Journal of Materials Science & Technology 214 (2025) 143–152
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Journal of Materials Science & Technology
journal homepage: www.elsevier.com/locate/jmst
Research Article
Restructuring the cell network of non-equiatomic CoCrFeNiMoC
medium-entropy alloy fabricated by laser powder bed fusion
Hyeonseok Kwon a,1, Eun Seong Kim b,1, Yoon-Uk Heo c, Jungho Choe d, Rae Eon Kim c,
Soung Yeoul Ahn b, Sang-Ho Oh a, Jeong Min Park d,e,∗, Byeong-Joo Lee b,
Hyoung Seop Kim b,c,f,g,∗
a
Center for Advanced Aerospace Materials, Pohang University of Science and Technology (POSTECH), Pohang, 37673, Republic of Korea
Department of Materials Science and Engineering, Pohang University of Science and Technology (POSTECH), Pohang, 37673, Republic of Korea
c
Graduate Institute of Ferrous & Eco Materials Technology, Pohang University of Science and Technology (POSTECH), Pohang, 37673, Republic of Korea
d
Department of 3D Printing Materials, Korea Institute of Materials Science (KIMS), Changwon 51508, Republic of Korea
e
Department of Advanced Future Convergence Materials, Korea University, Seoul 02841, Republic of Korea
f
Institute for Convergence Research and Education in Advanced Technology, Yonsei University, Seoul, 03722, Republic of Korea
g
Advanced Institute for Materials Research (WPI-AIMR), Tohoku University, Sendai, 980-8577, Japan
b
a r t i c l e
i n f o
Article history:
Received 25 April 2024
Revised 13 June 2024
Accepted 1 July 2024
Available online 16 July 2024
Keywords:
Laser powder bed fusion
Medium-entropy alloy
Cell structure
Strain hardening
Precipitation
Back stress hardening
a b s t r a c t
Metal additive manufacturing (MAM) enables near-net shape production of components with minimized
waste and excellent mechanical performance based on multi-scale microstructural heterogeneity. Especially, the dislocation cell network that often bears elemental segregation or precipitation of a secondary
phase contributes to enhancing the strength of additively manufactured materials. The cell boundaries
can also act as active nucleation sites for the formation of precipitates under post-MAM heat treatment,
as the chemical heterogeneity and profuse dislocations generate a driving force for precipitation. In this
work, we subjected a Co18 Cr15 Fe50 Ni10 Mo6.5 C0.5 (at%) medium-entropy alloy fabricated by laser powder
bed fusion (LPBF) to post-LPBF annealing at 900 °C for 10 min. Microstructural investigation revealed
that the cell boundaries of the as-built sample, which were decorated by Mo segregation, are replaced by
μ phase and M6 C type carbide precipitates during annealing while the grain structure and size remain
unaffected, indicating that the post-LPBF annealing delivered the proper amount of heat input to alter
only the cell structure. The yield strength slightly decreased with annealing due to a reduction in the
strengthening effect by the cell boundaries despite an increased precipitation strengthening effect. However, the post-LPBF annealing improved the strain hardenability and the ultimate tensile strength was
enhanced from ∼1.02 to ∼1.15 GPa owing to reinforced back stress hardening by the increased dislocation pile-up at the precipitates. Our results suggest that the cell structure with chemical heterogeneity
can be successfully controlled by careful post-MAM heat treatment to tailor the mechanical performance,
while also providing insight into alloy design for additive manufacturing.
© 2025 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science &
Technology.
1. Introduction
Over the last few years, multi-principal element alloys (MPEAs),
so-called high-entropy alloys (HEAs), and medium-entropy alloys
(MEAs) have been drawing a lot of attention for their unique design concept and promising properties [1–3]. Initial development
of the MPEAs began with the exploration of single-phase alloys
in equiatomic compositions, exemplified by the face-centered cu-
∗
Corresponding authors.
E-mail addresses: jmpark@kims.re.kr (J.M. Park), hskim@postech.ac.kr (H.S. Kim).
1
These authors contributed equally to this work.
bic (FCC) structured CoCrFeMnNi [4], CoCrNi [5], or the bodycentered cubic (BCC) structured TiZrHfNbTa [6]. In pursuit of enhanced mechanical properties, however, the focus has recently
expanded to a wider range of MPEAs that encompass multiphase, non-equiatomic alloys. As a part of such endeavors, ferrous
MEAs that possess both excellent mechanical properties and costeffectiveness compared to existing MPEAs are being actively investigated nowadays [7,8]. The ferrous MEAs are noted as materials that can provide the possibility for simultaneously exploiting
the advantages of the MPEAs and steels, by combining pronounced
solid solution strengthening in the MPEAs with conventional
strengthening mechanisms in steels [9]. A non-equiatomic Co-CrFe-Ni-Mo-C alloy system is a good example of these materials,
https://doi.org/10.1016/j.jmst.2024.07.010
1005-0302/© 2025 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.
H. Kwon, E.S. Kim, Y.-U. Heo et al.
Journal of Materials Science & Technology 214 (2025) 143–152
which shows superior strength and ductility over a wide temperature range of 0.5–298 K [10–13].
Laser-based metal additive manufacturing (MAM) technology
has been drawing wide attention as it enables the production of
components in near-net shapes with design freedom and reduced
material waste [14–17]. In general, it is known that MAM processing produces microstructures with multi-scale heterogeneity involving microscale anisotropic grain structure and nanoscale dislocation cell structure [17–20]. Especially, the dislocation of cell
structure often accompanies elemental segregation and the formation of a secondary phase at the cell boundaries [21–23]. MAM
processing of the MPEAs is under massive attention in the midst
of the search for enhanced mechanical performances owing to
the merits of such hierarchically heterogeneous microstructures
[24–26]. For example, laser powder bed fusion (LPBF) processed
Fe60 (CoNi)30 Cr10 (at%) MEA showed that the elemental heterogeneity at dislocation cell walls locally reduced the phase stability,
which promoted deformation-induced martensitic transformation
and improved work hardening of the MEA [20]. In another example, it has been reported that an Fe65 Ni15 Co8 Mn8 Ti3 Si1 (at%) MEA
fabricated by the direct energy deposition (DED) process exhibits η
nanoprecipitation at the cell boundaries and martensitic transformation at the cell interiors, which leads to improved mechanical
properties by precipitation strengthening and hetero-deformationinduced (HDI) strengthening effects [27].
Similarly, in our previous work [28], we discovered that the
LPBF fabrication of the Co18 Cr15 Fe50 Ni10 Mo6.5 C0.5 (at%) ferrous
MEA produces a microstructure where Mo atoms are heavily segregated and form precipitates along the cell boundaries. The heavy
elemental segregation created lattice misfits at the cell boundaries and enhanced their ability to hinder dislocation movement,
which significantly improved the strength of the material compared to its Mo-free counterpart. Such characteristics of the LPBFbuilt Co18 Cr15 Fe50 Ni10 Mo6.5 C0.5 MEA led us to conceptualize the
idea for this work: the chemical heterogeneity at the cell boundaries, combined with the inherent high dislocation density, would
generate enough driving force for sensitive microstructural changes
in response to heat treatment. In other words, we anticipated that
post-LPBF heat treatment could restructure the cellular network
of Co18 Cr15 Fe50 Ni10 Mo6.5 C0.5 MEA into nanoscale precipitates. Under a suitable condition of post-LPBF annealing, it was expected
that intragranular nanoprecipitates would form along the sub-grain
cell boundaries while suppressing excessive growth of grains or
precipitates, which would enhance the mechanical performance
through precipitation hardening. Following this strategy, the LPBFbuilt Co18 Cr15 Fe50 Ni10 Mo6.5 C0.5 MEA was subjected to an annealing heat treatment for a short duration of 10 min, and the impacts of the post-LPBF annealing on the microstructure and mechanical response were systematically assessed. The findings of
this study suggest the potential of controlling the cellular structure with chemical heterogeneity via post-LPBF heat treatment as
an effective approach to tailoring the mechanical properties of additively manufactured materials.
25 μm were applied. For reference, the scanning direction parallel
to the hatch spacing is defined as the X-direction, while the direction perpendicular to the hatch spacing is the Y-direction. The
building direction is defined as the Z-direction. For information
on the powders and scanning strategy, the reader is referred to
our previous work [28]. The printed sample was subsequently
annealed at 900 °C for 10 min, which is hereinafter denoted by
A900. The temperature was determined based on our previous
works on alloys with similar compositions [8,10], where annealing
at 900 °C promoted precipitation of carbides and μ phase. The
annealing duration was limited to 10 min to avoid grain growth or
excessive coarsening of precipitates.
2.2. Microstructural analysis
Scanning electron microscopy (SEM) analysis in backscattered
electron (BSE) mode was operated with a high-resolution fieldemission SEM (FE-SEM 7100, JEOL). Electron backscatter diffraction (EBSD) analysis was carried out with an FE-SEM (XL30 FEG,
Philips). The samples for the SEM analyses were mechanically polished with 60 0, 80 0, and 120 0 grit SiC papers and electropolished
in an 8 % HCl and 92 % CH3 COOH solution. Transmission electron
microscopy (TEM) and transmission Kikuchi diffraction (TKD) analyses were conducted for characterization at the nanoscale (JEM2100F, JEOL). The samples for the TEM and TKD analyses were
fabricated by a lift-out procedure with a dual-beam focused ion
beam (FIB; Helios Hikari UMSII, FEI). For the characterization of the
crystal structure, X-ray diffraction (XRD) analysis was performed
(D/MAX-2500, RIGAKU Co.) using a Cu Kα beam with a wavelength
of 0.1542 nm. The scans were conducted with a step size of 0.02°
and a holding time of 10 s per step. The dislocation density of the
samples was measured in the Williamson-Hall method [29]. The
instrumental effect was corrected with a standard LaB6 sample.
2.3. Uniaxial tensile test
Tensile test specimens with a gauge geometry of 1.5 mm (gauge
length) × 1 mm (width) × 0.7 (thickness) mm were taken from
the as-built and annealed samples, with the loading direction parallel to the X-direction [30]. The specimens were tested at a strain
rate of 1 × 10–3 s–1 (Instron 1361), and tensile tests were repeated
at least three times for each condition. A digital image correlation
(DIC) method was utilized to measure the tensile strain (ARAMIS
M12, GOM Optical Measuring Techniques).
3. Results
3.1. Microstructural evolution by the post-LPBF annealing
EBSD inverse pole figure (IPF) maps obtained from the X, Y,
and Z-planes of the samples are displayed in Fig. 1(a1 , b1 ). The
IPF maps show an anisotropic grain morphology typical of LPBFfabricated materials. The grain size distribution acquired from the
Z-planes indicates that the average grain sizes of the as-built
and A900 samples are estimated to be ∼17.5 ± 10.0 μm and
∼18.5 ± 10.1 μm, respectively. It is noteworthy that the general grain morphology and size do not vary with annealing. Grain
growth during heat treatment normally occurs through the absorption of small grains by larger grains. In LPBF-fabricated materials
that are built without lateral rotations during scanning, the epitaxial growth of grains along the building direction creates unique
grain structures composed of layers of epitaxially grown fine grains
surrounded by coarse grains formed at the overlapping regions of
neighboring melt pools. The current samples maintain similar grain
structures even after the annealing heat treatment, which implies
2. Experimental procedure
2.1. Fabrication of the samples
Spherical powders with a nominal composition of
Co18 Cr15 Fe50 Ni10 Mo6.5 C0.5 (at%) in a particle size range of 10–
55 μm were manufactured by gas atomization (Hot gas atomization system, PSI Ltd.). A commercial LPBF machine (Mlab, Concept
Laser GmbH) was utilized to build cuboidal samples from the
powders in a meander scanning pattern with a 180° rotation in a
layer-by-layer process. A laser power of 90 W, a hatch distance of
77 μm, a scanning speed of 500 mm/s, and a layer thickness of
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Journal of Materials Science & Technology 214 (2025) 143–152
Fig. 1. EBSD analyses of the (a) as-built and (b) A900 samples. (a1 , b1 ) 3-dimensional construction of IPF maps taken from the X, Y, and Z planes. (a2 , b2 ) phase maps and
(a3 , b3 ) KAM maps taken from the Z plane. A step size of 90 nm was used for the EBSD analyses.
that the heat input was not sufficient to lead to the absorption of
the fine grain layers by the coarse grains.
The phase maps show that both samples have a fully FCC structured matrix (Fig. 1(a2 , b2 )), and kernel average misorientation
(KAM) values are not significantly different between conditions
(Fig. 1(a3 , b3 )). The XRD patterns are presented in Fig. S1 in Supplementary materials. The samples exhibit peaks that correspond
to an FCC structure and minor precipitates, which will be dealt
with in more detail in the following paragraph. The shifting of
the FCC peaks to higher 2θ implies a precipitation-driven reduction in the lattice parameter of the matrix [9]. The dislocation
density calculated via the Williamson-Hall method [29] showed
∼1.83 × 1014 and ∼1.26 × 1014 m–2 for the as-built and A900 samples, respectively. The dislocation density has been reduced to approximately two-thirds of the original value during the post-LPBF
annealing. The unchanged KAM values in the EBSD analyses may
be attributed to an overestimation in A900 owing to the lattice
mismatch between the FCC matrix and the high density of precipitates [31]. Even with possible errors in calculation considered, the
dislocation density calculation results suggest that the MEA shows
a relatively slow recovery kinetics during annealing compared to
other materials such as LPBF-processed (CoCrFeMnNi)99 C1 (at%)
HEA [32], where the same condition of annealing heat treatment
(10 min at 900 °C) led to a significant decrease in the dislocation
density to less than a third. Furthermore, the (CoCrFeMnNi)99 C1
HEA showed an increase in grain size after the annealing, while
the MEA showed nearly no difference. Such different responses
to the heat treatment between the two alloys can be explained
based on the effect of alloy compositions, which will be discussed
shortly.
The SEM-BSE micrographs taken on the Z-plane of the as-built
and A900 samples are shown in Fig. 2. It is observed from the
low-magnification images that the distributions of grain sizes and
melt-pool boundaries (MPBs) along the X direction are quite similar in the as-built and annealed samples (Fig. 2(a1 , b1 )), as already
observed in Fig. 1. However, images at a higher magnification of
the marked areas (Fig. 1(a2 , b2 )) enable a deeper inspection and
reveal striking differences between the samples. The as-built sample forms a cellular structure commonly exhibited in LPBF-built
materials with nanoprecipitates which were identified as the Morich μ phase in our previous work [28]. On the other hand, A900
shows the destruction of the original cell structure. The previous
cell boundaries are completely lost and replaced by the bright precipitates. The bright color of the precipitates indicates that they are
rich in Mo which is heavier than the other elements [32,33]. Therefore, Mo atoms segregated at the cell boundaries are consumed by
the precipitation during annealing and promote the formation and
growth of the Mo-rich precipitates.
For a more detailed analysis of the nanostructures of the asbuilt and A900 samples, TEM analysis was employed (Fig. 3). The
scanning TEM (STEM) and energy dispersive spectroscopy (EDS) elemental maps of the as-built sample (Fig. 3(a1 )) show an enrichment of Mo and Cr atoms at the cell boundaries and precipitates.
The bright field (BF) image shows a precipitate at a vertex of a
hexagonal dislocation cell (Fig. 3(a2 )), which is characterized as a μ
phase by the high-resolution TEM (HR-TEM) image (Fig. 3(a3 )) and
a corresponding selected area electron diffraction (SAED) pattern
(Fig. 3(a4 )). In A900, the STEM image and EDS maps (Fig. 3(b1 ))
show that the cell boundaries with elemental segregation are less
discernible, and Mo-rich precipitates with coarser sizes are filling
their places. However, it should be noted that the dislocations have
not completely disappeared by annealing and some remain near
the cell boundaries or precipitates, as demonstrated by STEM images taken at a different region (Fig. S2). Some precipitates are enriched in both Mo and C, and the SAED and fast Fourier transform
(FFT) patterns (Fig. 3(b3 , b4 )) indicate that some Mo-rich M6 C type
carbides coexist with the μ phase. From the image analysis of the
micrographs, the area fraction of the μ phase and M6 C carbide
precipitates in A900 are calculated as 8.11 % and 3.26 %, respectively, which are much higher than the 1.92 % of the μ phase in
the as-built sample.
It is evident from the EDS maps that the μ phase of A900
shows a chemical composition that is rich in Mo but lean in Cr,
which is different from that of the as-built sample where high
concentrations of both Mo and Cr were contained. The EDS elemental profiles in Fig. S3 show the chemical compositions of the
μ phase particles of the as-built and A900 samples, demonstrating the difference in the compositions. This suggests that the μ
phase in A900 could have been precipitated by new nucleation
during annealing, rather than growing from those that existed beforehand. Furthermore, the μ phase in A900 exhibits (102̄ ) stacking faults (Fig. 4) which were not observed in the as-built sample (Fig. 3(a3 )). It is established that the growth of the Co7 Mo6
type rhombohedral μ phase can be mediated by the addition or
extension of planar defects such as micro-twins or stacking faults
[34–36]. To better understand the precipitation mechanism, we annealed the as-built sample at 800 °C and investigated the nanostructure. The STEM image, EDS elemental maps, and SAED pattern
in supplementary Fig. S4 demonstrate that all the precipitates in
the 800 °C annealed sample are identified as the (Mo, Cr)-rich M6 C
type carbides, with no trace of the original μ phase. It can be thus
elucidated that, during annealing at 900 °C, the original Cr-rich μ
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Journal of Materials Science & Technology 214 (2025) 143–152
Fig. 2. BSE micrographs of the (a) as-built and (b) A900 samples. Images at (a1 , b1 ) low magnifications and (a2 , b2 ) high magnifications.
Fig. 3. TEM analyses of the (a) as-built and (b) A900 samples. (a1 , b1 ) STEM images and EDS elemental maps. (a2 , b2 ) BF images. (a3 ) HR-TEM image of a μ phase precipitate
and (a4 ) corresponding SAED pattern. (b3 ) SAED pattern taken from an M6 C carbide precipitate and (b4 ) FFT pattern taken from a μ phase precipitate.
phase precipitates in the as-built sample completely transform into
the Cr-rich M6 C type carbides first, and some of them transform
again into the new Cr-lean μ phase by overcoming the energy barrier through further heat input. It is possible that the M6 C type
carbides have acted as an intermediate phase on the way to forming the new Cr-lean μ phase during the 900 °C annealing, which
is stable at higher temperatures than the carbide is: the carbides
may have acted as nucleation sites or directly transformed into the
μ phase by releasing C and Cr, to eventually form the microstructure of A900 that bears both the μ phase and M6 C type carbide. It
is also possible that the stacking faults in the μ phase have formed
from structural distortion during the ejection of elements. However, the detailed elucidation of the phase transformation mechanism requires in-situ characterization techniques, such as in-situ
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Journal of Materials Science & Technology 214 (2025) 143–152
Fig. 4. High-magnification observations of a μ phase precipitate in A900. (a) BF image. (b) HR-TEM image.
Fig. 5. Tensile properties of the as-built and A900 samples. (a) Engineering stress-strain curves. The inset in (a) shows the same engineering stress-strain curves drawn up
to uniform elongation. (b) True stress and SHR plotted against true strain.
Table 1
Tensile properties of the as-built and A900 samples.
SSamplesapmles
YS (MPa)
UTS (MPa)
Uniform
elongation (%)
Total
elongation (%)
aAs-built
A900
768 ± 12
656 ± 17
1016 ± 35
1153 ± 38
25.0 ± 1.0
20.5 ± 0.6
53.2 ± 2.9
39.8 ± 0.1
YS. The microstructural features where the original cell boundaries
have transformed into two types of precipitates must be responsible for the exceptional work hardening of A900. The reasoning
behind the mechanical behavior of the samples in relation to the
microstructural evolution during annealing is discussed in the following sections.
4. Discussion
TEM analysis during heating, as the current explanation relies on
ex-situ study.
4.1. The impact of alloy composition on the microstructural evolution
during post-LPBF annealing
3.2. Tensile properties
For the present LPBF-built Co18 Cr15 Fe50 Ni10 Mo6.5 C0.5 MEA, the
annealing heat treatment at 900 °C for 10 min does not make a
huge difference in the grain size, only altering the cell structure.
Such a microstructure of A900 has enabled striking a balance of
strength and ductility, as otherwise a reduction in dislocation density or an increase in grain size due to annealing may have harmed
strength. Several reasons for the retarded grain growth kinetics can
be given. The heat input during the annealing can be consumed
primarily in the formation and phase transformation of precipitates
which limits the generation of sufficient driving force for recrystallization or grain growth [11,37]. The newly formed precipitates
also assist in retarding the grain growth by pinning of the grain
boundary migration. However, even when compared with other
materials that form precipitates under similar conditions of heat
treatment, such as LPBF-processed C-added CoCrFeMnNi HEA, the
grain growth is clearly suppressed. The same parameters of heat
The engineering stress-strain curves of the as-built and A900
samples are shown in Fig. 5(a). The yield strength (YS) of A900 decreased to 656 MPa from 768 MPa of the as-built sample. However, the ultimate tensile strength (UTS) of A900 has increased
from 1016 to 1153 MPa with annealing. It is remarkable that an
increase in strength is observed in response to annealing at 900 °C
because recovery and grain growth usually occur during heat treatment at such high temperatures only to reduce strength. Ductility
has diminished with the increase in UTS, but uniform elongation
of 20.5 % remains at a similar level which was not greatly reduced
from 25.0 % before annealing. The tensile properties are summarized in Table 1. Fig. 5(b) shows the strain hardening rates (SHR) of
the samples plotted against true strain. A900 exhibits a higher SHR
than the as-built sample in the early stage of plastic deformation,
which gives rise to the enhanced UTS despite the slightly reduced
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Journal of Materials Science & Technology 214 (2025) 143–152
Table 2
Diffusion coefficient of each substitutional element in the FCC matrices of
Co18 Cr15 Fe50 Ni10 Mo6.5 C0.5 MEA and (CoCrFeMnNi)99 C1 HEA at 900 °C (m2 /s).
MEA
HEA
Co
Cr
Fe
Mn
Ni
Mo
6.66E-20
1.68E-17
1.10E-17
9.85E-17
9.67E-18
5.49E-17
–
2.27E-16
1.83E-18
1.54E-17
1.79E-18
–
treatment have led to an abrupt increase in grain size from 35.5 to
56.5 μm in (CoCrFeMnNi)99 C1 HEA, which is well-known for the
“sluggish diffusion” [32,38]. We attributed the further retardation
effect of grain growth to the slow diffusion kinetics owing to the
chemical composition of the MEA. For a quantitative account for
the diffusion kinetics, the diffusion coefficient of each element in
the present Co18 Cr15 Fe50 Ni10 Mo6.5 C0.5 MEA and (CoCrFeMnNi)99 C1
HEA at 900 °C was calculated using DICTRA software with an updated version of the TCFE20 0 0 thermodynamic database [39–41]
and atomic mobility data from the literature [42–46]. It is noted
that the bulk composition of each alloy and only the substitutional
elements were considered in the calculation due to the lack of
atomic mobility data for C. Table 2 comparatively shows the diffusion coefficients of the elements constituting each alloy. In the
Co18 Cr15 Fe50 Ni10 Mo6.5 C0.5 MEA, the diffusion coefficients are calculated to be generally 1–3 orders of magnitude smaller than those
of the (CoCrFeMnNi)99 C1 HEA, which strongly supports the experimental observations. This demonstrates that the inherent nature
of the present MEA, stemming from its chemical composition, facilitates deliberate control of the cell structure to achieve the desired microstructure. Such characteristics of the MEA can be taken
into consideration when designing alloys for additive manufacturing, as the deliberate retardation of the grain growth under heat
treatment by harnessing the precipitation and slow diffusion kinetics can be useful in improving mechanical performance or hightemperature stability of additively manufactured components.
Fig. 6. Multiple strengthening contributions to the YS of the as-built and A900
samples.
Hall-Petch strengthening is calculated as 148 MPa for the as-built
sample and 144 MPa for A900.
The hard μ phase and M6 C type carbides are known to be
non-shearable and contribute to precipitation strengthening by the
Orowan mechanism [11,12], which is expressed as follows [48]:
σppt = 0.538
(1)
(2)
where M (3.03 for the as-built sample, 3.00 for A900) is the Taylor
factor obtained from the EBSD analysis, α = 0.2 is a constant, G is
the shear modulus (∼87.6 GPa [12]), b = ∼0.256 nm is the magnitude of Burgers vector, and ρ is the dislocation density. The ρ values of the as-built and A900 samples are notably higher than those
of normal cast or annealed materials and induce substantial Taylor
strengthening of 173 MPa for the as-built sample and 141 MPa for
A900.
The grain boundaries generate a strengthening effect by the
Hall-Petch relationship written as follows:
σgb = k · d− 2 ,
1
(4)
1/2
where σ0 is the friction stress, and σdis , σgb , σppt , and σCE represent the contributions of dislocation strengthening, grain boundary strengthening, precipitation strengthening, and extra strengthening by the cell boundaries, respectively. For σ0 , 181 MPa of
Co17.5 Cr12.5 Fe55 Ni10 Mo5 (at%) MEA can be borrowed [47].
The σdis can be formulated by the Taylor equation as follows:
1
,
equation of D = ( 32 ) D0 . f and D0 were measured via 2D image analysis using ImageJ software. The σppt is calculated to be
150 MPa for the as-built sample and 210 MPa for A900.
The summation of the strengthening contributions listed so far
gives us YS of ∼652 MPa for the as-built sample and ∼676 MPa
for A900, leaving ∼116 MPa for the as-built sample. The extra
strengthening effect by the cell boundaries can account for the
remnant strengthening contribution for the as-built sample. The
cell boundaries can act as mild obstacles for dislocation motion
that generates a strengthening effect. In order to explain the effect
of the cell boundaries on mechanical properties, numerous studies have been performed on various types of materials fabricated
by LPBF [49–51]. Because the misorientation angle (< 1°–2°) across
cell boundaries is much lower than typical high-angle grain boundaries (> 15°), several studies have proposed that the cell boundaries are not as strong as grain boundaries [52,53]. Meanwhile, it
has been suggested that the ability of the cell boundaries to impede dislocation glide depends heavily on the types of segregated
elements at cell boundaries in a similar way to the concept of grain
boundary segregation engineering [21,28]. In the present as-built
sample, the cell boundaries are decorated with heavy segregation
of Mo, which further enhances their cohesion [28]. Therefore, it is
assumed that the cell boundaries of the as-built sample can create
an extra strengthening effect further than what can be explained
by the simple dislocation strengthening. The multiple strengthening contributions for the samples are visualized in Fig. 6.
The hierarchically heterogeneous microstructures of the samples activate multiple strengthening mechanisms that contribute to
the YS of the samples. The strengthening contributions for the asbuilt sample can be equated as follows:
σdis = Mα Gbρ 2 ,
D
where f is the volume fraction, and D is the real spatial diameter of the precipitates calculated from the average diameters (D0 ,
45.8 nm for the as-built sample, and 87.4 nm for A900) by the
4.2. Strengthening mechanisms
σYS = σ0 + σdis + σgb + σppt + σCE ,
Gb f 0.5 ln Db
4.3. Strain hardening mechanisms
The most interesting feature of A900 is the remarkable strain
hardening capacity, which enables the excellent UTS of ∼1.15 GPa
greatly enhanced from that of the as-built sample despite the
lower YS. Kinematic hardening related to back stress stemming
from the pile-up of geometrically necessary dislocations (GNDs)
at the interfaces between regions with different strengths can
(3)
where k is the Hall-Petch coefficient, and d is the average grain
size. Using the Hall-Petch coefficient of ∼621 MPa μm1/2 [47], the
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Journal of Materials Science & Technology 214 (2025) 143–152
Fig. 7. (a) LUR stress-strain curves of the as-built and A900 samples. (b) A magnified portion from (a). The evolutions of (c) back stress and effective stress, and (d) the
contribution to the flow stress with respect to true strain.
creasing ε loc , especially at the grain boundaries and along the epitaxially grown cell columns at the grain interiors. The GND density calculated from the EBSD analysis (Fig. 9) is higher in A900
than in the as-built sample, and the difference between the two
samples becomes greater as deformation proceeds. It can be thus
concluded that a larger amount of GND pile-up in A900 leads to
higher back stress hardening.
Generally, improvement in strain hardening in ferrous MEAs
with relatively low stacking fault energy has been often related
to the activation of additional deformation mechanisms other
than dislocation slip, such as twinning-induced plasticity (TWIP)
of transformation-induced plasticity (TRIP) [7]. The deformation
twins can block dislocation motion by decreasing the dislocation
mean free paths. Deformation-induced martensite can lead to dynamic stress-strain partitioning between the parent FCC phase
and newly born BCC or HCP martensite, generating TRIP. To investigate whether such additional deformation mechanisms come
into play for the present samples, we have obtained the EBSD
boundary maps from the deformation microstructures (Fig. S5).
The boundary maps show that the fraction of deformation twins
did not show a significant difference between the samples. The
deformation-induced martensite was not observed either, which
means that the dislocation slip is the main mechanism that governs the deformation behaviors of the present samples. Therefore,
the difference in strain hardening behaviors is relevant to the hindrance of dislocation movement by the precipitates or grain/cell
boundaries.
In our previous study on the as-built sample [28], we discovered that the cell boundaries decorated by Mo segregation and
μ phase nanoprecipitates greatly contribute to strain hardening
contribute largely to strain hardening, especially in materials
with heterogeneous microstructures. The LPBF-built materials with
multi-scale heterogeneous microstructures therefore tend to exhibit pronounced back stress hardening. The back stress, σb , can
be quantified by loading-unloading-reloading (LUR) tensile testing
using the following equation.
σb =
σu + σr
2
,
(5)
where σu and σr are the YS during unloading and reloading at a
specific true strain, respectively. The contribution of effective stress
can be calculated by subtracting the back stress from the flow
stress. Fig. 7(a) shows the LUR curves of the as-built and A900
samples. A magnified portion from the LUR hysteresis loops shows
that the loop of A900 is wider than that of the as-built sample,
which indicates higher back stress hardening (Fig. 7(b)). The fitting from the elastic regimes and yield stress calculation results
from the loops are also presented in Fig. 7(b). The evolution of
back stress attained from the LUR curves is presented with respect
to true strain (Fig. 7(c)). A900 shows higher back stress and its
contribution to the flow stress calculated by dividing back stress
by the effective stress is greater than that of the as-built sample
(Fig. 7(d)) throughout the entire plastic deformation. This indicates
that the excellent strain hardening capacity of A900 is supported
by the pronounced back stress hardening.
To explain the mechanism behind the back stress hardening behavior, an EBSD analysis of the deformation microstructures was
conducted. Fig. 8 shows the EBSD GND maps taken at regions that
correspond to different local strain (ε loc ) values of 5 %, 10 %, and
15 %. The GND maps show an increasing amount of GNDs with in149
H. Kwon, E.S. Kim, Y.-U. Heo et al.
Journal of Materials Science & Technology 214 (2025) 143–152
Fig. 8. EBSD analyses of the deformation microstructures of the (a) as-built and (b) A900 samples. GND maps are taken from regions with specific ε loc values based on the
DIC strain distribution maps. Note that high-angle grain boundaries (HAGBs) are marked by black lines in the EBSD maps. The color bar in the EBSD maps presents the range
of GND density from 1 to 400 × 1012 m−2 . A step size of 90 nm was used for the EBSD analyses.
when compared to the Mo-free counterpart. It is interesting that
A900, where most cell boundaries have transformed into precipitates, shows more pronounced strain hardening than the as-built
sample. The different initial dislocation densities can partly explain
the strain-hardening behavior, as the high dislocation density of
the as-built sample indicates a higher driving force for dynamic recovery. Another reason can be found in the precipitates and their
role in pinning the dislocation movement.
STEM images and TKD KAM maps taken from the region deformed to ε loc = 15 % are presented in Fig. 10 to enable nanoscale
examination of the GND pile-up. The STEM images (Fig. 10(a))
show profuse dislocation pile-up in both samples, mostly near
the cell boundaries and precipitates. Particularly, a higher amount
of dislocation accumulation at the precipitates is clearly observed
in A900 than in the as-built sample. The quantitative evaluation of the KAM values based on the TKD analysis confirms that
150
H. Kwon, E.S. Kim, Y.-U. Heo et al.
Journal of Materials Science & Technology 214 (2025) 143–152
where μ is the unit length (10–5 m [54]), and b is the magnitude
of the Burgers vector. The ρGND is calculated as ∼5.86 × 1014 m–2
for A900 and ∼4.84 × 1014 m–2 for the as-built sample. Therefore,
the higher fraction of precipitates is the main contributor to the
higher amount of the GND pile-up and corresponding back stress
hardening of A900. The μ phase and M6 C type carbides, which are
much harder than the FCC matrix [55,56], lead to severe strain
incompatibility under deformation. The GNDs accumulate at the
interfaces to accommodate the strain incompatibility and become
the source of back stress hardening. Additionally, it is also possible that the profuse cell boundaries confine space for dislocation
storage similar to the grain boundaries in ultrafine-grained materials [57,58], which limits the amount of dislocation accumulation in
the as-built sample as the deformation proceeds when compared
with A900 without the spatial confinement by the cell boundaries.
Fig. 9. Average GND density measured from the EBSD analyses at regions with specific ε loc values in the as-built and A900 samples.
5. Conclusion
A900 exhibits higher average KAM values of ∼0.75°, while the asbuilt sample shows ∼0.62°. The KAM maps also demonstrate a
high density of GND accumulation near the precipitates in A900
(Fig. 10(b)). The KAM values (θ ) can be converted directly into GND
density (ρGND ) by the following equation [54]:
LPBF-processed Co18 Cr15 Fe50 Ni10 Mo6.5 C0.5 MEA with cell
boundaries decorated by high-density dislocations and segregation of Mo was subjected to post-LPBF annealing at 900 °C
for 10 min. Due to the inherently slow diffusion kinetics of the
MEA, the post-LPBF annealing could restructure the nanoscale cell
structure by replacing the original cell boundaries with μ phase
and M6 C-type carbide precipitates without significantly affecting
the grain structure. The alteration of the cell structure resulted
ρGND = 2θ /μb,
(6)
Fig. 10. TEM and TKD micrographs of the deformed microstructures of the as-built and A900 samples at ε loc = 15 %. (a) STEM images of the (a1 ) as-built and (a2 ) A900
samples. (b) TKD KAM maps of the (b1 ) as-built and (b2 ) A900 samples. A step size of 10 nm was used for the TKD analyses.
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Journal of Materials Science & Technology 214 (2025) 143–152
in different mechanical responses under a tensile test. The YS
decreased from 768 to 656 MPa due to the reduction in the
strengthening effect by the cell boundaries, although the Orowan
strengthening effect by the precipitates increased with annealing.
On the other hand, the UTS increased from 1016 to 1153 MPa
supported by the greatly enhanced strain hardening ability, while
uniform ductility of 20.5 % was not significantly reduced from
25 %. The enhancement in the strain hardening by annealing is
attributed to the pronounced back stress hardening due to the
formation of hard precipitates that effectively accumulate the
GNDs to accommodate the strength difference between the FCC
matrix. Our work suggests that the dislocation cell structures with
chemical heterogeneity, which is characteristic of LPBF-processed
materials, can be successfully harnessed to control precipitation
and corresponding mechanical performance. We believe that our
strategy to reform the cell structure can provide insights into
tailoring microstructures to acquire desired properties in a wider
range of materials by applying post-LPBF heat treatment that is
pertinent to each material. Furthermore, our approach could offer
insight into designing HEAs and MEAs tailored for additive manufacturing, as an example where the slow grain growth kinetics and
chemical heterogeneity are exploited to the utmost.
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Declaration of competing interest
The authors declare that they have no known competing financial interests or personal relationships that could have appeared to
influence the work reported in this paper.
CRediT authorship contribution statement
Hyeonseok Kwon: Writing – original draft, Investigation, Conceptualization. Eun Seong Kim: Writing – review & editing,
Methodology, Formal analysis. Yoon-Uk Heo: Investigation, Formal
analysis. Jungho Choe: Methodology, Data curation. Rae Eon Kim:
Visualization, Methodology. Soung Yeoul Ahn: Software, Investigation. Sang-Ho Oh: Software. Jeong Min Park: Writing – review &
editing, Validation, Resources. Byeong-Joo Lee: Validation, Supervision. Hyoung Seop Kim: Writing – review & editing, Validation,
Supervision, Resources, Conceptualization.
Acknowledgments
This work was supported by the National Research Foundation of Korea (NRF) grant funded by the Korean government
(MSIT) (Nos. 2021R1A2C3006662 and RS-2023-00281246). This
study was also supported by the Principal R&D project (contract
no. PNK9950) of the Korean Institute of Materials Science (KIMS).
Supplementary materials
Supplementary material associated with this article can be
found, in the online version, at doi:10.1016/j.jmst.2024.07.010.
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