Study of microbanding in a Fe-22Mn-0.6C steel: implications on strain-hardening mechanisms

Recent investigation of the deformation structure in austenitic high-Mn steels, namely, Fe-Mn-C and Fe-Mn-Al-C alloys, have revealed that specific dislocation substructures play an important role on their strain hardening behavior [1-6]. At low deformation (true tensile strain <XY), microstructure refinement is characterized by the formation of complex dislocation substructures such as cells, cell blocks and Taylor lattices, depending on the grain orientation. Quantitatively, their contribution to strain hardening is described by t =GKb/D, where t is the resolved flow stress, G is the shear modulus, K is a constant, b is the magnitude of the Burgers vector and D is the dislocation substructure size. Dislocation structure hardening has been recently shown as a relevant strain hardening mechanism in Fe-Mn-Al-C alloys [4]. Specifically, the combination of dislocation substructure hardening (at low and medium true tensile strains  < XY) in conjunction with deformation twinning (at higher true tensile strains > XY) leads to multiple-stage strain hardening behavior resulting in permanent strain hardening reserves also at high deformations and hence, superior mechanical properties. The underlying mechanisms controlling the formation of dislocation structures are controlled by dislocation reactions that are stress-assisted mechanisms [7]. Accordingly, dislocation patterning is strongly dependent on crystallographic orientation. Microbands are dislocation substructures that are associated to strain localization phenomena [8, 9]. These dislocation configurations have been reported in Fe-Mn-Al-C alloys deformed in tension [4, 10-12]. However, their formation mechanism and contribution to strain hardening are still unclear. For instance, we have recently reported that in a Fe–30.5Mn–2.1Al–1.2C (wt. %) alloy [XY] deformed in tension, microbanding plays no role on strain hardening due to texture effects. Therefore, the present study aims at clarifying the role of the microbanding mechanism and its contribution to strain hardening in high-Mn steels. We have investigated the influence of strain path, namely, tension and shear, on microbanding in a Fe-22Mn-0.6C (wt. %) steel. Dislocation substructures were examined by combined electron channeling contrast imaging (ECCI) and electron backscatter diffraction (EBSD). The microbanding mechanism is analyzed in terms of dislocation configurations and its role on strain hardening is discussed.


The high-Mn steel used in this study had the chemical composition Fe-22Mn-0.6C (wt.%). Details on the alloy processing can be found in [1]. The hot-rolled material showed a fully austenitic structure with an average grain size of 50 mm which remained stable during deformation at room temperature. Tensile and shear tests were carried out at room temperature at an initial strain rate of 5 × 10-4 s-1 to an equivalent true strain of 0.1. At this strain level, the deformation structure mainly consists of dislocation substructures with few deformation twins [1]. The tensile bone-shaped samples had 8 mm gage length, 2 mm gage width and 1 mm gage thickness. Shear deformation tests were performed in an in-house shear test set-up [13]. The shear samples had a rectangular shape of 40 × 14 × 2 mm3. Dislocation substructures were investigated by combined electron channeling contrast imaging (ECCI) and electron back scatter diffraction (EBSD). Longitudinal section of the tensile deformed sample, i.e. the section perpendicular to the tensile axis, was examined. In the sample deformed in shear, the SD-ND section (SD: shear direction, ND: normal direction) was characterized. High-resolution EBSD maps (step size of 50-100 nm) were taken in a 6500 F JEOL field emission gun-scanning electron microscope (FEG-SEM) equipped with a TSL OIM EBSD system. Dislocation substructures of the grains mapped by EBSD were subsequently examined by ECCI under controlled diffraction conditions, as in previous works [1, 3, 4, 14-16]. ECCI images were obtained with optimum contrast by tilting the matrix crystal in Bragg condition for high intensity reflections and exciting the corresponding diffraction vector in a “two-beam” condition. ECCI observations were carried out in a Zeiss Crossbeam instrument (XB 1540, Carl Zeiss SMT AG, Germany) operated at 10 kV.


We have examined the dislocation substructures in about 30-40 individual grains of each individual sample, namely, deformed in tension and in shear, by combined ECCI and EBSD. The experimental crystal orientations of the analyzed grains in the sample deformed in tension are plotted in the TA-IPF of Fig. 1 (TA: tensile axis; IPF: inverse pole figure). The data provide good representation of the deformation texture at 0.1 true strain [1]. The dislocation substructure is formed by two types of dislocation patterns, namely, dislocation cells (DCs) and cell blocks (CBs). The formation of these dislocation patterns is associated to the multiple character of slip, i.e. wavy and planar [3]. CB is the most common type of dislocation pattern. It is delimited by geometrically necessary boundaries (GNBs), referred to as highly-dense dislocation walls (HDDWs), and incidental dislocation boundaries (IDB), termed as dislocation walls [more quotes here from RISO folks etc ?]. These dislocation configurations are imaged by ECCI as bright compact layers (HDDWs) subdivided by finer bright layers (dislocation walls), Fig. 1. CBs are formed in grains oriented along the line between the <001>//TA and <111>//TA crystallographic directions. Grains oriented close to <112>//TA directions develop HDDWs that are parallel within 1-2° to {111} plane traces, i.e. crystallographic boundaries. Grains close to <111>//TA directions build up non-crystallographic HDDWs that are deviated up to 10° from {111} planes. Dislocation substructures similar to those found here have been reported in medium-to-high stacking fault energy metals such as copper and aluminum deformed in tension [17, 18]. Equiaxed DCs are built up in grains oriented close to <001>//TA directions, Fig, 1. Detailed examination of the dislocation boundaries, in particular HDDWs, did not reveal any effect of strain localization phenomena on dislocation patterning such as microbanding.


Fig. 2(a) shows the SD-IPF (SD: shear direction; IPF: inverse pole figure) containing the experimental crystal orientations of the analyzed grains in the sample deformed in shear. The data provide good representation of the deformation texture at 0.1 true strain. At this strain level, the dislocation substructure mainly consists of CBs. DCs were observed in few grains. Analysis of dislocation boundary alignment by combined ECCI and EBSD revealed that most of the HDDWs delimiting the CBs are non-crystallographic boundaries, which are deviated up to 10° from the {111} planes. Fig. 2(b) shows an ECCI image of the CB structure developed in a grain oriented close to <213>//SD direction where HDDWs are formed almost parallel to {351} planes (plane trace analysis was carried out by combined ECCI and EBSD). Detailed examination of the CB structure by ECCI revealed that microbanding is strongly promoted in the sample deformed in shear when compared to the tensile deformed samples. Fig. 3(a) shows a high-resolution EBSD map of the CB structure developed in a grain oriented close to <101>//SD direction. The EBSD map displays the orientation gradients calculated with respect to the point in the map containing the lowest kernel average misorientation value (blue-red: 0-10°). The misorientation profile along several dislocation substructures (indicated by an arrow) is shown in Fig. 3(b). This figure reveals that the CB structure developed in shear at 0.1 true strain is formed by low-angle dislocation boundaries with misorientations smaller than 2°. Detailed examination of the dislocation substructure by ECCI revealed that this structure is formed by cell blocks and microbands (indicated by arrows), Fig. 3(c). Specifically, we observed lenticular-shaped microbands along the HDDWS of the existing CBs. These dislocation configurations are formed by pairs of dislocation layers spaced between 200 and 500 nm. The ECCI image in fig. 3(d) shows a detailed view of the internal dislocation structure of a microband. Under the current diffraction conditions, the crystal matrix appears dark and dislocations appear as sharp bright lines due to the electron channeling mechanism [14, 19]. This figure reveals that the dislocation configuration is mainly formed by straight dislocations lying along the [101] crystallographic direction. The analysis of the dislocation reactions produced by the interaction between the active slip systems shows that these dislocations are Lomer-Cottrell dislocations. Following Schmid and Boas’ labeling of the FCC slip systems [20], we obtain the following Schmid factors (SF) for the slip systems of the crystal containing the microband shown in Fig. 3(d) (only slip systems with SFs above 0.5 are considered): SF=1.0 (D6), SF=0.8 (B5), SF=0.6 (B2 and C5). According to Thomson’s tetrahedron [21], short range interaction among dislocations pertaining to slip systems D6 and B2 results in the formation of Lomer-Cottrell locks. The formation of Lomer-Cottrell locks at microband interfaces has been recently considered in discrete dislocation dynamic simulations of the internal dislocation substructure of microbands in a Cu single crystal deformed in shear [22].


There are two types of microbanding mechanisms proposed for fcc metals, namely, the double cross-slip model [23] and the dislocation boundary splitting mechanism [24, 25]. In the first mechanism, microbanding is ascribed to unstable glide on latent slip systems involving pronounced dislocation cross-slip activity. This mechanism enables microbands to accommodate only small lattice rotations and accordingly, their misorientation angles are small (misorientation angles below 1° are predicted) [23]. In the second mechanism, the driving force for microband nucleation is the high misorientation of the dislocation boundaries delimiting the existing dislocation substructure. As the misorientation of the dislocation walls forming the microbands is preserved during the boundary splitting mechanism, the resulting microbands contain high misorientation angles, typically larger than 1°. Our observations reveal that microbanding in the present Fe-Mn-C alloy deformed in shear can be explained in terms of the dislocation boundary splitting mechanism. Microbands are nucleated at HDDWs, which provide the dislocation sources to build up the dislocation walls required in the microband formation process. They subsequently grow by a dislocation boundary–type splitting mechanism involving the activation of  dislocation cross-slip [24, 25].


The examination of the dislocation patterns obtained in tension and shear reveals that in the Fe-22Mn-0.6C (wt.%) steel studied here microbanding is strongly dependent on the strain path. This can be explained in terms of Schmid’s law. Deformation in shear enhances slip concentration on single or coplanar slip. Stable crystal orientations upon deformation in shear such as <112>//SD and <101>//SD contain slip systems with Schmid factors ~ 1 resulting in enhanced shear localization. This result agrees with the intense microbanding activity observed in solid solution fcc alloys such as Ni-Co and Al-Mg deformed in torsion or cold rolling [24, 25]. In these studies, microbands similar to those studied here were observed, which were referred to as 'first generation microbands' [24, 25]. However, the internal dislocation configuration of microbands was not investigated. Our ECCI observations reveal that specific dislocation configurations inside microbands are created, namely, Lomer-Cottrell dislocations, Fig. 3(c). These sessile dislocations play an important role on strain hardening due to forest-type hardening associated to short-range dislocation interactions [26]. This hardening mechanism results in a greater critical stress to transfer plastic deformation across the microband and accordingly, strain hardening is enhanced. This hardening mechanism associated to microbanding can provide, under specific deformation conditions, enhanced strain hardening in FeMnC alloys. In particular, we envisage that this hardening mechanism could significantly contribute, under specific deformation conditions, to the strain hardening behavior of Fe-Mn-Al-C alloys where intense microbanding has been reported [4, 10, 11]


It is worth to underline the role of dislocation cross-slip on the microband mechanism. As stated above, microbanding requires the activation of dislocation cross-slip [9, 23-25]. The role of carbon on dislocation cross-slip has been recently addressed [4]. Carbon strongly inhibits dislocation cross-slip and hence, has strong influence on dislocation patterning [4, 27]. As cross-slip is a stress-assisted mechanism, it results in a pronounced dependence of the dislocation pattern on the macroscopic resolved stress. However, as microbanding requires the activation of cross-slip at a shorter scale, its dependence on the carbon content is expected to be lower than that of dislocation patterning, where significant variations in cross-slip frequency are required to modify the dislocation patterning.


In summary, we have investigated the influence of the strain path, namely, tension and shear, on microbanding in a Fe-22Mn-0.6C (wt. %) steel. Evaluation of the dislocation substructure by combined EBSD and ECCI mapping revealed that microbanding is dependent on the strain path. We explain this effect in terms of Schmid’s law. Our observations reveal that microbanding in the present FeMnC alloy can be explained in terms of the dislocation boundary splitting mechanism. Examination of the internal dislocation configuration of microbands provides new insights on their role on strain hardening behavior of high-Mn steels.





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