Scripta Materialia 215 (2022) 114706
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Scripta Materialia
journal homepage: www.journals.elsevier.com/scripta-materialia
In-situ observations of static recrystallization and texture formation in a
cold-rolled CoCrFeMnNi high entropy alloy
Baoqi Guo a, *, Ranjit K. Ray a, b, Shuhei Yoshida a, Yu Bai a, c, Nobuhiro Tsuji a, d
a
Department of Materials Science and Engineering, Kyoto University, Yoshida-Honmachi, Sakyo-ku, Kyoto 606-8501, Japan
Indian Institute of Engineering Science and Technology, Shibpur, 711103, India
c
School of Materials Science and Engineering, Dalian University of Technology, Dalian 116024, China
d
Elements Strategy Initiative for Structural Materials (ESISM), Kyoto University, Yoshida-Honmachi, Sakyo-ku, Kyoto 606-8501, Japan
b
A R T I C L E I N F O
A B S T R A C T
Keywords:
Texture
Recrystallization
High entropy alloy
In situ TEM
Static recrystallization and texture development in a CoCrFeMnNi high entropy alloy cold-rolled to various
strains and annealed were studied. Microstructural and crystallographic evolution was recorded by identical area
observations using quasi in-situ electron backscatter diffraction (EBSD). The deformation texture was preserved
after recrystallization, but the texture was weakened at this stage. Shear bands served as the preferential
nucleation sites for recrystallization. Thin foils of the deformed alloy were in-situ observed during heating in a
transmission electron microscope (TEM). It was observed that subgrains formed from deformed microstructures
by dislocation rearrangement and annihilation. The subgrains coarsened by boundary migration and coalescence.
These enlarged subgrains were the viable nuclei for the recrystallized grains.
High entropy alloys (HEAs) are composed of five or more principal
elements in equimolar ratios. The HEAs have attracted wide attentions
due to their excellent mechanical properties, such as high strength and
wear resistance as well as outstanding fracture toughness at cryogenic
temperatures [1–5]. Properties of metals and alloys are primarily
affected by microstructures and grain orientations [6,7]. Since industrial
manufacturing of HEAs will entail, almost invariably, thermomechan­
ical processing, it is of great significance to investigate the evolution of
microstructure and texture in HEAs during deformation and
recrystallization.
Evolution of texture during static recrystallization of a cold-rolled
CoCrFeMnNi high-entropy alloy has been described by Bhattacharjee
et al. [8]. It was demonstrated that brass-type texture formed after heavy
cold rolling by 90%, and annealing twins developed during annealing.
Other factors such as grain size and heating rate have been considered in
the study of annealing textures in a CoCrFeMnNi alloy [9]. It was shown
that heating rate could determine the grain size as well as its distribution
during annealing while they rarely affected the texture evolution during
annealing [10]. Recently, the work by Haase and Barrales-Mora [11]
showed that the microstructure and texture evolution during recrystal­
lization in a CoCrFeMnNi alloy were controlled by nucleation at grain
boundaries and annealing twin formation. Recrystallized grains were
found to show rolling texture components and annealing twins gave rise
to an overall texture randomization, although the mechanisms of texture
formation and recrystallization were still unknown at the time.
The previous works have focused on textural changes during
annealing, but the relationship between recrystallization texture and
deformation microstructures, such as the role of shear bands in recrys­
tallization texture formation, has been missed. Moreover, the afore­
mentioned studies with regard to recrystallization and textures were
carried out by means of ex-situ methods (i.e. observation of quenched
microstructure). Due to the limitation of ex-situ techniques, the precise
mechanisms of the recrystallization process in HEAs are still unclear.
The objective of the present work is to bridge the gaps by elaborating on
the nucleation and growth of recrystallized grains in the CoCrFeMnNi
alloy, and the concomitant texture evolution, using in-situ and identicalarea experimental techniques.
An as-cast equi-atomic CoCrFeMnNi HEA was forged at 1150◦ C and
homogenized at 1100◦ C for 24 h. The materials were subsequently coldrolled to 50% reduction in thickness followed by annealing in vacuum at
1100◦ C for 5 min. The annealed samples were further cold-rolled to
50%, 70% and 90% reductions. Subsequent annealing of the cold-rolled
sheets was carried out at 750◦ C for various periods of time, followed by
water quenching.
Microstructures on sections perpendicular to the transverse direction
(TD) of the cold-rolled sheets were observed by electron backscatter
* Corresponding author.
E-mail address: guo.baoqi.7z@kyoto-u.ac.jp (B. Guo).
https://doi.org/10.1016/j.scriptamat.2022.114706
Received 20 February 2022; Received in revised form 18 March 2022; Accepted 22 March 2022
Available online 29 March 2022
1359-6462/© 2022 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
B. Guo et al.
Scripta Materialia 215 (2022) 114706
measuring area fractions of texture components using orientation im­
aging microscopy (OIM) software by TSL co. Identical areas in the
specimens cold-rolled and annealed for different periods were observed
by EBSD, which is called quasi in-situ EBSD observations hereafter. In the
quasi in-situ EBSD observations of identical areas, fiducial marks were
firstly made on the surfaces of the cold-rolled samples. The samples were
diffraction (EBSD) in a field-emission scanning electron microscope (FESEM) and transmission electron microscopy (TEM). The EBSD specimens
were electropolished using electrolyte containing 90 vol.% ethanol and
10 vol.% perchloric acid at room temperature with a voltage of 30 V for
15 s. Texture components were identified using Euler angles determined
by the Bunge method. The texture intensities were obtained by
Fig. 1. EBSD micrographs containing (a) an inverse pole
figure (IPF) map, (b) image quality (IQ) map, and (c) Kernel
average misorientation (KAM) map showing deformation
microstructures in the 50% cold-rolled CoCrFeMnNi alloy. IPF
maps of an identical area in the 50% cold-rolled specimen
obtained by the quasi in-situ EBSD observations during
annealing at 750 ◦ C for (d) 0 min (as cold-rolled), (e) 8 min,
(f) 12 min, (g) 17 min, (h) 27 min, (i) 40 min, respectively.
Note that Fig. (d) is the same as Fig (a). Colors in the IPF map
indicate crystallographic orientation parallel to transverse
direction (TD) according to the key stereographic triangle
inserted. ND and RD indicate the normal direction and rolling
direction, respectively, in the cold-rolled specimens.
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Scripta Materialia 215 (2022) 114706
wrapped by titanium foils to minimize oxidation and repeated annealing
at the same annealing temperature was conducted in a vacuum furnace.
Water quenching was employed after each annealing operation.
Thin foils for TEM were prepared by twin-jet electro-polishing at a
voltage of 15 V and -30◦ C using electrolyte with a composition of 70 vol.
% methanol, 20 vol.% glycerin, and 10 vol.% perchloric acid. A TEM
(JEM-3200FSK) and a specimen holder equipped with a heating stage
were used in the in-situ TEM observations. The experiment was carried
out at an accelerating voltage of 200 kV. A low heating rate of 0.5◦ C s− 1
was used during annealing to reduce the drift of specimen caused by
Fig. 2. Texture components of the identical area (the same area as shown in Fig. 1) in the 50% cold-rolled and 750◦ C annealed specimen. The colors of texture
components are consistent with those in the table. The white color denotes random texture. The black color regions from (b) to (f) refer to unrecrystallized parts. The
specimen was annealed at 750◦ C for (a) 0 min (as cold-rolled), (b) 8 min, (c) 12 min, (d) 17 min, (e) 27 min and (f) 40 min.
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Scripta Materialia 215 (2022) 114706
heating. The samples were heated from room temperature (20◦ C)–750◦ C
and then isothermally held for 30 min.
Microstructures of the CoCrFeMnNi HEA after 50% cold-rolling are
shown in Fig. 1(a)–(c). In-grain shear bands frequently formed in the
cold-rolled specimens, as revealed in IPF map (a) and IQ map (b). The
Kernel average misorientation (KAM) map (c) indicated large local
misorientations and strains on the shear bands and heterogeneously
deformed local areas with large local misorientations. The shear bands
were more copious and well-delineated at the higher rolling reduction of
90%, as shown in the TEM image provided in the supplementary file
(Fig. S1). Here, it was confirmed that the shear bands penetrated twin­
ned areas (or twin-matrix lamellae) composed of very fine deformation
twins and matrices aligned parallel to the rolling direction (RD).
Fig. 1(d–i) shows EBSD micrographs of an identical area in the 50%
cold-rolled specimen (the same as Fig. 1(a–c)) taken in the quasi in-situ
observations during recrystallization annealing at 750◦ C for different
times. Sites “A”, “B” and “C” in the as-rolled (0 min. annealed) micro­
structure represent three shear bands and the region “D” surrounded by
broken lines indicates a heterogeneously deformed local area with large
local misorientations. After annealing for 3 min, there was no apparent
change in the microstructure. Then, recrystallization occurred on the
shear bands and in the heterogeneously deformed local areas, when the
sample was annealed for 8 min. The recrystallization nuclei became
more copious after annealing for 12 min. Further nucleation and growth
Fig. 3. EBSD maps obtained by the quasi in-situ observations of the specimens 70% cold-rolled and annealed at 750◦ C. IPF maps (a, d, f), texture component maps (b,
e, g) and KAM map (c) of the as-70%-rolled state (a-c) and those annealed at 750◦ C for 3 min (d, e) and 6 min (f, g). The colors in the texture components maps (b, e,
g) correspond to different texture components summarized in the table. The white color in texture component images denotes random texture.
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Scripta Materialia 215 (2022) 114706
of recrystallized grains occurred with increasing the annealing time.
The texture components in the 50% cold-rolled and 750◦ C annealed
specimen shown in Fig. 1 are displayed in Fig. 2. The areas of which
orientations were close to typical texture components in cold-rolled and
annealed alloys (e.g., Brass {110}<112> and Goss {110}<001> ori­
entations) were exhibited in different colors, as indicated in the table
inserted. Here texture characterization allows for a misorientation of
15 ◦ from the ideal orientation of the typical texture component. It
should be noted that the Goss/Brass in olive color refers to {110}<115>
orientation that is in the middle between the Brass and Goss
orientations. When the orientation of the area did not belong to any of
these typical texture components, it was categorized as “Random”
orientation and colored in white. It was found that the initial recrys­
tallization nuclei possessed orientation components which were present
in the cold-rolled microstructure. The shear bands “A” and “B” in the as
cold-rolled state were characterized by Goss (orange) and S (light blue)
orientation components. A magnified image of Fig 2(a) that shows
detailed texture components is provided in the supplementary file
(Fig. S2). As recrystallization progresses, some recrystallized grains
appeared on the shear bands “A” and “B” with Goss and S orientations,
Fig. 4. Images of an identical area obtained from the in-situ TEM observation of nucleation during annealing. Bulging of the dislocation boundary was detected at
450◦ C in Fig. (d).
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Scripta Materialia 215 (2022) 114706
In order to clarify the nucleation sites for recrystallization at early
stages, the quasi in-situ EBSD observations of the 70% cold-rolled ma­
terial were carried out during annealing at 750◦ C for various times. The
EBSD maps of an identical area obtained are shown in Fig. 3. The IPF
map Fig. 3(a) of the cold-rolled specimen showed a banded structure
composed of initial grains elongated to the rolling direction (RD) and
deformation bands. Specifically, the region “X” was an elongated initial
suggesting that local textures formed by cold rolling were preserved and
developed during recrystallization. It should be noted, however, that
many recrystallized grains had Copper (blue) and Random (white) ori­
entations. The heterogeneously deformed area “D” contained Brass (red)
and Goss as well as S orientations in the cold-rolled state. Apart from
these cold rolled textures, other orientations such as Copper were found
in some recrystallized grains formed during annealing.
Fig. 5. Images of an identical area obtained from the in-situ TEM observation during annealing. Three pre-existing cells with irregular shapes formed by cold rolling
are indicated by red arrows in (a). With increasing the temperature, the three sites grew from (a) to (e).
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Scripta Materialia 215 (2022) 114706
grain with small misorientations inside Fig. 3(c), having Brass orienta­
tion (Fig. 3(b)). The regions “Y” and “Z” represented heterogeneously
deformed local areas with large local misorientations (Fig. 3(c)), which
showed Brass, S, Goss and rotated Goss, along with partly Cu and BrassGoss orientation components (Fig. 3(b)). During annealing, very small
recrystallized grains were found to nucleate in the heterogeneously
deformed local areas (regions Y and Z). Some of new recrystallized
grains nucleated at boundaries of elongated initial grains. The recrys­
tallization nuclei within the heterogeneously deformed local areas
formed from or near the regions with similar orientations in the asdeformed state, comparing Fig. 3(b) and (e). As indicated by the black
arrows, two small areas “A” and “B” (blue in IPF image, Fig. 3(a)) had
Brass orientation (Fig. 3 (b)) in the as-deformed state. These two areas
acted as nucleation sites to form new recrystallized grains in annealing
for 3 min (d), and most grains nucleated and grown had Brass orienta­
tion (Fig. 3(e)). In a similar manner, heterogeneously deformed areas
with Goss and Copper orientation in the “Z” region were formed by 70%
cold rolling and these areas developed into recrystallized grains having
their original texture components. The major texture components found
in the 6 min annealed specimen were similar to those of the cold rolled
material.
In order to get a deeper insight into the early stages of nucleation of
recrystallized grains, in-situ heating experiments were conducted in the
heating stage of a TEM. A 90% cold-rolled specimen was heated from
room temperature to 750◦ C at a heating rate of 0.5◦ C/s. Fig. 4 (a)–(i)
showed a series of micrographs of an identical area, taken at different
stages of heating. The cold rolled thin foil showed a typical elongated
cell structure. The present observation focused on the area indicated by
the dashed red circle. Dislocation movement and interaction started
when the temperature reached 192◦ C. Annihilation of dislocations
continued during further annealing till 271◦ C. The bulging of the
dislocation boundary was activated from 450◦ C to 609◦ C. Finally, after
holding for 300s at 750◦ C, annihilation and rearrangement of disloca­
tions almost completed, forming a small recrystallization nucleus. Some
distinct movement of cell boundaries could also be observed.
Fig. 5(a)–(f) represent initial stages of the formation of recrystalli­
zation nucleus, adjacent to a heavily twinned area in the 90% cold rolled
alloy, obtained from the in-situ TEM observations. Fig. 5(a) shows three
small areas with nearly free of dislocation contrasts (which could be
described as pre-existing cells) close to the twin edges, marked as ‘X’, ‘Y’
and ‘Z’. When the temperature increased to 298◦ C, the boundaries of the
three sites became sharper and clearer due to dislocation rearrangement
and annihilation. At this temperature, the small subgrain at ‘X’ joined up
with another one below it, possibly by subgrain coalescence. With an
increase of temperature to 710◦ C, some growth of these subgrains could
be clearly seen, and they invaded adjacent twin-matrix bundles. The
subgrains at ‘X’, ‘Y’ and ‘Z’ showed some growth at this stage, and new
subgrains formed, typically in the area marked K, where a rather large
subgrain formed by dislocation annihilation and grain boundary
migration. The misorientation as well as dislocation density tended to
become high in the interface between twin bundle and deformed matrix,
which could facilitate the occurrence of nucleation of recrystallization.
What we have been describing as subgrains could be recrystallized
nuclei. When a subgrain enlarged, it built up sufficient misorientation
with surrounding areas, becoming a recrystallized grain.
The texture development in the current high entropy alloy during
annealing of the 70% cold-rolled alloy annealed at 750◦ C for different
periods of time is shown in supplementary material (Fig. S3(a)). Even
after 5 min of annealing, the texture intensity decreased drastically, and
this trend continued till recrystallization took place after annealing for
30 min. It should be noted that the volume fractions of texture com­
ponents of recrystallized grains remained approximately unchanged
during annealing (Fig. S3(b)). In Fig. S3(c), the fractions of Brass and
Goss components in the unrecrystallized initial grains increased during
annealing, which indicated that the grains with Brass and Goss com­
ponents were hard to recrystallize. In contrast, fractions of Copper and
Random textures in the unrecrystallized initial grains decreased with
increasing the annealing periods, suggesting that these components
were preferentially recrystallized and contributed to randomization of
the texture. Such a weak recrystallization texture has also been reported
previously in CoCrFeMnNi HEA by several researchers [8,11]. Their
works demonstrated that the basic characteristics of the textures of the
recrystallized grains and the un-recrystallized grains were essentially
similar. The texture of the recrystallized grains was the weakest, and
that of the un-recrystallized grains was the strongest. The overall texture
intensity of the entire sample located between these two extremes. This
analysis demonstrated that the deformation texture was preserved after
recrystallization, but the overall texture intensity was drastically
weakened, which could be further explained by the results obtained
from the in-situ TEM observations (Figs. 4 and 5). These experiments
have shed much light on the formation and growth of the subgrains from
the cold rolled dislocation structures. The subgrains formed were the
precursors to the new strain-free recrystallization nuclei. The current
investigations showed that dislocation rearrangement and annihilation
gave rise to the formation of nearly dislocation-free subgrains, which
coarsened by the classical processes like sub-boundary movement. The
reduced texture intensity during annealing is related to the recrystalli­
zation in shear bands. As shown in Figs. 1 and 2, cold rolling produced
numerous shear bands that acted as the preferential sites for recrystal­
lization. With few Goss and S orientations, the recrystallization textures
in shear bands appeared to be weak during annealing. Annealing twins
in FCC crystals contributes to weak textures during annealing as well
[12]. According to a recent work [13], annealing twins in CoCrFeMnNi
high entropy alloy generated new orientations such as Cube orientation
that weakened the textures during recrystallization.
In addition, the similarity in the nature of the texture from the
deformed, as well as from the recrystallized materials [8,11], indicated
that the process was oriented nucleation. That is, the nuclei of recrys­
tallized grains of particular orientations form at or near the deformed
regions of same or similar orientations. The oriented nucleation was
further illustrated during the recrystallization via evolution of
pre-existing cells indicated in Figs. 3 and Fig. 5. In-situ TEM has shown
that the processes of subgrain formation and their growth are rather
slow. It has been claimed that diffusion is sluggish in high entropy alloys
[2]. For example, Tsai et al. [14] concluded that the diffusion co­
efficients of the individual elements in the CoCrFeMnNi alloy were
indeed lower than those in conventional FCC metals. Although clarifi­
cation of diffusion mechanisms in HEAs requires further study, sluggish
diffusion could be closely related to the slow growth of recrystallized
grains in these alloys. Since the grain boundary migration involves
jumps of atoms (diffusion) across grain boundaries, elements with a slow
diffusion rate can dominate the recrystallization process. In fact, Liu
et al. [15] and Tsai et al. [14] reported that the activation energy of
grain growth in the same HEA was found to be comparable to that of the
lattice diffusion of Ni with the slowest diffusion rate among all the
constituent elements. The solute drag effect due to the interaction be­
tween elements and grain boundaries can also be another important
factor. As different elements occupy various lattice sites in the HEA,
some specific elements with very low potential energy can be trapped
and dragged by grain boundaries, leading to deceleration of grain
boundary migration kinetics. In conventional alloys, this solute drag
effect is evident in low-angle boundaries, but it is possible that the effect
in HEA can inhibit boundary migration even in high-angle boundaries.
The slow growth rate of the recrystallized grains will not allow grains of
any particular orientation or grains of several orientations to grow
preferentially at the expense of grains of other orientations. This will
result in the formation of a fine grain size and random texture in the
recrystallized material, as has been found to be the case.
In summary, in-situ TEM and quasi in-situ EBSD observations of
identical areas were conducted for clarifying static recrystallization and
texture formation during annealing of a CoCrFeMnNi equi-atomic high
entropy alloy subjected to various amounts of cold rolling. The general
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Scripta Materialia 215 (2022) 114706
characteristics of the cold rolling texture was preserved even after
recrystallization, but with a drastically decreased intensity. The initial
micro-scale mechanisms of the formation of dislocation-free subgrains
(which were the precursors of the viable recrystallization nuclei) were
rearrangement and annihilation of dislocations, sub-boundary migra­
tion and subgrain coalescence. Shear bands acted as the preferential
sites for recrystallization and randomized the recrystallization textures.
The mechanism of recrystallization in this alloy could be understood as
“oriented nucleation” followed by limited growth of recrystallized
grains, probably due to slow diffusion in the alloy.
was also financially supported by the Grant-in-Aid for Research Activity
Start-up (No. JP21K20487), and the Grant-in-Aid for JSPS Research
Fellow (No. JP18J20766). All the supports are greatly appreciated.
Supplementary materials
Supplementary material associated with this article can be found, in
the online version, at doi:10.1016/j.scriptamat.2022.114706.
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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.
[5]
[6]
[7]
Acknowledgment
[8]
The in-situ TEM observations in this work were conducted under
Advanced Characterization Platform of the Nanotechnology Platform
Japan (JPMXP09-A-19-KU-0324). This work was financially supported
by the Elements Strategy Initiative for Structural Materials (ESISM; No.
JPMXP0112101000), the Grant-in-Aid for Scientific Research on Inno­
vative Area "High Entropy Alloys" (No. JP18H05455), the Grant-in-Aid
for Scientific Research (A) (No. JP20H00306), through the Ministry of
Education, Culture, Sports, Science and Technology (MEXT), Japan. SY
[9]
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