1. Introduction
As global efforts to reduce dependence on fossil fuels intensify, the development of structural materials with low density and high strength has become a critical priority in materials science and engineering1-4). Among these, austenitic lightweight steels have emerged as highly promising candidates, drawing considerable attention due to their unique combination of reduced density and excellent mechanical performance. The primary advantages of austenitic lightweight steels originate from their alloy design, aluminum additions to reduce density and manganese and carbon additions to stabilize the austenitic phase and enhance strength. It is well understood that further strengthening can be achieved through the precipitation of κ-carbide, which is known to form rapidly when the material is exposed to temperatures between 400-600 °C5-7). This precipitation significantly improves tensile strength without decreasing ductility.
However, recent investigations have reported the detrimental influence of κ-carbide precipitation. Jeong et al. demonstrated that the rapid formation of κ-carbide particles can compromise structural safety by significantly reducing impact toughness8). The authors further demonstrated that the embrittlement mechanisms of κ-carbide differ between intergranular and intragranular precipitation in accordance with thermal histories in austenitic lightweight steels, confirming that the degree of embrittlement is closely governed by the volume fraction and size of the precipitates. Moreover, the presence of κ-carbide has been shown to significantly influence weldability, particularly within the heat-affected zone (HAZ), where the dynamic thermal cycles inherent to welding induce κ-carbide precipitation at relatively low temperatures and dissolution at elevated temperatures, resulting in complex and unpredictable metallurgical characteristics for structural service9). To mitigate the negative impact of κ-carbide, recent research has focused on suppressing κ-carbide precipitation while preserving the advantages of alloying elements. Among various strategies, several studies have demonstrated that Cr effectively suppresses unintended κ-carbide precipitation in the HAZ, showing great potential for improving weldability10). Based on these findings, it is anticipated that Cr addition may enable the combination of low density, high strength, and preserved impact toughness in austenitic lightweight steels in welded structure.
Despite these promising results, previous studies examining the influence of Cr have been limited to controlled laboratory conditions with fixed heat input and welding parameters11-13). In industrial applications, a broad range of welding processes such as GTAW, FCAW, SMAW, and SAW are employed with different heat input conditions14-16). Among various welding conditions, high heat-input processes expose the HAZ to prolonged high temperatures and complex thermal gradients, making it essential to evaluate κ-carbide behavior under realistic welding scenarios. Therefore, this study aims to investigate whether the addition of Cr can effectively suppress κ-carbide precipitation in the HAZ of austenitic lightweight steels under varying heat input conditions. Simulated HAZ specimens were fabricated using Gleeble thermal cycles, and κ-carbide formation was evaluated indirectly through hardness measurements, which reflect κ-carbide precipitation behavior.
2. Experimental Methods
To examine the metallurgical influence of chromium in austenitic lightweight steels, Fe-30.0Mn-10.4Al-0.93C (Steel A) and Fe-29.7Mn-10.4Al-3.1Cr-0.96C (Steel B) were manufactured using a vacuum induction melting furnace. After casting, each ingot was homogenized at 1,200 °C for 2 h, followed by hot rolling to a thickness of 13 mm. The final rolling temperature was 900 °C, and the plates were rapidly cooled by water quenching. Subsequently, a solution treatment was conducted at 1,050 °C for 2 h to avoid unintended microstructural transitions prior to water quenching.
HAZ simulations were performed using a Gleeble simulator (Gleeble 1500, Dynamic Systems Inc., USA) on round-bar shaped specimens with a diameter of 3 mm and a length of 90 mm. The representative peak temperature of the HAZ was set at 1,150 °C for every case, with heat input ranging from 10 kJ/cm to 300 kJ/cm. Fig. 1 illustrates a schematic of the thermal cycle for HAZ simulation. All thermal cycles were calculated based on Rosenthal’s heat flow equation.
Fig. 1
Schematic of the HAZ simulations with welding thermal cycles for specimens having a peak temperature of 1,150 °C and heat inputs from 10 to 300 kJ/cm
Micro Vickers hardness testing was performed on the austenite phase of each specimen using a micro-indentation device (HMV-2, Shimadzu, Japan) with a load of 1.96 N. Hardness values reported in this study correspond to the average of 20 measurements per specimen.
Feritscope and X-ray diffraction (XRD, SmartLab, Rigaku, Japan) with scanning angles from 20° to 80° under 200 mA and 45 kV conditions were used for phase identification, while optical microscopy (OM) and scanning electron microscopy (SEM, JSM-6360, JEOL, Japan) were utilized to obtain detailed microstructural images. For the microstructural analyses, the samples were mechanically polished using SiC papers ranging from 100 to 2000 grit, followed by mirror polishing with a 1 μm diamond suspension. To reveal detailed microstructural features, etching was conducted at room temperature using a solution composed of 6% nitric acid in ethanol. The grain size of each specimen was measured by analyzing SEM images. Thin foils were prepared for transmission electron microscopy (TEM) analysis, using a twin-jet electro-polisher (Tenupol-5, Struers, Denmark) with an electrolyte solution containing 5% perchloric acid in methanol at -40 °C. All the images obtained from each analysis were analyzed using commercial image-analysis software.
3. Results and Discussion
Fig. 2 presents representative microstructural images of the base metal and HAZ samples of Steel A and Steel B. For the purpose of metallurgical discussion, HAZ samples subjected to a heat input of 30 kJ/cm are referred to as L-HAZ, whereas those with a heat input of 300 kJ/cm are referred to as H-HAZ. The measured grain sizes and volume fractions of the ferrite phase under each condition are summarized in Table 1. As shown in the figure, both alloys exhibit typical austenitic microstructures containing annealing twins and a small fraction of elongated ferrite phases, indicated by colored arrows. The morphology of elongated ferrite suggests both alloys underwent a ferrite-to-austenite (F-A) solidification mode, subsequently deformed by the hot rolling process17). However, the microstructures of the HAZ samples exhibit notably different behaviors. As confirmed by Fig. 2 and Table 1, the HAZ of Steel A shows significant grain coarsening, which is a characteristic feature of the coarse-grained HAZ (CGHAZ). In contrast, Steel B displays a decrease in grain size in both HAZ samples.
Fig. 2
Representative micrographs of the base metals, L-HAZ, and H-HAZ of Steel A and Steel B, with colored arrows indicating elongated δ-ferrite within the microstructures
Table 1
Summary of grain size and ferrite fraction measured in the base metal, L-HAZ, and H-HAZ of Steel A and Steel B
| Grain size (μm) | Ferrite fraction (%) | |||
|---|---|---|---|---|
| Steel A | Steel B | Steel A | Steel B | |
| Base | 56.9 | 88.4 | 0.7 | 3.4 |
| L-HAZ | 70.1 | 71.2 | 0.3 | 7.6 |
| H-HAZ | 103.4 | 42.0 | 0.2 | 8.9 |
Additionally, the volume fraction of the ferrite phase decreases in Steel A following HAZ simulation, whereas Steel B demonstrates an increased ferrite fraction in the HAZ, with the extent of ferrite formation increasing as the heat input increases. This microstructural transition can be attributed to the higher Cr content in Steel B, which exhibits a strong ferrite-stabilizing effect and results in a higher ferrite fraction. Moreover, prolonged exposure to elevated temperatures during the welding thermal cycle, where δ-ferrite formation is favored, promotes additional formation of ferrite phase in the HAZ of Steel B. These findings are consistent with previous reports indicating that the δ-ferrite phase becomes stable in Fe-(25-31)Mn-(0.7-1)C-xAl austenitic lightweight steels when the Al content exceeds approximately 9.7%18). Furthermore, it can be inferred that the presence of the ferrite phase suppresses austenite grain coarsening during the welding thermal cycle, resulting in relatively small grain sizes even in the HAZ subjected to a heat input of 300 kJ/cm. Previous studies have also reported the grain-growth-inhibiting effect of δ-ferrite in various alloys19,20). Another important insight derived from the HAZ microstructure is the influence of alloying elements on κ-carbide precipitation behavior. It has been reported that Cr atoms can substitute for Al in the κ-carbide crystal structure, thereby increasing the thermodynamic energy required for κ-carbide precipitation and effectively suppressing its nucleation and growth21). In addition, considering that Al is also a well-known ferrite stabilizer, ferrite formation during the welding thermal cycle can consume key alloying elements required for κ-carbide formation, thereby providing an additional mechanism for inhibiting κ-carbide precipitation within the austenite matrix. In other words, ferrite formation in Steel B during the welding thermal cycle provides an additional metallurgical mechanism for suppressing κ-carbide formation during welding. It is noteworthy that similar effects may be achieved by introducing other ferrite-stabilizing alloying elements, offering various possibilities for alloy design in austenitic lightweight steel systems.
To clarify the relationship between phase transformation and heat input, Fig. 3 presents XRD data obtained from the base metal and H-HAZ samples of both alloys. As illustrated in the Fig, the normalized XRD intensity exhibits trends consistent with the microstructural analysis. In the case of Steel A, the diffraction profiles of the base metal and H-HAZ are largely comparable. Although distinct κ-carbide peaks are not clearly observed in the normalized XRD results, it is reasonable to assume that κ-carbide precipitation has likely occurred, as previous studies have demonstrated that increasing heat input directly promotes κ-carbide formation in Fe-30Mn-xAl-0.9C lightweight steels22). In contrast, Steel B displays a significant increase in ferrite peaks in the H-HAZ compared to the base metal, demonstrating preferential stabilization and enrichment of the ferrite phase during the welding thermal cycle. In addition, the H-HAZ of Steel B contains the DO3 phase, an ordered phase with the crystal structure of (Fe,Mn)3Al. It has been reported that the DO3 phase in austenitic lightweight steels can reduce overall ductility and embrittle the alloy23).
Fig. 3
Normalized XRD patterns illustrating the phase evolution of the base metal and H-HAZ specimens for Steel A and Steel B
For a more detailed investigation of κ-carbide precipitation behavior, which could not be sufficiently elucidated by XRD analysis, Fig. 4 presents the TEM results obtained from the L-HAZ samples of Steel A and Steel B. As shown in the figure, the selected-area diffraction patterns (SADPs) from both samples exhibit superlattice reflections corresponding to κ-carbide, confirming its presence within the austenite matrix. However, comparative analysis of the dark-field TEM images reveals a noticeable difference in κ-carbide distribution between the two alloys, despite both being subjected to the same welding heat input. These results suggest that the effect of Cr on κ-carbide formation in the present study is consistent with previous reports, which confirmed that Cr addition alters both the interfacial and elastic strain energies in austenitic lightweight steels24). Furthermore, as demonstrated in this study, Cr can also reduce the chemical driving force for κ-carbide formation during the welding thermal cycle due to its role as a ferrite stabilizer. The distinct κ-carbide precipitation behavior observed between Steel A and Steel B is believed to result from the combined effects of these mechanisms.
Fig. 4
Dark-field TEM images and corresponding SADPs showing the κ-carbide precipitation behavior within the L-HAZ of (a) Steel A and (b) Steel B
Fig. 5 shows the variation in surface microhardness of Steel A and Steel B within the same HAZ region as a function of welding heat input. As illustrated in the Fig, the two specimens exhibit clearly distinguishable hardness levels under all conditions. In the base metal, Steel A exhibits relatively high hardness, and the hardness behavior of the two alloys diverges more significantly following the HAZ simulation. For Steel A, the hardness in the HAZ increases relative to the base metal, and this increase becomes more pronounced with higher heat input. At a heat input of 300 kJ/cm, the hardness exceeds 294 HV. In contrast, Steel B displays no significant change in hardness after the HAZ simulation compared to the base metal, regardless of the applied heat input, including at 300 kJ/cm. Considering that previous studies have established microhardness in austenitic lightweight steels as a reliable indirect indicator of κ-carbide precipitation behavior, the absence of hardness change in the HAZ of Steel B can be interpreted as evidence that κ-carbide formation was effectively suppressed even under the applied thermal cycles.
4. Summary
In this study, the effects of chromium addition on the phase transformation and mechanical characteristics of the HAZ in austenitic lightweight steels were systematically investigated under various welding heat input conditions. The following conclusions can be drawn:
1) The addition of 3.1 wt% Cr to Fe-30Mn-10.4Al-0.9C austenitic lightweight steel promoted ferrite formation and inhibited austenite grain coarsening in the HAZ regardless of welding heat input, due to the strong ferrite-stabilizing effect of Cr.
2) TEM analysis confirmed that Cr addition effectively suppressed κ-carbide precipitation in the HAZ, which is attributed to both the increased energy barrier for κ-carbide nucleation induced by Cr and the consumption of κ-carbide-forming elements through ferrite formation.
3) As a result of κ-carbide suppression, the Cr-containing alloy exhibited nearly constant hardness in the HAZ regardless of heat input, suggesting that other alloying elements with similar characteristics may also provide comparable benefits in austenitic lightweight steels.
Overall, these findings significantly broaden the design flexibility and potential applications of Fe-Mn-Al-C austenitic lightweight steels with improved weldability for future structural materials.
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