FIELD OF THE INVENTION
The present investigation relates to a novel heat treatment method of stabilizing
austenite in high strength steels. The invention further relates to a method of
increasing austenite volume fraction at ambient temperature with leaner
chemistry in steels.
BACKGROUND OF THE INVENTION
Advanced high strength steels (AHSS) possesses combination of different phases
such as ferrite, austenite, martensite, bainite, carbides, etc. The right
combination, morphology, distribution, size and volume fraction of these phases
give desirable properties to these steels.
Amongst all the phases, austenite phase plays a very unique role in AHSS.
Presence of austenite in the final microstructure can increase strength and
ductility, simultaneously. For example, in Transformation Induced Plasticity (TRIP)
steels, austenite transform into martensite during deformation and thus increases
in-service strength as well as results in ductility during forming operation by
volume expansion due to austenite to martensite transformation. In contrast, in
Twinning Induced Plasticity (TWIP) steels, where the steels is 100% austenite at
ambient temperature, the high strength and high ductility arises because of
dynamic Hall-Petch kind of phenomenon due to formation of twins during
deformation operation. Both TRIP and TWIP steels show very high strength and
high elongation and MPa. % El can range from 20000 to 50000. These high
strength-high formable steels are always in the wish-list of auto manufacturers.
In regular plain carbon steels, austenite is not generally present at ambient
temperature. However, the demand of high strength-high formable steels drives
steel researchers across the world to focus on different methods of austenite
stabilization at room temperature.
The most common method of stabilising austenite at room temperature is
addition of strong austenite stabilizers; such as Ni and/or Mn in the steels. When
added in sufficient amount, austenite can be stabilised at room temperature and
steels becomes fully austenitic at ambient temperature. Addition of Ni and Mn
increases the austenite area of the phase diagram and when added sufficiently,

the austenite region can extend up to ambient temperature. Two such examples
are austenitic stainless steels and Twinning Induced Plasticity (TWIP) steels where
Ni and Mn are used to stabilised austenite at ambient temperature, respectively.
These two types of steels fall in the category of 2nd Generation of AHSS.
In steel, addition of 1-2 % Mn along with some carbon increases austenite stability
such a way that upon fast cooling some amount of austenite remains in the
structure. This austenite is not stable rather is metastable and is thus called
retained austenite. During deformation of this steels, austenite transform into
martensite and expands in volume which give rise to excellent ductility. In this
steels the retained austenite is actually distributed in a strong matrix like bainite
or mixture of bainite/martensite which gives the strength. The example of this
kind of steel is Transformation Induced Plasticity (TRIP) steels. This falls under the
gambit of 1st Generation AHSS.
The 1st generation AHSS can yield up to strength-ductility balance of 20000 to
30000 MPa.% whereas the 2nd generation AHSS can result in strength-ductility
balance close to 50000 MPa.%. 1st generation AHSS does not cost that much with
respect to alloying addition and processing. However, 2nd generation AHSS is quite
costly due to high alloying addition and expensive processing routes.
This leads to invention of 3rd generation AHSS which is supposed to yield a
property between 1st and 2nd generation AHSS but less costly than 2nd generation
AHSS.
The primary objective of 3rd generation AHSS is always to achieve higher strength-
ductility balance with leaner chemistry. As expected, this third-generation
Advanced High Strength Steels (AHSS) focuses on stabilizing austenite in fine
ferritic matrix which can be either martensite or bainite.
This is mainly done by manipulating chemical composition as well as processing
parameters simultaneously.
Austenite stabilization can occur through diffusion of interstitial carbon in most
processing strategies. There are many strategies to promote this interstitial
diffusion. In case of hot rolled product, after hot rolling the steel is isothermally
hold at bainitic temperature range followed by cooling to room temperature.
Bainite formation is accompanied by rejection of carbon in the remaining

austenite and enriches austenite with carbon, thus increasing its stability. The
enrichment is further aggravated by addition of Si which delays carbide
formation. As a result more carbon is rejected to austenite. Finally after cooling,
due to higher carbon concentration some part of the austenite become stable at
room temperature. In this process 7-15 volume % austenite can be retained in the
room temperature microstructure. In a similar fashion, normal hot rolled steels
can be cold rolled and subjected to isothermal bainitic treatment. Like earlier
case, this results in substantial amount of retained austenite in the
microstructure. In the above cases, the total amount of alloying elements are kept
around 3-3.5 wt%. The typical composition of such steels processed by the above
route is C ~0.2 -0.4, Mn ~0.8 – 2, Si ~0.8 – 1.5. These steels may also contain other
alloying elements such as Nb, Mo etc.
The other route to stabilise austenite is diffusion of substitutional alloying
elements. Amongst them Mn and Ni are very effective. Significant austenite
fractions resulting from prolonged holding at an Intercritical temperature allowing
for manganese partioning from ferrite into austenite have been reported. The
steels compositions using this study is way below than high Mn TWIP steels.
Miller1 reported austenite fraction of 40% in 0.1C-5.7 Mn steels. Miller1 used
three techniques to stabilise austenite in his hot rolled steels; (i) annealing at
Intercritical temperature for prolonged time, (ii) cold worked and annealed at
Intercritical temperature and (iii) warm worked at Intercritical temperature. He
had reported a suabtancial amount austenite at the final microstructure by
promoting rapid diffusion of Mn in austenite, making it stable at subsequent
processing.
Moor et al. 2 subjected hot rolled steels to batch annealing cycle for 20-40 hrs. He
reported retained austenite volume fraction up to 17% in a 0.1C-5.2Mn steel, up
to 28% in a 0.1C-5.8 Mn alloy and up to 38% in a 0.1C-7.1 Mn alloy. Huang3
reported similar amount of austenite in a 0.12C-5.10 Mn steel by holding the steel
in Intercritical region for close to an hour and above. Kim4 reported an austenite
level up to 48% in a 0.1-8Mn grade steel by keeping it in the Intercritical region
more than 6 hours or so.
Till date austenite stabilization has been done by partitioning of alloying elements
such as C and Mn. Processes that enhance the diffusivity of these elements are
employed. For example cold rolled or a fine martensitic structure provides easier

paths for these elements to diffuse through and ease of partitioning. Other way is
to give sufficiently long time so that these elements, especially Mn can diffuse
and partition to austenite. It is also said that not only enriched austenite but
smaller sized austenite also increase its stability. These two reasons together are
normally hold responsible for austenite stabilization.
However, all these methods, except Millers1 warm working method, are a time
consuming process and in order to get a substantial amount of austenite close to
an hour or higher time has to be given. On the contrary, in Miller’s warm working
process, he had shown that it is possible to achieve a substantial amount of
austenite by working the material at a constant Intercritical temperature in a
0.1C-5.7Mn steel1. In warm working, the specimens were heated in lead bath at
the desired temperature and returned to the bath between each deformation
step. A total of 75% deformation was given and cooled to room temperature.
Miller reported ~32-33 volume % austenite1.
However, in the warm working process of Miller1, it was not mention the post
deformation cooling rate. This means he did not mention whether fast cooling
such as water quenching or normal air cooling is required. He also did not
mention whether continuous deformation in the Intercritical region yield the
same effect or not. He also pointed out that the phenomenon responsible for
stabilization of austenite was enhance partitioning of alloying elements due to
high temperature deformation process.
In light of the said prior art, a method has been developed to obtain austenite at
room temperature which should be fast enough and at least comparable to
Miller’s warm working process with leaner chemistry.
OBJECTS OF THE INVENTION
It is therefore an object of the present invention to propose a novel and improved
austenite stabilization method for high strength steels.
Another object of the present invention is to propose a novel method to stabilise
austenite in a fast process.
A further object of this present investigation is to improve any existing process of
austenite stabilization.

Another object of the present invention is to propose a methodology of
stabilization of austenite other than alloying element partitioning.
SUMMARY OF THE INVENTION
The investigation discloses a methodology of stabilization of austenite in a high
strength Mn added steel by a simple hot working process. A steel ingot of the
composition comprising, in weight per cent 0.1-0.2 C, 4 - 5 Mn, 0.3 – 1.0 Si, 0.3-
0.5 Al and rest being substantially iron and incidental impurity was casted in the
laboratory. The cast ingot was homogenized at 1250 °C for 4 hours followed by
forging into 30 mm X 30 mm cross-section billet in the temperature range of 1200
°C to 900 °C, the forged ingot was subsequently air cooled to room temperature.
The forged ingots were austenitized at 1200 °C for 2 hours prior to laboratory hot
rolling to 6-7 mm. This material was further heated to Ae3 + 50°C and quenched
either in water or oil or in air. The material was further heated to the Intercritical
temperature and a deformation of 30-50% was given continuously. The starting
temperature of deformation was 50% austenite and 50% ferrite. After inter-
critical deformation, the material was cooled to room temperature with a cooling
rate varying from 1°C/sec to 60°C/sec. This process resulted in an austenite
fraction of 20-40% in the final structure.
BRIEF DESCRIPTION OF THE ACCOMPANYING DRAWINGS
Fig. 1: Schematic diagram of heat treatment of the invention.
Fig. 2: Optical micrographs of deformation heat treated followed by oil quenched
sample. The white areas look like austenite after tint etching.
Fig. 3: Transmission electron micrographs of deformation heat treated followed
by oil quenched sample with austenite and bcc phase identified by their
respective diffraction patterns; (a-c) an austenite region and (d-f) a ferritic region.
Fig. 4: X-Ray diffraction peak profile of the final structure.

DETAILED DESCRIPTION OF THE INVENTION
The hot rolled heat treated steel with 20-40% austenite in the final microstructure
comprising of C >= 0.1 wt. %, Mn between 4-5 wt. %, Si 0.3-1.0 wt. % and Al 0.3-
0.5 wt. % and rest is iron and incidental impurities in steel. The final
microstructure of the hot rolled heat treated steel according to present
investigation comprises of austenite and bcc martensite/ferrite/bainite. The hot
rolled heat treated steel according to present investigation contains austenite
which is not thermodynamically stable with the stated composition. The hot
rolled heat treated steel according to present investigation contains much higher
percentage of austenite in the final microstructure than what would have been a
simple Intercritical holding and quenching step at the said temperature and for
the said time of holding.
The method of manufacturing the hot-rolled steel sheet with a mixture of
austenite + bcc martensite/ferrite/bainite microstructure where austenite volume
fraction is way higher than the equilibrium predicted value or Intercritical holding
and quenching process includes casting and hot rolling in conventional mill or
casting in ingot, followed by forging and cooling to room temperature and then
hot rolling in laboratory mill. The hot rolled product obtained after hot rolling was
heated to austenitization temperature and cool (preferably fast cooling) to room
temperature in order to have a fine martensitic structure. The sheet is further
heated to Intercritical temperature and a deformation of 30-50% was given
continuously and then the sheet was cooled to room temperature.
According to the present invention, it is possible to produce a hot-rolled steel
sheet with an austenite fraction in the final microstructure 20-40 % with only 0.1C
and <5 wt. % Mn. Due to presence of such a high amount of austenite coupled
with martensite, this steel results in a minimum yield strength of 900 MPa, tensile
strength of ranging from 920- 1100 MPa and elongation ranging from 19 -28 %.
Such a steel sheet is adaptable in a wide spectrum of industrial fields including the
automobile industry, the electric industry and the machinery industry. It is
particularly suited to manufacture automotive parts and components and other
industrial parts and components which demand a very high strength, good crash
resistance, especially bumpers, etc.

The present invention relates to a hot-rolled steel sheet further heat treated to
produce a very high non equilibrium percentage of austenite with a Mn content
as low as 4 wt. % and C as low as 0.1 wt.% with hardly addition of any other costly
alloying elements such as Ti or Nb.
Alloying additions: The essential alloying elements for the said invention are
listed below.
C: 0.1-0.2 wt. %: Carbon is one of the most effective and economical austenite
stabilising elements. However, carbon can’t be added in plenty as it will affect
other properties such as Weldability in a great deal. The present invention of the
hot rolled heat treated steel has a carbon percentage ranging from 0.1 to 0.2 wt.
%.
Mn: Mn is the most important alloying elements for the present invention.
Normally a Mn content of above 13 wt% along with 1 wt% C gives a fully stabilise
austenite at room temperature in steel. This is because Mn is an austenite
stabiliser. However, a lower addition of Mn although does not stabilise austenite
at room temperature but certainly increases austenite stability. The hot rolled
heat treated steel of the present investigation has a Mn content below 5 wt.%.
Production process: The method of manufacturing the hot-rolled steel sheet
according to the present invention consists of an ingot casting step followed by
hot forging and hot rolling to 5-7 mm thick sheet. The hot rolled sheet was further
heat treated to obtain the final microstructure comprise of 20-40 % austenite and
bcc ferritic/martensitic/bainitic structure The various processing steps are
described in their respective order below:
Casting and forging: In the present invention, the steel of the specified
composition is first cast in a laboratory induction furnace and cooled to room
temperature. The cast ingot was homogenized at 1250 °C for 4 hours followed by
forging into 30 mm X 30 mm cross-section billet in the temperature range of 1200
°C to 900 °C. The forged ingot was subsequently air cooled to room temperature.
Hot Rolling: The forged ingots were austenitized at 1200 °C for 2 hours prior to
laboratory hot rolling to 6-7 mm. The Finish Rolling temperature was kept in fully
austenitic region.

Heat Treatment: This material was further heated to Ae3 + 50°C and quenched
either in water or oil or in air. The material was further heated to the Intercritical
temperature where 50% austenite and 50% ferrite are predicted by
thermodynamic calculation using thermo-chemical software like ThermoCalc. A
deformation of 30-50% was given continuously in 3-5 passes in the temperature.
The starting temperature of deformation was 50% austenite and 50% ferrite.
After inter-critical deformation, the material was cooled to room temperature
with a cooling rate varying from 1°C/sec to 60°C/sec. The heat treatment schedule
is given in Figure 1.
Examples
Samples as per the composition of present invention have been subjected to heat
treatment with varying cooling rate after deformation heat treatment. In order to
reveal the microstructure optically, the steel samples have been subjected to a
two stage etching process by 4% Picral + few drops of HCl followed by a 10%
aqueous sodium metabisulphate solution.5 The final microstructure after this two
stage etching process should tint martensite in straw colour, bainite in bluish
black, austenite white in colour. One such typical optical micrograph of the
processed steel subjected after the above etching process are shown in Figure 2.
The presence of austenite (white colour), martensite (straw colour) and bainitic
ferrite (blue colour) are quite clearly seen. The steel samples have been further
subjected to phase identification X-Ray diffraction study to find out the volume
fraction of austenite. The X-ray peak profile is given in Figure 3. XRD peak profile
clearly shows presence of austenite phase along with the bcc phase. The
measured volume fraction of austenite in the final microstructure has been found
to be varied between 20-40 %.
The steel samples have been subjected to transmission electron microscopy study
to reveal their microstructure in nanometer length scale. The typical TEM
photographs are shown in Figure 3. Figure 3 (a-c) shows presence of an austenite
region identified by selected area diffraction pattern (SADP) where zone axis is
austenite [011] whereas in Fig 3 (d-e) shows presence of a ferritic lath (probably
bainitic ferrite lath) identified by its SADP where zone axis is bcc [111].

Figure 3: Transmission electron micrographs of deformation heat treated
followed by oil quenched sample with austenite and bcc phase identified by their
respective diffraction patterns; (a-c) an austenite region and (d-f) a ferritic region.


Figure 4: XRD peak profile of final microstructure showing presence of austenite
in (a) deformation heat treated followed by oil quenched sample with austenite
fraction ~44% and (b) deformation heat treated followed by air quenched sample
with austenite fraction ~30%
The method as per the current invention provides a method of stabilising
austenite in high strength steels with addition of lesser amount of Mn than the
(~5 wt. % Mn or less) than TWIP steel (normally Mn above 17% is added). The
method as per the current invention can yield a high volume fraction of austenite
in the final microstructure. The manufactured steel sheet processed by proposed
invention can also yield ~900MPa YS, 920 and above ultimate tensile strength and
elongation 19-28 %. The steel manufactured by the proposed invention can be
used in automotive components where crash worthiness is important such as
bumper.
The present invention has the following advantages over the prior arts:
1) The time to stabilise austenite by the present invention is far too less than the
process followed in prior arts described in [2-4]. The said prior arts take at least a
time between 20 minutes to few hours to stabilise austenite. In contrast, the
present investigation takes few minutes (~2-3 minutes) to stabilise austenite at
room temperature.
2) The prior arts [2-4] described a process of annealing in two phase region
whereas the present investigation reveals a process of stabilising austenite by
working the material in two phase region.

3) A process of stabilising austenite by warm working in the two phase has been
described in prior art [1]. However, a high amount of deformation around 75%
reductions in thickness has been suggested in [1] whereas a 30-50% reduction has
been found suitable enough to stabilise austenite in the present investigation.
4) The present investigation has also revealed that a substantial amount of
austenite can be stabilised with addition of Mn even less than 5 wt.% whereas all
other prior arts [1-4] suggest a Mn content higher than 5 wt.%.
5) The present investigation also revealed that the final austenite fraction can
vary between 20-40 % depending on the cooling rate after deformation in two
phase region. However, effect of cooling rate after warm working was not
covered in prior art [1].
6) The present investigation revealed that it is possible to obtain a stabilised
austenite at room temperature by continuous working in two phase region while
temperature drops with deformation whereas the warm working process
mentioned in prior art [1] suggest an isothermal deformation process.

WE CLAIM :
1. A process of developing high strength steel with austenite fraction in the
range of 20-40%, the process comprising:
casting a steel ingot of composition comprising (wt. %) 0.1-0.2% of Carbon,
4.0-5.0% of Manganese, 0.5-1.0 % of Silicon, 0.3-0.5 % of Al, maximum
0.005% of Nitrogen, the remaining being iron and incidental impurities;
forging the steel ingot to a temperature greater than 1200oC and cool to
room temperature;
hot rolling the forged steel to produce a steel sheet such that finish rolling
is done at a temperature at fully austenitic region where the finish rolling
temperature should be at least Ae3 + 50 (°C), where Ae3 is the temperature
at which the transformation of austenite to ferrite starts at equilibrium and
cool to room temperature;
heating the hot rolled sheet to austenitization temperature and quenching
it in air or water or oil;
heating the hot rolled sheet to inter-critical annealing range at a
temperature where austenite and ferrite fractions are close to 50% each
and deforming it continuously by 30-50%; and
cooling to room temperature with a cooling rate in the range of 1°C/sec to
70°C/sec.
2. The process as claimed in claim 1, wherein microstructure of the high-
strength hot-rolled steel sheet comprises of 20-40% austenite and rest is
ferritic such as ferrite/martensite/bainite.
3. The process as claimed in claim 1, wherein the austenite lath size varies in
the range of 100-200 nano meter.
4. The process as claimed in claim 1, wherein the ferritic structure has a size in
the range of 100-300 nm.

5. The process as claimed in claim 1, wherein the hot rolled medium Mn steel
(Mn 4-5 wt. %) is kept in inter-critical region for few minutes and
deformation is given simultaneously.
6. The process as claimed in claim 1, wherein Mn is preferably present in the
range of 4-5 wt. %.
7. The high strength steel sheet as claimed in any of the preceding claims,
wherein the hot rolled medium Mn steel (4-5 wt.%) sheet has YS of at least
900 MPa, UTS more than 920 MPa with an elongation varying in the range
of 19-28 %.