M23C6 carbides in TWIP steel

Figure 8: (a) and (b) Typical TEM images of the TWIP steel after 10-revolutions HPT, showing the only austenite and multiple phase microstructures, respectively and the inserted images are corresponding SEAD patterns; (cee) typical morphology of ferrite, cementite and M23C6 carbide with their SEAD patterns, respectively. Scale bars: 100, 200, 50, 20 nm.

Figure 9: (a) and (b) Characteristic TEM images and the SEAD patterns of the TWIP steel microstructures after 12-revolutions HPT and the deformation twins aremarked by the arrows in (b); (cee) HRTEM images and corresponding FFT of austenite, ferrite and M23C6 carbide, respectively. Scale bars: 100, 200, 1 nm.

Carbide name: M23C6
Record No.: 1543
Carbide formula: M23C6
Carbide type: M23C6
Carbide composition in weight %: No data
Image type: TEM, SEAD
Steel name: TWIP steel
Mat.No. (Wr.Nr.) designation: No data
DIN designation: No data
AISI/SAE/ASTM designation: No data
Other designation: No data
Steel group: TWIP (twinning-induced plasticity) steels
Steel composition in weight %: Fe-17.37Mn-3.31C- 3.34Al-0.94Si in at.% or Fee18Mn-0.75C-1.7Al-0.5Si
Heat treatment/condition: The original TWIP steel used in this investigation is a commercial product produced by Ansteel Co. Ltd. in China. The nominal chemical composition of the TWIP steel is Fe-17.37Mn-3.31C- 3.34Al-0.94Si in at.% or Fee18Mn-0.75C-1.7Al-0.5Si in wt.%. Detailed information about the material is available in Refs. For HPT processing, the fully recrystallized steel was cut into disc samples with a diameter of 20 mm and a thickness of 1.5 mm. The disks were HPT processed through 1, 5, 10, and 12 revolutions (thereafter referred to as 1R, 5R, 10R and 12R samples, respectively) under a quasi-constrained condition with an imposed pressure of 4.0 GPa at 573 K. Before HPT processing, the anvils were preheated to 573 K and then held for 5 min to warm the samples. The temperature was kept at 573 10 K during the process. The samples were quenched into water immediately after HPT processing. One or two revolutions of HPT were carried out for one time and abovementioned procedure was repeated to reach the designed revolution numbers.
Note: The microstructural evolution of twinning-induced plasticity steel during high-pressure torsion (HPT) processing at 573 K was systematically evaluated. Due to the high processing temperature, the formation of a homogeneous nanostructure was primarily dominated by complicated dislocation and grain boundary activities in lieu of deformation twinning. Apart from the grain refinement process, phase transformation took place at late stages of deformation, resulting in the microstructural fingerprint of equaxied nanograins with multiple phases in the steel. On account of remarkable elemental redistribution, the diffusion-controlled nature of the transformation was convincingly identified. During the transformation, although the cementite also initially formed, austenite eventually decomposed into ferrite and Mn-riched M23C6 carbide, implying that Mn is the pivotal alloying element for the transformation kinetics. Owing to the sluggish bulk diffusivity of Mn, it is proposed that a high density of defects, nanostructures and the HPT processing play a crucial role in promoting the elemental diffusion and segregation and in stimulating the phase transformation.

TEM investigation on the microstructures of the 10R and 12R samples: Various dislocation configurations including dislocation cells and elongated subgrain boundaries are the characteristics of the inhomogeneous microstructures in many regions of 10R sample, interspersed with several isolated nanograins, as marked by arrows in Fig. 8 (a). Corresponding selected area electron diffraction (SAED) patterns confirm that only austenite exists in these areas. In contrast, many equiaxed ultrafine and nanoscale grains can be found in some local places, as exhibited in Fig. 8(b), while there are still apparent dislocation substructures in the ultrafine grains, most of which are the austenitic phase. SAED patterns obtained from these zones, as inserted in Fig. 8 (b), validate the existence of multiple phases, indicating that the phase transformation took place. According to the SAED patterns, besides the austenite phase, the ferrite phase can be readily recognized as well, while several isolated spots, originated from carbides, are difficult to be accurately indexed. Nevertheless, through indexing the SAED patterns of many individual grains, apart from the austenite and the ferrite, cementite and M23C6 carbide were efficaciously identified. Fig. 8 (cee) presents typical morphologies and SAED patterns of ferrites, cementite and M23C6, respectively. After HPT for 12 revolutions, NC grains with equiaxed morphology are the predominant microstructural signature of the steel in most regions, as shown in Fig. 9 (a) and (b). The corresponding SAED pattern inserted in Fig. 9 (a) identifies three phases, i.e., austenite, ferrite and M23C6 carbide. No cementite was found. Typical atomic-scale TEM images and their corresponding fast Fourier transformation (FFT) patterns of austenite, ferrite and M23C6 carbides are shown in Fig. 9 (c)e(e). While the GBs in SPD materials processed at the ambient temperature are usually blurred due to the significant internal stress and limited dynamic recovery, most GBs in present study are quite clear, as exhibited in Fig. 9 (a) and (b), implying that dynamic recovery or recrystallization caused by the high processing temperature, play an essential role in shaping the morphology of these fine grains. A few deformation twins were occasionally found in austenitic nanograins and M23C6 carbides. Some twinned grains are marked by arrows in Fig. 9 (b). These twins are expected to form via partial dislocation emissions from GBs. In a whole, the microstructure in the 12R sample is quite homogeneous in comparison to that in the 10R sample.
Links: No data
Reference: Not shown in this demo version.

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