Microstructure evolution and strengthening mechanisms
of Fe–23Mn–0.3C–1.5Al TWIP steel during cold rolling
Pavel Kusakin
a,
n
, Andrey Belyakov
a
, Christian Haase
b
, Rustam Kaibyshev
a
,
Dmitri A. Molodov
b
a
Belgorod State University, Pobeda 85, Belgorod 308015, Russia
b
Institute of Physical Metallurgy and Metal Physics, RWTH Aachen University, 52056 Aachen, Germany
article info
Article history:
Received 14 May 2014
Received in revised form
30 July 2014
Accepted 18 August 2014
Available online 28 August 2014
Keywords:
TWIP-steel
Electron microscopy
Microstructure
Fractography
Hardening
Mechanical characterization
abstract
The effect of cold rolling on the microstructure evolution and mechanical properties of Fe–23Mn–0.3C–
1.5Al twinning-induced plasticity (TWIP) steel was studied. The extensive mechanical twinning
subdivides the initial grains into nanoscale twin lamellas. In addition, the formation of deformation
micro bands at
ε
4 40% induces the formation of nanostructured bands of localized shear. It is
demonstrated that the mechanical twinning is notably important for dislocation storage within the
matrix, as the twin boundaries act as equally effective obstacles to dislocation glide as conventional
high-angle grain boundaries. However, the contribution of the grain size strengthening to the overall
yield stress (YS) is much smaller than that of the deformation strengthening, which plays a major role in
the superior work-hardening behavior of TWIP steels. A very high dislocation density of 2 10
15
m
2
is achieved after plastic deformation with moderate strains. The superposition of deformation
strengthening and grain boundary strengthening leads to an increase in the YS from 235 MPa in the
initial state to 1400 MPa after 80% rolling.
& 2014 Elsevier B.V. All rights reserved.
1. Introduction
There is a considerable demand for advanced high-strength
steels with high formability for structural parts in automobile
bodies [1]. Austenitic steels with high Mn content exhibit the most
attractive combination of tensile strength and superior ductility
due to their extraordinary strain-hardening rate, which is inter-
preted in terms of twinning-induced plasticity (TWIP) [2–6].Itis
known that in face-centered cubic (fcc) metals with low stacking
fault energies (SFEs),extensive deformation twinning results in the
formation of deformation twins with a plate-like shape as well as
nanometer spacing and thickness [2,7–9]. The twin boundaries,
which are special (low
Σ
coincidence site lattice) high-angle
boundaries [7] and contain a high density of sessile dislocations
[8], act as equally effective obstacles to dislocation gliding as
conventional grain boundaries in polycrystalline austenitic steels
[10]. The subdivision of initial grains into lamellas with thick-
nesses ranging from 10 to 40 nm [2–6,8] leads to a significant
decrease in the effective grain size and, therefore, results in
remarkable strengthening in accordance with the Hall–Petch
relationship [11–13]:
σ
0:2
¼σ
0
þK
H
d
0:5
ð1Þ
where
σ
0.2
is the offset yield strength; d is the effective grain size,
which is the distance between barriers to dislocation glide and can
be estimated to be twice the width of lamellae [14];
σ
0
is the
friction stress; and K
H
is the Hall–Petch coeffi cient. A “ dynamic
Hall–Petch effect” [2–6,8] attributed to the reinforcement of
austenite by deformation twins was exploited to explain the
extraordinary ductility of the TWIP steels [9,15]. However, it is
apparent that this interpretation of the very high strain-hardening
rate of TWIP steels in terms of the dynamic Hall–Petch effect is
oversimplified and open to debate.
Generally, an increase in the yield strength of TWIP steels by
cold working is associated with grain boundary strengthening,
strain hardening and dynamic strain aging (DSA) [2–6,8,9,15–17].
In high-Mn TWIP-steels, DSA is a type of solid solution strength-
ening attributed to the formation of interstitial C – substitutional
Mn dipoles interacting strongly with dislocations or stacking
faults. In turn, grain boundary strengthening may be caused by
different mechanisms [13,18] strongly related to dislocation
strengthening. Grain boundary strengthening due to mechanical
twinning is associated with the accumulation of lattice disloca-
tions in pile-ups to attain critical stress concentration, which is
Contents lists available at ScienceDirect
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Materials Science & Engineering A
http://dx.doi.org/10.1016/j.msea.2014.08.051
0921-5093/& 2014 Elsevier B.V. All rights reserved.
n
Corresponding author. Tel.: +7 4722 585409; fax: +7 4722 585406.
E-mail addresses: kusakin@bsu.edu.ru (P. Kusakin),
belyakov@bsu.edu.ru (A. Belyakov), haase@imm.rwth-aachen.de (C. Haase),
rustam_kaibyshev@bsu.edu.ru (R. Kaibyshev),
molodov@imm.rwth-aachen.de (D.A. Molodov).
Materials Science & Engineering A 617 (2014) 52–60
necessary to induce the operation of the dislocation sources in
neighboring grains or the boundary dislocation sources [13,18].In
the original Hall–Petch hypothesis, the dislocation pile-ups form at
the grain boundary [11–13,18]. The stress concentration at the
pile-up fronts, which is directly proportional to the number of
piled-up dislocations and dislocation storage within the crystal-
lites outlined by twin boundaries, increases the yield stress (YS)
with decreasing grain size. It is worth noting that materials with
nanocrystalline structures may exhibit a deviation from the Hall–
Petch relationship [19]. The Hall–Petch relationship was shown to
be valid for lamellar structures with an interboundary spacing of
20 nm in which the pile-ups involve approximately 10 disloca-
tions within a single lamella [14] and the length of the pile-up is
equal to the spacing between twin boundaries [15]. However, the
K
H
value tends to decrease with decreasing interlamellar spacing,
which can further lead to the so-called inverse Hall– Petch effect
when the effective grain size falls below 10 nm [20,21].
The dislocation strengthening provides a major contribution to
the overall strength of TWIP steels due to the storage of a very
high dislocation density of
ρ
10
15
m
2
during plastic deforma-
tion [2,22]. In general, the dislocation density in TWIP steels is one
order of magnitude higher than that in other fcc metals deformed
to the same strain [8]. The extensive mechanical twinning sub-
divides the initial grains into separate crystallites bounded by twin
boundaries, which reduces the mean free path of lattice disloca-
tion, leading to extraordinary accumulation of lattice dislocations
[2,5,8,22,23]. In addition, the DSA and the slip–twin interactions
hinder dynamic recovery [8]. The slip–slip and slip–twin interac-
tions enable the formation of numerous sessile dislocations, which
act as effective barriers for mobile dislocations and thus increase
dislocation hardening significantly [5,6,22,23]. Therefore, the
extensive mechanical twinning strongly promotes the dislocation
hardening through a progressive reduction of the mean free path
of dislocations. This builds up the dense dislocation pile-ups and
results in high long-range internal stresses [6], which act as back
stresses [9].
The main aim of the present work was to clarify the superior
strain-hardening capacity of Fe–23Mn–0.3C–1.5Al steel, which is a
typical representative of the advanced high-Mn TWIP steels. To
this end, the microstructure evolution during cold rolling and the
corresponding mechanical properties were studied. This steel
contains aluminum, which usually fully or partially suppresses
DSA at room temperature [17,24]. Therefore, only two strengthen-
ing mechanisms, grain boundary strengthening and dislocation
strengthening, may substantially contribute to the overall strain
hardening of this steel. Additionally, the aluminum additives
increase the twin thickness and suppress the secondary twinning
at high strains [4,6]. The twin thickness of approximately 30 nm is
well above the lower bound of the interboundary spacing that
allows the bowing of a dislocation loop across the twin. Therefore,
grain boundary strengthening is operative due to the increased
stress needed for dislocation propagation across the lamellar
section [14].
The major shortcoming of high-Mn austenitic steels is their
relatively low YS, which is associated with the recrystallized
microstructure that evolves after conventional thermo-mechanical
processing (TMP) [25,26]. Generally, the YS can be increased by an
appropriate cold working at the expense of ductility [25].The
development of technological routes involving cold working for
producing TWIP-steels with the beneficial combination of strength
and ductility requires detailed investigation of the mechanisms of
microstructure evolution during deformation and careful analysis of
the strain-hardening mechanisms. Recent studies on TWIP steels
with various manganese contents have revealed the common
sequence of structural changes during cold rolling [5,25,27,28].
Following a rapid increase in the dislocation density at an early
deformation, the deformation twinning progressively develops
throughout the deformation microstructures at low to medium
strains, whereas shear banding occurs at rather high strains. The
number of shear bands and their size gradually increase with
straining. The formation of shear bands at high strains is associated
with the rearrangement of twin lamellae along the rolling plane. In
the present work, the effect of strain on the deformation micro-
structure evolution and the mechanical behavior of Fe–23Mn–0.3C–
1.5Alsteelwasaddressed.
2. Experimental
The chemical composition of the investigated steel is given in
Table 1. The steel was melted in an air conduction furnace and cast
into a 30-kg ingot (140 mm in height), followed by homogeniza-
tion heat treatment at 1150 1C for 5 h. The ingot was forged at
1150 1C to a height of 50 mm and solution treated at 1150 1C for
5 h, followed by air cooling. The 50-mm-thick slab was hot rolled
at an initial temperature of 1150 1C to a thickness of 10 mm and
then annealed at the same temperature for 1 h. The obtained plate
was cold rolled to final thicknesses of 8, 6, 4 and 2 mm with rolling
reductions of 20, 40, 60 and 80%, respectively. The rolling direction
(RD) was the same under hot and cold rolling.
The structural investigations were carried out on the long-
itudinal section (RD: normal direction (ND)) of the rolled samples.
The sheets were ground on 1200, 2400 and 4000 grit SiC paper,
mechanically polished using 3- and 1-mm diamond suspensions,
and electropolished for 20 s at 22 V and 20 1C. The electrolyte
contained 700 ml ethanol (C
2
H
5
OH), 100 ml butyl glycol (C
6
H
14
O
2
)
and 78 ml perchloric acid (60%) (HClO
4
). Microstructure analysis
using electron backscatter diffraction (EBSD) was performed in a
LEO 1530 field emission gun electron microscope equipped with
an EBSD detector and operated at a 20-kV accelerating voltage.
A step size of 0.8 mm was used. For optical microscopy, the
specimens were etched at room temperature using an etching
solution consisting of 2 g potassium disulfide (K
2
S
2
O
5
) and 100 ml
cold saturated Klemm's I solution (Na
2
S
2
O
3
þ5H
2
O). For micro-
structural characterization by transmission electron microscopy
(TEM), thin foils of 3-mm diameter were cutout parallel to the RD-
ND plane and ground to a 0.1-mm thickness before being polished
using a double-jet TENUPOL-5 electrolytic polisher in a mixed
solution of 90% acetic acid and 10% perchloric acid at room
temperature and an applied potential of 20 V. The foils were
examined using a JEOL JEM-2100 TEM operated at an acceleration
voltage of 200 kV. The distance between deformation twins was
measured by the mean linear intercept method, in which the
direction of the measuring line was normal to the twin bound-
aries. The volume fraction of shear bands was calculated by the
systematic point count method using a square grid. A least 10 TEM
micrographs per specimen were used to obtain the relevant
statistics. Average values of the crystallite size d
0
and microstrains
ε
2
50
were estimated on the basis of the Williamson–Hall plot [29]
given by Eq. (2):
β
s
cos Θ
λ
¼
2 ε
2
50
sin
Θ
λ
þ
K
d
0
ð2Þ
where
Θ
is Bragg angle, K is the Scherrer constant and
β
s
is the full
width at the half maximum height (FWHM) of K
α
1
line with the
Table 1
Chemical composition of the investigated steel.
Element С Mn Al Si Cr S P Fe
[Wt%] 0.304 23.1 1.5 0.09 0.08 0.006 0.017 Bal.
P. Kusakin et al. / Materials Science & Engineering A 617 (2014) 52 – 60 53
correction of instrumental line broadening. In this study, the
FWHM values were measured for
γ
-Fe (111) and (222) reflections.
The instrumental line broadening was obtained from the FWHM
value (
β
r
)ofK
α
1
line of an annealed silicon powder and removed
from measured FWHM value (
β
m
)ofK
α
1
line of TWIP-steel on the
following Eq. [30]:
β
2
s
¼β
2
m
β
2
r
ð3Þ
The dislocation density was evaluated by X-ray diffraction profile
analysis using an ARL-Xtra diffractometer operated at 45 kV and
35 mA with Cu K
α
radiation. The value of the dislocation density
ρ
was calculated from the average values of the crystallite size D and
microstrain 〈
ε
2
〉 by the following relationship [31]:
ρ ¼
3
ffiffiffiffiffiffi
2π
p
o
ε
2
50
4
Db
ð4Þ
where b is the Burgers vector (b¼2.5 10
10
m).
The tensile tests were carried out at room temperature using an
Instron 5882 testing machine. Tensile specimens with a gauge
length of 16 mm and cross section of 1.5 mm 3 mm were cut out
with the tensile axis parallel to the rolling axis. The samples were
tested at an initial strain rate of 2 10
3
s
1
. Two tensile speci-
mens were used to obtain each experimental point. The fracture
surfaces after tensile tests were observed using Quanta 200 3D
scanning electron microscope.
3. Results
3.1. Microstructure evolution
Hot rolling and annealing resulted in the development of a
uniform microstructure (Fig. 1) with an average grain size of
24
μ
m containing numerous annealing twins and an average
dislocation density of
ρ
¼10
14
m
2
.
At a reduction of 20%, the initial grains elongated along the
rolling direction and deformation twinning occurred in some
favorably oriented grains (Fig. 2a). Relatively fine initial grains
promoted mechanical twinning [5] and hindered the formation of
a low-energy dislocation structure in the form of planar disloca-
tion boundaries [32]. To elucidate the characteristics of the long-
itudinal features observed in Fig. 2a, EBSD analysis was performed
on the 20% cold-rolled material. On one hand, grains with low
misorientation angles across these longitudinal features indicate
intragranular slip lines (grain 2 in Fig. 3). On the other hand, grains
including a high number of 601 misorientation angles along the
misorientation line profile prove the presence of twin bundles
containing
Σ
3 grain boundaries (601{111}), as exemplarily shown
by the profile of grain 1 in Fig. 3. With respect to the operation of
deformation twinning, three types of grains can be distinguished
(Fig. 2a). Type I grains are almost free of mechanical twins. Type II
grains contain one active primary twinning system. In these
grains, the bundles of straight and thin twins alternate with the
non-twinned matrix, and band-like regions of extensive mechan-
ical twinning aligned with {111} plane were observed (Fig. 4a).
Type III grains were characterized by more than one primary
twinning system or secondary twinning system. Most of the grains
were of types I or II, whereas the fraction of type III grains was
minor. In the type II and III grains, the average twin thickness was
Fig. 1. Initial microstructure of the investigated Fe–23Mn–0.3C–1.5Al TWIP steel.
Fig. 2. Optical micrographs showing the microstructure of specimens after cold rolling: (a) 20%, (b) 40%, (c) 60% and (d) 80%.
P. Kusakin et al. / Materials Science & Engineering A 617 (2014) 52 – 6054
approximately 20 nm, and the average distance between the twins
was 570 nm. The dislocation density at a 20% rolling degree
increased substantially to approximately 2 10
15
m
2
(Fig. 5),
providing high internal stress.
As the rolling degree increases from 20 to 40%, the number of
deformation twins increased considerably (Figs. 2b, 4b), while the
dislocation density remained nearly unchanged (Fig. 5). Almost all
grains experienced deformation twinning. Therefore, the deforma-
tion microstructure consisted of type II and III grains. The bundles
of deformation twins became thicker due to the increased number
of twins within these bundles (Fig. 4b). Deformation twins
belonging to several twinning systems appeared in many grains
(Fig. 2b). The secondary twins crossed over the previous twins. An
example of multiple twining with twins oriented along (1
11) and
(11
1) planes is presented in Fig. 4 b. Up to three activated twinning
systems were observed, with primary twinning systems being
dominant. The angle between the primary and secondary twin
planes, as depicted in Fig. 4b, is approximately 751, which is close
to the angle of 70.51 between the {111} twinning planes in an fcc
lattice. In the type III grains, the multiple deformation twinning
resulted in the subdivision of the non-twinned matrix into
rectangular crystallites due to transverse deformation twins and
dislocation boundaries (Figs. 2b and 4b). In the type II grains, the
mechanical twinning led to the development of nanoscale layered
structures because of the small spacing between twin boundaries.
Furthermore, the first micro shear bands with a typical angle of
401 with the rolling direction were observed (Fig. 2b).
Further rolling to 60% reduction increased the density of twins
and lattice dislocations concurrently (Figs. 2c, 4c, 5). It is worth
noting that there were still many type II grains, in which only
primary twin families could be observed. These twins also have
orientation parallel to a {111} plane. A specific feature of the
microstructural evolution is that the previously formed twin
boundaries tended to rearrange along rolling direction and thus
promoted the appearance of micro shear bands to accommodate
further strain (Fig. 2c). The micro shear bands sheared the
previous deformation twins at acute angles below about 301 as
illustrated by Fig. 4c, where the micro shear band aligned along
(
113) crystal plane. This corresponds to other studies reported
that micro shear bands in FCC crystals form as non-
crystallographic band-like deformation regions of highly concen-
trated plastic flow [33,34]. The dislocation density at 60% rolling
degree increased to
ρ
3.5 10
15
m
2
. Upon further processing to
80% rolling reduction, the number of shear bands and their
thickness increased significantly (Figs. 2d, 4d and 5). The mutual
intersection of shear bands and twins led to the formation of a
spatial net of shear bands consisting of greatly misoriented
crystallites with a size of 40 nm. The evolution of nanocrystal-
line bands was accompanied by a further increase of the disloca-
tion density to approximately 4.5 10
15
m
2
.
Fig. 5 summarizes the effect of cold rolling on the distance
between deformation twins, twin thickness, dislocation density
and volume fraction of shear bands. The dislocation density
rapidly attained a high value of 2 10
15
m
2
just after 20% rolling
reduction. The dislocation density then gradually increased with
straining and approached 4.5 10
15
m
2
at the highest rolling
reduction of 80%. The deformation twins appeared with a thick-
ness of 20 nm, which remained almost constant during cold
rolling. In general, this thickness is a characteristic feature of the
primary twins. The thickness of the secondary twins was found to
be less than 10 nm.
However, an increase in the strain suppressed the secondary
mechanical twinning; as a result, the overall twin thickness was
mostly attributed to the thickness of the primary twins. On the
other hand, the distance between twins decreased rapidly from
570 nm to 180 nm as the rolling reduction increased from 20% to
40%, followed by a continuous decrease to 40 nm as the
subsequent rolling reached 80%. It should be noted that the
decrease in the distance between the twins from 180 nm at a
rolling reduction of 40% to 100 nm and 40 nm after rolling to 60%
and 80%, respectively, roughly corresponds to the overall change in
the sample thickness. Thus, almost no new twins developed
during rolling with reductions above 40%. Therefore, extensive
deformation twinning readily developed during rolling with
reductions up to approximately 40%. After that deformation
twinning as a deformation mechanism seemed to be exhausted.
This saturation of deformation twinning led to the development of
shear banding at large rolling reductions. The volume fraction of
shear bands increased from 7% to 21% with increasing rolling
reduction from 60% to 80%. Hence, the crystallites composing the
shear bands may contribute to grain boundary strengthening in
combination with the lamellas outlined by the twin boundaries
[14].
3.2. Mechanical properties
The true stress–true strain curves and the change in the work-
hardening rate with true strain are presented in Fig. 6. No serrated
flow attributed to DSA was found; thus, this phenomenon is not
important for the mechanical behavior of the Fe–23Mn–0.3C–1.5Al
Fig. 3. EBSD band contrast mapping of the specimen after 20% cold rolling. The misorientation profiles measured along the arrows in grains 1 and 2 indicate the existence
and absence of intragranular grain boundaries with a misorientation of 601, i.e., twin boundaries.
P. Kusakin et al. / Materials Science & Engineering A 617 (2014) 52 – 60 55
steel. This steel exhibited a relatively low YS of 235 MPa in the
initial recrystallized state (Fig. 6a). Extensive strain hardening
initially occurred at a rate of 1750 MPa, after which the rate
increased slightly to 2000 MPa during tension, providing excep-
tionally high necking resistance (Fig. 6b). As a result, a superior
ductility of 96% was attained. It is worth noting that these values of
elongation-to-failure and the strain-hardening rate are typical for
Al-containing TWIP steels [2,24]. No well-defined yielding point
could be found (Fig. 6a). This is in contrast to the data reported by
Kim et al. [24] and may be attributed to the relatively coarse initial
grain size [35].
The cold rolling with reductions of 20, 40, 60, 80% provided
increases of 193, 138, 427 and 495% in the YS and 26, 48, 60 and
66% in the UTS, respectively (Table 2). After a 20% rolling reduc-
tion, the UTS/YS ratio decreased from 2.81 in the initial
state to 1.22, and the ductility decreased by a factor of 3. The
rate of initial strain hardening decreased slightly to 1500 MPa
and tended to decrease gradually with strain to a value of
1150 MPa at the end of the plastic stability stage. An increase
in the rolling reduction led to a significant drop in the UTS/YS ratio
down to 1.12. The
σ
–
ε
curves show a quite small uniform
elongation followed by strain-softening to failure (Fig. 6). An
apparent steady-state flow could be observed in the samples
subjected to cold rolling with reductions below approximately
40% (Fig. 6). In the samples rolled with reductions above 40%, the
onset of localized necking occurred at peak stress, and the flow
Fig. 4. TEM images of specimens after cold rolling: (a) 20%, (b) 40%, (c) 60% and (d) 80%. TW1 and TW2 denote twins belonging to different twinning systems. MSB and SB
denote micro shear and shear bands, respectively.
P. Kusakin et al. / Materials Science & Engineering A 617 (2014) 52 – 6056
