The ductility behavior of a high-Mn twinning induced plasticity steel during cold-to-hot deformation

The ductility behavior of a high-Mn TWIP steel (containing 30% wt Mn) has been studied using tensile testing method in a wide range of temperature (100–1000 °C) under the strain rate of 10⁻⁴ s⁻¹. The hot compression characteristics of the experimental alloy are considered to assist in explaining the related deformation mechanisms. The results indicate that the ductility decreases with temperature; however, two regions of moderately improved ductility have also been realized. The former is attributed to the reduction of twinning activity by increasing the temperature. On the other hand, the activation of dynamic recovery at 400 °C causes the ductility to increase. The fracture surface observations denote the occurrence of grain boundary sliding at temperatures above 500 °C. As the dominant restoration process alters to partial dynamic recrystallization at 800 °C, the tensile ductility continues to decrease. By increasing the temperature to 1000 °C, the fraction of dynamically recrystallized grains is significantly increased and the ductility is improved.

EXPERIMENTAL 

The experimental steel was supplied in as-cast condition with the chemical composition of 29.7 Mn–2.5Al–0.6Si–0.17C, wt%. In order to eliminate the dendritic structure and to achieve a uniform microstructure, the as-received material was forged at 1150 °C. This was followed by subsequent annealing at 1000 °C for 90 min to remove any segregation of alloying elements in particular that of Mn. The initial microstructure of the experimental alloy is shown in Fig. 1. As is seen the microstructure is fully composed of an austenitic structure characterized by annealing twins. The initial grain size, measured by linear-interception method, is 80 μm.

The tensile tests were conducted according to ASTM E8M standard [16] using cylindrical specimens with a reduced section diameter of 6mm and a gauge length of 30 mm. The isothermal tensile tests were carried out in the temperature range of 100–1000 °C with intervals of 100 °C under the strain rate of 10⁻⁴ s⁻¹. The hot compression tests were also carried out in the aforementioned deformation conditions to study the involved micro-mechanism. The hot compression testing specimens were machined according to ASTM E209 standard [17] using cylindrical specimens in the sizes of Φ8 mm×H12 mm. The specimens were first heated up to the deformation temperature and held isothermally for 5 min prior to straining. The tension and compression tests were carried out using an Instron-4208 universal testing machine, equipped with a contact extensometer and resistance furnace. All thermo-mechanical cycles were ceased by quenching the specimens in water just after straining. The elongation-to-failure was measured from the gauge length of the fractured specimens. To investigate the final microstructures, the specimens were sectioned along deformation axis, mounted using cold curing resin, ground and polished step by step up to the final polishing by 0.05 μm Al₂O₃ powders. The related fracture surfaces were also examined using scanning electron microscopy (SEM) to clarify the ductility behavior.

Results and discussion

COMPRESSION BEHAVIOR

To assist determining the dominant mechanisms operating under the specified deformation condition, the warm-to-hot compression characteristics of the alloy were considered. The obtained true stress–true strain curves are shown in Fig. 2. This figure reveals three characteristic flow curves. The first category (in the range of 100–300 °C) includes a continuous work hardening region, which may imply the prevalence of twinning induced plasticity effect at these relatively medium temperatures. The twinning causes a high value of instantaneous hardening rate (n value). This is commonly attributed to the reduction of the dislocation mean free path by increasing the fraction of deformation twins as strong obstacles to dislocation glide [18]. The second category of the flow curves, achieved at 400–700 °C, shows a steady flow stress plateau. This appears to be originating from the balance of dynamic recovery (DRV) and work hardening. In fact, the dislocation activity and DRV are markedly enhanced above 400 °C, so that the rates of work hardening and recovery reach a dynamic equilibrium. The third category (in the temperature range of 800–1000 °C) exhibits typical dynamic recrystallization (DRX) behavior with a single peak stress followed by a gradual fall toward a steady-state stress 

As the dislocation density is directly related to the flow stress, some indication of the change in dislocation content with strain may often be inferred from the stress–strain behavior. Accordingly the dislocation density increases progressively during deformation. The increase in dislocation density is due to the continued trapping of newly created mobile dislocations by existing dislocations and their incorporation into the various microstructural features (twins and grains boundaries) that are characteristic of the deformed state. This would result in a continuous work hardening region in Type A deformation regime. In the other deformation conditions (Type B and Type C deformation regimes), which are usually of importance for the occurrence of restoration processes, the dislocation density decreases in a way to achieve a balance between dislocation generation and annihilation at the larger imposed strains. Recovery generally involves a partial restoration because the dislocation structure is not completely removed, but reaches a metastable state. In contrast, during dynamic recrystallization new dislocation-free grains are formed within the deformed or recovered structure. These then grow and consume the old grains, resulting in a new grain structure with a low dislocation density.

TENSILE DEFORMATION BEHAVIOR

Fig.shows the different types of true stress–strain curves obtained from tensile tests in the temperature range of 100–1000 °C. Different deformation mechanisms may be realized from various tensile behaviors. The corresponding flow curves in the range of 700–1000 °C (Type A) include a short work hardening region up to the ultimate tensile strength (UTS) followed by a long post-UTS region. The usual flow softening in tensile curve is invariably associated with geometric instability related to the necking; this generally results in a short post-UTS region. However, the lower rate of stress drop beyond the necking (i.e. lower rates of work softening in post-UTS regions) of type A curves may be described relying on the occurrence of DRX in this range of temperature (800–1000 °C, ). In fact the cavities and discontinuities, which are being formed in the microstructure, are isolated from the grain boundaries through grain boundary migration taking place during DRX. Consequently, the growth and coalescence of these cavities is not readily achieved away from grain boundaries so the strain to fracture beyond the necking is increased. In fact it is anticipated that the occurrence of DRX assists in obtaining an advanced flow localization, which in turn may postpone the necking.

The second type of tensile flow behavior (Type B) exhibits a shorter post-UTS strain portion with respect to that of type A. The type B curves, which are achieved at 100–600 °C, reveal a very long pre-UTS region with higher rate of work hardening where a large amount of deformation may be driven through the mechanical twinning. This was followed by a short range of work softening region after UTS, ultimately leading to fracture.

DUCTILITY BEHAVIOR

The variation of elongation-to-fracture as a function of test temperature is shown in Fig. . The curves indicate that there is a general trend of decreasing ductility with temperature. But, as is seen there are two regions (regions II and IV, denoted in Fig. 4) where the general trend changes and the ductility increases. The details of ductility variation at different deformation conditions have been described point by point and are as follows.

FUENTES 

http://www.sciencedirect.com.bdatos.usantotomas.edu.co:2048/science/article/pii/S0921509312014864