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MATERIALS PERSPECTIVE
State-of-the-knowledge on TWIP steel
B. C. De Cooman*
1
, O. Kwon
2
and K.-G. Chin
3
High Mn twinning induced plasticity (TWIP) steel is a new type of structural steel, characterised by
both high strength and superior formability. TWIP steel offers an extraordinary opportunity to
adjust the mechanical properties of steel by modifying the strain hardening. The use of TWIP steel
may therefore lead to a considerable lightweighting of steel components, a reduction of material
use and an improved press forming behaviour. These key advantages will help implement current
automotive vehicle design trends which emphasise a reduction of greenhouse gas emissions and
lowering of fuel consumption. In addition, high strength TWIP steel will effectively contribute to
weight containment in vehicles equipped with hybrid and electric motors, as these are
considerably heavier than conventional motors. The present review addresses all aspects of
the physical metallurgy of the high strength TWIP steel with a special emphasis on the properties
and key advantages of TWIP sheet steel products relevant to automotive applications.
Keywords: Twinning induced plasticity, Stacking fault energy, Automotive steel
Introduction
Increased overall vehicle quality improvements have
resulted in passenger cars which have steadily gained in
weight. This weight spiral is a direct result of improved
vehicle safety and increased space combined with
enhanced performance, reliability and passenger com-
fort. Carmakers are increasingly building passenger cars
with body-in-white (BIW) designs which emphasise
passenger safety. The safety issue directly related to
BIW materials is passive safety. High impact energy
absorption is required for frontal crash and rear
collision, and anti-intrusion properties are required in
situations when passenger injury must be avoided, i.e.
during a side impact and in case of a rollover, with its
associated roof crush. The weight issue is also high on
the agenda of BIW designers, as it is directly related to
environmental concerns, i.e. CO
2
emission and the
efficient use of non-renewable fuels. Steel based part
design using advanced high strength steel (AHSS) offers
the potential for a combination of increased passenger
safety and vehicle mass containment at lower production
costs. Hence, when material specific properties are
considered, there is an increasing interest in very high
and ultra high strength (ultimate tensile strength,
.1 GPa) materials. Dual phase and transformation
induced plasticity (TRIP) steels are now well established
as AHSS, with major applications in BIW parts related
to crash energy management. In the case of dual phase
grades, the emphasis is on front–end applications and
exterior panels. The use of dual phase and TRIP steels
has been reported to result in a weight saving in the
range of 10–25%. The potential for weight reductions
becomes very important when ultra high strength steels
are considered, and weight reductions in the range of
30–40% are possible for 1300–1500 MPa steels. Strength
levels in the range of 1800–2000 MPa have been
mentioned as future requirements for anti-intrusion
parts. An ultra high strength, a high stiffness, and only
very low levels of collision related deformations,
typically less than 5%, are allowed for these parts. Hot
press forming 22MnB5 grades are increasingly being
used for B pillar and front–rear reinforcements.
Strength is often achieved at the expense of form-
ability, and a global steel research effort is currently
underway focusing on TWIP steel, which offers an
extraordinary opportunity to adjust the mechanical
properties by modifying the strain hardening. Figure 1
compares the mechanical properties of 1 GPa TWIP
steel with the properties of common ferritic and
austenitic sheet steel grades routinely used to manufac-
ture press formed parts. Figure 2 compares the strain
hardening behaviour of 1 GPa TWIP, high strength
interstitial free (IF) and TRIP steels. TWIP steel has
twice the uniform elongation than that of TRIP steel
and considerably higher ultimate strength. This unu-
sually high strain hardening leads to superior press
forming performance, which is very convincingly illu-
strated in Fig. 3. The strain hardening mechanism in
TWIP steel, illustrated in the schematic of Fig. 4, is best
described as a ‘dynamical Hall–Petch effect’. The
dominant deformation mode in TWIP steel is still
dislocation glide but, in the dynamical Hall–Petch effect,
mechanical twins are also continuously being formed
during straining. As the formation of mechanical twins
involves the creation of new crystal orientations, the
twins progressively reduce the effective ‘mean free path’
or effective glide distance of dislocations and increase
the flow stress. The gradually increasing twin density
results in the very high strain hardening observed in
1
Graduate Institute of Ferrous Technology, Pohang University of Science
and Technology, Pohang, Korea
2
POSCO Center, Seoul, Korea
3
POSCO Technical Research Laboratories, Gwangyang, Korea
*Corresponding author, email decooman@postech.ac.kr
ß2012 Institute of Materials, Minerals and Mining
Published by Maney on behalf of the Institute
Received 5 August 2011; accepted 22 September 2011
DOI 10.1179/1743284711Y.0000000095 Materials Science and Technology 2012 VOL 28 NO 5513
TWIP steel. The key to twin formation lies in the precise
control of the stacking fault energy (SFE): if the SFE is
,20 mJ m
22
, strain induced transformation is more
likely to take place; if the SFE is .50 mJ m
22
, the
formation of twins is suppressed.
The focus of the present review is on high Mn
austenitic TWIP steel with a Mn content in the range of
12–30 mass%.
1–4
These relatively high Mn ferrous alloys
are characterised by strength–ductility products in the
range of 40?000–60?000 MPa %. The main alloying
additions to TWIP steel are C, Mn, Al and, in some
cases, Si. The C, Mn, Al and Si contents are precisely
controlled to obtain a SFE of 20–50 mJ m
22
. The C
additions, typically in the range 0?4–1?0 mass-%, also
result in the stabilisation of the austenite phase and solid
solution strengthening.
Note that TWIP steel has a different alloy design
concept from the one used for high Mn austenitic
specialty steels which are currently already being
produced industrially, such as wear resistant steel (e.g.
X120Mn12 with 12 mass-%Mn), cryogenic steel (e.g.
X40MnCr22 with 22 mass-%Mn), Ni free stainless steel
(e.g. AISI 205 with 14–15?5 mass-%Mn), non-magnetic
steel (23–25 mass-%Mn), high damping steel (17 mass-%
Mn) and shape memory steel (Fe–Mn–Si alloys with
28 mass-%Mn).
TheearlyscientificworkonhighMnferrous
alloys
1–3,5–8
did not receive much industrial attention
originally. In addition, there were considerable difficul-
ties during the first attempts to process TWIP steel
industrially a few years later. The perception that TWIP
steel was prone to delayed fracture and dynamic strain
aging (DSA), made the initial industry efforts to develop
TWIP steel even more challenging. Two factors have
recently resulted in a renewed industry wide interest in
high Mn TWIP steel:
(i) the discovery of the beneficial effects of Al
additions on the suppression of the delayed
fracture phenomenon
(ii) the strong demand for formable ultra high
strength steels from the automotive industry.
By 2008, three types of TWIP steel composition had
been extensively investigated: Fe–25–30Mn–3Si–3Al,
1
Fe–22Mn–0?6C,
9
and Fe–18Mn–1?5Al–0?6C.
10
At present, cost considerations and the marked
improvement of the mechanical properties of TWIP
steel with a reduced Mn content are the main drivers in
TWIP steel research. New, leaner, compositional designs
for TWIP steel have been proposed and the current
worldwide efforts in the development of TWIP steel
suggest that industrial production of TWIP sheet steel
will very likely focus on TWIP steel with y15 mass-
%Mn. Examples of industrial TWIP steel with 15 mass-
%Mn which had been developed by early 2009, include
the following alloys: Fe–15Mn–2?0Al–0?6C–0?5Si,
11
Fe–
15Mn–3?0Al–0?7C,
12,13
Fe–15Mn–2?5Al–0?7C,
14
Fe–
15Mn–2?5Si–2?5Al–0?6C, Fe–15Mn–2?5Al–0?6C,
15
and
Fe–16Mn–0?6C–0?2Si–0?2Al.
16
Research efforts are
currently underway to evaluate the performance of
various 12 mass-%Mn TWIP steel variants. Examples
include the following alloys: Fe–12Mn–0?8C,
17
and Fe–
12Mn–2Si–0?9C.
18
The potential applications for Mn TWIP steel will
very likely not be limited to the automotive industry,
i.e. for applications in passenger cars, light trucks and
2 Comparison of strain hardening behaviour of high
strength IF, TRIP and TWIP steels
3 Comparison of Nakashima test samples for IF and
TWIP steels
4 Schematic of strain hardening mechanism in TWIP
steel
1 Comparison of mechanical properties of 1 GPA TWIP
steel with properties of common ferritic and austenitic
sheet steel grades
De Cooman et al. State-of-the-knowledge on TWIP steel
514 Materials Science and Technology 2012 VOL 28 NO 5
commercial vehicles such as vans and heavy trucks.
Additional applications may exist in other transporta-
tion areas (street cars, buses, trains and railcars), line
pipe production, ship building and in special applica-
tions which require non magnetic panels.
Thermodynamic properties of TWIP steel
The Fe rich side of the Fe–Mn equilibrium phase
diagram is relatively simple, with an open cloop,
19
but
the microstructures of actual Fe–Mn alloys shown in
Fig. 5 reveal that in the range of 5–25 mass-%Mn, the
microstructure of Fe–Mn alloys is dominated by the
presence of a9martensite, at a low Mn content, and e
martensite, at a higher Mn content. Approximately
27 mass-%Mn is required to obtain metastable austenite
at room temperature. Relatively small C additions of
0?6 mass-% result in martensite free, austenitic micro-
structures for Mn contents as low as 12 mass-% (Fig. 5).
The microstructure of TWIP steel is single phase
austenitic, with relatively coarse grains, which often
contain wide recrystallisation twins. Al is added to
control the SFE of TWIP steel. Al also suppresses the
formation of the carbide Fe
3
C. The Al additions,
typically less than 3 mass-%, also result in a slight
reduction in density. This is due to the combined effect
of the lower molecular weight of Al and the increase of
the lattice parameter by Al additions. An alternative
approach to obtain TWIP steel with a uniform, carbide
free, austenitic microstructure is to use a higher Mn
content and avoid C additions. This TWIP steel
composition concept typically requires Si and Al
additions to control the SFE. The importance of the
Al additions to TWIP steel cannot be underestimated as
it results in much improved TWIP properties. Jung
et al.
20
have shown that small additions of Al facilitated
the TWIP effect and the formation of emartensite was
effectively suppressed by the addition of 1?5 mass-%Al
to Fe–15Mn–0?6C TWIP steel. Nitrogen has been found
to have similar effects as Al.
22
Stacking fault energy
The SFE is a key factor controlling the mechanical
properties of the high Mn alloys. The SFE plays an
essential role in the occurrence of the TWIP effect, and
the SFE which is most often reported to be required for
an TWIP effect is 20–50 mJ m
22
. It is still unclear why
this minimum ‘magic number’ is essential for the
occurrence of the strain induced twinning, but it appears
to be related to the suppression of the athermal cRe
martensitic transformation (Fig. 6). As the SFE is an
essential parameter, there has been a considerable
interest in determining its value for TWIP steels. SFE
values for TWIP steels, taken from the literature, are
listed in Table 1.
22–27
From a theoretical point of view, the SFE is
proportional to the fcc and hcp free energies difference
DG
cRe
. The contributions from the interfacial energy
DGc?e
surface and the magnetic energy DGc?e
magnetic, to the SFE
need to be taken into account as they may have a
significant influence
5 Pseudobinary phase diagrams for various Fe–Mn alloys (top). In contrast to what is predicted by the equilibrium phase
diagram, the room temperature microstructure of binary Fe–Mn alloys is dominated by the formation of a9and emar-
tensite. These transformations are suppressed by alloying additions of C (below)
De Cooman et al. State-of-the-knowledge on TWIP steel
Materials Science and Technology 2012 VOL 28 NO 5515
c~1
8V2=3(DGc?e
bulkzDGc?e
surface zDGc?e
magnetic)
The interfacial energy can be taken as the coherent twin
boundary energy and the energy of the twinning
dislocations. Most authors report that stable, fully
austenitic microstructures with TWIP properties have
a SFE in the range 20–30 mJ m
22
. C additions are
required to obtain a low SFE. Some data on the effect of
the C content in Fe–22Mn–C alloys have been reported
by Yakubtsov et al.
28
They report that the SFE of a Fe–
22Mn alloy is y30 mJ m
22
. C additions less than
1 mass-% reduce the SFE to y22 mJ m
22
. At higher C
contents, the SFE is reported to increase. Frommeyer
et al.
3
indicated that whereas a SFE larger than
y25 mJ m
22
would result in the twinning effect in a
stable cphase, a SFE smaller than y16 mJ m
22
resulted
in ephase formation. Alain et al.
29
gave a much
narrower range. According to them, the SFE should
be at least 19 mJ m
22
to obtain mechanical twinning.
They mentioned that a SFE less than 10 mJ m
22
resulted in ephase formation. Dumay et al.
30
mentioned
that below a SFE of 18 mJ m
22
, twinning tended to
disappear and was replaced by eplatelets. They
mentioned that a SFE of y20 mJ m
22
was needed for
the best hardening rate.
Al increases the SFE and very effectively suppresses
the cRetransformation. Jin et al.
31
mentioned that a
SFE value of 33 mJ m
22
is required to obtain twinning
in Fe–18Mn–0?6C–1?5Al. Recently, Kim et al.
27
reported
that the SFE of Fe–18Mn–0?6C–1?5Al TWIP steel
was 30 mJ m
22
. The SFE for Fe–Mn–Si–Al TWIP
steel has been also studied by Huang et al.
32
They have
also studied the effect of 0?011–0?052% nitrogen on the
SFE of Fe–(20?24–22?57)Mn–(2–3)Si–(0?69–2?46)Al
containing 100 ppm C, by means of XRD. Their results
indicate that Al and N increase the SFE and decrease the
stacking fault formation probability. Dumay et al.
30
have shown that Al increases the SFE by about
z5mJm
22
per added mass-%Al. Si also increases the
SFE by about z1mJm
22
per mass-%Si. These results
are contradicted by Tian et al.
33
who measured the SFE
Table 1 Review of reported SFE values for TWIP steels
Composition/wt-%
Microstructure
SFE/mJ m
22
Method ReferenceDeformation mechanism
Fe–19Mn–5Cr–0.25C–1Al cze20.90 TEM 21
Mechanical twinning
a9martensite
Fe–19Mn–5Cr–0.25C–2.5Al c30.50
Mechanical twinning
Dislocation glide
Fe–19Mn–5Cr–0.25C–3.5Al 39.40
Fe–19Mn–5Cr–0.25C–4Al cza47.50
Mechanical twinning
Dislocation glide
Fe–22Mn–0.6C (T5673 K) c80.00 Thermodynamic calculations 22
Dislocation glide
Fe–22Mn–0.6C (T5293 K) c19.00
Mechanical twinning
Dislocation glide
Fe–22Mn–0.6C (T577 K) c10.00
emartensite formation
Dislocation glide
Fe–25Mn–0.15C–0.6Al cze7.75 XRD 23
Mechanical twinning
Dislocation glide
Fe–25Mn–0.15C–1.5Al c10.67
Fe–25Mn–0.15C–2.2Al 15.12
Fe–25Mn–0.15C–3.1Al Mechanical twinning 14.95
Fe–25Mn–0.15C–4.8Al 54.74
Fe–31Mn–0.17C Dislocation glide 17.53 XRD 24
Fe–18Mn–0.6C 13.00 TEM 25
Fe–18Mn–0.6C–1.5Al 26.40 TEM 26
30.00 TEM
6 Schematic showing relationships between composition,
SFE and phases present at room temperature after
quenching from annealing temperature of 700
u
C
De Cooman et al. State-of-the-knowledge on TWIP steel
516 Materials Science and Technology 2012 VOL 28 NO 5
for Al added Fe–25Mn–0?7C steel. They report a much
smaller SFE increase of about z1?4mJm
22
per added
mass-%Al.
Wang et al.
34
have remarked that although there is a
general consensus that the SFE is an essential para-
meter, it is by no means proven that it is the single most
important parameter controlling the TWIP mechanism.
They noticed that only a small SFE difference of 5–
10 mJ m
22
caused an apparently sharp transition from
strain induced emartensite formation to strain induced
twinning.
The question is further complicated because there is
no clear relationship between the SFE and the
twinning mechanism operating in TWIP steel. Recent
experimental measurements of the nature of the
stacking faults have resulted in the suggestions that e
martensite formation and mechanical twinning is me-
diated by the extrinsic stacking fault (ESF) and the
intrinsic stacking fault (ISF) respectively. Idrissi et al.
35
studied the deformation mechanism of a two phase
azcFe–19?7Mn–3?1Al–2?9Si steel. Deformation at
86 and 160uCresultedinemartensite and twinning at
lower temperatures, and exclusively mechanical twin-
ning at the higher temperatures. At room tempera-
ture, only emartensite was observed. They argued
that this was due to the presence of ESFs at lower
temperatures acting as precursors to emartensite
formation and ISFs at higher temperatures acting as
twin precursors.
Based on this, thermodynamic models are available
which predict the SFE correctly for the quaternary Fe–
Mn–Al–C alloy system.
36
A comparison of calculated
and experimentally determined SFEs for Fe–Mn–Al–C
TWIP steel is shown in Fig. 7.
Mechanical properties of TWIP steel
Fe–Mn alloys are characterised by the occurrence of
complex deformation phenomena: strain induced a9
martensite formation, emartensite formation, deforma-
tion twinning, pseudotwinning, extended dislocation
glide and perfect dislocation glide. The dominant
deformation mode is dislocation glide but the process
is strongly influenced by these complex phenomena.
The best way to gain insight into the deformation
behaviour of TWIP steels is by TEM studies, as the
twinning microstructure is very fine. This is illustrated
in Fig. 8. At low strains, the dislocation density in-
creases and the grain boundaries seem to be particularly
effective source of isolated stacking faults. The onset of
twinning requires multiple slip within deformed grains.
The strain induced twins have a high aspect ratio and
cross the entire grain. At 20% strain, the high dis-
location density between the deformation twins clearly
shows that twin boundaries act as effective barriers for
dislocation movement. High resolution TEM has
clearly shown that the twins are very thin and there
may be continuous nucleation of new deformation
twins of decreasingly smaller size. Consequently, the
twin volume fraction does not represent a large portion
of the total volume.
Although the role of deformation induced twins will
mainly be discussed in the following paragraphs, it
must not be forgotten that the rate of dislocation
accumulation will automatically increase when an alloy
has a low SFE, independently of twin formation, as the
larger dissociation width will more effectively reduce
the cross-slip and result in a higher rate of dislocation
accumulation.
Whereas the TWIP steel has a high strain hardening,
it also has a relatively low yield stress. Grain size
reduction and solid solution hardening can be used to
obtain higher yield strengths for applications where this
is required. The grain size dependence of the yield
strength s(MPa) has been reported to follow the Hall–
Petch relationship
37
s~206:9z704:3
d1=2
where dis the grain diameter in mm. Solid solution
hardening data has recently been made available for Fe–
Mn–C TWIP steel
38
s~228z187|%C{2|%Mn
The data shown in Fig. 9
39
are for the solid solution
hardening of Fe–Mn–Al–C TWIP steel. The yield
strength is increased by Al solid solution hardening.
The decrease of the tensile strength in Al added TWIP
steel is due to the reduction of the strain hardening,
resulting from a decrease in the rate of strain induced
twin formation.
The following Swift and Voce equations, which have
been derived by Barlat and co-workers,
40
clearly point to
the pronounced work hardening of TWIP steel
s~2260|0:072zeðÞ
0:626
s~2380{1940e{1:71e
Various models have been proposed to explain the
pronounced work hardening of TWIP steel. Note that
none of them uses the SFE as an input parameter,
illustrating the fact that the relationship between the
TWIP effect and the SFE is still not clearly understood.
Bouaziz et al.
41
and Allain et al.
42,43
proposed a physical
model for the effect of strain induced twinning on the
mechanical properties of TWIP steel using the Kocks–
Mecking
44
approach. In their model, the tensile stress–
strain curve is calculated on the basis of the evolution of
the dislocation density dr/de, and the twin volume
fraction
7 Comparison of calculated and experimentally deter-
mined SFE for Fe–xMn–0?6C and Fe–xMn–1?5Al–0?6C
TWIP steels
De Cooman et al. State-of-the-knowledge on TWIP steel
Materials Science and Technology 2012 VOL 28 NO 5517
dr
de
~M1
Lb zk
br1=2{fr
1
L~1
dz1
t
s~aMGbb1=2
In these equations, ais a numerical parameter, Mis the
Taylor factor, Gis the shear modulus, bis the Burgers
vector, ris the dislocation density, kis the forest
hardening coefficient, fis the coefficient for dynamic
recovery, dis the grain size and tis the twin spacing. The
model leads to an isotropic hardening behaviour. Their
results support the fact that the total volume fraction of
the twins is very low and that plastic deformation is
mainly achieved by dislocation glide. In a first stage,
nanometre thick twins move until they reach a strong
boundary. In a second stage, the twins thicken. Two
twinning systems are sequentially activated, whereby
twins of the first system develop across the entire grain,
and twins of the second system develop between the
primary twins and are much shorter and thinner. The
thickening of the twins is still controversial. In their
study of the strain hardening of Fe–18Mn–0?6C–1?5Al
TWIP steel, Jin et al.
31
reported that the amount of
twinned volume was controlled by the increase in the
number of new deformation twins and not by their
lateral growth.
A micromechanical model incorporating elasto-visco-
plasticity was developed by Shiekhelsouk et al.
45
for Fe–
22Mn–0?6C TWIP steel. They reported that the twinned
volume fraction was dependent on the grain orientation,
and less than 0?08 for a macroscopic strain of 0?4. This is
in disagreement with the large twin volume fraction of
0?56 at a strain of 0?4 reported by Dini et al.,
37
who
analysed the dislocation density evolution in Fe–31Mn–
3Al–3Si TWIP steel during straining by means of XRD.
In the model of Kim et al.,
46
based on the Kubin–
Estrin approach,
47
the strain hardening is computed
from the evolution of the coupled densities of the mobile
dislocations and immobile forest dislocations. A typical
result for Fe–18Mn–0?6C–1?5Al TWIP steel is shown in
Fig. 10.
Although Xu and Barlat
48
have recently shown that
the work hardening of 980 MPa Fe–Mn–Al–C TWIP
steel is almost isotropic up to a longitudinal plastic
strain of 0?24, the strain hardening of TWIP steel is very
likely not isotropic over the full deformation path, as
there are clear indications of kinematic hardening at
larger strains. Sevillano
49
has argued that the observa-
tion of the Bauschinger effect in TWIP steel is due to the
simultaneous deformation of the grains and their
twinned part. This requires a forward internal stress
operating on the twin and a backward internal stress
operating on the untwinned matrix, as both matrix and
twins must share similar strain components. Bouaziz
8 Images (TEM) of TWIP steel after 5, 10 and 20% of strain. In the initial stages of deformation (left), the dislocation den-
sity increases and there is no formation of twins. In addition, some grain boundaries emit bundles of stacking faults.
At a higher strain (middle), narrow strain induced twins cross the entire grain. At a very high strain (right), TWIP steel
develop a dislocation cell structure and arrays of stacking faults in the regions between twins
De Cooman et al. State-of-the-knowledge on TWIP steel
518 Materials Science and Technology 2012 VOL 28 NO 5
et al.
50
have linked the back-stress to dislocations of a
given slip system being stopped at the twin boundaries
and developing a stress which prevents similar disloca-
tions from moving.
Figure 11 illustrates the wide variety of properties one
can obtain from TWIP steel starting from the Fe–
18Mn–0?6C composition. A reduction of Mn from 18 to
12% results in fracture at lower strengths. A slight
increase in the C content to 0?7 mass-% and the addition
of 1 mass-%Al gives rise to a substantial recovery of the
original ultimate strength level. A further increase in
elongation is achieved by Cu addition. These experi-
mental observations, very likely the result of changes
in both SFE and twinning kinetics, highlight the com-
plexity of the plasticity for TWIP steel and the need
for further fundamental work on their deformation
mechanisms.
Strain induced twinning
In a metastable austenitic steel, the deformation selec-
tion rules can in principle be controlled by the SFE, the
martensite transformations in the alloy system (M
s
,Ms
s
and M
d
), the critical resolved shear stress for each
deformation mode, the stress level and the strain path.
The SFE is an important parameter, but it cannot by
itself predict the plasticity mechanisms. Even in TWIP
steel, mechanical twinning only takes place if strict
conditions related to the grain orientation, the stress
direction (i.e. tension or compression), the amount of
deformation, the SFE and the temperature are met. The
transition from slip only deformation to slip and
twinning deformation occurs when the slip stress reaches
the twinning stress. As there is no agreed model for twin
formation, the stress required to nucleate a twin is
difficult to compute without making some essential
simplifications. The growth of a twin usually requires a
much lower stress than what is usually computed by
models, and it is therefore assumed that the high
nucleating stress is due to a local stress concentration,
as the applied stress results in homogeneous stresses too
low to nucleate twins. The twinning stress increases with
increasing SFE, and the stress required to nucleate a
twin is related to the intrinsic SFE in a quadratic or
linear manner. Various critical twinning stress formulas
have been proposed and they are listed in Table 2,
together with critical shear stress values for an SFE of
20 mJ m
22
.
51–55
9 Solid solution hardening of TWIP steel by C, Mn and Al. The solution hardening coefficient for C in the table is based
onthedatafromFe–15Mn–2Al–xC alloys. The coefficient for Mn is based on the data from Fe–xMn–2Al–0?7C alloys.
The coefficient for Al is based on the data from Fe–15Mn–xAl–0?9C alloys
10 Comparison of experimental and calculated stress–
strain curves for twin free grains, twinning grains and
‘phase mixture’ model for Fe–18Mn–1?5Al–0?6C TWIP
steel
De Cooman et al. State-of-the-knowledge on TWIP steel
Materials Science and Technology 2012 VOL 28 NO 5519
The grain size Dplays a role in the value of the
twinning stress and larger grains tend to expand the
twinning domain
53
sT~sT0zkT=D1=2
The k
T
value is usually much larger than the k
y
value for
dislocation slip in the standard Hall–Petch relation.
In TWIP steel, the nucleation of twins is not a
homogeneous process: in the nucleation stage the
formation of twins is closely related to prior dislocation
activity, and strain induced twin formation occurs only
after some amount of prior dislocation generation and
dislocation–dislocation interactions. Twins are initiated
in special dislocation configurations created by these
interactions generally resulting in multilayer stacking
faults which can act as twin nuclei. Three types of strain
induced twinning nucleation mechanisms have been
proposed: pole mechanisms, glide dislocation sources
and cross-slip mechanisms. A number of these mechan-
isms are illustrated in Fig. 12. In the pole mechanism
proposed by Venables,
56–58
a jog is created on a
dislocation by dislocation intersection (Fig. 12a). This
jog dissociates in a sessile Frank partial dislocation and
a Shockley partial dislocation. When the partial
dislocation moves under the influence of an externally
applied force, it trails an ISF and it rotates repeatedly
around the pole dislocations, generating a twin in the
process.
Glide type deformation induced twinning mechanisms
have also been proposed. Glide sources are less probable
sources of twins, but Bracke et al.,
59
who studied
twinning in Fe–22Mn–0?5C TWIP steel by means of
TEM, found support for the Mahajan and Chin
60
model
for the creation of a three layer stacking fault acting as
twin embryos (Fig. 12b). These twin embryos are
formed by the interaction between two coplanar glide
dislocations. The interaction leads to the formation of
an ESF configuration. Larger twins form by the growth
of twin embryos into each other.
In the cross-slip models for twinning,
61–65
a twinning
partial is formed by a double cross-slip mechanism
(Fig. 12c) or by the interaction of a primary dislocation
with a Frank dipole, a faulted dipole or a Lomer–
Cottrell lock (Fig. 12d).
Although mechanical twins could in principle be
nucleated at dislocations, grain boundaries or other pre-
existing defects, it appears that some predeformation,
considerable dislocation interactions and the nucleation
of ESFs are always required to nucleate twins.
Grain orientation and crystallographic texture in
general, can also have a pronounced effect on twinning.
For example, whereas n100moriented grains deformed
in tension do not twin, grains oriented with a n111mtype
axis parallel to the tensile axis are heavily twinned.
53,66
Press forming properties
Compared to a uniaxial tensile test, deformation in press
forming can be very complicated, involving complex and
11 Stress–strain curves illustrating the effect of the reduction of Mn content of TWIP steel and the improvement of
mechanical properties by minor additions of C, Al and Cu
Table 2 Critical stress for twinning (SFE520 mJ m
22
,
G560 GPa, b50?257 nm, b
p
50?148 nm, d520 mm,
mean twin width L
0
550 nm, K56 GPa, M53?06)
Critical stresses for twinning
t
T
References
T
5Mt
T
tT~c=bt
T
571 MPa 50
s
T
5238 MPa
tT~2c=bpt
T
5270 MPa 51
s
T
5827 MPa
tT~c=bzGb=dt
T
579 MPa 52
s
T
5240 MPa
sT~6:14c=bpt
T
5271 MPa 53
s
T
5830 MPa
sT~Kc=GbðÞ
1=2t
T
571 MPa 54
s
T
5216 MPa
De Cooman et al. State-of-the-knowledge on TWIP steel
520 Materials Science and Technology 2012 VOL 28 NO 5
inhomogeneous conditions of strain and stress. The yield
surface is often needed for numerical simulations and
the shape of the yield surface is known to have a
significant effect on the forming limit of anisotropic
sheet metals. Figure 13 shows phenomenological yield
functions for 1 GPa TWIP steel. In addition, typical
values for the angular dependence of the normal
anisotropy are listed.
40
The planar anisotropy of
TWIP steel is limited.
In sheet forming, the forming limit diagram (FLD) is
obtained by determining experimentally the strains at
which necking occurs for different deformation paths.
Figure 14 compares the FLD of TWIP, IF, low C Extra
Deep Drawing Quality (EDDQ) and TRIP steels
currently used for automotive BIW applications. The
necking strain in the plane strain deformation mode, the
Forming Limit Curve Minimum (FLC
0
) value of TWIP
steel is very high (45%) compared to those for 590 MPa
dual phase steel (30%) and 780 MPa TRIP steel (28%).
The press forming properties of TWIP steel have
therefore proven to be excellent as illustrated by the
example of the shock absorber housing in Fig. 15. This
makes TWIP steel ideally suited for the press forming of
high strength parts with a complex shape.
The stretch flanging properties of TWIP steel,
typically evaluated by means of a hole expansion test,
are considerably better than those of the other AHSS of
a similar strength level, but not as good as those of fer-
ritic deep drawing steel grades. The combination of low
normal anisotropy and low strain rate sensitivity results
in lower hole expansion ratios. This is very pronounced
when a less appropriate hole edge preparation procedure
leading to considerable hole edge deformation, such as
in the case of hole punching, is used.
67
This is illustrated
in Fig. 16.
12 Schematic showing different stages in process of mechanical twin formation for aVenables pole mechanism, b
Pirouz cross-slip mechanism and glide source mechanisms according to cVergnol and dMahajan
13 Yield surfaces for 980 MPa Fe–Mn–Al–C TWIP steel
De Cooman et al. State-of-the-knowledge on TWIP steel
Materials Science and Technology 2012 VOL 28 NO 5521
High strain rate properties
The dynamic energy absorption of different types of
automotive steels when tested at a strain rate of 10
2
s
21
are compared in Fig. 17. Frommeyer et al.
3
have
reported high strain rate properties for Fe–25Mn–3Si–
3Al–0?03C TWIP steel. Extensive twinning was reported
to occur during high strain rate deformation, and no
brittle fracture was observed even at low temperatures.
Ueji et al.
68
studied the high strain rate deformation of
Fe–31Mn–3Si–3Al TWIP steel for a grain size in the
range 1?1–35?5mm. In contrast to ferritic steels, there
was still a large elongation at small grain sizes. Sahu
et al.
69
have studied the high strain rate behaviour of two
Fe–24Mn–0?5Si–(0?11–0?14)C TWIP steels with 0?91 and
3?5%Al additions. The transformation of austenite to
martensite was reported to take place up to a strain rate
of 10
3
s
21
. TWIP steel alloyed with 3?5%Al had higher
stability, and the transformation of this TWIP steel was
limited to the strain rate range 10
23
–720 s
21
. Irrespective
of the Al content, the transformation of the austenite
phase is suppressed during high strain rate deformations
due to the adiabatic heating of the sample. Based on the
observation of serrated grain boundaries, they also
argued that dynamic recrystallisation may be taking
place during the high strain rate tests.
Strain localisation and point defects in
TWIP steel
Dynamic strain aging occurs at room temperature in
Fe–22Mn–0?6C and Fe–18Mn–0?6C TWIP steel. This
15 Example illustrating use of TWIP steel for press form-
ing of a shock absorber housing
16 Hole Expansion Ratio (HER) for TWIP steel compared to
the HER–ultimate tensile strength relationship observed
for a large number of automotive materials indicated by
grey band (top). Illustration of the difference in TWIP
steel hole expansion performance for a low quality
punched hole (below, left) and a high quality drilled hole
(below, right)
17 Comparison of energy absorption in J mm
23
,during
high strain deformation (strain rate: 10
3
s
21
) for TWIP
steel and common types of automotive steels
14 Comparison of FLD curve of IF, EDDQ, TRIP and
TWIP steel
De Cooman et al. State-of-the-knowledge on TWIP steel
522 Materials Science and Technology 2012 VOL 28 NO 5
phenomenon manifests itself as serrations in the stress–
strain curve. During DSA, the mechanical deformation
of TWIP steel is entirely localised in deformation bands
which cross the tensile testing sample. The properties of
these bands are illustrated in Fig. 18, which shows the
hopping of the strain state at two points in the tensile
sample during the motion of these Portevin–Lechatelier
(PLC) bands. The properties of the PLC bands have
been analysed in detail.
70–72
The band velocity decreases
with strain and the band strain rate is 15–100 times the
applied value. Localisation may in principle result in
press forming difficulties, but the occurrence of PLC
bands in uniaxial tensile testing has not been reported to
lead to the poor press forming performance for TWIP
steel in practice. This is very likely due to the relatively
high strain rates used in press forming. Other aspects of
DSA should, however, not be overlooked. Dynamic
strain aging results in negative strain rate sensitivity and
it is the cause of a limited post-uniform elongation.
In C alloyed fcc alloys, the room temperature DSA
cannot be explained by long range diffusion of C. Instead,
it results from the short range order due to the presence of
point defect complexes which can re-orient themselves
in the stress field of dislocations or in the stacking
faults. Possible defect complexes in TWIP steels are the
following: C–vacancy complex, C–C complex and C–Mn
complex. The most likely complex has one C atom in an
octahedral interstice and one Mn atom. Direct evidence
for the presence of these point defect complexes, and their
interaction with dislocations, comes from internal friction
measurements (Fig. 19). The DSA mechanism is similar to
the one proposed by Curtin et al.
73
The re-orientation of
the point defect complex does not require long range
diffusion, and only a single diffusional hop of the
interstitial C in the complex is needed to achieve a suitable
orientation with respect to the strain field of the partial
dislocation or the stacking fault.
Dynamic strain aging and the associated serrated
stress–strain curves can be avoided by addition of Al as
illustrated in Fig. 20. As Al increases the stacking fault
energy, but does not interact with the point defect
complex and decreases the C diffusivity in TWIP steels,
the data clearly suggest that the interaction giving rise to
the flow localisation is the interaction between the C–
Mn point defect complexes and the stacking faults. The
study of fundamental aspects of DSA is ideally suited
for analysis by ab initio modelling. Mn rich octahedra
are preferred locations for C atoms. The probability of a
pure Mn octahedron is low. The Mn–C bond is stronger
than the Fe–C bond. Preliminary results show that the
Fe–Fe and Mn–Mn bonds are influenced differently by
the presence of C. In the presence of C, the Fe–Fe bond
strength is slightly destabilised and the Fe–Fe bond
length becomes slightly longer. In contrast, the Mn–Mn
bond length is reduced. This leads to the stabilisation of
C containing Mn rich octahedra, and hence short range
ordering.
74,75
18 Dynamic strain aging in TWIP steel: type A serrations
on the stress–strain curve of C alloyed TWIP steel are
associated with the propagation of isolated PLC bands.
Local strain analysis (left) shows that the sample defor-
mation is entirely by the propagation of PLC bands
(right), which are visible by infrared thermography
19 Internal friction measurement showing the presence of a damping peak associated with point defect complexes invol-
ving C in TWIP steel (left). The increase in the internal friction peak amplitude after straining is indicative of the inter-
action between the dislocations and the point defect complexes (middle). Annealing restores the peak amplitude
measured before deformation (right)
De Cooman et al. State-of-the-knowledge on TWIP steel
Materials Science and Technology 2012 VOL 28 NO 5523
Delayed fracture
The delayed fracture phenomenon is illustrated in
Fig. 21. It was originally identified as a major problem
in Fe–22Mn–0?6C TWIP steel, where the effect appears
in fully deep drawn cups a relatively short time after cup
drawing. Note that high residual tensile hoop stresses
are known to be present in the edge of fully drawn cups.
Kim et al.
76
have suggested that it is related to
martensitic transformation in the presence of residual
stresses and hydrogen. They investigated the influence of
martensite formed during the deformation of Fe–18Mn–
0?6C and Fe–18Mn–0?6C–1?5Al TWIP steel. Al added
TWIP steel was free of martensite after tensile testing,
but both types of TWIP steel contained martensite after
cup drawing. The amount of the embrittling martensite
phase was lower in the Al added TWIP steel.
Jung et al.
77
compared the hydrogen embrittlement of
TRIP and TWIP steels after cathodic hydrogen char-
ging. They report that Fe–15Mn–0?45C–1Al and Fe–
18Mn–0?6C TWIP steels, with and without Al additions,
contained less hydrogen and were much more resistant
to embrittlement than TRIP steel after U bend and cup
drawing tests.
The delayed fracture issue has prompted interest in
testing the sensitivity of TWIP steel to hydrogen induced
embrittlement.
78
The absence of any noticeable degra-
dation of the mechanical properties after H charging,
except for a relatively small reduction of the total
elongation, suggests that TWIP steel may actually be
insensitive to hydrogen induced fracture (Fig. 22).
Delayed fracture in TWIP steel is therefore very likely
a form of stress corrosion cracking as three conditions
must apparently be satisfied simultaneously: high
residual stresses, presence of a strong hydrogen trap
site and hydrogen absoption. If any of these conditions
is not satisfied, delayed fracture is unlikely to occur.
Fatigue properties
In the as received state, TWIP steels experience cyclic
softening during fatigue testing. During cyclic tests the
dislocation density decreases and the existing twins
become wider. The resulting lack of dislocation–twin
interactions and the absence of nucleation of new twins
are the major causes of the observed softening.
Predeformation has a positive effect on the fatigue life
of Fe–22Mn–0?52C TWIP steel,
79
as the presence of
20 Stress–strain curves of TWIP steel illustrating suppression of serrations due to DSA by Al additions
21 Example of delayed fracture deep drawn Fe–22Mn–0?6C TWIP steel (left) and suppression of delayed fracture in deep
drawn Fe–15Mn–0?6C TWIP steel by alloying additions of 2?5%Al (right)
De Cooman et al. State-of-the-knowledge on TWIP steel
524 Materials Science and Technology 2012 VOL 28 NO 5
twins formed during predeformation leads to a stable
response in cyclic loading and a longer fatigue life. High
cycle bending fatigue results for Fe–22?3Mn–0?6C, Fe–
17?8Mn–0?6C and Fe–16?4Mn–0?29C–1?54Al have been
reported by Hamada et al.
80
The steels had a
2610
6
cycles fatigue stress limit of 400 MPa which
was well above the yield stress. The ratio of fatigue limit
to tensile strength was comparable to the ratio for
austenitic steels, 0?42–0?48.
Welding
Resistance spot welding (RSW) is the most important
joining method for automotive sheet steel. The RSW of
TWIP steel is more challenging than the RSW of plain C
steels, and special precautions need to be taken when
welding TWIP steels to standard low C grades. As TWIP
steel has a fully austenitic microstructure, care must be
taken to avoid solidification cracks in the welds. TWIP
steel has a smaller welding current range, and the weld
expulsion occurs at lower welding currents. In addition,
the weld nugget is lower in hardness than the base metal
hardness. Common TWIP steel weld defects included grain
boundary liquation crack in the heat affected zone and
void formation.
81
Welding TWIP steel to conventional C
steel may result in inhomogeneous temperature profiles
and weld nugget properties as the thermal (e.g. electrical
resistivity) and physical (e.g. melting temperature) proper-
ties of the two materials may be significantly different.
Al addition also decreases the welding current range
of TWIP steel, but new welding conditions have been
proposed recently.
Ultra fine grained TWIP steel
Ultra fine grained austenitic Fe–31Mn–3Al–3Si TWIP
steel has considerable ductility in contrast to ultra fine
grained Al or IF steel.
82,83
Nanostructured Fe–22Mn–
0?6C TWIP steel, obtained by a combination of cold
deformation and recovery annealing, has been reported
to have a very high yield strength and an adequate
elongation.
84
The material preparation process decreases
the dislocation density and retains the very dense
nanoscale twin microstructure.
Industrialisation of TWIP steel
The considerable interest in high Mn TWIP steels is due
to their superior mechanical properties. Compared to
conventional low C steels, high Mn TWIP steels have
high C and Mn contents. When Al is added, the content
also tends to be high. It is clear that the cost issue, partly
related to the alloy cost, will be of prime importance in
addition to the remaining technical problems related to
the manufacturing of TWIP steels. In early 1990s, the
industrial production of high Mn TWIP sheet steel faced
a large array of technical difficulties. However, the
South Korean integrated steel producer POSCO
85
has
recently shown that high Mn TWIP sheet steel could be
produced in conventional steel strip production facilities
which usually process ferritic low C sheet steel, if some
key precautions and working methods are used. This
achievement is illustrated with the schematic flowchart
shown in Fig. 23, which is used to produce industrially
TWIP steel. One of the challenges is the ferromanganese
alloy required for the steelmaking process. Standard
ferromanganese has a high P content and is unsuitable
for the production of quality TWIP sheet steel products.
Ultra low P content ferromanganese alloys, with P
contents of 0?05%, have been developed for the
production of TWIP steel.
86
Steel plant practice also
requires alternative working methods for the production
of TWIP steel, such as the premelting of the ferroman-
ganese additions before alloying, and the use of liquid
rather than powder type fluxes during continuous
casting to reduce the Al
2
O
3
pickup, ensure edge quality
and avoid cracks. The oxide scale formation during
reheating should be controlled to avoid internal grain
boundary oxidation which results in surface defects or
edge crack of hot coil. The presence of Al in Fe–Mn–Al–
C TWIP steel results in the formation of some limited
amounts of (Fe,Mn)Al
2
O
4
intergranular oxides and
plate shaped hexagonal AlN precipitates during the slab
reheating stage. No specific problems are encountered
during hot rolling, coiling and pickling. Cold rolling
reductions are slightly limited due to the high strain
hardening behaviour of TWIP steels. During the
recrystallisation annealing after cold rolling the recrys-
tallisation takes 10
3
s at 560uC, and only 10 s at 630uC.
The application of Zn and Zn alloy coatings by hot
dip galvanising requires special care as there are clear
indications that a MnO and Al
2
O
3
surface layer is
formed during continuous annealing and processing in a
hot dip galvanising line. These MnO and Al
2
O
3
surface
layers will very likely influence coating adhesion, and
electrolytic Zn deposition may be the preferred route for
coating TWIP steel in certain cases. Both hot dip
galvanising and electrolytic coating of TWIP steel have
been attempted and defect free Zn coatings have been
produced. The hot dip galvanising of TWIP steel
requires the use of a very low dew point to avoid bare
spot defects on the hot dip Zn coated sheet steel. If the
dew point range is 0–20uC during the annealing process,
a thick oxide layer will form on the strip surface and
bare spot defects will present in the coating. Recent
reports have shown that at a very low dew point, the
bare steel is present between MnO selective oxide
particles
87,88
rather than a continuous MnO layer.
Lowering the Mn content from 18 to 15%, will also
facilitate the hot dip coating with Zn.
22 Influence of cathodic hydrogen charging on stress–strain
curve of Fe–18Mn–1?5Al–0?6C TWIP steel: low current den-
sity hydrogen charging was carried out at 40
u
Cfor24h
De Cooman et al. State-of-the-knowledge on TWIP steel
Materials Science and Technology 2012 VOL 28 NO 5525
Alternative Zn coating routes have been suggested.
One approach is the ‘heat to coat’ method in which the
pickled cold rolled strip is annealed a first time and
repickled before a low temperature heat to coat galvanis-
ing process. The heat to coat process involves annealing
in an initially oxidising gas atmosphere in a direct fired
furnace heating section containing 1 vol.-%O
2
.The
reduction is carried out in the soaking section which
contains a 15 vol.-%H
2
atmosphere. Successful heat to
coat pilot tests have been reported for a Fe–15Mn–2?5Al–
0?7C TWIP steel with ,0?25 mass-%Si.
14
Conclusions
The present review of the properties of high Mn TWIP
steel clearly shows that high Mn ferrous alloys with
additions of C, Al and/or Si to fully stabilise the fcc
phase and control the SFE within the narrow range of
20–50 mJ m
22
, have a very wide range of mechanical
properties, making this relatively new class of steel of
interest for many automotive applications.
The physical metallurgy of TWIP steels is still relatively
unexplored and the following aspects need to receive an in
depth analysis: the twinning mechanism, texture evolu-
tion and delayed fracture. The determination of the
twinned volume fraction remains a challenge and is
needed to evaluate the different models proposed to
explain the mechanical behaviour of TWIP steels. The
distribution of the twinning as it is related to the
formation of texture components must also be given a
clear analysis. The precise mechanism of delayed fracture
is still not known. In particular, the complex interaction
of factors related to transformation, residual stresses and
the influence of hydrogen, has made the issue particularly
difficult to address. It is, however, clear that Al added
TWIP steels may be considered immune to the problem.
Some fundamental aspects remain poorly understood,
such as the precise nature of the relation between the strain
induced twinning and the SFE. The relatively limited
results based on ab initio calculations obtained up to now,
have exposed the fact that random transition metal alloys
with complex magnetic properties remain very challenging
material systems for first principles modelling.
The technological progress has, however, been
impressive, and it has been shown that TWIP steel can
be processed successfully in existing industrial facilities.
Acknowledgements
This research was supported by WCU (World Class
University) program through the National Research
Foundation of Korea funded by the Ministry of
Education, Science and Technology (grant no. R32-
10147). The authors also gratefully acknowledge the help
of Dr F. Barlat, Dr Yuri Estrin (Monash University,
Australia), Dr Seok-Jae Lee, Dr Chen Lei, Dr Xu Le, Mr
Jinkyung Kim and Mr Sangwon Lee.
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Materials Science and Technology 2012 VOL 28 NO 5527
