1. Introduction
The global demand for advanced structural materials for reducing carbon emissions has increased markedly. In particular, researchers and engineers have conducted extensive efforts for the development of high strength steels due to their diverse industrial applications, leading to the commercialization of various next generation alloy systems, including twinning induced plasticity (TWIP), transformation induced plasticity (TRIP), and dual-phase (DP) steels1-4). Among these, austenitic Fe-Mn-Al-C lightweight steels have attracted considerable attention owing to their unique combination of high strength, excellent ductility, and reduced density5-7). In austenitic lightweight steels, Al provides density reduction, while Mn and C control the microstructure by acting as austenite stabilizers. These alloying elements are also known to exhibit a unique κ-carbide precipitation strengthening behavior in the actual alloy state8,9). κ-carbide, which is (Fe, Mn)3AlC, is a nanometer-sized precipitate that forms a coherent interface with the austenite matrix. It has been well reported that κ-carbide can precipitate rapidly within the temperature range of 450-600 °C through appropriate heat treatment. Most studies have focused on the effect of κ-carbide in markedly enhancing the strength and hardness of lightweight steels. However, recent studies have suggested that κ-carbide may have a negative influence in various conditions, such as welds and the heat-affected zone (HAZ). Jeong et al. confirmed that inter-granular κ-carbide formed along austenite grain boundaries severely deteriorates the impact toughness of lightweight steels, particularly in the HAZ during welding, where the precipitation of κ-carbide is difficult to control10). More recently, it has been further revealed that even intra-granular κ-carbide within the austenite matrix can negatively affect impact toughness, and the associated crack propagation mechanisms have been systematically discussed11). Based on these findings, scientists have focused on advanced alloy design of austenitic Fe-Mn-Al-C lightweight steels, while retaining advantageous density reduction characteristics but suppressing κ-carbide precipitation to prevent unintended microstructural behavior in welds. For instance, the addition of Cr has been identified as an effective strategy in this context. Moon et al. confirmed that the presence of Cr increases the thermodynamic barrier for κ-carbide formation, thereby retarding its precipitation12). Jeong et al. further demonstrated that this effect remains valid under welding thermal cycles, effectively preventing mechanical property discontinuities in the HAZ13). The authors emphasized the importance of systematic investigations since there can be many other potential alloying elements that have beneficial influences on austenitic lightweight steels, either through similar or distinct metallurgical mechanisms. In particular, further research is essential since the detrimental effect of κ-carbide is more significant during welding, which is an inevitable process in actual product fabrication, as previous studies confirmed. However, studies related to clarifying the roles of various alloying elements in austenitic lightweight steels remain relatively limited. To overcome these limitations, in the present study, we investigated the effects of several alloying elements, including Mo, V, and Nb, on the metallurgical and mechanical behavior of austenitic Fe-Mn-Al-C lightweight steels with the aim of further exploring and expanding the potential of this alloy system reported in previous research. Three kinds of alloys were fabricated. For each alloy, HAZ simulations were conducted using a Gleeble thermal simulator, and detailed microstructure and mechanical properties were analyzed on both the base metal and the simulated HAZ.
2. Experimental Methods
The chemical compositions of the three austenitic Fe-Mn-Al-C alloys used in this study are listed in Table 1. All ingots were originally fabricated using a vacuum induction melting furnace. The manufactured ingots were homogenized at 1,200 °C for 2 h and subsequently hot-rolled to plates with a thickness of 13 mm followed by water quenching. After hot rolling, solution treatment was conducted at 1,050 °C for 2 h, followed by water quenching.
Table 1
Chemical compositions of austenitic lightweight steels in weight percent
| Steel | C | Mn | Al | Mo | V | Nb | Fe |
|---|---|---|---|---|---|---|---|
| A | 0.93 | 29.8 | 10.4 | - | - | - | Bal. |
| B | 0.98 | 29.8 | 10.5 | 1.9 | |||
| C | 0.94 | 29.8 | 10.5 | 2.0 | 0.5 | 0.04 |
To compare the base metal and HAZ of each steel, HAZ simulations were carried out using a Gleeble simulator (Gleeble 1500, Dynamic Systems Inc., USA). Based on a previous study that measured the zero ductility temperature (ZDT) and zero strength temperature (ZST) of Fe-30Mn-9Al-0.9C lightweight steel, the peak temperature of the HAZ was set to 1,150 °C for all alloys, with an applied heat input of 30 kJ/cm to enable a direct comparison with other reports investigating the influence of alloying elements13,14). The actual thermal cycle was calculated using Rosenthal’s heat flow equation. Rectangular bars with dimensions of 10 × 5 × 55 mm and 106 × 10 × 2 mm were utilized for HAZ simulations, and then the samples were machined into tensile and sub-sized V-notched impact specimens. Schematic diagrams of the applied HAZ thermal cycle and specimen geometries are shown in Fig. 1 and Fig. 2, respectively.
Fig. 1
Schematic of the thermal cycle for HAZ simulations with peak temperatures of 1150 °C, with a heat input of 30 kJ/cm
Fig. 2
Schematics of (a) the tensile test specimen and (b) the sub-sized impact toughness test specimen, with boxes indicating the HAZ simulation region
To evaluate the mechanical properties, tensile tests and impact toughness tests were conducted. Tensile tests were performed using a universal testing machine (Z100, Zwick Roell Group, Germany). Both base metal and HAZ specimens were tested under identical conditions with a crosshead speed of 2.4 mm/min.
Impact toughness was evaluated by a Charpy impact tester. The impact direction was set perpendicular to the rolling direction. To conservatively account for the fracture mode transition behavior of austenitic lightweight steels, the test temperature was set to -40 °C15). Prior to testing, the specimens were immersed in an ethanol bath maintained at -40 °C for 20 min, and three tests were conducted for each condition within 5 seconds after being taken out of the ethanol bath.
For microstructural analysis, specimens were mechanically polished using SiC papers followed by fine polishing with a 1 μm diamond suspension. After mirror polishing, the samples were etched with a 6% nitric acid solution at room temperature. Detailed microstructural observations were performed using optical microscopy (OM).
3. Results and Discussion
Fig. 3 shows the tensile test results of the base metal and simulated HAZ for the three austenitic lightweight steels. As shown in the figure, all three steels exhibit the typical high-strength and high-ductility behavior of austenitic lightweight steels. Among the base metals, Steel C shows the highest strength, with a yield strength of 850 MPa and a tensile strength over 1040 MPa. On the other hand, Steel A exhibits the highest elongation, recording 52% total elongation. In the HAZ specimens, it is noticeable that all three steels present strength levels comparable to those of the corresponding base metals, while the ductility levels are slightly reduced. Since all alloys underwent identical manufacturing processes, the differences observed in tensile properties can be attributed to the effects of the alloying elements, such as Mo, V, and Nb.
Fig. 3
The tensile characteristics of three austenitic lightweight steels, with solid lines for base metal and dashed lines for HAZ
To gain metallurgical insight into these tensile test results, Fig. 4 presents macro images of the fractured base metal and HAZ specimens after tensile tests. As shown in the figure, the base metal specimens underwent uniform elongation, and the final fracture occurred near the center of the gauge section in each alloy. In contrast, for the HAZ specimens, all the fractures occurred in the base metal area, which is outside of the simulated HAZ, as indicated by the yellow lines in the figure. While the fracture locations can be varied in tensile tests, it is noteworthy that no fracture occurred in the HAZ despite multiple tests, providing meaningful insight into the microstructural transition behavior that occurred in the HAZ. From the perspective of physical metallurgy, it is reasonable to say that extending and fracturing behaviors tend to initiate in the region with relatively lower strength. When the HAZ was simulated at the center of the specimen and fracture consistently occurred only in the base metal, it is conceivable to say that the HAZ region possessed relatively higher strength than the adjacent base metal. Considering that the HAZ of austenitic lightweight steels can experience κ-carbide precipitation during the welding thermal cycle, as reported in previous studies, the base metal fracturing behaviors in all HAZ specimens are indirectly associated with κ-carbide precipitation during the simulated thermal cycle16,17).
Fig. 4
Macro images of all specimens after tensile tests with yellow marks indicating HAZ simulated area
For a more precise comparison, Fig. 5 summarizes the sub-sized Charpy impact toughness test results. As shown in the figure, the toughness of the base metal is in the order of Steel A, Steel B, and Steel C, while that of the HAZ is in the order of Steel B, Steel A, and Steel C. One notable insight is that Steel B exhibits base metal toughness comparable to that of Steel A, whereas the base metal of Steel C shows significantly lower toughness than Steel A. In other words, Steel B demonstrates excellent mechanical properties in both the base metal and the welded condition, including the HAZ.
To clarify the understanding of the results of impact toughness, Fig. 6 presents the microstructural observations through OM for each steel. As shown in the figure, Steel A exhibits a typical austenitic lightweight steel microstructure containing annealing twins, and the ferrite fraction is measured to be less than 1%. In contrast, Steel B and Steel C show a significantly increased ferrite fraction, with island-shaped ferrite phases being observed. Furthermore, Steel C contains carbides distributed throughout the microstructure. According to previous studies, the κ-carbide precipitation in austenitic lightweight steels has a direct relationship with impact toughness. In other words, the difference in impact toughness between the base metal and HAZ infers κ-carbide formation behavior during the welding thermal cycle. Therefore, it is reasonable to say that the addition of Mo, V, and Nb suppressed κ-carbide formation during the simulated welding thermal cycle. In Steel B, the suppression mechanism is understood to be caused by an increased thermodynamic barrier for κ-carbide formation, induced by Mo addition. Moon et al. have reported that Mo atoms can substitute into the κ-carbide structure and consequently raise the formation energy, thereby retarding both nucleation and growth of κ-carbide18). In Steel C, κ-carbide precipitation is considered to be suppressed not only by Mo but also by V and Nb. V and Nb are representative alloying elements that preferentially form stable carbides that are distributed throughout the microstructure, thereby reducing the amount of carbon content available for κ-carbide formation.
Fig. 6
Representative microstructures of the base metal in (a) Steel A, (b) Steel B, and (c) Steel C
According to a comprehensive interpretation of the mechanical properties and microstructural analyses, it is evident that κ-carbide precipitation is suppressed in Steel B and Steel C compared to Steel A through the addition of Mo, V, and Nb. Although the current interpretation is primarily based on evaluations of mechanical properties, it provides a reasonable basis for engineering solutions related to κ-carbide precipitation behavior, such as alloying strategies for suppressing κ-carbide formation. However, this effect corresponds to suppression rather than complete elimination, which makes it reasonable to assume that κ-carbide may still precipitate in the HAZ during welding. However, the presence of κ-carbide cannot be regarded as inherently detrimental. In many previous studies, the emphasis on suppressing κ-carbide precipitation mainly arose from the severe deterioration of impact toughness, particularly under welding conditions where uncontrolled degradation occurred locally in the HAZ and led to discontinuous HAZ mechanical properties19-21). From this perspective, the degradation of impact toughness in the HAZ is significantly reduced in the Mo, V, and Nb added alloys compared to the base metal, suggesting that κ-carbide formation is limited to a tolerable level during the welding process. However, considering that both the base metal and the HAZ of Steel C exhibited the lowest impact toughness among the investigated specimens, it is evident that additional microstructural mechanisms are responsible for the reduced toughness in Steel C and remain to be clarified. Accordingly, since the current study primarily focused on κ-carbide behavior, further in-depth studies on alloy design and steelmaking processes are suggested to clarify the underlying mechanisms and complex interactions among the alloying elements.
4. Summary
In this study, the effects of Mo, V, and Nb additions on the metallurgical characteristics and mechanical properties of austenitic Fe-Mn-Al-C lightweight steels were systematically investigated, with particular emphasis on κ-carbide behavior in the HAZ under simulated welding thermal cycles. Based on the experimental results and discussions, the following conclusions can be drawn:
1) The addition of Mo, V, and Nb effectively suppressed κ-carbide precipitation during welding thermal cycles, by increased thermodynamic barrier for κ-carbide formation and carbon consumption through the formation of stable carbides instead of κ-carbide.
2) As a result of suppressed κ-carbide precipitation, the degradation behavior of impact toughness in the HAZ was significantly mitigated in the additionally alloyed steels compared to the base alloy, indicating improved mechanical stability under welding conditions.
3) Although κ-carbide precipitation was not completely eliminated, resulting in tensile fracture consistently occurring in the base metal rather than in the HAZ due to localized strengthening, such a limited level of κ-carbide precipitation does not appear to have a significantly negative influence from a metallurgical perspective.
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