1. Introduction
Owing to the increasing severity of global warming, studies have been conducted to enhance energy efficiency and reduce fuel consumption in various industrial fields to meet the environmental regulations1-7). For instance, commercializing electric vehicles has become one of the most important objectives in the automotive industry to completely eliminate carbon dioxide emissions8-10). Scientists have been steadily researching the optimization of electric vehicles by boosting battery efficiency and developing advanced, performance-oriented, and eco-friendly steels. The use of high strength materials reduces the amount of steel required for vehicles, thereby reducing weight and increasing energy efficiency11-13). In particular, the development of Fe-Mn- Al-C steel, with a chemical composition range of Fe-(18-28)Mn-(9-12)Al-(0.7-1.2)C, has garnered considerable attention owing to its exceptional mechanical properties combined with low material density14). Scientists have reported that austenitic lightweight steels have significant potential for achieving low density because of their high Al content. Additionally, the rapid precipitation of κ-carbide through spinodal decomposition mechanisms, with a composition of (Fe, Mn)3AlC, contributes to the enhancement of their mechanical properties15,16). It was also confirmed that optimization of the κ-carbide precipitation can be achieved through heat treatments, such as age hardening processes, involving specific time and temperature ranges aligned with their chemical compositions, resulting in excellent mechanical characteristics17,18). As outlined in the literature, it is conceivable to say that austenitic Fe-Mn-Al-C lightweight steels hold significant potential for the next generation of the automotive industry.
While researches have extensively reported the metallurgical characteristics of austenitic lightweight steels, there remains a need for further studies that focus on the weldability since welding is a critical and inevitable process in every industrial field. During the welding process, the materials undergo drastic thermal cycles of heating and cooling in accordance with welding methods and welding parameters, resulting in the creation of a heat affected zone (HAZ) characterized by microstructural transition and thermal stress. In addition, potential degradation of the mechanical properties of the HAZ has been reported in austenitic lightweight steels, induced by κ-carbide precipitation. Jeong et al. investigated the effect of welding thermal cycles on the HAZ in various austenitic lightweight steels and confirmed that the dissolution and precipitation behaviors of κ- carbide during the welding process significantly influence the microstructural and mechanical characteristics19). Kim et al. also reported that welding thermal cycles have a notable influence in promoting local HAZ embrittlement in some austenitic lightweight steels, induced by both nano-sized intra granular and inter granular κ-carbide precipitation20). Therefore, it is conceivable to say that a comprehensive investigation for the microstructural and mechanical characteristics of the HAZ should be conducted to improve the usability of austenitic lightweight steel systems.
Given the prevailing situation, where the significance of metallurgical understanding of austenitic lightweight steels at high temperatures is recognized, we investigated hot ductility behavior of Fe-30Mn-10.5Al-0.9C lightweight steel to further understand the characteristics of the HAZ and ensure structural safety. Cr added alloy has also been analyzed to understand whether κ- carbide precipitation has influence on mechanical characteristics at elevated temperature, since Cr is known as an inhibitor of κ-carbide precipitation in austenitic lightweight steel system21,22). The hot ductility test was conducted using a Gleeble simulator to determine the microstructural behavior during a thermal cycle of heating and cooling. To comprehensively investigate the influence of microstructural transitions, we used scanning electron microscopy (SEM).
2. Experimental Procedure
The chemical compositions of the austenitic lightweight steels used in this study are listed in Table 1. Both original ingots were manufactured through a vacuum induction melting furnace. Each ingot was homogenized at 1,200 °C for 2 h prior to hot rolling into plates with 13 mm thickness. The rolling process was completed at 900 °C and the plates were subsequently water quenched. Additional solution treatment was conducted on each plate for 2 h at 1,050 °C prior to water quenching.
Table 1
Chemical compositions of austenitic lightweight steels
| Fe | Mn | Al | C | Cr | |
|---|---|---|---|---|---|
| Steel A | Bal. | 30.0 | 10.4 | 0.93 | - |
| Steel B | 29.7 | 10.4 | 0.96 | 3.1 |
To perform hot ductility tests, a commercial Gleeble simulator (Gleeble 1500, Dynamic Systems Inc., USA) was employed. The samples for simulation were fabricated into round bar-shaped specimens (3 mm dia.×90 mm). Heating rate and cooling rate were set as 350 °C/sec and 50 °C/sec, respectively. For the on-cooling test, temperature slightly over the zero ductility temperature (ZDT) was set as peak temperature, and reduction of area of each sample was measured for understanding hot ductility after test. The stroke rate upon reaching the intended temperature was 50 mm/sec. Fig. 1 illustrates the thermal cycle simulated in this study.
Microstructural investigation was conducted through commercial optical microscope and SEM (JSM-6360, JEOL, Japan). All specimens were polished with a silicon carbide paper up to 2,000 grit, and subsequently micro polished with a 1 μm diamond suspension. Nitric acid solution (6%) was used to etch samples at room temperature. The collected images were analyzed using a commercial image analysis software.
3. Results and Discussion
The representative microstructural images of each alloy after solution treatment are shown in Fig. 2. As shown in the Fig, both alloys exhibited a typical microstructure of austenitic steel, composed of coarse austenite grains, annealing twins, and a minor proportion of elongated δ-ferrite phases, as indicated by the yellow arrows. The average values of austenite grain size in steel A and steel B were 57 and 71 μm, respectively. The volume fraction of the ferrite phase was measured as 0.7 and 5.4% in the steel A and steel B, respectively.
Fig. 2
Representative microstructures of (a) steel A and (b) steel B, obtained using an optical microscope
Fig. 3 shows the hot ductility test results during heating and cooling thermal cycles for each alloy. As shown in the Fig, steel B had slightly lower ductility than steel A in every temperature condition. However, both steels exhibited similar hot ductility behavior on heating and cooling. The similar ductility trend of the two lightweight steels indicates that the influence of Cr on hot ductility behavior was insignificant in the current study. Previous studies reported that the addition of Cr in austenitic lightweight steels suppressed the precipitation mechanism of κ-carbide by increasing the energy for formation and growth, while the Fe-30Mn- 10.5Al-0.9C alloy system was confirmed to have rapid κ-carbide formation during the welding thermal cycle22,23). Therefore, it is reasonable to say that κ-carbide has a negligible influence on hot ductility tests in the experimented alloys. In the viewpoint of hot ductility behavior, during the heating thermal cycle ductility declined between 600 and 700 °C and subsequently recovered in the temperature range from 900 to 1,100 °C. Thereafter, significant ductility deterioration was observed as the temperature increased. The rapid decrease in ductility was induced by grain boundary liquation at elevated temperatures near the melting point in which the ZDT, since the liquefied grain boundaries break under tensile loading immediately24). The measured ZDT of steel A and steel B were 1,210 and 1,215 °C, respectively. In the case of the cooling thermal cycle, on the other hand, ductility recovery occurred in both alloys as temperatures decreased. During the initial stage, the solidification of the molten microstructure promoted significant ductility recovery behavior at temperatures ranging from 1,200 to 1,000 °C, wherein the specimens regained their mechanical characteristics. With further decrease in temperature, ductility dipped within the temperature range of 700-900 °C, and subsequently recovered at 600 °C.
Fig. 3
On-heating and cooling hot ductility test results in accordance with temperature for steel A and steel B
To understand the transition behavior of ductility in accordance with temperature for both lightweight steels, microstructural analysis was conducted through SEM. Fig. 4 shows representative images of cross-sectional microstructures near the fracture surface in steel A after on-heating hot ductility tests at 600 and 1,000 °C. According to the Fig. 4(a), surface slip traces are evident in the cross-sectional microstructure after test at 600 °C. On the other hand, refined austenite grain size indicating dynamic recrystallization (DRX) can be found after test at 1,000 °C as shown in Fig. 4(b). This deformation behavior in austenitic lightweight steels has been extensively studied and several underlying mechanisms have been identified. Yoo et al. reported that the high stacking fault energy of austenitic lightweight steels causes micro band induced plasticity, resulting in excellent mechanical properties, including high elongation characteristic25). Choi et al. observed the planar slip behavior due to the κ-carbide shearing by dislocations during tensile tests26). Therefore, it is certain that the high ductility exhibited in the studied lightweight steels and the slight difference of hot ductility behavior between steel A and steel B can be reasonably understood from the study, even though more investigation need to be conducted to completely understand the deformation mechanism of the austenitic lightweight steels. In the Fig. 4(b), the cross-section image exhibited fairly refined austenite grains and significant traces of void coalescence, indicating ductile fracturing at elevated temperature27). It is evident to say that the observed ductility recovery in the temperature range of 900-1,100 °C was induced by the exceeding combination of thermal and mechanical energy, causing the activation of the DRX mechanism. The softening effect of DRX persisted until partial melting occurred.
Fig. 4
SEM micrographs showing the cross section of steel A after on-heating ductility tests at (a) 600 and (b) 1,000 °C
In the on-cooling thermal cycle, the solidification of the melted microstructure promoted significant ductility recovery behavior at the temperature range of 1,000- 1,200 °C at the initial stage. After the specimens regained their mechanical properties, the behavior of ductility transition during the cooling thermal cycle exhibited similar tendency in the corresponding on-heating test results, but all the values were relatively low. To understand this on-cooling behavior, Fig. 5 shows a representative fractograph after on-cooling hot ductility tests at 1,100 °C. As shown in the Fig, the fractured surface included melting traces on the grain boundary, which must be a re-solidified structure during cooling after reached at ZDT. The regenerated grain boundaries are known to degrade mechanical properties owing to the segregation of the alloying elements28). In addition, according to previous researches, segregated alloying elements in the austenitic lightweight steels might promote the formation of brittle intermetallic compounds such as DO3 phase which has (Fe, Mn)3Al structure on grain boundaries when the cooling progresses, resulting in overall loss of ductility. Scientists have reported that the DO3 phase can be formed on grain boundaries in the temperature range from 550 to 825 °C in the Fe-29.5 Mn-7.8Al-1.5Si-1.05C alloy29). Jeong et al. reported that DO3 phase causes hardening and embrittling in austenitic lightweight steel steels30). It was also confirmed that the formation of DO3 phase causes ductility loss at high temperatures ranging from 800 to 900 °C in the Fe-30Mn-9Al-0.9C alloy31).
4. Summary
In this study, on-heating and on-cooling hot ductility behaviors of Fe-30Mn-10.5Al-0.9C-Cr austenitic lightweight steels were evaluated. All tests were performed using a Gleeble simulator, and the following insights were obtained according to the microstructural analysis.
1) While addition of Cr is known to retard the precipitation of κ-carbide, both alloys exhibited similar hot ductility behaviors both on-heating and on-cooling thermal cycles because κ-carbide exhibited no significant influence on hot ductility behavior.
2) Microstructural analysis demonstrated that the main reason of ductility recovery at temperatures higher than 900 °C during on-heating thermal cycle was DRX mechanism, induced by excessive thermal and mechanical energy inputs, until the grain boundary started to melt at the ZDT.
3) Both alloys exhibited evidently lower hot ductility characteristics on-cooling thermal cycles than those of on-heating thermal cycles, due to segregation of alloying elements and formation of DO3 phase during the re-generation of grain boundaries.
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