Final_Report.pdf

DEVELOPMENT OF LOW DENSITY & HIGH
STIFFNESS STEEL, (Fe-Mn-Al-C) ALLOYS.
SUBMITTED IN PARTIAL FULFILMENT FOR THE AWARD OF DEGREE
OF
MASTER OF TECHNOLOGY
(INTEGRATED DUAL DEGREE)
IN
METALLURGICAL ENGINEERING
BY
VAIBHAV JANARDHAN WAGHMARE
16144018
UNDER THE SUPERVISION OF
PROF. R. MANNA
DEPARTMENT OF METALLURGICAL ENGINEERING
INDIAN INSTITUTE OF TECHNOLOGY
(BANARAS HINDU UNIVERSITY) VARANASI-
221005

D e v e l o p m e n t o f L o w D e n s i t y & h i g h S t i f f n e s s S t e e l
Roll Number – 16144018 Year-2021
CERTIFICATE OF APPROVAL
This is to certify that the present work “Development of low density & high
stiffness steels, (Fe-Mn-Al-C) Alloys” has been carried out by VAIBHAV
JANARDHAN WAGHMARE (Roll No – 16144018) in the Department of
Metallurgical Engineering, Indian Institute of Technology (Banaras Hindu
University) under my supervision during the academic session 2020-2021.
Prof. R. Manna
Department of Metallurgical Engineering
Indian Institute of Technology
(Banaras Hindu University)
Varanasi-221005
Prof. S. Mohan
(Head of Department)

Candidate's Declaration
I hereby declare that the dissertation work entitled " Development of low density & high
stiffness steels, (Fe-Mn-Al-C) Alloys" being submitted to the Department of
Metallurgical Engineering, Indian Institute of Technology (Banaras Hindu University),
Varanasi for the partial fulfillment of the requirement for the award of Master of
Technology(IDD) in Metallurgical Engineering, carried out under the supervision of Prof.
R. Manna, Department of Metallurgical Engineering, Indian Institute of Technology
(Banaras Hindu University). All information in this report had been obtained and
presented in accordance with the academic rules and ethical conduct. I also declare that,
as required by the rules and conduct, I have fully cited and referenced all materials and
revealed that it is not original to this work. I have not submitted the same in part or in full
to this university or to any other university for the award of degree or diploma.
VAIBHAV JANARDHAN WAGHMARE DATE:
Roll No. - 16144018
IDD (Part - V)
Department of Metallurgical Engineering
Indian Institute of Technology (BHU)
Varanasi - 221005, India

COPYRIGHT TRANSFER CERTIFICATE
Title of the Thesis: Development of low density & high stiffness steesl, (Fe-Mn-Al-C) Alloys
Name of the Student: VAIBHAV JANARDHAN WAGHMARE
Copyright Transfer
The undersigned hereby assigns to the Indian Institute of Technology (Banaras Hindu
University) Varanasi all rights under copyright that may exist in and for the above thesis
submitted for the award of the " INTEGRATED DUAL DEGREE. "
Date: Signature of the Student
Place: VARANASI ("VAIBHAV J. WAGHMARE")
Note: However, the author may reproduce or authorize others to reproduce material
extracted verbatim from the thesis or derivative of the thesis for author's personal use
provided that the source and the Institute's copyright notice are indicated.

Acknowledgment
First and foremost, I would like to express my sincere gratitude and thankfulness to my
supervisor, Prof. R.Manna, Department of Metallurgical Engineering, Indian Institute of
Technology (Banaras Hindu University). His suggestions and thoughts helped me throughout
the period of my project work. It has been a great pleasure to be acquainted with his diligent
effort and professional competencies, which has inspired me to do this project.
I would also like to acknowledge the Mechanical Metallurgy Staff of the Department of
Metallurgical Engineering without whose assistance this work would have been difficult.
Lastly, I thank all those who are directly or indirectly involved in the successful completion
of this Project work.
Special thanks go to Prof. S. Mohan head of the department of Metallurgical engineering IIT
(BHU), who provided an opportunity to work in this field. It was a pleasure working at the
department of Metallurgical engineering IIT (BHU).

LIST OF FIGURES/ TABLES:
Sr.No. FIGURES
1 Process variants for producing hot rolled and cold rolled austenitic Fe-
Mn-Al-C steel strips.
2 XRD results of the three hot-rolled steels
3,4 High-temperature equilibrium phase distribution calculated by using
ThermoCal
5 Diffraction patterns
6 TE as a function of UTS in Fe-Mn-Al-C alloys for a uniform specimen
geometry corresponding to ASTM E-8 standard
7 2θ Vs Intensity for ( Fe-1.5C-9Al-30Mn )
Sr. No. Tables
1 Volume fractions of ferrite matrix, ferrite and κ -carbide colony existing
within κ -carbide band and grain boundary κ -carbide in the three hot-
rolled lightweight steels
2 Thermal Treatments for the Three Microstructural Studies: (i) Grain Size
Effect, (ii) Carbon Effect, (iii) Ferrite Effect
3 Tensile properties of steels in hot rolled condition
4 List of compounds from various sources. (Red)-Feasible for
reinforcement in the base alloy.

CONTENTS:
 CERTIFICATE ……………………………………………………..(2)
 CANDIDATE DECLARATION …………………………………..(3)
 COPYRIGHT TRANSFER CERTIFICATE ………………………(4)
 ACKNOWLEDGEMENT ………………………………………...(5)
 LIST OF FIGURES/ TABLES ……………....................................(6)
CHAPTERS
1. ABSTRACT ……………………………………………………............(8)
2. INTRODUCTION , REVIEW OF THE LITERATURE
AND OBJECTIVES……………………………….................................(9)
3. ALLOY FORMATION
3.1. OVERALL VIEW ………………………………………………..(12)
3.2. EFFECTS OF ALUMINUM CONTENT ON MICROSTRUCTURAL
MODIFICATION …………………………………………………(15)
3.3. EFFECTS OF MANGANESE CONTENT ON MICROSTRUCTURAL
MODIFICATION …………………………………………...........(19)
4. POST ALLOY PREPARATION AND STRENGTHENING MECHANISMS
4.1 SOLID SOLUTION HARDENING ……………………................(22)
4.2 GRAIN REFINEMENT ……………………………………...........(25)
4.3 PRECIPITATION HARDENING & WORK HARDENING …….(26)
5. ANALYSIS OF TENSILE PROPERTIES
5.1 FE-MN-AL-C STEELS ……………………………………...........(26)
5.2 ADDITION OF OTHER ELEMENTS IN THE BASE ALLOY….(29)
6. RESULT AND DISCUSSION ………………………………………..(32)
7. CONCLUSION ………………………………………………..............(34)
8. REFERENCES ………………………………………………………..(36)

1. ABSTRACT:
Different types of steels are used in manufacturing of automotive parts. Steels have got
attractive properties like formability, recyclability, strength which makes it completely viable
to be economical to be used in the automotive industry. There has always been a conflict in
developing an efficient automotive that has a material of high tensile strength and lower weight
is that low density. The Fe-Mn-Al-C alloys have the ability to age harden, absorb energy in a
crash and high strength all summed up makes them a suitable fit for the required alloy
development. Melting techniques, age hardening, balance between the strength and density are
the major attributes to be looked upon during alloy designing. Austenitic low density steels
contain a higher Mn content (12-30)%, Al 12%, C (0.6-2)%. This type of alloy can have a fully
equiaxed-austenitic microstructure at hot working temperatures and the austenite is metastable
after fast cooling. Also the austenitic Fe-Mn-Al-C steel provides better combination of density
and tensile strength as compared to any other form of steels present.

2. INTRODUCTION:
The physical metallurgy of Fe-Mn-Al-C is largely different from that of the conventional steels.
Literature has established the basic understanding of this alloy system through experiments and
from theoretical calculations based on the thermodynamics. The current available databases and
softwares could provide reasonable information about the phase constitution in the Fe- Mn-Al-
C system but could not give accurate predictions on phase boundaries and temperature range,
especially when the Al and Mn contents become higher, although the calculations from the
different approaches are similar. A few experimental data show that the 𝛾 region is extended to
the direction of the high Al concentration as Mn is increased. However, the calculations show
that the 𝛾 region expands towards a higher Al concentration with increasing Mn up to 20%, but
shrinks with further increasing Mn.
REVIEW OF THE LITERATURE:
The austenite based low density steels can be placed in the space provided for the second
generation AHSS in the traditional strength vs. elongation ‘‘banana diagram”, very often quite
similar to those of Mn TRIP and Mn TWIP steels. Austenitic Fe-Mn-Al-C steels have a rather
robust response to the minor changes in composition and processing. Austenitic Fe-Mn-Al-C
steels can be produced as hot rolled and cold rolled strips by using the conventional production
lines for HSLA steels. The final conditions can be solution treated or annealed after cold rolling,
or aged. Age hardening is not a necessary step for automotive applications but can be a
consequence of overaging regime found in many continuous annealing lines.

Al brings two main effects in the solution treated austenitic steels: increasing the SFE, and
producing short range orders (SRO) or 𝜅′-carbides. It can be generalised that dislocation planar
gliding is a dominant deformation mechanism in austenitic Fe-Mn-Al-C alloys regardless of
the SFE and aging conditions. In austenitic Fe-Mn-Al-C steels with a lower Al content, TWIP
is still an active deformation mechanism in addition to the dislocation planar glide. Thus, strain
hardening behaviour very similar to high Mn TWIP steels can be obtained. When an increasing
amount of Al is added, the TWIP effect is suppressed due to the increase in the SFE. The
existence of SROs or nano-sized 𝜅′-carbides and its interaction with dislocations play an
important role on planar glide mode deformation of the high Al austenitic steels. Thus, the high
solute elements contribute to the deformation gliding and the high SFE suppresses the TWIP
effect. The precipitation of 𝜅′-carbides and the resulting age-hardening has been obtained in
some alloys, no 𝜅 precipitation and therefore hardening has been reported in some other alloys.
All the results, however, indicate clearly that C and Al contents of the steels play a very
important role in the formation of 𝜅-carbides. It appears that Al > 7% and C > 0.7% are
necessary to form intragranular 𝜅′-precipitates in austenitic Fe-Mn-Al-C alloys. Another form
of precipitation may occur by way of formation of the 𝛼-ferrite by transformation from the 𝛾
matrix. These 𝛼-ferrite particles may transform into the B2 or DO3 intermetallic phases, leading
to significant strengthening.
The tensile properties of Mn TWIP steels generally degrade when a medium level of Al (2–6%)
is added because of the change in the deformation mechanisms. Austenitic Fe-Mn-Al-C alloys
can only exhibit excellent combinations of mechanical properties when the Al content is higher

than 7%, and this is attributed to the formation of SROs or shearable nano-sized 𝜅′-carbides
and their role in the development of planar dislocation substructures. When Al is higher than
10%, 𝛿-ferrite is introduced. Therefore, the effective Al composition for austenitic Fe-Mn-Al-
C alloys for the best combination of mechanical properties is quite narrow and limited between
7 and 10%. Typical compositions for austenitic low density steels should be in the range of 15–
30% Mn, 7–10% Al, and 0.7–1.2% C, with the balance being Fe.
OBJECTIVES:
1. Reviewing and Studying the effect of the composition of Fe, Mn, Al, C and
reinforcement materials on the density and tensile strength of the material primarily.
2. Post processing techniques for leveraging the properties of the alloy.
3. Best fit for the reinforcement material

3. ALLOY FORMATION:
3.1 Overall View:
The matrix phase of low-density steels can be either ferrite, austenite or a mixture of ferrite and
austenite depending on the relative contents of the alloying elements such as Al, C and Mn and
temperature. As summarized, according to the matrix phases at high (hot working)
temperatures, where it is assumed that equilibrium conditions are approached, Fe-Mn-Al-C
lightweight alloys can be classified into four categories: (1) ferritic steels, (2) ferrite based
duplex steels, (3) austenite based duplex steels, and (4) austenitic steels. It is to be noted, the
ferrite produced from the liquid phase is denoted as 𝛼-ferrite, while the ferrite precipitated from
austenite is denoted as 𝛿-ferrite. The 𝛿-ferrite has different elemental composition/partitioning
and intrinsic chemical gradients from the 𝛼 –ferrite although they have the same crystal
structure (BCC).
To date, the Fe–Mn-Al–C alloys reported in literature were prepared as small ingots (~50 kg)
by using standard melting and ingot casting route on laboratory scale. Usually, the ingots are
homogenised at a temperature in the range of 1100–1250 ℃ for 1 to 3 h and then hot rolled to
2–5 mm thickness at a finishing temperature of about 850–1000 ℃. In some cases, slabs are
reheated a few times between rolling passes to avoid cracking during hot rolling. After hot
rolling, the hot rolled strips are cooled to a temperature between 500 and 650 ℃ for 1–5 h to
simulate the coiling procedure for conventional HSLA steels or directly water-cooled, or air-
cooled to room temperature. [1]

A cost-effective ingot metallurgy route comprising air induction melting with a flux cover for
production of clean and sound ingots of Fe-Al alloys containing 7 wt.% has been developed.
Alloy produced by AIMFC(Air induction melting of Fe-Al alloys with Flux Cover) exhibited
superior properties compared with alloy produced by the AIM(Air induction melting) process.
The properties of the alloy produced by AIMFC and VIM(Vaccum Induction Melting) were
comparable. Fe-9 wt.% Al alloy produced by AIM exhibited gas porosity, whereas alloy
produced by AIMFC and VIM were free from gas porosity. However, the alloy produced by
AIMFC still exhibited microcracks. It has been demonstrated that the alloy Fe-9 wt.% Al can
be produced by VIM only. Increasing carbon in Fe-7 wt.% Al alloy resulted in significant
improvement in strength but adversely affected ductility. [2]
For austenite based low density steels, there are many process variants, as shown in Fig.1 . For
hot rolled products, after hot rolling, the strips can be directly fast cooled to a temperature in
the range of 500–750 ℃, then slow cooled or isothermally held, as indicated by process 1 in
Fig.1 Process variants for producing hot rolled and cold rolled austenitic Fe-Mn-Al-C
steel strips. The numbers identify process routes as described in the text. [1]
casda

Fig. 1. Alternatively, the hot rolled strips can be fast cooled to room temperature, followed by
isothermal annealing, as indicated by process 2. The cooling rate after hot rolling should be
high enough to avoid the formation of the intergranular κ carbides. Coiling temperature and
cooling rate during coiling can be aimed at obtaining nanoscale κ precipitates but overaging of
the κ precipitates is to be avoided. [1]
To obtain cold rolled products, for age-hardenable austenitic Fe-Mn-Al-C steels, the cold rolled
strips are solution treated in the temperature range of 900–1100 ℃ in the single austenite phase
area, and then quenched in water, oil or other quench media. Subsequent aging treatments can
be performed to produce precipitation hardening, as indicated by processes 3, 4 and 5 in Fig. 1.
Process 3 has no isothermal holding and processes 4 and 5 are applied to maximize the age
hardening effects of the 𝜅′ phase precipitation. The common practice for age hardening is to
hold isothermally for 5–20 h in the temperature range 500–700 ℃ [1]. For non-age hardenable
austenitic Fe-Mn-Al-C steels, after cold rolling, recovery or recrystallization annealing in the
temperature range of 600 to 900 ℃ for a short time (1–5 min) can be applied to restore the
ductility or to obtain the grain refinement. In this case, the continuous annealing line for
producing conventional steel strips can be used. Tempering at 600 ℃ triggered slight/incipient
transformations: discontinuous precipitation and precipitation transformation. Differently,
tempering at 800 ℃ triggered complete and widespread transformations like cellular
transformation coupled with spinodal decomposition [3]. On one hand discontinuous
precipitation developed γ + κ-carbide fine lamellae, on the other hand precipitation
transformation produced microstructures featuring fine/coarse γ + κ-carbide lamellae and

carbon-depleted γ-bands + κ-carbide-nodules. Cellular transformation developed α + κ lamellar
structures and two different morphologies were detected: wide-inter-spaced thick lath-like κ-
carbide strips within a ferrite matrix and α + κ wide lamellae with high inter-lamellar spacing
that does not affect the improvement of hardness.
3.2 Effects of aluminum content on microstructural modification:
Microstructures of the present hot-rolled lightweight steels are basically composed of ferrite
grains and κ -carbides in a banded shape, but the cracking behavior is affected mainly by the
morphology, volume fraction and distribution κ -carbides. κ -carbide bands are inevitably
formed by the segregation of C, Mn and Al during casting and solidification [4], and their
formation behavior depends on hot-rolling conditions and alloying elements.
In the XRD data in Fig. 2,
Fig.2 XRD results of the three hot-rolled steels. [4]

however, peaks of κ -carbide are lowest in the A4 steel, although the A4 steel shows the highest
volume fraction of the κ -carbide band. According to the equilibrium phase diagram of the A4
steel (Fig. 3a), the fraction of austenite is highest at 900 ℃, and the fraction of κ -carbide is
lowest at the coiling temperature of 650 ℃. In this case, a considerable amount of austenite is
retained because of the low fraction of κ -carbide at 650 ℃, and the retained austenite is
transformed to ferrite, which surrounds the κ -carbide colonies already formed within κ -carbide
bands (Fig. 3d). Thus, in the A4 steel, the volume fraction of the κ -carbide colony is lowest in
spite of the highest volume fraction of κ -carbide band, as shown in Table 1.
In the A6 steel, in contrast, the volume fraction of austenite at 900 ℃ is lowest, and the fraction
of κ -carbide is highest at 650 ℃. Since a considerable number of κ -carbides are formed without
retained austenite, the fraction of ferrite located within κ -carbide bands is lowest among the
the three steels.
Table 1. Volume fractions of ferrite matrix, ferrite and κ -carbide colony existing within κ -
carbide band and grain boundary κ -carbide in the three hot-rolled lightweight steels (unit %).[4]

Fig.3 .High-temperature equilibrium phase distribution calculated by using
ThermoCalc for (a) A4, (b) A5 and (c) A6 steels. [6]

The Al content also affects the morphology of κ -carbide as well as the volume fraction. The
eutectoid reaction temperature of κ -carbide increases from 670 to 780 ℃ with increasing Al
content, as shown in Fig. 3a–c. This increased eutectoid reaction temperature changes the
morphology of κ -carbide. [5] compared the thickness or width of κ -carbide with annealing
temperatures of 500 and 600 ℃ with TEM analysis. They reported that the average widths of κ
-carbide increased about three times because of the nucleation–growth process. In the A6 steel,
κ -carbides are formed at relatively high temperatures, and their diffusion rate is fast. Thus, the
shape of κ -carbides located inside κ -carbide bands becomes long and thick [6], and κ -carbides
located inside the ferrite matrix are formed continuously along the ferrite grain boundaries [6].

3.3 Effects of Manganese content on microstructural modification:
The fractions of phases existing at high temperatures were determined by thermodynamic
calculations in order to identify high-temperature phases of the lightweight steels. ThermoCalc,
Fig. 4—High-temperature equilibrium phase distribution calculated using
ThermoCalc[7] for the (a) M3, (b) M6, (c) M9, and (d) M12 steels. [7]

which is a commercial thermodynamic calculating program, was used for the calculation, and
an upgraded version of TCFE2000 was used for the thermodynamic database. The various
fractions of equilibrium phases in the M3 through M12 steels are shown in Figures 4(a) through
(d). The M3 steel consists of similar amounts of ferrite and austenite at 1473 K (1200 ℃) (Figure
4(a)). The amount of austenite decreases, while that of ferrite increases, as the temperature
decreases. At 993 K (720 ℃), the amount of austenite abruptly decreases, and κ -carbides begin
to precipitate. The fraction of precipitated κ -carbide is about 7 pct at the coiling temperature
of 923 K (650 ℃). According to the present thermodynamic database,[7] other phases, e.g.,
cementite, M5C2, and M7C3, originating from the Fe-Mn-C ternary system and graphite might
appear as stable phases during the equilibrium calculations for these steels, but these phases are
not considered in the calculations.
The M6 steel consists mostly of austenite at 1473 K (1200 ℃), and a small amount (26 pct) of
ferrite (Figure 4(b)). The amount of austenite decreases, while that of ferrite increases, as the
temperature decreases. The fraction of precipitated κ -carbide is about 2 pct at the coiling
temperature of 923 K (650 ℃), which is lower than that in the M3 steel, while the fraction of
austenite remains at about 20 pct.
In the M9 steel, about 80 pct of austenite is found at 1473 K (1200 ℃), and the amount of
austenite smoothly decreases with decreasing temperature, while the amount of ferrite increases
(Figure 4(c)). When the temperature decreases to 883 K (610 ℃), κ -carbides begin to
precipitate, but their amount is relatively small at 0.7 pct. At this temperature, about 36 pct of
austenite is retained.

In the M12 steel, the formation behavior of equilibrium phases of austenite, ferrite, and κ -
carbide is similar to that in the M9 steel, while κ -carbides are not precipitated even at 873 K
(600 ℃) (Figure 4(d)). At the coiling temperature of 923 K (650 ℃), the amount of retained
austenite is about 47 pct, which is highest among the four steels. The phase transformation
temperatures can be estimated from these equilibrium phase diagrams.
Considering that the hot-rolled steel sheet was elongated along the longitudinal direction during
the cold rolling, plate-type sub-sized tensile specimens having a gage length of 12.6 mm, gage
width of 5 mm, and gage thickness of 1 mm were prepared in the same longitudinal direction.
The specimens were obtained from the 1/2 thickness location of the 3-mm-thick hot-rolled sheet
or from the full thickness of the 1-mm-thick cold-rolled sheet. [7]
As a summary, in the duplex steels, the microstructural constituents are more diverse and the
microstructural evolution is more complicated than those in the ferritic or austenitic low density
steels. Because of the co-existence of the 𝛿 phase and the 𝛾 phase, the difference in
recrystallization behaviour between the 𝛿 and the 𝛾 phases, the partially reverse transformation
of 𝛾 to 𝛼 and the element partitioning between the austenite and ferrite during intercritical
annealing, many variants of microstructures can be obtained. The final microstructure of duplex
steels may comprise of a combination of 𝛿, 𝛾, 𝛼, κ martensite (M) and bainite (B) phases or
some of them. The phase constitution, size, volume fraction and the distribution of phases can
be controlled by adjusting the annealing thermal parameters in combination with the preceding
cold rolling and the austempering that follows. One key aspect is to control the mechanical
stability of the austenite through optimizing the steel composition and process.

4. POST ALLOY PREPARATION AND STRENGTHENING MECHANISMS:
Austenitic low density steels contain a higher Mn content, typically between 12 and 30%, Al
up to 12% and C between 0.6 and 2.0%. This type of alloys can have a fully equiaxed-austenitic
microstructure at hot working temperatures and the austenite is metastable after fast cooling.
Some of the strengthening mechanisms that can be incorporated are 1.) Solid solution hardening
2.) Grain Refinement 3.) Precipitation hardening 4.) Work hardening.
4.1 Solid Solution Hardening:
Solid solution hardening plays an important role in the strengthening of Fe-Mn-Al-C steels due
to the high amounts of the alloying elements C, Al and Mn present in these steels. Especially,
the interstitially dissolved C increases the strength of any Fe-Mn-Al-C alloy substantially. In
the solutionized and quenched condition, all the alloying elements Mn, C and Al dissolves in
the FCC austenite. The effect of C on the yield strength of austenite in Fe-Mn-Al-C steels has
been reported to be 187–300 MPa/wt% C (187 MPa/wt% C in Mn TWIP steels) [8]. The
presence of Al in solid solution of austenite slightly increases the yield strength (< 10 MPa/wt%
Al) [9,10].
Four laboratory steel grades with different carbon contents, referred to as steel A, B, C and D
were prepared by vacuum induction melting. After being homogenized at 1443 K (1170 ℃) for
2 hours, ingots of laboratory grades were hot rolled to produce steel sheets of 2.4 to 2.9-mm
thickness. All steel sheets were then cold rolled with a reduction ratio of 55 pct for all samples
(except one type of sample which was cold rolled with a reduction ratio of 85 pct). In order to

study the effects of the austenite grain size, the austenite carbon content, and the ferrite volume
fraction on the tensile properties, the four laboratory steel sheets underwent different solution
heat treatments at high temperature [between 1123 K and 1473 K (850℃ and 1200℃)]. After
this thermal treatment, samples were recrystallized. There were no residual stresses in samples,
except stresses induced by the water quench. Microstructures were first examined by using
optical microscope after polishing with colloidal silica or etching with 2 pct Nital or Villela
solution. From optical observations, samples and heat treatments were selected as reported in
Table 2.
For all samples used for studying the grain size effect or the carbon effect, it should be noted
that no 𝛽-Mn precipitates were detected by EBSD. In addition, no intergranular κ′-carbides
were observed by TEM. However, TEM selected area diffraction patterns in austenite grains
exhibited diffuse superlattice reflections as shown in Figure 5(a). Using intensity profile (Figure
5(b)), it was observed that {100} reflections are stronger than {110} reflections. As pointed out
Table 2. Thermal Treatments for the Three Microstructural Studies: (i) Grain Size Effect, (ii)
Carbon Effect, (iii) Ferrite Effect [11-13]

in References [11] through [13], this indicates the presence of L’12 ordering in austenite, and
thus the presence of intragranular κ -carbides. However, it was not possible to clearly image
these κ -carbides or modulations in dark-field imaging. This supports the idea that the volume
fraction of κ -carbides is probably very low, due to the small size of the carbides and/or to their
low number per unit volume. Finally, it has to be highlighted that the presence of intragranular
κ -carbides after quenching of the samples of steel A (Fe-28.8Mn-9.4Al-0.92C) is not surprising
as the same had already been observed by [14] and [15] respectively, in a Fe-30Mn-7.8Al-1.3C
steels solution treated at 1253 K (980 ℃) and ice water quenched, and in a Fe-28Mn-9Al-0.8C
steel solution treated at 1273 K (1000 ℃) and water quenched.
Fig. 5—(a) Diffraction pattern in [001]c zone axis obtained in an austenite grain of the steel A
solution treated at 1148 K (875 C) for 10 min. Some superlattice reflections are marked with
white circles. (b) Intensity profile plotted along the white dashed arrow. {110} reflections are
weaker than {100} reflections. [14-15]

4.2 Grain Effect
Grain refinement is another strengthening mechanism that can be applied for non-age hardenable
Fe-Mn-Al-C alloys. The austenitic grain size can be refined by a combination of cold rolling and
recrystallization annealing. The effects of austenitic grain size on the tensile properties have been
determined on the basis of the Hall-Petch relationship, using the following values: an initial
strength of 𝜎y0= 288 MPa and Hall-Petch parameter K = 461 MPa𝜇m1/2
for yield strength, and an
initial strength of 𝜎uts0= 742 MPa and K = 351 MPa𝜇m1/2
for ultimate tensile strength.
4.3 Precipitation and Work hardening.
Precipitation hardening is a significant strengthening mechanism in highly alloyed Fe-Mn-Al-C
compositions. The first form of precipitation is that of fine (nano-scaled) and homogeneously
distributed 𝜅′-carbides. The presence of the 𝜅′-carbide has a significant influence on the movement
and arrangement of dislocations during deformation. The second form of precipitation is that of
the 𝛼-ferrite, which can be present as fine stringers in the 𝛾 matrix of the higher Al-alloyed
compositions [1]. The 𝛼-ferrite will transform to B2 or DO3 (intermetallic phases), which may also
lead to strengthening of Fe-Mn-Al-C steels.
In the plastically deformed state, the Al-free Fe-18Mn-0.6C steel possesses a high density of twin
planes where twinned regions exhibit additional microtwins. The Fe–28Mn–9Al–0.8C steel in the
solid solution state was fully austenitic with abundant annealing twins. The stacking fault energy
of the steel was estimated as ∼85 mJ/m2 and its austenite stability was firm with ΔG( 𝛾 → 𝜀)
of

∼1130 J/mol. The fully austenitic steel exhibited an excellent combination of strength and
ductility over 80,000 MPa% with fairly high post-necking elongation of 14% at room temperature
[16]. The strain hardening rate of the steel during tensile deformation continuously increased with
increasing strain before yielding plastic instability. The deformed microstructures of the steel
exhibited the typical planar glide characteristics such as dislocation pile ups on a single slip plane
and formation of a Taylor lattice structure at low strains, crystallographic microband formation at
medium strains, and microband intersections at high strains. Continuous strain hardening causing
an excellent combination of strength and ductility of the steel is to be originated by plasticity
induced by microbands which represent a geometrically necessary structure with high dislocation
density.
5. ANALYSIS OF TENSILE PROPERTIES
5.1 Fe-Mn-Al-C Steels
Depending on the alloy composition and the process route, Fe-Mn-Al-C steels provide a wide
range of tensile properties. The tensile properties of Fe-Mn-Al-C alloys have been intensively
reviewed by Howell et al. [17] and Kim et al. [18]. Plots of total elongation (TE) against ultimate
tensile strength (UTS) have become a key tool in claiming success in alloy design and property
differentiation. There is a large variation in the TE of Fe-Mn-Al-C steels reported in literature as
tensile samples with various geometries were tested. The tensile strength and yield strength are
reasonably independent of sample geometry, but the elongation, especially TE, is dependent on

the gauge length and other geometrical parameters, such as the ratio of the gauge length to parallel
length, and the width/thickness ratio. The tensile properties also depend on the surface condition
of the specimens, especially for hot rolled strips. Fig. 6 is a so-called ‘‘banana diagram” showing
the TE vs. the UTS of Fe–Mn–Al–C lightweight alloys in the solution treated and quenched
condition. The positions of the conventional steels are indicated in dotted areas for comparison.
The tensile properties collected from literature were measured from specimens with tensile
direction parallel to the strip rolling direction. It has been reported that the tensile properties of
duplex and austenitic Fe-Mn-Al-C steels have very low planar anisotropy [1]. As seen, the tensile
properties (UTS, TE) of Fe-Al ferritic alloys are similar to those of conventional low-alloyed C-
Mn and HSLA steels. The ferrite based duplex steels are located on the upper bound of the first
generation advanced high strength steels (1G-AHSS). The single-phase austenitic steels cover the
location of Mn TWIP steels in the banana diagram but with a larger area. The austenite based
duplex steels show a lower TE than the single-phase austenitic steels, while the strength levels are
similar. With the formation of coarse 𝜅-carbides in the austenitic matrix, the alloys display very
poor tensile elongation irrespective of the strength.
Fig. 6. TE as a function
of UTS in Fe-Mn-Al-C
alloys for a uniform
specimen geometry
corresponding to ASTM
E-8 standard (50 mm _
12.5 mm _ 1 mm). [1]
The tensile data were
taken from solution-
treated and water-
quenched strips [1]

The addition of Al helps in decreasing the density but the specific stiffness of the steels goes down
as Al widens the composition range and temperature range of the primary 𝛿-phase during
solidification, which promotes the formation of cavities. A relatively low E value is observed in
the as-cast state, and it is increased by hot rolling. After cold rolling and annealing, E modulus is
slightly decreased again from the as-hot rolled value. This happens due to the casting defects which
decreases the specific stiffness. Optimized approach has exploited the content of Al on hot rolled
steels and below are the observed values of such steels.
Ferritic
steels
Ferrite based duplex
steels
Austenite based duplex
steels
Austenitic steels
Al 5–9% Al 3–7% Al 5–10% Al 5–12%
Mn <5% Mn 2–12% Mn 5–30% Mn 12–30%
C<0.05% C 0.05–0.5% C 0.4–0.7% C 0.6–2.0%
UTS 200–
600 MPa
UTS 400–900 MPa UTS 800–1300 MPa UTS 800–1500 MPa
TE
10–40%
TE 10–40% TE 10–40% TE 30–80%
Table 3. Tensile properties of steels in hot rolled condition. [1]

5.2 Addition of other elements in the base alloy.
Regression analysis of the influence of composition on the elastic modulus of austenitic steels,
Nitrogen as an interstitial solute in austenite deforms the lattice and increases the elastic modulus.
Probable cause of the strong effect might be clustering to Cr dissolved in the lattice. It has been
demonstrated in [19] that the elastic modulus increases with increasing N-content while the
composition otherwise was held virtually constant, indicating a coefficient of 34 GPa/N mass
content in % in austenite at RT. This result is reproduced within the large standard deviation in
our regression analysis. The small amount of data on N-containing steels in the study (4 batches,
3 batches from source [19]) causes this deviation. The effect of Cr and Ni on the elastic modulus
of austenite has been qualitatively known. NbC has a reported elastic modulus of 350-500 GPa
and TiC of 460 GPa. Elastic modulus data for Cr23C6 was not available. The effect of precipitates
like carbides on the elastic modulus must be related to the modulus of the precipitate, so this
regression result seems to be reasonable. Conclusions from the elastic modulus data of pure
elements to the effect of that element when in solid solution with iron cannot be drawn.
Other possible elements which would help in increasing the stiffness are listed but every element
has its’ own effects and may not show the same results at different conditions. Some compounds
like NbC and TiC have contributed to the elastic modulus. For austenitic steels, the phase
compositions and the elemental compositions of the phases were computed in a thermodynamic
analysis;

 For pure 𝛾-iron at room temperature, the elastic modulus of 195 GPa is found to be lower
than that of pure 𝛼-iron,
 N, Cr, NbC, and TiC increase and - Ni and Mo reduce the elastic modulus of austenitic
steels;
 Cr23C6 appears with the same regression coefficient as y-iron in the study, thus having
no effect on the elastic modulus of austenitic steels. [19]

Listed are the compounds that would increase the stiffness with their Young’s modulus but not
all are feasible in many cases.
Material Density(kg/m
3
) Elastic Modulus(N/m
2
) Melting point (˚C)
AlN (hex) 3.05 315 2200
Al2
O3
(hex) 3.99 400 2063
B4
C (rhomb) 2.52 450 2450
BN (hex) 3.48 720 2300
NbC (cub-B1) 7.6 580 3500
SiC (hex) 3.2 480 2200
Si3
N4
(hex) 3.44 210 1750
TiC (hex) 4.93 400 3067
TiB2
(hex) 4.50 370 2900
ZrC (fcc) 6.5 380 3535
Table 4. List of compounds from various sources. (Red)-Feasible for reinforcement in the base alloy.
Various source

6. Results And Discussion.
 A- (2𝜃), B-Intensity, Sample-(Fe-1.5C-9Mn-30Al) expected.
Expected phases: 𝛾−(𝑎𝑢𝑠𝑡𝑒𝑛𝑖𝑡𝑒) and 𝜅′𝑐𝑎𝑟𝑏𝑖𝑑𝑒.
Observed phase: Fe with HCP Crystal structure according to JCPDS database with
Reference codes: 98-063-1723, 98-063-1726 for the significant peaks as visible in the
Graph.
Fig.7 - 2𝜃 Vs Intensity for ( Fe-1.5C-9Al-30Mn ), Water Quenched to RT from (900-1100 ℃),
aged (at 550 ℃ for 24 h in air atmosphere).

 Addition of Al more than 7% gives rise to brittle nature of alloy. The 𝛿- ferrite formed
increases the tensile properties but decreases the ductility of the alloy. [1]
 In order to prevent cracking occurring during cold rolling, it is necessary to avoid the
lengthening or thickening of lamellar κ-carbides. It is recommended that the steels are
rapidly cooled from the finish rolling temperature to coiling temperature through the
formation temperature of κ-carbides. As practical methods for preventing the cracking,
an Al content <6% or the appropriate control of C and Mn contents is suggested to
decrease the formation temperature of κ-carbides. [6]
 Fe-Mn-Al-C Austenitic steels strengthened by nanoscale ordered 𝜅-carbides exhibit high
strength (yield stress 0.8–1.2 GPa) and good ductility (elongation: 30–50%). Dislocation–
𝜅 -carbide interactions analysed by weak beam TEM reveal that at the current
deformation conditions and 𝜅 -carbide morphology, size and distribution, the
predominant deformation mechanisms are Orowan bypassing of rods of 𝜅 -carbides and
subsequent expansion of dislocation loops assisted by cross-slip and, to a less extent,
shearing of 𝜅 -carbides.[17]
 Solution-treated alloys with sufficient C and Mn have a fully austenitic microstructure.
Aging produces equilibrium phase constituents of austenite, ferrite, κ-carbide and β-Mn.
κ-carbide, an E21 perovskite precipitate, is responsible for strength increases during
aging. [19]
 Aging below 650°C precipitates κ-carbide in a homogeneous fashion throughout
austenite, and over-aging produces heterogeneous grain boundary phases deleterious to
mechanical properties. Mechanical properties range from a solution-treated 350 MPa

yield stress, with elongation to failure as high as 77%, to alloys with tensile strengths
greater than 2 GPa. [19]
 Fe-Mn-AI-C alloys can also be engineered to have a higher work hardening rate than
AISI 304 stainless steel. The Fe-Mn-AI-C alloys are susceptible to stress corrosion
cracking and hydrogen embrittlement. Alloy additions of silicon and refractory elements
are beneficial to both tensile and fatigue performance. The addition of silicon has been
shown to prevent the formation of β-Mn during age hardening, improving age hardening
behavior and strength, and increased castability over non-silicon-alloyed Fe-Mn-AI-C
steels. [19]
7. Conclusion
The lightweight Fe-Mn-Al-C steels are a new type of alloy systems. Due to the addition of high
amounts of Al, together with Mn and C, the physical metallurgy, general processing,
microstructural evolutions and deformation mechanisms of these steels are largely different from
those of the conventional steels. Aluminium plays the central role in reducing the density of
steels; a 1.3% reduction in density is obtained per 1 wt% Al addition. Al also brings a large
change in the physical metallurgy and deformation mechanisms of steels. Aluminium is a ferrite
stabiliser and introduces a disordered (BCC_A2) to ordered (BCC_B2) transition when the Al
content is higher than 6%. To make phase transformation feasible in the low density steels, the
austenite stabilisers Mn and C are needed. Al also brings other two main effects in the solution
treated austenitic steels: increasing the SFE, producing SRO or 𝜅′-carbides.

The Fe-Mn-Al-C low density alloys can achieve a wide range of tensile properties. The austenite
based low density steels can be placed in the space provided for the second generation AHSS in
the traditional strength vs. elongation ‘‘banana diagram”, very often quite similar to those of the
Mn-TRIP and the Mn-TWIP steels. The ferrite based duplex low density steels can be located at
the upper bound of the first generation AHSS in the ‘‘banana diagram”, while the ferritic Fe-Al
steels show behaviour typical of HSLA steels of 400–500 MPa level. Based on the current
understanding of the low density steels, austenite based Fe-Mn-Al-C alloys are the most
promising due to their potentials to have superior strain hardening, high energy absorption, and
the robust response to minor changes in composition and processing. The toughness of the
solution treated austenitic Fe-Mn-Al-C steels is slightly lower than that of Cr-Ni stainless steels
but is higher than that of the conventional high strength steels. The dynamic fracture toughness
is equivalent to or higher than that of 4130 steels, indicating an improved crashworthiness. The
ferrite based duplex low-density steels is another promising alternative; a bimodal microstructure
can be obtained through process control for steels with lower alloying contents, in which the
deformation of the ferrite and the TRIP/TWIP effects from the retained austenite can be
advantageously used.

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[16] Yoo, J. D., & Park, K.-T. (2008). Microband-induced plasticity in a high Mn–Al–C light
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