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Microstructure evolution in dual-phase stainless steel during severe deformation
Acta Materialia 54 (2006) 2521–2532 www.actamat-journals.com
Microstructure evolution in dual-phase stainless steel during severe deformation A. Belyakov
*,1,
Y. Kimura, K. Tsuzaki
Steel Research Center, National Institute for Materials Science, Sengen 1-2-1, Tsukuba, Ibaraki 305-0047, Japan Received 2 August 2005; received in revised form 24 January 2006; accepted 25 January 2006 Available online 27 March 2006
Abstract Deformation microstructures of an Fe–27% Cr–9% Ni dual-phase stainless steel, which was bar rolled/swaged to a total strain of 6.9 at ambient temperature, were studied. After a rapid increase in the hardness during early deformation, the rate of strain hardening slowed and produced a steady-state-like deformation behaviour at strains above 4. The severe deformation resulted in the evolution of similar microstructures in both austenite and ferrite consisting of elongated (sub)grains with a final transverse size of about 0.1 lm and about 70% of high-angle (sub)boundaries. However, the different phases were characterised by different structural change kinetics. The ferrite transverse (sub)grain size decreased continuously, approaching its minimum at large strains above 5.0, while the distinct grain subdivision in the austenite reduced the transverse (sub)grain size to its final value quickly at an early processing stage. The main mechanism of microstructure evolution during the large strain processing was considered to be micro-shearing with dynamic recovery. 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Severe plastic deformation; EBSP; TEM; Stainless steel; Grain refining
1. Introduction Severe plastic deformations, i.e., large strain processing at relatively low temperatures, have been the subject of renewed interest recently for their use in the production of structural metallic materials with superior mechanical properties [1–7]. A beneficial combination of mechanical properties has been achieved by a specific structural state with ultrafine grains on a submicrometer scale [8–12]. The structural changes leading to the submicrocrystalline states during severe deformations can be associated with a kind of strain-induced continuous reaction, which includes the formation of dislocation subboundaries such as a dense dislocation wall and a gradual increase in the subboundary misorientations to typical values for conventional high-angle grain boundaries [3,4,13–17]. *
Corresponding author. Tel.: +81 298 59 2185; fax: +81 298 59 2101. E-mail address: [email protected] (A. Belyakov). 1 On leave from the Institute for Metals Superplasticity Problems, Ufa, Russia.
However, in contrast to plastic working at elevated temperatures, the relationships between the deformation conditions and the mechanisms that are responsible for the new fine-grain evolution during cold working have not been developed well. The effects of the nature of the processed material, its phase composition and deformation method on the final grain size and the grain refinement kinetics have not been studied in sufficient detail. The difficulties in studying the peculiarities of microstructure evolution upon large-strain cold working could result from the very complicated deformation methods, such as mechanical milling, torsion under high pressure, and equal-channel angular pressing, that are frequently used to provide severe deformations [18–25]. Some conventional processing methods like rolling and drawing could also result in rather large strains during cold working. These methods are quite simple for applications; therefore, they could be utilised easily for a simulation of severe deformation. An alloy with finer initial grains has been shown to demonstrate much faster kinetics of grain refinement, i.e., decreasing the (sub)grain size and increasing the fraction of high-angle
1359-6454/$30.00 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2006.01.035
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(sub)boundaries, than those of a material with coarser initial grains [26,27]. The initial fine-grained microstructure could be used to obtain steady-state-like deformation behaviour, when the main microstructural parameters do not change during cold working. However, the structural mechanisms that are responsible for such deformation behaviour have not been studied in sufficient detail. The aim of this study is to clarify the deformation mechanisms and the structural changes that are associated with the evolution of strain-induced (sub)grains during largestrain cold working by bar rolling/swaging. A dual-phase ductile alloy seems to be very useful for clarifying the structural mechanisms operating upon severe deformation, since the interphase boundaries serving like markers can provide a simultaneous structural analysis on different scale levels. The changes in phase shapes and the mutual arrangement of the different phases are related to the change in the shape of the whole sample during plastic working and provide important information on the plastic flow on macro- to mesoscales. Therefore, an austenite–ferrite stainless steel was selected as a typical representative of dual-phase ductile metallic materials. 2. Experimental A dual-phase stainless steel, Fe–0.017C–0.01Mn–0.01P– 0.001S–26.8Cr–9.03Ni–0.002N (mass%), was vacuum melted and cast into a 20 kg ingot. The steel was then hot forged and homogenised at 1200 C and warm rolled to 21.3 · 21.3 mm2 square bars at 700 C. The ferrite/austenite ratio was about 3/2 in the initial warm-rolled state. The transverse phase size for the austenite and the ferrite was measured crosswise to the rolling axis and was about 8 and 14 lm, respectively. The transverse (sub)grain size was about 0.4 lm, and the percentage of high-angle (sub)grain boundaries within the austenite phases was about 17%. In the ferrite phases, almost all of the subgrains with a transverse size of 0.7 lm were composed of low-angle subboundaries. The phase and (sub)grain sizes were evaluated by the linear intercept method, by counting the interphase boundaries and (sub)boundaries. The critical misorientation between the low- and high-angle (sub)grain boundaries was chosen at 15. Severe plastic deformation was carried out at ambient temperature by rolling to 7.8 · 7.8 mm2 square bars (strain of 2.0) followed by swaging from B7.0 to B0.6 mm, providing a total strain of 6.9. Structural investigations were performed on sections parallel to the rolling/swaging axis, using a JSM-6500 scanning electron microscope equipped with an electron back scattering diffraction pattern (EBSP) analyser incorporating an orientation imaging microscopy (OIM) system and a JEM-2010F transmission electron microscopy (TEM) instrument. The OIM images were subjected to a cleanup procedure with a minimal confidence index of 0.1. The phase and (sub)grain sizes were measured perpendicular to the rolling/swaging axis from the OIM and TEM images, respectively. Misorientations across the (sub)grain
boundaries were analysed by the conventional TEM Kikuchi line method [28], collecting about 100 (sub)grain boundaries per strain level, and by an EBSP technique with analytical software provided by TexSEM Lab. Inc. All the clearly defined (sub)grain boundaries were taken into account to determine the (sub)grain sizes and the (sub)boundary misorientations on the TEM images. The phase fractions were studied using X-ray diffraction [29], with a RINT-2500 X-ray diffractometer equipped with RINT 2000 series application software (developed by Rigaku Co.). Strain hardening was studied using Vickers hardness tests with a load of 3 N. 3. Results 3.1. Strain hardening The effect of the strain on hardness is shown in Fig. 1. The rate of strain hardening increases rapidly early in the deformation and gradually slows to some constant value at a strain of around 1.0. The hardness continues to increase with an almost constant increment with further cold deformation to a strain of about 5. The hardness then approaches a saturation value with little change during subsequent deformation and shows steady-state-like behaviour at larger strains. Similar deformation behaviour has also been observed quite often in metals and alloys with high stacking fault energies during deformation at elevated temperatures [30–32], when recovery processes resulted in a dynamic equilibrium between the generation and annihilation of strain-induced crystallographic defects, mainly dislocations. However, the dislocation rearrangements were suppressed upon cold working. In most studies, cold deformations were commonly characterised by the strengthening of materials during processing [33–36]. Therefore, the apparent steady-state deformation behaviour that was observed in this study with severe deformation may suggest a specific deformation mechanism operating at sufficiently large strains.
Fig. 1. Effect of cold bar rolling/swaging on the hardness of the dualphase stainless steel.
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3.2. Structural response 3.2.1. OIM Fig. 2 shows typical deformation microstructures that were developed in the steel under bar rolling/swaging to various strains. Note that the rolling/swaging direction is horizontal in all the micrographs that are shown in this
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paper. Generally, cold rolling/swaging results in the development of ribbon-like microstructures that are aligned along the deformation axis, which is quite typical for cold working by unidirectional deformation. A detailed analysis using the EBSP technique reveals a clear difference in the deformation structures between the austenite and the ferrite. As expected, the cold rolling/swaging results in the
Fig. 2. OIM images of the dual-phase stainless steel processed by bar rolling to strains (e) of: (a) e = 1.0; (b) e = 2.0 followed by swaging; (c) e = 3.2; (d) e = 4.4; (e) e = 5.8; (f) e = 6.9. (Sub)grain boundaries with misorientations above 15 are shown. The inverse pole figure relates to the rolling/swaging axis. The white arrows in (b) indicate the kinks on the interphase boundaries that resulted from micro-shearing.
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development of characteristic fibre textures with a rolling/ swaging axis corresponding to Æ1 1 0æ in the ferrite and Æ1 1 1æ or Æ2 0 0æ in the austenite. Therefore, the austenite and ferrite phases can be recognised easily by the different colours in the orientation microscopy images: the ferrite is green and the austenite is red or blue (Fig. 2). The austenite and ferrite phases can be clearly differentiated by their internal microstructures, especially at relatively small strains below 3 (Fig. 2(a) and (b)). The austenite has a much higher density of strain-induced (sub)boundaries with misorientations above 15 than the ferrite. The development of micro-shear bands, which can be recognised easily by the serration of interphase boundaries (some of the kinks on the interphase boundaries are shown by the white arrows in Fig. 2(b)), is considered to be another interesting feature of the deformation microstructure in the sample processed to a strain of 2.0. The micro-shearing in the austenite is frequently connected with clusters of fine (sub)grains that appear as chains along the shear bands. In contrast, the micro-shearing in the ferrite phases cannot be characterised by the associated straininduced (sub)boundaries, and results in an extremely limited number of large crystallographic rotations. The micro-shearing becomes less distinguishable upon further processing to larger strains (Fig. 2(c)–(f)). The homogeneous deformation microstructures consist of fine (sub)grains, which are highly elongated in the swaging direction, and the transverse sizes of austenite and ferrite phases gradually decrease with straining. The volume fraction of the austenite decreases considerably after processing to large strains above 5 and results in the appearance of elongated austenite phases as individual grains, which are surrounded by ferrite (Fig. 2(f)). The effect of strain on the austenite volume fraction obtained using EBSP and X-ray diffraction is shown in Fig. 3. The inset in this figure represents equilibrium austenite fractions, which were evaluated for various tempera-
Fig. 3. Variation of the austenite volume fraction in the dual-phase stainless steel with cold bar rolling/swaging obtained from EBSP and Xray diffraction. The austenite fraction evaluated by thermodynamic analysis is shown in the inset for reference.
tures by thermodynamic analysis, using the Thermo-Calc software TCW 3.1 (developed by the Foundation for Computational Thermodynamics). Thermodynamics predicts a decrease in the equilibrium austenite fraction on cooling to room temperature. However, the real phase fraction could barely follow the equilibrium one during cooling because the diffusivity of the alloying elements, mainly Cr and Ni, essentially decreases with temperature. Therefore, the austenite in the starting material can be considered as a metastable phase (the gain in the Gibbs energy is about 1.8 kJ/mol, if the alloying composition is invariant under the transformation). The austenite fraction of about 40% in the steel samples is almost constant during cold working to strains of 4–5. However, further processing to a total strain of 6.9 results in a rapid decrease in the percentage of austenite to a thermodynamically equilibrium value of about 10%, i.e., most of the austenite undergoes a c ! a martensitic transformation upon severe deformation. The critical strain for the strain-induced martensitic transformation in this study is considered to be around 4.5. Interestingly, the martensitic transformation at strains of 5–7 does not lead to a strengthening of the steel, but corresponds to the apparent steady-state behaviour (cf. Figs. 1 and 3).
Fig. 4. Typical substructures developed in the dual-phase stainless steel after cold bar rolling to a strain of 0.4.
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3.2.2. Fine substructures The strain-induced substructural changes depend on the intensity of plastic working. The deformation substructures that were developed in the austenite and ferrite phases at different strain levels will now be considered in more detail. The TEM images of the fine substructures that evolved at the early stage of deformation clearly show the development of multiple twinning in the austenite phase (Fig. 4). On the other hand, the ferrite phase is characterised by conventional cold-worked substructures, which consist of dislocation cell blocks that are separated by dense dislocation walls. The deformation twinning provides a rapid grain subdivision in the austenite phase, leading to the evolution of a fine fragmented substructure with high misorientations between fragments. The twin boundaries lose their distinguishing characteristics such as straight shapes, close parallel locations, and misorientations quickly with straining as shown in Fig. 5. The dislocation substructures in the austenite appear as very fine irregular fragments with rather thick diffuse (sub)boundaries having a wide spec-
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trum of low- to high-angle misorientations. In contrast, the ferrite substructures do not change much with an increase in the strain to 1.0 and are composed of distinguishable subgrains that are outlined by low-angle dislocation subboundaries. The dislocation density in the ferrite is also much lower than that in the austenite. An increase in the strain to 2.0 is accompanied by the development of micro-shearing leading to the serration of the interphase boundaries (the latter is shown by the dotted line in Fig. 6). The white arrows indicate the probable arrangement of the micro-shears resulting in the phase boundary serration. The austenite substructures that have evolved in the micro-shear bands look like ultrafine cells with a high density of dislocations. The boundaries between such substructural elements cannot be distinguished precisely. Contrary to the ordinary cell substructure, the present one is characterised by large misorientations, which are more typical of a grain structure. In the ferrite phase, the micro-shearing hardly occurs in the evolution of a number of largely misoriented fragments. The micro-shearing
Fig. 5. Deformation substructures developed in the dual-phase stainless steel after cold bar rolling to a strain of 1.0. The numbers indicate the (sub)grain misorientations in degrees. The dotted line shows an interphase boundary.
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mainly corrugates the ribbon-like ferrite subgrains and leads to the formation of a wavy lamellar substructure. Further processing to large strains above 4 eliminates the differences that were present between the substructures of ferrite and austenite (Figs. 7 and 8). Both phases are characterised by the development of highly elongated (sub)grains with a large fraction of high-angle (sub)boundaries. The austenite in Figs. 7 and 8 could only be recognised by the electron diffraction pattern. The severe deformation results in a martensitic transformation, causing the amount of original ferrite to increase by straininduced ferrite. The strain-induced ferrite in Fig. 8 was determined by X-ray energy-dispersive spectrometry as ferrite (sub)grains having a larger Ni/Cr ratio than the nominal one of 0.33. Fig. 8 shows that the substructure of the strain-induced ferrite and the original one do not differ. Therefore, there is no significant difference between the fine substructures of austenite and the ferrite including the strain-induced ferrite after severe deformation. Figs. 9 and 10 show the quantitative results for the misorientations that evolved in the fine substructures for the
austenite and the ferrite, respectively. The dashed lines in these figures indicate data that were obtained from OIM. The results obtained using the TEM and OIM techniques match well within experimental scatter. Deformation twinning, which starts to operate in the austenite at relatively low strains, results in the development of a large number of high-angle (sub)boundaries. The corresponding misorientation distribution is bimodal and shows high peaks at low angles below 15 which are associated with dislocation subboundaries and at high angles of around 60 resulting from the deformation twins. The twinning effect on the misorientations in the deformation substructures becomes less distinguishable upon further processing. After severe deformation, the evolution of strain-induced (sub)boundaries that are parallel to the rolling/swaging axis can be characterised by almost equal fractions of various misorientations. Misorientations among the ferrite substructures increase gradually with cold working. The sharp peak from the low-angle dislocation subboundaries decreases and spreads out towards the larger misorientations. The misorientation distribution is also characterised by a maximum at
Fig. 6. Deformation substructures developed in the dual-phase stainless steel after cold bar rolling to a strain of 2.0. The numbers indicate the (sub)grain misorientations in degrees. The dotted line corresponds to a serrated interphase boundary. The white arrows indicate the directions of micro-shears leading to the serrations of the interphase boundary.
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of 0.1 lm at large strains above 5.0. The refinement of the (sub)grains is accompanied by an increase in the fraction of high-angle (sub)boundaries. The latter also approaches a saturation of about 70% for large strains above 5. Such variations in the transverse (sub)grain size and the (sub)boundary misorientations are quite similar to those reported in many studies of the cold rolling/drawing of a single-phase ferrite with large initial grains [27,33– 35]. In contrast, the deformation twinning results in a sudden decrease in the austenite (sub)grain size to about 0.1 lm and a corresponding increase in the fraction of HAB to above 50% at an early processing stage. The transverse (sub)grain size in the austenite does not change much during subsequent deformation, while the HAB fraction slowly increases to the same saturation value as that in the ferrite. The increase in the dislocation density during the processing also follows the hardness curve. The value of q for the austenite is several times larger than that for the ferrite in all processing stages. 3.3. Deformation textures
Fig. 7. Deformation substructures developed in the dual-phase stainless steel after cold bar rolling/swaging to a strain of 4.4. The numbers indicate the (sub)grain misorientations in degrees. The dotted lines correspond to interphase boundaries.
large angles of around 60 at large strains above 4, which can be associated with an evolution of a specific deformation texture [27,37] as well as a strain-induced martensitic transformation occurring under severe deformation. Fig. 11 summarises the effect of strain on three microstructural parameters, the transverse (sub)grain size (d), the number fraction of high-angle (sub)boundaries (HAB) and the dislocation density in (sub)grain interiors (q), for the cold-worked steel samples that were obtained from TEM. Variations of these microstructural parameters during processing generally correlate with the strain hardening; in other words, after rapid changes occur at an early processing stage, gradual saturations take place after sufficient strains. However, the structural changes in the ferrite and the austenite are characterised by different kinetics. The transverse (sub)grain size in the ferrite decreases rapidly to about 0.3 lm at relatively low strains below 1.0, then decreases continuously and approaches a saturation
The bar rolling/swaging results in the development of strong fibre textures in both the austenite and ferrite phases. Fig. 12 shows the inverse pole figures for the rolling/swaging direction (RD). The data were obtained using the EBSP technique. A rather homogeneous deformation of the ferrite is characterised by the evolution of a unique Æ0 1 1æ i RD texture. The texture intensity increases with plastic working to strains of about 6 then becomes somewhat randomised at the largest strain studied. In the austenite phase, the operation of localised deformation modes, consisting of twinning and micro-shearing, is accompanied by the development of two texture components, Æ1 1 1æ i RD and Æ0 0 1æ i RD, and their intensities increase during straining to moderate strains of about 4. Strain localisation in the austenite is suppressed on processing to larger strains. Correspondingly, the Æ0 0 1æ i RD fibre intensity decreases, whereas the Æ1 1 1æ i RD texture component randomises only after severe deformation. The apparent randomisation in the austenite and ferrite textures after processing to a large strain of 6.9 may be associated with the sudden decrease in the austenite volume fraction (initially from 40% to about 10%) caused by the strain-induced martensitic transformation. 4. Discussion Cold working was accompanied by continuous strain hardening in many studies of cold rolling/drawing [33– 36]. The deformation microstructures were shown to consist of highly elongated subgrains, where the transverse size decreased with increasing strain. Steady-state deformation behaviour was only observed for cold working in the samples with an original fine-grained martensitic microstructure [27]. In the present dual-phase stainless steel, variations of the main structural parameters in the different
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Fig. 8. Deformation substructures developed in the dual-phase stainless steel after cold bar rolling/swaging to a strain of 6.9. The numbers indicate the (sub)grain misorientations in degrees. The white, grey and black areas correspond to ferrite, strain-induced ferrite and austenite, respectively. The white arrows bound the continual sequences of (sub)boundaries, which are skewed with respect to the swaging axis and can originate from the micro-shears.
phases were peculiar to both conventional work hardening and steady-state deformation behaviour. The ferrite showed a standard sequence for the structural changes during the cold working. The transverse (sub)grain size gradually decreased and the (sub)grain misorientations increased, and approached their apparent saturation at large strains above 5. In contrast, the transverse (sub)grain size in the austenite phase quickly decreased to its final minimal value, while the (sub)grain misorentations rapidly increased after surprisingly small strains. This resulted in a
steady-state-like deformation behaviour during subsequent straining. Generally, the transverse (sub)grain sizes during uniaxial cold deformation should decrease due to grain subdivision by the strain-induced (sub)boundaries as well as to a change in the shape of the whole sample [3,27,38–41]. However, the operation of specific structural mechanisms including shear banding [42–45], subgrain coalescence by the plastic hinge mechanism [34], the collapse of a ribbon-like substructure with local migration and the merger
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Fig. 9. Misorientation distributions for the austenite (sub)boundaries in the dual-phase stainless steel subjected to various strains (e). The bars and the dashed lines represent the TEM and OIM data, respectively.
of adjacent subboundaries [46,47] may delay the reduction of transverse (sub)grain size significantly during cold working. The mechanisms of microstructure evolution during large-strain deformation can be clarified using the ductile dual-phase material with various deformation behaviours for different phases. Fig. 13 shows the strain dependence for the structural parameters of transverse (sub)grain size and transverse phase size. Assuming that the reduction of the (sub)grains and the phases during the cold bar rolling/swaging corresponds to that of the whole sample, the transverse sizes (d) should vary with the strain (e) as ln d 0.5e. In the austenite phase, grain subdivision by multiple twinning at the early stage of deformation results in a rapid decrease of the transverse (sub)grain size to some value, which does not change with further processing. In the ferrite phase, the transverse (sub)grain size basically follows the change in the shape of the sample during processing to strains of about 2, whereas further deformation results in a (sub)grain size larger than that predicted by the change in the shape of the sample. In other words, the grain subdivision is more evident in the phase that is characterised by a higher hardenability. A rapid increase in the dislocation density limits the homogeneous deformation of the austenite and activates some localised deformation modes including twinning and micro-shearing. The ability of ferrite to deform homogeneously is greater. Therefore, the
Fig. 10. Misorientation distributions for the ferrite (sub)boundaries in the dual-phase stainless steel subjected to various strains (e). The bars and the dashed lines represent the TEM and OIM data, respectively.
microstructure evolution in the ferrite to moderate strains is mainly associated with the elongation of the initial (sub)grains along the metal flow direction which is accompanied by increasing (sub)boundary misorientations. Fig. 13 shows that the dependence of the transverse austenite/ferrite phase sizes on the strain and the change in the shape of the samples correspond well during processing to strains of about 4.5, where the phase composition no longer shows a marked change. Since the shear bands definitely develop in the same strain range (see Figs. 2 and 6), a rather high density of micro-shears is suggested; in other words, the presence of a relatively high level of deformation homogeneity on a macroscopic scale. In fact, if the spacing between the intersecting shears is much smaller than the phase size, the micro-shearing will affect the change in the transverse phase size in the same way as conventional dislocation motion, although it can retard the change in the transverse (sub)grain size. The clear correlation of the change in transverse phase sizes in the austenite and the ferrite is indicative of a uniform deformation across the samples. The micro-shears propagate over large distances (at least larger than the transverse dimensions of
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Fig. 11. Effect of strain on the transverse (sub)grain size (d), the number fraction of HAB and the dislocation density (q) in (sub)grain interiors that developed in the cold-worked dual-phase stainless steel.
individual phases); thus, the ferrite and the austenite experience the same strains, regardless of the different levels of strengthening. If the strain distribution is assumed to be uniform among all the (sub)grains across the sample, (sub)boundary migration that leads to grain coarsening and/or a micro-shearing can be used to show that the transverse (sub)grain size is constant during processing. The microshearing seems to be more adequate for the cold working in this study. Therefore, the scheme for the mechanism of the microstructure evolution that operates in ribbon-like substructures at large strains during cold rolling/swaging [27] can be clarified with an emphasis on micro-shearing and is shown in Fig. 14. Local sliding along some transverse (sub)boundaries can be considered as a source of stress concentration, thus creating conditions for strain localisation. The sequential micro-shearing on a microscale level (individual subgrains) that is associated with local (sub)boundary sliding results in a strain localisation on a mesoscale level (several subgrains). Upon unidirectional deformation, the shear planes tend to turn towards the metal flow, which is the deformation axis. A high density of such intersecting micro-shears provides an apparently homogeneous deformation throughout the sample and does not decrease the transverse (sub)grain size. The micro-shears that are involved in sliding along the (sub)boundaries can hardly be recognised on the micrographs, since the shear directions are close to the deformation axis. The arrows in Fig. 8 indicate the possible micro-shears. Various metallic materials subjected to severe deformations were characterised by a rapid recovery releasing internal stresses associated with strain-induced (sub)boundaries [4,48–52]; thus dynamic recovery can play
Fig. 12. Inverse pole figures for the rolling/swaging direction of the dualphase stainless steel. The data were obtained separately for the austenite and ferrite phases using the EBSP method.
Fig. 13. Reduction of the transverse (sub)grain and phase sizes in the dual-phase stainless steel during cold bar rolling/swaging.
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fraction of high-angle (sub)boundaries increased to more than 50%. No more large changes occurred in the transverse (sub)grain size during further processing up to the largest strain studied. Such a structural response was associated with deformation twinning resulting in fast grain subdivision at the early stage of deformation and micro-shearing, which provided uniform macroscopic deformation without any changes in the transverse (sub)grain size at large strains. 4. Finally, severe deformation to a total strain of about 6 resulted in the development of similar microstructures in both the austenite and ferrite phases. The microshearing assisted by the recovery processes was considered as the structural mechanism operating during large-strain cold working.
Acknowledgements
Fig. 14. Schematic drawing of the micro-shears operating in the ribbonlike substructures upon severe unidirectional deformation.
an important role in the deformation behaviour even at relatively low temperatures. In this discussion, the dynamic recovery promoted the micro-shears that propagated partially along the (sub)boundaries and is considered as the mechanism that is responsible for the steady-state-like deformation behaviour at large strains. 5. Conclusions The microstructure evolution in an Fe–27% Cr–9% Ni dual-phase stainless steel during large-strain deformation by bar rolling/swaging to a total strain of 6.9 was studied. The main results are summarised as follows: 1. After a rapid strain hardening at the early stage of deformation, the rate of strain hardening gradually decreased to almost zero at large strains of about 4, leading to a steady-state-like deformation behaviour during large-strain cold working. 2. In the ferrite phase, the transverse (sub)grain size decreased gradually during processing, approaching approximately 0.1 lm at strains above 6.0. The (sub)grain reduction was accompanied by an increase in the (sub)boundary misorientations. Correspondingly, the number fraction of high-angle (sub)boundaries increased to a saturation of about 70%. 3. In the austenite phase, the transverse (sub)grain size decreased rapidly to its minimal value of about 0.1 lm at relatively small strains of about 1, while the number
The authors are grateful to Drs. N. Sakuma, T. Hibaru, S. Kuroda and M. Kobayashi of the Steel Research Center, National Institute for Materials Science, for their assistance in the materials processing and to Ms. J. Hono, National Institute for Materials Science, for improving the language of the paper. One of the authors (A.B.) would like to express his thanks to the National Institute for Materials Science for providing a scientific fellowship. References [1] Wadsworth J, Sherby OD. Progr Mater Sci 1980;25:35. [2] Mishin OV, Gottstein G. Philos Mag A 1998;78:373. [3] Humphreys FJ, Prangnell PB, Bowen JR, Gholinia A, Harris C. Phil Trans R Soc Lond 1999;357:1663. [4] Valiev RZ, Islamgaliev RK, Alexandrov IV. Progr Mater Sci 2000;45:103. [5] Wang Y, Chen M, Zhou F, Ma E. Nature 2002;419:912. [6] Mughrabi H, Hoppel HW, Kautz M, Valiev RZ. Z Metallkd 2003;94:1079. [7] Ohsaki S, Hono K, Hidaka H, Takaki S. Script Mater 2005;52:271. [8] Suryanarayana C. Int Mater Rev 1995;40:41. [9] Morris DG. In: Dinesen AR, Eldrup M, Juul Jensen D, Linderoth S, Pedersen TB, Pryds NH, Pedersen SA, Wert JA, editors. Science of metastable and nanocrystalline alloys. Roskilde, Denmark: Riso National Laboratory; 2001. p. 89. [10] Umemoto M, Liu ZG, Masuyama K, Hao XJ, Tsuchiya K. Script Mater 2001;44:1741. [11] Valiev RZ, Alexandrov IV, Zhu YT, Lowe TC. J Mater Res 2002;17:5. [12] Tsuji N, Ueji R, Minamino Y, Saito Y. Script Mater 2002;46:305. [13] Kaibyshev R, Sitdikov O. Z Metallkd 1994;85:738. [14] Belyakov A, Sakai T, Miura H. Mater Trans JIM 2000;41:476. [15] Belyakov A, Sakai T, Miura H, Tsuzaki K. Philos Mag A 2001;81:2629. [16] Ivanisenko Yu, Lojkowski W, Valiev RZ, Fecht H-J. Acta Mater 2003;51:5555. [17] Xu C, Furukawa M, Horita Z, Langdon TG. Mater Sci Eng A 2005;A398:66. [18] Segal VM, Reznikov VI, Drobyshevskiy AE, Kopylov VI. Russian Metall 1981;1:115. [19] Saunders I, Nutting J. Met Sci 1984;18:571. [20] Richert J, Richert M. Aluminium 1986;62:604.
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Microstructure evolution in dual-phase stainless steel during severe deformation A. Belyakov
*,1,
Y. Kimura, K. Tsuzaki
Steel Research Center, National Institute for Materials Science, Sengen 1-2-1, Tsukuba, Ibaraki 305-0047, Japan Received 2 August 2005; received in revised form 24 January 2006; accepted 25 January 2006 Available online 27 March 2006
Abstract Deformation microstructures of an Fe–27% Cr–9% Ni dual-phase stainless steel, which was bar rolled/swaged to a total strain of 6.9 at ambient temperature, were studied. After a rapid increase in the hardness during early deformation, the rate of strain hardening slowed and produced a steady-state-like deformation behaviour at strains above 4. The severe deformation resulted in the evolution of similar microstructures in both austenite and ferrite consisting of elongated (sub)grains with a final transverse size of about 0.1 lm and about 70% of high-angle (sub)boundaries. However, the different phases were characterised by different structural change kinetics. The ferrite transverse (sub)grain size decreased continuously, approaching its minimum at large strains above 5.0, while the distinct grain subdivision in the austenite reduced the transverse (sub)grain size to its final value quickly at an early processing stage. The main mechanism of microstructure evolution during the large strain processing was considered to be micro-shearing with dynamic recovery. 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Severe plastic deformation; EBSP; TEM; Stainless steel; Grain refining
1. Introduction Severe plastic deformations, i.e., large strain processing at relatively low temperatures, have been the subject of renewed interest recently for their use in the production of structural metallic materials with superior mechanical properties [1–7]. A beneficial combination of mechanical properties has been achieved by a specific structural state with ultrafine grains on a submicrometer scale [8–12]. The structural changes leading to the submicrocrystalline states during severe deformations can be associated with a kind of strain-induced continuous reaction, which includes the formation of dislocation subboundaries such as a dense dislocation wall and a gradual increase in the subboundary misorientations to typical values for conventional high-angle grain boundaries [3,4,13–17]. *
Corresponding author. Tel.: +81 298 59 2185; fax: +81 298 59 2101. E-mail address: [email protected] (A. Belyakov). 1 On leave from the Institute for Metals Superplasticity Problems, Ufa, Russia.
However, in contrast to plastic working at elevated temperatures, the relationships between the deformation conditions and the mechanisms that are responsible for the new fine-grain evolution during cold working have not been developed well. The effects of the nature of the processed material, its phase composition and deformation method on the final grain size and the grain refinement kinetics have not been studied in sufficient detail. The difficulties in studying the peculiarities of microstructure evolution upon large-strain cold working could result from the very complicated deformation methods, such as mechanical milling, torsion under high pressure, and equal-channel angular pressing, that are frequently used to provide severe deformations [18–25]. Some conventional processing methods like rolling and drawing could also result in rather large strains during cold working. These methods are quite simple for applications; therefore, they could be utilised easily for a simulation of severe deformation. An alloy with finer initial grains has been shown to demonstrate much faster kinetics of grain refinement, i.e., decreasing the (sub)grain size and increasing the fraction of high-angle
1359-6454/$30.00 2006 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2006.01.035
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(sub)boundaries, than those of a material with coarser initial grains [26,27]. The initial fine-grained microstructure could be used to obtain steady-state-like deformation behaviour, when the main microstructural parameters do not change during cold working. However, the structural mechanisms that are responsible for such deformation behaviour have not been studied in sufficient detail. The aim of this study is to clarify the deformation mechanisms and the structural changes that are associated with the evolution of strain-induced (sub)grains during largestrain cold working by bar rolling/swaging. A dual-phase ductile alloy seems to be very useful for clarifying the structural mechanisms operating upon severe deformation, since the interphase boundaries serving like markers can provide a simultaneous structural analysis on different scale levels. The changes in phase shapes and the mutual arrangement of the different phases are related to the change in the shape of the whole sample during plastic working and provide important information on the plastic flow on macro- to mesoscales. Therefore, an austenite–ferrite stainless steel was selected as a typical representative of dual-phase ductile metallic materials. 2. Experimental A dual-phase stainless steel, Fe–0.017C–0.01Mn–0.01P– 0.001S–26.8Cr–9.03Ni–0.002N (mass%), was vacuum melted and cast into a 20 kg ingot. The steel was then hot forged and homogenised at 1200 C and warm rolled to 21.3 · 21.3 mm2 square bars at 700 C. The ferrite/austenite ratio was about 3/2 in the initial warm-rolled state. The transverse phase size for the austenite and the ferrite was measured crosswise to the rolling axis and was about 8 and 14 lm, respectively. The transverse (sub)grain size was about 0.4 lm, and the percentage of high-angle (sub)grain boundaries within the austenite phases was about 17%. In the ferrite phases, almost all of the subgrains with a transverse size of 0.7 lm were composed of low-angle subboundaries. The phase and (sub)grain sizes were evaluated by the linear intercept method, by counting the interphase boundaries and (sub)boundaries. The critical misorientation between the low- and high-angle (sub)grain boundaries was chosen at 15. Severe plastic deformation was carried out at ambient temperature by rolling to 7.8 · 7.8 mm2 square bars (strain of 2.0) followed by swaging from B7.0 to B0.6 mm, providing a total strain of 6.9. Structural investigations were performed on sections parallel to the rolling/swaging axis, using a JSM-6500 scanning electron microscope equipped with an electron back scattering diffraction pattern (EBSP) analyser incorporating an orientation imaging microscopy (OIM) system and a JEM-2010F transmission electron microscopy (TEM) instrument. The OIM images were subjected to a cleanup procedure with a minimal confidence index of 0.1. The phase and (sub)grain sizes were measured perpendicular to the rolling/swaging axis from the OIM and TEM images, respectively. Misorientations across the (sub)grain
boundaries were analysed by the conventional TEM Kikuchi line method [28], collecting about 100 (sub)grain boundaries per strain level, and by an EBSP technique with analytical software provided by TexSEM Lab. Inc. All the clearly defined (sub)grain boundaries were taken into account to determine the (sub)grain sizes and the (sub)boundary misorientations on the TEM images. The phase fractions were studied using X-ray diffraction [29], with a RINT-2500 X-ray diffractometer equipped with RINT 2000 series application software (developed by Rigaku Co.). Strain hardening was studied using Vickers hardness tests with a load of 3 N. 3. Results 3.1. Strain hardening The effect of the strain on hardness is shown in Fig. 1. The rate of strain hardening increases rapidly early in the deformation and gradually slows to some constant value at a strain of around 1.0. The hardness continues to increase with an almost constant increment with further cold deformation to a strain of about 5. The hardness then approaches a saturation value with little change during subsequent deformation and shows steady-state-like behaviour at larger strains. Similar deformation behaviour has also been observed quite often in metals and alloys with high stacking fault energies during deformation at elevated temperatures [30–32], when recovery processes resulted in a dynamic equilibrium between the generation and annihilation of strain-induced crystallographic defects, mainly dislocations. However, the dislocation rearrangements were suppressed upon cold working. In most studies, cold deformations were commonly characterised by the strengthening of materials during processing [33–36]. Therefore, the apparent steady-state deformation behaviour that was observed in this study with severe deformation may suggest a specific deformation mechanism operating at sufficiently large strains.
Fig. 1. Effect of cold bar rolling/swaging on the hardness of the dualphase stainless steel.
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3.2. Structural response 3.2.1. OIM Fig. 2 shows typical deformation microstructures that were developed in the steel under bar rolling/swaging to various strains. Note that the rolling/swaging direction is horizontal in all the micrographs that are shown in this
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paper. Generally, cold rolling/swaging results in the development of ribbon-like microstructures that are aligned along the deformation axis, which is quite typical for cold working by unidirectional deformation. A detailed analysis using the EBSP technique reveals a clear difference in the deformation structures between the austenite and the ferrite. As expected, the cold rolling/swaging results in the
Fig. 2. OIM images of the dual-phase stainless steel processed by bar rolling to strains (e) of: (a) e = 1.0; (b) e = 2.0 followed by swaging; (c) e = 3.2; (d) e = 4.4; (e) e = 5.8; (f) e = 6.9. (Sub)grain boundaries with misorientations above 15 are shown. The inverse pole figure relates to the rolling/swaging axis. The white arrows in (b) indicate the kinks on the interphase boundaries that resulted from micro-shearing.
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development of characteristic fibre textures with a rolling/ swaging axis corresponding to Æ1 1 0æ in the ferrite and Æ1 1 1æ or Æ2 0 0æ in the austenite. Therefore, the austenite and ferrite phases can be recognised easily by the different colours in the orientation microscopy images: the ferrite is green and the austenite is red or blue (Fig. 2). The austenite and ferrite phases can be clearly differentiated by their internal microstructures, especially at relatively small strains below 3 (Fig. 2(a) and (b)). The austenite has a much higher density of strain-induced (sub)boundaries with misorientations above 15 than the ferrite. The development of micro-shear bands, which can be recognised easily by the serration of interphase boundaries (some of the kinks on the interphase boundaries are shown by the white arrows in Fig. 2(b)), is considered to be another interesting feature of the deformation microstructure in the sample processed to a strain of 2.0. The micro-shearing in the austenite is frequently connected with clusters of fine (sub)grains that appear as chains along the shear bands. In contrast, the micro-shearing in the ferrite phases cannot be characterised by the associated straininduced (sub)boundaries, and results in an extremely limited number of large crystallographic rotations. The micro-shearing becomes less distinguishable upon further processing to larger strains (Fig. 2(c)–(f)). The homogeneous deformation microstructures consist of fine (sub)grains, which are highly elongated in the swaging direction, and the transverse sizes of austenite and ferrite phases gradually decrease with straining. The volume fraction of the austenite decreases considerably after processing to large strains above 5 and results in the appearance of elongated austenite phases as individual grains, which are surrounded by ferrite (Fig. 2(f)). The effect of strain on the austenite volume fraction obtained using EBSP and X-ray diffraction is shown in Fig. 3. The inset in this figure represents equilibrium austenite fractions, which were evaluated for various tempera-
Fig. 3. Variation of the austenite volume fraction in the dual-phase stainless steel with cold bar rolling/swaging obtained from EBSP and Xray diffraction. The austenite fraction evaluated by thermodynamic analysis is shown in the inset for reference.
tures by thermodynamic analysis, using the Thermo-Calc software TCW 3.1 (developed by the Foundation for Computational Thermodynamics). Thermodynamics predicts a decrease in the equilibrium austenite fraction on cooling to room temperature. However, the real phase fraction could barely follow the equilibrium one during cooling because the diffusivity of the alloying elements, mainly Cr and Ni, essentially decreases with temperature. Therefore, the austenite in the starting material can be considered as a metastable phase (the gain in the Gibbs energy is about 1.8 kJ/mol, if the alloying composition is invariant under the transformation). The austenite fraction of about 40% in the steel samples is almost constant during cold working to strains of 4–5. However, further processing to a total strain of 6.9 results in a rapid decrease in the percentage of austenite to a thermodynamically equilibrium value of about 10%, i.e., most of the austenite undergoes a c ! a martensitic transformation upon severe deformation. The critical strain for the strain-induced martensitic transformation in this study is considered to be around 4.5. Interestingly, the martensitic transformation at strains of 5–7 does not lead to a strengthening of the steel, but corresponds to the apparent steady-state behaviour (cf. Figs. 1 and 3).
Fig. 4. Typical substructures developed in the dual-phase stainless steel after cold bar rolling to a strain of 0.4.
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3.2.2. Fine substructures The strain-induced substructural changes depend on the intensity of plastic working. The deformation substructures that were developed in the austenite and ferrite phases at different strain levels will now be considered in more detail. The TEM images of the fine substructures that evolved at the early stage of deformation clearly show the development of multiple twinning in the austenite phase (Fig. 4). On the other hand, the ferrite phase is characterised by conventional cold-worked substructures, which consist of dislocation cell blocks that are separated by dense dislocation walls. The deformation twinning provides a rapid grain subdivision in the austenite phase, leading to the evolution of a fine fragmented substructure with high misorientations between fragments. The twin boundaries lose their distinguishing characteristics such as straight shapes, close parallel locations, and misorientations quickly with straining as shown in Fig. 5. The dislocation substructures in the austenite appear as very fine irregular fragments with rather thick diffuse (sub)boundaries having a wide spec-
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trum of low- to high-angle misorientations. In contrast, the ferrite substructures do not change much with an increase in the strain to 1.0 and are composed of distinguishable subgrains that are outlined by low-angle dislocation subboundaries. The dislocation density in the ferrite is also much lower than that in the austenite. An increase in the strain to 2.0 is accompanied by the development of micro-shearing leading to the serration of the interphase boundaries (the latter is shown by the dotted line in Fig. 6). The white arrows indicate the probable arrangement of the micro-shears resulting in the phase boundary serration. The austenite substructures that have evolved in the micro-shear bands look like ultrafine cells with a high density of dislocations. The boundaries between such substructural elements cannot be distinguished precisely. Contrary to the ordinary cell substructure, the present one is characterised by large misorientations, which are more typical of a grain structure. In the ferrite phase, the micro-shearing hardly occurs in the evolution of a number of largely misoriented fragments. The micro-shearing
Fig. 5. Deformation substructures developed in the dual-phase stainless steel after cold bar rolling to a strain of 1.0. The numbers indicate the (sub)grain misorientations in degrees. The dotted line shows an interphase boundary.
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mainly corrugates the ribbon-like ferrite subgrains and leads to the formation of a wavy lamellar substructure. Further processing to large strains above 4 eliminates the differences that were present between the substructures of ferrite and austenite (Figs. 7 and 8). Both phases are characterised by the development of highly elongated (sub)grains with a large fraction of high-angle (sub)boundaries. The austenite in Figs. 7 and 8 could only be recognised by the electron diffraction pattern. The severe deformation results in a martensitic transformation, causing the amount of original ferrite to increase by straininduced ferrite. The strain-induced ferrite in Fig. 8 was determined by X-ray energy-dispersive spectrometry as ferrite (sub)grains having a larger Ni/Cr ratio than the nominal one of 0.33. Fig. 8 shows that the substructure of the strain-induced ferrite and the original one do not differ. Therefore, there is no significant difference between the fine substructures of austenite and the ferrite including the strain-induced ferrite after severe deformation. Figs. 9 and 10 show the quantitative results for the misorientations that evolved in the fine substructures for the
austenite and the ferrite, respectively. The dashed lines in these figures indicate data that were obtained from OIM. The results obtained using the TEM and OIM techniques match well within experimental scatter. Deformation twinning, which starts to operate in the austenite at relatively low strains, results in the development of a large number of high-angle (sub)boundaries. The corresponding misorientation distribution is bimodal and shows high peaks at low angles below 15 which are associated with dislocation subboundaries and at high angles of around 60 resulting from the deformation twins. The twinning effect on the misorientations in the deformation substructures becomes less distinguishable upon further processing. After severe deformation, the evolution of strain-induced (sub)boundaries that are parallel to the rolling/swaging axis can be characterised by almost equal fractions of various misorientations. Misorientations among the ferrite substructures increase gradually with cold working. The sharp peak from the low-angle dislocation subboundaries decreases and spreads out towards the larger misorientations. The misorientation distribution is also characterised by a maximum at
Fig. 6. Deformation substructures developed in the dual-phase stainless steel after cold bar rolling to a strain of 2.0. The numbers indicate the (sub)grain misorientations in degrees. The dotted line corresponds to a serrated interphase boundary. The white arrows indicate the directions of micro-shears leading to the serrations of the interphase boundary.
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of 0.1 lm at large strains above 5.0. The refinement of the (sub)grains is accompanied by an increase in the fraction of high-angle (sub)boundaries. The latter also approaches a saturation of about 70% for large strains above 5. Such variations in the transverse (sub)grain size and the (sub)boundary misorientations are quite similar to those reported in many studies of the cold rolling/drawing of a single-phase ferrite with large initial grains [27,33– 35]. In contrast, the deformation twinning results in a sudden decrease in the austenite (sub)grain size to about 0.1 lm and a corresponding increase in the fraction of HAB to above 50% at an early processing stage. The transverse (sub)grain size in the austenite does not change much during subsequent deformation, while the HAB fraction slowly increases to the same saturation value as that in the ferrite. The increase in the dislocation density during the processing also follows the hardness curve. The value of q for the austenite is several times larger than that for the ferrite in all processing stages. 3.3. Deformation textures
Fig. 7. Deformation substructures developed in the dual-phase stainless steel after cold bar rolling/swaging to a strain of 4.4. The numbers indicate the (sub)grain misorientations in degrees. The dotted lines correspond to interphase boundaries.
large angles of around 60 at large strains above 4, which can be associated with an evolution of a specific deformation texture [27,37] as well as a strain-induced martensitic transformation occurring under severe deformation. Fig. 11 summarises the effect of strain on three microstructural parameters, the transverse (sub)grain size (d), the number fraction of high-angle (sub)boundaries (HAB) and the dislocation density in (sub)grain interiors (q), for the cold-worked steel samples that were obtained from TEM. Variations of these microstructural parameters during processing generally correlate with the strain hardening; in other words, after rapid changes occur at an early processing stage, gradual saturations take place after sufficient strains. However, the structural changes in the ferrite and the austenite are characterised by different kinetics. The transverse (sub)grain size in the ferrite decreases rapidly to about 0.3 lm at relatively low strains below 1.0, then decreases continuously and approaches a saturation
The bar rolling/swaging results in the development of strong fibre textures in both the austenite and ferrite phases. Fig. 12 shows the inverse pole figures for the rolling/swaging direction (RD). The data were obtained using the EBSP technique. A rather homogeneous deformation of the ferrite is characterised by the evolution of a unique Æ0 1 1æ i RD texture. The texture intensity increases with plastic working to strains of about 6 then becomes somewhat randomised at the largest strain studied. In the austenite phase, the operation of localised deformation modes, consisting of twinning and micro-shearing, is accompanied by the development of two texture components, Æ1 1 1æ i RD and Æ0 0 1æ i RD, and their intensities increase during straining to moderate strains of about 4. Strain localisation in the austenite is suppressed on processing to larger strains. Correspondingly, the Æ0 0 1æ i RD fibre intensity decreases, whereas the Æ1 1 1æ i RD texture component randomises only after severe deformation. The apparent randomisation in the austenite and ferrite textures after processing to a large strain of 6.9 may be associated with the sudden decrease in the austenite volume fraction (initially from 40% to about 10%) caused by the strain-induced martensitic transformation. 4. Discussion Cold working was accompanied by continuous strain hardening in many studies of cold rolling/drawing [33– 36]. The deformation microstructures were shown to consist of highly elongated subgrains, where the transverse size decreased with increasing strain. Steady-state deformation behaviour was only observed for cold working in the samples with an original fine-grained martensitic microstructure [27]. In the present dual-phase stainless steel, variations of the main structural parameters in the different
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Fig. 8. Deformation substructures developed in the dual-phase stainless steel after cold bar rolling/swaging to a strain of 6.9. The numbers indicate the (sub)grain misorientations in degrees. The white, grey and black areas correspond to ferrite, strain-induced ferrite and austenite, respectively. The white arrows bound the continual sequences of (sub)boundaries, which are skewed with respect to the swaging axis and can originate from the micro-shears.
phases were peculiar to both conventional work hardening and steady-state deformation behaviour. The ferrite showed a standard sequence for the structural changes during the cold working. The transverse (sub)grain size gradually decreased and the (sub)grain misorientations increased, and approached their apparent saturation at large strains above 5. In contrast, the transverse (sub)grain size in the austenite phase quickly decreased to its final minimal value, while the (sub)grain misorentations rapidly increased after surprisingly small strains. This resulted in a
steady-state-like deformation behaviour during subsequent straining. Generally, the transverse (sub)grain sizes during uniaxial cold deformation should decrease due to grain subdivision by the strain-induced (sub)boundaries as well as to a change in the shape of the whole sample [3,27,38–41]. However, the operation of specific structural mechanisms including shear banding [42–45], subgrain coalescence by the plastic hinge mechanism [34], the collapse of a ribbon-like substructure with local migration and the merger
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Fig. 9. Misorientation distributions for the austenite (sub)boundaries in the dual-phase stainless steel subjected to various strains (e). The bars and the dashed lines represent the TEM and OIM data, respectively.
of adjacent subboundaries [46,47] may delay the reduction of transverse (sub)grain size significantly during cold working. The mechanisms of microstructure evolution during large-strain deformation can be clarified using the ductile dual-phase material with various deformation behaviours for different phases. Fig. 13 shows the strain dependence for the structural parameters of transverse (sub)grain size and transverse phase size. Assuming that the reduction of the (sub)grains and the phases during the cold bar rolling/swaging corresponds to that of the whole sample, the transverse sizes (d) should vary with the strain (e) as ln d 0.5e. In the austenite phase, grain subdivision by multiple twinning at the early stage of deformation results in a rapid decrease of the transverse (sub)grain size to some value, which does not change with further processing. In the ferrite phase, the transverse (sub)grain size basically follows the change in the shape of the sample during processing to strains of about 2, whereas further deformation results in a (sub)grain size larger than that predicted by the change in the shape of the sample. In other words, the grain subdivision is more evident in the phase that is characterised by a higher hardenability. A rapid increase in the dislocation density limits the homogeneous deformation of the austenite and activates some localised deformation modes including twinning and micro-shearing. The ability of ferrite to deform homogeneously is greater. Therefore, the
Fig. 10. Misorientation distributions for the ferrite (sub)boundaries in the dual-phase stainless steel subjected to various strains (e). The bars and the dashed lines represent the TEM and OIM data, respectively.
microstructure evolution in the ferrite to moderate strains is mainly associated with the elongation of the initial (sub)grains along the metal flow direction which is accompanied by increasing (sub)boundary misorientations. Fig. 13 shows that the dependence of the transverse austenite/ferrite phase sizes on the strain and the change in the shape of the samples correspond well during processing to strains of about 4.5, where the phase composition no longer shows a marked change. Since the shear bands definitely develop in the same strain range (see Figs. 2 and 6), a rather high density of micro-shears is suggested; in other words, the presence of a relatively high level of deformation homogeneity on a macroscopic scale. In fact, if the spacing between the intersecting shears is much smaller than the phase size, the micro-shearing will affect the change in the transverse phase size in the same way as conventional dislocation motion, although it can retard the change in the transverse (sub)grain size. The clear correlation of the change in transverse phase sizes in the austenite and the ferrite is indicative of a uniform deformation across the samples. The micro-shears propagate over large distances (at least larger than the transverse dimensions of
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Fig. 11. Effect of strain on the transverse (sub)grain size (d), the number fraction of HAB and the dislocation density (q) in (sub)grain interiors that developed in the cold-worked dual-phase stainless steel.
individual phases); thus, the ferrite and the austenite experience the same strains, regardless of the different levels of strengthening. If the strain distribution is assumed to be uniform among all the (sub)grains across the sample, (sub)boundary migration that leads to grain coarsening and/or a micro-shearing can be used to show that the transverse (sub)grain size is constant during processing. The microshearing seems to be more adequate for the cold working in this study. Therefore, the scheme for the mechanism of the microstructure evolution that operates in ribbon-like substructures at large strains during cold rolling/swaging [27] can be clarified with an emphasis on micro-shearing and is shown in Fig. 14. Local sliding along some transverse (sub)boundaries can be considered as a source of stress concentration, thus creating conditions for strain localisation. The sequential micro-shearing on a microscale level (individual subgrains) that is associated with local (sub)boundary sliding results in a strain localisation on a mesoscale level (several subgrains). Upon unidirectional deformation, the shear planes tend to turn towards the metal flow, which is the deformation axis. A high density of such intersecting micro-shears provides an apparently homogeneous deformation throughout the sample and does not decrease the transverse (sub)grain size. The micro-shears that are involved in sliding along the (sub)boundaries can hardly be recognised on the micrographs, since the shear directions are close to the deformation axis. The arrows in Fig. 8 indicate the possible micro-shears. Various metallic materials subjected to severe deformations were characterised by a rapid recovery releasing internal stresses associated with strain-induced (sub)boundaries [4,48–52]; thus dynamic recovery can play
Fig. 12. Inverse pole figures for the rolling/swaging direction of the dualphase stainless steel. The data were obtained separately for the austenite and ferrite phases using the EBSP method.
Fig. 13. Reduction of the transverse (sub)grain and phase sizes in the dual-phase stainless steel during cold bar rolling/swaging.
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fraction of high-angle (sub)boundaries increased to more than 50%. No more large changes occurred in the transverse (sub)grain size during further processing up to the largest strain studied. Such a structural response was associated with deformation twinning resulting in fast grain subdivision at the early stage of deformation and micro-shearing, which provided uniform macroscopic deformation without any changes in the transverse (sub)grain size at large strains. 4. Finally, severe deformation to a total strain of about 6 resulted in the development of similar microstructures in both the austenite and ferrite phases. The microshearing assisted by the recovery processes was considered as the structural mechanism operating during large-strain cold working.
Acknowledgements
Fig. 14. Schematic drawing of the micro-shears operating in the ribbonlike substructures upon severe unidirectional deformation.
an important role in the deformation behaviour even at relatively low temperatures. In this discussion, the dynamic recovery promoted the micro-shears that propagated partially along the (sub)boundaries and is considered as the mechanism that is responsible for the steady-state-like deformation behaviour at large strains. 5. Conclusions The microstructure evolution in an Fe–27% Cr–9% Ni dual-phase stainless steel during large-strain deformation by bar rolling/swaging to a total strain of 6.9 was studied. The main results are summarised as follows: 1. After a rapid strain hardening at the early stage of deformation, the rate of strain hardening gradually decreased to almost zero at large strains of about 4, leading to a steady-state-like deformation behaviour during large-strain cold working. 2. In the ferrite phase, the transverse (sub)grain size decreased gradually during processing, approaching approximately 0.1 lm at strains above 6.0. The (sub)grain reduction was accompanied by an increase in the (sub)boundary misorientations. Correspondingly, the number fraction of high-angle (sub)boundaries increased to a saturation of about 70%. 3. In the austenite phase, the transverse (sub)grain size decreased rapidly to its minimal value of about 0.1 lm at relatively small strains of about 1, while the number
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