- Email: [email protected]
High temperature deformation behavior of two as-cast high-manganese TWIP steels
Materials Science & Engineering A 582 (2013) 15–21
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
High temperature deformation behavior of two as-cast high-manganese TWIP steels A. Khosravifard a, A.S. Hamada b,c, M.M. Moshksar a, R. Ebrahimi a, D.A. Porter b, L.P. Karjalainen b,n a
Department of Materials Science and Engineering, School of Engineering, Shiraz University, Shiraz 71348, Iran Centre for Advanced Steels Research, University of Oulu, PO Box 4200, FI-90014 Oulu, Finland c Metallurgical and Materials Engineering Department, Faculty of Petroleum and Mining Engineering, Suez Canal University, Box 43721, Suez, Egypt b
ar t ic l e i nf o
a b s t r a c t
Article history: Received 24 May 2013 Received in revised form 5 June 2013 Accepted 8 June 2013 Available online 15 June 2013
The high-temperature behavior of two as-cast high-manganese steels with different levels of carbon (0.49 and 0.07 wt%) has been studied by employing hot compression tests at different temperatures (900, 1000 and 1100 1C) and strain rates (0.01, 0.1, 1 and 10/s). Microstructures of the deformed specimens have been examined using SEM–EBSD. The steels are compared in terms of their flow stress level, activation energy of deformation, critical stress and strain for the initiation of softening, and extent of grain refinement. The two steels behave quite differently: flow stress levels at small strains are higher for the high-carbon steel than for the low-carbon one but softening, starting at very small strains, is very pronounced in the former, whereas only slight softening is observed for the low-carbon steel. This peculiar behavior of the high-carbon steel is due to the localization of strain along segregation bands and possibly the presence of ferrite at high temperatures in the highly segregated regions of the cast structure. Effective grain refinement occurs by dynamic recrystallization in both the steels. & 2013 Elsevier B.V. All rights reserved.
Keywords: High-manganese steel As-cast structure Micro-segregation Flow stress Softening Dynamic recrystallization
1. Introduction Among the various austenitic steels, those showing twinning induced plasticity (TWIP) are well known for their excellent ductility at room temperature due to intense strain hardening up to large values of strain [1]. A high-manganese content of about 15–30 wt% is necessary to maintain austenite as the stable phase down to room temperature. Additionally, certain amounts of Al and Si may also be added to adjust the stacking fault energy of the steel to make twinning the dominant deformation mechanism [2]. The generated twins are responsible for the very high hardening rate of TWIP steels during their plastic deformation. Physical and mechanical metallurgy of these steels are comprehensively reviewed in a recent work [3]. As with other metals and alloys, hot working is an integral part of the production of TWIP steels and the investigation of its hot deformation behavior is of industrial importance. Li et al. [4] and Dobrzanski et al. [5–7] have studied the high-temperature behavior of high-manganese steels extensively. Li et al. considered the
n
Corresponding author. Tel.: +358 294 482140; fax: +358 8 553 2165. E-mail addresses: [email protected] (A. Khosravifard), atef.hamada@oulu.fi (A.S. Hamada), [email protected] (M.M. Moshksar), [email protected] (R. Ebrahimi), david.porter@oulu.fi (D.A. Porter), pentti.karjalainen@oulu.fi (L.P. Karjalainen). 0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.06.014
flow stress and dynamic recrystallization (DRX) characteristics of a low-carbon 20Mn–3Si–3Al steel (all concentrations are hereafter in wt%). Dobrzanski et al. investigated the flow stress levels and the conditions for the occurrence of static recrystallization (SRX), DRX and metadynamic recrystallization (MDRX), and microstructure evolution during multi-pass deformation in a low-carbon 26Mn TWIP steel with Ti–Nb microalloying and various Al and Si additions. The flow stress was observed to be quite high but it was effectively reduced by DRX. Hamada et al. [8] confirmed that Al alloying increases the high-temperature flow resistance of the steel and retards the onset of DRX. However, in another work, only a minor influence of Al on the SRX kinetics was found [9]. By conducting double-hit compression tests, Hajkazemi et al. [10] showed that in a low-carbon 29Mn-Si-Al TWIP steel, SRX results in softening of the steel during the inter-pass time, while simultaneously, static strain aging causes a considerable hardening effect. In some very recent work [11,12], the flow resistance of lowcarbon Si–Al-alloyed TWIP steels at a variety of temperatures and deformation rates was modeled by the Arrhenius-type hyperbolic sine equation. All the above-mentioned studies concerned wrought highmanganese steels with low carbon contents (less than 0.15%). Hamada et al. [13] investigated the high-temperature flow stress and recrystallization behavior of several wrought TWIP steels with somewhat higher carbon contents up to 0.27%. In another research
16
A. Khosravifard et al. / Materials Science & Engineering A 582 (2013) 15–21
work, a series of high-carbon wrought TWIP steels with 0.3 and 0.6%C were studied [14]. Similar behavior was reported for these wrought steels, in spite of their different carbon contents. Even though industrial hot rolling starts from a cast structure, there are few studies on the hot deformation behavior of as-cast austenitic steels. Ryan and McQueen [15] investigated the hot deformation behavior of as-cast and wrought Type 316 stainless steel and found a higher hardening rate and lower hot ductility for the as-cast material compared to its worked samples. Fujita et al. [16] and Hotta et al. [17] studied the grain refinement in as-cast austenite of high strength low alloy (HSLA) and a 9% Ni steel, respectively. Higher flow stresses were observed in the as-cast HSLA steel compared to its wrought form, whereas the flow curves were similar for as-cast and wrought 9% Ni steel. Mandal et al. [18] confirmed the occurrence of DRX in as-cast Type 304 stainless steel even without distinct peak stresses in the flow curves. Very recently, Han et al. [19] studied the hot compression of as-cast 904L austenitic stainless steel and found that the process of DRX was sluggish with the increase of strain and occurred inhomogeneously. However, to the best of authors' knowledge, hot deformation of as-cast TWIP steels has not yet been reported in the literature. In the present work, the hot deformation behavior of two as-cast high-manganese TWIP steels with different carbon concentrations (0.49 and 0.07 wt%) are investigated using compression tests conducted at different temperatures and strain rates. The deformed specimens are characterized using light optical and electron microscopy to follow the microstructure evolution. Very different flow behavior was observed in these two steels. Also, the high-carbon steel behaved very differently from corresponding wrought steels and the possible reasons for this are discussed.
3. Results 3.1. Flow curves The compressive flow curves at two strain rates of 0.1 and 1/s are shown in Fig. 1. Both steels show typical effects of temperature and strain rate, but they exhibit very different shapes of flow stress curves at high temperatures. The flow stress of the HC steel is high at small strains and the curves reveal an early peak and intense continuous softening following the peak without a subsequent steady state. Contrary to this, the flow stress curves of the LC steel are similar to those reported for TWIP steels in the wrought condition [4,12], showing a flat shape after initial strain hardening, so that a distinct peak stress is not apparent. 3.2. As-cast microstructures Before testing, the room temperature microstructure of HC steel was fully austenitic with the grain size of approximately 700 μm, while that of LC steel exhibited a duplex mixture of ε-martensite and austenite phases with a grain size of about 800 μm. Neither microstructure contained any ferrite. Of course, at the reheating temperatures, no epsilon martensite exists. However, the level of micro-segregation in the cast structure can influence the hightemperature microstructure. Thus, in order to obtain a measure of micro-segregation, EDS analyses were conducted on both steels. Fig. 2 shows micrographs of the two cast steels and the points where the EDS analyses were taken. A dendritic solidification pattern is clearly seen in the microstructure of HC steel, whereas such was hardly distinguishable in the LC steel. The concentrations of elements at the designated points are given in Table 2.
2. Experimental procedure The chemical compositions of the steels used in this study are given in Table 1. The steels were received in the form of 350 80 60 mm cast ingots. Cylindrical specimens with a length of 12 mm and diameter 10 mm were cut with their axes parallel to the length of the ingot taking care to avoid center-line segregation. Hot compression tests were conducted using a Gleeble 3800 thermo-mechanical simulator. The specimens were heated at the rate of 10 1C/s to the test temperature (900, 1000 and 1100 1C). After soaking for 2 min at the test temperature, the specimens were compressed in a single hit at a constant true strain rate (0.01, 0.1, 1 and 10/s) to the true strain of 0.8. Graphite foil was used as a lubricant on the surfaces of the specimens to minimize friction and barreling effects during compression, while tantalum foil was employed to prevent carbon diffusion into the specimen. The deformed specimens were then cut along the axis of compression and studied using a scanning electron microscope with an electron backscatter diffraction detector (SEM–EBSD). For EBSD examinations, the specimens were polished using a 0.05 mm colloidal suspension of silica after mechanical polishing down to 1 mm by a Struers automatic polisher. In addition, energy dispersive spectroscopy (EDS) was employed to evaluate the microsegregation in the as-cast steel specimens etched in 4% nital.
Table 1 Chemical composition of the investigated steels (wt%). Alloy
C
Mn
Al
Si
Fe
HC steel LC steel
0.49 0.07
21.6 20.1
0.8 0.6
2.7 2.1
Bal. Bal.
Fig. 1. True stress–strain curves of the investigated steels at two strain rates of (a) 0.1/s and (b) 1/s.
A. Khosravifard et al. / Materials Science & Engineering A 582 (2013) 15–21
17
Fig. 2. As-cast solidification structures of (a) HC steel and (b) LC steel showing points where EDS elemental analyses were performed.
Table 2 Concentrations (wt%) of elements in the areas designated in Fig. 2. Analyzed area
Mn
Al
Si
HC steel Zone A Zone B
20.07 25.73
2.24 0.69
2.80 2.67
LC steel Zone A Zone B Zone C
20.02 19.32 20.89
0.91 0.76 0.45
2.32 2.53 2.16
Fig. 3. Optical micrograph of the HC steel, reheated to 1100 1C for 2 min, deformed at 1/s to the strain of 0.8. Compression axis shown by arrows.
Obviously, there exists a higher level of micro-segregation of Mn and Al in the HC steel. Fig. 3 shows severe flow localization and possible traces of ferrite along the segregation bands in a specimen of the HC steel which had been reheated to 1100 1C for 2 min and deformed at 1/s to the strain of 0.8. The arrows in the figure show the direction of compression. The shear localization is similar as recently reported for 904L type stainless steel [19]. Areas with a lower manganese and higher aluminum and silicon contents might be more prone to ferrite formation at high temperatures, even though it should be noted that calculations conducted using the Thermocalc software do not predict this for any of the compositions given in Table 2.
Fig. 4. Critical stresses. (a) Definition (inflection point of the hardening rate vs. stress curve) and (b) effect of temperature at different strain rates for the two steels.
for both steels based on the method presented by Najafizadeh and Jonas [20]. For each test, the strain hardening rate ðds=dεÞ is plotted versus the true stress and the inflection point is selected as the critical stress. This is depicted in Fig. 4(a) for two of the LC steel specimens. The values of critical stress obtained for the two steels at different temperatures and strain rates are presented in Fig. 4(b). It is interesting to compare the above behavior in terms of the Zener-Hollomon parameter (Z).1 At the higher strain rates of 1 and 10/s and the lower temperature of 900 1C (high values of Z), the
3.3. Critical stress/strain and deformation activation energy Critical stresses correspond to the initiation of microstructural changes in the tested material [18]. These values were determined
1 Z ¼ ε_ expðQ =RTÞ, where ε_ is the strain rate, Q, is the activation energy of deformation, R is the universal gas constant, and T is the absolute temperature.
18
A. Khosravifard et al. / Materials Science & Engineering A 582 (2013) 15–21
critical stresses of LC steel are slightly higher than those of HC steel. However, at a lower strain rate of 0.1/s or higher temperature of 1100 1C, the critical stress of the LC steel falls below that of the HC steel (Fig. 4b). Using the flow stress curves at different temperatures and strain rates, the activation energy of deformation for both steels was calculated based on the well-known Arrhenius-type hyperbolic sine equation according to the method described in [21]. The average values of activation energy over the entire range of applied strains are 429 and 411 kJ/mol for the HC and LC steels, respectively. A slightly lower value of 388 kJ/mol has been reported for a wrought Fe–20Mn–3Si–3Al steel [4], whereas values in the range of 390–424 kJ/mol are calculated for a wrought Fe– 21Mn–2.7Si–1.6Al steel [12], and 406 kJ/mol for a Fe–25Mn–3Si–
2.9Al wrought steel [11]. It can be concluded that the values of the activation energy of deformation for the present as-cast highmanganese steels are close to those of similar wrought steels. The same conclusion has been drawn by Fujita et al. [16] in comparing as-cast and wrought HSLA steels. Critical strains for the onset of microstructural evolution can be obtained from the critical stresses determined in the way described above. Based on the calculated activation energy values, critical strains for the HC and LC steels have been calculated and are plotted as a function of ln (Z) in Fig. 5. For comparison, the figure also shows critical strains determined from flow stress curves for two low-carbon wrought TWIP steels (0.04C–27.5Mn– 4.18Si–1.96Al [6] and 0.14C–25.8Mn–0.02Al–0.4Si [8]) and two higher carbon wrought TWIP steels (0.41C–21.2Mn–1.5Al–1.5Si [22] and 0.22C–24Mn–5.7Al–0.48Si [13]). The critical strains of the HC steel are obviously lower than those of the LC steel. The difference increases as Z increases. The dependence of the critical strain on Z follows two different trends. This will be discussed below. 3.4. Grain size evolution
Fig. 5. Linear dependence of the critical strain on ln (Z). Some data from the literature are also included for comparison. Solid and dashed lines refer to the highcarbon and low-carbon steels, respectively.
Examples of EBSD maps corresponding to specimens of the HC steel deformed at different strain rates and temperatures are shown in Fig. 6. It is seen that DRX has occurred for all the specimens deformed under various conditions. DRX in the specimen deformed at 900 1C and 0.01/s (Z ¼ 1:3 1017 ) has effectively refined the grain size to an average diameter of about 6 μm. However, some clearly larger grains containing substructure as indicated by arrows in Fig. 6(a) still remain. Thus, DRX is not yet complete and a uniform grain structure could not be obtained at this temperature, even at the lowest applied strain rate of 0.01/s. At the same low strain rate of 0.01/s, the higher temperature of 1100 1C (Z ¼ 2:1 1014 ) has led to more homogeneous DRX
Fig. 6. EBSD maps corresponding to the HC steel deformed at (a) 900 1C and 0.01/s (Z ¼ 1:3 1017 ), (b) 1100 1C and 0.01/s (Z ¼ 2:1 1014 ), (c) 1100 1C and 1/s (Z ¼ 2:1 1016 ), and (d) 1100 1C and 10/s (Z ¼ 2:1 1017 ).
A. Khosravifard et al. / Materials Science & Engineering A 582 (2013) 15–21
19
the HC steel than for the LC steel (Fig. 1). Similarly, Iza-Mendia et al. [23] also observed a higher initial flow stress for as-cast duplex stainless steel compared to that of the wrought steel. They explained their observation by a constraint effect of the austeniteferrite semi-coherent interface with a K–S orientation relationship in the cast structure. We have no evidence to confirm this for the HC steel in the present instance. However, another mechanism to explain the high flow stress level might be connected with the higher carbon content of the HC steel and thereby a higher number of Mn–C dipoles available for short range ordering, as observed and discussed by Saeed-Akbari et al. [24] in the case of the ambient temperature behavior of TWIP steels. The theoretical ordering index (TOI), as defined in [24], is 0.11 and 0.02 for the HC and LC steels, respectively. It has been stated that the probability of formation of CMn6 clusters in highmanganese steels is considerable when 0.1 oTOI o0.3. Thus, the development of these clusters, which act as obstacles for dislocations, will be effective for the HC steel, but negligible for the LC steel. Interestingly, Hamada et al. [13] also observed higher initial flow stress levels for two steels which were thought to contain ferrite at high temperatures. They attributed the higher flow stress levels to the higher Al concentrations. However, those two steels also had higher carbon contents, about double, compared to the steels showing lower flow stresses. This suggests that some mechanism connected with the carbon content of a TWIP steel could be present increasing the flow resistance at small strains. 4.2. Softening of the as-cast steels Fig. 7. EBSD maps corresponding to LC steel deformed at 1100 1C and (a) 0.01/s (Z ¼ 4:3 1013 ) and (b) 1/s (Z ¼ 4:3 1015 ).
(Fig. 6b). However, a relatively larger average grain size of 29 μm was obtained in this condition. Annealing twins which are normally generated during DRX of austenitic steels are also seen at this temperature (shown by arrows in the figure). At the same high temperature of 1100 1C, but a higher strain rate of 1/s (Z ¼ 2:1 1016 ), an average grain size of 17 μm was obtained. This grain structure also seems quite uniform suggesting that DRX was almost completed. Further increase of the strain rate to 10/s (Z ¼ 2:1 1017 ), as shown in Fig. 6(d), had no significant impact on the final grain size. The high temperature flow curves in Fig. 1 showed that a peak stress could be hardly seen for the LC steel up to the true strain of 0.8. This is consistent with the observation that only partial DRX could be found in the deformed specimens of this steel. EBSD maps for the LC steel deformed at 1100 1C and strain rates of 0.01 and 1/s (Z ¼ 4:3 1013 and Z ¼ 4:3 1015 ) are illustrated in Fig. 7, demonstrating that very different structures can be obtained after cooling to room temperature in this steel. For the specimen deformed at 0.01/s, as shown in Fig. 7(a), a dual-phase structure of austenite and epsilon-martensite is present (the red areas in the phase map represent epsilon-martensite, while the yellow color is for austenite). At the higher strain rate of 1/s, very tiny areas of the specimen were dynamically recrystallized, and no epsilonmartensite could be found in these areas (Fig. 7b).
4. Discussion 4.1. Initial flow stresses of the as-cast steels Regarding the high-temperature deformation behavior of the studied TWIP steels, it is seen that the flow stress level at small strains (before the initiation of softening) was markedly higher for
The flow curves of the HC steel showed distinct peak stress at very small strains after which pronounced softening was observed (Fig. 1). In some of the tests, especially at higher temperatures, no steady state is reached up to the strain of 0.8 used in the tests. Contrary to this, compression tests performed on a wrought 22Mn–0.6C TWIP steel have shown only slight softening with much larger peak strains [14]. Thus, it can be inferred that the hot deformation behavior of the as-cast HC steel is different from that of wrought TWIP steels of similar compositions. Continuous softening after very small strain during hot compression of an as-cast 904L stainless steel was observed by Han et al. [19], and localization of strain due to chemical inhomogeneity was suggested to be the reason for this. Similarly, strain localization has been observed in the deformed specimen of the HC steel (Fig. 3) and this can be responsible for the peculiar intense softening. No such strain localization was present in the LC steel, where the microsegregation was less severe. The flow curves reported by Reyes-Calderon et al. [22] for a wrought 21Mn–1.5Al–1.5Si–0.41C TWIP steel also showed peaks at very small strains. No explanation was, however, suggested for the early softening. As presented in Fig. 5, Hamada et al. [13] also observed softening at very small strains for a 0.22C–24Mn–5.7Al– 0.48Si TWIP steel, and the presence of ferrite was suggested as the reason. Similar shapes of flow stress curves have been observed for a cast duplex austenite–ferrite stainless steel [23]. Also, Ryan and McQueen [13] suggested that segregated delta-ferrite areas led to more rapid dislocation accumulation and hence lower peak strains in an as-cast Type 316 stainless steel. On the other hand, it is well established that dynamic recovery (DRV) occurs more readily in ferrite than in austenite [23]. Accordingly, in addition to the localization of strain which was discussed above, the peculiar softening behavior of the present as-cast HC steel might also result from the existence of ferrite at high temperatures due to the intense micro-segregation. According to the EDS data represented in Table 2, a higher level of micro-segregation of manganese and aluminum existed in the HC steel than in the LC steel. Higher carbon content in the
20
A. Khosravifard et al. / Materials Science & Engineering A 582 (2013) 15–21
Fe–Mn–C alloy system leads to a lower solidus temperature [25]. On the other hand, it is well-known that as the solidus temperature decreases, the elemental segregation level increases. Therefore, the degree of micro-segregation is expected to be higher in the steel with a higher carbon content [26]. Segregated areas with a lower manganese and higher aluminum and silicon contents are prone to ferrite formation at high temperatures. Even though, only minor fraction of ferrite along the segregated bands of the microstructure may be seen at room temperature, it has been suggested that this can due to transformation to austenite on cooling and that a considerable amount of the ferrite might be present in the high-temperature structure [13]. Unlike for the HC steel, the flow stress curves of the LC steel do not exhibit pronounced softening and a distinct peak is rarely observed. However, it should be noted that non-existence of a peak stress does not necessarily imply the absence of DRX. There are many studies reporting partial or even complete DRX in lowcarbon austenitic steels without a distinct peak stress and steady state, e.g. [4,12,18]. Thus, critical stresses for the initiation of DRX were determined by the method shown in Fig. 4(a). At the highest values of Z, softening for the LC steel initiates after relatively high strains (see Figs. 1b and 5). As a result, the critical stresses of the LC steel are slightly higher than those of the HC steel (Fig. 4b). As seen in Fig. 5, Hamada et al. [8] and Dobrzanski et al. [6] have also reported high critical strains (close to those of the present LC steel) for wrought low-carbon TWIP steels, especially at large Z values. From Fig. 5, it is obvious that quite a big scatter is present concerning critical strains, so that for instance at ln Z ¼35, the critical strain varies from 0.15 to 0.50, being below 0.10 for the HC steel. Interestingly, it seems that there are two trends: three sets of data (LC and Refs. 6 and 8) show a much higher dependence on Z than the other three (HC and Refs. 13 and 22). Anyhow, it is seen that the critical strains for the HC steel are drastically lower than those for the LC steel, and the difference is more noticeable at high Z values (i.e. high strain rates as seen in Fig. 1b). However, because the localization of strain along the segregation bands and the possible presence of ferrite in the HC steel at high temperatures seem to contribute to the softening of the steel, the critical stresses and strains of this steel are not necessarily related to the onset of DRX. In addition to the localization of strain and the possible presence of ferrite, another reason for the difference in the critical strains of these HC and LC steels might be connected to the higher density of carbon-manganese dipoles in the steels with higher carbon contents. It is well known that these dipoles effectively increase the work hardening rate of a high-manganese steel by pinning its dislocations [27]. Thus, the C–Mn dipoles could lead to a larger work hardening rate upon deformation of steels with higher carbon contents and thereby enhance the driving force for the occurrence of softening. However, this cannot be the only reason for the lower critical strains of the HC steel, since it should work as well in wrought steels, but the flow curves which are obtained for a wrought 22Mn–0.6C TWIP steel show clearly larger critical strains [14]. From Fig. 1, it is seen that the flow stress levels at typical rolling pass strains of 0.2–0.3 [6] are about 100–270 MPa depending on the temperature (strain rate 1/s), i.e. quite high and somewhat higher for the HC steel compared to the LC steel. However, the critical strains for the HC steel are small enough for pronounced softening to be obtained at strains of 0.3 and beyond. 4.3. Grain refinement by DRX The dynamically recrystallized average grain size is not dependent on the initial grain size if the imposed strain is sufficient [28]. Therefore, despite the very large initial grain sizes of the investigated
as-cast steels, the refined grain sizes are expected to be comparable to those of similar wrought steels. However, it has been shown that the strain distribution in the hot-deformed as-cast steel is relatively heterogeneous [23]. The occurrence of DRX in specimens of the HC steel, which were deformed at different values of Z, was confirmed by the EBSD maps shown in Fig. 6(a–d). It is known that the DRX grain size is inversely related to Z [29]. In agreement with this, smaller DRX grains were obtained for this steel at higher strain rates and lower temperatures. However, at the same time, increasing Z led to greater inhomogeneity of the structure. It is worth noting that although the specimens corresponding to Fig. 6(a and d) have been tested at close values of Z, a much more uniform structure was obtained in the latter condition, i.e. at the higher temperature. Thus, it can be inferred that uniformity of the hot deformed structure, i.e. the extent of DRX, is much more influenced by temperature than strain rate. The hot deformed microstructure of the LC steel was different from that of the HC steel. High-temperature flow curves (Fig. 1) showed that a distinct peak stress could be hardly seen for LC steel up to the true strain of 0.8. In consistence with this, only partial DRX could be found in some of the hot-deformed specimens of this steel. Fernandez et al. [29] have shown that the initial austenite grain size affects the kinetics of DRX. Thus, the extremely coarse grain structure of the LC steel in the as-cast condition ( 800 mm) can explain the slow kinetics of microstructural developments in this steel. However, in the HC steel, the DRX kinetics seems to be much faster. In addition to the effect of strain localization and possible presence of ferrite on softening, it has been suggested that carbon interstitial atoms cause lattice expansion, so that diffusion-controlled microstructural evolutions are more enhanced in a steel with a higher carbon content [30]. Due to the lower carbon content in the LC steel, the roomtemperature stacking fault energy (SFE) for this steel is lower than that of the HC steel [31]. As a result, a considerable amount of epsilon martensite is present in the room temperature microstructure of the LC steel (Figs. 2 and 7a). However, Hamada et al. [32] showed that the formation of epsilon martensite in a lowcarbon TWIP steel (Fe–0.14C–26Mn) is dependent on the grain size. The grain boundaries act as obstacles to the transformation of austenite to epsilon martensite so that below a certain grain size, this transformation may be totally suppressed [32,33]. As was expected, the DRX grain size after deformation at 0.01/s was larger than that obtained at the strain rate of 1/s, when the average grain size was about 8 μm. Martensite formation was suppressed only in the latter case (Fig. 7b).
5. Summary The hot deformation behavior of two as-cast coarse-grained TWIP steels with different levels of carbon (0.49% and 0.07%) has been studied by conducting hot compression tests at different temperatures and strain rates. The following conclusions can be drawn from the work: ▪ At lower strains, flow stress levels for the 0.49% C steel were higher than those of the 0.07% C steel. The reason is still unclear, but might be connected with higher carbon content in the former providing a larger number of Mn–C dipoles resulting in short range ordering. ▪ In the cast steel with 0.49% C, softening was very pronounced starting at very small strains and not reaching a steady state up to the maximum applied strain of 0.8, whereas in the cast steel with 0.07% C, the flow stress curves showed no clear peaks. ▪ The level of interdendritic micro-segregation was higher for the high-carbon steel and led to the formation of strain localization
A. Khosravifard et al. / Materials Science & Engineering A 582 (2013) 15–21
▪
▪
▪
▪
bands. Formation of ferrite at high temperatures in the segregation bands with lower Mn and higher Al and Si might also be possible, though not predicted by Thermocalc software. The localization of strain and the possible presence of ferrite are suggested as the reasons for the early and pronounced softening of this steel. The activation energy of deformation of the high-carbon steel is only slightly higher than that of the low-carbon steel. However, both values are comparable to those reported for wrought TWIP steels. In the high-carbon steel deformed at various strain rates and temperatures, the occurrence of dynamic recrystallization (DRX) was proven by microstructural studies. However, only partial DRX could be observed in some of the low-carbon steel specimens. The large grain size of the as-cast structures was effectively refined by DRX in both steels. With increasing strain rate, the DRX grain size became smaller. In the low-carbon steel, the formation of epsilon-martensite during cooling was suppressed after the attainment of a sufficiently fine austenite grain size.
References [1] O. Bouaziz, S. Allain, Y. Estrin, Scr. Mater. 62 (2010) 713–715. [2] I. Gutierrez-Urrutia, S. Zaefferer, D. Raabe, Mater. Sci. Eng. A 527 (2010) 3552–3560. [3] L. Chen, Y. Zhao, X. Qin, Acta Metall. Sin. (Engl. Lett.) 26 (2013) 1–15. [4] D. Li, Y. Feng, Z. Yin, F. Shangguan, K. Wang, Q. Liu, F. Hu, Mater. Des. 34 (2012) 713–718. [5] L.A. Dobrzanski, A. Grajcar, W. Borek, J. Achiev. Mater. Manuf. Eng. 29 (2008) 139–142. [6] L.A. Dobrzanski, A. Grajcar, W. Borek, Arch. Mater. Sci. Eng. 37 (2009) 69–76. [7] L.A. Dobrzanski, W. Borek, J. Achiev. Mater. Manuf. Eng. 46 (2011) 71–78.
21
[8] A.S. Hamada, L.P. Karjalainen, M.C. Somani, Mater. Sci. Eng. A 467 (2007) 114–124. [9] A.S. Hamada, M.C. Somani, L.P. Karjalainen, ISIJ Int. 47 (2007) 907–912. [10] J. Hajkazemi, A. Zarei-Hanzaki, M. Sabet, S. Khoddam, Mater. Sci. Eng. A 530 (2011) 233–238. [11] D. Li, Y. Feng, Z. Yin, F. Shangguan, K. Wang, Q. Liu, F. Hu, Mater. Sci. Eng. A 528 (2011) 8084–8089. [12] A. Marandi, A. Zarei-Hanzaki, N. Haghdadi, M. Eskandari, Mater. Sci. Eng. A 554 (2012) 72–78. [13] A.S. Hamada, L.P. Karjalainen, M.C. Somani, Can. Metall. Quart. 46 (2007) 47–56. [14] A. Ferraiuolo, A. Smith, J.G. Sevillano, F. de las Cuevas, G. Pratolongo, H. Gouveia, M. Mendes Rodrigues, P. Karjalainen, Metallurgical design of high strength austenitic Fe–C–Mn steels with excellent formability, Contract RFSRCT-00030, Research Programme of the Research Fund for Coal and Steel, European Union (2005-2008). [15] N.D. Ryan, H.J. McQueen, J. Mater. Process. Technol. 21 (1990) 177–199. [16] N. Fujita, T. Narushima, Y. Iguchi, C. Ouchi, ISIJ Int. 43 (2003) 1063–1072. [17] S. Hotta, T. Murakami, T. Narushima, Y. Iguchi, C. Ouchi, ISIJ Int. 45 (2005) 338–346. [18] G.K. Mandal, N. Stanford, P. Hodgson, J.H. Beynon, Mater. Sci. Eng. A 556 (2012) 685–695. [19] Y. Han, G. Liu, D. Zou, R. Liu, G. Qia, Mater. Sci. Eng. A 565 (2013) 342–350. [20] A. Najafizadeh, J.J. Jonas, ISIJ Int. 46 (2006) 1679–1684. [21] X. Xiao, G.Q. Liu, B.F. Hu, X. Zheng, L.N. Wang, S.J. Chen, A. Ullah, Comput. Mater. Sci. 62 (2012) 227–234. [22] F. Reyes-Calderon, I. Mejia, J.M. Cabrera, Mater. Sci. Eng. A 562 (2013) 46–52. [23] A. Iza-Mendia, A. Pinol-Juez, J.J. Urcola, I. Gutierrez, Metall. Mater. Trans. A 29 (1998) 2975–2986. [24] A. Saeed-Akbari, L. Mosecker, A. Schwedt, W. Bleck, Metall. Mater. Trans. A 43 (2012) 1688–1704. [25] G.I. Sil’man, Met. Sci. Heat Treat. 47 (2005) 397–401. [26] M. Wolf, W. Kurz, Metall. Trans. B 12 (1981) 85–93. [27] B. Hutchinson, N. Ridley, Scr. Mater. 55 (2006) 299–302. [28] Z. Jin, Z. Cui, J. Mater. Sci. Technol. 26 (2010) 1063–1070. [29] A.I. Fernandez, P. Uranga, B. Lopez, J.M. Rodriguez-Ibabe, Mater. Sci. Eng. A 361 (2003) 367–376. [30] S. Serajzadeh, A. Karimi Taheri, Mech. Mater. 35 (2003) 653–660. [31] A. Saeed-Akbari, J. Imlau, U. Prahl, W. Bleck, Metall. Mater. Trans. A 40 (2009) 3076–3090. [32] A.S. Hamada, P. Sahu, S. Ghosh Chowdhury, L.P. Karjalainen, J. Levoska, T. Oittinen, Metall. Mater. Trans. A 39 (2008) 462–465. [33] S. Takaki, H. Nakatsu, Y. Tokunaga, Mater. Trans. 34 (1993) 489–495.
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
High temperature deformation behavior of two as-cast high-manganese TWIP steels A. Khosravifard a, A.S. Hamada b,c, M.M. Moshksar a, R. Ebrahimi a, D.A. Porter b, L.P. Karjalainen b,n a
Department of Materials Science and Engineering, School of Engineering, Shiraz University, Shiraz 71348, Iran Centre for Advanced Steels Research, University of Oulu, PO Box 4200, FI-90014 Oulu, Finland c Metallurgical and Materials Engineering Department, Faculty of Petroleum and Mining Engineering, Suez Canal University, Box 43721, Suez, Egypt b
ar t ic l e i nf o
a b s t r a c t
Article history: Received 24 May 2013 Received in revised form 5 June 2013 Accepted 8 June 2013 Available online 15 June 2013
The high-temperature behavior of two as-cast high-manganese steels with different levels of carbon (0.49 and 0.07 wt%) has been studied by employing hot compression tests at different temperatures (900, 1000 and 1100 1C) and strain rates (0.01, 0.1, 1 and 10/s). Microstructures of the deformed specimens have been examined using SEM–EBSD. The steels are compared in terms of their flow stress level, activation energy of deformation, critical stress and strain for the initiation of softening, and extent of grain refinement. The two steels behave quite differently: flow stress levels at small strains are higher for the high-carbon steel than for the low-carbon one but softening, starting at very small strains, is very pronounced in the former, whereas only slight softening is observed for the low-carbon steel. This peculiar behavior of the high-carbon steel is due to the localization of strain along segregation bands and possibly the presence of ferrite at high temperatures in the highly segregated regions of the cast structure. Effective grain refinement occurs by dynamic recrystallization in both the steels. & 2013 Elsevier B.V. All rights reserved.
Keywords: High-manganese steel As-cast structure Micro-segregation Flow stress Softening Dynamic recrystallization
1. Introduction Among the various austenitic steels, those showing twinning induced plasticity (TWIP) are well known for their excellent ductility at room temperature due to intense strain hardening up to large values of strain [1]. A high-manganese content of about 15–30 wt% is necessary to maintain austenite as the stable phase down to room temperature. Additionally, certain amounts of Al and Si may also be added to adjust the stacking fault energy of the steel to make twinning the dominant deformation mechanism [2]. The generated twins are responsible for the very high hardening rate of TWIP steels during their plastic deformation. Physical and mechanical metallurgy of these steels are comprehensively reviewed in a recent work [3]. As with other metals and alloys, hot working is an integral part of the production of TWIP steels and the investigation of its hot deformation behavior is of industrial importance. Li et al. [4] and Dobrzanski et al. [5–7] have studied the high-temperature behavior of high-manganese steels extensively. Li et al. considered the
n
Corresponding author. Tel.: +358 294 482140; fax: +358 8 553 2165. E-mail addresses: [email protected] (A. Khosravifard), atef.hamada@oulu.fi (A.S. Hamada), [email protected] (M.M. Moshksar), [email protected] (R. Ebrahimi), david.porter@oulu.fi (D.A. Porter), pentti.karjalainen@oulu.fi (L.P. Karjalainen). 0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.06.014
flow stress and dynamic recrystallization (DRX) characteristics of a low-carbon 20Mn–3Si–3Al steel (all concentrations are hereafter in wt%). Dobrzanski et al. investigated the flow stress levels and the conditions for the occurrence of static recrystallization (SRX), DRX and metadynamic recrystallization (MDRX), and microstructure evolution during multi-pass deformation in a low-carbon 26Mn TWIP steel with Ti–Nb microalloying and various Al and Si additions. The flow stress was observed to be quite high but it was effectively reduced by DRX. Hamada et al. [8] confirmed that Al alloying increases the high-temperature flow resistance of the steel and retards the onset of DRX. However, in another work, only a minor influence of Al on the SRX kinetics was found [9]. By conducting double-hit compression tests, Hajkazemi et al. [10] showed that in a low-carbon 29Mn-Si-Al TWIP steel, SRX results in softening of the steel during the inter-pass time, while simultaneously, static strain aging causes a considerable hardening effect. In some very recent work [11,12], the flow resistance of lowcarbon Si–Al-alloyed TWIP steels at a variety of temperatures and deformation rates was modeled by the Arrhenius-type hyperbolic sine equation. All the above-mentioned studies concerned wrought highmanganese steels with low carbon contents (less than 0.15%). Hamada et al. [13] investigated the high-temperature flow stress and recrystallization behavior of several wrought TWIP steels with somewhat higher carbon contents up to 0.27%. In another research
16
A. Khosravifard et al. / Materials Science & Engineering A 582 (2013) 15–21
work, a series of high-carbon wrought TWIP steels with 0.3 and 0.6%C were studied [14]. Similar behavior was reported for these wrought steels, in spite of their different carbon contents. Even though industrial hot rolling starts from a cast structure, there are few studies on the hot deformation behavior of as-cast austenitic steels. Ryan and McQueen [15] investigated the hot deformation behavior of as-cast and wrought Type 316 stainless steel and found a higher hardening rate and lower hot ductility for the as-cast material compared to its worked samples. Fujita et al. [16] and Hotta et al. [17] studied the grain refinement in as-cast austenite of high strength low alloy (HSLA) and a 9% Ni steel, respectively. Higher flow stresses were observed in the as-cast HSLA steel compared to its wrought form, whereas the flow curves were similar for as-cast and wrought 9% Ni steel. Mandal et al. [18] confirmed the occurrence of DRX in as-cast Type 304 stainless steel even without distinct peak stresses in the flow curves. Very recently, Han et al. [19] studied the hot compression of as-cast 904L austenitic stainless steel and found that the process of DRX was sluggish with the increase of strain and occurred inhomogeneously. However, to the best of authors' knowledge, hot deformation of as-cast TWIP steels has not yet been reported in the literature. In the present work, the hot deformation behavior of two as-cast high-manganese TWIP steels with different carbon concentrations (0.49 and 0.07 wt%) are investigated using compression tests conducted at different temperatures and strain rates. The deformed specimens are characterized using light optical and electron microscopy to follow the microstructure evolution. Very different flow behavior was observed in these two steels. Also, the high-carbon steel behaved very differently from corresponding wrought steels and the possible reasons for this are discussed.
3. Results 3.1. Flow curves The compressive flow curves at two strain rates of 0.1 and 1/s are shown in Fig. 1. Both steels show typical effects of temperature and strain rate, but they exhibit very different shapes of flow stress curves at high temperatures. The flow stress of the HC steel is high at small strains and the curves reveal an early peak and intense continuous softening following the peak without a subsequent steady state. Contrary to this, the flow stress curves of the LC steel are similar to those reported for TWIP steels in the wrought condition [4,12], showing a flat shape after initial strain hardening, so that a distinct peak stress is not apparent. 3.2. As-cast microstructures Before testing, the room temperature microstructure of HC steel was fully austenitic with the grain size of approximately 700 μm, while that of LC steel exhibited a duplex mixture of ε-martensite and austenite phases with a grain size of about 800 μm. Neither microstructure contained any ferrite. Of course, at the reheating temperatures, no epsilon martensite exists. However, the level of micro-segregation in the cast structure can influence the hightemperature microstructure. Thus, in order to obtain a measure of micro-segregation, EDS analyses were conducted on both steels. Fig. 2 shows micrographs of the two cast steels and the points where the EDS analyses were taken. A dendritic solidification pattern is clearly seen in the microstructure of HC steel, whereas such was hardly distinguishable in the LC steel. The concentrations of elements at the designated points are given in Table 2.
2. Experimental procedure The chemical compositions of the steels used in this study are given in Table 1. The steels were received in the form of 350 80 60 mm cast ingots. Cylindrical specimens with a length of 12 mm and diameter 10 mm were cut with their axes parallel to the length of the ingot taking care to avoid center-line segregation. Hot compression tests were conducted using a Gleeble 3800 thermo-mechanical simulator. The specimens were heated at the rate of 10 1C/s to the test temperature (900, 1000 and 1100 1C). After soaking for 2 min at the test temperature, the specimens were compressed in a single hit at a constant true strain rate (0.01, 0.1, 1 and 10/s) to the true strain of 0.8. Graphite foil was used as a lubricant on the surfaces of the specimens to minimize friction and barreling effects during compression, while tantalum foil was employed to prevent carbon diffusion into the specimen. The deformed specimens were then cut along the axis of compression and studied using a scanning electron microscope with an electron backscatter diffraction detector (SEM–EBSD). For EBSD examinations, the specimens were polished using a 0.05 mm colloidal suspension of silica after mechanical polishing down to 1 mm by a Struers automatic polisher. In addition, energy dispersive spectroscopy (EDS) was employed to evaluate the microsegregation in the as-cast steel specimens etched in 4% nital.
Table 1 Chemical composition of the investigated steels (wt%). Alloy
C
Mn
Al
Si
Fe
HC steel LC steel
0.49 0.07
21.6 20.1
0.8 0.6
2.7 2.1
Bal. Bal.
Fig. 1. True stress–strain curves of the investigated steels at two strain rates of (a) 0.1/s and (b) 1/s.
A. Khosravifard et al. / Materials Science & Engineering A 582 (2013) 15–21
17
Fig. 2. As-cast solidification structures of (a) HC steel and (b) LC steel showing points where EDS elemental analyses were performed.
Table 2 Concentrations (wt%) of elements in the areas designated in Fig. 2. Analyzed area
Mn
Al
Si
HC steel Zone A Zone B
20.07 25.73
2.24 0.69
2.80 2.67
LC steel Zone A Zone B Zone C
20.02 19.32 20.89
0.91 0.76 0.45
2.32 2.53 2.16
Fig. 3. Optical micrograph of the HC steel, reheated to 1100 1C for 2 min, deformed at 1/s to the strain of 0.8. Compression axis shown by arrows.
Obviously, there exists a higher level of micro-segregation of Mn and Al in the HC steel. Fig. 3 shows severe flow localization and possible traces of ferrite along the segregation bands in a specimen of the HC steel which had been reheated to 1100 1C for 2 min and deformed at 1/s to the strain of 0.8. The arrows in the figure show the direction of compression. The shear localization is similar as recently reported for 904L type stainless steel [19]. Areas with a lower manganese and higher aluminum and silicon contents might be more prone to ferrite formation at high temperatures, even though it should be noted that calculations conducted using the Thermocalc software do not predict this for any of the compositions given in Table 2.
Fig. 4. Critical stresses. (a) Definition (inflection point of the hardening rate vs. stress curve) and (b) effect of temperature at different strain rates for the two steels.
for both steels based on the method presented by Najafizadeh and Jonas [20]. For each test, the strain hardening rate ðds=dεÞ is plotted versus the true stress and the inflection point is selected as the critical stress. This is depicted in Fig. 4(a) for two of the LC steel specimens. The values of critical stress obtained for the two steels at different temperatures and strain rates are presented in Fig. 4(b). It is interesting to compare the above behavior in terms of the Zener-Hollomon parameter (Z).1 At the higher strain rates of 1 and 10/s and the lower temperature of 900 1C (high values of Z), the
3.3. Critical stress/strain and deformation activation energy Critical stresses correspond to the initiation of microstructural changes in the tested material [18]. These values were determined
1 Z ¼ ε_ expðQ =RTÞ, where ε_ is the strain rate, Q, is the activation energy of deformation, R is the universal gas constant, and T is the absolute temperature.
18
A. Khosravifard et al. / Materials Science & Engineering A 582 (2013) 15–21
critical stresses of LC steel are slightly higher than those of HC steel. However, at a lower strain rate of 0.1/s or higher temperature of 1100 1C, the critical stress of the LC steel falls below that of the HC steel (Fig. 4b). Using the flow stress curves at different temperatures and strain rates, the activation energy of deformation for both steels was calculated based on the well-known Arrhenius-type hyperbolic sine equation according to the method described in [21]. The average values of activation energy over the entire range of applied strains are 429 and 411 kJ/mol for the HC and LC steels, respectively. A slightly lower value of 388 kJ/mol has been reported for a wrought Fe–20Mn–3Si–3Al steel [4], whereas values in the range of 390–424 kJ/mol are calculated for a wrought Fe– 21Mn–2.7Si–1.6Al steel [12], and 406 kJ/mol for a Fe–25Mn–3Si–
2.9Al wrought steel [11]. It can be concluded that the values of the activation energy of deformation for the present as-cast highmanganese steels are close to those of similar wrought steels. The same conclusion has been drawn by Fujita et al. [16] in comparing as-cast and wrought HSLA steels. Critical strains for the onset of microstructural evolution can be obtained from the critical stresses determined in the way described above. Based on the calculated activation energy values, critical strains for the HC and LC steels have been calculated and are plotted as a function of ln (Z) in Fig. 5. For comparison, the figure also shows critical strains determined from flow stress curves for two low-carbon wrought TWIP steels (0.04C–27.5Mn– 4.18Si–1.96Al [6] and 0.14C–25.8Mn–0.02Al–0.4Si [8]) and two higher carbon wrought TWIP steels (0.41C–21.2Mn–1.5Al–1.5Si [22] and 0.22C–24Mn–5.7Al–0.48Si [13]). The critical strains of the HC steel are obviously lower than those of the LC steel. The difference increases as Z increases. The dependence of the critical strain on Z follows two different trends. This will be discussed below. 3.4. Grain size evolution
Fig. 5. Linear dependence of the critical strain on ln (Z). Some data from the literature are also included for comparison. Solid and dashed lines refer to the highcarbon and low-carbon steels, respectively.
Examples of EBSD maps corresponding to specimens of the HC steel deformed at different strain rates and temperatures are shown in Fig. 6. It is seen that DRX has occurred for all the specimens deformed under various conditions. DRX in the specimen deformed at 900 1C and 0.01/s (Z ¼ 1:3 1017 ) has effectively refined the grain size to an average diameter of about 6 μm. However, some clearly larger grains containing substructure as indicated by arrows in Fig. 6(a) still remain. Thus, DRX is not yet complete and a uniform grain structure could not be obtained at this temperature, even at the lowest applied strain rate of 0.01/s. At the same low strain rate of 0.01/s, the higher temperature of 1100 1C (Z ¼ 2:1 1014 ) has led to more homogeneous DRX
Fig. 6. EBSD maps corresponding to the HC steel deformed at (a) 900 1C and 0.01/s (Z ¼ 1:3 1017 ), (b) 1100 1C and 0.01/s (Z ¼ 2:1 1014 ), (c) 1100 1C and 1/s (Z ¼ 2:1 1016 ), and (d) 1100 1C and 10/s (Z ¼ 2:1 1017 ).
A. Khosravifard et al. / Materials Science & Engineering A 582 (2013) 15–21
19
the HC steel than for the LC steel (Fig. 1). Similarly, Iza-Mendia et al. [23] also observed a higher initial flow stress for as-cast duplex stainless steel compared to that of the wrought steel. They explained their observation by a constraint effect of the austeniteferrite semi-coherent interface with a K–S orientation relationship in the cast structure. We have no evidence to confirm this for the HC steel in the present instance. However, another mechanism to explain the high flow stress level might be connected with the higher carbon content of the HC steel and thereby a higher number of Mn–C dipoles available for short range ordering, as observed and discussed by Saeed-Akbari et al. [24] in the case of the ambient temperature behavior of TWIP steels. The theoretical ordering index (TOI), as defined in [24], is 0.11 and 0.02 for the HC and LC steels, respectively. It has been stated that the probability of formation of CMn6 clusters in highmanganese steels is considerable when 0.1 oTOI o0.3. Thus, the development of these clusters, which act as obstacles for dislocations, will be effective for the HC steel, but negligible for the LC steel. Interestingly, Hamada et al. [13] also observed higher initial flow stress levels for two steels which were thought to contain ferrite at high temperatures. They attributed the higher flow stress levels to the higher Al concentrations. However, those two steels also had higher carbon contents, about double, compared to the steels showing lower flow stresses. This suggests that some mechanism connected with the carbon content of a TWIP steel could be present increasing the flow resistance at small strains. 4.2. Softening of the as-cast steels Fig. 7. EBSD maps corresponding to LC steel deformed at 1100 1C and (a) 0.01/s (Z ¼ 4:3 1013 ) and (b) 1/s (Z ¼ 4:3 1015 ).
(Fig. 6b). However, a relatively larger average grain size of 29 μm was obtained in this condition. Annealing twins which are normally generated during DRX of austenitic steels are also seen at this temperature (shown by arrows in the figure). At the same high temperature of 1100 1C, but a higher strain rate of 1/s (Z ¼ 2:1 1016 ), an average grain size of 17 μm was obtained. This grain structure also seems quite uniform suggesting that DRX was almost completed. Further increase of the strain rate to 10/s (Z ¼ 2:1 1017 ), as shown in Fig. 6(d), had no significant impact on the final grain size. The high temperature flow curves in Fig. 1 showed that a peak stress could be hardly seen for the LC steel up to the true strain of 0.8. This is consistent with the observation that only partial DRX could be found in the deformed specimens of this steel. EBSD maps for the LC steel deformed at 1100 1C and strain rates of 0.01 and 1/s (Z ¼ 4:3 1013 and Z ¼ 4:3 1015 ) are illustrated in Fig. 7, demonstrating that very different structures can be obtained after cooling to room temperature in this steel. For the specimen deformed at 0.01/s, as shown in Fig. 7(a), a dual-phase structure of austenite and epsilon-martensite is present (the red areas in the phase map represent epsilon-martensite, while the yellow color is for austenite). At the higher strain rate of 1/s, very tiny areas of the specimen were dynamically recrystallized, and no epsilonmartensite could be found in these areas (Fig. 7b).
4. Discussion 4.1. Initial flow stresses of the as-cast steels Regarding the high-temperature deformation behavior of the studied TWIP steels, it is seen that the flow stress level at small strains (before the initiation of softening) was markedly higher for
The flow curves of the HC steel showed distinct peak stress at very small strains after which pronounced softening was observed (Fig. 1). In some of the tests, especially at higher temperatures, no steady state is reached up to the strain of 0.8 used in the tests. Contrary to this, compression tests performed on a wrought 22Mn–0.6C TWIP steel have shown only slight softening with much larger peak strains [14]. Thus, it can be inferred that the hot deformation behavior of the as-cast HC steel is different from that of wrought TWIP steels of similar compositions. Continuous softening after very small strain during hot compression of an as-cast 904L stainless steel was observed by Han et al. [19], and localization of strain due to chemical inhomogeneity was suggested to be the reason for this. Similarly, strain localization has been observed in the deformed specimen of the HC steel (Fig. 3) and this can be responsible for the peculiar intense softening. No such strain localization was present in the LC steel, where the microsegregation was less severe. The flow curves reported by Reyes-Calderon et al. [22] for a wrought 21Mn–1.5Al–1.5Si–0.41C TWIP steel also showed peaks at very small strains. No explanation was, however, suggested for the early softening. As presented in Fig. 5, Hamada et al. [13] also observed softening at very small strains for a 0.22C–24Mn–5.7Al– 0.48Si TWIP steel, and the presence of ferrite was suggested as the reason. Similar shapes of flow stress curves have been observed for a cast duplex austenite–ferrite stainless steel [23]. Also, Ryan and McQueen [13] suggested that segregated delta-ferrite areas led to more rapid dislocation accumulation and hence lower peak strains in an as-cast Type 316 stainless steel. On the other hand, it is well established that dynamic recovery (DRV) occurs more readily in ferrite than in austenite [23]. Accordingly, in addition to the localization of strain which was discussed above, the peculiar softening behavior of the present as-cast HC steel might also result from the existence of ferrite at high temperatures due to the intense micro-segregation. According to the EDS data represented in Table 2, a higher level of micro-segregation of manganese and aluminum existed in the HC steel than in the LC steel. Higher carbon content in the
20
A. Khosravifard et al. / Materials Science & Engineering A 582 (2013) 15–21
Fe–Mn–C alloy system leads to a lower solidus temperature [25]. On the other hand, it is well-known that as the solidus temperature decreases, the elemental segregation level increases. Therefore, the degree of micro-segregation is expected to be higher in the steel with a higher carbon content [26]. Segregated areas with a lower manganese and higher aluminum and silicon contents are prone to ferrite formation at high temperatures. Even though, only minor fraction of ferrite along the segregated bands of the microstructure may be seen at room temperature, it has been suggested that this can due to transformation to austenite on cooling and that a considerable amount of the ferrite might be present in the high-temperature structure [13]. Unlike for the HC steel, the flow stress curves of the LC steel do not exhibit pronounced softening and a distinct peak is rarely observed. However, it should be noted that non-existence of a peak stress does not necessarily imply the absence of DRX. There are many studies reporting partial or even complete DRX in lowcarbon austenitic steels without a distinct peak stress and steady state, e.g. [4,12,18]. Thus, critical stresses for the initiation of DRX were determined by the method shown in Fig. 4(a). At the highest values of Z, softening for the LC steel initiates after relatively high strains (see Figs. 1b and 5). As a result, the critical stresses of the LC steel are slightly higher than those of the HC steel (Fig. 4b). As seen in Fig. 5, Hamada et al. [8] and Dobrzanski et al. [6] have also reported high critical strains (close to those of the present LC steel) for wrought low-carbon TWIP steels, especially at large Z values. From Fig. 5, it is obvious that quite a big scatter is present concerning critical strains, so that for instance at ln Z ¼35, the critical strain varies from 0.15 to 0.50, being below 0.10 for the HC steel. Interestingly, it seems that there are two trends: three sets of data (LC and Refs. 6 and 8) show a much higher dependence on Z than the other three (HC and Refs. 13 and 22). Anyhow, it is seen that the critical strains for the HC steel are drastically lower than those for the LC steel, and the difference is more noticeable at high Z values (i.e. high strain rates as seen in Fig. 1b). However, because the localization of strain along the segregation bands and the possible presence of ferrite in the HC steel at high temperatures seem to contribute to the softening of the steel, the critical stresses and strains of this steel are not necessarily related to the onset of DRX. In addition to the localization of strain and the possible presence of ferrite, another reason for the difference in the critical strains of these HC and LC steels might be connected to the higher density of carbon-manganese dipoles in the steels with higher carbon contents. It is well known that these dipoles effectively increase the work hardening rate of a high-manganese steel by pinning its dislocations [27]. Thus, the C–Mn dipoles could lead to a larger work hardening rate upon deformation of steels with higher carbon contents and thereby enhance the driving force for the occurrence of softening. However, this cannot be the only reason for the lower critical strains of the HC steel, since it should work as well in wrought steels, but the flow curves which are obtained for a wrought 22Mn–0.6C TWIP steel show clearly larger critical strains [14]. From Fig. 1, it is seen that the flow stress levels at typical rolling pass strains of 0.2–0.3 [6] are about 100–270 MPa depending on the temperature (strain rate 1/s), i.e. quite high and somewhat higher for the HC steel compared to the LC steel. However, the critical strains for the HC steel are small enough for pronounced softening to be obtained at strains of 0.3 and beyond. 4.3. Grain refinement by DRX The dynamically recrystallized average grain size is not dependent on the initial grain size if the imposed strain is sufficient [28]. Therefore, despite the very large initial grain sizes of the investigated
as-cast steels, the refined grain sizes are expected to be comparable to those of similar wrought steels. However, it has been shown that the strain distribution in the hot-deformed as-cast steel is relatively heterogeneous [23]. The occurrence of DRX in specimens of the HC steel, which were deformed at different values of Z, was confirmed by the EBSD maps shown in Fig. 6(a–d). It is known that the DRX grain size is inversely related to Z [29]. In agreement with this, smaller DRX grains were obtained for this steel at higher strain rates and lower temperatures. However, at the same time, increasing Z led to greater inhomogeneity of the structure. It is worth noting that although the specimens corresponding to Fig. 6(a and d) have been tested at close values of Z, a much more uniform structure was obtained in the latter condition, i.e. at the higher temperature. Thus, it can be inferred that uniformity of the hot deformed structure, i.e. the extent of DRX, is much more influenced by temperature than strain rate. The hot deformed microstructure of the LC steel was different from that of the HC steel. High-temperature flow curves (Fig. 1) showed that a distinct peak stress could be hardly seen for LC steel up to the true strain of 0.8. In consistence with this, only partial DRX could be found in some of the hot-deformed specimens of this steel. Fernandez et al. [29] have shown that the initial austenite grain size affects the kinetics of DRX. Thus, the extremely coarse grain structure of the LC steel in the as-cast condition ( 800 mm) can explain the slow kinetics of microstructural developments in this steel. However, in the HC steel, the DRX kinetics seems to be much faster. In addition to the effect of strain localization and possible presence of ferrite on softening, it has been suggested that carbon interstitial atoms cause lattice expansion, so that diffusion-controlled microstructural evolutions are more enhanced in a steel with a higher carbon content [30]. Due to the lower carbon content in the LC steel, the roomtemperature stacking fault energy (SFE) for this steel is lower than that of the HC steel [31]. As a result, a considerable amount of epsilon martensite is present in the room temperature microstructure of the LC steel (Figs. 2 and 7a). However, Hamada et al. [32] showed that the formation of epsilon martensite in a lowcarbon TWIP steel (Fe–0.14C–26Mn) is dependent on the grain size. The grain boundaries act as obstacles to the transformation of austenite to epsilon martensite so that below a certain grain size, this transformation may be totally suppressed [32,33]. As was expected, the DRX grain size after deformation at 0.01/s was larger than that obtained at the strain rate of 1/s, when the average grain size was about 8 μm. Martensite formation was suppressed only in the latter case (Fig. 7b).
5. Summary The hot deformation behavior of two as-cast coarse-grained TWIP steels with different levels of carbon (0.49% and 0.07%) has been studied by conducting hot compression tests at different temperatures and strain rates. The following conclusions can be drawn from the work: ▪ At lower strains, flow stress levels for the 0.49% C steel were higher than those of the 0.07% C steel. The reason is still unclear, but might be connected with higher carbon content in the former providing a larger number of Mn–C dipoles resulting in short range ordering. ▪ In the cast steel with 0.49% C, softening was very pronounced starting at very small strains and not reaching a steady state up to the maximum applied strain of 0.8, whereas in the cast steel with 0.07% C, the flow stress curves showed no clear peaks. ▪ The level of interdendritic micro-segregation was higher for the high-carbon steel and led to the formation of strain localization
A. Khosravifard et al. / Materials Science & Engineering A 582 (2013) 15–21
▪
▪
▪
▪
bands. Formation of ferrite at high temperatures in the segregation bands with lower Mn and higher Al and Si might also be possible, though not predicted by Thermocalc software. The localization of strain and the possible presence of ferrite are suggested as the reasons for the early and pronounced softening of this steel. The activation energy of deformation of the high-carbon steel is only slightly higher than that of the low-carbon steel. However, both values are comparable to those reported for wrought TWIP steels. In the high-carbon steel deformed at various strain rates and temperatures, the occurrence of dynamic recrystallization (DRX) was proven by microstructural studies. However, only partial DRX could be observed in some of the low-carbon steel specimens. The large grain size of the as-cast structures was effectively refined by DRX in both steels. With increasing strain rate, the DRX grain size became smaller. In the low-carbon steel, the formation of epsilon-martensite during cooling was suppressed after the attainment of a sufficiently fine austenite grain size.
References [1] O. Bouaziz, S. Allain, Y. Estrin, Scr. Mater. 62 (2010) 713–715. [2] I. Gutierrez-Urrutia, S. Zaefferer, D. Raabe, Mater. Sci. Eng. A 527 (2010) 3552–3560. [3] L. Chen, Y. Zhao, X. Qin, Acta Metall. Sin. (Engl. Lett.) 26 (2013) 1–15. [4] D. Li, Y. Feng, Z. Yin, F. Shangguan, K. Wang, Q. Liu, F. Hu, Mater. Des. 34 (2012) 713–718. [5] L.A. Dobrzanski, A. Grajcar, W. Borek, J. Achiev. Mater. Manuf. Eng. 29 (2008) 139–142. [6] L.A. Dobrzanski, A. Grajcar, W. Borek, Arch. Mater. Sci. Eng. 37 (2009) 69–76. [7] L.A. Dobrzanski, W. Borek, J. Achiev. Mater. Manuf. Eng. 46 (2011) 71–78.
21
[8] A.S. Hamada, L.P. Karjalainen, M.C. Somani, Mater. Sci. Eng. A 467 (2007) 114–124. [9] A.S. Hamada, M.C. Somani, L.P. Karjalainen, ISIJ Int. 47 (2007) 907–912. [10] J. Hajkazemi, A. Zarei-Hanzaki, M. Sabet, S. Khoddam, Mater. Sci. Eng. A 530 (2011) 233–238. [11] D. Li, Y. Feng, Z. Yin, F. Shangguan, K. Wang, Q. Liu, F. Hu, Mater. Sci. Eng. A 528 (2011) 8084–8089. [12] A. Marandi, A. Zarei-Hanzaki, N. Haghdadi, M. Eskandari, Mater. Sci. Eng. A 554 (2012) 72–78. [13] A.S. Hamada, L.P. Karjalainen, M.C. Somani, Can. Metall. Quart. 46 (2007) 47–56. [14] A. Ferraiuolo, A. Smith, J.G. Sevillano, F. de las Cuevas, G. Pratolongo, H. Gouveia, M. Mendes Rodrigues, P. Karjalainen, Metallurgical design of high strength austenitic Fe–C–Mn steels with excellent formability, Contract RFSRCT-00030, Research Programme of the Research Fund for Coal and Steel, European Union (2005-2008). [15] N.D. Ryan, H.J. McQueen, J. Mater. Process. Technol. 21 (1990) 177–199. [16] N. Fujita, T. Narushima, Y. Iguchi, C. Ouchi, ISIJ Int. 43 (2003) 1063–1072. [17] S. Hotta, T. Murakami, T. Narushima, Y. Iguchi, C. Ouchi, ISIJ Int. 45 (2005) 338–346. [18] G.K. Mandal, N. Stanford, P. Hodgson, J.H. Beynon, Mater. Sci. Eng. A 556 (2012) 685–695. [19] Y. Han, G. Liu, D. Zou, R. Liu, G. Qia, Mater. Sci. Eng. A 565 (2013) 342–350. [20] A. Najafizadeh, J.J. Jonas, ISIJ Int. 46 (2006) 1679–1684. [21] X. Xiao, G.Q. Liu, B.F. Hu, X. Zheng, L.N. Wang, S.J. Chen, A. Ullah, Comput. Mater. Sci. 62 (2012) 227–234. [22] F. Reyes-Calderon, I. Mejia, J.M. Cabrera, Mater. Sci. Eng. A 562 (2013) 46–52. [23] A. Iza-Mendia, A. Pinol-Juez, J.J. Urcola, I. Gutierrez, Metall. Mater. Trans. A 29 (1998) 2975–2986. [24] A. Saeed-Akbari, L. Mosecker, A. Schwedt, W. Bleck, Metall. Mater. Trans. A 43 (2012) 1688–1704. [25] G.I. Sil’man, Met. Sci. Heat Treat. 47 (2005) 397–401. [26] M. Wolf, W. Kurz, Metall. Trans. B 12 (1981) 85–93. [27] B. Hutchinson, N. Ridley, Scr. Mater. 55 (2006) 299–302. [28] Z. Jin, Z. Cui, J. Mater. Sci. Technol. 26 (2010) 1063–1070. [29] A.I. Fernandez, P. Uranga, B. Lopez, J.M. Rodriguez-Ibabe, Mater. Sci. Eng. A 361 (2003) 367–376. [30] S. Serajzadeh, A. Karimi Taheri, Mech. Mater. 35 (2003) 653–660. [31] A. Saeed-Akbari, J. Imlau, U. Prahl, W. Bleck, Metall. Mater. Trans. A 40 (2009) 3076–3090. [32] A.S. Hamada, P. Sahu, S. Ghosh Chowdhury, L.P. Karjalainen, J. Levoska, T. Oittinen, Metall. Mater. Trans. A 39 (2008) 462–465. [33] S. Takaki, H. Nakatsu, Y. Tokunaga, Mater. Trans. 34 (1993) 489–495.