1. Introduction
Owing to the increasing requirements of energy efficiency in transportation systems, advanced steels with superior mechanical characteristics have been extensively studied in various industrial fields to meet the environmental regulations1-7). Among the different alloy systems explored, the austenitic Fe-Mn-Al-C lightweight steels, covering the chemical compositional range of Fe-(15-30)Mn-(2-12)Al-(0.5-2.0)C, have been developed with pricise controlling the alloying elements because of their unique combination of advanced mechanical properties and density reduction8-10). Res- earchers have focused on the formation mechanism and usability of nano-sized κ-carbide precipitation in austenitic lightweight steels, which plays a critical role in enhancing the mechanical properties of austenitic lightweight steels with a crystal structure of (Fe, Mn)3 AlC11-13). Ley et al. suggested an optimized heat treatment process to simultaneously obtain superior strength and toughness by controlling the size of κ-carbide precipitation in Fe-30Mn-9Al-1C-1Si-0.7Mo austenitic lightweight steel14). It has been also established that the addition of alloying elements in austenitic lightweight steels directly affects the mechanical characteristics owing to the transition of thermodynamic stability of phases15-17).
As interest in austenitic lightweight steel has steadily increased, the investigation into the weldability of the Fe-Mn-Al-C alloy system has become significantly more important since welding is a crucial process in the manufacturing of products across most industrial fields. Previous studies have provided valuable insights into the metallurgical behavior of austenitic lightweight steels during welding thermal cycles. Jeong et al. investigated the influence of precipitation behavior in the heat affected zone (HAZ), where significant microstructural and mechanical transitions can occur18). Scientists also confirmed that alloy design can effectively control the precipitation behavior of κ-carbide in the HAZ19). Kim et al. reported factors affecting ductility of the austenitic Fe-30Mn-9Al-0.9C lightweight steel at elevated temperature, which could be potential risk in welding processes20). However, further information is needed to confirm the suitability of these steels for industrial applications. In particular, solidification cracking susceptibility, which can lead to critical defects and severe failures of products, should be analyzed in detail to improve the weldability of the high manganese steel system21).
Given that the Fe-Mn-Al-C alloy system contains a significant amount of alloying elements including potentially detrimental elements with respect to hot cracking, it is essential to investigate the weld solidification cracking susceptibility of austenitic lightweight steels to ensure structural safety. In the current study, we examined the weld solidification cracking susceptibility of various Fe-Mn-Al-C austenitic lightweight steels in relation to solidification behavior and microstructural characteristics. To perform the analysis, Varestraint weldability tests on Fe-30Mn-10.5Al lightweight steels with varying carbon content were carried out. Micro- structural analyses were performed using a scanning electron microscopy (SEM) and a transmission electron microscopy (TEM) to understand the relationship between microstructure and solidification cracking susceptibility in each alloy.
2. Experimental Methods
The chemical compositions of lightweight alloys used in the current study are listed in Table 1. The original ingots were manufactured using a vacuum induction melting furnace. Each ingot was homogenized at 1,200 °C for 2 hours and then hot-rolled into plates with a thickness of 13 mm prior to water quenching. The finish rolling temperature for each plate was 900 °C. After the hot-rolling, solution treatment was conducted at 1,050 °C for 2 hours and then the plates were quenched in an ice water bath.
Table 1
Chemical compositions of the austenitic light-weight steel samples
| Steel (wt.%) | C | Mn | Al | Fe |
|---|---|---|---|---|
| A | 0.74 | 29.9 | 10.5 | Bal. |
| B | 0.92 | 29.8 | 10.4 | |
| C | 1.15 | 29.7 | 10.4 |
To investigate the solidification cracking susceptibility of the alloys, a longitudinal Varestraint test was carried out using a machined and ground rectangular shaped specimen (127 mm × 25.4 mm × 3 mm). Autogenous gas tungsten arc welding with Ar shielding gas was performed from one side of the specimen to the other. The detailed welding conditions are listed in Table 2. During the welding process, the augment strain was applied via die block with different radius to achieve the targeted curvature, using a pneumatically actuating yoke. The applied strain, ranging from 1% to 4%, was calculated as t/2R, where the t is the thickness of specimen and R is the radius of die block22).
Table 2
Welding parameters for longitudinal varestraint test
| Welding current | 100 A |
|---|---|
| Welding voltage | 12 V |
| Travel speed | 4 mm/sec |
| Shielding gas | Argon |
| Flow rate of shielding gas | 15 L/min |
Microstructural analyses were performed using a field emission scanning electron microscope (FE-SEM, JSM-6700F, JEOL, Japan) and the transmission electron microscope (TEM, JEOL 2100, JEOL, Japan). For FE-SEM observation, specimens were prepared by mechanical polishing with SiC papers up to 2,000 grit followed by micro-polishing with 1 µm diamond suspension. The samples were then etched with a 6% nitric acid solution in ethanol, at room temperature. To conduct TEM analysis, focused ion beam (FIB, Scios, FEI, USA) sampling method was employed for site-specific observation. All the images collected in microstructural observations were utilized for detailed investigation, using a commercial image analyzing software.
3. Results and Discussion
Fig. 1 exhibits representative microstructural images of base steel in each alloy. As shown in the figure, the microstructures of base steel in experimented alloys exhibit typical austenitic characteristics, featuring angled grains with twin boundaries. The average effective austenite grain size and ferrite fraction were measured as 53.4 µm and 6.6 % in the steel A. As the carbon content increased, the austenite grain size expanded while the fraction of ferrite phase decreased. The grain size and ferrite fraction were 72.5 µm and 0.81 % in the steel B, and the steel C exhibited fully austenitic matrix with the effective grain size of 91.1 µm.
Fig. 1
Representative base steel microstructure images of the steel A, steel B, and steel C samples with yellow arrows indicating ferrite phase
Fig. 2 illustrates the transition behavior of maximum crack length (MCL) in accordance with strain in the longitudinal Varestraint test. The results indicate that the carbon content significantly influences the solidification cracking susceptibility in the Fe-30Mn-10.5 Al-xC lightweight alloys. As shown in the figure, MCL steadily increased in all specimens as the augmented strain increased. In the steel A, no crack was observed in the 1% augmented strain condition, and the MCL was measured as 88 µm in the strain condition of 4%. In the steel B, the solidification crack lengthened from 34 µm to 171 µm as applied strain increased. In the steel C, the MCL increased drastically compared to the others. It was measured from 130 µm to 373 µm in accordance with applied strain condition. While much data should be required for verification, the current results exhibit a linear relationship between the ferrite fraction of fusion zone (F) and the MCL in the strain condition of 4% (L4%) as follow:
Fig. 2
Varestraint test results of maximum crack length (MCL) at the fusion zone as a function of augmented strain in Fe-30Mn-10.5Al-xC austenitic lightweight steels
It should be noteworthy that while there was evident increment of MCL values in accordance with experimental conditions. In addition, all the experimented alloys had great resistance to solidification crack susceptibility despite the large amount of alloying elements, demonstrating less susceptibility than austenitic stainless steel welds23). Yoo et al. performed a varestraint test on Fe-18Mn-0.6C-xSi alloys, which exhibited cracks exceeding 600 μm at 4% augmented strain24). In the case of Fe-15Mn-0.5C-3.5Al-xCr alloys, cracks exceeding 700 μm were observed depending on the Cr content25). Furthermore, in comparison with previous studies on austenitic stainless steel welds, the Fe-30Mn-10.5Al-xC lightweight steel demonstrates excellent resistance to solidification cracking26).
To investigate the relationship between chemical composition and solidification cracking mechanism in the experimented lightweight alloys, Fig. 3 presents SEM micrographs of fusion zone microstructure for each alloy, which were etched to reveal the distribution of the ferrite phase. According to the figures, it was confirmed that all the lightweight alloys exhibited F-A solidification mode, as the fusion zones displayed a vermicular δ-ferrite dendritic structure. The fraction of ferrite phase decreased from 19.8% (Steel A) to 5.9% (Steel B) and 2.1% (Steel C) as the carbon content increment, as indicated in the figures. It is reasonable to state that the stability of δ-ferrite is primarily influenced by the addition of aluminum, while variations in its content among different steels are generally attributed to differences in carbon content, as reported in various studies27-30). It is obvious that the microstructure of fusion zone directly affected solidification crack resistance in each alloy, as the reduction in primary δ-ferrite is one of the significant factors in accelerating crack generation and propagation. The δ-ferrite phase has not only a higher solubility for elements with low melting temperatures but also greater elongation at elevated temperatures compared to the austenite phase31-33). In addition, it was reported that F-A solidification mode reduces both the wettability of the liquid phase and the shrinkage ratio during solidification compared to the other modes34). It is also believed that the primary δ-ferrite phase has a crack reducing effect as ferrite-austenite interface energy is lower than that of austenite-austenite interface35). Consequently, it is reasonable to conclude that the high content of Al contributes to suppress overall cracking behavior, as aluminum is a strong ferrite stabilizer. It is well known that the addition of aluminum can decrease the activity and diffusivity of carbon in austenitic steel system36). It has also been reported that aluminum inhibited the formation of the eutectic phase by maintaining the liquid phase at relatively low temperatures37). Considering the metallurgical roles of high aluminum content and the associated solidification mechanism in the lightweight steels, it is evident that the increase in MCL in the Varestraint test is directly related to δ-ferrite reduction.
Fig. 3
Representative SEM micrographs showing ferrite phase of the weld in (a) steel A, (b) steel B, and (c) steel C
For a more detailed understanding of solidification cracking behavior, Fig. 4 shows a representative SEM micrograph of a crack formed in the steel C under a 4% strain condition. According to the figure, steel C exhibited enlarged cracks with a continuous secondary phase along the grain boundaries, which was observed exclusively in the steel C. The precipitation at the crack tip can be evidence of eutectic phase which widens the solidification temperature range during the welding thermal cycle. This terminal stage of solidification may be an important factor increasing solidification crack susceptibility, promoting generation and propagation of cracks along the grain boundaries38,39). In addition, the trace of κ-carbide precipitation was also observed in the austenitic phase. Considering that the temperature range of κ-carbide formation mechanism is significantly lower than that of the fusion zone solidification, it is reasonable to conclude that the effect of κ-carbide precipitation on the solidification crack susceptibility is likely negligible40-43). In other words, κ-carbide affects the metallurgical characteristics of austenitic Fe-Mn-Al-C lightweight steels only after solidification has been completed during the welding process. In the final product stage, on the other hand, the presence of κ-carbide may have a significant impact, as its formation can enhance tensile strength and hardness while reducing ductility and impact toughness. Previous study confirmed that κ-carbide can rapidly form during the welding thermal cycle in austenitic Fe-Mn-Al-C lightweight steels, leading to increased strength and toughness loss in the simulated HAZ44).
Fig. 4
Representative solidification crack appeared in the steel C with (a) low magnification showing overall crack appearance and (b) high magnification showing crack tip with κ-carbide and grain boundary precipitates
To understand the increase of solidification cracking susceptibility with respect to carbon content, Fig. 5 illustrates the results of TEM analysis for the nano-sized precipitation at the crack tip of steel C. Through the TEM micrograph and selected area electron diffraction (SAED) patterns, it was confirmed that the particles formed along the crack tip were eutectic M3C carbide45). Ishida et al. confirmed that M₃C carbide can form at elevated temperatures in various Fe-Mn-Al-C alloys46). It was also reported that the formation of M₃C carbide affects the mechanical properties of Fe-Mn- Al-C lightweight steels47). It is to say that the M3C carbide at the dendritic and grain boundaries formed due to the segregation of austenite stabilizing elements such as Mn and C during cooling thermal cycle, since the tested alloys had F-A solidification mode. As the segregation and precipitation mechanisms are inevitably accelerated by the carbon content, the temperature range of solidification during welding expands thereby cracking susceptibility increment. Lee et al. showed that the eutectic M3C precipitation causes increasing the temperature range of mushy zone, where the solid and liquid phases coexist because of its low melting point48). Similar experimental results were reported in the austenitic high manganese steels, increased solidification cracking behavior because of M3C eutectic phase formation at the boundaries in the welding process25). Therefore, considering the drastic increase of solidification cracking susceptibility from steel B to steel C, it is reasonable to conclude that the formation of eutectic phase in the fusion zone should be avoided in designing Fe-Mn-Al-C austenitic lightweight steels.
Fig. 5
(a) TEM micrograph of grain boundary precipitation obtained by FIB extraction with a SAED pattern confirming eutectic M3C phase, (b) TEM micrograph of fusion zone with a SAED pattern confirming the austenite matrix and κ-carbide precipitation
Based on the experimental results, we found that carbon content possesses significant influence on solidification crack susceptibility in the Fe-30Mn-10.5Al austenitic lightweight steels. The crack propagation mechanism can be explained by the transition in fusion zone microstructure and presence of nano-sized eutectic phase, which are caused by increasing carbon content. However, while the excessive carbon content caused negative influence on the solidification crack susceptibility, all the experimented alloys exhibited excellent hot ductility characteristics. Therefore, to improve the usability of austenitic lightweight steels in industrial fields, it is suggested that further investigation showing solidification crack susceptibility should be conducted as researchers continue to develop various types of austenitic lightweight steels with different alloying elements and manufacturing methods49-51).
4. Conclusions
In the current study, the solidification cracking susceptibilities of Fe-30Mn-10.5Al-(0.7-1.1)C alloys were evaluated via longitudinal Varestraint test. We found several insights from the experiments, summarized as follows:
1) The experimental results confirmed that the Fe- 30Mn-10.5Al-xC austenitic lightweight steels commonly had a great resistance against solidification crack susceptibility in the Varestraint test, compared to austenitic stainless steels.
2) While the generation and propagation of cracks were suppressed by the high content of Al, microstructural analysis confirmed that the increasing carbon content directly affected solidification crack susceptibility, causing deterioration of δ-ferrite phase and increasing MCL.
3) The reason of significantly accelerated crack propagation in the steel with 1.1 wt.% C was explained by the formation of eutectic (Fe, Mn)3C phase at the grain boundary, caused continuous vulnerable grain boundary at the final stage of solidification.
4) While κ-carbide particles were observed in the fusion zone microstructure, it had only negligible influence on solidification crack susceptibility due to its precipitation temperature range. However, the precipitates might have an additional effect on the product because of their mechanical characteristics.
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