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Article

Thermodynamic Modeling of Microstructural Design of Lightweight Ferritic Steels

by
Tamiru Hailu Kori
1,
Adam Skowronek
2,
Jarosław Opara
3,
Ana Paula Domingos Cardoso
4 and
Adam Grajcar
1,*
1
Department of Engineering Materials and Biomaterials, Faculty of Mechanical Engineering, Silesian University of Technology, Konarskiego, 44-100 Gliwice, Poland
2
Materials Research Laboratory, Faculty of Mechanical Engineering, Silesian University of Technology, Konarskiego, 44-100 Gliwice, Poland
3
Łukasiewicz Research Network—Upper Silesian Institute of Technology, K. Miarki, 44-100 Gliwice, Poland
4
International Zinc Association, Avenue de Tervueren, 1150 Brussels, Belgium
*
Author to whom correspondence should be addressed.
Metals 2025, 15(8), 912; https://doi.org/10.3390/met15080912 (registering DOI)
Submission received: 11 July 2025 / Revised: 8 August 2025 / Accepted: 14 August 2025 / Published: 16 August 2025
(This article belongs to the Special Issue Thermodynamic Modeling of Phase Equilibrium in Metallic Materials)
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Abstract

Ferritic lightweight steels are an emerging class of low-density steels (LDSs) with promising mechanical properties. The study aimed to develop two ferritic lightweight steels with different Mn concentrations. Al was incorporated to achieve the lightweighting effect due to its relatively low atomic mass of substitutional solutions. The C concentration was kept at a minimum level to avoid the precipitation of carbides and the Mn addition was intended to increase solid solution strengthening. Thermodynamic calculations (Thermo-Calc) were employed to design the composition, analyze the phase constituents, and predict the phase transformation behavior. Microstructural investigation and hardness tests were conducted to experimentally verify the calculations. Both produced alloys exhibited a fully ferritic microstructure. Compared to industrially produced DP980 steel, a density reduction of about 7.2% and 8.3% was attained for the Fe-0.04C-5.5Al-1.6Mn-0.075Nb and Fe-0.04C-5.6Al-5.5Mn-0.08Nb steels, respectively. The steel with the higher Mn content showed increased hardness attributed to its solution strengthening effect. An increase in the hardness values was also measured with the progress in hot-rolling thickness reductions for both alloys. The alloying elements influenced the microstructural characteristics, phase transformation behavior, density, and hardness of the newly designed lightweight steels.

1. Introduction

Over the past four decades, significant emphasis on the development of alternative new lightweight materials has prompted a reduction in fuel consumption. Since the automotive sector is the second-largest contributor to greenhouse gas emissions, globally, firm environmental regulations have been enacted. Low-density steels (LDSs), developed in the 1950s as a substitute for Fe-Cr-Ni stainless steels, have gained research interest due to their comprehensive mechanical properties and lightweight effect [1]. The integration of medium-to-high Mn concentrations along with Al alloying techniques is often employed in advanced high-strength steels (AHSSs). The density of steel decreases linearly with the addition of the elements Al, Si, C, and Mn [2]. Al plays an important role in weight reduction in lightweight steels [3], often without compromising its mechanical properties [4]. Hence, lightweight, thin-walled automotive components can be produced with improved crashworthiness [5]. Light chemical elements, due to their low atomic masses, alter the lattice parameter of steels and decrease their density simultaneously. For instance, adding 1 wt% Al results in approximately a 1.3% decrease in the density of steel [6]. Based on microstructural features, lightweight steels are categorized into four groups: ferritic, ferrite-based, austenite-based, and fully austenitic steels [7].
Ferritic lightweight steels are one of the promising advanced LDSs, consisting of the alloying elements (<0.05C-(5–9)Al-<5Mn) and they exhibit an ultimate tensile strength ranging from 200 to 600 MPa and total elongation between 10 and 40% [1]. In these steels, a high Al addition causes an ordering reaction and reduces dislocation mobility due to large solid solution strengthening [8,9]. In addition to this, a high Al content leads to κ-carbide precipitation. Hence, the Al content must be below 6.5 wt% [10]. Moreover, Al has a higher affinity for oxygen than nitrogen, and so in the presence of oxygen and nitrogen, some Al2O3 and AlN can be formed. The study of [11] reported that at a high aluminum content, a continuous oxide layer can be formed, which can completely prevent nitrogen penetration into the bulk of the samples. To improve the vulnerability of sheet steel to hot-dip galvanizing, Al fully replaced the Si addition [12]. Usually, the C concentration is kept as low as possible to avoid the formation of κ-carbide, (Fe, Mn)3AlC. It precipitates mainly at ferrite grain boundaries, which has a deleterious effect on the ductility of these steels [8,13], and often results in cracking during hot working and cold working [14]. The lattice misfit and interface shape between κ-carbides and the matrix affect the steel’s formability [15]. On the other hand, the addition of Mn decreases the formation temperature and κ-carbide besides promoting the occurrence of thermally stable austenite [16]. However, these steels are still characterized by coarse ferritic grain in the whole temperature range, resulting in poor mechanical properties, due to a lack of grain refinement during phase transformation. Thus, adding Nb to lightweight steels can promote the precipitation hardening of NbC and improve the tensile strength, while maintaining ductility [17], and enhances the strength via grain refinement resulting from delaying recovery and recrystallization [18]. In medium-Mn steels, a microaddition of Nb can improve the hardness [12]. Thus, it is essential to optimize the alloying elements’ concentration to obtain the desired mechanical properties.
Depending on the content of the alloying elements, various microstructural evolutions are observed in steel. To guide the proper selection of alloy compositions and processing strategies, modern approaches to alloy design progressively exploit the framework of computational thermodynamics [19,20]. Recently, thermodynamic calculations have emerged as a powerful tool in steel design and optimization. They enable predictions of phase compositions, fractions, and equilibria in steels [21]. Interestingly, a thermodynamic parameter set for the Fe-Mn-Al-C system has been assessed, indicating agreement with phase structures [22]. Thermo-Calc software has proven itself by enabling thermodynamic studies, phase diagrams, crystallization, and critical temperature determination [23,24]. Although several thermodynamic calculation results have been reported in works [22,23,25] for LDSs, such as phase diagrams and equilibrium phases at specific temperatures, none of these precisely align with the specific chemical compositions and temperature ranges investigated in the current research.
This study mainly focused on designing the chemical composition and thermodynamic modeling of ferritic lightweight steels. Thermo-Calc software is employed to analyze the existing phases for different Mn concentrations and their impact on phase transformation temperatures. The significance of this study is to contribute novel lightweight alloys to automotive sheets that fulfill the mechanical performance requirements as well as hot-dip galvanization processes.

2. Materials and Methods

Ferritic lightweight steels are characterized by a single fully ferritic microstructure and there is no α → γ transformation in the whole temperature range [13]. Al incorporation into steels is intended for weight reduction associated with lattice expansion and the relatively low atomic mass of substitutional solutions [26]. Additionally, Al extends the ferrite phase region into high temperatures, and increases the recrystallization temperature and grain size of ferrite [27]. During designing the chemical composition, the Al concentration is limited below 6.5% to minimize ordered lattice structures, to reduce κ-carbide formation, and in order to avoid dislocation mobility issues [28], while Si is excluded despite its ductility benefits due to its negative influence on galvanizing [29,30]. The concentration of C was kept at a low level to prevent κ-carbide formation [31], with Mn added to suppress this carbide and to improve solid solution strengthening [32,33], and Nb was included for potential grain refinement and precipitation strengthening [34]. The function of Nb addition is to form potential NbC, which acts as a fine precipitate that pins the grain boundaries and refines the microstructure, leading to a higher strength of steel while maintaining a low density [35]. The precipitation occurs after solution treatment during cooling and aging [36]. It depends on the C concentration in the alloy (the tendency to form precipitates). The addition of Nb at a low C content also aims to form precipitates but of a much smaller fraction, the purpose of which is only to refine the microstructure or most probably to act through solid solution strengthening (a solute drag effect of Nb). The limiting effect of Nb microaddition in high Al-added steel is that Al bonds with N, decreasing the driving force for Nb (C,N) precipitation [17]. The low C content in the investigated steels further decreases the probability to form NbC. However, in general, Nb should increase YS and UTS at the expense of UEl and TEl [37].
The thermodynamic calculations were performed to obtain the optimum desired phase constituents by varying the alloying element content. The calculations for designing the chemical compositions as well as phase constituent analysis were conducted using the Thermo-Calc (TCFE9 database) software (2023a) and are described in the later part of the paper. The density of each designed alloy was calculated using Thermo-Calc by applying the rule of mixtures, combining the equilibrium phase fractions with their respective molar masses and molar volumes from the thermodynamic database, which ensures that the influence of the alloying elements and temperature on the phase constitution is accurately reflected. Based on the thermodynamic calculations, two lightweight ferritic steels were designed (Table 1). The alloy melts were prepared using a vacuum induction furnace (Balzers VSG-50, PVA Industrial Vacuum Systems GmbH, Wettenberg, Germany) and the ingots were cast into hot-topped closed-bottom wide-end-up cast-iron molds with a capacity of 25 kg in the Ar atmosphere. The ingots were hot-forged using a hydraulic press at a temperature range of 1200 to 900 °C to the final cross-section of 100 × 40 mm2 billets. The billets were homogenized at 1180 °C for 30 min and hot-rolled to 22, 8, and 4 mm thickness plates between the temperature range of 1100 and 850 °C using the LPS/B semi-industrial line (Łukasiewicz Research Network—Upper Silesian Institute of Technology, Gliwice, Poland), followed by air cooling to room temperature.
The samples were cut parallel to the rolling direction and embedded in a thermosetting resin. The sample surfaces were ground using SiC paper up to a grit of 1200 and polished with diamond paste containing particles of 3 μm and 1 μm in size. Subsequently, the samples were etched with a 3% nital etchant solution (3% HNO3 and 97% ethanol) for 30 s at room temperature for microstructural investigation. The microstructural examination was conducted with a Zeiss Observer Z1m light microscope (Carl Zeiss AG, Jena, Germany).
The grain size was measured according to the methodology used for AHSS steels exhibiting elongated and lath-like microstructures [38,39,40]. As the ferritic grains exhibit highly elongated morphology, their thickness was measured instead of the diameter.
The hardness of the samples was measured using a Future-Tech FM-700 Vickers hardness tester (Future-Tech Corp., Kawasaki, Japan) with a 0.5 kgf load. Ten measurements were conducted for each sample, and any outliers were excluded. The final hardness values were calculated as the average of the remaining measurements.
Archimedes’ principle is a common method used for measuring density due to its simplicity and low cost as well as repeatability and high accuracy [41]. The actual density of the alloys was measured using Archimedes’ principle and Thermo-Calc was used for theoretical calculations.

3. Results

3.1. High-Temperature Equilibrium Phase Diagrams

In the present study, thermodynamic calculations were performed to determine the fractions of phases and phase constituents present at high temperatures. The diagrams in Figure 1, Figure 2, Figure 3, Figure 4 and Figure 5 show how the concentration of Al, Mn, and C affects the phase constituent state as a function of the temperature. The microstructure evolution diagrams were created by varying the Al content from 3 to 10 wt%, the C content from 0.03 to 0.1 wt%, and the Mn content from 1 to 7 wt%. The fraction of the phases can be controlled by varying the alloy element content. As the amount of Al increased from 3 wt% to 10 wt%, there was a significant increment in κ-carbide (KAPPA_E21) precipitation, along with the existence of its stability temperature range, extending up to 1000 °C, as shown in Figure 1a–c, since Al promotes the precipitation of κ-carbides [42] and stabilizes the ferrite (BCC_A2). At 3 wt% Al addition, cementite (CEMENTITE_D011) was observed. Further increment of Al suppressed the formation of cementite. The presence of intergranular κ-carbides at hot rolling temperatures has a detrimental effect since it induces hot cracking and is unfavorable for better strength–ductility requirements [43]. Therefore, the presence of this carbide should be avoided at hot rolling temperatures by keeping the Al content ≤ 6 wt%. However, intragranular κ-carbide formation in the matrix enhances the strength through precipitation strengthening [42].
In steels containing elements that increase hardenability, such as C, carbide precipitates can form in the presence of Mn and Al. These precipitates expose the steels to cracking during cold rolling [44]. Ferritic lightweight steels are susceptible to κ-carbide unless the C content is kept at a minimal level. The thermodynamic calculations indicated that the volume fraction of κ-carbide increases with the carbon content ranging from 0.03 to 0.1 wt%, as shown in Figure 1d–f. However, the simulation clearly indicated the C increment does not affect the stability temperature of this carbide, 800 °C. Furthermore, adding more C slightly increased the volume fraction of the austenite phase at high temperatures, since C is an austenite stabilizer. To achieve the intended microstructure and desired properties for hot working and cold working, the C and Al contents in both steels were limited to 0.03 wt% and 6 wt%, respectively.
The addition of a higher Mn concentration caused the formation of secondary intermetallic phases like BCC_B2 (iron aluminide), as indicated in Figure 2, for the Fe-0.03C-6Al-xMn system. This phase started appearing at 900 °C and increased abruptly with the increasing temperature and Mn concentration. The stringer morphology of iron aluminide affects forming by reducing the ductility in the transverse direction [45]. Mishra et al. [46] reported there was a crack nucleation in iron aluminide, which resulted from strain accumulation around the neighboring hard-oriented of this phase and brought deformation heterogeneity.
The optimization of the alloy content, particularly Mn, is crucial in ferritic lightweight steels in terms of solid solution strengthening and hindering the formation of carbides. A Mn content ranging from 1 to 7 wt% was included in the calculations to see the effects on κ-carbide formation by keeping C and Al to 0.03 and 6 wt%, respectively. The results, as shown in Figure 3a–d, implied that the further addition of more Mn to the alloys hampered the precipitation of κ-carbides. Park et al. [47] confirmed that the formation of κ-carbides was inhibited by the addition of Mn to a certain extent. Increasing the Mn content beyond 5 wt% stabilized the austenite phase more, as shown in Figure 3d. The higher Mn content increased the γ-austenite volume fraction, particularly at medium to high concentrations, which is essential for stabilizing the desired microstructure in lightweight steels [48].
The addition of Nb to alloys related to the grain refinement, as mentioned previously, increased the strength while maintaining ductility. This element leads to the precipitation of different carbides, such as FCC_A1#2 (NbC) and κ-carbides, which serve as obstacles to dislocation movement, thereby strengthening the steel [49]. The microstructural refinement and NbC carbide precipitations resulted in a strength increase. Bai et al. [50] and Cai et al. [51] stated the addition of this element increases the stacking fault energy (SFE), hardness, elongation, and strength primarily through precipitation. In addition, 0.08 wt% Nb addition to the Fe-0.03C-6Al-5Mn alloy showed decreases in the austenite phase between 500 and 1000 °C, as shown in Figure 4. The study [52] confirmed that Nb delayed recrystallization and reduced the stability of austenite.
After assessing the influences of each element, the two steels with chemical compositions of Fe-0.03C-6Al-1.5Mn-0.08Nb and Fe-0.03C-6Al-5Mn-0.08Nb were designed. Based on the calculations performed, the low-Mn lightweight steel showed a fully ferritic microstructure in the entire temperature range (see Figure 5a). The medium-Mn lightweight steel should contain a small (below 7%) fraction of austenite up to 1100 °C caused by the austenitic nature of Mn, as shown in Figure 5b.
Figure 6 presents the pseudo-binary Fe-C phase diagrams for the investigated steels (low-Mn lightweight steel and medium-Mn lightweight steel), illustrating the influence of the carbon content on the stability and distribution of phases. The equilibrium phases in the Fe-Mn-Al-C system include ferrite, austenite, κ-carbide, M23C6_D04 (M23C6), M7C3_D101 (M7C3), cementite, and NbC. These phases coexisted under conditions that were dependent on both the compositions and temperature. The addition of Al in the Fe-Mn-C steels enlarges the phase field of ferrite and shifts the austenite single-phase region toward higher carbon contents. The pseudo-binary phase diagram, as seen in Figure 6b, shows that at lower C contents, a higher Mn concentration suppresses the formation of κ-carbide at low temperatures and also contributes to widening the intercritical region. Increasing the C content facilitates the formation of the austenite phase at lower temperatures in the case of the low-Mn lightweight steel, as seen in Figure 6a, where κ-carbide appears at lower carbon concentrations and temperatures. In contrast, the increased Mn content in the medium-Mn lightweight steel suppresses κ-carbide formation. The study of [53] supports this statement. In the high-temperature range (1100–1550 °C), the low-Mn steel exhibits a slightly broader stability range of δ-ferrite compared to the steel with the higher Mn content. The maximum carbon content at which δ-ferrite is present in the low-Mn lightweight steel is approximately 0.16 wt% C at around 1410 °C, whereas in the 5Mn steel, it is 0.14 wt% C. The maximum stability temperature of this phase is also higher in the low-Mn lightweight steel (1534 °C) than in the medium-Mn lightweight steel (1517 °C), which results from the stronger ferrite-stabilizing effect of aluminum and the lower manganese content in the low-Mn lightweight steel. In the temperature range of 1380–1430 °C, a much broader three-phase field of LIQUID (liquid) + ferrite + austenite is observed in the medium-Mn steel (approx. 0.14–0.22 wt% C) compared to the low-Mn steel (approx. 0.16–0.20 wt% C). This range is critical during casting processes—the coexistence of three phases at similar temperatures and compositions may lead to segregation, microstructural heterogeneity, and an increased risk of casting defects such as microporosity or grain boundary cracking, due to differences in phase density and thermal expansion.
In the range of 700–1400 °C, a notable narrowing of the ferrite + austenite two-phase region is observed in the medium-Mn lightweight steel compared to the low-Mn steel. At a representative temperature of 800 °C, the difference in carbon content is about 0.02 wt% C. This is attributed to the strong austenite-stabilizing effect of manganese in the medium-Mn lightweight steel, which leads to an expansion of the austenite field at the expense of the ferrite + austenite region. Significant differences are also seen in the austenite + κ-carbide field. In the low-Mn lightweight steel, this region begins at approximately 0.025 wt% C and 750 °C, while in the medium-Mn steel, it begins only at about 0.13 wt.% C and at a lower temperature (around 630 °C). This shift in the κ-carbide phase toward higher carbon contents and lower temperatures is associated with the increased manganese content, which stabilizes alternative carbides and suppresses κ-carbide precipitation. In the low-temperature equilibrium regime, the medium-Mn lightweight steel exhibits a more complex phase assemblage but a much simpler phase constituent is expected under real cooling conditions.

3.2. Microstructural Results

Ferritic lightweight steels have a ferritic matrix and varying amounts of secondary phases like κ-carbide, which significantly impact their mechanical performance [54]. The microstructural results showed a fully ferritic microstructure with large and elongated grains in both alloys in the hot-rolled condition, as shown in Figure 7. The addition of a high amount of Al stabilized the ferrite phase. However, from the thermodynamic calculations, the medium-Mn lightweight steel should also consist of a small fraction of austenite. Thus, in this case, the phases predicted using Thermo-Calc for the steel containing the higher Mn are not consistent with the microstructural result. This suggests that the existing databases are not precisely aligned with the real complex Fe-Mn-Al-C system [55]. The plate thickness related to a different cooling rate following hot rolling affected the microstructure size and morphology of the as-hot-rolled samples, as indicated in Figure 8. Larger grains, even up to ~1 mm in some cases, were observed in thicker plates compared to ~200–400 μm in thinner plates. Moreover, the medium-Mn steel contains smaller grains compared to the low-Mn steel in the thicker plate (Figure 7 and Figure 8).

3.3. Density

The theoretical (using Thermo-Calc software) and the real (Archimedes’ principle) density measurements of the two alloys were conducted and are summarized in Table 2. Theoretically, a density of 7.297 g/cm3 was found for the low-Mn lightweight steel, which is almost similar to that of the 7.306 g/cm3 experimentally measured. In the medium-Mn lightweight steel, a 7.273 g/cm3 density was calculated and experimentally, 7.212 g/cm3 was found. The theoretical results are well consistent with the experimental results. The low-Mn and medium-Mn lightweight steels achieved a density reduction of about 7.23% and 8.31%, respectively, compared to DP980 automotive sheet steel [56]. The medium-Mn lightweight steel acquired a slightly higher density reduction from the Mn content (5.5 wt%) than the low-Mn lightweight steel, which contains the smaller Mn content (1.6 wt%).

3.4. Hardness

Hardness tests were used to assess the mechanical properties of the developed alloys [57]. The hardness tests were conducted for 4, 8, and 22 mm thicknesses. For the low-Mn lightweight steel, 202, 203, and 193 HV results were measured for the 4, 8, and 22 mm plates, respectively. In the same way, for the medium-Mn steel, 223, 212, and 208 HV hardness values were obtained corresponding to the 4, 8, and 22 mm plates, respectively. The slight increment in hardness was found in the medium-Mn lightweight steel due to the reason that this steel has the higher Mn content, which contributes to solid solution strengthening [32]. Rahnama et al. [58] employed nanoindentation to assess the mechanical behavior of Fe-15Mn-10Al-0.8C steels, revealing that the presence of Mn in the κ-carbide structure and the Mn-Al synergy in achieving nano-scale precipitation strengthening within the ferritic matrix notably enhanced the hardness. In both steels, there is a tendency for a reduction in the hardness values as the plate thickness increases, as shown in Figure 9, which corresponded to a different cooling rate from a finishing rolling temperature of ~850 °C. A significant density reduction was achieved at the minimum expense of the mechanical property (hardness) compared to DP980 steel, which has about 320 HV [59].

4. Discussion

The results of the thermodynamic calculations indicated the Al, C, and Mn concentrations alter the phase composition, phase fraction, and temperature stability. Al and C contributed to the precipitation of κ-carbide. Increasing the Al content contributed to the precipitation and temperature stability of κ-carbide. The studies of [22,25,30] reported Al stabilizes carbide at high temperatures. Also, the increase in the C content resulted in more precipitation of κ-carbide and the increased stability of austenite at higher temperatures. Increasing the C content enhances the austenite stability, leading to a higher volume fraction of austenite [60], and affects the primary carbide fraction and dissolution temperature, with a higher carbon content requiring higher temperatures for complete dissolution [61]. The formation of κ-carbides is influenced by the Mn content, as Mn replaces Fe in the carbide structure [62]. The thermodynamic calculations revealed that Mn hampered the precipitation of κ-carbides. Precipitation occurs through spinodal decomposition or classical nucleation and growth mechanisms, depending on the alloy composition and thermal treatment conditions [63]. Mn increases the chemical potential of C, promoting carbide formation, and suppresses C atom occupation at the L12 structure of the body-centered site and delays intragranular κ-carbide formation [47,64]. Refining coarse ferrite grains is important to improve the mechanical properties of ferritic lightweight steel, since these steels are characterized by coarse grains. Nb significantly impacts the ferritic Fe-Mn-Al-C system, enhancing its strength via grain refinement. It alters the transformation kinetics and microstructural characteristics [65]. However, while Nb enhances the strength of ferritic steels through precipitation strengthening, it may also reduce the work-hardening capacity [51].
The thermodynamic calculations under equilibrium conditions predicted the presence of κ-carbide, as seen in Figure 6. However, as clearly indicated by these diagrams, κ-carbide formation occurs only outside of the targeted composition–temperature window selected for our designed steels. Due to the deliberately low carbon content and relatively high aluminum concentrations aimed at maintaining a stable, fully ferritic microstructure, κ-carbides were predicted to be thermodynamically unstable under the investigated processing and service conditions. The experimental observations confirmed this prediction, as no κ-carbide precipitation was detected in the investigated alloys.
The microstructural results of both alloys showed a fully ferritic microstructure. The thermodynamic calculations were consistent with the experimental microstructural results in the low-Mn lightweight steel. However, in the medium-Mn lightweight steel, the thermodynamic calculation results did not fully agree with the microstructural investigation as they revealed the absence of the austenite phase. This implies that the developed databases that currently exist have limitations in predicting the thermodynamic equilibrium of the complex Fe-Mn-Al-C system with increased Al and Mn additions. The discrepancy also may relate to the potential impact of insufficient heat treatment duration, which could prevent the system from reaching equilibrium, resulting in limited diffusion and hindering the phase transformation kinetics [66].
Apart from microstructural influences, Al, Mn, and C contributed to density reduction. Each of these elements contributes exceptionally to the overall properties and density of the steel, making it suitable for the intended applications. The addition of Al is especially crucial as it reduces the density of the steel significantly [67]. The presence of Al increased the SFE, thereby suppressing the twinning effect and enhancing the ductility through dislocation slip [68]. Beyond enhancing the strength and ductility of the steel, Mn contributes to its lightweight characteristics [69]. The main concern of this study was to achieve density reduction in steels without compromising the mechanical properties required in automotive applications. A density reduction up to 8.3% was obtained in reference to the DP980 industrially produced steel. The medium-Mn lightweight steel achieved a relatively higher density reduction due to its higher Mn content. The theoretical approach and Archimedes’ principle density calculation showed excellent agreement, as shown in Figure 10.
A comparison of the hardness results showed slight increments (average of about 13 HV) for the steel containing 5.5 wt%. Mn. The high SFE of these alloys allows for unique deformation mechanisms, contributing to their hardness [7]. Rawat et al. [70] analyzed phase transformations in high-Al Fe-Al-Mn-C ferritic low-density steels, and the findings showed that the Mn content controls the formation of the primary and secondary phases, thereby affecting the hot workability and final mechanical properties. The hardness values of about 15 HV increased as the plate thickness decreased from 22 mm to 4 mm. Thicker plates generally exhibit lower hardness compared to thinner plates [71]. This trend is related to the cooling rates, as the thinner plates cool faster compared to the thicker plates, which influences the grain size as is evidenced in Figure 6, especially for the low-Mn steel.

5. Conclusions

In this study, ferritic lightweight steels were designed and analyzed thermodynamically using Thermo-Calc software. The influence of the alloying elements on the phase transformation and phase volume fraction was thoroughly examined. Based on the findings of this study, several conclusions can be drawn:
(1)
The thermodynamic calculations implied increasing the concentration of Al and C stabilizes the precipitation of harmful κ-carbides. Exceptionally, Al increases the κ-carbide phase stability temperature. Conversely, the increment in Mn content hampered the κ-carbide formation and stabilized the austenitic phase.
(2)
In the case of the low-Mn lightweight steel, the thermodynamic calculations’ prediction on the phase constituents was consistent with the microstructural results. On the other hand, for the medium-Mn lightweight steel, the calculation showed a small amount of austenite, but the microstructural results showed only the ferrite phase. This may imply that the existing databases are not precisely aligned with the real complex Fe-Mn-Al-C system. Additionally, it may be related to limited diffusion and hindered phase kinetics due to insufficient heat treatment longevity.
(3)
The analysis of pseudo-binary Fe-C phase equilibrium systems for the low-Mn and medium-Mn lightweight steels revealed a significant influence of the increased manganese content on the expansion of the austenite field, the stabilization of and shift in carbides (M23C6, M7C3), and the suppression or displacement of traditional phase fields involving ferrite and cementite. The high aluminum content in both steels promotes the presence of δ-ferrite at elevated temperatures, with a broader stability range observed in the medium-Mn lightweight steel. The three-phase field (Liquid + δ + γ), which is critical in casting processes, is more pronounced in the medium-Mn lightweight steel, potentially increasing the risk of casting defects. The medium-Mn lightweight steel also exhibits a more complex phase transformation sequence at lower temperatures, which may significantly affect the final microstructure and in-service performance.
(4)
A prominent density reduction was achieved in both steels. The low-Mn steel has a density of 7.30 g/cm3, while the medium-Mn steel achieved a relatively higher density reduction, with a density of 7.21 g/cm3. About a 7.2% and 8.3% density reduction was achieved, in comparison to DP980 automotive steel, respectively.
(5)
The low-Mn steel shows a hardness between 193 and 203 HV, whereas it increases to a range of 208–223 for the medium-Mn steel, clearly showing the solid solution strengthening effect of Mn. Higher hardness values in both steels are received for thinner plates due to faster cooling rates but the effect is only minor.

Author Contributions

Conceptualization, A.G. and T.H.K.; methodology, T.H.K. and A.S.; software, A.S.; validation, A.G., A.S. and J.O.; formal analysis, A.G. and A.P.D.C.; investigation, T.H.K.; resources, A.G.; data curation, A.G.; writing—original draft preparation, T.H.K.; writing—review and editing, A.G., A.S., J.O. and A.P.D.C.; visualization, J.O.; supervision, A.G.; project administration, A.G.; funding acquisition, A.G. and A.P.D.C. All authors have read and agreed to the published version of the manuscript.

Funding

This research was financially supported by the International Zinc Association Galvanized Autobody Partnership (IZA-GAP) program (ZCO-90 project). Tamiru Hailu Kori acknowledges financial support through the 10/010/BKM25/1243 project, Faculty of Mechanical Engineering, Silesian University of Technology, Gliwice, Poland.

Data Availability Statement

The original contribution of this research is included in the paper. For further inquiries, please contact the corresponding author.

Conflicts of Interest

The authors declare no conflicts of interest.

Abbreviations

The following abbreviations are used in this manuscript:
BCCBody-centered cube
FCCFace-centered cube
LDSLow-density steel
SFEStacking fault energy

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Metals 15 00912 g001
Figure 1. Effects of alloying elements (Al and C) on carbide formation in ferritic steels: (a) Fe-0.03C-3Al; (b) Fe-0.03C-6Al; (c) Fe-0.03C-10Al; (d) Fe-0.03C-6Al; (e) Fe-0.06C-6Al; (f) Fe-0.1C-6Al.
Figure 1. Effects of alloying elements (Al and C) on carbide formation in ferritic steels: (a) Fe-0.03C-3Al; (b) Fe-0.03C-6Al; (c) Fe-0.03C-10Al; (d) Fe-0.03C-6Al; (e) Fe-0.06C-6Al; (f) Fe-0.1C-6Al.
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Figure 2. Effects of Mn on phase composition calculated using Thermo-Calc for the Fe-0.03C-6Al-xMn system.
Figure 2. Effects of Mn on phase composition calculated using Thermo-Calc for the Fe-0.03C-6Al-xMn system.
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Metals 15 00912 g003
Figure 3. The influence of Mn on carbide formation of ferritic steel: (a) Fe-0.03C-6Al-1Mn; (b) Fe-0.03C-6Al-3Mn; (c) Fe-0.03C-6Al-5Mn; (d) Fe-0.03C-6Al-7Mn.
Figure 3. The influence of Mn on carbide formation of ferritic steel: (a) Fe-0.03C-6Al-1Mn; (b) Fe-0.03C-6Al-3Mn; (c) Fe-0.03C-6Al-5Mn; (d) Fe-0.03C-6Al-7Mn.
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Figure 4. Effect of Nb on carbide formation of ferritic steel: (a) Fe-0.03C-6Al-5Mn; (b) Fe-0.03C-6Al-5Mn-0.08Nb.
Figure 4. Effect of Nb on carbide formation of ferritic steel: (a) Fe-0.03C-6Al-5Mn; (b) Fe-0.03C-6Al-5Mn-0.08Nb.
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Figure 5. The equilibrium volume fraction evolution: (a) low-Mn lightweight steel; (b) medium-Mn lightweight steel.
Figure 5. The equilibrium volume fraction evolution: (a) low-Mn lightweight steel; (b) medium-Mn lightweight steel.
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Figure 6. Pseudo-binary Fe-C phase diagrams: (a) low-Mn lightweight steel; (b) medium-Mn lightweight steel.
Figure 6. Pseudo-binary Fe-C phase diagrams: (a) low-Mn lightweight steel; (b) medium-Mn lightweight steel.
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Figure 7. Light microscope images of produced steels: low-Mn lightweight steel (a) 4 mm; (b) 8 mm; (c) 22 mm plates; medium-Mn lightweight steel (d) 4 mm; (e) 8 mm; (f) 22 mm plates.
Figure 7. Light microscope images of produced steels: low-Mn lightweight steel (a) 4 mm; (b) 8 mm; (c) 22 mm plates; medium-Mn lightweight steel (d) 4 mm; (e) 8 mm; (f) 22 mm plates.
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Figure 8. The average grain size of 4, 8, and 22 mm plates for both alloys.
Figure 8. The average grain size of 4, 8, and 22 mm plates for both alloys.
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Metals 15 00912 g009
Figure 9. Comparison of hardness results for steel plates of different thicknesses.
Figure 9. Comparison of hardness results for steel plates of different thicknesses.
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Figure 10. Density comparison of experimental measurements and theoretical calculations.
Figure 10. Density comparison of experimental measurements and theoretical calculations.
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Table 1. Chemical composition (wt%) of developed steels.
Table 1. Chemical composition (wt%) of developed steels.
Chemical Composition/wt%CAlMnNbPmaxSmaxO, ppmN, ppm
Low-Mn lightweight steel0.045.51.60.075<0.010<0.0101625
Medium-Mn lightweight steel0.045.65.50.080<0.010<0.010620
Table 2. Comparison of theoretical and experimental density approaches and density reduction with respect to DP980 steel (7.85 g/cm3).
Table 2. Comparison of theoretical and experimental density approaches and density reduction with respect to DP980 steel (7.85 g/cm3).
ApproachElement (wt%)Density (g/cm3)Density Reduction (%)
CAlMnNbArchimedes’ (A)Thermo-Calc (TC)
Theoretical0.036.01.60.080-7.3067.00
Experimental0.045.51.60.0757.297-7.23
Theoretical0.036.05.00.080 7.2737.00
Experimental0.045.65.50.0807.212-8.31
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Kori, T.H.; Skowronek, A.; Opara, J.; Cardoso, A.P.D.; Grajcar, A. Thermodynamic Modeling of Microstructural Design of Lightweight Ferritic Steels. Metals 2025, 15, 912. https://doi.org/10.3390/met15080912

AMA Style

Kori TH, Skowronek A, Opara J, Cardoso APD, Grajcar A. Thermodynamic Modeling of Microstructural Design of Lightweight Ferritic Steels. Metals. 2025; 15(8):912. https://doi.org/10.3390/met15080912

Chicago/Turabian Style

Kori, Tamiru Hailu, Adam Skowronek, Jarosław Opara, Ana Paula Domingos Cardoso, and Adam Grajcar. 2025. "Thermodynamic Modeling of Microstructural Design of Lightweight Ferritic Steels" Metals 15, no. 8: 912. https://doi.org/10.3390/met15080912

APA Style

Kori, T. H., Skowronek, A., Opara, J., Cardoso, A. P. D., & Grajcar, A. (2025). Thermodynamic Modeling of Microstructural Design of Lightweight Ferritic Steels. Metals, 15(8), 912. https://doi.org/10.3390/met15080912

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