Study on the Deformation Behavior and Mechanical Properties of Lightweight Economic Stainless Steels with Varying Al and Mn Contents

Abstract

In order to reduce the density and alloy cost of austenitic stainless steel, this study designed Fe-0.35C-12Cr-5Ni-(0,2,4)Al-(6,10)Mn (wt.%) stainless steels with different Al and Mn contents. The effects of Al and Mn contents on the microstructure, deformation behavior, and mechanical properties were investigated using microstructural analyses, quasi-static tensile tests, and Charpy impact tests. The results showed that an increase in Al content led to the formation of austeniteferrite duplex microstructure, while an increase in Mn content reduced the ferrite fraction. In the Al-free steel, the deformation mechanism was deformation-induced α′-martensitic transformation. When the Al content increased to 2 wt.%, the deformation mechanism was primarily mechanical twinning due to the increased stacking fault energy caused by Al. This resulted in a lower tensile strength but better toughness. When the Al content was further increased to 4 wt.%, the proportion of mechanical twinning decreased. The presence of ferrite led to cleavage at the fracture surface. The cleavage fracture explained the low elongation and toughness of duplex stainless steels. However, the elongation and toughness were enhanced with the increase in Mn content.

1. Introduction

Stainless steels are widely used in many industries, especially in the transportation sector [1,2,3,4], thanks to their excellent corrosion resistance [5]. However, their high alloy content leads to a relatively higher density, which in turn results in heavier components. This not only increases energy consumption, but also carbon emissions. With the continuous introduction of environmental protection policies, the high density of stainless steels has become a significant challenge in the efforts to reduce energy consumption and carbon dioxide emissions. Therefore, it is essential to develop lighter stainless steels. At present, the main strategy to achieve a lightweight material is to reduce material thickness [6]. However, this approach is limited by safety requirements, which restrict the extent of the lightweight characteristic [7]. Consequently, it is of great importance to optimize the chemical composition of stainless steel by adding lightweight elements, thereby achieving effective weight reduction while maintaining essential properties.

A considerable amount of studies [8,9,10,11] have been conducted on low-density Fe-Mn-Al-C steels, which achieved a weight reduction by adding sufficient Al. Similarly, there have been studies that have reduced the density of stainless steels through adding Al [12,13,14]. For austenitic stainless steels, which are characterized by their low stacking fault energy (SFE), deformation-induced martensitic transformation (DIMT) is widely recognized as the primary deformation mechanism at room temperature [13,15,16]. However, the addition of Al increases the SFE, thereby altering the deformation mechanism at room temperature. For instance, in the authors’ previous research [17] on Fe-0.3C-3Mn-14Cr-9Ni-(0,4)Al (wt.%) austenitic stainless steel, DIMT was identified as the dominant deformation mechanism. However, the addition of 4 wt.% Al significantly suppressed DIMT. Further, as Al is regarded as a ferrite stabilizer, it causes the microstructure to transition from a fully austenitic to a duplex austeniteferrite structure [12,18]. This is a critical factor influencing the mechanical properties of the material [19]. For Fe-Mn-Al-C low-density steels, Chen et al. [20] reported that the addition of 3–5 wt.% Al did not lead to ferrite formation in Fe-0.6C-0.5Si-18Mn (wt.%) steel. However, an excessive Al content (8 wt.%) induced the appearance of ferrite. Similarly, for austenitic stainless steels, Scherbring et al. [18] investigated the effects of Al addition ranging from 0 wt.% to 6 wt.% on the microstructure and properties of Fe-0.3C-12Cr-9Ni-3Mn (wt.%) austenitic stainless steel. The results indicate that excessive Al addition (4.5 wt.%) led to a duplex microstructure, with the presence of ferrite enhancing the material’s strength but reducing its ductility. Additionally, the addition of austenite-stabilizing elements, such as Mn, could reduce the potential of ferrite. Khorrami et al. [21] observed a positive correlation between Mn content (4–8 wt.%) and austenite phase fraction in Fe-20Cr-0.05Ni-0.02C-0.3N-0.3Si-Mn (wt.%) duplex stainless steels and demonstrated that a higher Mn content enhanced austenite stability and improve the ductility. Meanwhile, Mn also influences the deformation mechanism of austenite by affecting the SFE.

Most of the existing studies are focused on the separate effect of Al and Mn content on the microstructure and mechanical properties. In order to achieve a desired balance of mechanical properties, it is of great importance to accurately control the microstructure in stainless steel. Therefore, it is necessary to study the microstructure evolution and mechanical properties of stainless steels with varying Al and Mn contents. In the present study, the influences of Al and Mn on the microstructure of as-forged stainless steels were investigated, followed by quasi-static tensile tests and impact tests. The deformed microstructure was analyzed to clarify the deformation mechanism. This study aims to provide a theoretical basis for the research and development of lightweight Al-containing stainless steels.

2. Materials and Methods

To study the influence of different Mn and Al content on the microstructure and mechanical properties of lightweight stainless steels, four stainless steels with different alloying compositions, as presented in

Table 1. Alloying compositions of the experimental steels (wt.%).

Steel ID	C	Al	Mn	Cr	Ni	Fe
6Mn0Al	0.35	0.01	5.81	11.58	4.82	Bal.
6Mn2Al	0.34	2.12	6.17	12.19	4.82	Bal.
6Mn4Al	0.32	3.70	5.80	12.10	5.01	Bal.
10Mn4Al	0.35	4.24	10.07	11.72	4.89	Bal.

, were adopted for the experimental steels. The experimental steels were melted in a 15 kg induction furnace and cast into ingots with a height of 350 mm, a top diameter of 180 mm, and a bottom diameter of 130 mm. The ingots were then reheated to 1250 °C and held for 2 h for homogenization, followed by cooling to a start temperature of 1100 °C for forging. The final forging temperature was 950 °C. The ingots were forged into round bars with a diameter of 60 mm and finally cooled to room temperature in the furnace.

To analyze the microstructure after forging, the experimental steels were mechanically ground and polished, followed by etching using Beraha II etchant, which consists of 100 mL water + 50 mL concentrated hydrochloric acid + 6 g ammonium bifluoride + 1 g potassium metabisulfite. The etched microstructure was observed on a ZEISS optical microscope (OM, Carl Zeiss AG, Oberkochen, Germany). The equilibrium phase diagram of the experimental steels was additionally calculated using Thermo-Calc 2023b thermodynamic software with the TCFE 12: Steel/Fe-Alloys v12.0 database.

In order to test the mechanical properties of the experimental steels, quasi-static tensile tests and impact tests were carried out. According to GB/T 228.1-2021 standard [22], the experimental steels were machined to flat samples with a parallel length of 32 mm, width of 6 mm, and thickness of 1.2 mm for the tensile tests. Tensile tests were conducted at room temperature using flat tensile samples on an INSTRON-3382 (Instron, Norwood, MA, USA) electronic universal tester. Tensile tests were conducted at a strain rate of 10⁻³ s⁻¹ until fracture. For each experimental steel, the tensile test was repeated three times. The Charpy impact experiments were conducted on a PST752H-4 WANCE (WanTest Testing Equipment Coporation, Shenzhen, China) impact tester. According to the ASTM E23 standard [23], samples with a size of 10 mm × 10 mm × 55 mm with a 2 mm depth V-notch were used for Charpy impact experiments. Three impact tests were carried out for four experimental steels. The deflection−load curves were obtained using an impact tester.

In order to analyze the deformed microstructure and deformation mechanism, the deformed microstructure of the uniformly deformed section (the area not affected by necking) for tensile samples was analyzed. After mechanical polishing, the samples were vibrationally polished in a silica suspension to eliminate the deformed layer caused by mechanical polishing. Deformed microstructure analyses were performed using the electron backscatter diffraction (EBSD) method on an Apreo S HiVac SEM (Thermo Fisher Inc., Waltham, MA, USA). The step size for EBSD analyses was 0.7 μm. For impact test samples, the cross-sectional microstructure near the fracture was analyzed by EBSD. The adopted step size for EBSD analyses was 0.2 μm. The fracture surface after tensile tests and Charpy impact tests were additionally observed using a Nova 400 Nano SEM (Thermo Fisher Scientific Inc., Waltham, MA, USA).

3. Results

3.1. Initial Microstructure

The densities were measured using water displacement method to be 7.68 g/cm³, 7.45 g/cm³, 7.24 g/cm³, and 7.20 g/cm³ for 6Mn0Al, 6Mn2Al, 6Mn4Al, and 10Mn4Al steels, respectively. This indicates that adding every 1 wt.% Al and 1 wt.% Mn reduced the density by about 0.11 g/cm³ and 0.0096 g/cm³. Figure 1 presents the as-forged microstructure of four experimental steels. As shown in Figure 1a, the initial microstructure of 6Mn0Al steel consists of equiaxed austenite grains, with carbides distributing due to the high C content (0.35 wt.%). According to the equilibrium phase diagram of 6Mn0Al steel calculated by Thermo-calc software in Figure 2a, the carbides are likely to be M₂₃C₆-type carbides, which not only precipitate directly from austenite, but also transform from M₇C₃ carbides [24]. The OM micrograph in Figure 1b indicated a fully austenitic structure with carbide. Additionally, sub-grain boundaries can be observed in the microstructure. In sharp contrast, the OM micrographs of 6Mn4Al and 10Mn4Al steels (Figure 1c,d) show a mixed microstructure consisting of austenite, ferrite, and carbides. By using ImageProPlus 6.0 software, the ferrite fractions of 6Mn4Al and 10Mn4Al steels were estimated to be about 7.25% and 3.36%, respectively. As indicated by the equilibrium phase diagram in Figure 2a, ferrite potential increases with Al content, resulting in the formation of ferrite in the steels with 4 wt.% Al addition. Further, because Mn increases the stability of austenite (Figure 2b), the higher Mn addition of 10Mn4Al steel contributes to a lower ferrite fraction compared with 6Mn4Al steel.

According to the author’s previous researches [12,13], in molten Al-containing austenitic stainless steel (Fe-0.3C-3.29Mn-13.74Cr-9.11Ni-3.52Al, in wt.%), dendritic segregation was observed significantly, which was not the case for the Al-free steel (Fe-0.3C-3.39Mn-14.09Cr-9.52Ni, in wt.%). This is attributed to the fact that Al altered the austenite cooling mode (L → L + γ → γ) to a ferrite cooling mode (L → L + δ → L + δ + γ → L + γ → γ) [12,25]. Although the experimental steel subsequently underwent a thermomechanical process, it is very likely that the microstructure still inherited this local segregation of alloying elements [12]. Further, the low solubility of Al may also lead to a local segregation.

3.2. Tensile Tests

3.2.1. Deformed Microstructure

In order to compare the influences of Al and Mn content on the mechanical properties, tensile tests were conducted for the experimental steels at room temperature until fracture. Figure 3 presents the EBSD inverse pole figure (IPF) maps and phase maps for the deformed microstructure of four experimental steels. The initial microstructure of 6Mn0Al and 6Mn2Al steels is fully austenitic (Figure 1a,b). The EBSD analyses for the deformed microstructure suggested that both steels exhibited a typical characteristic of planar glide. After tensile deformation until fracture, both DIM and mechanical twins formed in the deformed microstructure of 6Mn0Al steel (as shown in Figure 3a,e), indicating the occurrence of transformation-induced plasticity (TRIP) and twinning-induced plasticity (TWIP) effects during tensile deformation. For 6Mn2Al steel, due to the increasing SFE by 2 wt.% Al addition, DIMT was effectively suppressed while mechanical twinning became the dominant deformation mechanism. Consequently, the fraction of mechanical twins in the deformed microstructure was significantly elevated, as indicated in Figure 3b,f. Compared with the two austenitic stainless steels (6Mn0Al and 6Mn2Al), the SFE of two austenite−ferrite duplex steels (6Mn4Al and 10Mn4Al) are higher due to the addition of higher Al content (4 wt.%), thereby inhibiting mechanical twinning (Figure 3c). Additionally, the higher Mn concentration of 10Mn4Al steel compared with 6Mn4Al steel further increased SFE, which further inhibited the occurrence of mechanical twinning [26]. Moreover, due to the pre-existing ferrite in both duplex steels, ferrite elongated along the tensile direction was observed in the deformed microstructure, as presented in Figure 3g,h.

Figure 4 presents the tensile fracture surfaces of four experimental steels obtained by SEM. Due to the occurrences of DIMT and mechanical twinning in 6Mn0Al and 6Mn2Al steels during tensile deformation, both experimental steels exhibit ductile fractures, which are characterized by the presence of dimples (Figure 4a,b). In contrast, Figure 4c,d show the coexistence of dimples and cleavage. This indicates that in 6Mn4Al and 10Mn4Al steels, the ferrite fractured first due to lower capacity for plastic deformation. Subsequently, stress was redistributed and the austenite underwent plastic deformation until fractures occurred.

3.2.2. StressStrain Curve Analyses

Figure 5 presents the engineering stressstrain curves, true stressstrain curves, and work hardening rate curves obtained by tensile tests. The mechanical properties, including yield strength (YS), ultimate tensile strength (UTS), uniform elongation (UE), total elongation (TE), and product of strength and elongation (PSE) are summarized in

Table 2. Mechanical properties of the experimental steels. The errors represent the mean difference.

Steel ID	YS/MPa	UTS/MPa	UE/%	TE/%	PSE/MPa·%
6Mn0Al	552 ± 8	1102 ± 12	38 ± 3	52 ± 3	57,946
6Mn2Al	512 ± 15	786 ± 20	18 ± 5	54 ± 2	43,480
6Mn4Al	626 ± 14	844 ± 21	15 ± 6	32 ± 3	25,964
10Mn4Al	554 ± 10	722 ± 16	16 ± 5	38 ± 4	27,238

.

The 6Mn0Al steel exhibited DIMT and mechanical twinning during tensile deformation. The deformed microstructures such as DIM and mechanical twins, which are the products of planar glide, can act as strong barriers for the slipping of dislocations during plastic deformation [27,28]. This enhances the work hardening rate (Figure 5b), resulting in the highest UTS (Table 2). Additionally, the relatively high level of work hardening is conducive to resisting the occurrence of necking. Therefore, the necking point of 6Mn0Al steel, where σtrue is equal to dσ_true/dε_true, reaches a large strain (Figure 5b), leading to a larger UE (Table 2). For 6Mn2Al steel, due to the absence of DIMT, the work hardening rate and UTS are lower than that of 6Mn0Al steel. Although there are studies suggesting the contribution of mechanical twins to work hardening in TWIP steels [29,30], the work hardening behavior is weaker than that caused by the TRIP effect. Meanwhile, the necking presents at a relatively small strain, resulting in an obviously lower UE. Further, for duplex steels (6Mn4Al and 10Mn4Al), the presence of ferrite significantly reduces the elongation compared to the austenitic steels. For 10Mn4Al steel with a higher Mn concentration, the ductility is improved significantly due to less ferrite compared to 6Mn4Al steel. However, because the fraction of harder ferrite structure decreased, the work hardening rate and UTS of 10Mn4Al steel decreased. Regarding the PSE of experimental steels, it is found that when the Mn content is constant, PSE decreases with the increasing Al content because of the weaker work hardening of 6Mn2Al steel and the cleavage occurrence of 6Mn4Al steel. However, as the Mn content increases, the decreasing ferrite fraction optimizes the fracture mode, resulting in an increase in PSE.

3.3. Impact Toughness Tests

3.3.1. LoadDeflection Curve Analyses

Figure 6 illustrates the load-deflection curves during Charpy impact tests. These curves were subsequently analyzed using the “ISO evaluation” procedure introduced in refs. [31,32] to identify characteristic events and their corresponding values, such as the general yield load (F_gy), the maximum Charpy impact load (F_m), the crack initiation load (F_in), the start of unstable crack growth load (F_iu), and the end of unstable crack growth load (F_a). F_gy marks the transition from the elastic to the plastic deformation stage [10,32]. As shown in Figure 6, the changes in F_gy values for four experimental steels are in good agreement with the changes in YS by tensile tests. The F_m value corresponding to the peak load reflects the capacity of the experimental steels withstanding impact loads, which varies with the Al and Mn contents. The determination of F_in is based on the compliance changing rate (CCR) method described in Ref. [32]:

△C/C = (C − C_el)/C_el

(1)

where C and C_el represent compliance and elastic compliance, respectively. The units are mm/N. C and C_el can be calculated using Equations (2) and (3), respectively:

C = ds/dF

(2)

C = s_el/F_el

(3)

where s and F represent the deflection and load of the load−deflection curves, respectively. The elastic compliance C_el is defined as the inverse of the slope of the linear elastic segment of the load−deflection curve. Crack initiation leads to a slope change in CCR, where the corresponding load is referred to as F_in. F_iu is corresponds to the deflection where the slope decreases rapidly.

The impact absorbed energy of the experiment steels was calculated based on the curves obtained from the “ISO evaluation”, as shown in Figure 6e. W_all represents the total impact absorbed energy, which includes three parts: Crack initiation energy (W_initiation), stable crack growth energy (W_stable) and unstable crack growth energy (W_unstable). For austenitic steels (6Mn0Al and 6Mn2Al), the W_initiation is almost the same, but it is significantly lower for the duplex steels. This is because ferrite is very likely to become crack initiation sites, leading to the premature formation of cracks. Due to the less ferrite in 10Mn4Al steels, the W_initiation is larger than that of 6Mn4Al steel. Regarding the crack growth energy (W_stable + W_unstable), 6Mn2Al steel exhibits a substantially higher value than 6Mn0Al steel. For 6Mn0Al steel, DIMT occurred during impact deformation. On the one hand, the TRIP effect could absorb the impact energy, thereby enhancing toughness [33]. On the other hand, the formation of brittle DIM provides pathways for crack propagation, thus reducing the crack growth energy [34]. Different from 6Mn0Al steel, the TWIP effect in 6Mn2Al steel could absorb impact energy. Additionally, the presence of mechanical twin boundaries could impede crack propagation, thereby increasing the crack growth energy. Owing to its significantly higher crack growth energy, 6Mn2Al steel exhibits a superior impact toughness. For duplex steels, the crack growth energy is much lower due to the pre-existing ferrite. The γ/α phase boundaries not only provide nucleation sites, but also propagation pathways for cracks [35], therefore resulting in a low crack growth energy.

3.3.2. Impact Fracture Surface

Figure 7 presents the SEM micrographs for the impact fracture surface of the experimental steels. The 6Mn0Al and 6Mn2Al steels display high-density dimples (Figure 7a,b), which are characteristic of ductile fracture. The quasi-cleavage observed in Figure 7a is very likely to be attributed to the presence of DIM [34]. In sharp contrast, both dimples and cleavages are evident on the fracture surface of 6Mn4Al and 10Mn4Al steels, as shown in Figure 7c,d. The dimples correspond to the austenite regions, while the cleavage regions are associated with the ferrite phase in the duplex steels [36]. This explains the lower impact absorbed energy of duplex steels.

4. Conclusions

The mechanical properties of lightweight economic stainless steels with varying Al and Mn contents were studied using quasi-static tensile tests and Charpy impact tests, followed by microstructure analyses. The following conclusions can be obtained:

(1)

A high C content led to the formation of carbides. Increasing the Al content led to the formation of the austenite−ferrite duplex microstructure. Increasing the Mn content reduced the ferrite fraction.

(2)

The occurrence of DIMT and mechanical twinning improved the work hardening rate, leading to high UTS and UE. The addition of 2 wt.% Al suppressed DIMT but promoted mechanical twinning by increasing SFE, resulting in a lower work hardening rate, UTS, and UE. As Al content increased to 4 wt.%, the ferrite induced cleavage, deteriorating elongation. Increasing the Mn content improved ductility by reducing the ferrite fraction.

(3)

The 6Mn0Al steel exhibited a lower crack growth energy compared to 6Mn2Al steel due to the formation of brittle DIM. TRIP still provided a high toughness for 6Mn0Al. The high toughness and crack growth energy of 6Mn2Al were attributed to the formation of mechanical twins. The occurrence of ferrite in duplex steels resulted in a low crack initial energy and crack growth energy, consequently a low toughness. The higher Mn content increased toughness by reducing the ferrite fraction.