band propagation and increase the flow
stress, e.g., from point 2 to 3 in Fig. 2f, triggering more
shear bands to bear further deformation. Therefore, a
stable and prolonged plastic deformation process was
attained in smaller pillars due to the development of
numerous small and slowly propagated shear bands.
However, in larger pillars, instantaneous occurrence of
the shear event suppresses the strain hardening, which
enables the softening effect to be dominant, inducing the
sudden fracture of the pillar by the fast propagation of the
shear bands. The transition from brittle to ‘ductile’ de-
formation (i.e., the increase of the deformability) with
decreasing pillar diameters can thus be attributed to the
decreased propagation ability and resulting increased
amount of the shear bands.
Pillar geometry may influence the shear banding be-
havior, thus, the geometric factors, such as tapering,
should be seriously considered when ascertaining the size
effect on the deformation mode of pillars [63–65]. As
shown in Fig. 7, with comparable tapering angles of ~4.3°,
the larger pillar with a diameter of 1.25 µm still exhibits a
distinct catastrophic failure mode (Fig. 7b), while the
steady shearing events can be apparently detected in the
small pillar with a diameter of 0.4 µm (Fig. 7d). There-
fore, the brittle to ‘ductile’ transition observed in this
study is governed by the intrinsic size effect, rather than
an artifact.
Phase transformation from HCP to FCC phase
In the bulk CrCoNi-based HEAs and MEAs, previous
studies have revealed that the phase transformation from
FCC to HCP is, indeed, possible [16,17,43,66]. Theore-
tical investigations have proposed that the HCP phase is
more thermodynamically stable than the FCC phase
[67,68] due to its low free energy and thus the formation
of the HCP phase in the FCC phase is energetically fa-
vorable. However, experimental studies have revealed that
the sophisticated generalized SFE curve of HEAs and
MEAs due to the compositional complexity may kineti-
cally hinder the phase transformation [16,17,69], leading
to only a limited quantity of the HCP phase that can
transform from the FCC matrix. In striking contrast, the
current study verified the unexpected reverse phase
transformation from HCP to FCC in a dual-phased
CrCoNi thin film at room temperature.
For all the pillars, a complete transformation from HCP
to FCC is accomplished if a Shockley partial dislocation
Figure 6 Schematic illustration of the transition from catastrophic un-
steady shear banding to stable and slow shear banding with decreasing
pillar diameter. The minimum stress required for a shear band to
fracture the pillar (blue solid line) crosses the experimental-measured
flow stress (~4.2 GPa) when shear band initiates (green solid line) at a
critical diameter, dc, where the transition between catastrophic and
stable shear banding occurs. When the diameter is smaller than the
critical size of ~0.93 µm, the strain energy raised by the flow stress is not
high enough to allow for the shear band to fracture the pillar, thus,
showing stable shear banding.
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that creates a SF occurs on alternating close-packed
planes in the HCP phase [70]. However, the driving force
for the phase transformation in larger pillars and small
(submicron) pillars could be different. In large pillars, an
appreciable localized heating is estimated to occur if the
stored elastic energy is catastrophically released in an
avalanche configuration [71]. The catastrophic failure and
the localized deformation would lead to an acceleration of
strain rate, which might induce heat concentration and
thermal softening in the shear band [72]. Unlike the vein-
like patterns usually observed on the fracture plane of a
metallic glass sample [73], the sheared surface of the large
CrCoNi pillar (Fig. 2b) was marked with a striated pat-
tern, which might be formed by softened metal surfaces
rubbing against each other during shear deformation
under an extremely high strain rate. This observation
indicates that the significant rise of the temperature may
result from fast, localized deformation within shear
bands. Based on the theoretical calculations, for CrCoNi,
HCP is the stable structure at low temperature, while FCC
becomes more stable when the temperature increases to
above ~500 K [74]. Therefore, the HCP to FCC phase
transformation observed in the catastrophically failed
large pillars could be induced by the localized heating in
the shear band.
For sub-micron pillars with steady deformation, the
local temperature increase induced by dissipative heating
during straining may not be high enough to drive the
HCP to FCC phase transformation. The phase transfor-
mation in small pillar could be promoted by the high
shear stress. At the early stage of deformation, multiple
dislocations are nucleated at intersections between the
planar defects and GBs/free surfaces [51,75], and then
glide on the slip planes until they pile up against local
obstacles (SFs and TBs). As the deformation proceeds,
these piled-up dislocations start to react with the TBs and
SFs when the shear stress is sufficiently large (Fig. 8a, b).
Extensive dislocation-planar defects interactions enable
the formation of shear bands with apparent crystal rota-
tion (Fig. 8c)[51,75]. As shown in Fig. 8, the (0001) basal
plane in the HCP phase is normal to the loading direc-
tion, thus there is no resolved shear stress available for
driving dislocation motion on the basal plane. The crys-
tallographic rotation caused by shear banding results in
an increase in Schmidt factor and resolved shear stresses,
thus facilitating the movement of partial dislocations on
the basal plane of the HCP phase (Fig. 8d). Successive
gliding of partials on every other basal plane of the HCP
Figure 7 The size effect on the deformation mode remains even if the
tests were conducted on larger pillars with the same degree of tapering as
the smaller ones. SEM images of undeformed and post-compressed
pillars with the same tapering angle of 4.3°, but different diameters of
(a, b) 1.25 µm, and (c, d) 0.4 µm.
Figure 8 Schematic illustrations showing the structural evolution under shear banding in a nanostructured CrCoNi pillar.
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phase induces the phase transformation inside the shear
band [70]. This rotation corresponds to geometric soft-
ening since it causes a lattice rotation into a geometrically
softer orientation [76]. The large localized resolved shear
stress caused by the lattice rotation inside the shear band
is critical to the HCP to FCC phase transformation
(Fig. 8e).
The HCP to FCC phase transformation in the nanos-
tructured CrCoNi film might also be promoted by the
high density of pre-existing SFs in the HCP phase, which
act as nucleation sites for the formation of FCC segments
at the beginning of deformation, overcoming the initial
barrier seen from a planar-defect-free HCP phase tran-
sitioning to the FCC phase [77]. As reported by previous
studies, the reverse martensitic transformation tempera-
ture can be lowered by the introduction of mechanical
energy by deformation [78,79]. Noting that the original
CrCoNi alloy has a high density of obstacles (GBs, TBs
and SFs) for dislocation activities, thus, the mechanical
energy cannot be easily released by means of dislocation
propagation. A large amount of strain energy stored in
the pillar might raise the free energy, providing the
thermodynamic driving force for the HCP to FCC phase
transformation, while further investigations are still nee-
ded to clarify the underlying mechanisms. All these re-
sults show that the HCP to FCC phase transformation in
nanostructured CrCoNi is possible, and the HCP phase is
mechanically unstable at room temperature. The HCP to
FCC phase transformation leads to the formation of
transparent HCP/FCC phase interface and softening. It is
understood that strain softening would contribute to ac-
commodate plasticity in the CrCoNi pillars.
CONCLUSIONS
The size effect on the deformation behavior of nanos-
tructured, dual-phase CrCoNi alloy was examined by
using in-situ compression testing, assisted by STEM
analysis. The nanostructured CrCoNi MEA with hier-
archal microstructures exhibits a ultrahigh size-
independent yield strength of ~4 GPa, which has its ori-
gins in the densely spaced planar defects including SFs,
TBs and PBs. The compression tests of micro-pillars with
different diameters showed that the pillars deformed
mainly via shear banding due to the presence of na-
noscale planar defects that limited the formation of mo-
bile dislocations. As the pillar diameter decreased from 2
to 0.23 µm, the pillars showed transitions from highly
localized shear banding, through steady shear banding,
and to homogeneous deformation that promoted sig-
nificant plasticity. An unusual phase transformation from
the HCP to FCC phase was observed in the heavily de-
formed regions mediated by Shockley partial dislocations.
The high density SFs presented in the CrCoNi alloy are
believed to facilitate the HCP to FCC phase transforma-
tion.
Received 28 March 2020; accepted 29 April 2020;
published online 29 July 2020
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Acknowledgements This work was supported by the Australian Re-
search Council Discovery Projects Grant, and partly supported by the
Fundamental Research Funds for the Central Universities
(SWU118105). An X acknowledges the financial support from Australia
Research Council (DE170100053) and the Robinson Fellowship Scheme
of the University of Sydney (G200726). The authors acknowledge the
facilities and the scientific and technical assistance of the Australian
Microscopy and Microanalysis Research Facility (ammrf.org.au) node at
the University of Adelaide: Adelaide Microscopy. In particular, the au-
thors thank Dr Animesh Basak and Dr Ashley Slattery of Adelaide
Microscopy for their support and expertise.
Author contributions Chen Y, Xie Z, An X and Zhang S conceived
the project. Chen Y conducted the FIB, microcompression and TEM
experiments. Zhou Z fabricated the samples. Chen Y, An X, Liao X and
Xie Z interpreted the results and wrote the manuscript. All authors
contributed to the discussion of the results, and comments on the
manuscript.
Conflict of interest The authors declare that they have no conflict of
interest.
Yujie Chen obtained her BEng degree (first-class
honors) in 2011 and PhD degree in materials
science in 2016 from The University of Sydney.
Upon completion of her PhD, she was employed
as a postdoc in the School of Mechanical En-
gineering in the University of Adelaide in Aus-
tralia. She is currently a research fellow in the
Southwest University in China. Her current re-
search involves microstructure optimization and
mechanical properties enhancement of alloys,
and calcified tissues.
Xianghai An received his PhD degree from the
Institute of Metal Research, Chinese Academy of
Sciences in 2012. After receiving his PhD degree,
he commenced to work as a research fellow at
The University of Sydney. He is currently a
Lecture/Robinson Fellow at The University of
Sydney. His research mainly focuses on materials
design, mechanical behavior, and structure-
property relationship of advanced metallic ma-
terials, nanomechanics and nanoplasticity, me-
tallic additive manufacturing and advanced
materials processing.
SCIENCE CHINA Materials. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .ARTICLES
13
© Science China Press and Springer-Verlag GmbH Germany, part of Springer Nature 2020
Sam Zhang received his PhD degree (1991) in
ceramic materials at the University of Wisconsin-
Madison, USA. He joined Nanyang Technologi-
cal University as an associate professor and was
promoted to full professor in 2006. He is cur-
rently a professor and head of the Center for
Advanced Thin Films and Devices in the
Southwest University in China. He is also Fellow
of the Institute of Materials, Minerals and
Mining, Fellow of the Royal Society of Chemistry
and Fellow of the Thin Films Society. His re-
search interests include preparation and characterization of hard yet
tough ceramic nanocomposite coatings, and functional thin films.
双相纳米结构CrCoNi中熵合金变形行为的尺寸
效应
陈玉洁1,2,3,安祥海3*,周志烽4, Paul Munroe5,张善勇1*,
廖晓舟3,谢宗翰2,6
摘要 本文结合原位扫描电子显微镜微柱压缩与透射电子显微镜
技术,研究了具有双相多级纳米结构的CrCoNi中熵合金变形行为
的尺寸效应.研究表明,该合金的屈服强度高达~4 GPa, 这主要归
因于其多级微观结构特征,包括柱状纳米尺寸晶粒、织构、高密
度的层错、孪晶界、晶界和相界.在变形过程中,该合金的屈服强
度基本与微米尺度样品的尺寸无关,但其变形行为却强烈依赖于
样品大小.具体来说,随着微柱直径减小,材料主要的变形模式从
突发的局部剪切带变为具有明显塑性的均匀变形.这种转变是由
剪切带穿过微柱所需的临界应力与样品尺寸紧密相关所决定的,
剪切诱导的软化和随后的硬化机制之间的竞争也起了重要作用.
此外,变形引起了六方密排结构到面心立方结构的相变,该相变导
致的应变软化对材料变形中的塑性有重要贡献.这些发现揭示了
样品尺寸对可用于微纳机电系统的纳米结构金属材料的力学行为
有着重要影响.
ARTICLES . . . . . . . . . . . . . . . . . . . . . . . . . SCIENCE CHINA Materials
14 © Science China Press and Springer-Verlag GmbH Germany, part of Springer Nature 2020
