Effect of Tempering and Strain on Decomposition of Metastable Austenite in X210CrW12 Thixo-Cast Steel

The thixoformed X210CrW12 steel consisted of austenitic globular grains (average size 46 μm) surrounded by the eutectic mixture of M₇C₃ carbides, γ-Fe and α-Fe (Ref 9). The compression test results of X210CrW12 steel after SSM processing are shown in Fig. 2. The relatively low yield strength (YS—810 MPa) was mainly due to a large fraction of austenite, which began to deform at low stress. It shows, however, a strong hardening effect resulting in very high final compression strength of 4.85 GPa. Its hardness also increased with increasing plastic strain due to the occurrence of tension. Rogal et al. (Ref 11) precisely studied the deformation of the X210CrW12 steel, which proceeded by dislocation slip and mechanical twining. Figure 3 shows the x-ray analysis of steel in the un-deformed state (directly after SSM) and after compression at 4.8 GPa. The changes in the intensity and width of austenite peaks were observed. It may suggest that the deformation led to a decrease of the crystal size. Small shifts of austenite peaks to higher 2θ values were observed due to the stresses accumulated within the deformed material, which contributed to hardening of the material. The presence of ε-Fe (Ref 11) was confirmed using the HRTEM analysis of highly strained area, which did not have large influence on hardening. Such a result led to the conclusion that the X210CrW12 thixo-cast austenite was rather mechanically stable.

The sample of the X210CrW12 steel formed with thixotropic technique revealed intensive broadening of crystallite orientation in both, the austenitic (Fig. 4a) and ferrite phases (Fig. 4b). It seems that the thixo-cast might have had properties typical for isotropic material. Figure 1c shows the sample after compression at strain 4850 MPa, which is the limit value of compression strength. The analysis of pole figures (Fig. 4a, b) indicated distinct texturization of microstructure in comparison with the un-deformed sample. The main texture components for austenite were {101}, 〈010〉, while ferrite did not reveal clearly formed orientation. The failure of material occurred by shear mechanism. It can be seen in (111) pole figure of austenite (Fig. 4c) with increased density of poles located close to the direction normal to the shear plane. The crystallographic slip occurred in (111) plane in [−101] and [01−1] direction, which is defined in Fig. 1d. However, due to weakly defined texture in the ferrite phase (Fig. 4b) its contribution to the mechanism of destruction has been difficult to interpret based on pole figures.

The austenitic structure of X210CrW12 steel after the SSP was metastable and its decomposition led to the improvement of mechanical properties (Ref 1). Positive effects of stress introduced during warm and hot plastic deformation of tool steels were reported in numerous papers (Ref 13, 17-19). In order to study the influence of stress on decomposition mechanisms of austenite at tempering temperature range, thixo-cast samples were subjected to cold plastic deformation. The strengthening effect of strain-induced transformation was observed in a Hadfield or stainless steel deformed at room temperature (Ref 20-23).

The DSC studies were conducted in order to analyze the influence of strain on temperature of austenite decomposition of the X210CrW12 thixo-cast steel. Figure 5 shows calorimetric heat flow curves for the samples: unstressed, marked as (1) and compressed up to 4.8 GPa, marked as (4) in the temperature range of 50 to 685 °C. The magnified temperature peak in range (575 to 655 °C) for various stressed samples: (0)—unstressed, (2)—compressed to 2 GPa, (3)—compressed to 3.8 GPa, and (4)—to 4.8 GPa are presented in the dotted frame. The heat flow curve of the un-deformed sample (green curve with dots) showed a small exothermic effect in the temperature range from 150 to 350 °C, which was probably connected with tempering of martensite. The hardness of grains slightly decreased during heating the sample up to 400 °C, what could be due to the release of tension, thus suggesting that the positive effect was not related to the transformations in the austenitic grains (Ref 9). The exothermic effect of transformation of austenite globular grains into ferrite, carbides and perlite is visible in temperature range 572 to 668 °C. It should be noted that the inclination angle of heat flow curve of unstressed sample is lower in comparison with the curves after deformation. It suggests that the phase transformation rate of the un-deformed sample was lower. The compression strain of the X210CrW12 thixo-cast steel brought about a shift of the beginning of small exothermic reaction down to range 100 to 250 °C. The large positive effect connected with austenite decomposition that started at above 558 °C and disappeared at about 677 °C is shown in Fig. 5. The thermal effect in range 100 to 250 °C could be also connected with the precipitation of transition carbides such as χ-Fe₅C₂, η-Fe₂C (Ref 24) from the ε-Fe martensite (Ref 11). With increasing stress, huge peak movements to lower temperatures: unstressed sample—to 634 °C and compressed ones to 2, 3.8 and 4.8 GPa to 621, 611, and 599 °C, respectively, was observed. It was confirmed in (Ref 16), that with increasing level of deformation of the supersaturated solid solution, the peak temperatures moved to lower ones. Similar effects of the movement of γ-Fe transformation temperature were visible in the hot deformed M2 tool steel, in which a high density of defects in the strained austenite resulted in the acceleration of diffusion-controlled phase transformations (Ref 17). Additionally, straining the austenitic phases raised the final hardness by 10 to 20% through the precipitation of very fine carbides (Ref 17).

Figure 6a shows a SEM microstructure of SSM processed X210CrW12 steel heated up to 634 °C (temperature peak according to DSC analysis visible in Fig. 5) at rate 20 °C/min, cooled next with the furnace. One can observe globular grains with α-Fe and carbides nodules resulting from the decomposition of γ-Fe. The reaction started inwards from the grain boundaries surrounded by the eutectic mixture. The observed changes in the microstructure is somewhat similar to the discontinuous precipitation reaction in steel reported by Ainsley et al. in (Ref 25), who showed that the transformation γ-Fe → γ₁-Fe +VC proceeded through the discontinuous precipitation, which started from the grain boundary. According to Bhadeshia (Ref 26) the decomposition of austenite can occur at temperatures above Bs (if the holding time is short enough) or below Bs (if the holding time is long). Such a transformation product has been classified as pearlite consisting of ferrite and M₇C₃ carbides.

The microhardness analysis (Fig 6b) shows that the hardness of dark-etched-area was about 700 HV_0.02, while that of light places (un-etched) in the center part of grains reached around 780 HV_0.02. It suggested, that in the center part of grains, a mixture of martensite and residual austenite (un-etched area) existed in front of lamellar pearlite. The hard α′-Fe phase formed probably by the depletion of austenite in carbon, leading to the decrease of Ms temperature. The γ-Fe to α′-Fe transformation probably took place during cooling. The x-ray analysis of the investigated steel confirmed the occurrence of ferrite, M₇C₃ carbides and residual austenite in the amount below 5 vol.%. Similar results were obtained for the X210CrW12 steel after SSM and tempered for 2 h at 525 °C (Ref 9).

In the TEM bright field micrograph (Fig. 7a), the areas similar to those appeared due to the discontinuous precipitation reaction were visible, in which a supersaturated γ-Fe matrix (right side of BF image) transformed into more thermodynamically stable mixture of very fine ferrite and carbides (the top left corner). The microstructure with morphology close to perlite could be also seen in the bottom left corner of Fig. 7a. According to Thompson (Ref 27), the composition of the matrix changed discontinuously across the reaction front. The electron diffraction pattern (Fig. 7b) from the light area shows reflections from the γ-Fe (1−1−1) and (02−2), which were indexed according to austenite zone axis [211]. The γ-Fe solid solution without carbides indicated a large supersaturation with alloying elements and carbon. Figure 7c shows SAEDP of the decomposed austenite giving reflections from M₇C₃ (0111) and (−1100) of tetragonal lattice and corresponding zone axis [112−3] and α-Fe [100]. The obtained phases are stable at room temperature according to the phase diagram of X210Crw12 steel (Ref 7). The typical, commonly known discontinuous reaction in steel is responsible for the transformation of austenite into perlite during slow cooling in eutectoid steels (Ref 28). The discontinuous precipitation in austenitic phase of Fe13Mn2V0.8C steel produces the γ-Fe phase and VC precipitates with either a fibrous or particulate morphology and (Ref 25).

The bright field micrograph (Fig 8) taken of a grain of X210CrW12 thixo-cast after compression stress to 4.8 GPa showed frequent twin bands within the austenitic grains, which revealed high density of defects. Additionally, the presence of ε-martensite of hexagonal lattice with numerous random stacking faults was confirmed in (Ref 11). According to the selected area electron diffraction pattern SAEDP (shown as the insert), strong reflections of the γ-Fe at zone axis orientation [110] with twins on (1−11) plane are visible.

The decomposition caused by tempering of severely strained and supersaturated steel can lead to obtaining an uncommon product of transformation, which have influence on mechanical properties.

Heating the deformed sample up to 599 °C (temperature peak according to DSC analysis Fig. 5), resulted in a partial decomposition of residual austenite. The transformation started simultaneously from grain boundaries, twins and slip bands (Fig. 9a), which generated a high number of defects in the austenite. As can be seen in Fig. 9b, the lateral areas of the image showed the morphology of microstructure similar to lamellar. A high magnification obtained in the TEM analysis (Fig. 9c) revealed the pearlite structure in the selected part of middle area grains. In the central part of Fig. 9b the areas, in which carbides (marked by arrows) started nucleating at slip bands as well as at twins (marked by ring) are visible. Figure 9d shows the TEM-BF image with corresponding selected area DF of central parts of Fig 9b. It presents ferrite bands with carbide particles along the previous austenitic grain boundaries. The strong contrast suggests the presence of defects and carbides, which probably formed in the original orientation of deformed γ-Fe band. The SAEDP from plates and areas of their boundaries (marked with dotted ring) shows strong reflections from α-Fe [11−1] matrix and Fe₃C carbides at zone axis [−1−13]. Such an area in steel may be the starting point for the discontinuous transformation due to the coagulation of carbides, which can lead to breaking the coherence with the matrix. It suggests that M₃C carbides precipitated mainly at strongly deformed areas as a first product of strained austenite. The appearance of the M₃C carbide might have been preceded by the precipitation of transient carbides (χ-Fe₅C₂, η-Fe₂C) probably present in the structure. The x-ray analysis confirmed the presence of α-Fe and M₇C₃ carbides in the sample strained 4.8 GPa/599 °C (Fig. 3). The content of M₃C carbides was probably too small to be identified by x-ray examination. Higher density of defects in the hot deformed austenite increased the heterogeneous nucleation rate and supported the carbide precipitation which has been shown in numerous works (Ref 17, 19, 29).

The microhardness measurements of 4.8 GPa/599 °C sample (Fig. 10) showed various hardness values across the grain. It was connected with the progress of γ-Fe decomposition. At places, where the front of discontinuous transformation passed forming the mixture of ferrite and M₇C₃ carbides, the hardness was relatively low; about 660 HV_0.02, while in the area with ferrite plates and M₃C carbides it reached 900 HV_0.02. Slightly higher hardness in comparison with the unstressed sample was probably caused by the strain-induced transformation of grains.

The TEM-BF image (Fig. 11a) shows the interface between the front of discontinuous precipitation, in which previously decomposed austenite into mixture of M₇C₃ carbides and α-Fe solid solution (right side of micrograph) reacted with α-Fe plate and M₃C formed as a result of plastic deformation (left side). It suggests that such a microstructure was not stable at 599 °C and transformed due to high diffusivity of carbon. Additionally, a new discontinuous precipitation was observed in comparison with the un-deformed material. The electron diffraction pattern (Fig 11b) from the left side of Fig 11b (area marked with dotted ring) shows reflections from γ-Fe (211) and (011), which were indexed according to ferrite zone axis [0−11]. Reflections from M₇C₃ carbides (01−10) and (−1−122) with hexagonal lattice and [20−2−1] zone axis are also visible in the SAEDP. Figure 11c shows the SAEDP from the area of microstructure marked with dotted ring giving reflections from the orthorhombic MC₃ carbide (11−2) and (010) corresponding zone axis [112−3] and α-Fe [302]. It suggests that the decomposition of austenite by discontinuous precipitation mechanism played the crucial role in the deformed steel. In contrast to non-deformed state they cover the area of grains. An additional front of reaction α-Fe plate +M₃C → mixture of α-Fe and M₇C₃ carbides was also observed. Similar mechanisms were observed by Shtansky (Ref 30) in the Fe17Cr0.5C steel tempered at 735 °C/10 s, in which the martensite transformed by discontinuous-like precipitation into the mixture of cementite, α-Fe, and carbides: M₂₃C₆ and M₇C₃. Various carbide morphologies were observed close to the reaction front: rod-like, spherical, or lamellar.