# Design of a High Strength, High Ductility 12 wt% Mn Medium Manganese Steel
With Hierarchical Deformation Behaviour

T. W. J. Kwok  K. M. Rahman  X. Xu  I. Bantounas  J. F. Kelleher  S. Dasari
T. Alam  R. Banerjee  D. Dye  Department of Materials, Royal School of Mines,
Imperial College London, Prince Consort Road, London, SW7 2BP, United Kingdom
ISIS Neutron and Muon Source, Rutherford Appleton Laboratory, Didcot, OX11
0QX, United Kingdom  Department of Materials Science and Engineering,
University of North Texas, 1155 Union Circle #305310, Denton, TX 75203-5017

###### Abstract

A novel medium Mn steel of composition Fe-12Mn-4.8Al-2Si-0.32C-0.3V was
manufactured with 1.09 GPa yield strength, 1.26 GPa tensile strength and 54%
elongation. The thermomechanical process route was designed to be industrially
translatable and consists of hot and then warm rolling followed by a 30 min
intercritical anneal. The resulting microstructure comprised of coarse
elongated austenite grains in the rolling direction surrounded by necklace
layers of fine austenite and ferrite grains. The tensile behaviour was
investigated by in-situ neutron diffraction and the evolution of
microstructure studied with Electron Backscattered Diffraction (EBSD). It was
found that the coarse austenite grains contributed to the first stage of
strain hardening by transforming into martensite and the fine austenite
necklace grains contributed to the second stage of strain hardening by a
mixture of twinning and transformation induced plasticity (TWIP and TRIP)
mechanisms. This hierarchical deformation behaviour contributed to the
exceptional ductility of this steel.

††journal: Materials Science and Engineering: A

##  1 Introduction

Mn-containing steels exploiting twinning to provide hardening and therefore
ductility date back to Hadfield steels, discovered by Sir Robert Hadfield in
1888 [1]. Current high Mn steels or TWIP steels are fully austenitic steels
that deform by forming nano-scale twins in addition to dislocation glide [2,
3]. The deformation twins are thought to provide a dynamic Hall-Petch effect
where the twins continuously subdivide the grain [4] and provide high strain
hardening rates of up to 3 $\usk\mathrm{G}\mathrm{Pa}$ and exceptional
ductility, often in the excess of 50% [5].

These mechanical properties made TWIP steels very attractive for energy
absorbing applications such as automotive crash assemblies and armour plate
[6]. However, the high Mn content, up to 30 wt% Mn [7], poses many problems in
secondary steel melting and downstream processing such as (1) macrosegregation
in slabs resulting in poor hot roll-ability [8], (2) hydrogen cracking [3] and
(3) poor weldability, especially in Zn-coated sheets [3, 7, 9]. Medium Mn
steels containing $6-12$ wt% Mn have been recently gaining significant
attention in an attempt to solve this problem. The alloying concept has
therefore been to significantly reduce the amount of Mn while retaining the
exceptional mechanical properties of high Mn steel.

Many researchers have been succesful in retaining the TWIP effect with
significantly lowered Mn content [10, 11, 12, 13]. Lee et. al. [11] reported
twins in a medium Mn steel with only 6 wt% Mn. However, the unique feature of
medium Mn steels is not only the ability to retain the TWIP effect with less
alloyed Mn but the ability to exhibit both TWIP and TRIP effects [12, 11, 14].
TWIP and TRIP effects have been observed to occur either in succession (i.e.
TWIP then TRIP), or simultaneously. Sohn et. al. [15] showed that twinning and
martensitic transformation were able to occur simultaneously within the same
austenite grain. However, simultaneous TWIP and TRIP is not as commonly
reported as the successive TWIP then TRIP mechanism. In the successive TWIP
then TRIP mechanism, the steel initially deforms by the TWIP mechanism and
with increasing strain, twin intersections serve as potent nucleation sites
for strain induced martensitic transformation and the TRIP mechanism is
activated [12].

Whilst the successive TWIP+TRIP effect has been studied in metastable
austenitic stainless steels [16], the conditions in which this effect is
active in medium Mn steel are less well understood. It is gathered that the
TWIP+TRIP effect depends strongly on both alloy chemistry and microstructure
[12, 17, 18, 10], which is unsurprising since these two factors significantly
influence stacking fault energy (SFE) and austenite stability which govern the
TWIP and TRIP effects respectively. In turn, thermomechanical processing
influences the microstructure and phase chemistry through element partitioning
during processing and heat treatment in the intercritical $\gamma+\alpha$
temperature region.

Most thermomechanical processing methods for medium Mn steel in the literature
aim to overcome the problem of slow Mn diffusivity in both austenite and
ferrite [19, 20, 21]. Incomplete partitioning of Mn may lead to undesired SFE,
austenite stability, non-equilibrium phase fractions and ultimately poor
mechanical properties. To overcome this problem, one of the most common
processing approaches are of hot roll - quench - cold roll - intercritical
anneal (IA) type [12, 22, 15]. Cold rolling prior to IA is beneficial as it
introduces a large density of defects which allows for recrystallisation and
rapid partitioning of elements within a short period of time during
intercritical annealing. Lee et. al. [19] showed that in a 6 wt% Mn alloy,
partitioning of Mn was nearly complete, i.e. reaching thermodynamic
equilibrium, after cold rolling and IA for only 180 s. Without cold rolling,
the desired extent of partitioning may only be reached after several hours of
annealing or not at all [23, 20]. Warm rolling as an alternative to cold
rolling of medium Mn steel has been explored [13, 24]. Hu et. al. [13]
developed a 7 wt% Mn steel with 63% ductility through warm rolling at 700 °C
and a 5 hr IA also at 700 °C. The Mn content in austenite grains was found to
be between 9 and 10.9 wt%, which is similar to the prediction by Thermo-Calc
(11 wt%) under equilibrium conditions, demonstrating the potential of warm
rolling to replace cold rolling.

However, continuous thermomechanical processing without the need for quenching
or cold rolling has not received much attention although it is the most
economical and widespread method of producing sheet steel in industry. In this
study on a 12 wt% Mn medium Mn steel, we present a novel continuous
thermomechanical processing route and examine the evolution of microstructure
with each processing step. The tensile behaviour and deformed microstructures
were also examined with in-situ neutron diffraction and EBSD.

##  2 Experimental Procedures

The steel used in this work was of nominal composition
Fe-12Mn-4.8Al-2Si-0.3C-0.3V (wt%) prepared using pure elements and vacuum arc
melted to produce 60$\times$10$\times$10 mm bars. The bars were quartz
encapsulated and homogenised in low pressure Ar at 1200 °C for 24 h and water
quenched. Thermomechanical processing was conducted at 900 °C in five passes
(hot rolling), 700 °C in four passes (warm rolling) and annealed at 700 °C for
30 min before air cooling to room temperature. All reductions were
approximately 20% per pass and the final thickness of the strip was
approximately 1 mm. To chart the evolution of microstructure along the
processing schedule, several bars were quenched after the 5th, 7th and 9th
pass before IA. A schematic of the rolling schedule and specimens used for
microstructual analysis are shown in Figure 1a.

Sub-sized tensile specimens with gauge dimensions of 19$\times$1.5$\times$1 mm
were cut via Electric Discharge Machining (EDM) such that the gauge was
parallel to the rolling direction. Tensile testing was conducted on an Instron
5960 load frame equipped with a 30 kN load cell at a nominal strain rate of
$10^{-3}\>s^{-1}$. To investigate the microstructural changes during uniaxial
deformation, interrupted tensile tests were conducted to 0.05, 0.18 and 0.28
true strain ($\epsilon_{t}$).

Scanning Electron Microscopy (SEM), EBSD and Energy Dispersive Spectroscopy
(EDS) measurements were performed on a Zeiss Sigma FE-SEM equipped with a
Bruker EBSD detector and Bruker XFlash 6160 EDS detector. Transmission
Electron Microscopy (TEM) was conducted on a JEOL JEM-2100F. TEM samples were
obtained by Focused Ion Beam (FIB) milling and lifting out a foil in a FEI
Helios Nanolab 600 focused ion beam workstation. Specimens for Atom Probe
Tomography (APT) were prepared by FIB milling and lifting out a foil before
mounting the small sections of the samples on suitable holders for analysis.

The APT experiments were conducted on a CAMECA local electrode atom probe
5000XS instrument. All experiments were performed at a temperature of 30 K
with a target evaporation of 0.6% and laser power of 30 nJ with a steady-state
applied DC voltage. APT data reconstruction and analysis were carried out
using CAMECA IVAS 3.8.4 software. The mass spectrum for the APT reconstruction
was calibrated based on the average phase compositions obtained from SEM-EDS.
The mass-to-charge ratio peaks were assigned as follows: C2+ (6, 6.5 Da), C+
(12, 13 Da), CC+ (24, 26 Da), CCC+ (18, 18.5, 36 Da), Al3+ (9.0 Da), Al2+
(13.5 Da), Al+ (27 Da), Si2+ (14, 14.5, 15 Da), V3+ (17 Da), V2+ (25.5 Da),
Mn2+ (27.5 Da), Mn+ (55.0 Da), Fe2+ (27, 28, 28.5, 29 Da), Fe+ (54, 56, 57, 58
Da) and CCFe2+ (40 Da). The overlap in Fe2+ peaks with Si+ was ignored as the
intensities matched with the relative abundance. The 27 Da peak was
proportionally split between Al+ and Fe2+ based on the relative intensities of
remaining three Fe2+ peaks. While it has been shown that lower specimen
temperature is preferred for better quantitative analysis [25], Takahashi et.
al. [26] and Miyamoto et. al. [27] have observed an increase in carbon
concentration in steels with decrease in specimen temperature. In this study,
a relatively lower specimen temperature (30 K) was used. With these
considerations in mind, it is acknowledged that there will be some uncertainty
in the elemental compositions during the APT reconstruction of the specimen.
Nevertheless, APT provides superior spatial resolution and the results are
still useful in the investigation of relative element partitioning between the
different phases in medium Mn steel, which is the principal focus.

Figure 1: (a) Thermomechanical processing schedule showing extracted samples,
i.e. stages 1-4. Backscattered electron (BSE) micrographs of the steel after
(b) hot rolling at 900 °C where dotted lines indicate elongated prior
austenite grain boundaries, (c) 2 passes at 700 °C where white arrows point to
the first necklace layers which have started to form, (d) 4 passes at 700 °C
and (e) IA at 700 °C for 30min. (f) Etched secondary electron micrograph of
(d) and (g) EBSD IQ+PM scan of (e). Figure 2: (a) Schematic of the in-situ
neutron diffraction tensile experiment. (b) Secondary electron micrograph of
the neutron tensile specimen at 0% strain. White arrow indicates a large
precipitate.

In order to further study the tensile behaviour of the steel, another ingot of
the same composition measuring 60$\times$23$\times$23 mm was homogenised in a
vacuum furnace at 1200 °C for 24 h, furnace cooled and processed according to
a similar rolling schedule. The ingot was hot rolled at 900 °C from 23 mm to
12.5 mm in 4 passes of approximately $15-20$% reduction per pass and warm
rolled at 700 °C from 12.5 mm to 6 mm in 4 passes of approximately $15-20$%
reduction per pass before a 30 min IA at 700 °C.

A tensile specimen for neutron diffraction with gauge dimensions of
50$\times$6$\times$6 mm and sub-sized tensile specimens were similarly cut via
EDM with the gauge parallel to the rolling direction from the same strip. The
in-situ neutron diffraction experiment was conducted at the ENGIN-X
diffractometer at the ISIS Neutron and Muon Source, Rutherford Appleton
Laboratory (Oxfordshire, UK). A tensile test was performed at room temperature
using a 100 kN load frame with the loading axis placed on the bisection of the
incident and diffracted beams (Figure 2). The tensile test was performed in
load control in the elastic region and displacement control mode during
plastic elongation. Loading was paused for each measurement. Diffracted peaks
obtained in the detector correspond to the loading of grains with scattering
plane normal to the loading direction and were simultaneously collected by the
north and south detector banks, both positioned at the horizontal scattering
angles of $\pm\>90\mathrm{\SIUnitSymbolDegree}$. In this study, only the
patterns recorded by the north detector bank were used, which depicts the
longitudinal lattice strains of the lattice planes with the normal in the
tensile direction. The elastic lattice strain, $\epsilon^{hkl}$, of grains in
scattering condition {hkl} can be obtained from shift in peak positions
according to $\epsilon^{hkl}=(d^{hkl}-d_{0}^{hkl})/{d_{0}^{hkl}}$, where
$d_{0}^{hkl}$ and $d^{hkl}$ are the unstrained and strained lattice parameters
respectively.

##  3 Results

###  3.1 Microstructural evolution during processing

Medium Mn steels have great flexibility in microstructural design and can
exhibit very different microstructures depending on the processing route. At
900 °C, the steel was fully austenitic and at stage 1 in Figure 1b, the
microstructure consisted of elongated grains, suggesting that
recrystallisation of austenite was limited at this temperature. When the
temperature was reduced to 700 °C and after the first 2 warm passes (stage 2),
a single layer of very fine grains was observed to have nucleated on the Prior
Austenite Grain Boundaries (PAGB). After stage 3, it was observed that
additional layers of fine grains had formed along the PAGB, partially
consuming the prior austenite grain. Etching in 2% nital solution clearly
revealed that the layers consisted of two phases (Figure 1f) which were likely
austenite and ferrite, with ferrite also nucleating along shear bands in the
deformed parent austenite grain. The final microstructure after the full
processing schedule could not easily be discerned in the BSE micrograph but
the EBSD image quality and phase map (IQ+PM) clearly showed bands of elongated
austenite and also bands of fine austenite and ferrite grains. The grains in
the fine grained region will be termed necklace grains as they formed in
layers or necklaces, closely resembling the necklace-type recrystallisation
mechanism [28]. The retained elongated austenite grains will be termed the
core austenite grains because the necklace recrystallisation mechanism was
incomplete, leaving an austenite “core” in the middle of a prior austenite
grain.

Table 1: Composition of the bulk steel measured by ICP and individual phases
measured by SEM-EDS and APT. Uncertainties shown in parentheses. $\dagger$The
C content of the bulk steel was measured by Time of Flight Mass Spectrometry
(TOFMS). The C content in necklace and core austenite EDS measurements was
assumed to be the equal and calculated by the lever rule on the basis that C
has minimal solubility in ferrite.

Mass percent (%) SFE Md Mn Al Si V C$\dagger$ mJ m-2 °C Bulk (ICP) 11.99 4.84
1.98 0.28 0.32 - - Bulk (EDS) 10.3 (0.1) 4.9 (0.1) 1.3 (0.1) 0.2 0.32 - -
$\gamma$ Thermo-Calc 13.90 4.16 1.40 0.003 0.57 36.6 - $\gamma$ core (EDS)
10.0 (0.5) 4.8 (0.3) 1.3 (0.1) 0.2 (0.1) 0.52 36.0 250 $\gamma$ necklace (EDS)
11 (2) 4.0 (0.7) 1.0 (0.4) 0.3 (0.3) 0.52 31.1 225 $\alpha$ necklace (EDS) 7.8
(0.8) 4.6 (0.4) 1.3 (0.2) 0.3 (0.3) 0 - - Atomic percent (%) SFE Md Mn Al Si V
C$\dagger$ mJ m-2 °C $\gamma$ Thermo-Calc 13.07 7.96 2.57 0.003 2.45 36.6 -
$\gamma$ necklace (EDS) 11(1) 8 (1) 1.9 (0.7) 0.3 (0.3) 2.25 31.1 225
$\gamma_{1}$ necklace (APT) 13.28 (0.07) 10.20 (0.06) 2.77 (0.03) 0.05 1.63
(0.02) 36.3 273 $\gamma_{2}$ necklace (APT) 13.31 (0.07) 9.48 (0.06) 2.97
(0.03) 0.04 1.37 (0.02) 30.1 298 $\alpha$ necklace (EDS) 7.5 (0.8) 9.0 (0.4)
2.4 (0.2) 0.3 (0.3) 0 - - $\alpha$ necklace (APT) 8.28 (0.06) 14.53 (0.08)
3.78 (0.04) 0.03 0 - -

Figure 3: (a) EBSD IQ+PM and corresponding (b) qualitative SEM-EDS map of a
specimen processed according to Figure 1a strained to $\epsilon_{t}=0.05$.
Locations of strain induced martensite are indicated by the white arrows.

As the deformation mechanisms active in medium Mn steels are heavily dependent
on composition, the composition of the studied steel was measured with three
different methods at different length scales and is shown in Table 1. The bulk
composition was measured by Inductively Coupled Plasma Spectroscopy (ICP) and
SEM-EDS. SEM-EDS was also used to map the distribution of Mn in the
microstructure (Figure 3) and to probe the composition of individual core
austenite, necklace ferrite and austenite grains. Finally, APT was used to
probe the local composition within one necklace ferrite and two necklace
austenite grains.

Compositions measured by SEM-EDS were averaged over a minimum of 30 grains.
Due to the difficulty in measuring C content in SEM-EDS, the C content in
ferrite was assumed to be negligible and C content in austenite was determined
by the lever rule. C content in a necklace ferrite grain was indeed confirmed
to be negligible through APT as shown in Table 1. The C content in both core
and necklace austenite grains were assumed to be equal due to the fast
diffusion kinetics of C during partitioning [29] but some slight differences
between necklace austenite grains was observed in the APT measurement. The SFE
and Md temperature were calculated according to the method developed by Pierce
et. al. [30] and Nohara et. al. [31, 32, 18] accordingly. The thermodynamic
equilibrium composition of austenite at 700 °C was determined through Thermo-
Calc using the TCFE 7.0 database based on the bulk composition obtained by
SEM-EDS.

Figure 4: (a) Mn ion map with 9.3 at% iso-surface obtained from the necklace
region. (b) Proxigram analysis of both the top and bottom interfaces. Far
field analyses of the compositions of the $\alpha$, $\gamma_{1}$ and
$\gamma_{2}$ necklace austenite grains are given at the edges of (b). Figure
5: EBSD IQ+PM of interrupted tensile specimens obtained parallel to the
transverse direction at (a) $\epsilon_{t}=0$, (b) $\epsilon_{t}=0.05$, (c)
$\epsilon_{t}=0.18$, (d) $\epsilon_{t}=0.28$ and (e) $\epsilon_{t}=0.43$.
White arrows indicate strain induced blocky $\alpha$’ martensite. (f) EBSD
IQ+PM map at $\epsilon_{t}=0.43$ but obtained parallel to the normal
direction. White arrows indicate locations of completely transformed core
austenite grains. (g) True stress-strain curve and hardening rate. (h) Change
in austenite, ferrite/martensite and non-indexed (N.I.) area fraction obtained
from (a-e), the dotted line and open symbol represents the volume fractions
obtained in (f). Figure 6: Engineering stress strain curve for the steel
studied.

When comparing the EBSD IQ+PM map and Mn EDS map in Figure 3, it can be seen
that the fine necklace austenite grains were enriched in Mn and the ferrite
grains were Mn depleted. The core austenite can be distinguished as having an
intermediate level of Mn and from Table 1, the core austenite grains had a
very similar composition to the bulk composition. This suggests that
partitioning did not occur significantly in these grains. SEM-EDS measurements
of the necklace austenite grains showed that they were slightly enriched in Mn
and depleted in Al and Si, suggesting that partitioning occurred between the
necklace austenite and ferrite grains to a greater degree. However, the
composition of the necklace austenite grains measured by APT revealed that the
extent of partitioning was much greater than that measured by SEM-EDS. Due to
the small size of the austenite and ferrite grains within the necklace region,
the electron interaction area associated with SEM-EDS may be larger than a
single grain and the composition of any carbides and neighbouring ferrite
grains could have been sampled during the SEM-EDS measurement of individual
necklace austenite grains. The SEM-EDS measurement of necklace austenite
grains can therefore be interpreted as an “average” composition of these
grains. APT also revealed that very little V remained in solution in both
phases and likely precipitated out as carbides as shown in Figure 4a of
approximate stoichiometry V2C.

As a result of partitioning, core austenite grains and necklace austenite
grains will have different SFE. From the SEM-EDS measurements, the core
austenite grains had a larger SFE than the necklace austenite grains due to
the higher Al content [30, 33]. Md temperature of the necklace austenite
grains was lower and therefore more stable against transformation. However,
the APT measurements showed that there was some variation in SFE between two
necklace austenite grains, which may indicate a spread of SFE amongst necklace
austenite grains. Nevertheless, the core austenite grains and necklace
austenite grains still possess SFEs within the twinning regime [34, 3].

###  3.2 Mechanical behaviour and deformation microstructure

EBSD phase maps from interrupted tensile tests were obtained to investigate
the microstructural evolution during tensile deformation in Figure 5, and the
engineering stress-strain tensile curve is shown in Figure 6. The steel
possessed a yield strength of 1090 $\usk\mathrm{M}\mathrm{Pa}$, tensile
strength of 1260 $\usk\mathrm{M}\mathrm{Pa}$ and elongation of 54%. It is
evident from the true stress - true strain tensile curve (Figure 5g) that the
steel underwent two stages of hardening. In the first stage, Stage A
($\epsilon_{t}=0.02\,-\,0.25$), the tensile curve was characterised by
unsteady flow and a fluctuating strain hardening rate. In the second stage,
Stage B ($\epsilon_{t}=0.25\,-\,0.43$), there was steady hardening with a near
constant hardening rate of 2.8 $\usk\mathrm{G}\mathrm{Pa}$ up to the point of
fracture.

Due to the difficulty in resolving ferrite and $\alpha$’ martensite in EBSD,
both ferrite and $\alpha$’ martensite are coloured green in Figure 5. However,
$\alpha$’ martensite can be distinguished from ferrite to some degree from
correlative EBSD and SEM-EDS maps [35] as shown in Figure 3. Strain induced
$\alpha$’ martensite has a blocky morphology and possesses the same
composition as the core austenite from which it transformed from as compared
to ferrite which is equiaxed in morphology and is Mn depleted. Selected
regions which could be confidently identified as strain induced $\alpha$’
martensite are indicated by the white arrows in Figure 5.

Unfortunately, due to the large strains involved, a significant area ($<$0.2)
was not indexed in EBSD and so microstructural results will be interpreted and
discussed qualitatively. From $\epsilon_{t}=0\,-\,0.28$, the austenite volume
fraction steadily decreased while ferrite content steadily increased. This
suggests that TRIP was occurring within the austenite phase and from Figure
5b-c, the TRIP effect was observed to be mostly confined within the core
austenite grains. At $\epsilon_{t}=0.43$, the indexing rate was very poor
parallel to the transverse direction but improved drastically when scanned
parallel to the normal direction. In Figure 5f, it can be seen that the core
austenite regions have completely transformed and any untransformed austenite
was confined to within the necklace zones. When comparing the volume fractions
obtained between Figures 5d and f, it is evident that the ferrite content
continued to increase, indicating continued TRIP.

In the specimens deformed to fracture and at $\epsilon_{t}=0.28$, twin-like
features were observed in some of the necklace grains under BSE imaging,
presumed to be necklace austenite. Twinning was indeed confirmed within the
necklace austenite grains when a FIB foil was lifted out from the necklace
region of a specimen deformed to $\epsilon_{t}=0.28$. Figure 7 shows the
evidence of twinning under BSE imaging and also confirmation of twins under
TEM imaging.

Figure 7: (a) BSE micrograph showing evidence of twinning in grains marked by
the white arrow. (b) Diffraction pattern with a [110]γ beam direction showing
matrix and twin spots obtained from a necklace austenite grain when deformed
to $\epsilon_{t}=0.28$. Streaks were observed normal to the twin plane. Centre
beam dark field micrograph obtained from the (c) austenite matrix spot and (d)
twin reflection spot. Figure 8: Engineering stress-strain curves of the
tensile specimen used for in-situ neutron diffraction and of the sub-sized
tensile specimen obtained from the same strip.

###  3.3 In-situ neutron diffraction measurements

Figure 8 shows the tensile curve obtained from the in-situ neutron diffraction
tensile test, the time-averaged measurements and the tensile curve from the
sub-sized specimen from the same ingot. It is noted that while the tensile
properties obtained in the full sized and sub-sized neutron tensile specimens
differed from those shown in Figure 6, the two-stage strain hardening
behaviour and the necklace-core microstructure was preserved (Figure 2b).
There are two likely reasons for the difference in tensile properties. The
first reason was the difference in cooling rates after the homogenisation heat
treatment. The ingot used to produce the tensile specimen for the in-situ
neutron diffraction tensile test was furnace cooled, instead of water
quenched, which allowed V to precipitate as carbides and coarsen significantly
(Figure 2b). The coarse precipitates would not be able to aid in grain
refinement nor precipitation strengthening and possibly act as crack
initiation sites. Secondly, warm rolling reductions were lower due to
limitations of the laboratory rolling mill and the necklace layers were
therefore not as developed. As a result, the neutron diffraction experiment
was only able to proceed to approximately 9% engineering strain which captured
stage A type hardening and a small amount of stage B type hardening
($\epsilon=7.7$ to $8.9$%).

Figure 9 shows the stress dependence of lattice strain in the austenite (FCC)
and ferrite (BCC) phases. The austenite peaks showed a linear increase in
lattice strain with increasing stress up to approximately 430 MPa, whereupon
the (220)γ peak showed a continuous increase in stress dependence with
increasing strain up to the macroscopic yield stress. At 530MPa, the stress
dependence of the (111)γ and (220)γ peaks decreased while the other peaks
continued to increase linearly. At the macroscopic yield stress, all the FCC
peaks experienced a decrease in lattice strain at constant stress before
increasing again, with the exception of the (220)γ peak where the lattice
strain continuously increased at the macroscopic yield stress and then
decreased with increasing stress.

Figure 9: Lattice strain dependence on engineering stress of (a) FCC (b) BCC
phases. Figure 10: Normalised neutron diffraction integrated peak of (a) FCC
and (b) BCC phases within the plastic strain region.

The BCC peaks showed a linear increase in lattice strain with increasing
stress up to 490 MPa where the stress dependence of the (110)α peak increased
and at 590 MPa, the stress dependence of the (200)α peak increased
dramatically. At the macroscopic yield stress, the lattice strain of the
(110)α and (222)α peak fluctuated at constant stress before increasing
slightly at higher stresses. The (200)α peak however, continued to increase in
lattice strain dramatically at the macroscopic yield stress before increasing
again with increasing stress at higher stresses. The ferrite produced by the
martensitic TRIP will be of different composition and d-spacing to the ferrite
existing after IA, therefore limited reliance should be placed on the
interpretation of the BCC lattice strains beyond the macroscopic yield point.

The elastic regions of the diffraction elastic constants (DECs) calculated
from the curves in Figure 9 are shown in Table 2. It is evident that the
austenite is stiffer than the ferrite phase overall, and therefore the
austenite will shield the ferrite from stress to a certain extent.

From Figure 10, the integrated intensity of FCC peaks exhibited a slight
decrease during initial plastic strain. While the maximum intensity has
decreased (Figure 11c), the increase in width resulted in a near constant
integrated intensity. The BCC peaks exhibited more varied behaviour. The
integrated intensity of the (200)α peak increased slightly just before 2%
strain and increased gradually up to 5% and remained fairly constant. Most
notably, the (110)α peak intensity increased significantly at 1.7% strain and
increased sharply again at 7.7% strain, which corresponds to the beginning of
stage B hardening.

Table 2: Table of DECs and anisotropy of resolved planes. The anisotropy of cubic crystals scale with 3$\Gamma$ where $3\Gamma=3(h^{2}k^{2}+h^{2}l^{2}+k^{2}l^{2})/(h^{2}+k^{2}+l^{2})^{2}$ [36]. Uncertainties shown in parantheses. Plane | Modulus (GPa) | 3$\Gamma$
---|---|---
(200)γ | 139 (3) | 0
(311)γ | 164 (3) | 0.47
(220)γ | 176 (4) | 0.75
(111)γ | 224 (4) | 1
(222)γ | 216 (9) | 1
(200)α | 119 (9) | 0
(110)α | 169 (7) | 0.75
(222)α | 182 (16) | 1
Figure 11: (a) Full neutron diffraction pattern and measured peaks when, from
bottom to top, unstrained, at macroscopic yield point and fully deformed
respectively. (b) BCC (200) peak. (c) BCC (110)α and FCC (111)γ peak.

##  4 Discussion

###  4.1 Processing and microstructure

The investigated necklace-core type microstructure arose from a continuous hot
roll - warm roll - anneal process. Necklacing or partial recrystallisation is
common in the hot or warm deformation of stainless steel [37] but has not been
commonly reported in medium Mn steel. He et. al. [38] and Hu et. al. [24]
obtained a similar partially recrystallised necklace-core type microstructure
in a Fe-10Mn-2Al-0.5C-0.7V steel that was hot rolled, cooled to room
temperature and subsequently warm rolled at various temperatures and
reductions followed by a long IA of 5 h. While the mechanical properties were
different, the resulting strain hardening behaviour was remarkably similar
with stage A and B-type strain hardening, suggesting the importance of
microstructure on the tensile behaviour of medium Mn steels with necklace-core
type microstructures.

The key processing step was the warm rolling step at 700
$\mathrm{\SIUnitSymbolDegree}$C where several processes were observed to
occur: nucleation of ferrite and recrystallisation of austenite, i.e.
formation of the necklace and element partitioning. It is known that ferrite
transformation from austenite is sluggish when austenite is held isothermally
within the intercritical temperature regime [20]. Warm rolling introduces
defects as nucleation sites for ferrite formation, which explains why ferrite
was largely observed to be located in the necklace layer (i.e. PAGB) and along
shear bands within the prior austenite grain. Because austenite grains were
also found in the necklace layer, recrystallisation of austenite was thought
to be occuring simultaneously with ferrite transformation via the necklace
recrystallisation mechanism. As warm rolling progressed, more necklace layers
were added, consistent with the findings by Ponge and Gottstein [28]. However,
recrystallisation was incomplete, resulting in austenite cores which are
likely highly deformed. The core austenite grains were likely able to recover
during the IA at 700 °C but are still presumed to have a much higher
dislocation density compared to the necklace grains, contributing to the high
yield strength of the steel.

Partitioning was also thought to be enhanced by warm rolling. From Figure 3
and Table 1, the necklace austenite grains were observed to have a
significantly higher Mn content than the core austenite grains. APT
measurements also showed that the Mn content in the necklace austenite grains
were very close to the thermodynamic equilibrium. A possible reason was due to
the small grain sizes of the necklace grains which considerably shortens the
diffusion distance between neighbouring ferrite grains [39]. Another reason
could be due to pipe diffusion of Mn [40, 41]. As the necklace grains formed
as warm rolling progressed, it is likely that these grains also experienced a
certain degree of deformation. While not enough to trigger recrystallisation
within the necklace, the dislocation density is held to be sufficient to
enhance partitioning between austenite and ferrite within the necklace via
pipe diffusion. Necklace grains which formed early during warm rolling would
have experienced a larger amount of warm deformation as compared to the
necklace grains which formed later on or during the final IA. As a
consequence, it is likely that partitioning would have occurred to a greater
extent in the early necklace austenite grains. This effect may explain the
slight difference in compositions between necklace austenite grains measured
by APT shown in Table 1 and Figure 4.

###  4.2 Tensile behaviour and microstructure

The two-stage tensile behaviour of the steel was investigated through an in-
situ neutron diffraction tensile test and the evolution of microstructure
during tensile deformation was investigated with EBSD through a series of
interrupted tensile tests.

Stage A of the strain hardening curve is characterised by negligible increase
in engineering stress with applied strain and a serrated strain hardening
rate. Many researchers associate this post-yielding serration with the TRIP
effect [42, 43, 24]. Zhang et. al. [43] demonstrated that before uniform
elongation, retained austenite undergoes the TRIP effect caused by the passage
of a Lüders band. The increase in BCC (110)α peak intensity in Figure 10b also
indicates that martensitic transformation was active.

Microstructurally, it can be seen from Figure 5 that the TRIP effect in stage
A was largely confined within the core austenite grains. This can be explained
by the difference in austenite stability between the core and necklace
austenite grains (SEM-EDS measurement in Table 1). The core austenite having
less Mn and a significantly larger grain size is less stable and more likely
to TRIP. The core austenite also possess a grain size of $>$10 $\mu$m, and
will therefore be less strong than the necklace austenite which, like the
necklace ferrite, is much smaller than 1 $\mu$m in size. It should be noted
that Md calculations from Table 1 were based on empirical calculations based
on austenitic stainless steels dating back to Angel [44] in 1954 and Nohara
et. al. [31] in 1977 and may not be fully applicable in medium Mn steels. Md
temperatures in Table 1 should therefore be interpreted conceptually.

It is also noteworthy that while the SFE of the austenite core was within the
twinning region, TRIP rather than TWIP was observed. Several factors could be
at work. The SFE calculation may be erroneous as the empirical model [30] used
in this study was not developed for use in medium Mn steel but more
importantly, the deformation mode in medium Mn steels cannot be determined by
SFE alone, unlike TWIP steels which do not transform [45]. In this scenario,
the austenite stability of the core austenite grains was such that strain
induced martensitic transformation was preferable to twinning.

During stage A hardening, the FCC peaks in Figure 9a, with the exception of
(220)γ experienced a decrease in lattice strain, indicating that a stress-
relieving mechanism, i.e. TRIP, was active. This result was also observed by
Zhang et. al. [43] in a medium Mn TRIP steel. It is likely that load
partitioning after yielding took place with the martensite which formed from
the core austenite grains taking on the majority of the load. It is worth
noting from Figure 10b that the intensity of BCC (110)α peak increased rapidly
at 1.7% strain but slowed at 7.7% strain which likely indicates the decreasing
transformation rate of the core austenite grains.

Stage B of the strain hardening curve is characterised by a more uniform but
slightly fluctuating strain hardening rate. The core austenite having
transformed to core martensite now strengthens the steel which may explain why
the strain hardening rate in stage B is on average higher than stage A. It is
likely that in stage B, the load is now partitioned from the core martensite
grains to the necklace grains which now begin to deform. The necklace
austenite grains are of interest as both TWIP and TRIP behaviour were observed
in stage B. The twins observed in a necklace austenite grain in Figure 7
support the suggestion that TWIP was active while the increase in
ferrite/$\alpha$’ martensite fraction in Figure 5h and the second increase in
intensity of the BCC (110)α peak at 7.7% strain in Figure 10b support the
suggestion that TRIP was active. It is therefore reasonable to conclude that
both TWIP and TRIP mechanisms were active in the necklace austenite grains
during stage B. However it was not determined if the successive TWIP+TRIP
effect was active.

Several reasons may be possible for the difference in deformation mechanisms
in the necklace austenite grains. From Table 1, there is a spread in
composition in the necklace austenite grains resulting in dfferences in SFE
and austenite stability. It has also been reported that grain size could be a
determining factor whether an austenite grain deforms by TWIP or TRIP in
medium Mn steels [17, 15], however other factors such as orientation of
austenite grains relative to the loading direction [46] may affect which
deformation mechanism activates in any given grain. As the necklace austenite
composition was not at the thermodynamic equilibrium as seen in Table 1,
differences in composition between necklace austenite grains may also account
for variations in SFE and austenite stability and therefore influence the type
of active deformation mechanism.

###  4.3 Industrial relevance

Researchers and automakers are aiming to develop new steel grades with $>$1000
$\usk\mathrm{M}\mathrm{Pa}$ tensile strength and 20% elongation termed as the
3rd generation Advanced High Strength Steels (AHSS) [47]. WorldAutoSteel [48]
postulates a TWIP550/900 grade ($500-550\usk\mathrm{M}\mathrm{Pa}$ yield
strength, $980-990\usk\mathrm{M}\mathrm{Pa}$ tensile strength and $50-60$%
elongation) in the design of the Future Steel Vehicle (FSV). Many medium Mn
steels in the literature exceed this target, demonstrating the potential of
this family of steels. However, medium Mn steels will still face industrial
concerns such as cost, formability and weldability [47].

At this stage of alloy development, we introduce a comparator for performance,
U30, which represents the energy absorbed during tensile deformation which is
determined by the area under the tensile curve up to $\epsilon_{t}=0.3$, where
0.3 is determined to be the limit of useful ductility. As medium Mn steels are
designed for energy absorbing applications, U30 would serve as a useful
estimate. In Figure 12, several commercial steels used in the automotive and
armour industries are compared to medium Mn steels in the literature based on
price and U30. Steel prices were determined through prices of elements,
obtained from www.metalbulletin.com and The London Metal Exchange. Prices are
correct as of May 2019 and calculated on the basis of steel scrap with the
addition of ferroalloys or pure elements wherever appropriate. Cost of
production was not factored but is acknowledged to be a significant factor in
the overall price. Medium Mn steels are subdivided depending on whether cold
rolling was utilised in their processing. The issue with cold rolling followed
by heat treatment is that this necessitates a batch production process which
is expensive to realise commercially.

It can be seen that most commercial steels lie within the shaded region and
most medium Mn steels lie above, which indicates that most medium Mn steels
have better energy absorbtion capabilities for the same alloying price. Many
medium Mn steels also have U30 values greater than the FSV TWIP550/900
benchmark, which demonstrates the readiness of medium Mn steels in terms of
mechanical properties. In the current study, U30 was relatively high but also
in price. In order to lower the price, the use of expensive elements such as V
needs to be reduced. Mn content also needs to be lowered to $\leq 8$ wt% for
cost and processability reasons. The aim of our future work will be to reduce
the price of the alloy to $400-500$ $/t, while keeping the same U30 and a
simple thermomechanical process.

Figure 12: Comparison of cost and energy absorbtion (U30) between commercial
steels and medium Mn steels in the literature. CR - Cold Rolling.
*ThyssenKrupp dual phase steel product catalogue.

##  5 Conclusions

In conclusion, a 12 wt% Mn medium Mn steel with an engineering yield stress of
1.09 GPa and true ultimate tensile strength of 1.9 GPa was manufactured and
processed in a continuous method which consisted of hot and warm rolling
followed by a 30 min intercritical anneal. The steel possessed a partially
recrystallised or necklace-core microstructure with coarse austenite grains
forming the core and fine-grained ferrite and austenite grains forming the
necklace. It was found using EBSD, TEM and in-situ neutron diffraction that
the steel had a two-stage hierarchical deformation behaviour where the core
austenite grains first transform to martensite and the necklace austenite
grains subsequently deform by transformation and twinning-induced plasticity.
Its deformation behaviour resulted in a ductility of 54% and sustained
hardening throughout the tensile curve, with an energy absorbtion capability
per unit alloy cost competitive with the best model steels in the literature.

##  6 Acknowledgements

The authors would like to acknowledge the support from A*STAR, the Royal
Society, UK Ministry of Defence and EPSRC grant number EP/L025213/1. We would
also like to thank Felicity Dear and Felix Wolfgang Liwerski for their
contributions.

##  7 References

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