Effects of Intercritical Annealing Temperature on the Tensile Behavior of Cold Rolled 7Mn Steel and the Constitutive Modeling

Abstract

Medium Mn steel is composed of sub-micron grained ferrite and austenite, the unstable austenite may transform to martensite during plastic straining. Although the mechanical properties of medium Mn steel could be easily tested by tensile test, it is quite difficult to directly measure the influences of different constituent phases on the tensile and work hardening behavior. Thus, at the present work, EBSD, TEM, XRD and a constitutive model based on dislocation density have been used to study the effects of intercritical annealing (IA) temperature on the tensile properties and work hardening behavior of a newly designed medium Mn steel, Fe-7%Mn-0.3%C-2%Al (mass fraction). Experimental results showed that with the increase of IA temperature, the mechanic stability of reverted austenite decreased gradually and the kinetics of strain induced martensite rose rapidly. The stability of the reverted austenite was moderate when intercritically annealed at 700 ℃, this led to the best plasticity and the optimal mechanical properties. Simulated results exhibited that the mechanic stability of austenite has a decisive influence on the tensile behavior of the material. The austenite stability will be too high if the IA temperature is lower, and this will lead to the lower work hardening rate and uniform elongation; when the IA temperature is moderate, the stability of austenite will be optimum, consequently strain-induced martensite would be progressively produced during straining and result in the higher work hardening rate and prolonged uniform elongation; the stability of austenite will be too lower if the IA temperature is higher, thus larger volume fraction of strain-induced martensite would be formed in a short period, and this would result in the higher tensile strength but the inferior uniform elongation.

Key words: medium Mn steel    austenite stability    TRIP effect
YANG Feng1, LUO Haiwen2, DONG Han3

1 Central Iron and Steel Research Institute, Beijing 100081, China
2 School of Metallurgical and Ecological Engineering, University of Science and Technology Beijing, Beijing 100083, China

3 School of Materials Science and Engineering, Shanghai University, Shanghai 200072, China


With the continuous improvement of the energy-saving, environmental protection and safety requirements of the automotive industry, high-strength-plastic automotive structural steel has increasingly become the focus of attention. In the past two or three decades, steels for automobiles have developed rapidly. American researchers have classified automotive steels into three generations based on the strong plastic product (namely, the product of tensile strength and elongation after breaking) of automotive steels [1]. The first generation automobile steels are represented by gap-free atomic steel, dual-phase steel and low alloy phase transformation induced plasticity (TRIP), and the strong plastic product is 10-20 GPa. It is difficult to meet the automotive industry at present. The dual requirements of lightweight and high security. The second-generation automotive steels are represented by twinning induced plasticity (TWIP) austenitic steels and TWIP steels. The strong plastic products reach 50 to 70 GPa%, but a large amount of Cr, Ni, Mn, etc. are added. Elements, its high cost and smelting production there are some difficulties. The Mn content in the third-generation automotive steels developed in recent years is about 3.5% to 12.0% (mass fraction, the same applies hereinafter), and the use of reverse-transformed austenite in the annealing process of hot-rolled or cold-rolled steel sheets to form submicron Grade austenitic and ferritic duplex structure, TRIP effect of austenite in the strain process to improve the plasticity and strength of steel, its excellent comprehensive mechanical properties can not only meet the requirements of automotive lightweight and collision safety, It also ensures the formability of automotive parts.
At present, the research of medium manganese steel has progressed from the early Fe-Mn-C system [2,3,4,5] to the Fe-Mn-C-Al(-Si) system [6,7,8,9,10] ]. Reasonable alloy composition design can greatly improve the comprehensive mechanical properties of medium manganese steel. In [11], the authors discussed the effect of annealing temperature on the microstructure and mechanical properties of 7%Mn-0.3%C-2%Al (hereinafter referred to as cold-rolled 7Mn steel) in a new type of cold-rolled manganese steel. The experimental results show that as the annealing temperature increases, the grain size of the reversed austenite gradually increases, and the content of solid solution Mn and C gradually decreases, so that the mechanical stability of austenite significantly decreases. Tensile experiments show that austenite with moderate mechanical stability can be obtained after cold rolling 7Mn steel annealed at 700 °C for 1 h, and it is transformed into martensite during deformation to optimize the mechanical properties of the material. Although the experimental results show that the mechanical properties of cold-rolled 7Mn steel are mainly affected by the austenite stability, it is still unclear whether the ferrite, austenite and strain-induced martensite are stretched and work hardened during the deformation process. The behavior and how these three phases affect the mechanical properties of the material, because these data are difficult to measure directly by experimental means. Sun Chaoyang et al [12] successfully simulated the true stress-true strain curve of Fe-22Mn-0.6C TWIP steel at different strain rates using a constitutive model based on dislocation density, and analyzed the twinning and slip mechanisms of TWIP steel. Interactions and effects on macroscopic deformation. The simulation data showed that there was a negative correlation between the twinning rate and the slip rate, ie the early stage of deformation, the twinning rate was larger and the slipping rate was smaller. When the twinning trend became saturated, the twinning rate decreased and the slip rate rose rapidly. Lee et al. [7,8] also used the constitutive model to analyze the flow stress and work hardening behavior of the various phases of the Fe-10Mn-0.3C-3Al-2Si type manganese steel during the deformation process, and simulated the true stress- The true strain curve and the work hardening rate curve agree well with the measured results. The simulation results also show that the effect of strengthening the TRIP effect in metastable austenite grains is more significant than that of the TWIP effect after Fe-10Mn-0.3C-3Al-2Si is cold rolled and annealed at 800 °C. The work of Sun Chaoyang et al [12] and Lee et al [7,8] shows that the constitutive model can be used to analyze in depth the mechanical behavior of TWIP steel or medium manganese steel during deformation. Therefore, this work will use the dislocation density-based constitutive model to study the tensile and work hardening behavior of cold-rolled 7Mn steel, and combine the experimental results to clarify the intrinsic mechanism of the difference in mechanical behavior of manganese steel after annealing at different temperatures.
1 Experimental method
7Mn steel is smelted by a vacuum induction furnace and its chemical composition (mass fraction, %) is: Mn 7, C 0.3, Al 2, Fe balance. The ingots are forged at 1200 °C for 2 h and then forged at 1200 °C to 850 °C. The furnace is then cooled to room temperature. The billet was heated to 1200 °C and held for 1 h. Hot rolling was performed between 1200 °C and 850 °C. After rolling, it was water-cooled to room temperature. The final thickness of the hot rolled sheet was about 4 mm. The hot rolled sheet was cold-rolled at 700°C for 30 minutes, and the cold rolling reduction was 70%. Subsequently, the cold rolled strips were kept at 680, 700, and 720 °C for 1 h and cooled to room temperature. For ease of description, the samples annealed at different temperatures are referred to as S680, S700, and S720, respectively.
Specimens used for electron backscatter diffraction (EBSD) observation on a Quanta 650 SEM were mechanically ground and electrolyzed using a 10% perchloric acid solution. The EDSD scanning step was 0.05 μm. Scan data were processed using HKL Channel5 software. Specimens used for H-800 transmission electron microscopy (TEM) observations were ground to 35 μm and then electro-polished with 6% perchloric alcohol solution at -20 °C for TEM observation of fine structures. Standard stretched specimens with parallel and segment lengths and widths of 50 and 10 mm, respectively, were subjected to room temperature tensile tests on a WDW-300E tensile machine with a chuck movement rate of 2 mm/min. X-ray diffraction (XRD) experiments were performed on a SmartLab diffractometer. The scanning area of the sample was 40° to 100° and the scanning rate was 1°/min. The sample was loaded to a predetermined amount of deformation and then unloaded to measure the austenite content and dislocation density at different deformation amounts; for the fractured specimen, a parallel section far from the fracture was sampled. The austenite content (Vγ) in the material is calculated using the following formula [13]:

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In the formula, Iγ is the average of the integrated intensity of the {200}, {220}, and {311} diffraction peaks of austenite, and Iα is the average of the integrated intensity of the diffraction peaks of ferrite {200} and {211}. The dislocation density of austenite was measured using the modified Williamson-Hall equation [14,15]. The measurement method is described in [16]. Jade6.0 software was used to peak-process XRD data and calculate the full width at half maximum and integral intensity of the diffraction peak.
2 experimental results
Figure 1 shows the EBSD images of samples annealed at different temperatures. The white region is ferrite, the gray region is austenite, the black solid line shows the large-angle grain boundary, and the gray solid line is the small-angle grain boundary. It can be seen that the ferrite and austenite grains are basically equiaxed after cold-rolling annealing, which is very different from the strip-like morphology after hot-rolled annealing[2,3]. In addition, most of the ferrite grain boundaries are small-angle grain boundaries, ie, ferrite still does not recrystallize after annealing. Cao et al. [2] believed that this is due to the higher Mn content in the ferrite grains that drags grain boundaries and inhibits grain boundary migration. The grain size of the reversed austenite and ferrite after annealing was counted, as shown in Table 1. It can be seen that as the annealing temperature increases, the average size of the reverse-transformed austenite grains increases while the ferrite average grain size decreases slightly. The chemical composition of austenite after annealing of cold-rolled 7Mn steel was measured in [11]. The data shows that the content of Mn and C in solid solution in austenite grains decreases as the annealing temperature increases. Since austenite grain size and composition determine its mechanical stability, the stability of austenite gradually decreases with increasing annealing temperature. Figure 2 shows the TEM images of S680~S720. It can be seen that although no recrystallization has occurred, the dislocation density inside the ferrite grains is still very low, and Han et al. [17] described this phenomenon as a large recovery.

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Fig.1 EBSD phase maps of S680 (a), S700 (b) and S720 (c) (The white phase is ferrite and the gray phase is austenite, the black lines are high-angle grain boundaries with misorientation angles of over 15°, the gray lines are low-angle grain boundaries with misorientation angles of 2°~15°)

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Fig.2 TEM images of S680 (a), S700 (b) and S720 (c) (γ denotes austenite grains and the rest are ferrite grains)

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Table 1 The average grain sizes of ferrite and austenite and parameters for calculating martensite volume fraction (VM)

Figure 3 shows the martensite content of S680~S720 at different deformations [16]. It can be seen that the transformation kinetics of martensite is related to the stability of austenite, and the lower the austenite stability, the faster the formation rate of martensite. In addition, the martensite content increases in a S-shape with the strain, so the transition kinetics of martensite is simulated using the Olson-Cohen model [18]:

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In the formula, VM and Vγ0 are respectively the martensite volume fraction and the initial austenite content; ε is the true strain; α, β and n are the parameters related to the martensitic transformation kinetics [18], the values are listed in the table 1. Figure 4 shows the true stress-true strain curve of S680~S720 at room temperature. It can be seen that as the annealing temperature increases, the yield strength of cold-rolled 7Mn steel gradually decreases, while the tensile strength monotonously increases. The elongation of the material increases first and then decreases with the annealing temperature. It reaches a maximum at 700°C, and the corresponding engineering strain is about 68%. At this time, the strong plastic product of the material also reaches an optimum, about 65 GPa%[11] , far higher than the Fe-Mn-C series of medium Mn steel, even without the loss of the general TWIP steel. The literature [11] compared the Fe-Mn-C and Fe-Mn-C-Al manganese steels. It is believed that the Al addition increases the stability of the reversed austenite, making it slow during deformation. The continuous transformation into martensite improves the material’s elongation. It can also be seen from Fig. 4 that the tensile curve of S680 is very smooth in the uniform deformation zone, and no stress sawing occurs; and S700 and S720 appear different types of stress saw teeth during plastic deformation. This indicates that the dynamic strain aging of the cold-rolled 7Mn steel was observed during the annealing process at 700 and 720 °C for 1 h. The influence of annealing temperature on the dynamic strain aging of manganese steel during cold rolling is discussed in detail in [16]. It is believed that the austenite grain size and stacking fault energy determine whether the stress saw teeth are generated during the austenizing process. Stability determines the type of stress saw tooth.

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Fig.3 VM of S680~S720 after deformed to various strains[16] and the corresponding fitted curves

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Fig.4 True stress-true strain curves of cold-rolled 7Mn steel after annealed at different temperatures

Figure 5 shows the TEM image of the S700 at different strain rates. When the true strain is 0.095, there is a large amount of stacking faults inside the austenite grains, as shown in Figure 5a. Selected area electron diffraction (SAED) shows that no martensite is produced at this time. When the strain increase to 0.35, a large amount of strain-induced martensite can be observed inside the austenite grains, as shown in Figure 5b, and the dark field image of austenite is shown in Figure 5c. It can be seen that the strain-induced lateral expansion of martensite gradually encroaches on the parent phase austenite. The SAED spectrum shows that there is twinning of martensite, which is consistent with the observations in [19].

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Fig.5 TEM images of deformed microstructures in S700 after the strain of 0.095 (a) and 0.35 (b) and the dark-field image of austenite in Fig.5b (c) (Insets show the SAED patterns of the circles in Figs.5a and b. SF—stacking fault, γ—austenite, M—strain induced martensite)

3 Constitutive model
The influence of annealing temperature on the microstructure and properties of cold-rolled 7Mn steels was discussed in the previous section. This section will introduce a constitutive model based on dislocation density to simulate the calculation of tensile and work hardening behaviors of 7Mn cold-rolled steels. Since the cold-rolled 7Mn steel is annealed by ferrite and metastable austenite 2 phases, and the metastable austenite is transformed into martensite during the deformation process, the flow stress and ferrite of the cold-rolled 7Mn steel are changed. The flow stress and volume fraction of the three phases of body, austenite, and strain-induced martensite are related. The volume fraction of ferrite remains unchanged during deformation, the volume fraction of martensite can be described by equation (2), and the content of austenite Vγ=Vγ0-VM. Therefore, the flow stress of cold rolled 7Mn steel can be expressed by the mixing law [7,8]:
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In the formula, σt, σf, σγ, and σM are total stress, ferrite stress, austenite stress, and martensite stress, respectively. Vf, Vγ, and VM are ferrite, austenite, and horse, respectively. The volume fraction of the spheroids.
The microstrain dεt of the material can also be expressed by the mixing law, ie:
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In the formula, dεf, dεγ, and dεM are the micro strains of ferrite, austenite, and martensite, respectively. Due to the different stresses of ferrite, austenite and martensite, the strain distribution of the 3 phases during the deformation process will also be different. Bouaziz et al. [20] considered that the distribution of stress and strain in multiphase materials is in accordance with the principle of equal power, that is, the work done by the applied loads during deformation is evenly distributed among the phases and can be expressed by the following equation:

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The simultaneous expressions (4) and (5) can be used to obtain the expression of each phase’s micro-strain.
The flow stress of each phase is calculated by the following formula:
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In the formula, the subscript i represents ferrite, austenite, or martensite, and the σsiσis, σgbiσigb, and σdisiσidis are solid solution strengthening terms, grain boundary strengthening terms, and dislocation strengthening terms, respectively. It is worth noting that when Jian et al. [21] studied low-level fault-tolerant magnesium alloys, it was found that the dislocation and fault-tolerant interactions occurred during the deformation of the material, resulting in an increase in material strength. Strictly speaking, when calculating the flow stress of metastable austenite, the influence of faults should also be considered. However, due to the fact that the effect of faults on the strength of the matrix has not been extensively studied and there is no mature physical model, the contribution of faults to flow stress is not considered in this work. Since the effect of solid solution elements and dislocation density is generally considered in the fitting of martensite strength, and the effect of martensite size is not considered, the grain boundary strengthening of martensite is also not considered in this work. The solid solution strengthening of ferrite, austenite and martensite can be represented by the following 3 formulas respectively [7, 22]:

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In the formula, XαCXCα, XαMnXMnα and XαAlXAlα are the mass fractions of C, Mn and Al in ferrite, XγCXCγ and XγMnXMnγ are the mass fractions of C and Mn in austenite, and XMCXCM is the C content in solid solution in martensite. The study in [11] shows that the content of Mn and Al in ferrite and austenite grains after annealing at 680~720 °C for 1 mm of cold-rolled 7Mn steel is similar to that calculated by ThermalCalc software, so it is used in simulation calculations. The balanced C, Mn, and Al contents account for the solid solution strengthening of ferrite, austenite, and martensite. Note that the C content of martensite is equal to that of austenite.
The grain boundary strengthening of ferrite can be expressed by the following formula:
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In the formula, Kα is the Hall-Petch coefficient of ferrite, and dα is the grain size of ferrite. Since austenite has strain-induced martensite at the time of deformation (as shown in Fig. 5c), they will separate the austenite grains and gradually refine the austenite during deformation, thus promoting the so-called austenite transformation. Dynamic Hall-Petch effect. Therefore, grain boundary strengthening of austenite can be expressed by formula (11):

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In the formula, K γ is the Hall-Petch coefficient of austenite, and L is the effective interfacial distance of austenite, which can be expressed by the following formula [23]:

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In the formula, dγ is the grain size of austenite, λM is the average spacing of martensite laths, and iM is the adjustment coefficient. λM can be calculated by the following formula [24]:

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In the formula, cM is the average thickness of the martensite lath and is taken as 0.2 μm.
The dislocations of ferrite, austenite and martensite are strengthened as follows:
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In the formula, αi is a constant, and Mi, Gi, and bi are respectively the Taylor factor, shear modulus, and Burgers vector mode of each phase. The values are shown in Table 2; ρi is the dislocation density of each phase and is a dependent variable. function. The initial dislocation density ρ0 of ferrite and austenite is related to the heat treatment temperature, and the specific values are shown in Table 2. The initial dislocation density ρ0 of the strain-induced martensite was set to 1×1015 m-2. The variation law of ρi-dependent variables can be expressed by the Kocks-Mecking model [25,26]:

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In the formula (15), the first and second terms in the parentheses on the right side represent the proliferation of dislocation densities caused by grain boundaries and Lin dislocations, respectively, and the third term represents the decrease in dislocation density caused by dynamic recovery. Λi is the effective inter-area spacing for each phase, Λα and ΛM are equal to dα and cM, respectively, and Λγ can be expressed by (12) (Λγ=L). At this time, iM=1. K1iki1 and k2iki2 represent the proliferation coefficient and dynamic recovery coefficient of each phase error, respectively. It is worth noting that the values of k1αkα1 and k1γkγ1 in ferrite and austenite will cause volume expansion due to martensitic transformation, and the resulting deformation induced martensite will simultaneously compress the surrounding austenite and ferrite. Therefore, the values of k1αkα1 and k1γkγ1 should be related to the volume fraction of martensite. The specific values are shown in Table 2. Pi is a coefficient related to the grain size, and Bouaziz et al. [27] define it as the probability of dislocations from being absorbed by grain boundaries. Its expression is as follows:

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Where dcidic is the critical grain size. The growth of dislocation density has a great relationship with the grain size [28]. When the actual grain size is smaller than the critical size, the dislocation quenching rate increases, so that the smaller the grain size is, the less dislocations are prone to be propagated. The lower the hardening rate.

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Table 2 Parameters used in the constitutive model calculations for cold rolled 7Mn steel at room temperature

4 Simulation Results and Discussion
Figure 6 shows simulated calculations and measured austenite dislocation densities. Since some of the austenite diffraction peaks of the tensile specimens disappear after the deformation increases, S700 and S720 only measure dislocation densities when the true strain is less than 0.22. It can be seen from Fig. 6 that the simulated austenite dislocation density is basically consistent with the measured value in terms of numerical values and changing trends. Figures 7a and b are the true stress curve and the work hardening rate curve for the simulation calculation and the actual measurement, respectively. From Fig. 7a, it can be seen that for the S680 and S700, the simulated and measured true stress-true strain curves agree well, but the true stress of the simulated S720 has a certain deviation from the measured value after the true strain is greater than 0.2. It may be related to two factors: one is the accuracy of the model itself, and the other is that it is difficult to accurately measure the kinetics of martensite transformation during the experiment. Because the volume fraction of martensite is not measured in situ, the macroscopic deviation of the chemical composition and the slight temperature difference during heat treatment may cause a certain difference in the kinetics of martensite transformation for different tensile specimens. Affected by the measured results. It can also be seen from Figure 7b that because the model does not take into account the influence of static strain aging (Lüders strain) and dynamic strain aging (stress sawing), there is also a certain deviation between the simulated value and the measured work hardening rate. However, in general, the simulated true stress and work hardening rate values and the variation trend with true strain agree well with the measured data. Looking at Figures 6 and 7, it can be considered that the constitutive model used in this work can better describe the evolution of the microstructure and mechanical properties of the cold-rolled 7Mn steel during the drawing process.
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Fig.6 Measured[16] and calculated dislocation densities (ρ) of austenite

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Fig.7 Measured and calculated true stress-true strain curves (a) and corresponding curves of work hardening rate (WHR) (b)

From equation (3), it can be seen that the flow stress of cold-rolled 7Mn steel is jointly provided by ferrite, austenite and martensite. Differentiating the left and right sides of equation (3) yields dσt/dε=d(σfVf)/dε+d(σγVγ)/dε+d(σMVM)/dε, ie, the total work hardening rate of cold rolled 7Mn steel. Provided by ferrite, austenite, and martensite. The constitutive model introduced in this paper will be used to analyze the flow stress and work-hardening rate changes of the various compositional phases in the cold-rolled 7Mn steel after annealing at different temperatures. The total flow stress of the simulation calculation and the flow stress of each component phase are shown in Figures 8a, c, and e. The corresponding total work-hardening rates and the work-hardening rates of the phases are shown in Figures 8b, d, and f.

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Fig.8 The calculated true stress-true strain curves (a, c, e) and curves of work hardening rate of the composite and each constituent phase (b, d, f) in S680 (a, b), S700 (c, d) and S720 (e, f)

It can be seen from Fig. 8a that the yield strength of cold-rolled 7Mn steel is jointly provided by ferrite and austenite after annealing at 680 °C. The ferrite provides a yield strength of about 510 MPa and austenite provides about 370 MPa. Considering that the volume fraction of austenite in the initial stage of deformation is only about 30%, it can be considered that the yield strength of austenite is larger than that of ferrite, which is consistent with the description in [17]. As the amount of deformation increases, the flow stress of ferrite slightly increases, while the flow stress of austenite increases slightly. The study [29,30] showed that in the single-phase ferrite materials, the work hardening ability of the finer grains is lower, so the lower work hardening rate of the ferrite in the S680 is mainly due to the small grain size. . The grain size of austenite in S680 is also small (as shown in Table 1), but its work hardening rate is higher than that of ferrite, indicating that the response of ferrite and metastable austenite is not the same. In the middle and later stages of deformation, a small amount of martensite is formed in S680, which can also provide part of the flow stress. The work hardening rate of the total and each constituent phase at the time of drawing of S680 is shown in Fig. 8b. When the strain amount is less than 0.1, the work hardening rate of the material is mostly provided by austenite, and a small part is provided by ferrite; after the strain amount is greater than 0.1, the work hardening provided by the deformation induced martensite gradually rises, in the true strain. It peaks at around 0.17 and then slowly declines. Under the joint action of ferrite, austenite and martensite, the total work hardening rate of S680 slowly decreases from about 1700 MPa in the initial stage of deformation. When the true strain is between 0.10 and 0.17, it basically remains unchanged, and then continues to decrease until When it is equal to true stress, S680 begins to neck down.
The flow stress of the S700 overall and each component phase is shown in Figure 8c. Comparing Figures 8c and a, it can be seen that although the yield strength of S700 is lower than that of S680, the slope of the true stress curve of S700 is significantly higher than that of the latter, ie the work hardening rate of S700 is higher than that of S680. The lower yield strength of the S700 is due to its larger ferrite and austenite grain sizes than the S680, as shown in Table 1. There are two main reasons for the higher work hardening rate of the S700. The first is that the work hardening rate of austenite is higher in the early stage of deformation; the second is that the austenite produces a significant TRIP effect in the later stage of deformation, resulting in a total work hardening rate. Stay at a higher level, as shown in Figure 8d. Due to the large austenite grain size in S700, the rate of proliferation of dislocations at the initial stage of deformation is higher, resulting in a higher work hardening rate than austenite in S680. In addition, due to the reduced stability of austenite in S700 [11], the austenite will continuously transform into martensite during the deformation and continue to produce the TRIP effect. As the volume fraction continues to decrease, the flow stress provided by austenite gradually decreases after the true strain is greater than 0.2, and the work hardening rate is reduced to a negative value when the true strain is approximately 0.2; at the same time, the rheology of the newly formed martensite is also reduced. The stress and work hardening rate increase significantly, making up for the effect of a reduction in austenite volume fraction. Under the combined action of three phases, the work hardening rate of S700 decreased from about 2000 MPa in the initial stage of deformation to 1350 MPa at a true strain of 0.13. Since the rapid formation of martensite, the work hardening rate was increased again to 0.28 in true strain. At 1600 MPa, the work hardening rate drops again until it intersects with true stress. It can be seen that due to the obvious TRIP effect of austenite, the strain amount when the work hardening rate of S700 intersects with true stress is greatly increased, ie, the uniform elongation of the material is significantly increased.
When the annealing temperature is increased to 720 °C, the size of the reversed austenite becomes larger, and the content of Mn and C in the solid solution is lower than those in S680 and S700, so its stability is further reduced. If ΔV/ΔεΔV/Δε is used to represent the martensite transformation rate (ΔVΔV and ΔεΔε represent the martensite content and the true strain variation, respectively), the average martensite transformation rate in S700 is 0.57 based on the data in FIG. 3 . The S720 is as high as 0.91, which shows that in S720, the volume fraction of martensite is rapidly increased due to the reduction of austenite stability, resulting in a rapid increase of the flow stress of the material, and finally making the flow stress of the S720 true strain. Between 0.1 and 0.2, there are obvious S-shaped features, as shown in Figure 8e. The work hardening rate of the S720 is followed by a distinct three-stage feature [3]. That is, the work hardening rate decreases first with the increase of the deformation amount, then rises and then decreases again, as shown in I, II, and III in Fig. 8f. It is worth pointing out that the S700’s work-hardening rate curve also has a three-stage characteristic, but due to the slower rate of martensite transformation, the difference between the highest and lowest values of the second segment is not large, resulting in a true stress curve S-type The characteristics are not obvious.
5 Conclusion
(1) As the annealing temperature increases, the mechanical stability of the reversed austenite in the cold-rolled 7Mn steel gradually decreases, resulting in a rapid increase in the strain-induced martensitic transformation rate. After annealing at 700 °C, the stability of the reversed austenite is moderate. At this time, the comprehensive mechanical properties of the material are optimal.
(2) Using the dislocation density-based constitutive model to study the tensile and work hardening behaviors of cold-rolled 7Mn steel, the calculated values are in good agreement with the measured results. Therefore, the constitutive model can be used to analyze the cold-rolled 7Mn steel. The reason for the difference in mechanical behavior after annealing at different temperatures.
(3) Austenite stability has a decisive influence on the true stress and work hardening rate curves of cold-rolled 7Mn steels. After annealing at 680 °C, the stability of the reversed austenite is too high, and no obvious TRIP effect can be produced during deformation. The work hardening rate and uniform elongation of the material are both low; after annealing at 700 °C, the stability of austenite is moderate. During the deformation process, the TRIP effect is continuously generated, so that the work hardening rate of the material is maintained at a relatively high level within a relatively large strain range, and the ultimate mechanical properties of the material are optimized; after the annealing temperature is increased to 720 °C, the reverse transformation occurs. The stability of austenite is too low, and almost all of it transforms into martensite in a small strain range, resulting in a material with a higher tensile strength but a lower uniform elongation.
The authors have declared that no competing interests exist.

Source: China Pipe Fittings Manufacturer – Yaang Pipe Industry (www.metallicsteel.com)

(Yaang Pipe Industry is a leading manufacturer and supplier of nickel alloy and stainless steel products, including Super Duplex Stainless Steel Flanges, Stainless Steel Flanges, Stainless Steel Pipe Fittings, Stainless Steel Pipe. Yaang products are widely used in Shipbuilding, Nuclear power, Marine engineering, Petroleum, Chemical, Mining, Sewage treatment, Natural gas and Pressure vessels and other industries.)

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