# Ordered B2 phase in a disordered FCC matrix: Current state of research on the development of low-density B2-strengthened Fe-Mn-Al-Ni-C steels

Mainak Saha

Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Madras, Chennai-600036, India

*Currently at Nanostructure Analysis Group, Research Centre for magnetic and Spintronic Materials, National Institute for Materials Science, Sengen, Tsukuba, Ibaraki: 305-0047, Japan

Email address: mainaksaha1995@gmail.com

ORCID ID: 0000-0001-8979-457X

## Abstract

Over the last five years, the automotive industry has primarily focussed on utilising a number of extraordinary properties of B2-strengthened Fe-Mn-Al-Ni-C steels, especially high specific strength, and high strain hardening. While reduction in density may be attributed to the addition of Al, high toughness and strain hardening may be ascribed to the precipitation of non-shearable B2 precipitates which suppresses the formation of shearable L'₁₂-κ-carbides ((Fe, Mn)₃AlC carbides). Considering the growing interest in these steels, the present review is aimed at highlighting the present state of research and associated challenges (both fundamental and industrial) with a special emphasis on the influence of B2 precipitate distribution and morphology on the strengthening effect in these steels. In addition, the necessity of obtaining both structural and chemical information about ordered B2 precipitates using the newly emerged methodology of correlative microscopy has also been highlighted.

**Keywords:** Ordered precipitate, B2 morphology, thermomechanical processing, strain hardening, correlative microstructural characterisation.

## 1. Introduction

Until 2015, lightweight Fe-Mn-Al-C steels (also known as TRIPLEX steels) with high specific strength (density ~18% lower than that of conventional steels and strength ~1 GPa) were the main centres of attraction in the community of automotive-grade low-density steels (Park 2013; Raabe et al. 2014b; Kim et al. 2015; Yao et al. 2016a; Park et al. 2019a). Their microstructures comprised of α-Fe(Al, Mn) based BCC phase and finely dispersed L'₁₂-κ-carbides ((Fe,Mn)₃AlCₓ (x ≤ 1) carbides) in FCC-austenite matrix (Yao et al. 2016b; Wang et al. 2020; Gutierrez-Urrutia 2021). However, the primary limitation of these steels was the low strain hardening owing to the precipitation of shearable κ-carbides. Besides, as reported in Ref. (Kim
et al. 2015), the specific strength of these steels was not comparable to those of Al and Ti
alloys. These limitations (of TRIPLEX steels) were addressed by Kim et al. (Kim et al. 2015),
wherein addition of ~5wt.% Ni to TRIPLEX steels was shown to (i) lead to the formation of
non-shearable B2-(Ni,Fe)Al based precipitates and (ii) suppress the precipitation of κ-carbides
(Frommeyer and Brüx 2006; Springer and Raabe 2012a; Springer and Raabe 2012b; Seol et al.
2013a; Sohn et al. 2013; Kim et al. 2015; Rahnama et al. 2017c; Rahnama and Sridhar 2019;
Wang et al. 2020; Mozumder et al. 2021). This is in contrast to the conventional wisdom of
FCC stabilisation by Ni (Kim et al. 2015). Besides, addition of C and Mn has also been reported
to stabilise the FCC matrix in Ni-added TRIPLEX steels (or, Fe-Mn-Al-Ni-C steels) (Kim et
al. 2015). In the context of B2-strengthened (B2s) Fe-Mn-Al-Ni-C steels (B2-strengthened Fe-
Mn-Al-Ni-C steels will be referred to as B2s Fe-Mn-Al-Ni-C steels, hereafter, in the
manuscript), FCC phase provides ductility whereas B2-(Ni,Fe)Al based precipitates account
for high strength (Kim et al. 2015; Rahnama et al. 2018a; Mozumder et al. 2020b; Gao et al.
2020). In other words, the high overall toughness (total elongation ~20-35% and strength ~1.6
GPa) in B2s Fe-Mn-Al-Ni-C steels may be attributed to an outstanding combination of strength
and ductility combination provided by the FCC matrix and B2 precipitates respectively (Kim
et al. 2015; Kim 2019; Zargaran et al. 2021, Saha 2022, 2023).

B2 precipitate morphology has been reported to influence strengthening effect in B2s Fe-Mn-Al-Ni-C steels (Park et al. 2019a; Zargaran et al. 2021). In addition, distribution of ordered B2 phase in a disordered FCC matrix hugely influences the strengthening effect in these steels (Kim et al. 2015; Rahnama et al. 2018a; Kim 2019; Park et al. 2019a). For instance, one of the most common ways of enhancing strength-ductility combination in B2s Fe-Mn-Al-Ni-C steels is through an appropriate utilisation of potent nucleation sites such as FCC grain boundaries (GBs), grain edges and shear bands (formed due to plastic deformation) (Kim et al. 2015; Rahnama and Sridhar 2019; Hwang et al. 2020). At present, correlative microscopy has emerged as a powerful methodology in terms of providing both structural and chemical information from the same region in a given microstructure (Raabe et al. 2014a). However, the aforementioned methodology has not been utilised (till date) to understand the microstructure evolution of B2 precipitates in FCC matrix (in B2s Fe-Mn-Al-Ni-C steels). To this end, the present review article is aimed at highlighting some of the developments (till date) and associated challenges (both fundamental and industrial) in the avenue of B2s Fe-Mn-Al-Ni-C lightweight steels with a special emphasis on the influence of B2 precipitate morphology, and distribution on the deformation characteristics of these steels. In addition, the necessity of
utilising correlative microscopy methodology towards understanding the microstructure evolution of ordered B2 precipitates in a disordered FCC matrix (in B2s Fe-Mn-Al-Ni-C steels) has also been highlighted.

## 2. B2s Fe-Mn-Al-Ni-C steels: Microstructure

Ni and Fe are the most common B2 forming elements with Al owing to their highly negative enthalpy of mixing with Al (Noebe et al. 1993; Noebe et al. 1996; Takeuchi and Inoue 2000). Ni has been widely reported as stabilise FCC phase in steels, similar to Mn and C (Rana 2014a; Rana 2014b; Rana et al. 2014; Kim et al. 2015; Chen et al. 2017a). However, the first report of B2-(Ni,Fe)Al phase stabilisation by addition of ~5wt.% Ni to Fe-16Mn-9Al-Ni (wt.%) TRIPLEX steels (with a microstructure comprising of FCC and κ phases) was first obtained from the work of Kim et al. (Kim et al. 2015). While Al accounts for low density (~5.5-6.5 g/cc), the room temperature microstructure of B2s Fe-Mn-Al-Ni-C steels typically comprises of B2 precipitates in FCC matrix. B2 phase is rich in Fe, Al and Ni (Kim et al. 2015). On the other hand, C and Mn have been reported to undergo preferential partitioning to FCC matrix (Kim et al. 2015). In addition to the above, a number of literatures also report the presence of (i) Ni and Al-rich BCC phase in FCC matrix (Rahnama et al. 2017d; Rahnama et al. 2018b; Mozumder et al. 2018; Mozumder et al. 2020a; Mozumder et al. 2020b; Mozumder et al. 2021), (ii) B2-(Ni,Fe)Al precipitates in BCC phase (Rahnama et al. 2018b; Mozumder et al. 2018; Rahnama and Sridhar 2019; Mozumder et al. 2020a; Mozumder et al. 2020b; Mozumder et al. 2021), and (iii) L'₁₂ structured κ-carbides in BCC phase and FCC matrix (Rahnama et al. 2017a; Rahnama et al. 2017d; Rahnama et al. 2017c; Rahnama et al. 2018b; Wang et al. 2020; Zargaran et al. 2021). However, there is still a limited understanding on the correlation between overall alloy composition and phase evolution in these steels.

B2 phase is non-shearable by gliding dislocations unlike the κ-carbides present in TRIPLEX steels (Gutierrez-Urrutia and Raabe 2012; Gutierrez-Urrutia and Raabe 2013; Gutierrez-Urrutia and Raabe 2014; Kim et al. 2015; Gutierrez-Urrutia 2021). An optimisation of B2 phase morphology, and distribution through appropriate thermomechanical processing leads to high strain hardening which is responsible for an excellent combination of high specific ultimate tensile strength (SUTS ~200-230 MPa g⁻¹cm³) and high specific yield strength (SYS ~150-200 MPa g⁻¹cm³) in B2s Fe-Mn-Al-Ni-C steels (Kim et al. 2015; Hwang et al. 2020). For instance, cold rolling of Fe-16Mn-9Al-Ni-0.86C (wt.%) steel leads to the formation of B2 phase in the form of stringers whereas annealing at 900 °C (after cold deformation) leads to the nucleation
of B2 phase with three possible morphologies: (i) stringer bands, (ii) fine precipitates (size ~200-1000 nm), and (iii) finer precipitates (size ~50-300 nm) (Kim et al. 2015). In the context of B2s Fe-16Mn-9Al-5Ni-0.9C (wt.%) steel, the annealed microstructure comprising of B2 stringers in FCC matrix has been described as a Type 1 microstructure (Fig. 1) (Kim et al. 2015). Nucleation of B2 phase in the form of precipitates (size ~200-1000 nm) at recrystallised FCC grain edges and GBs leads to a Type 2 microstructure (Fig. 1) (Kim et al. 2015). On the other hand, precipitation of B2 (size ~50-300 nm) at shear bands in deformed (or non-recrystallised) FCC grains leads to a Type 3 microstructure (Fig. 1) (Kim et al. 2015). While, Type 2 and Type 3 microstructures enhance the strength, Type 1 microstructure is detrimental for the same (Rahnama et al. 2018b; Hwang et al. 2020).

As already discussed in the previous section, κ-carbides have been reported to show a stoichiometry of (Fe,Mn)₃AlCₓ type (x ≤ 1) (Yao et al. 2016b). It is interesting to note that in the context of off-stoichiometric κ-carbides (x<1), C vacancies were shown to lead to low formation energies of MnAlγ anti-sites and these anti-sites were shown to result in the off-stoichiometry of κ-carbides (Yao et al. 2016b). This was contradictory to the previous report where C depletion and elastic coherency strains were identified as the primary reasons for Al depletion in Al-alloyed Fe-Mn-C steels (Dey et al. 2017). One of the other reports by Seol et al. (Seol et al. 2013b) suggests that the off-stoichiometry of κ-carbides may be explained in terms of C deficiency at body-centered positions and partial replacement of Al by Fe or Mn in the corner positions of L1₂'-structured κ-carbide unit cell.
<figure><img src="image_1.png" /><figcaption>**Fig. 1** Schematic of different microstructures classified on the basis of different morphologies of B2 phase in B2s Fe-Mn-Al-Ni-C lightweight steels (based on the proposition of Kim et al. (Kim et al. 2015)).</figcaption></figure>

**3. B2s Fe-Mn-Al-Ni-C steels: Micromechanical properties**

<figure><img src="image_2.png" /></figure>
**Fig. 2** Fe-15Mn-9Al-5Ni-0.8C (wt.%) steel: Electron Backscatter Diffraction (EBSD) based phase distribution maps in (a1) as-quenched condition, and (a2-a4) after solution treatment at 1250 °C for 35 mins followed by subsequent heat treatment at 1050 °C for 5 mins with heating rates: (a2) 2.5 °C/s, (a3) 5 °C/s. and (a4) 7.5 °C/s, and (b) Transmission Electron Microscopy-Bright Field (TEM-BF) image showing a triangular-shaped B2 precipitate (designated as B2γ) in FCC matrix for a heating rate of 2.5 °C/s. In parts (a1-a4), the red and blue phases (in EBSD phase distribution maps) correspond to B2 and FCC phases respectively (Rahnama et al. 2017d).

In the context of Fe-15Mn-9Al-5Ni-0.8C (wt.%) steel after solution treatment at 1250 °C for 35 mins followed by subsequent annealing at 1050 °C for 5 mins, heating rate has been reported to influence the morphology and distribution of B2 precipitates in FCC matrix (Rahnama et al. 2017d). It was shown that a low heating rate (~2.5 °C/s) leads to (i) an increase in the volume fraction, and (ii) a uniform distribution of B2 precipitates in FCC matrix and GBs (in Fe-15Mn-9Al-5Ni-0.8C (wt.%) steel) (Fig. 2(a1-a4)) (Rahnama et al. 2017d). On the contrary, high heating rate (~7.5 °C/s) was reported to reduce the volume fraction of B2 precipitates in FCC matrix and GBs (as shown in Fig. 2(a1-a4)). Fig. 2(b) shows the Transmission Electron Microscope-Bright Field (TEM-BF) image of triangular B2 precipitate (with flat interfaces and an average size of ~1.4 µm) in FCC matrix (Rahnama et al. 2017d). The precipitate morphology of B2 phase may be explained in terms of the competing energetic contributions of (i) strain energy, and (ii) interfacial energy during the growth (of the precipitate). For instance, the point at which the interfacial energy contribution predominates over strain energy contribution, is largely system-dependent (Park 2013; Zargaran et al. 2014; Rahnama et al. 2017b; Rahnama et al. 2017c; Zambrano 2018; Park et al. 2019b). This is highly essential for designing microstructures where precipitate morphology largely influences the mechanical properties of the material (Rahnama et al. 2017d; Rahnama et al. 2017c; Rahnama and Sridhar 2019; Kies et al. 2021).

<table><caption>**Table. 1** Fe-16Mn-8Al-5Ni-1C and Fe-15Mn-10Al-5Ni-0.8C (wt.%) steels: Tensile properties along rolling direction (RD) and transverse direction (TD) (The numerical values in brackets denote the average size for different morphologies of B2 phase) (Park et al. 2019a).</caption><thead><tr><td>Steel composition (wt.%)</td><td>B2 morphology</td><td>Direction</td><td>Yield strength (MPa)</td><td>Ultimate tensile strength (MPa)</td><td>Total elongation (%)</td></tr></thead><tbody><tr><td>Fe-16Mn-8Al-5Ni-1C</td><td></td><td>RD</td><td>971 ± 14</td><td>1215 ± 4</td><td>35.9 ± 0.1</td></tr></tbody></table>
<table><tr><td></td><td>Polygonal precipitates (~0.35 µm)</td><td>TD</td><td>947 ± 2</td><td>1231 ± 11</td><td>35.5 ± 0.6</td></tr><tr><td rowspan="2">Fe-15Mn-10Al-5Ni-0.8C</td><td rowspan="2">Polygonal precipitates (~1.9 µm), Stringer (~ 1.4 µm)</td><td>RD</td><td>965 ± 2</td><td>1278 ± 4</td><td>31.7 ± 0.3</td></tr><tr><td>TD</td><td>1001 ± 6</td><td>1234 ± 4</td><td>25.2 ± 1.3</td></tr></table>

<figure><img src="image_3.png" /><figcaption>**Fig. 3** Fe-16Mn-8Al-5Ni-1C (wt.%) steel: Scanning Electron Microscope (SEM) images showing B2 phase nucleated in the form of (a) fine precipitates and stringers in FCC matrix, Fe-15Mn-10Al-5Ni-0.8C (wt.%) steel: (b) coarse B2 precipitates and bands (stringers), Band-free Fe-16Mn-8Al-5Ni-1C (wt.%) steel (c) EBSD-based phase distribution map, and (d) TEM-BF image showing B2 precipitates formed at FCC GBs (highlighted using white arrows in part (c)) and within FCC grains. In parts (a) and (b), dark and bright areas correspond to FCC and B2 phases respectively. In part (d), the bright and dark regions correspond to FCC and B2 phases respectively (Park et al. 2019a).</figcaption></figure>
In the context of B2-strengthened Fe-15Mn-10Al-5Ni-0.8C and Fe-16Mn-8Al-5Ni-1C (wt.%) steels, Park et al. (Park et al. 2019a) have reported that B2 stringers exhibit texture hardening which, in turn, leads to a strong anisotropy in terms of both tensile properties (along rolling and transverse directions) and formability of the steels (Table. 1). It was highlighted that a significant improvement in the formability of these steels may be achieved (without the loss of tensile strength) by the replacement of B2 stringers with uniformly dispersed fine polygonal precipitates of B2 at FCC GBs and grain interiors (Park et al. 2019a). Fig. 3 shows the (i) presence of polygonal B2 precipitates in FCC matrix and GBs and that of (ii) B2 stringers in FCC matrix (Park et al. 2019a). Fig. 3(a) shows a homogeneous distribution of fine polygonal B2 precipitates in Fe-16Mn-8Al-1C-5Ni (wt.%) steel (Park et al. 2019a). On the other hand, SEM image of Fe-15Mn-10Al-0.8C-5Ni (wt.%) steel (Fig. 3(b)) shows a microstructure comprising of comparatively coarser B2 precipitates and B2 stringers (bands) (Park et al. 2019a). Fig. 3(c) shows an EBSD-based GB map overlaid on image quality (IQ) map of band-free Fe-16Mn-8Al-1C-5Ni (wt.%) steel. B2 precipitates of ~350 nm size are observed at FCC GBs (Fig. 3(c)) (Park et al. 2019a). On the other hand, extremely fine B2 precipitates of average size ~60 nm are observed at FCC grain interior, as shown in the TEM-BF image of Fe-16Mn-8Al-1C-5Ni (wt.%) steel (Fig. 3(d)) (Park et al. 2019a). Table. 2 shows the contribution of mechanisms such as solid solution strengthening, GB strengthening, and precipitation strengthening (due to the precipitation of B2 phase) of FCC phase on the overall yield strength of Fe-15Mn-10Al-5Ni-0.8C and Fe-16Mn-8Al-5Ni-1C (wt.%) steels (Park et al. 2019a). Besides, the influence of B2 morphology on the strengthening contribution of B2 phase has also been highlighted in Table. 2 (Park et al. 2019a). Comparing the overall yield strength of Fe-15Mn-10Al-5Ni-0.8C and Fe-16Mn-8Al-5Ni-1C (wt.%) steels with the morphologies of B2 phase in the aforementioned steel compositions, it is evident that B2 morphology largely influences the yield strength of B2s Fe-Mn-Al-Ni-C steels. The overall yield strength of these steels may be expressed using the rule of mixtures, as highlighted in Refs. (Yang et al. 2016; Park et al. 2019a):

$$ \sigma_{YS} = (\sigma_{FCC-SS/GB} + \sigma_{FCC-ppt})f_{FCC} + f_{B2-ppt}\sigma_{B2-ppt} + f_{B2-stringer}\sigma_{B2-stringer} \quad (1) $$

where, $f_{FCC}$, $f_{B2-ppt}$, and $f_{B2-stringer}$ are the volume fractions of FCC matrix, B2 precipitate, and B2 stringer bands respectively (Park et al. 2019a). $\sigma_{FCC-SS/GB}$ is the strength of FCC phase due to solid solution and GB strengthening and $\sigma_{FCC-ppt}$ denotes the magnitude of precipitation strengthening in FCC phase (Park et al. 2019a). $\sigma_{B2-ppt}$ and $\sigma_{B2-stringer}$ denote the strength
of B2 precipitates and stringers respectively (Park et al. 2019a). Dependence of $\sigma_{B2-ppt}$ on the size of B2 precipitates is obtained from Refs. (Crimp and Vedula 1991; Crimp and Vedula 1993; Baker 1995). In equation (1), $\sigma_{FCC-SS/GB}$ is calculated using the relation proposed by Choi et al. (Choi et al. 2015):

$$ \sigma_{FCC-SS/GB} = \sigma_0 + \Sigma \sigma_i X_i + k_{HP} d^{-0.5} \quad (2) $$

where, $\sigma_0$ is the lattice friction strength, $\sigma_i$ denotes the strengthening contribution of individual alloying elements towards solid solution strengthening of FCC phase. $X_i$ is the concentration (in wt.%) of individual alloying elements, $k_{HP}$ is the Hall-Petch coefficient, and $d$ is the average grain size of FCC phase (Choi et al. 2015). In equation (2), $\Sigma \sigma_i X_i$ accounts for solid-solution strengthening whereas $k_{HP} d^{-0.5}$ accounts for GB strengthening of FCC phase (Choi et al. 2015). In equation (1), $\sigma_{FCC-ppt}$ (for non-shearable B2 precipitates) is determined on the basis of Orowan strengthening mechanism, using the following relation: (Gladman 1999; Park et al. 2019a)

$$ \sigma_{FCC-ppt} = \frac{0.538 Gbf^{0.5}}{D} \ln\left(\frac{D}{2b}\right) \quad (3) $$

where, G and b are the shear modulus and Burgers vector of FCC phase respectively (Gladman 1999). D and f are the average size and volume fraction of B2 precipitates respectively (Gladman 1999). In Table. 2, for banded Fe-15Mn-10Al-5Ni-0.8C (wt.%) steel, due to the presence of coarse B2 precipitates (very large value of D in equation (3)) with very low fraction ($f_{B2}$) in FCC matrix, as illustrated in Ref. (Park et al. 2019a), the contribution of $\sigma_{FCC-ppt}$ towards strengthening in the banded Fe-15Mn-10Al-5Ni-0.8C (wt.%) steel is considered to be negligible (in Table. 2) unlike that in band-free Fe-16Mn-8Al-5Ni-1C (wt.%) steel where a large fraction of fine B2 precipitates were shown to be present in FCC matrix (Park et al. 2019a). Gao et al. (Gao et al. 2020) have reported that B2 phase plays a vital role in influencing the strength and ductility of Fe-20Mn-8Al-5Ni-0.8C (wt.%) steel. It was shown that strain partitioning between FCC matrix and B2 precipitates at the initial stages of tensile deformation is associated with Luders Band formation (Gao et al. 2020). In addition, it was highlighted that the dislocation density in B2 phase increases with progress of tensile deformation (Gao et al. 2020). Hwang et al. (Hwang et al. 2020) have devised two alloy design criterion for optimising strength-ductility balance in B2s Fe-Mn-Al-Ni-C lightweight steels. Firstly, Al content should be increased in order to increase the volume fraction of B2 phase. The basis for this criterion is obtained from the Park et al. (Park et al. 2019a) and Wang et al. (Wang et al. 2017) wherein
B2 stringer-free Fe-16Mn-8Al-5Ni-1C (wt.%) lightweight steel has been reported to show a smaller fraction of B2 particles than that of B2 stringer-containing steels. Secondly, the content of FCC stabilizers, especially Mn, should be adjusted for the purpose of avoiding the formation of B2 stringers. Based on the two aforementioned criterion, a new alloy composition, viz. Fe-21Mn-10Al-5Ni-1C (wt.%) was devised (Hwang et al. 2020). Ageing (at 550 °C for 10 mins) after annealing (at 900 °C for 2 and 15 mins) was reported to induce strengthening through enhanced precipitation of B2 phase (at FCC grains and GBs) in Fe-21Mn-10Al-5Ni-1C (wt.%) steel (Hwang et al. 2020). As observed in **Fig. 4** (obtained from Ref. (Hwang et al. 2020)), the FCC and B2 phases (in Fe-21Mn-10Al-5Ni-1C (wt.%) steel) show both Kurdjumov-Sachs (KS) and Nishiyama-Wassermann (NW) orientation relationships with each other. **Fig. 5** shows the TEM images of deformed microstructures of FCC and B2 phases after straining at two different levels, viz. 3% and 10% in Fe-21Mn-10Al-5Ni-1C (wt.%) steel (Hwang et al. 2020). It may be observed that the microstructure of 3% deformed specimen (after annealing at 900 °C for 2 and 15 mins) shows the presence of planar slip bands inside FCC grains (**Fig. 5(a and e)**) (Hwang et al. 2020). The black arrows in **Fig. 5(c and g)** show that deformation of B2 precipitates (in as-annealed condition) mostly occurs near FCC/B2 interfaces with a low dislocation density inside B2 (Hwang et al. 2020). On increasing the strain to 10%, there is a decrease in spacing between planar slip bands within FCC grains (**Fig. 5(b and f)**) with an increase in dislocation density inside B2 precipitates (**Fig. 5(d and h)**) (Hwang et al. 2020). **Fig. 5(i-l)** highlights the same for FCC and B2 phases in aged specimen (aged at 550 °C for 10 mins after annealing at 900 °C for 15 mins) (Hwang et al. 2020).

**Table. 2** Fe-16Mn-8Al-5Ni-1C and Fe-15Mn-10Al-5Ni-0.8C (wt.%) steels: Contribution of FCC and B2 phases to the overall yield strength ($\sigma_{YS}$ is the overall yield strength, $f_i$ and $\sigma_i$ (i: FCC, B2) are the volume fraction and strength of constituent phases, $\sigma_{FCC-SS/GB}$ is the strength of FCC phase due to solid solution and GB strengthening, and $\sigma_{FCC-ppt}$ denotes the magnitude of precipitation strengthening in FCC phase) (Park et al. 2019a). Strength values have been expressed in MPa.

<table><thead><tr><th rowspan="2">Steel<br>composition<br>(wt.%)</th><th colspan="2">FCC matrix</th><th colspan="4">B2 precipitates</th><th colspan="3">B2 stringers</th><th rowspan="2">σ<sub>YS</sub><br>(MPa)</th></tr><tr><th>σ<sub>FCC-SS/GB</sub><br>(MPa)</th><th>σ<sub>FCC-ppt</sub><br>(MPa)</th><th>f<sub>FCC</sub></th><th>f<sub>FCC</sub>·σ<sub>FCC</sub></th><th>σ<sub>B2</sub><br>(MPa)</th><th>f<sub>B2</sub></th><th>f<sub>B2</sub>·σ<sub>B2</sub></th><th>σ<sub>B2</sub><br>(MPa)</th><th>f<sub>B2</sub></th><th>f<sub>B2</sub>·σ<sub>B2</sub></th></tr></thead><tbody></tbody></table>
<table><tr><td>Fe-16Mn-8Al-5Ni-1C</td><td>747</td><td>190</td><td>0.97</td><td>909</td><td>2373</td><td>0.03</td><td>71</td><td>-</td><td>-</td><td>-</td><td>980</td></tr><tr><td>Fe-15Mn-10Al-5Ni-0.8C</td><td>761</td><td>Negligible</td><td>0.72</td><td>548</td><td>1404</td><td>0.18</td><td>253</td><td>1525</td><td>0.10</td><td>153</td><td>954</td></tr></table>

<figure><img src="image_4.png" /><figcaption>**Fig. 4** Fe-21Mn-10Al-5Ni-1C (wt.%) steel: SEM images of (a) 900°C, 2 min and (b) 900°C, 15 min annealed samples (yellow arrows in part (b) represent B2 precipitates within FCC matrix), (c) TEM-BF image and corresponding Selected Area Diffraction Patterns (SADPs) of recrystallized FCC matrix and B2 precipitates (SADPs highlight the Kurdjumov-Sachs (K-S) and Nishiyama-Wassermann (N-W) orientation relationships between FCC and B2 phases) in 900°C, 2 min sample, and (d) TEM-BF, SADPs and corresponding dark-field (DF) images of B2 precipitates within non-recrystallised FCC matrix in 900°C, 2 min sample. In parts (a) and (b), B2 appears as bright particles in a relatively darker FCC matrix. In parts (c) and (d1-d3), the SADP spots in FCC and B2 phases are represented using dashed and solid lines respectively. In part (d), some of the B2 precipitates formed along and between shear bands, have been marked using red and blue lines respectively (Hwang et al. 2020).</figcaption></figure>
<figure><img src="image_5.png" /><figcaption>**Fig. 5** TEM images of deformed microstructures of (a, b, e, f, i, j) FCC matrix and (c, d, g, h, k, l) B2 phase (in Fe-21Mn-10Al-5Ni-1C (wt.%) steel) after straining at 3% and 10% (ε represents the magnitude of strain): (a-h) annealed, and (i-l) aged conditions. In parts (a), (b), (e), (f), (i) and (j), dashed lines indicate the slip traces in FCC phase. In parts (d), (h), (k) and (l), dotted lines represent slip traces in B2 phase. In parts (a-l), A2 and A15 represent 2 and 15 mins of annealing (at 900 °C) respectively (Hwang et al. 2020).</figcaption></figure>

A primary requirement in the avenue of B2s Fe-Mn-Al-Ni-C steels is to control the morphology of B2 precipitates in hot-rolled products where there are limited nucleation sites for a homogeneous distribution of B2 phase in FCC matrix. This is unlike the scenario in cold-rolled products where FCC GBs and shear bands act as heterogeneous nucleation sites for B2 precipitation. In particular, dynamic recovery and/or recrystallisation leading to high dislocation annihilation rate and large FCC grain sizes in hot-rolled products, limit the number of nucleation sites for precipitation of B2 phase. Recently, Zargaran et al. (Zargaran et al. 2021) have used κ-carbide particles for promoting the nucleation of B2 precipitates (in FCC matrix)
in a hot-rolled Fe-21Mn-10Al-5Ni-1C (wt.%) steel. It was highlighted that thermomechanical
processing (in post hot-rolled condition) may be used as a potential tool towards engineering
(i) the nucleation sites and (ii) morphology of B2 phase in hot-rolled B2s Fe-Mn-Al-Ni-C
steels. Besides, it was shown that elastic strain energy (at FCC/B2 interfaces) hardly influences
the nucleation of B2 precipitates in presence of κ-carbide particles (in a hot-rolled Fe-21Mn-
10Al-5Ni-1C (wt.%) steel) (Zargaran et al. 2021). **Fig. 6** highlights the influence of
thermomechanical processing on the size distribution and morphology of B2 precipitates
(Zargaran et al. 2021). From **Fig. 6(a and b)** (showing the microstructure of Fe-21Mn-10Al-
5Ni-1C (wt.%) steel after solutionising (at 1200 °C) and subsequent annealing treatment (at
900 °C)), it may be observed that the B2 precipitates (of average size ~2 µm) are
inhomogeneously distributed in FCC matrix (Zargaran et al. 2021). However, from **Fig. 6(c
and d)** (representing the microstructure after solutionising (at 1200 °C) followed by
intermediate annealing (at 750 °C) and final annealing (at 900 °C)), a uniform distribution of
B2 precipitates (with average size ~1 µm) in FCC matrix may be observed (Zargaran et al.
2021). Furthermore, **Fig. 6(c)** shows the presence of cuboidal κ-carbides (average size ~50 nm)
within FCC matrix (Zargaran et al. 2021). High resolution Transmission Electron Microscopy
(HRTEM) based investigations (**Fig. 6(d)**) reveal the presence of disc/plate-shaped
nanoprecipitates of B2 phase adjacent to the κ-carbides, highlighting the nucleation and growth
of B2 phase on κ-carbide particles during annealing (Zargaran et al. 2021). Atom Probe
Tomography (APT) based 1D concentration profile across κ-carbides (**Fig. 6(e)**) demonstrates
that rejection of Ni from κ-carbides (during annealing) to the narrow FCC channels (present
between κ-carbide particles) leads to the enrichment of Ni at FCC matrix (Zargaran et al. 2021).
This provides a chemical driving force for the nucleation of B2 precipitates on κ-carbide
particles (Zargaran et al. 2021).
<figure><img src="image_6.png" /></figure>

**Fig. 6** Fe-21Mn-10Al-5Ni-1C (wt.%) steel: (a, b) SEM images after solution-treatment (at 1200 °C) followed by annealing at 900 °C, (c, d) SEM and high resolution TEM (HRTEM) images after intermediate annealing at 750 °C: (e) SEM image showing the presence of κ-carbides within FCC grains, and (f) HRTEM image showing a triple junction between FCC, B2 and κ-carbide phases. The corresponding Fast Fourier Transformation (FFT) pattern (inset of part (d)) shows K-S orientation relationship between B2 and κ-carbide phases, and Atom Probe Tomography (APT) analysis: (e) 1D concentration profile across κ-carbide particles (shown by a black line) within FCC grain in hot-rolled condition. In parts (a) and (b), B2 appears as bright particles in a relatively darker FCC matrix. In the inset image of part (d), the yellow and red arrows represent the diffracted spots of B2 and κ-carbide phases respectively (Zargaran et al. 2021).

Yang et al. (Yang et al. 2016) have used in-situ X-Ray Diffraction (XRD) technique to investigate the evolution of lattice strains in FCC and B2 phases in Fe-16Mn-10Al-5Ni-0.86C (wt.%) steel. It was shown that the pronounced strain partitioning between the softer γ and the relatively harder BCC phases leads to back stresses which further leads to enhanced strain hardening of the steel (Yang et al. 2016). In the context of Fe-26Mn-16Al-5Ni-5C (at. %) steel, Wang et al. (Wang et al. 2020) have shown that dislocation pinning by B2 precipitates leads to local strain hardening at FCC/B2 interfaces. This minimises the probability of strain localization and hence, damage initiation (during plastic deformation) at the aforementioned interfaces (Wang et al. 2020). It was also highlighted that pile-up of dislocations at B2 particles leads to back-stresses, further resulting in strain hardening (Wang et al. 2020). Similar
observations have been reported in the context of Fe-10Mn-9Al-4.5Ni-0.7C (wt.%) steel in Refs. (Mozumder et al. 2018; Mozumder et al. 2020a; Mozumder et al. 2021). Kim (Kim 2019) has reported that the extraordinary strain hardening observed in Fe-16Mn-10Al-5Ni-0.86C (wt.%) steel primarily arises due to the short-range ordering (SRO) of the constituent atoms present in the FCC matrix. It was shown that the back stresses arising due to strain incompatibility between FCC matrix and B2 phase play a secondary role in influencing strain hardening of the aforementioned steel (Kim 2019). This is unlike the major observation of Yang et al. (Yang et al. 2016), where it was initially reported that back stresses arising due to strain incompatibility between FCC matrix and B2 phase play a primary role in influencing strain hardening in B2s Fe-Mn-Al-Ni-C steels.

<figure><img src="image_7.png" /><figcaption>**Fig. 7** Fe-10Mn-9Al-4.5Ni-0.7C (wt.%) steel: TEM-BF images of deformed microstructure (true strain ~0.69) at 1100 °C highlighting (a, b) the non-shearable nature of nano-scale GB B2 precipitates, and (c) high dislocation density near the GB-B2 precipitate (Mozumder et al. 2021).</figcaption></figure>

Mozumder et al. (Mozumder et al. 2021) have investigated the compressive deformation behaviour of Fe-10Mn-9Al-4.5Ni-0.7C (wt.%) steel between 950 to 1150 °C at a high strain rate (~10 s⁻¹). The microstructure at the aforementioned deformation conditions was reported to comprise of fine intragranular B2 and GB B2 precipitates in BCC phase and FCC matrix respectively (Mozumder et al. 2021). **Fig. 7** shows the presence of dislocation cell structures near B2 precipitates in BCC GBs (in Fe-10Mn-9Al-4.5Ni-0.7C (wt.%) steel), highlighting that B2 precipitates (at BCC GBs) are non-shearable by the gliding dislocations (Mozumder et al. 2021). In addition, κ-carbides (shearable by gliding dislocations) were observed to form within FCC grains and at FCC/BCC interfaces. It was reported that the GB B2 precipitates aid discontinuous dynamic recrystallisation (DDRX) phenomenon (through the introduction of high stored energy near FCC GBs) but retard the growth of FCC grains by retarding GB
migration (of FCC GBs) (Mozumder et al. 2021). On the contrary, BCC phase shows DDRX mechanism at FCC/BCC interfaces (Mozumder et al. 2021). In addition, the fine intragranular B2 precipitates in BCC phase facilitate continuous dynamic recrystallisation (CDRX) in the aforementioned deformation conditions (Mozumder et al. 2021). A similar investigation (with alloy composition similar to that reported in Ref. (Mozumder et al. 2021)) at a much lower strain rate (~0.001 s⁻¹) by Mozumder et al. (Mozumder et al. 2020b) revealed that the mechanism of CDRX in both FCC and BCC phases is associated with progressive transformation of low-angle GBs (or sub-grain boundaries) into high-angle GBs. It was further highlighted that BCC shows DDRX (facilitated by FCC/BCC interfaces) at all temperatures ranging from 950-1150 °C whereas FCC phase exhibits DDRX at initial deformation stages (between 950 and 1150 °C) and dominant CDRX at later stages of deformation (between 950 and 1050 °C) (Mozumder et al. 2020b). Most importantly, it was reported that the dissolution of B2 precipitates (present near FCC GBs) in FCC matrix promotes GB migration, leading to DDRX accompanied by twinning in FCC phase (Mozumder et al. 2020b). **Fig. 8(a and b)** shows the presence of high dislocation density walls (HDDWs) and substructures within the deformed BCC grain at 950 and 1050 °C in Fe-10.8Mn-9Al-4.5Ni-0.7C (wt.%) steel (Mozumder et al. 2020b). From **Fig. 8(c)**, it is observed that HDDWs and substructures are not formed during hot compression at 1150 °C (in the aforementioned steel composition) (Mozumder et al. 2020b). This indicates that DDRX mechanism is operative during compressive deformation at 1150 °C (Mozumder et al. 2020b). **Fig. 8(d)** shows the presence of dislocation networks in FCC matrix (at 1050 °C). **Fig. 8(e)** shows a substructure (in FCC matrix) formed by the consumption of dislocations by the HDDWs (Mozumder et al. 2020b).

<figure><img src="image_8.png" /></figure>
**Fig. 8** Fe-10.8Mn-9Al-4.5Ni-0.7C (wt.%) steel: HRTEM image of deformed microstructure (strained at 0.69): (a–c) BCC and the (d–e) FCC at different temperatures (in °C): (a, d) 950; (b, e) 1050; and (c) 1150. The yellow arrows in part (c) represent the dislocations. In part (d), the white dashed circle encloses dislocation networks (Mozumder et al. 2020b).

**4. B2s Fe-Mn-Al-Ni-C steels: Nanomechanical response**

Rahnama et al. (Rahnama et al. 2018b) have investigated the nanomechanical properties of FCC and B2-NiAl phases in Fe-15Mn-10Al-5Ni-0.8C (wt.%) steel as a function of different annealing temperatures between 500 and 1050 °C using nanoindentation and Density Functional Theory (DFT). Three different morphologies of B2 phase were observed: (i) coarse precipitates (~2-7 µm) (at FCC matrix and GBs), (ii) stringer bands (in FCC matrix), and (iii) nano-sized discs (~200 nm) (in BCC phase) (Rahnama et al. 2018b). B2 stringers were observed at temperatures ≤ 700 °C (Rahnama et al. 2018b). On the contrary, B2-nano-sized discs (in BCC phase) were observed to precipitate only at temperatures >700 °C (Rahnama et al. 2018b). It was reported that an increase in annealing temperature leads to an increase in Ni content of B2 phase thereby increasing its hardness (Rahnama et al. 2018b). **Fig. 9** shows the variation of hardness (of B2-NiAl phase in Fe-15Mn-10Al-5Ni-0.8C (wt.%) steel) as a function of three different parameters, viz. (i) B2 morphology, (ii) Ni content of B2 phase, and (iii) annealing temperature (between 500 and 1050 °C) (Rahnama et al. 2018b). **Fig. 9(a and b)** shows the nanoindentation-based load-depth plots for two different morphologies of B2 phase, viz. (a) nano-sized discs (in BCC phase) and (b) stringers (in FCC matrix) (Rahnama et al. 2018b). It was specifically highlighted that B2 stringers show a lower hardness as compared to that of nano-sized discs (of B2 phase) owing to the lower Ni content of the former (**Fig. 9(c and d)**) (Rahnama et al. 2018b; Rahnama and Sridhar 2019). Nevertheless, the hardness of B2 precipitates was observed to be much higher than that of FCC matrix suggesting that B2 precipitates account for the strength whereas FCC matrix accounts for ductility in Fe-Mn-Al-Ni-C steel (Rahnama et al. 2018b).

The nanomechanical response of Fe-15Mn-10Al-5Ni-0.8C (wt.%) steel was compared with that of Fe-15Mn-10Al-0.8C (wt.%) TRIPLEX steel (for the same annealing conditions) for the purpose of understanding the influence of Ni on the nano-scale deformation behaviour of the individual phases present in the aforementioned steels. The nano-sized B2 discs were observed to be completely replaced by the micron-sized κ-carbides in FCC matrix in Fe-15Mn-10Al-0.8C (wt.%) TRIPLEX steel at temperatures ≤ 700 °C. At temperatures > 700 °C, coarse
precipitates (~2-4 µm) and nano-sized discs (~200 nm) of B2 phase were observed in FCC matrix and BCC phase respectively. Needle-shaped κ-carbides (width ~200-500 nm, formed at temperatures ≤ 700 °C in FCC matrix and BCC phase) in Fe-15Mn-10Al-0.8C (wt.%) TRIPLEX steel have been reported to be shearable by gliding dislocations, unlike that of B2 precipitates in Fe-15Mn-10Al-5Ni-0.8C (wt.%) steel (as highlighted in **section 2**) (Rahnama et al. 2017d; Rahnama and Sridhar 2019). Most interestingly, it was shown that (Ni, Al)-rich B2 precipitates (in B2s Fe-Mn-Al-Ni-C steels) exhibit a higher hardness as compared to that of (Fe, Al)-rich B2 precipitates (in Fe-Mn-Al-C steels). This may be explained on the basis of the higher bond strength of Ni-Al bond as compared to that of Fe-Al bond. The aforementioned observations (from Ref. (Rahnama et al. 2018b)) demonstrate that Ni not only influences phase evolution but also influences the nanomechanical properties of B2 phase.

<figure><img src="image_9.png" /><figcaption>**Fig. 9** Fe-15Mn-10Al-5Ni-0.8C (wt.%) steel: (a, b) Nanoindentation-based load vs depth curves of B2 (a) nano-sized discs (in BCC phase) and (b) stringers (in FCC matrix) and (c, d) the variation of hardness values and Ni content in B2 (c) nano-sized discs and (d) stringers as a function of annealing temperature between 500 and 1050 °C (Rahnama et al. 2018b).</figcaption></figure>
## 5. B2s Fe-Mn-Al-Ni-C steels: Challenges and future perspectives

At present, most of the investigations in the area of Fe-Mn-Al-Ni-C lightweight steels are primarily focussed on enhancing the strength-ductility combination through an understanding of the microstructure evolution of B2 phase. More particularly, in the context of experimental research, the avenues such as alloy design, precipitation behaviour, and strengthening mechanisms of the aforementioned steels are being extensively investigated. In addition, the present state of computation-based research (in B2s Fe-Mn-Al-Ni-C steels) is primarily focussed towards developing a database to address the thermodynamics and kinetics of B2 precipitation in these steels (as briefly highlighted in **section 2**). However, there still remain a number of unexplored avenues which need to be addressed before utilising the structural properties of Fe-Mn-Al-Ni-C lightweight steels in industrial applications (especially, automotive applications). This necessitates a collaboration of both fundamental and industrial research for the purpose of overcoming the present challenges in these steels. In the following section, challenges in both fundamental and industrial research have been highlighted.

### 5.1 Fundamental research

From an experimental viewpoint, SRO of constituent atoms in FCC phase has been reported to play the most crucial role in influencing the strain hardening behaviour of B2s Fe-Mn-Al-Ni-C steels (Kim 2019). In addition, SRO clusters largely influence plasticity (as reported elsewhere (Ref. (Gutierrez-Urrutia 2021))). Hence, understanding the influence of SRO on the deformation behaviour of the aforementioned steels is critical for controlling the microstructure evolution of B2 precipitates in FCC matrix and is presently one of the greatest challenges which needs to be addressed. B2s Fe-Mn-Al-Ni-C lightweight dual-phase contain two types of microstructural interfaces viz. homophase interfaces (such as twins and dislocation configurations) and heterophase interfaces (such as FCC/B2 interfaces). Thermal properties of microstructural interfaces is highly essential in terms of understanding interface mobility and its associated influence on the thermodynamics and kinetics of recrystallisation and grain growth. On the other hand, mechanical behaviour of microstructural interfaces is critical towards understanding the influence of plastic deformation on a number of properties such as crack nucleation and propagation, defect interaction (such as precipitate/dislocation and twin/dislocation interaction), strain partitioning, and cohesive strength of the interfaces. At present, there is a strong requirement towards understanding the influence of B2 precipitate composition and morphology on the strain hardening behaviour of B2s Fe-Mn-Al-Ni-C steels.
This avenue is highly challenging due to the possibility of local chemical and strain gradients at FCC/B2 interfaces and is also presently unexplored. It is worth highlighting that in the context of metallic materials, elemental segregation largely influences cohesive strength at different microstructural interfaces such as GBs, precipitate/matrix interphase interfaces etc. (Seah 1975; Seah 1976; Seah 1980b; Seah 1980a; Lejcek 2010). However, in the context of B2s Fe-Mn-Al-Ni-C steels, this is an unexplored avenue and shows huge potentials for future investigations.

## 5.2 Industrial research

A number of casting defects originate during the steelmaking process of B2s Fe-Mn-Al-Ni-C steels. Large additions of Mn and Al are difficult to handle during the steelmaking process of B2s Fe-Mn-Al-Ni-C steels (Gutierrez-Urrutia 2021). A number of complex chemical reactions occur between the melt and refractory metals. These may lead to the formation of impurities and a significant deviation from the desired steel composition. A very common example of this is the formation of alumina (Al₂O₃) in liquid state (during continuous casting) which leads to the clogging of nozzles (Gutierrez-Urrutia 2021). These necessitate the implementation of novel casting techniques. In terms of processing, cold-rolling and subsequent recrystallization have been reported to be the starting conditions for almost every experimental investigation in as-cast B2s Fe-Mn-Al-Ni-C steels (Kim et al. 2015; Rahnama et al. 2017a; Rahnama et al. 2017d; Rahnama et al. 2017c; Rahnama et al. 2018b; Mozumder et al. 2018; Kim 2019; Rahnama and Sridhar 2019; Park et al. 2019a; Mozumder et al. 2020a; Hwang et al. 2020; Mozumder et al. 2020b; Wang et al. 2020; Mozumder et al. 2021; Zargaran et al. 2021). Although, a recent investigation has demonstrated the formation of B2 phase (through κ-carbide assisted nucleation) in hot rolled Fe-Mn-Al-Ni-C steel (Zargaran et al. 2021), however, evolution of texture and subsequent recrystallization kinetics are the major avenues which have not been explored in detail till date.

It is worth highlighting that only the tensile (Kim et al. 2015; Park et al. 2019a; Hwang et al. 2020; Wang et al. 2020; Zargaran et al. 2021) and compressive (Mozumder et al. 2018; Mozumder et al. 2020a; Mozumder et al. 2020b; Mozumder et al. 2021) properties of B2s Fe-Mn-Al-Ni-C steels have been studied so far. Although, Park et al. (Park et al. 2019a) have discussed the influence of B2 morphology on the formability of these steels, however, structural properties critical for industrial applications such as fracture and fatigue properties, and mechanical anisotropy have not been investigated till date. The other industrial challenge
is to overcome the phenomenon of Hydrogen embrittlement in high-strength steels. This is
crucial for application of B2s Fe-Mn-Al-Ni-C steels in industrial scale since hydrogen leads to
a catastrophic and unprecedented failure of industrial components (Chen et al. 2020). In the
context of Hydrogen embrittlement, the primary challenge is to detect the position of hydrogen
in a material (Cairney 2017; Chen et al. 2017b). This may be attributed to the high atomic
mobility arising from the extremely small atomic radius (~0.037 nm) of hydrogen (Koyama et
al. 2012). At present, there are limited literatures on understanding the influence of hydrogen
on the mechanical behaviour of TRIPLEX steels. In the avenue of B2s Fe-Mn-Al-Ni-C steels,
the aforementioned phenomenon is completely unexplored. This necessitates extensive
investigations towards determining the hydrogen trapping affinities of microstructural features
(present in B2s Fe-Mn-Al-Ni-C steels) such as FCC GBs, dislocation configurations, FCC/B2
interfaces, and B2 precipitates.

**6. The need for correlative microstructural characterisation**

The emergence of the novel correlative microscopy methodology enables a simultaneous
determination of both structural and chemical information from the same region (in a
microstructure) using different characterisation tools at different length scales. The
aforementioned methodology has been widely utilised to investigate a number of phenomenon
such as GB segregation, precipitation (both at grain interiors and GBs), phase transformation,
and clustering in a number of metallic materials. For instance, Seol et al. (Seol et al. 2013c)
have investigated the mechanism of κ-carbide formation using correlative Transmission
Electron Microscopy (TEM)-Atom Probe Tomography (APT) methodology. Herbig (Herbig
2018) has used correlative Scanning Transmission Electron Microscopy (STEM)-TEM-APT
approach to obtain information such as distribution of shear bands in the near vicinity of κ-
carbides in Fe-Mn based austenitic (FCC) steel. Besides, correlative EBSD-APT and
Transmission Kikuchi Diffraction (TKD)-APT methodologies have been used to study GB
segregation in steels. In other words, the emergence of correlative microscopy has enabled the
characterisation of microstructural features which are otherwise, hard to characterise using
conventional methodology. For instance, during GB segregation studies, a separate
characterisation of GB crystallography and local chemical composition using different electron
microscopy and spectroscopy techniques, cannot be used to correlate GB structure with the
extent of segregation at GBs. This correlation is only possible when both structural and
chemical information are obtained from the same GB region. As a characterisation
methodology, correlative microscopy may be utilised to establish the aforementioned
correlation. However, in the context of B2s Fe-Mn-Al-Ni-C steels, there is hardly any correlative microstructural investigation till date. In terms of addressing the fundamental challenges, especially associated with (i) the characterisation of FCC/B2 interfaces (discussed in **section 5.1**), (ii) influence of B2 precipitate composition and morphology on the strain hardening behaviour of B2s Fe-Mn-Al-Ni-C steels (discussed in **section 5.1**), and (iii) hydrogen trapping affinities of different microstructural features (discussed in **section 5.2**), there is a huge scope of future research using correlative microscopy methodology.

## 7. Conclusions

In the context of advanced high-strength steels, addition of Ni to Fe-Mn-Al-C TRIPLEX steels has provided a pathway towards the development of B2s Fe-Mn-Al-Ni-C steels. Recent investigations on the evolution of B2 precipitates in FCC matrix and its associated influence on the mechanical properties (both micro and nanoscales) of B2s Fe-Mn-Al-Ni-C lightweight steels have highlighted that precipitation of Fe, Ni and Al-rich B2 phase leads to strengthening and excellent strain hardening in B2s Fe-Mn-Al-Ni-C lightweight steels. However, at present, there still remain a number of fundamental challenges (mainly associated with understanding the microstructure evolution of B2 phase) and industrial challenges (associated with alloy design, thermomechanical processing, and mechanical deformation behaviour) for establishing these steels as structural materials for industrial applications (especially the automotive sector). Nevertheless, these challenges render B2s Fe-Mn-Al-Ni-C lightweight steels as attractive candidates for further fundamental and industrial research.

## Acknowledgement

MS would like to thank the Department of Metallurgical and Materials Engineering, IIT Madras for its constant support during the documentation of the review article.

## Funding

The author hereby declares that no external funding was received for this work.

## Conflict of interest

The author hereby declares no conflict of interest.

## References

Baker I (1995) A review of the mechanical properties of B2 compounds. Materials Science and Engineering A 192–193(PART 1):1–13. https://doi.org/10.1016/0921-5093(94)03200-9
Cairney J (2017) Atoms on the move - Finding the hydrogen. Science (1979) 355:1128-1129

Chen S, Rana R, Haldar A, Ray RK (2017a) Current state of Fe-Mn-Al-C low density steels. Prog Mater Sci 89:345-391

Chen YS, Haley D, Gerstl SSA, London AJ, Sweeney F, Wepf RA, Rainforth WM, Bagot PAJ, Moody MP (2017b) Direct observation of individual hydrogen atoms at trapping sites in a ferritic steel. Science (1979) 355(6330):1196-1199.
https://doi.org/10.1126/science.aal2418

Chen YS, Lu H, Liang J, Rosenthal A, Liu H, Sneddon G, McCarroll I, Zhao Z, Li W, Guo A, Cairney JM (2020) Observation of hydrogen trapping at dislocations, grain boundaries, and precipitates. Science (1979) 367(6474):171-175.
https://doi.org/10.1126/science.aaz0122

Choi WS, de Cooman BC, Sandlöbes S, Raabe D (2015) Size and orientation effects in partial dislocation-mediated deformation of twinning-induced plasticity steel micro-pillars. Acta Mater 98:391-404. https://doi.org/10.1016/J.ACTAMAT.2015.06.065

Crimp MA, Vedula K (1991) Room-temperature deformation of single-crystal B2 Fe-Al alloys: The effect of stoichiometry and cooling rate. Philosophical Magazine A: Physics of Condensed Matter, Structure, Defects and Mechanical Properties 63(3):559-570.
https://doi.org/10.1080/01418619108213899

Crimp MA, Vedula KM (1993) The relationship between cooling rate, grain size and the mechanical behavior of B2 FeAl alloys. Materials Science and Engineering A 165(1):29-34. https://doi.org/10.1016/0921-5093(93)90623-M

Dey P, Nazarov R, Dutta B, Yao M, Herbig M, Friák M, Hickel T, Raabe D, Neugebauer J (2017) Ab initio explanation of disorder and off-stoichiometry in Fe-Mn-Al-C κ carbides. Phys Rev B 95(10):104108. https://doi.org/10.1103/PhysRevB.95.104108

Frommeyer G, Brüx U (2006) Microstructures and mechanical properties of high-strength Fe-Mn-Al-C light-weight TRIPLEX steels. Steel Res Int 77(9-10):627-633.
https://doi.org/10.1002/srin.200606440

Gao S, Yoshimura T, Mao W, Bai Y, Gong W, Park M, Shibata A, Adachi H, Sato M, Tsuji N (2020) Tensile Deformation of Ultrafine-Grained Fe-Mn-Al-Ni-C Alloy Studied by In Situ Synchrotron Radiation X-ray Diffraction. Crystals 2020, Vol 10, Page 1115 10(12):1115. https://doi.org/10.3390/CRYST10121115

Gladman T (1999) Precipitation hardening in metals. Materials Science and Technology 15(1):30-36. https://doi.org/10.1179/026708399773002782

Gutierrez-Urrutia I (2021) Low Density Fe-Mn-Al-C Steels: Phase Structures, Mechanisms and Properties. ISIJ International 61(1):16-25.
https://doi.org/10.2355/ISIJINTERNATIONAL.ISIJINT-2020-467

Gutierrez-Urrutia I, Raabe D (2013) Influence of Al content and precipitation state on the mechanical behavior of austenitic high-Mn low-density steels. Scr Mater 68(6):343-347.
https://doi.org/10.1016/j.scriptamat.2012.08.038
Gutierrez-Urrutia I, Raabe D (2014) High strength and ductile low density austenitic FeMnAlC steels: Simplex and alloys strengthened by nanoscale ordered carbides. Materials Science and Technology (United Kingdom) 30(9):1099-1104.
https://doi.org/10.1179/1743284714Y.0000000515

Gutierrez-Urrutia I, Raabe D (2012) Multistage strain hardening through dislocation substructure and twinning in a high strength and ductile weight-reduced Fe-Mn-Al-C steel. Acta Mater 60(16):5791-5802. https://doi.org/10.1016/j.actamat.2012.07.018

Herbig M (2018) Spatially correlated electron microscopy and atom probe tomography: Current possibilities and future perspectives. Scr Mater 148:98-105.
https://doi.org/10.1016/J.SCRIPTAMAT.2017.03.017

Hwang JH, Trang TTT, Lee O, Park G, Zargaran A, Kim NJ (2020) Improvement of strength - Ductility balance of B2-strengthened lightweight steel. Acta Mater 191:1-12.
https://doi.org/10.1016/j.actamat.2020.03.022

Kies F, Wu X, Hallstedt B, Li Z, Haase C (2021) Enhanced precipitation strengthening of multi-principal element alloys by κ- and B2-phases. Mater Des 198:109315.
https://doi.org/10.1016/J.MATDES.2020.109315

Kim H (2019) Strain hardening of novel high Al low-density steel consisting of austenite matrix and B2-ordered intermetallic second phase in the perspective of non-cell forming face-centered-cubic alloy with high stacking fault energy. Scr Mater 160:29-32.
https://doi.org/10.1016/j.scriptamat.2018.09.044

Kim SH, Kim H, Kim NJ (2015) Brittle intermetallic compound makes ultrastrong low-density steel with large ductility. Nature 518(7537):77-79. https://doi.org/10.1038/nature14144

Koyama M, Akiyama E, Tsuzaki K (2012) Effect of hydrogen content on the embrittlement in a Fe-Mn-C twinning-induced plasticity steel. Corros Sci 59:277-281.
https://doi.org/10.1016/j.corsci.2012.03.009

Lejcek P (2010) Grain Boundaries: Description, Structure. Grain boundary Segregation in Metals: 239. https://doi.org/10.1007/978-3-642-12505-8

Mozumder YH, Arun Babu K, Mandal S (2020a) Compressive Flow Behaviour and Deformation Instabilities of Fe-Mn-Al-Ni-C Lightweight Duplex Steel. Transactions of the Indian National Academy of Engineering 5:465-474. https://doi.org/10.1007/s41403-020-00105-x

Mozumder YH, Arun Babu K, Saha R, Mandal S (2018) Flow characteristics and hot workability studies of a Ni-containing Fe-Mn-Al-C lightweight duplex steel. Mater Charact 146:1-14. https://doi.org/10.1016/J.MATCHAR.2018.09.036

Mozumder YH, Babu KA, Saha R, Mandal S (2021) Deformation mechanism and nano-scale interplay of dual precipitation during compressive deformation of a duplex lightweight steel at high strain rate. Materials Science and Engineering: A 823:141725.
https://doi.org/10.1016/J.MSEA.2021.141725

Mozumder YH, Babu KA, Saha R, Sarma VS, Mandal S (2020b) Dynamic microstructural evolution and recrystallization mechanism during hot deformation of intermetallic-
hardened duplex lightweight steel. Materials Science and Engineering A 788.
https://doi.org/10.1016/j.msea.2020.139613

Noebe RD, Bowman RR, Nathal M v. (1993) Physical and mechanical properties of the B2 compound nial. International Materials Reviews 38(4):193-232.
https://doi.org/10.1179/imr.1993.38.4.193

Noebe RD, Bowman RR, Nathal M v. (1996) The Physical and Mechanical Metallurgy of NiAl. Physical Metallurgy and processing of Intermetallic Compounds :212-296.
https://doi.org/10.1007/978-1-4613-1215-4_7

Park G, Nam CH, Zargaran A, Kim NJ (2019a) Effect of B2 morphology on the mechanical properties of B2-strengthened lightweight steels. Scr Mater 165:68-72.
https://doi.org/10.1016/j.scriptamat.2019.02.013

Park KT (2013) Tensile deformation of low-density Fe-Mn-Al-C austenitic steels at ambient temperature. Scr Mater 68(6):375-379. https://doi.org/10.1016/j.scriptamat.2012.09.031

Park SW, Park JY, Cho KM, Jang JH, Park SJ, Moon J, Lee TH, Shin JH (2019b) Effect of Mn and C on Age Hardening of Fe-Mn-Al-C Lightweight Steels. Metals and Materials International 25(3):683-696. https://doi.org/10.1007/s12540-018-00230-x

Raabe D, Herbig M, Sandlöbes S, Li Y, Tytko D, Kuzmina M, Ponge D, Choi PP (2014a) Grain boundary segregation engineering in metallic alloys: A pathway to the design of interfaces. Curr Opin Solid State Mater Sci 18:253-261

Raabe D, Springer H, Gutierrez-Urrutia I, Roters F, Bausch M, Seol JB, Koyama M, Choi PP, Tsuzaki K (2014b) Alloy Design, Combinatorial Synthesis, and Microstructure-Property Relations for Low-Density Fe-Mn-Al-C Austenitic Steels. JOM 66(9):1845-1856.
https://doi.org/10.1007/s11837-014-1032-x

Rahnama A, Dashwood R, Sridhar S (2017a) A phase-field method coupled with CALPHAD for the simulation of ordered κ-carbide precipitates in both disordered γ and α phases in low density steel. Comput Mater Sci 126:152-159.
https://doi.org/10.1016/j.commatsci.2016.09.015

Rahnama A, Dashwood R, Sridhar S (2017b) A phase-field method coupled with CALPHAD for the simulation of ordered κ-carbide precipitates in both disordered γ and α phases in low density steel. Comput Mater Sci 126:152-159.
https://doi.org/10.1016/J.COMMATSCI.2016.09.015

Rahnama A, Kotadia H, Clark S, Janik V, Sridhar S (2018a) Nano-mechanical properties of Fe-Mn-Al-C lightweight steels. Sci Rep 8(1):9065. https://doi.org/10.1038/s41598-018-27345-w

Rahnama A, Kotadia H, Clark S, Janik V, Sridhar S (2018b) Nano-mechanical properties of Fe-Mn-Al-C lightweight steels. Scientific Reports 2018 8:1 8(1):1-12.
https://doi.org/10.1038/s41598-018-27345-w

Rahnama A, Kotadia H, Sridhar S (2017c) Effect of Ni alloying on the microstructural evolution and mechanical properties of two duplex light-weight steels during different
annealing temperatures: Experiment and phase-field simulation. Acta Mater 132:627-643. https://doi.org/10.1016/j.actamat.2017.03.043

Rahnama A, Spooner S, Sridhar S (2017d) Control of intermetallic nano-particles through annealing in duplex low density steel. Mater Lett 189:13-16. https://doi.org/10.1016/j.matlet.2016.11.020

Rahnama A, Sridhar S (2019) Multiscale micromechanical model of duplex low density steel with B2 ordered precipitates. Mechanics of Materials 131:136-140. https://doi.org/10.1016/j.mechmat.2019.02.003

Rana R (2014a) Low-Density Steels. JOM 66:1730-1733

Rana R (2014b) High modulus steels. Canadian Metallurgical Quarterly 53:241-242

Rana R, Liu C, Ray RK (2014) Evolution of microstructure and mechanical properties during thermomechanical processing of a low-density multiphase steel for automotive application. Acta Mater 75:227-245. https://doi.org/10.1016/j.actamat.2014.04.031

Saha, Mainak. "Order within disorder." *Mainak Saha. Order within disorder* (2022).

Saha M (2023) On the phase evolution and the associated mechanical response of 5Ni-alloyed Fe-Mn-Al-C low density steel: A Correlative Microscopy Study. Ph.D thesis. Department of Metallurgical and Materials Engineering, Indian Institute of Technology Madras, India. http://hdl.handle.net/10603/602042

Seah MP (1975) Interface adsorption, embrittlement and fracture in metallurgy. A review. Surf Sci 53:168-212

Seah MP (1980a) Adsorption-induced interface decohesion. Acta Metallurgica 28(7):955-962. https://doi.org/10.1016/0001-6160(80)90112-1

Seah MP (1976) Segregation and the Strength of Grain Boundaries. Proc R Soc Lond Ser A 349(1659):535-554. https://doi.org/10.1098/rspa.1976.0088

Seah MP (1980b) Adsorption-induced interface decohesion. Acta Metallurgica 28(7):955-962. https://doi.org/10.1016/0001-6160(80)90112-1

Seol JB, Raabe D, Choi P, Park HS, Kwak JH, Park CG (2013a) Direct evidence for the formation of ordered carbides in a ferrite-based low-density Fe-Mn-Al-C alloy studied by transmission electron microscopy and atom probe tomography. Scr Mater 68(6):348-353. https://doi.org/10.1016/j.scriptamat.2012.08.013

Seol JB, Raabe D, Choi P, Park HS, Kwak JH, Park CG (2013b) Direct evidence for the formation of ordered carbides in a ferrite-based low-density Fe-Mn-Al-C alloy studied by transmission electron microscopy and atom probe tomography. Scr Mater 68(6):348-353. https://doi.org/10.1016/j.scriptamat.2012.08.013

Seol JB, Raabe D, Choi P, Park HS, Kwak JH, Park CG (2013c) Direct evidence for the formation of ordered carbides in a ferrite-based low-density Fe-Mn-Al-C alloy studied by transmission electron microscopy and atom probe tomography. Scr Mater 68(6):348-353. https://doi.org/10.1016/j.scriptamat.2012.08.013

Sohn SS, Lee BJ, Lee S, Kwak JH (2013) Effects of aluminum content on cracking phenomenon occurring during cold rolling of three ferrite-based lightweight steel. Acta Mater 61(15):5626-5635. https://doi.org/10.1016/j.actamat.2013.06.004

Springer H, Raabe D (2012a) Rapid alloy prototyping: Compositional and thermo-mechanical
high throughput bulk combinatorial design of structural materials based on the example
of 30Mn-1.2C-xAl triplex steels. Acta Mater 60(12):4950-4959.
https://doi.org/10.1016/j.actamat.2012.05.017

Springer H, Raabe D (2012b) Rapid alloy prototyping: Compositional and thermo-mechanical high throughput bulk combinatorial design of structural materials based on the example of 30Mn-1.2C-xAl triplex steels. Acta Mater 60(12):4950-4959.
https://doi.org/10.1016/j.actamat.2012.05.017

Takeuchi A, Inoue A (2000) Calculations of Mixing Enthalpy and Mismatch Entropy for Ternary Amorphous Alloys. Materials Transactions, JIM 41(11):1372-1378.
https://doi.org/10.2320/MATERTRANS1989.41.1372

Wang W, Zhang H, Yang M, Jiang P, Yuan F, Wu X (2017) Shock and spall behaviors of a high specific strength steel: Effects of impact stress and microstructure. J Appl Phys 121(13):135901. https://doi.org/10.1063/1.4979346

Wang Z, Lu W, Zhao H, Liebscher CH, He J, Ponge D, Raabe D, Li Z (2020) Ultrastrong lightweight compositionally complex steels via dual-nanoprecipitation. Sci Adv 6(46):9543-9556. https://doi.org/10.1126/sciadv.aba9543

Yang MX, Yuan FP, Xie QG, Wang YD, Ma E, Wu XL (2016) Strain hardening in Fe-16Mn-10Al-0.86C-5Ni high specific strength steel. Acta Mater 109:213-222.
https://doi.org/10.1016/J.ACTAMAT.2016.02.044

Yao MJ, Dey P, Seol JB, Choi P, Herbig M, Marceau RKW, Hickel T, Neugebauer J, Raabe D (2016a) Combined atom probe tomography and density functional theory investigation of the Al off-stoichiometry of κ-carbides in an austenitic Fe-Mn-Al-C low density steel. Acta Mater 106:229-238. https://doi.org/10.1016/j.actamat.2016.01.007

Yao MJ, Dey P, Seol JB, Choi P, Herbig M, Marceau RKW, Hickel T, Neugebauer J, Raabe D (2016b) Combined atom probe tomography and density functional theory investigation of the Al off-stoichiometry of κ-carbides in an austenitic Fe-Mn-Al-C low density steel. Acta Mater 106:229-238. https://doi.org/10.1016/j.actamat.2016.01.007

Zambrano OA (2018) A general perspective of Fe-Mn-Al-C steels. J Mater Sci 53(20):14003-14062. https://doi.org/10.1007/s10853-018-2551-6

Zargaran A, Kim HS, Kwak JH, Kim NJ (2014) Effects of Nb and C additions on the microstructure and tensile properties of lightweight ferritic Fe-8Al-5Mn alloy. Scr Mater 89:37-40. https://doi.org/10.1016/j.scriptamat.2014.06.018

Zargaran A, Trang TTT, Park G, Kim NJ (2021) κ-Carbide assisted nucleation of B2: A novel pathway to develop high specific strength steels. Acta Mater 220:117349.
https://doi.org/10.1016/J.ACTAMAT.2021.117349