Influence of C content and annealing temperature on the microstructures and tensile properties of Fe–13Mn–8Al–(0.7, 1.2)C steels

Abstract

In the present work, the Fe-13Mn-8Al-(0.7, 1.2)C steels were subjected to different post-cold rolling annealing treatments in a temperature range of 800–1000 °C for 15 min to investigate the microstructural evolution and the mechanical properties. Both steels annealed at 800 °C consisted of fine austenite and ferrite grains along with intergranular κ-carbides formed by eutectoid reaction, which contributed to the high strength but deteriorated ductility. With annealing temperature increasing, the volume fraction of ferrite and intergranular κ-carbides progressively decreased. Meanwhile, the precipitation of intragranular κ′-carbides enhanced the yield strength, whereas coarse austenite was harmful to both strength and elongation. The increase in C content increased the volume fraction of intergranular κ-carbides and austenite as well as retarded the austenite recrystallization, which further increased strength and strain hardening rate but deteriorate ductility. The deformation mechanism in austenite was massive planar slip. The microbands were widely observed in both steels at an annealing temperature above 900 °C.

Introduction

Due to the needs of safety, cost reduction and energy saving, many efforts have been made to reduce the weight of automobile and consequently decrease the fuel consumption and greenhouse gas emission [[1], [2], [3], [4]]. Among various steel categories, Fe-Mn-Al-C steels have received much attention due to the combination of low density and excellent mechanical properties [[5], [6], [7], [8], [9], [10]].

Fe-Mn-Al-C steels can be classified into three categories by their phase constitution: ferrite based duplex steels [[11], [12], [13]], austenite based duplex steels [[14], [15], [16]] and austenitic steels [9,10,17]. The mechanical properties of Fe-Mn-Al-C steels are not only controlled by the phase constitution but also significantly influenced by the morphology and volume fraction of κ-carbides. κ-carbide, which has a perovskite crystal structure, can form in the grain interior and at the grain boundaries in an isothermal temperature range from 450 to 1000 °C in Fe-Mn-Al-C steels [[18], [19], [20], [21]]. The intragranular κ′-carbides can increase yield strength as well as ultimate tensile strength of the materials while the good ductility is preserved [3,22,23]. However, these precipitates can be sheared by accumulated dislocations during deformation, which reduces the strain hardening rate [24]. The intergranular κ-carbides were reported to contribute to the strength by secondary phase strengthening but deteriorate the ductility [25].

The increase in C content would significantly increase the cold-workability and ductility of the alloys due to the decrease in the ferrite fraction [26]. Meanwhile, the high hardness and strength are obtained through the precipitation of intragranular κ′-carbides. However, the excessive addition of C may lead to the deterioration of ductility due to the formation of intergranular κ-carbides [27].

To overcome the disadvantage in the fabrication, forming and welding of high-Mn low density steels, medium-Mn duplex (γ+α) or triplex (γ+α+κ) steels were recently proposed. Cheng et al. [28] investigated the phase transformation of Fe-13.5Mn-6.3Al-0.78C and found that no grain boundary precipitates were observed in austenite at the temperatures above 800 °C. With a decrease in the isothermal temperature, ferrite and intergranular κ-carbides were found in the microstructure. At the temperatures below 650 °C, the austenite was decomposed by a eutectoid reaction into lamellar phases (i.e., γ → α + κ + M₂₃C₆). The intragranular κ′-carbides were observed in Fe-11Mn-10Al-1.25C steel under 1000 °C annealing condition [18]. By controlling the fraction of phase and precipitates, a wide range of mechanical properties can be achieved, showing that the yield strength, ultimate strength and elongation were in a range of 495- 1140 MPa, 610-1210 MPa and 5-46%, respectively [14,18,25].

In view of the aforementioned findings, the mechanical properties of Fe-Mn-Al-C steels are closely related to the phase constituents and the precipitates, which can be controlled by annealing temperatures. Thus, in the present work, a novel medium-Mn system with a composition of Fe-13Mn-8Al-(0.7, 1.2)C was designed to investigate the microstructural evolution, the mechanical properties and the effect of carbon content on them. This work would provide theoretical guide for designing the alloy compositions as well as controlling the mechanical properties in medium Mn low density steels.

The chemical composition of Fe-Mn-Al-C low density steels in the present study was listed in Table 1. The densities of 0.7C and 1.2C steels were calculated [26] as 7.13 g/cm³ and 7.03 g/cm³, which were 9.3% and 10.6% lower than a typical TWIP steel with a composition of Fe-22Mn-0.6C (in wt.%), respectively.

50 kg ingots were prepared by an induction melting under an argon atmosphere. The ingots were initially subjected to the homogenization treatment at 1200 °C for 3 h. Then, they were hot rolled for multiple passes in a temperature range of 1150-850 °C to a thickness of 4 mm and subsequently water quenched. The hot-rolled strips were further annealed at 1000 °C for 1 h and water quenched. Afterwards, they were cold rolled to a thickness of 1 mm.

To determine the most appropriate annealing temperature region, the equilibrium phase diagrams were calculated using the Thermo-Calc [29] software along with TCFE9 Database. According to Fig. 1, Fig. 12, 1.2C steel is expected to exhibit a dual-phase microstructure (consisting of austenite and ferrite) in a temperature range of 800-900 °C, while 0.7C steel appears to be in a fully austenitic state. Based on these calculations, the cold-rolled strips were subjected to the annealing treatment within the temperature range from 800 °C to 1000 °C to obtain diverse microstructures. The annealing time of Fe-Mn-Al-C low density steels is usually in the range of 10–30 min [5,7,20,24,[30], [31], [32]]. According to our previous work [32], the highest UTS × EL in Fe-12Mn-8Al-0.8C steel was achieved at an annealing temperature of 900 °C. Tensile testing was conducted on both 0.7C and 1.2C steels annealed at 900 °C for 5, 15, 30 and 60 min. Best mechanical properties of two steels are both achieved at an annealing time of 15 min (see Supplementary Data). The annealed specimens were cooled at different speeds, including water quenching, air cooling and furnace cooling to determine the proper cooling rate. The results revealed that with the decrease in cooling speed, the grain size and volume fraction of intergranular κ-carbides increased (see Supplementary Data), which would deteriorate the ductility of Fe-Mn-Al-C steels. Therefore, the cold-rolled strips were annealed for 15 min and subsequently water quenched.

The microstructure characterization was carried out by using scanning electron microscope (SEM), electron back-scattered diffraction (EBSD) and transmission electron microscope (TEM). The SEM specimens were electro-polished at room temperature in a solution of 10% perchloric acid and 90% ethanol at an operating voltage of 25 V and then etched by picric acid. The SEM was performed using a Zeiss UltraPlus FE-SEM equipped with energy dispersive X-ray spectroscopy (EDX). The EBSD samples were prepared by a standard grinding and mechanical polishing operation including a final polishing with a 0.04 μm colloidal silica suspension. The EBSD was conducted using a Zeiss LEO 1530 FEG SEM operated at 20 kV. The step size of 0.05 μm was used to acquire the EBSD data and HKL Channel 5 software was utilized for the EBSD data post-processing. Thin foils for TEM were mechanically ground to a thickness of about 0.06 mm and then twin-jet electropolished using a solution of 10% perchloric acid and 90% ethanol at a temperature of −25 °C and a voltage of 31 V. TEM analysis was performed in a Tecnai G² 20 TEM and a JEOL JEM 2100F TEM at an acceleration voltage of 200 kV.

The tensile specimen had flat dog-bone geometry with a 25 mm gauge length, 12.5 mm width and 1 mm thickness. The tensile axis was parallel to the rolling direction. Tensile tests were carried out at room temperature using a universal testing machine (SANSCMT5000) at a constant crosshead displacement velocity (equivalent to a strain rate of 1.33 × 10⁻³ s⁻¹).

The volume fraction of coarse intergraular κ-carbides was measured by image analysis technique. The volume fraction of austenite and ferrite was determined by X-ray diffraction (XRD) with CuK_α radiation using the direct comparison method. For specimens with intergranular κ-carbides distributed along the grain boundaries, the volume fraction of austenite, V_γ, was calculated by the following equation [33]:

$$V_{\text{r}}=\left(1-V_{\text{c}}\right)\left[\left(\frac{1}{q}\sum _{\text{j}=1}^{\text{q}}\left(I_{\mathrm{\gamma j}}/R_{\mathrm{\gamma j}}\right)\right)\right]/\left[\frac{1}{p}\sum _{\text{i}=1}^{\text{p}}\left(I_{\mathrm{\alpha i}}/R_{\mathrm{\alpha i}}\right)+\frac{1}{q}\sum _{\text{j}=1}^{\text{q}}\left(I_{\mathrm{\gamma j}}/R_{\mathrm{\gamma j}}\right)\right]$$

Where V_C is the volume fraction of intergraular κ-carbide, R is proportional to the theoretical integrated intensity. I_γ and I_α are the integrated intensity of austenite diffraction lines and ferrite diffraction lines, respectively. Here, the integrated intensities of the (110)_α, (200)_α and (211)_α peaks and those of the (111)_γ, (200)_γ, (220)_γ and (311)_γ peaks were used.