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Recovery and recrystallization in ferritic stainless steel after large strain deformation
Materials Science and Engineering A 403 (2005) 249–259
Recovery and recrystallization in ferritic stainless steel after large strain deformation A. Belyakov ∗,1 , Y. Kimura, K. Tsuzaki Steel Research Center, National Institute for Materials Science, Tsukuba 305-0047, Japan Received in revised form 28 April 2005; accepted 6 May 2005
Abstract Recovery and recrystallization were studied in a Fe–22% Cr–3% Ni ferritic stainless steel processed by bar rolling/swaging at an ambient temperature. The annealing behaviour significantly depended on the preceding cold strain. The samples processed to relatively small strain of 2.0 did not show remarkable softening upon annealing at temperatures below 600 ◦ C (0.5Tm ), while the heating to higher temperatures resulted in the development of conventional primary (discontinuous) recrystallization leading to a rapid annealing softening. On the other hand, the annealing behaviour of the samples processed to strain of 4.4 was characterised by a continuous recrystallization involving a rapid polygonization followed by a normal grain growth that resulted in a gradual softening upon annealing at temperatures above 400 ◦ C (0.38Tm ). The continuously recrystallized microstructures consisted of much finer grains than those resulted from the discontinuous recrystallization. © 2005 Elsevier B.V. All rights reserved. Keywords: Ferritic stainless steel; Severe plastic deformation; Transmission electron microscopy; Ultra fine-grained microstructure; Recovery; Recrystallization
1. Introduction Interest in structural metallic materials processed by large strain deformations has quickened in the past years. Recently, several processing methods including mechanical milling, multiple forging, accumulative roll-bonding, equal channel angular pressing, etc. were successfully developed to achieve severe deformations [1–9]. In certain cases, the large strain deformations resulted in the evolution of submicrocrystalline structures, which could be characterised by very fine grains and high internal stresses. The later ones were frequently discussed to be associated with a non-equilibrium state of straininduced grain boundaries [10–15]. These ultra fine-grained materials are generally believed to have some benefit combination of their properties. An annealing behaviour is one of such properties, which differentiate the materials developed by severe deformations from conventional metals and alloys. ∗
Corresponding author. Tel.:+81 298592185; fax: +81 298592101. E-mail address: [email protected] (A. Belyakov). 1 On leave from the Institute for Metals Superplasticity Problems, Ufa, Russia. 0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.05.057
The fine-grained microstructures developed by large strain deformations were shown to be essentially stable against any discontinuous recrystallization (e.g. primary recrystallization) during subsequent heat treatment [16–22]. Note here, similar peculiarities of annealing behaviour were also reported for materials processed to large strains by conventional methods like rolling and drawing [23–27]. Such annealing behaviour was related to the large fraction of high-angle grain boundaries among the all straininduced (sub)grain boundaries [24,27–30]. Namely, the annealing structural mechanism operating in the deformation microstructures with the fraction of high-angle grain boundaries more than about 60% was considered as a continuous recrystallization. The later could be subdivided into recovery and transient recrystallization followed by normal grain growth [20,31]. Another interesting point to be noted for the severe deformed materials was an ability to fast recovery, which was attributed to the non-equilibrium deformation (sub)grain boundaries [13,32,33]. The rapid recovery of the strain-induced grain boundaries was discussed to operate at an early annealing stage, resulting in the evolution of a number of recrystallization nuclei that finally led to the recrystal-
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lization by continuous mechanism [20]. However, the most of research works was focussed on studying the late annealing stages corresponding to fully recrystallized microstructures with a relatively large grain size. Contrary to discontinuous recrystallization, the nucleation stage of continuous one was not studied in sufficient detail. The role of recovery processes and, especially, polygonization in the formation of continuously recrystallizing microstructures remains unclear. The aim of the present work is to study the annealing softening mechanisms operating in ferritic steel after large strain processing by conventional bar rolling and swaging. In particular, the paper is aimed to clarify the effect of cold deformation on the structural mechanisms responsible for nucleation of recrystallizing grains that takes place at early annealing stages.
2. Experimental procedure A 22Cr–3Ni ferritic stainless steel (0.002% C, 0.01% Mn, 0.005% P, 0.001% S, 21.99% Cr, 3.12% Ni, 0.002% N, all in mass%, and the balance Fe) was vacuum melted and cast into the 20 kg ingot. Then, the steel was hot forged and homogenized at 1200 ◦ C followed by bar rolling at 700 ◦ C. The cold deformation was carried out at ambient temperatures by bar rolling/swaging. The deformation behaviour of the steel during large strain processing was described in detail elsewhere [34]. Effect of the cold working on the hardness and the transverse (sub)grain size are represented in Fig. 1. Two strain levels were selected for comparative study in the present work, namely, strain of 2.0 (cross-section area of about 60 mm2 ) and strain of 4.4 (cross-section area of about 3.5 mm2 ). The deformation microstructures consisted of fine (sub)grains, which were elongated along the deformation axis. The transverse (sub)grain size (d) and the number fraction of high-angle (sub)grain boundaries (HAB) were about 0.27 m and 25% and those of 0.2 m and 75% in the samples processed to strains of 2.0 and 4.4, respectively. The cold-worked steel rods were cut into 10 mm sample pieces, which were annealed in an argon atmosphere at various temperatures from 400 to 700 ◦ C, followed by air-cooling. Structural analysis was performed on sections parallel to the rolling/swaging axis using conventional light microscopy and a JEM-2010F transmission electron microscope. Almost all (sub)grain size data presented in the paper were measured perpendicularly to the rolling/swaging axis by a linear intercept method. The fraction recrystallized was evaluated by point-count technique. Misorientations across the (sub)grain boundaries were analysed by the Kikuchi-line technique [35], collecting 120–140 (sub)boundaries for each specimen. The annealing softening was studied by means of the fractional softening, X = (Hvε − HvT )/(Hvε − Hv0 ), where Hvε and HvT are the hardness for the cold-worked and annealed samples, respectively, and Hv0 is the hardness for a sample annealed at 1200 ◦ C. The hardness tests were carried out with a load of 5N.
Fig. 1. The hardness (Hv), the transverse (sub)grain size (d) and the distributions of (sub)boundary misorientations that developed in the steel by cold rolling/swaging [34]. The arrows point the samples, which were selected for the present study.
3. Results 3.1. Isochronal annealing Effect of the temperature and the cold strain on the softening kinetics during annealing for 30 min is shown in Fig. 2. The hardness of cold-worked samples generally decreases with increasing the annealing temperature; however, the rate of annealing softening in different temperature ranges depends on the intensity of previous cold deformation. This strain dependence of the annealing behaviour is clearly illustrated by the plot of fractional softening versus temperature. The samples cold-rolled to the total strain of 2.0 do not soften remarkably during annealing at temperatures below 500 ◦ C, while the heating to higher temperatures results in rapid softening that leads to approaching the fully annealed state at 700 ◦ C. On the other hand, the samples processed to larger strain of 4.4 are characterised by gradual softening at temperatures above 400 ◦ C. Contrary to the samples processed to strain of 2.0, the samples strained to 4.4 demonstrate faster softening kinetics at relatively low annealing temperatures about 500–600 ◦ C. At higher temperatures, however, the deformation effect on the annealing softening becomes opposite. The softening kinetics of the steels strained to 4.4 slow down when compared to the 2.0-strained samples, leading to an incomplete softening even at relatively high annealing temperature of 700 ◦ C.
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to 700 ◦ C (Fig. 3d and f). However, detailed observations of the fine structures evolved in these samples upon annealing reveal that some grain coarsening certainly takes place at temperatures above 500 ◦ C, leading to the development of almost equiaxed ultra fine-grained microstructures (Fig. 4f). The effect of annealing temperature on the transverse (sub)grain size, i.e. that measured crosswise to the rolling/swaging axis, is quantitatively represented in Fig. 5. The initial as-processed (sub)grain size of 0.27 m in the sample cold-rolled to strain of 2.0 slightly increases to 0.33 m with increase in the annealing temperature to 600 ◦ C. Then, further increase of the annealing temperature results in a drastic grain coarsening. Such annealing behaviour is typical of the primary recrystallization, which develops in conventional cold-worked metals and alloys [36]. For the present case the temperature about 600 ◦ C can be considered as a critical recrystallization temperature. Alternatively, the samples cold swaged to strain of 4.4 demonstrate quite different temperature dependence for the annealed grain size. Namely, the initial as-processed (sub)grain size of 0.2 m does not show any significant jumps with increasing the annealing temperature but gradually grows to 0.73 m at 600 ◦ C and then to 1.5 m at 700 ◦ C. This is unusual recrystallization behaviour and sometimes discussed as a continuous recrystallization [20,24,36] in contradistinction to the primary recrystallization, which is discontinuous in its nature. 3.2. Isothermal annealing
Fig. 2. Effect of cold strain on softening in a ferritic stainless steel during annealing for 30 min.
The variation of softening behaviour in the samples deformed to different strains suggests the difference in the mechanisms of microstructure evolution that operate upon annealing. Typical annealed microstructures are shown in Figs. 3 and 4. These figures also represent the as-processed deformation microstructures for comparison. (Note here, the rolling/swaging axes are horizontal in the all micrographs presented in the paper.) In the as-processed state, the deformation microstructures in the samples strained to 2.0 consisted of cell blocks separated by strain-induced geometrically necessary (sub)boundaries such as dense dislocation walls, microbands, etc. (Fig. 4a) and the samples strained to 4.4 were characterised by ribbon-like substructures that arranged parallel to the swaging axes (Fig. 4d). The samples processed to strain of 2.0 do not demonstrate any remarkable change in their microstructures during 30 min annealing at temperatures of T ≤ 600 ◦ C, while the heating to 700 ◦ C changes the microstructure completely (Fig. 3c). The later can be considered as a perfectly recrystallized microstructure with an average grain size about 10 m. On the other hand, the samples strained to 4.4 apparently keep their submicrocrystalline deformation structures under annealing at temperatures up
Let us consider the sequences of structural changes taking place during annealing of the steels cold-worked to different levels of strain, 2.0 and 4.4, in more detail. Since recovery and recrystallization in the both sets of samples readily develop at T > 600 ◦ C, the temperature of 625 ◦ C is selected here for comparative study on annealing mechanisms of microstructure evolution. 3.2.1. Annealed microstructures in samples strained to 2.0 Fig. 6 illustrates the structural changes taking place in the steel samples cold-rolled to total strain of 2.0 during annealing at 625 ◦ C. A short time annealing does not lead to any significant structural changes that associate with recrystallization. The microstructure annealed for 10 min is almost identical with the cold-worked one (cf. Figs. 3a and 6a). After 20 min annealing, the microstructure is partially recrystallized (Fig. 6b). The recrystallizing grains rapidly grow, consuming cold-worked surroundings. Further annealing leads to the completion of recrystallization and then, a normal grain growth becomes a main mechanism of microstructure evolution upon a long time annealing. The recrystallization kinetics is presented in Fig. 7. The time dependence of the fraction recrystallized displays the characteristic sigmoid shape, where recrystallization nucleation, progress and completion can be clearly defined. The slope of Avrami plot is about 1.7 in Fig. 7. Although this value is smaller than originally theoreti-
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Fig. 3. Annealed microstructures in a ferritic stainless steel cold-worked to strains of 2.0 (a–c) and 4.4 (d–f): (a) and (d) represent as-processed states; (b) and (e) annealed for 30 min at 600 ◦ C; (c) and (f) annealed for 30 min at 700 ◦ C.
cally predicted one, the similar results were reported in many other experimental studies [36]. Variations in the recrystallization kinetics may be associated with specific nucleation sites that lead to a non-random nucleation. Substructures corresponding to the nucleation stage of recrystallization are shown in Fig. 8. Several (sub)grains without dislocations in their interiors that arrange a chain
along the rolling axes are clearly distinguished from the coldworked matrix substructures containing high dislocation densities. Such dislocation free (sub)grains can be considered as nuclei for recrystallizing grains. They are located between the closely spaced high-angle (sub)boundaries, which separate them from cold-worked surroundings. It should be noted that the transverse (sub)boundaries in these dislocation free
Fig. 4. Annealed substructures in a ferritic stainless steel cold-worked to strains of 2.0 (a–c) and 4.4 (d–f): (a) and (d) correspond to as-processed states; (b) and (e) annealed for 30 min at 500 ◦ C; (c) and (f) annealed for 30 min at 600 ◦ C.
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Fig. 5. Temperature effect on the transverse (sub)grain size in a ferritic stainless steel after annealing for 30 min.
chains are frequently low-angle (sub)boundaries. This confirms that the recrystallization nucleation frequently occurs at the strain-induced geometrically necessary (sub)boundaries, when the later ones are close together [37]. 3.2.2. Annealed microstructures in samples strained to 4.4 Typical annealed microstructures that evolve in the samples after cold swaging to total strain of 4.4 are shown in Fig. 9. It is clearly seen that the deformation ribbon-like substructures quickly change just after the heating. The transverse (sub)grain size increases accompanied with the formation of many (sub)boundaries crosswise to the swaging axes, leading to the evolution of almost equiaxed ultra fine-grained microstructure even after short time annealing for 5 min (Fig. 9b). In contrast to the primary recrystallization, however, further annealing does not result in the development of relatively coarse-grained microstructure (cf. Figs. 6c and 9c). That is to say, the grain coarsening rate apparently decreases during annealing. The time dependence for the (sub)grain size is presented in Fig. 10. After fast (sub)grain growth at an early annealing stage, the grain coarsening rate decreases to some value and then does not change under long time annealing. Therefore, the annealing behaviour looks similar to a normal grain coarsening with a grain growth exponent about 6, which is consistent with other research works on normal grain growth behaviour [36,38–40]. Since the recrystallizing grains homogeneously evolve in place of deformation (sub)grains, such process is generally discussed as a continuous recrystallization. In the present study, the nucleation stage for continuous recrystallization seems to be very short. A detailed micrograph of fine structure evolved at an early annealing is shown in Fig. 11. This micrograph represents the upper portion
Fig. 6. Development of recrystallized grains in a ferritic stainless steel coldworked to strain of 2.0 and then annealed at 625 ◦ C for various times: (a) 10 min; (b) 20 min; (c) 120 min annealing.
of Fig. 9a. Compared to the as-processed state, the shortly annealed microstructure is characterised by the evolution of rather sharp transverse dislocation (sub)boundaries (Fig. 9a), which modify ribbon-like deformation (sub)grains with rows of somewhat elongated (sub)grains (Fig. 4d). The most interesting feature of this fine structure, however, is the appearance of bulging (sub)boundaries leading to the evolution of fine (sub)grains along the longitudinal (sub)boundaries
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Fig. 7. Primary recrystallization behaviour in a ferritic stainless steel strained to 2.0 and then annealed at 625 ◦ C.
of ribbon-like substructures. Some of such fine (sub)grains are pointed by asterisks in Fig. 11. It is evident that the bulging (sub)boundaries can merge with opposite longitudinal (sub)boundaries, resulting in the formation of spatial network of (sub)boundaries with low- to high-angle misorientations. This process operating simultaneously with a local migration and merging of the strain-induced longitudinal (sub)boundaries seems to be responsible for the rapid development of ultra fine-grained microstructure upon annealing. The annealed microstructures are characterised by the specific distributions of (sub)grain boundary misorientations (θ) as shown in Fig. 12. The misorientation distributions for the cold-worked sample and that annealed at 500 ◦ C for 30 min are also shown in this figure for comparison. (The misorientation distribution for the cold-worked sample was re-evaluated here with improved statistics compared to that in [34].) The fractions of low-angle (sub)boundaries (θ < 15◦ ) and grain boundaries with misorientations close to 60◦ (θ > 50◦ ) increase rapidly at an early annealing stage. It should be noted that the annealing at lower temperature of 500 ◦ C gives the same result for the misorientation distribution, while the transverse (sub)grain size changes insignif-
Fig. 8. Recrystallization nucleation in a ferritic stainless steel cold-worked to strain of 2.0 and annealed at 625 ◦ C for 10 min: (b) represents the enlarged central portion of (a); the numbers in (b) indicate the (sub)boundary misorientations in degrees.
icantly. Therefore, the flat-type misorientation distribution with almost equal fractions of (sub)boundaries with different misorientations that evolved by cold working is modified by the bimodal one with two widely separated peaks similar to those observed during heavy warm deformation [41]. Further annealing is characterised by the normal grain growth following the progress in continuous recrystallization. It is clearly seen in Fig. 12 that the normal grain growth operating at a late stage of continuous recrystallization does not change the character of grain boundary misorientation distribution. 3.2.3. Softening behaviour The different levels of cold deformation cause the variation in the softening behaviour (Fig. 13). The hardness of the samples processed to strain of 2.0 gradually decreases during annealing. Note here that in case of partially recrystallized microstructures, the hardness of recrystallized volumes was measured separately from unrecrystallized (recovered) ones. In spite of some difference between recrystallized and recovered volumes in their hardness, the hardness roughly demonstrates a power low function of annealing time. On the
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Fig. 10. Grain coarsening upon annealing at 625 ◦ C of a ferritic stainless steel strained to 4.4.
The effect of the different annealing mechanisms, i.e. continuous and discontinuous recrystallizations operating in the different samples, on the softening kinetics can be more clearly illustrated with the fractional softening (X). The hardness for the partially recrystallized samples (those processed to strain of 2.0) was averaged through the hardness of recrystallized (HvReX ) and recovered (HvRCV ) volumes as follows: Hv = HvReX FReX + HvRCV (1 − FReX ), where FReX is the fraction recrystallized. The primary recrystallization development in the samples strained to 2.0 leads the fractional softening to unique time dependence that similar to the hardness change. It is clearly seen in Fig. 13 that the continuous recrystallization in the samples strained to 4.4 is characterised by a remarkable change in the softening kinetics. The high softening rate at an early annealing corresponds to the drastic structural changes leading to the formation of very fine-
Fig. 9. Continuous recrystallization in a ferritic stainless steel cold-worked to strain of 4.4 and then annealed at 625 ◦ C: (a) 2 min; (b) 5 min; (c) 120 min annealing.
other hand, the samples strained to 4.4 are characterised by a quite rapid softening just after heating. These samples have a rather high level of the initial strain hardening; however, they quickly lose their strength at an early annealing and become even softer that those processed to smaller strain. Then, the hardness of the large strained samples shows very weak time dependence at later annealing stage.
Fig. 11. Polygonization development during 2 min annealing at 625 ◦ C of a ferritic stainless steel cold-worked to strain of 4.4. The picture represents the enlarged upper portion of Fig. 9a.The numbers indicate the (sub)boundary misorientations in degrees.
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Fig. 13. Effect of cold strain on softening in a ferritic stainless steel during annealing at 625 ◦ C.
Fig. 12. Misorientation distributions of (sub)boundaries in a ferritic stainless steel cold-worked to strain of 4.4 and annealed at 500 and 625 ◦ C.
equiaxed grains instead of highly elongated strain-induced (sub)grains. Such fine-grained microstructure is rather stable upon further annealing; the structural changes are associated with a normal grain growth. Compared to the primary recrystallization, the normal grain growth is a very slow process wherein the softening depends on the grain coarsening.
4. Discussion 4.1. Recovery Recovery is a common softening process operating in cold-worked materials just after heating. The recovery development results in the annihilation and rearrangement of deformation defects, mainly dislocations, leading to the formation of low energy substructures, polygonization, followed by the (sub)grain coalescence. Contrary to recrystallization, recovery does not lead to the complete softening. When the primary recrystallization was suppressed, the fractional softening by
the recovery was shown to depend on the previous cold deformation and, in certain cases, the recovery required a critical pre-strain to develop [42,43]. When the primary recrystallization developed, the recovery was considered as a structural mechanism responsible for the development of recrystallizing nuclei at an early annealing stage [36]. In this case, the fractional softening by the recovery was negligible, and the strain hardening (i.e. stored energy) could be completely removed by the following recrystallization. In the present study, recovery preceding recrystallization in the samples processed to strain of 2.0 results in the evolution of dislocation free fine (sub)grains with low- to high-angle boundaries (Fig. 8). These (sub)grains are surrounded by dislocation substructures, which look like typical cold-worked ones. Such local formation of perfectly recovered (sub)grains suggests some heterogeneity in the recovery development. The fast recovery occurs in the local volumes with a high density of the deformation (sub)boundaries with high-angle misorientations. The high-angle (sub)boundaries may act as preferable absorption sites for the lattice dislocations. In fact, the increase in the density of grain boundary dislocations does not change the surface energy of high-angle grain boundaries significantly as compared to that in low-
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angle (sub)boundaries [36,44]. Therefore, the inhomogeneity of deformation substructures results in the heterogeneous recovery. Since the volume fraction of perfectly recovered (sub)grains in the samples strained to 2.0 is small at early annealing, these samples are characterised by relatively small fractional softening (Figs. 2 and 13). On the other hand, the samples processed to strain of 4.4 are characterised by the large fraction of high-angle (sub)grain boundaries of about 75%. In other words, almost all (sub)boundaries are high-angle grain boundaries. Such homogeneous deformation substructures that consist of highangle (sub)boundaries are responsible for the rapid recovery development throughout the sample. The high internal stresses in the severely deformed materials are associated with strain-induced (sub)grain boundaries, in turn, the latter ones are easy to be recovered [13,32]. The large strained samples, therefore, shows remarkable fractional softening at the recovery stage during annealing, i.e. just after heating to high temperatures as well as during annealing at relatively low temperatures. 4.2. Recrystallization Recrystallization in the samples strained to 2.0 develops in conventional manner of the primary recrystallization, which is common annealing mechanism operating frequently in cold-worked metals and alloys. Let us discuss, therefore, recrystallization in the samples processed to rather large strain of 4.4, the annealing behaviour of which suggests the continuous recrystallization mechanism. The continuous recrystallization was discussed as an annealing mechanism operating in some alloys containing a relatively large fraction of hard second-phase particles [24,29,45]. The particle growth upon annealing was accompanied by the coarsening of deformation substructures, leading to the equiaxed grained microstructures like recrystallized ones. The present results suggest that the structural changes taking place at the beginning of the continuous recrystallization are qualitatively similar to those occurring at the nucleation stage of the primary recrystallization. Namely, the fast recovery in the ribbon-like deformation substructures at an early annealing results in the dislocation rearrangement within the strain-induced continuously curved crystallites, leading to the break up of the highly elongated (sub)grains into the chains of almost equiaxed (sub)grains (cf. Figs. 4d and 9a). However, compared to the primary recrystallization, where the polygonization leading to recrystallizing nuclei develops heterogeneously, the continuous recrystallization is characterised by the homogeneous nucleation. Also, the rapid increase of the transverse (sub)grain size (Fig. 10) suggests that the recovery of the ribbon-like deformation substructures is accompanied by a merger of the longitudinal (sub)boundaries that results from their local migration. The polygonization and the local (sub)boundary migration during annealing can be accelerated by high internal stresses
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that develop in severe deformed materials. The stress sources are frequently associated with specific lattice defects, disclinations, that evolve at the strain-induced (sub)grain boundaries and their junctions [13,34,46–48]. Release of such internal stresses may be realised by the motion and then annihilation of disclinations. The disclination motion is necessary provided by the generation and/or rearrangement of dislocations. It was theoretically predicted that the disclination motion along the grain boundary associated with the emission of lattice dislocations from the boundary may be energetically favourable [49]. The similar consideration can be applied to discuss the structural changes upon early annealing of the samples strained to 4.4. The bulging (sub)boundaries (Fig. 11) may be composed from the emitted grain boundary dislocations. This process looks like a splitting of the deformation (sub)grain boundaries, leading to the evolution of many low-angle (sub)boundaries and may be considered as a special structural mechanism assisting polygonization. Further annealing results in a coarsening of the polygonized structure much similar to the normal grain growth. It may be concluded, finally, that the continuous recrystallization in the present study involves two simple sequential steps, i.e. polygonization followed by normal grain growth. In the present discussion, the term polygonization is used to emphasize the phenomenon, namely, the quick in situ transformation of ribbon-like deformation substructures into equiaxed (sub)grains. This transformation is accompanied with a sharp increase of the (sub)grain sizes in the transverse direction (Fig. 10) as well as a rapid rise of the number fraction of low-angle (sub)boundaries (Fig. 12). Here, the polygonization stage is essentially identical to the transient recrystallization discussed in the previous study on continuous recrystallization after large strain deformation [20,31]. A relatively fast development kinetics in the present study can be caused by a larger difference between the processing and annealing temperatures. Considering the conditions for continuous recrystallization, the homogeneity of deformation substructures seems to be a key parameter controlling the recrystallization mechanism. In the present case, the uniaxial large strain processing provided the highly uniform deformation substructures, i.e. (sub)grain sizes, dislocation densities, (sub)boundary misorientations, etc. Therefore, the probability for appearance of recrystallizing nuclei is the same for almost all (sub)grains throughout the deformation structures. Then, the simultaneous development of recrystallizing nuclei in place of deformation substructures inhibits any selective growth of separate nuclei, resulting in the homogeneous, continuous recrystallization.
5. Conclusions The annealing behaviour of a cold bar-rolled/swaged Fe–22% Cr–3% Ni ferritic stainless steel was studied at temperatures of 400–700 ◦ C. The main results could be sum-
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marised as follows: 1. The mechanisms of microstructure evolution that operated in cold-worked substructures during annealing depended significantly on the intensity of previous deformation; 2. The samples cold-worked to strain of 2.0 demonstrated conventional primary (discontinuous) recrystallization behaviour during annealing with a critical recrystallization temperature of about 600 ◦ C. Recovery at temperatures below 600 ◦ C did not lead to remarkable softening, while heating to higher temperatures resulted rapidly in complete softening according to the recrystallization development; 3. The annealing behaviour of the samples cold-worked to strain of 4.4 was considered as a continuous recrystallization, which involved recovery followed by normal grain growth. The recovery developed quickly just after the heating and resulted in the evolution of polygonized microstructures in place of ribbon-like deformation substructures. Such polygonized microstructures consisting of low- to high-angle (sub)boundaries were essentially stable against discontinuous coarsening and, therefore, only the normal grain growth operated upon further annealing. The samples were characterised by a gradual softening with increasing the temperature in the all temperature range studied, but did not approach perfect softening even at a rather high temperature of 700 ◦ C; 4. The continuously recrystallized microstructures consisted of much finer grains than those resulted from the discontinuous recrystallization.
Acknowledgements The authors are grateful to Drs. N. Sakuma, T. Hibaru, S. Kuroda, M. Kobayashi, T. Kanno, Steel Research Center, National Institute for Materials Science, for their assistance in the materials processing. One of the authors, A.B., would like to express his thanks to the National Institute for Materials Science for providing a scientific fellowship.
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Recovery and recrystallization in ferritic stainless steel after large strain deformation A. Belyakov ∗,1 , Y. Kimura, K. Tsuzaki Steel Research Center, National Institute for Materials Science, Tsukuba 305-0047, Japan Received in revised form 28 April 2005; accepted 6 May 2005
Abstract Recovery and recrystallization were studied in a Fe–22% Cr–3% Ni ferritic stainless steel processed by bar rolling/swaging at an ambient temperature. The annealing behaviour significantly depended on the preceding cold strain. The samples processed to relatively small strain of 2.0 did not show remarkable softening upon annealing at temperatures below 600 ◦ C (0.5Tm ), while the heating to higher temperatures resulted in the development of conventional primary (discontinuous) recrystallization leading to a rapid annealing softening. On the other hand, the annealing behaviour of the samples processed to strain of 4.4 was characterised by a continuous recrystallization involving a rapid polygonization followed by a normal grain growth that resulted in a gradual softening upon annealing at temperatures above 400 ◦ C (0.38Tm ). The continuously recrystallized microstructures consisted of much finer grains than those resulted from the discontinuous recrystallization. © 2005 Elsevier B.V. All rights reserved. Keywords: Ferritic stainless steel; Severe plastic deformation; Transmission electron microscopy; Ultra fine-grained microstructure; Recovery; Recrystallization
1. Introduction Interest in structural metallic materials processed by large strain deformations has quickened in the past years. Recently, several processing methods including mechanical milling, multiple forging, accumulative roll-bonding, equal channel angular pressing, etc. were successfully developed to achieve severe deformations [1–9]. In certain cases, the large strain deformations resulted in the evolution of submicrocrystalline structures, which could be characterised by very fine grains and high internal stresses. The later ones were frequently discussed to be associated with a non-equilibrium state of straininduced grain boundaries [10–15]. These ultra fine-grained materials are generally believed to have some benefit combination of their properties. An annealing behaviour is one of such properties, which differentiate the materials developed by severe deformations from conventional metals and alloys. ∗
Corresponding author. Tel.:+81 298592185; fax: +81 298592101. E-mail address: [email protected] (A. Belyakov). 1 On leave from the Institute for Metals Superplasticity Problems, Ufa, Russia. 0921-5093/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2005.05.057
The fine-grained microstructures developed by large strain deformations were shown to be essentially stable against any discontinuous recrystallization (e.g. primary recrystallization) during subsequent heat treatment [16–22]. Note here, similar peculiarities of annealing behaviour were also reported for materials processed to large strains by conventional methods like rolling and drawing [23–27]. Such annealing behaviour was related to the large fraction of high-angle grain boundaries among the all straininduced (sub)grain boundaries [24,27–30]. Namely, the annealing structural mechanism operating in the deformation microstructures with the fraction of high-angle grain boundaries more than about 60% was considered as a continuous recrystallization. The later could be subdivided into recovery and transient recrystallization followed by normal grain growth [20,31]. Another interesting point to be noted for the severe deformed materials was an ability to fast recovery, which was attributed to the non-equilibrium deformation (sub)grain boundaries [13,32,33]. The rapid recovery of the strain-induced grain boundaries was discussed to operate at an early annealing stage, resulting in the evolution of a number of recrystallization nuclei that finally led to the recrystal-
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lization by continuous mechanism [20]. However, the most of research works was focussed on studying the late annealing stages corresponding to fully recrystallized microstructures with a relatively large grain size. Contrary to discontinuous recrystallization, the nucleation stage of continuous one was not studied in sufficient detail. The role of recovery processes and, especially, polygonization in the formation of continuously recrystallizing microstructures remains unclear. The aim of the present work is to study the annealing softening mechanisms operating in ferritic steel after large strain processing by conventional bar rolling and swaging. In particular, the paper is aimed to clarify the effect of cold deformation on the structural mechanisms responsible for nucleation of recrystallizing grains that takes place at early annealing stages.
2. Experimental procedure A 22Cr–3Ni ferritic stainless steel (0.002% C, 0.01% Mn, 0.005% P, 0.001% S, 21.99% Cr, 3.12% Ni, 0.002% N, all in mass%, and the balance Fe) was vacuum melted and cast into the 20 kg ingot. Then, the steel was hot forged and homogenized at 1200 ◦ C followed by bar rolling at 700 ◦ C. The cold deformation was carried out at ambient temperatures by bar rolling/swaging. The deformation behaviour of the steel during large strain processing was described in detail elsewhere [34]. Effect of the cold working on the hardness and the transverse (sub)grain size are represented in Fig. 1. Two strain levels were selected for comparative study in the present work, namely, strain of 2.0 (cross-section area of about 60 mm2 ) and strain of 4.4 (cross-section area of about 3.5 mm2 ). The deformation microstructures consisted of fine (sub)grains, which were elongated along the deformation axis. The transverse (sub)grain size (d) and the number fraction of high-angle (sub)grain boundaries (HAB) were about 0.27 m and 25% and those of 0.2 m and 75% in the samples processed to strains of 2.0 and 4.4, respectively. The cold-worked steel rods were cut into 10 mm sample pieces, which were annealed in an argon atmosphere at various temperatures from 400 to 700 ◦ C, followed by air-cooling. Structural analysis was performed on sections parallel to the rolling/swaging axis using conventional light microscopy and a JEM-2010F transmission electron microscope. Almost all (sub)grain size data presented in the paper were measured perpendicularly to the rolling/swaging axis by a linear intercept method. The fraction recrystallized was evaluated by point-count technique. Misorientations across the (sub)grain boundaries were analysed by the Kikuchi-line technique [35], collecting 120–140 (sub)boundaries for each specimen. The annealing softening was studied by means of the fractional softening, X = (Hvε − HvT )/(Hvε − Hv0 ), where Hvε and HvT are the hardness for the cold-worked and annealed samples, respectively, and Hv0 is the hardness for a sample annealed at 1200 ◦ C. The hardness tests were carried out with a load of 5N.
Fig. 1. The hardness (Hv), the transverse (sub)grain size (d) and the distributions of (sub)boundary misorientations that developed in the steel by cold rolling/swaging [34]. The arrows point the samples, which were selected for the present study.
3. Results 3.1. Isochronal annealing Effect of the temperature and the cold strain on the softening kinetics during annealing for 30 min is shown in Fig. 2. The hardness of cold-worked samples generally decreases with increasing the annealing temperature; however, the rate of annealing softening in different temperature ranges depends on the intensity of previous cold deformation. This strain dependence of the annealing behaviour is clearly illustrated by the plot of fractional softening versus temperature. The samples cold-rolled to the total strain of 2.0 do not soften remarkably during annealing at temperatures below 500 ◦ C, while the heating to higher temperatures results in rapid softening that leads to approaching the fully annealed state at 700 ◦ C. On the other hand, the samples processed to larger strain of 4.4 are characterised by gradual softening at temperatures above 400 ◦ C. Contrary to the samples processed to strain of 2.0, the samples strained to 4.4 demonstrate faster softening kinetics at relatively low annealing temperatures about 500–600 ◦ C. At higher temperatures, however, the deformation effect on the annealing softening becomes opposite. The softening kinetics of the steels strained to 4.4 slow down when compared to the 2.0-strained samples, leading to an incomplete softening even at relatively high annealing temperature of 700 ◦ C.
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to 700 ◦ C (Fig. 3d and f). However, detailed observations of the fine structures evolved in these samples upon annealing reveal that some grain coarsening certainly takes place at temperatures above 500 ◦ C, leading to the development of almost equiaxed ultra fine-grained microstructures (Fig. 4f). The effect of annealing temperature on the transverse (sub)grain size, i.e. that measured crosswise to the rolling/swaging axis, is quantitatively represented in Fig. 5. The initial as-processed (sub)grain size of 0.27 m in the sample cold-rolled to strain of 2.0 slightly increases to 0.33 m with increase in the annealing temperature to 600 ◦ C. Then, further increase of the annealing temperature results in a drastic grain coarsening. Such annealing behaviour is typical of the primary recrystallization, which develops in conventional cold-worked metals and alloys [36]. For the present case the temperature about 600 ◦ C can be considered as a critical recrystallization temperature. Alternatively, the samples cold swaged to strain of 4.4 demonstrate quite different temperature dependence for the annealed grain size. Namely, the initial as-processed (sub)grain size of 0.2 m does not show any significant jumps with increasing the annealing temperature but gradually grows to 0.73 m at 600 ◦ C and then to 1.5 m at 700 ◦ C. This is unusual recrystallization behaviour and sometimes discussed as a continuous recrystallization [20,24,36] in contradistinction to the primary recrystallization, which is discontinuous in its nature. 3.2. Isothermal annealing
Fig. 2. Effect of cold strain on softening in a ferritic stainless steel during annealing for 30 min.
The variation of softening behaviour in the samples deformed to different strains suggests the difference in the mechanisms of microstructure evolution that operate upon annealing. Typical annealed microstructures are shown in Figs. 3 and 4. These figures also represent the as-processed deformation microstructures for comparison. (Note here, the rolling/swaging axes are horizontal in the all micrographs presented in the paper.) In the as-processed state, the deformation microstructures in the samples strained to 2.0 consisted of cell blocks separated by strain-induced geometrically necessary (sub)boundaries such as dense dislocation walls, microbands, etc. (Fig. 4a) and the samples strained to 4.4 were characterised by ribbon-like substructures that arranged parallel to the swaging axes (Fig. 4d). The samples processed to strain of 2.0 do not demonstrate any remarkable change in their microstructures during 30 min annealing at temperatures of T ≤ 600 ◦ C, while the heating to 700 ◦ C changes the microstructure completely (Fig. 3c). The later can be considered as a perfectly recrystallized microstructure with an average grain size about 10 m. On the other hand, the samples strained to 4.4 apparently keep their submicrocrystalline deformation structures under annealing at temperatures up
Let us consider the sequences of structural changes taking place during annealing of the steels cold-worked to different levels of strain, 2.0 and 4.4, in more detail. Since recovery and recrystallization in the both sets of samples readily develop at T > 600 ◦ C, the temperature of 625 ◦ C is selected here for comparative study on annealing mechanisms of microstructure evolution. 3.2.1. Annealed microstructures in samples strained to 2.0 Fig. 6 illustrates the structural changes taking place in the steel samples cold-rolled to total strain of 2.0 during annealing at 625 ◦ C. A short time annealing does not lead to any significant structural changes that associate with recrystallization. The microstructure annealed for 10 min is almost identical with the cold-worked one (cf. Figs. 3a and 6a). After 20 min annealing, the microstructure is partially recrystallized (Fig. 6b). The recrystallizing grains rapidly grow, consuming cold-worked surroundings. Further annealing leads to the completion of recrystallization and then, a normal grain growth becomes a main mechanism of microstructure evolution upon a long time annealing. The recrystallization kinetics is presented in Fig. 7. The time dependence of the fraction recrystallized displays the characteristic sigmoid shape, where recrystallization nucleation, progress and completion can be clearly defined. The slope of Avrami plot is about 1.7 in Fig. 7. Although this value is smaller than originally theoreti-
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Fig. 3. Annealed microstructures in a ferritic stainless steel cold-worked to strains of 2.0 (a–c) and 4.4 (d–f): (a) and (d) represent as-processed states; (b) and (e) annealed for 30 min at 600 ◦ C; (c) and (f) annealed for 30 min at 700 ◦ C.
cally predicted one, the similar results were reported in many other experimental studies [36]. Variations in the recrystallization kinetics may be associated with specific nucleation sites that lead to a non-random nucleation. Substructures corresponding to the nucleation stage of recrystallization are shown in Fig. 8. Several (sub)grains without dislocations in their interiors that arrange a chain
along the rolling axes are clearly distinguished from the coldworked matrix substructures containing high dislocation densities. Such dislocation free (sub)grains can be considered as nuclei for recrystallizing grains. They are located between the closely spaced high-angle (sub)boundaries, which separate them from cold-worked surroundings. It should be noted that the transverse (sub)boundaries in these dislocation free
Fig. 4. Annealed substructures in a ferritic stainless steel cold-worked to strains of 2.0 (a–c) and 4.4 (d–f): (a) and (d) correspond to as-processed states; (b) and (e) annealed for 30 min at 500 ◦ C; (c) and (f) annealed for 30 min at 600 ◦ C.
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Fig. 5. Temperature effect on the transverse (sub)grain size in a ferritic stainless steel after annealing for 30 min.
chains are frequently low-angle (sub)boundaries. This confirms that the recrystallization nucleation frequently occurs at the strain-induced geometrically necessary (sub)boundaries, when the later ones are close together [37]. 3.2.2. Annealed microstructures in samples strained to 4.4 Typical annealed microstructures that evolve in the samples after cold swaging to total strain of 4.4 are shown in Fig. 9. It is clearly seen that the deformation ribbon-like substructures quickly change just after the heating. The transverse (sub)grain size increases accompanied with the formation of many (sub)boundaries crosswise to the swaging axes, leading to the evolution of almost equiaxed ultra fine-grained microstructure even after short time annealing for 5 min (Fig. 9b). In contrast to the primary recrystallization, however, further annealing does not result in the development of relatively coarse-grained microstructure (cf. Figs. 6c and 9c). That is to say, the grain coarsening rate apparently decreases during annealing. The time dependence for the (sub)grain size is presented in Fig. 10. After fast (sub)grain growth at an early annealing stage, the grain coarsening rate decreases to some value and then does not change under long time annealing. Therefore, the annealing behaviour looks similar to a normal grain coarsening with a grain growth exponent about 6, which is consistent with other research works on normal grain growth behaviour [36,38–40]. Since the recrystallizing grains homogeneously evolve in place of deformation (sub)grains, such process is generally discussed as a continuous recrystallization. In the present study, the nucleation stage for continuous recrystallization seems to be very short. A detailed micrograph of fine structure evolved at an early annealing is shown in Fig. 11. This micrograph represents the upper portion
Fig. 6. Development of recrystallized grains in a ferritic stainless steel coldworked to strain of 2.0 and then annealed at 625 ◦ C for various times: (a) 10 min; (b) 20 min; (c) 120 min annealing.
of Fig. 9a. Compared to the as-processed state, the shortly annealed microstructure is characterised by the evolution of rather sharp transverse dislocation (sub)boundaries (Fig. 9a), which modify ribbon-like deformation (sub)grains with rows of somewhat elongated (sub)grains (Fig. 4d). The most interesting feature of this fine structure, however, is the appearance of bulging (sub)boundaries leading to the evolution of fine (sub)grains along the longitudinal (sub)boundaries
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Fig. 7. Primary recrystallization behaviour in a ferritic stainless steel strained to 2.0 and then annealed at 625 ◦ C.
of ribbon-like substructures. Some of such fine (sub)grains are pointed by asterisks in Fig. 11. It is evident that the bulging (sub)boundaries can merge with opposite longitudinal (sub)boundaries, resulting in the formation of spatial network of (sub)boundaries with low- to high-angle misorientations. This process operating simultaneously with a local migration and merging of the strain-induced longitudinal (sub)boundaries seems to be responsible for the rapid development of ultra fine-grained microstructure upon annealing. The annealed microstructures are characterised by the specific distributions of (sub)grain boundary misorientations (θ) as shown in Fig. 12. The misorientation distributions for the cold-worked sample and that annealed at 500 ◦ C for 30 min are also shown in this figure for comparison. (The misorientation distribution for the cold-worked sample was re-evaluated here with improved statistics compared to that in [34].) The fractions of low-angle (sub)boundaries (θ < 15◦ ) and grain boundaries with misorientations close to 60◦ (θ > 50◦ ) increase rapidly at an early annealing stage. It should be noted that the annealing at lower temperature of 500 ◦ C gives the same result for the misorientation distribution, while the transverse (sub)grain size changes insignif-
Fig. 8. Recrystallization nucleation in a ferritic stainless steel cold-worked to strain of 2.0 and annealed at 625 ◦ C for 10 min: (b) represents the enlarged central portion of (a); the numbers in (b) indicate the (sub)boundary misorientations in degrees.
icantly. Therefore, the flat-type misorientation distribution with almost equal fractions of (sub)boundaries with different misorientations that evolved by cold working is modified by the bimodal one with two widely separated peaks similar to those observed during heavy warm deformation [41]. Further annealing is characterised by the normal grain growth following the progress in continuous recrystallization. It is clearly seen in Fig. 12 that the normal grain growth operating at a late stage of continuous recrystallization does not change the character of grain boundary misorientation distribution. 3.2.3. Softening behaviour The different levels of cold deformation cause the variation in the softening behaviour (Fig. 13). The hardness of the samples processed to strain of 2.0 gradually decreases during annealing. Note here that in case of partially recrystallized microstructures, the hardness of recrystallized volumes was measured separately from unrecrystallized (recovered) ones. In spite of some difference between recrystallized and recovered volumes in their hardness, the hardness roughly demonstrates a power low function of annealing time. On the
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Fig. 10. Grain coarsening upon annealing at 625 ◦ C of a ferritic stainless steel strained to 4.4.
The effect of the different annealing mechanisms, i.e. continuous and discontinuous recrystallizations operating in the different samples, on the softening kinetics can be more clearly illustrated with the fractional softening (X). The hardness for the partially recrystallized samples (those processed to strain of 2.0) was averaged through the hardness of recrystallized (HvReX ) and recovered (HvRCV ) volumes as follows: Hv = HvReX FReX + HvRCV (1 − FReX ), where FReX is the fraction recrystallized. The primary recrystallization development in the samples strained to 2.0 leads the fractional softening to unique time dependence that similar to the hardness change. It is clearly seen in Fig. 13 that the continuous recrystallization in the samples strained to 4.4 is characterised by a remarkable change in the softening kinetics. The high softening rate at an early annealing corresponds to the drastic structural changes leading to the formation of very fine-
Fig. 9. Continuous recrystallization in a ferritic stainless steel cold-worked to strain of 4.4 and then annealed at 625 ◦ C: (a) 2 min; (b) 5 min; (c) 120 min annealing.
other hand, the samples strained to 4.4 are characterised by a quite rapid softening just after heating. These samples have a rather high level of the initial strain hardening; however, they quickly lose their strength at an early annealing and become even softer that those processed to smaller strain. Then, the hardness of the large strained samples shows very weak time dependence at later annealing stage.
Fig. 11. Polygonization development during 2 min annealing at 625 ◦ C of a ferritic stainless steel cold-worked to strain of 4.4. The picture represents the enlarged upper portion of Fig. 9a.The numbers indicate the (sub)boundary misorientations in degrees.
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Fig. 13. Effect of cold strain on softening in a ferritic stainless steel during annealing at 625 ◦ C.
Fig. 12. Misorientation distributions of (sub)boundaries in a ferritic stainless steel cold-worked to strain of 4.4 and annealed at 500 and 625 ◦ C.
equiaxed grains instead of highly elongated strain-induced (sub)grains. Such fine-grained microstructure is rather stable upon further annealing; the structural changes are associated with a normal grain growth. Compared to the primary recrystallization, the normal grain growth is a very slow process wherein the softening depends on the grain coarsening.
4. Discussion 4.1. Recovery Recovery is a common softening process operating in cold-worked materials just after heating. The recovery development results in the annihilation and rearrangement of deformation defects, mainly dislocations, leading to the formation of low energy substructures, polygonization, followed by the (sub)grain coalescence. Contrary to recrystallization, recovery does not lead to the complete softening. When the primary recrystallization was suppressed, the fractional softening by
the recovery was shown to depend on the previous cold deformation and, in certain cases, the recovery required a critical pre-strain to develop [42,43]. When the primary recrystallization developed, the recovery was considered as a structural mechanism responsible for the development of recrystallizing nuclei at an early annealing stage [36]. In this case, the fractional softening by the recovery was negligible, and the strain hardening (i.e. stored energy) could be completely removed by the following recrystallization. In the present study, recovery preceding recrystallization in the samples processed to strain of 2.0 results in the evolution of dislocation free fine (sub)grains with low- to high-angle boundaries (Fig. 8). These (sub)grains are surrounded by dislocation substructures, which look like typical cold-worked ones. Such local formation of perfectly recovered (sub)grains suggests some heterogeneity in the recovery development. The fast recovery occurs in the local volumes with a high density of the deformation (sub)boundaries with high-angle misorientations. The high-angle (sub)boundaries may act as preferable absorption sites for the lattice dislocations. In fact, the increase in the density of grain boundary dislocations does not change the surface energy of high-angle grain boundaries significantly as compared to that in low-
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angle (sub)boundaries [36,44]. Therefore, the inhomogeneity of deformation substructures results in the heterogeneous recovery. Since the volume fraction of perfectly recovered (sub)grains in the samples strained to 2.0 is small at early annealing, these samples are characterised by relatively small fractional softening (Figs. 2 and 13). On the other hand, the samples processed to strain of 4.4 are characterised by the large fraction of high-angle (sub)grain boundaries of about 75%. In other words, almost all (sub)boundaries are high-angle grain boundaries. Such homogeneous deformation substructures that consist of highangle (sub)boundaries are responsible for the rapid recovery development throughout the sample. The high internal stresses in the severely deformed materials are associated with strain-induced (sub)grain boundaries, in turn, the latter ones are easy to be recovered [13,32]. The large strained samples, therefore, shows remarkable fractional softening at the recovery stage during annealing, i.e. just after heating to high temperatures as well as during annealing at relatively low temperatures. 4.2. Recrystallization Recrystallization in the samples strained to 2.0 develops in conventional manner of the primary recrystallization, which is common annealing mechanism operating frequently in cold-worked metals and alloys. Let us discuss, therefore, recrystallization in the samples processed to rather large strain of 4.4, the annealing behaviour of which suggests the continuous recrystallization mechanism. The continuous recrystallization was discussed as an annealing mechanism operating in some alloys containing a relatively large fraction of hard second-phase particles [24,29,45]. The particle growth upon annealing was accompanied by the coarsening of deformation substructures, leading to the equiaxed grained microstructures like recrystallized ones. The present results suggest that the structural changes taking place at the beginning of the continuous recrystallization are qualitatively similar to those occurring at the nucleation stage of the primary recrystallization. Namely, the fast recovery in the ribbon-like deformation substructures at an early annealing results in the dislocation rearrangement within the strain-induced continuously curved crystallites, leading to the break up of the highly elongated (sub)grains into the chains of almost equiaxed (sub)grains (cf. Figs. 4d and 9a). However, compared to the primary recrystallization, where the polygonization leading to recrystallizing nuclei develops heterogeneously, the continuous recrystallization is characterised by the homogeneous nucleation. Also, the rapid increase of the transverse (sub)grain size (Fig. 10) suggests that the recovery of the ribbon-like deformation substructures is accompanied by a merger of the longitudinal (sub)boundaries that results from their local migration. The polygonization and the local (sub)boundary migration during annealing can be accelerated by high internal stresses
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that develop in severe deformed materials. The stress sources are frequently associated with specific lattice defects, disclinations, that evolve at the strain-induced (sub)grain boundaries and their junctions [13,34,46–48]. Release of such internal stresses may be realised by the motion and then annihilation of disclinations. The disclination motion is necessary provided by the generation and/or rearrangement of dislocations. It was theoretically predicted that the disclination motion along the grain boundary associated with the emission of lattice dislocations from the boundary may be energetically favourable [49]. The similar consideration can be applied to discuss the structural changes upon early annealing of the samples strained to 4.4. The bulging (sub)boundaries (Fig. 11) may be composed from the emitted grain boundary dislocations. This process looks like a splitting of the deformation (sub)grain boundaries, leading to the evolution of many low-angle (sub)boundaries and may be considered as a special structural mechanism assisting polygonization. Further annealing results in a coarsening of the polygonized structure much similar to the normal grain growth. It may be concluded, finally, that the continuous recrystallization in the present study involves two simple sequential steps, i.e. polygonization followed by normal grain growth. In the present discussion, the term polygonization is used to emphasize the phenomenon, namely, the quick in situ transformation of ribbon-like deformation substructures into equiaxed (sub)grains. This transformation is accompanied with a sharp increase of the (sub)grain sizes in the transverse direction (Fig. 10) as well as a rapid rise of the number fraction of low-angle (sub)boundaries (Fig. 12). Here, the polygonization stage is essentially identical to the transient recrystallization discussed in the previous study on continuous recrystallization after large strain deformation [20,31]. A relatively fast development kinetics in the present study can be caused by a larger difference between the processing and annealing temperatures. Considering the conditions for continuous recrystallization, the homogeneity of deformation substructures seems to be a key parameter controlling the recrystallization mechanism. In the present case, the uniaxial large strain processing provided the highly uniform deformation substructures, i.e. (sub)grain sizes, dislocation densities, (sub)boundary misorientations, etc. Therefore, the probability for appearance of recrystallizing nuclei is the same for almost all (sub)grains throughout the deformation structures. Then, the simultaneous development of recrystallizing nuclei in place of deformation substructures inhibits any selective growth of separate nuclei, resulting in the homogeneous, continuous recrystallization.
5. Conclusions The annealing behaviour of a cold bar-rolled/swaged Fe–22% Cr–3% Ni ferritic stainless steel was studied at temperatures of 400–700 ◦ C. The main results could be sum-
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marised as follows: 1. The mechanisms of microstructure evolution that operated in cold-worked substructures during annealing depended significantly on the intensity of previous deformation; 2. The samples cold-worked to strain of 2.0 demonstrated conventional primary (discontinuous) recrystallization behaviour during annealing with a critical recrystallization temperature of about 600 ◦ C. Recovery at temperatures below 600 ◦ C did not lead to remarkable softening, while heating to higher temperatures resulted rapidly in complete softening according to the recrystallization development; 3. The annealing behaviour of the samples cold-worked to strain of 4.4 was considered as a continuous recrystallization, which involved recovery followed by normal grain growth. The recovery developed quickly just after the heating and resulted in the evolution of polygonized microstructures in place of ribbon-like deformation substructures. Such polygonized microstructures consisting of low- to high-angle (sub)boundaries were essentially stable against discontinuous coarsening and, therefore, only the normal grain growth operated upon further annealing. The samples were characterised by a gradual softening with increasing the temperature in the all temperature range studied, but did not approach perfect softening even at a rather high temperature of 700 ◦ C; 4. The continuously recrystallized microstructures consisted of much finer grains than those resulted from the discontinuous recrystallization.
Acknowledgements The authors are grateful to Drs. N. Sakuma, T. Hibaru, S. Kuroda, M. Kobayashi, T. Kanno, Steel Research Center, National Institute for Materials Science, for their assistance in the materials processing. One of the authors, A.B., would like to express his thanks to the National Institute for Materials Science for providing a scientific fellowship.
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