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The ductility behavior of a high-Mn twinning induced plasticity steel during cold-to-hot deformation
Materials Science & Engineering A 561 (2013) 411–418
Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
The ductility behavior of a high-Mn twinning induced plasticity steel during cold-to-hot deformation A. Haft Baradaran, A. Zarei-Hanzaki n, H.R. Abedi, S.M. Fatemi-Varzaneh, A. Imandoust School of Metallurgy and Materials Engineering, College of Engineering, University of Tehran, Tehran, Iran
a r t i c l e i n f o
abstract
Article history: Received 21 July 2012 Received in revised form 7 October 2012 Accepted 15 October 2012 Available online 22 October 2012
The ductility behavior of a high-Mn TWIP steel (containing 30% wt Mn) has been studied using tensile testing method in a wide range of temperature (100–1000 1C) under the strain rate of 10 4 s 1. The hot compression characteristics of the experimental alloy are considered to assist in explaining the related deformation mechanisms. The results indicate that the ductility decreases with temperature; however, two regions of moderately improved ductility have also been realized. The former is attributed to the reduction of twinning activity by increasing the temperature. On the other hand, the activation of dynamic recovery at 400 1C causes the ductility to increase. The fracture surface observations denote the occurrence of grain boundary sliding at temperatures above 500 1C. As the dominant restoration process alters to partial dynamic recrystallization at 800 1C, the tensile ductility continues to decrease. By increasing the temperature to 1000 1C, the fraction of dynamically recrystallized grains is significantly increased and the ductility is improved. & 2012 Elsevier B.V. All rights reserved.
Keywords: Mechanical characterization Steel Thermomechanical processing Fracture Twinning Recrystallization
1. Introduction The twinning induced plasticity (TWIP) steels, with an excellent combination of strength and ductility, have been attracting great attention as a new group of materials in automotive industries [1,2]. The compromise between high strength and desired ductility in these particular austenitic steels results from their high potential for strain hardening. The dynamic formation of mechanical twins in TWIP steels reduces the mean free path of dislocations thereby increasing the rate of strain hardening [3,4]. As a matter of fact, the austenite is to be stabilized in the steel microstructure to render sufficient twin formation inside the austenite grains as deformation proceeds. The high level of manganese in these steels not only stabilizes the austenite but also decreases the related stacking fault energy (SFE) [5,6]. The high addition of Mn, however, may often cause problems such as cracking during hot rolling, and easy surface oxidation of rolled steel products [7,8]. Therefore, if the high-Mn TWIP steels should be well industrialized their manufacturing processes needto be optimized. Accordingly, a thorough investigation of their ductility behavior is highly necessary to forecast any probable ductility trough thereby providing a proper ductility map. Recently, Hamada et al. [9] studied the hot ductility behavior of four high-Mn TWIP steels (with 20% wt Mn) in the temperature
n
Corresponding author. Tel.: þ98 21 61114167; fax: þ 98 21 88006076. E-mail address: [email protected] (A. Zarei-Hanzaki).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.10.051
range of 700–1200 1C under the high strain rate of 1 s 1. They have reported a ductility trough at 700–900 1C, which was attributed to the grain boundary sliding (GBS) phenomena. Moreover, it was shown that the addition of Al (to 3% wt) would lead to the formation of ferrite grains at the austenite grain boundaries at higher temperatures and this in turn would render a detrimental effect to the hot ductility behavior. The other systematic study was conducted by Mintz et al. [10], where adding S and N to a 22% Mn TWIP steel made ductility worse in the temperature range of 700–900 1C due to the precipitation of AlN, VN and MnS. In general, the hot ductility may be subjected to a dramatic decrease due to the various mechanisms which would be activated at higher temperatures such as grain boundary sliding followed by cracking (W and R type cracks formation), strain concentration and flow localization in a more ductile phase, incipient melting reactions and diffusional transformations. The occurrence of restoration processes was also speculated to improve the hot ductility [11,12]. In addition, the ductility behavior of TWIP steels is strongly related to their stacking fault energy [13–15], which is the most crucial parameter determining whether twinning, martensite transformation or dislocation glide alone would occur during deformation of the material. Considering the researches involved, there has been a lack of organized study dealing with the hot ductility behavior of TWIP steels especially with higher Mn content. In the present work, the ductility evolution of TWIP steel containing 30% Mn (different to those employed by Hamada and Karjalainen [9] and Mintz [10]) has been studied using the conventional tensile testing method in
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3. Results and discussion 3.1. Compression behavior To assist determining the dominant mechanisms operating under the specified deformation condition, the warm-to-hot compression characteristics of the alloy were considered. The obtained true stress–true strain curves are shown in Fig. 2.
600 100°C 200°C 300°C
Fig. 1. The as hot forged–annealed experimental TWIP steel microstructure, which composes of equiaxed grains with annealing twins.
True Stress (MPa)
500 400 300 200 100
a wide range of temperature (100–1000 1C), where lower temperatures were included to obtain a better understanding of the possible SFE effects. An attempt has been also made to characterize the fracture surfaces and to justify them through the associated deformation mechanisms.
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The experimental steel was supplied in as-cast condition with the chemical composition of 29.7 Mn–2.5Al–0.6Si–0.17C, wt%. In order to eliminate the dendritic structure and to achieve a uniform microstructure, the as-received material was forged at 1150 1C. This was followed by subsequent annealing at 1000 1C for 90 min to remove any segregation of alloying elements in particular that of Mn. The initial microstructure of the experimental alloy is shown in Fig. 1. As is seen the microstructure is fully composed of an austenitic structure characterized by annealing twins. The initial grain size, measured by linear-interception method, is 80 mm. The tensile tests were conducted according to ASTM E8M standard [16] using cylindrical specimens with a reduced section diameter of 6mm and a gauge length of 30 mm. The isothermal tensile tests were carried out in the temperature range of 100–1000 1C with intervals of 100 1C under the strain rate of 10 4 s 1. The hot compression tests were also carried out in the aforementioned deformation conditions to study the involved micro-mechanism. The hot compression testing specimens were machined according to ASTM E209 standard [17] using cylindrical specimens in the sizes of F8 mm H12 mm. The specimens were first heated up to the deformation temperature and held isothermally for 5 min prior to straining. The tension and compression tests were carried out using an Instron-4208 universal testing machine, equipped with a contact extensometer and resistance furnace. All thermo-mechanical cycles were ceased by quenching the specimens in water just after straining. The elongation-tofailure was measured from the gauge length of the fractured specimens. To investigate the final microstructures, the specimens were sectioned along deformation axis, mounted using cold curing resin, ground and polished step by step up to the final polishing by 0.05 mm Al2O3 powders. The related fracture surfaces were also examined using scanning electron microscopy (SEM) to clarify the ductility behavior.
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True Strain Fig. 2. Typical true stress–true strain curves of the experimental alloy obtained by hot compression tests.
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larger imposed strains. Recovery generally involves a partial restoration because the dislocation structure is not completely removed, but reaches a metastable state. In contrast, during dynamic recrystallization new dislocation-free grains are formed within the deformed or recovered structure. These then grow and consume the old grains, resulting in a new grain structure with a low dislocation density.
3.2. Tensile deformation behavior Fig. 3 shows the different types of true stress–strain curves obtained from tensile tests in the temperature range of 100–1000 1C. Different deformation mechanisms may be realized from various tensile behaviors. The corresponding flow curves in the range of 700–1000 1C (Type A) include a short work hardening region up to the ultimate tensile strength (UTS) followed by a long post-UTS region. The usual flow softening in tensile curve is invariably
80 I
Elongation to Fracture (%)
This figure reveals three characteristic flow curves. The first category (in the range of 100–300 1C) includes a continuous work hardening region, which may imply the prevalence of twinning induced plasticity effect at these relatively medium temperatures. The twinning causes a high value of instantaneous hardening rate (n value). This is commonly attributed to the reduction of the dislocation mean free path by increasing the fraction of deformation twins as strong obstacles to dislocation glide [18]. The second category of the flow curves, achieved at 400–700 1C, shows a steady flow stress plateau. This appears to be originating from the balance of dynamic recovery (DRV) and work hardening. In fact, the dislocation activity and DRV are markedly enhanced above 400 1C, so that the rates of work hardening and recovery reach a dynamic equilibrium. The third category (in the temperature range of 800–1000 1C) exhibits typical dynamic recrystallization (DRX) behavior with a single peak stress followed by a gradual fall toward a steady-state stress [19]. As the dislocation density is directly related to the flow stress, some indication of the change in dislocation content with strain may often be inferred from the stress–strain behavior. Accordingly the dislocation density increases progressively during deformation. The increase in dislocation density is due to the continued trapping of newly created mobile dislocations by existing dislocations and their incorporation into the various microstructural features (twins and grains boundaries) that are characteristic of the deformed state. This would result in a continuous work hardening region in Type A deformation regime. In the other deformation conditions (Type B and Type C deformation regimes), which are usually of importance for the occurrence of restoration processes, the dislocation density decreases in a way to achieve a balance between dislocation generation and annihilation at the
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Dislocation Slip
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Fig. 4. Variation of elongation-to-fracture as a function of test temperature for the experimental TWIP steel.
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Table 1 The effect of SFE values on the dominant deformation mechanism in TWIP steels [1].
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True Strain Fig. 3. Different types of flow curves obtained by hot tensile tests.
Fig. 5. Effect of deformation temperature on SFE values for the experimental TWIP steel.
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associated with geometric instability related to the necking; this generally results in a short post-UTS region. However, the lower rate of stress drop beyond the necking (i.e. lower rates of work softening in post-UTS regions) of type A curves may be described relying on the occurrence of DRX in this range of temperature (800–1000 1C, Fig. 2). In fact the cavities and discontinuities, which are being formed in the microstructure, are isolated from the grain boundaries through grain boundary migration taking place during DRX. Consequently, the growth and coalescence of these cavities is not readily achieved away from grain boundaries so the strain to fracture
beyond the necking is increased [20,21]. In fact it is anticipated that the occurrence of DRX assists in obtaining an advanced flow localization, which in turn may postpone the necking. The second type of tensile flow behavior (Type B) exhibits a shorter post-UTS strain portion with respect to that of type A. The type B curves, which are achieved at 100–600 1C, reveal a very long pre-UTS region with higher rate of work hardening where a large amount of deformation may be driven through the mechanical twinning. This was followed by a short range of work softening region after UTS, ultimately leading to fracture.
3.3. Ductility behavior
Serrated boundaries
The variation of elongation-to-fracture as a function of test temperature is shown in Fig. 4. The curves indicate that there is a general trend of decreasing ductility with temperature. But, as is seen there are two regions (regions II and IV, denoted in Fig. 4) where the general trend changes and the ductility increases. The details of ductility variation at different deformation conditions have been described point by point and are as follows.
3.4. Ductility drop in region I (25–300 1C)
Fig. 6. Optical microstructure of the tensile specimens elongated to fracture at 400 1C.
The observed ductility drop below 300 1C is justified relying on the inhibition of twinning plasticity during deformation. As is well demonstrated one variant of the {111} o1124 twin system is activated within grains at the early stages of plastic deformation, followed by the nucleation of twins of different variants of the
Fig. 7. Fracture surface of the tensile specimens elongated to fracture at: (a) 200 1C, (b) 300 1C and (c) 400 1C.
A. Haft Baradaran et al. / Materials Science & Engineering A 561 (2013) 411–418
{111} o1124 twin system between the boundaries of the first set of twins [22]. Also, interaction of {111} o1104 slip and {111} o1124 twinning has been directly observed by previous researchers [3], by which a favorable balance of strength and ductility is attained. Accordingly the appearance of a ductility drop through this temperature range could be attributed to the change of dominant deformation mechanism from mechanical twinning to dislocation slip. The latter results from an increase in the alloy stacking fault energy by increasing the temperature. The stacking fault energy is known to be the main factor in driving the different plasticity mechanisms (twinning, martensite transformation or dislocation glide) during deformation of TWIP steels (see Table 1). As is well established the plasticity is governed by SFE values in the following disciplines: (i) the slipping and martensitic transformation (r18 mJ/m2), (ii) the mechanical twinning and partial dislocation slipping (18–60 mJ/m2), and (iii) the perfect dislocations slipping (60–85 mJ/m2). To calculate the SFE values for the current testing temperatures, the Grassel model [1] was employed and the results are given in Fig. 5. As is expected the magnitude of SFE is increased by the deformation temperature. It is interesting to note that the slope obviously changes as the temperature increases to 300 1C. Furthermore, the related SFE value corresponds to the boundary magnitude beyond which the dominant deformation mechanism switches to partial dislocation glide. Thus, with increasing temperature in the region I, the share of twinning in deformation plasticity vanishes. Accordingly the work hardenability of the experimental steel is exhausted and the overall ductility is reduced.
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3.4.1. Ductility enhancement in region II (300–500 1C) The moderate ductility improvement at 400 1C is attributed to the occurrence of an enhanced dynamic recovery (DRV). On going back to the compression results (Fig. 2) it is seen that as the temperature increased to 400 1C, the occurrence of DRV could successfully restore the deforming microstructure. Indeed, at the current applied low strain rates, the ductility stops decreasing due to the recovery capacity of steel to modify high-energy zones (like crack tips) into zones of lower energy by dislocation sliding, diffusion, and stress relieving mechanisms [23]. The development of serrated boundaries which are found in the microstructure after tensile deformation at 400 1C (Fig. 6) imply the predominance of DRV. The fracture surfaces of the specimens tested at temperatures of 200, 300 and 400 1C are shown in Fig. 7(a–c). The ductile dimples with different sizes and depths are seen without any cleavage facet on the presented fracture surfaces. This indicates a ductile fracture mode. The contribution of dimpling at 200 and 400 1C is much lower and the dimples are shallower than that of 300 1C. As is expected, the observed changes in surface morphology coincide with the ductility drop in the ductility–temperature curve.
3.4.2. Ductility drop in region III (500–900 1C) The fractographical observations are taken into consideration to explore the origin of ductility decline recorded in region III. As is observed in Fig. 8, a detailed view of the fracture surface at 500 1C demonstrates the appearance of small voids. This suggests that the relative resistance of the austenite grain interior to
c
b
Step
Void
Fig. 8. Fracture surface of the tensile specimen elongated to fracture at 500 1C.
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ductility continues to decrease. Thus, the occurrence of GBS associated with DRV may promote the nucleation and growth of R-type cavities [24]. The void formation at this temperature starts from any small decohesion at the grain boundary which has been developed by plane sliding. With further deformation, new decohesions are generated alongside the original ones, allowing the void growth as well as further plane sliding [24]. This explanation is applicable for any temperature in the range of 500–700 1C; however, the size of the steps and the magnitude of the sliding are dependent on the test temperature. Those steps can be readily observed in Fig. 8b. This may explain the reason for the decrease in hot ductility with increase in temperature from 500 1C to 700 1C. As was shown in Fig. 2, the material undergoes restoration via dynamic recrystallization as the deformation temperature increases to 800 1C. The sub-fracture surface microstructures of the specimens elongated to fracture at 800 and 900 1C (Fig. 9) include a low fraction of DRXed grain (10% and 15%, respectively). At these deformation conditions, GBS does not have a great contribution to total deformation because it may not be well accommodated by diffusional flow processes and dislocation slip. Thus, any stress concentration at certain sites (triple junction of grains and even ledges on grain) may cause the development of a w-type crack thereby reducing the elongation-to-fracture [23]. The fracture surfaces of the specimens tested at temperatures of 800 and 900 1C are shown in Fig. 10(a–d). As is seen, at 800 and 900 1C the character of the fracture is not ductile. The observed facet grain surfaces with shallow dimples are characteristics of failure due to GBS, i.e., intergranular decohesion. Similarly, in C–Mn–Nb steel, Carpenter et al. [25] found that the fracture surfaces obtained from hot tensile specimens tested in the single phase austenite region exhibited intergranular decohesion, displaying facet, due to the GBS. 3.4.3. Ductility enhancement in region IV (900–1000 1C) As the temperature increases to 1000 1C, the volume fraction of DRXed grains in the microstructure is significantly increased to about 80% (Fig. 9c). In a recrystallized microstructure with a large fraction of high angle grain boundaries, GBS can be more readily operated during deformation. Moreover, considering the relatively high diffusion rate (the lattice and grain boundary diffusion) in this condition in TWIP steels [11], the diffusion controlled GBS may lead to a higher formability at current elevated temperature and low strain rate. Thus, the high ductility observed at 1000 1C is attributed to the thermal activation of two simultaneous phenomena, namely GBS and the extended DRX [9]. These two mechanisms move the grain boundaries away from micro-cracks, keeping them isolated and preventing in turn their coalescence [23]. As is indicated in Fig. 10(e–f), the fracture mode at 1000 1C is ductile rupture, and refined grain is observed at the fracture surface. Consistently, Crowther and Mintz [11] believed that where the temperature is high enough, the DRX occurs and the cracks become displaceable from the prior boundaries as a result of grain boundary migration [11]. A similar role of DRX was discussed by Hamada and Karjalainen [9] to explain the ductility increase in a 22%Mn TWIP steel. Fig. 9. Optical microstructure (sub-fracture surface) of the tensile specimens elongated to fracture at: (a) 800 1C, (b) 900 1C and (c) 1000 1C.
plastic flow decreases and consequently the austenite grain boundary sliding (GBS) leads to an early void formation [24]. On the other hand, the DRV is operative during deformation of the experimental alloy at this temperature. However, the DRV may not effectively restore the stress concentrations anymore and the
4. Conclusion The cold-to-hot ductility of a TWIP steel containing 30% Mn has been studied in a wide range of temperature (100–1000 1C). The following conclusions have been drawn: 1. The overall ductility of 30%Mn decreases as the temperature increases.
A. Haft Baradaran et al. / Materials Science & Engineering A 561 (2013) 411–418
417
Wedge Crack
Wedge Crack
Refined grains
Fig. 10. Fracture surface of the tensile specimens elongated to fracture at: (a, b) 800 1C, (c, d) 900 1C and (e, f) 1000 1C.
2. The experimental alloy shows a ductility drop below 300 1C and this is related to the twinning inhibition through an increase in SFE value. 3. The moderate ductility improvement at 400 1C is attributed to the activation of DRV process. 4. As the dominant restoration process alters to partial DRX at 800 1C, the tensile ductility continues to decrease. 5. By increasing the temperature to 1000 1C, the dynamically recrystallized new grains are well developed in the microstructure thereby improving the ductility via isolation of the micro-cracks.
Acknowledgment Authors would like to acknowledge the financial support of University of Tehran for this research under the grant number of 8107001/1/01.
References ¨ [1] O. Grassel, L. Kruger, G. Frommeyer, L.W. Meyer, Int. J. Plast. 16 (2000) 1391–1409. [2] D. Hua, T. Zheng-You, L. Wei, J. Iron. Steel Res. Int. 13 (2006) 66–70. [3] S. Vercammen, B. Blanpain, B.C. De Cooman, Acta Mater. 52 (2004) 2005–2012. ¨ P. Neumann, ISIJ Int. 43 (2003) 438–446. [4] G. Frommeyer, U. Brux, [5] S. Allain, J.P. Chateau, O. Bouaziz, S. Migot, N. Guelton, Mater. Sci. Eng. A 387–389 (2004) 158–162. [6] S. Curtze, V.T. Kuokkala, Acta Mater. 58 (2010) 5129–5141. [7] U. Brux, G. Frommeyer, O. Grassel, L.W. Meyer, A. Weise, Steel Res. 73 (2002) 294–298. [8] M. Iker, D. Gaude-Fugarolas, P.J. Jacques, F. Delannay, Adv. Mater. Res. 15–17 (2007) 852–857. [9] A.S. Hamada, L.P. Karjalainen, Mater. Sci. Eng. A 528 (2011) 1819–1827. [10] S.E. Kang, A. Tuling, J.R. Banerjee, W.D. Gunawardana, B. Mintz, Mater. Sci. Technol. 27 (2011) 95–100. [11] B. Mintz, D.N. Crowther, Int. Mater. Rev. 55 (2010) 168–196. [12] B. Mintz, R. Abushosha, Ironmaking Steelmaking 20 (1993) 445–452. [13] A. Dumay, J.P. Chateau, S. Allain, S. Migot, O. Bouaziz, Mater. Sci. Eng. A 483–484 (2008) 184–187. [14] Y.N. Petrov, Z. Metallkunde, Mater. Res. Adv. Tech. 94 (2003) 1012–1016. [15] J. Kliber, K. Drozd, I. Mamuzic, Metal 5 (2009) 19–21.
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[16] Standard test methods for tension testing of metallic materials, Annual Book of ASTM Standards, vol. 03.01, ASTM E8, 2010. [17] Standard practice for compression tests of metallic materials at elevated temperatures with conventional or rapid heating rates and strain rates, Annual Book of ASTM Standards, vol. 03.01, ASTM E209, 2010. [18] A.S. Hamada, L.P. Karjalainen, M.C. Somani, Mater. Sci. Eng. A 464 (2007) 114–124. [19] M. Sabet, A. Zarei-Hanzaki, Sh. Khoddam, Eng. Mater. Technol. 131 (2009) 1–5.
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Contents lists available at SciVerse ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
The ductility behavior of a high-Mn twinning induced plasticity steel during cold-to-hot deformation A. Haft Baradaran, A. Zarei-Hanzaki n, H.R. Abedi, S.M. Fatemi-Varzaneh, A. Imandoust School of Metallurgy and Materials Engineering, College of Engineering, University of Tehran, Tehran, Iran
a r t i c l e i n f o
abstract
Article history: Received 21 July 2012 Received in revised form 7 October 2012 Accepted 15 October 2012 Available online 22 October 2012
The ductility behavior of a high-Mn TWIP steel (containing 30% wt Mn) has been studied using tensile testing method in a wide range of temperature (100–1000 1C) under the strain rate of 10 4 s 1. The hot compression characteristics of the experimental alloy are considered to assist in explaining the related deformation mechanisms. The results indicate that the ductility decreases with temperature; however, two regions of moderately improved ductility have also been realized. The former is attributed to the reduction of twinning activity by increasing the temperature. On the other hand, the activation of dynamic recovery at 400 1C causes the ductility to increase. The fracture surface observations denote the occurrence of grain boundary sliding at temperatures above 500 1C. As the dominant restoration process alters to partial dynamic recrystallization at 800 1C, the tensile ductility continues to decrease. By increasing the temperature to 1000 1C, the fraction of dynamically recrystallized grains is significantly increased and the ductility is improved. & 2012 Elsevier B.V. All rights reserved.
Keywords: Mechanical characterization Steel Thermomechanical processing Fracture Twinning Recrystallization
1. Introduction The twinning induced plasticity (TWIP) steels, with an excellent combination of strength and ductility, have been attracting great attention as a new group of materials in automotive industries [1,2]. The compromise between high strength and desired ductility in these particular austenitic steels results from their high potential for strain hardening. The dynamic formation of mechanical twins in TWIP steels reduces the mean free path of dislocations thereby increasing the rate of strain hardening [3,4]. As a matter of fact, the austenite is to be stabilized in the steel microstructure to render sufficient twin formation inside the austenite grains as deformation proceeds. The high level of manganese in these steels not only stabilizes the austenite but also decreases the related stacking fault energy (SFE) [5,6]. The high addition of Mn, however, may often cause problems such as cracking during hot rolling, and easy surface oxidation of rolled steel products [7,8]. Therefore, if the high-Mn TWIP steels should be well industrialized their manufacturing processes needto be optimized. Accordingly, a thorough investigation of their ductility behavior is highly necessary to forecast any probable ductility trough thereby providing a proper ductility map. Recently, Hamada et al. [9] studied the hot ductility behavior of four high-Mn TWIP steels (with 20% wt Mn) in the temperature
n
Corresponding author. Tel.: þ98 21 61114167; fax: þ 98 21 88006076. E-mail address: [email protected] (A. Zarei-Hanzaki).
0921-5093/$ - see front matter & 2012 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2012.10.051
range of 700–1200 1C under the high strain rate of 1 s 1. They have reported a ductility trough at 700–900 1C, which was attributed to the grain boundary sliding (GBS) phenomena. Moreover, it was shown that the addition of Al (to 3% wt) would lead to the formation of ferrite grains at the austenite grain boundaries at higher temperatures and this in turn would render a detrimental effect to the hot ductility behavior. The other systematic study was conducted by Mintz et al. [10], where adding S and N to a 22% Mn TWIP steel made ductility worse in the temperature range of 700–900 1C due to the precipitation of AlN, VN and MnS. In general, the hot ductility may be subjected to a dramatic decrease due to the various mechanisms which would be activated at higher temperatures such as grain boundary sliding followed by cracking (W and R type cracks formation), strain concentration and flow localization in a more ductile phase, incipient melting reactions and diffusional transformations. The occurrence of restoration processes was also speculated to improve the hot ductility [11,12]. In addition, the ductility behavior of TWIP steels is strongly related to their stacking fault energy [13–15], which is the most crucial parameter determining whether twinning, martensite transformation or dislocation glide alone would occur during deformation of the material. Considering the researches involved, there has been a lack of organized study dealing with the hot ductility behavior of TWIP steels especially with higher Mn content. In the present work, the ductility evolution of TWIP steel containing 30% Mn (different to those employed by Hamada and Karjalainen [9] and Mintz [10]) has been studied using the conventional tensile testing method in
412
A. Haft Baradaran et al. / Materials Science & Engineering A 561 (2013) 411–418
3. Results and discussion 3.1. Compression behavior To assist determining the dominant mechanisms operating under the specified deformation condition, the warm-to-hot compression characteristics of the alloy were considered. The obtained true stress–true strain curves are shown in Fig. 2.
600 100°C 200°C 300°C
Fig. 1. The as hot forged–annealed experimental TWIP steel microstructure, which composes of equiaxed grains with annealing twins.
True Stress (MPa)
500 400 300 200 100
a wide range of temperature (100–1000 1C), where lower temperatures were included to obtain a better understanding of the possible SFE effects. An attempt has been also made to characterize the fracture surfaces and to justify them through the associated deformation mechanisms.
0 0.0
0.1
0.2
0.3
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0.7
0.8
True Strain
500
2. Experimental
400°C
True Stress (MPa)
400 600°C
300 700°C
200
800°C
100
0 0.0
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True Strain 200 180 160
True Stress (MPa)
The experimental steel was supplied in as-cast condition with the chemical composition of 29.7 Mn–2.5Al–0.6Si–0.17C, wt%. In order to eliminate the dendritic structure and to achieve a uniform microstructure, the as-received material was forged at 1150 1C. This was followed by subsequent annealing at 1000 1C for 90 min to remove any segregation of alloying elements in particular that of Mn. The initial microstructure of the experimental alloy is shown in Fig. 1. As is seen the microstructure is fully composed of an austenitic structure characterized by annealing twins. The initial grain size, measured by linear-interception method, is 80 mm. The tensile tests were conducted according to ASTM E8M standard [16] using cylindrical specimens with a reduced section diameter of 6mm and a gauge length of 30 mm. The isothermal tensile tests were carried out in the temperature range of 100–1000 1C with intervals of 100 1C under the strain rate of 10 4 s 1. The hot compression tests were also carried out in the aforementioned deformation conditions to study the involved micro-mechanism. The hot compression testing specimens were machined according to ASTM E209 standard [17] using cylindrical specimens in the sizes of F8 mm H12 mm. The specimens were first heated up to the deformation temperature and held isothermally for 5 min prior to straining. The tension and compression tests were carried out using an Instron-4208 universal testing machine, equipped with a contact extensometer and resistance furnace. All thermo-mechanical cycles were ceased by quenching the specimens in water just after straining. The elongation-tofailure was measured from the gauge length of the fractured specimens. To investigate the final microstructures, the specimens were sectioned along deformation axis, mounted using cold curing resin, ground and polished step by step up to the final polishing by 0.05 mm Al2O3 powders. The related fracture surfaces were also examined using scanning electron microscopy (SEM) to clarify the ductility behavior.
140
850°C
120
900°C
100
950°C
80
1000°C
60 40 20 0 0.0
0.1
0.2
0.3
0.4
0.5
0.6
0.7
0.8
True Strain Fig. 2. Typical true stress–true strain curves of the experimental alloy obtained by hot compression tests.
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larger imposed strains. Recovery generally involves a partial restoration because the dislocation structure is not completely removed, but reaches a metastable state. In contrast, during dynamic recrystallization new dislocation-free grains are formed within the deformed or recovered structure. These then grow and consume the old grains, resulting in a new grain structure with a low dislocation density.
3.2. Tensile deformation behavior Fig. 3 shows the different types of true stress–strain curves obtained from tensile tests in the temperature range of 100–1000 1C. Different deformation mechanisms may be realized from various tensile behaviors. The corresponding flow curves in the range of 700–1000 1C (Type A) include a short work hardening region up to the ultimate tensile strength (UTS) followed by a long post-UTS region. The usual flow softening in tensile curve is invariably
80 I
Elongation to Fracture (%)
This figure reveals three characteristic flow curves. The first category (in the range of 100–300 1C) includes a continuous work hardening region, which may imply the prevalence of twinning induced plasticity effect at these relatively medium temperatures. The twinning causes a high value of instantaneous hardening rate (n value). This is commonly attributed to the reduction of the dislocation mean free path by increasing the fraction of deformation twins as strong obstacles to dislocation glide [18]. The second category of the flow curves, achieved at 400–700 1C, shows a steady flow stress plateau. This appears to be originating from the balance of dynamic recovery (DRV) and work hardening. In fact, the dislocation activity and DRV are markedly enhanced above 400 1C, so that the rates of work hardening and recovery reach a dynamic equilibrium. The third category (in the temperature range of 800–1000 1C) exhibits typical dynamic recrystallization (DRX) behavior with a single peak stress followed by a gradual fall toward a steady-state stress [19]. As the dislocation density is directly related to the flow stress, some indication of the change in dislocation content with strain may often be inferred from the stress–strain behavior. Accordingly the dislocation density increases progressively during deformation. The increase in dislocation density is due to the continued trapping of newly created mobile dislocations by existing dislocations and their incorporation into the various microstructural features (twins and grains boundaries) that are characteristic of the deformed state. This would result in a continuous work hardening region in Type A deformation regime. In the other deformation conditions (Type B and Type C deformation regimes), which are usually of importance for the occurrence of restoration processes, the dislocation density decreases in a way to achieve a balance between dislocation generation and annihilation at the
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II
III
IV
60
40
20
Twinning
Dislocation Slip
0 0
100
200
300
400
500
600
700
800
900
1000 1100
Temperature (°C) 250
Fig. 4. Variation of elongation-to-fracture as a function of test temperature for the experimental TWIP steel.
True Stress(MPa)
200
150
Table 1 The effect of SFE values on the dominant deformation mechanism in TWIP steels [1].
800°C
100
900°C
SFE (mJ/m2)
18
18–38
38–60
60–85
1000°C
50
Deformation mechanism e-Martensite Twining Partial dislocation Slip 0 0
0.1
0.2
0.3 140
Stacking Fault Energy (mJ/m2)
True Strain
600 100°C
200°C
500
True Stress(MPa)
400°C 300°C
400
500°C
300
600°C 700°C
200
120 100
Dislocation Slip
80 60 40
Twinning 20
100
E-martensite 0 0
0 0
0.1
0.2
0.3
0.4
0.5
100
200
300
400
500
600
700
800
Temperature (°C)
True Strain Fig. 3. Different types of flow curves obtained by hot tensile tests.
Fig. 5. Effect of deformation temperature on SFE values for the experimental TWIP steel.
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associated with geometric instability related to the necking; this generally results in a short post-UTS region. However, the lower rate of stress drop beyond the necking (i.e. lower rates of work softening in post-UTS regions) of type A curves may be described relying on the occurrence of DRX in this range of temperature (800–1000 1C, Fig. 2). In fact the cavities and discontinuities, which are being formed in the microstructure, are isolated from the grain boundaries through grain boundary migration taking place during DRX. Consequently, the growth and coalescence of these cavities is not readily achieved away from grain boundaries so the strain to fracture
beyond the necking is increased [20,21]. In fact it is anticipated that the occurrence of DRX assists in obtaining an advanced flow localization, which in turn may postpone the necking. The second type of tensile flow behavior (Type B) exhibits a shorter post-UTS strain portion with respect to that of type A. The type B curves, which are achieved at 100–600 1C, reveal a very long pre-UTS region with higher rate of work hardening where a large amount of deformation may be driven through the mechanical twinning. This was followed by a short range of work softening region after UTS, ultimately leading to fracture.
3.3. Ductility behavior
Serrated boundaries
The variation of elongation-to-fracture as a function of test temperature is shown in Fig. 4. The curves indicate that there is a general trend of decreasing ductility with temperature. But, as is seen there are two regions (regions II and IV, denoted in Fig. 4) where the general trend changes and the ductility increases. The details of ductility variation at different deformation conditions have been described point by point and are as follows.
3.4. Ductility drop in region I (25–300 1C)
Fig. 6. Optical microstructure of the tensile specimens elongated to fracture at 400 1C.
The observed ductility drop below 300 1C is justified relying on the inhibition of twinning plasticity during deformation. As is well demonstrated one variant of the {111} o1124 twin system is activated within grains at the early stages of plastic deformation, followed by the nucleation of twins of different variants of the
Fig. 7. Fracture surface of the tensile specimens elongated to fracture at: (a) 200 1C, (b) 300 1C and (c) 400 1C.
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{111} o1124 twin system between the boundaries of the first set of twins [22]. Also, interaction of {111} o1104 slip and {111} o1124 twinning has been directly observed by previous researchers [3], by which a favorable balance of strength and ductility is attained. Accordingly the appearance of a ductility drop through this temperature range could be attributed to the change of dominant deformation mechanism from mechanical twinning to dislocation slip. The latter results from an increase in the alloy stacking fault energy by increasing the temperature. The stacking fault energy is known to be the main factor in driving the different plasticity mechanisms (twinning, martensite transformation or dislocation glide) during deformation of TWIP steels (see Table 1). As is well established the plasticity is governed by SFE values in the following disciplines: (i) the slipping and martensitic transformation (r18 mJ/m2), (ii) the mechanical twinning and partial dislocation slipping (18–60 mJ/m2), and (iii) the perfect dislocations slipping (60–85 mJ/m2). To calculate the SFE values for the current testing temperatures, the Grassel model [1] was employed and the results are given in Fig. 5. As is expected the magnitude of SFE is increased by the deformation temperature. It is interesting to note that the slope obviously changes as the temperature increases to 300 1C. Furthermore, the related SFE value corresponds to the boundary magnitude beyond which the dominant deformation mechanism switches to partial dislocation glide. Thus, with increasing temperature in the region I, the share of twinning in deformation plasticity vanishes. Accordingly the work hardenability of the experimental steel is exhausted and the overall ductility is reduced.
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3.4.1. Ductility enhancement in region II (300–500 1C) The moderate ductility improvement at 400 1C is attributed to the occurrence of an enhanced dynamic recovery (DRV). On going back to the compression results (Fig. 2) it is seen that as the temperature increased to 400 1C, the occurrence of DRV could successfully restore the deforming microstructure. Indeed, at the current applied low strain rates, the ductility stops decreasing due to the recovery capacity of steel to modify high-energy zones (like crack tips) into zones of lower energy by dislocation sliding, diffusion, and stress relieving mechanisms [23]. The development of serrated boundaries which are found in the microstructure after tensile deformation at 400 1C (Fig. 6) imply the predominance of DRV. The fracture surfaces of the specimens tested at temperatures of 200, 300 and 400 1C are shown in Fig. 7(a–c). The ductile dimples with different sizes and depths are seen without any cleavage facet on the presented fracture surfaces. This indicates a ductile fracture mode. The contribution of dimpling at 200 and 400 1C is much lower and the dimples are shallower than that of 300 1C. As is expected, the observed changes in surface morphology coincide with the ductility drop in the ductility–temperature curve.
3.4.2. Ductility drop in region III (500–900 1C) The fractographical observations are taken into consideration to explore the origin of ductility decline recorded in region III. As is observed in Fig. 8, a detailed view of the fracture surface at 500 1C demonstrates the appearance of small voids. This suggests that the relative resistance of the austenite grain interior to
c
b
Step
Void
Fig. 8. Fracture surface of the tensile specimen elongated to fracture at 500 1C.
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ductility continues to decrease. Thus, the occurrence of GBS associated with DRV may promote the nucleation and growth of R-type cavities [24]. The void formation at this temperature starts from any small decohesion at the grain boundary which has been developed by plane sliding. With further deformation, new decohesions are generated alongside the original ones, allowing the void growth as well as further plane sliding [24]. This explanation is applicable for any temperature in the range of 500–700 1C; however, the size of the steps and the magnitude of the sliding are dependent on the test temperature. Those steps can be readily observed in Fig. 8b. This may explain the reason for the decrease in hot ductility with increase in temperature from 500 1C to 700 1C. As was shown in Fig. 2, the material undergoes restoration via dynamic recrystallization as the deformation temperature increases to 800 1C. The sub-fracture surface microstructures of the specimens elongated to fracture at 800 and 900 1C (Fig. 9) include a low fraction of DRXed grain (10% and 15%, respectively). At these deformation conditions, GBS does not have a great contribution to total deformation because it may not be well accommodated by diffusional flow processes and dislocation slip. Thus, any stress concentration at certain sites (triple junction of grains and even ledges on grain) may cause the development of a w-type crack thereby reducing the elongation-to-fracture [23]. The fracture surfaces of the specimens tested at temperatures of 800 and 900 1C are shown in Fig. 10(a–d). As is seen, at 800 and 900 1C the character of the fracture is not ductile. The observed facet grain surfaces with shallow dimples are characteristics of failure due to GBS, i.e., intergranular decohesion. Similarly, in C–Mn–Nb steel, Carpenter et al. [25] found that the fracture surfaces obtained from hot tensile specimens tested in the single phase austenite region exhibited intergranular decohesion, displaying facet, due to the GBS. 3.4.3. Ductility enhancement in region IV (900–1000 1C) As the temperature increases to 1000 1C, the volume fraction of DRXed grains in the microstructure is significantly increased to about 80% (Fig. 9c). In a recrystallized microstructure with a large fraction of high angle grain boundaries, GBS can be more readily operated during deformation. Moreover, considering the relatively high diffusion rate (the lattice and grain boundary diffusion) in this condition in TWIP steels [11], the diffusion controlled GBS may lead to a higher formability at current elevated temperature and low strain rate. Thus, the high ductility observed at 1000 1C is attributed to the thermal activation of two simultaneous phenomena, namely GBS and the extended DRX [9]. These two mechanisms move the grain boundaries away from micro-cracks, keeping them isolated and preventing in turn their coalescence [23]. As is indicated in Fig. 10(e–f), the fracture mode at 1000 1C is ductile rupture, and refined grain is observed at the fracture surface. Consistently, Crowther and Mintz [11] believed that where the temperature is high enough, the DRX occurs and the cracks become displaceable from the prior boundaries as a result of grain boundary migration [11]. A similar role of DRX was discussed by Hamada and Karjalainen [9] to explain the ductility increase in a 22%Mn TWIP steel. Fig. 9. Optical microstructure (sub-fracture surface) of the tensile specimens elongated to fracture at: (a) 800 1C, (b) 900 1C and (c) 1000 1C.
plastic flow decreases and consequently the austenite grain boundary sliding (GBS) leads to an early void formation [24]. On the other hand, the DRV is operative during deformation of the experimental alloy at this temperature. However, the DRV may not effectively restore the stress concentrations anymore and the
4. Conclusion The cold-to-hot ductility of a TWIP steel containing 30% Mn has been studied in a wide range of temperature (100–1000 1C). The following conclusions have been drawn: 1. The overall ductility of 30%Mn decreases as the temperature increases.
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Wedge Crack
Wedge Crack
Refined grains
Fig. 10. Fracture surface of the tensile specimens elongated to fracture at: (a, b) 800 1C, (c, d) 900 1C and (e, f) 1000 1C.
2. The experimental alloy shows a ductility drop below 300 1C and this is related to the twinning inhibition through an increase in SFE value. 3. The moderate ductility improvement at 400 1C is attributed to the activation of DRV process. 4. As the dominant restoration process alters to partial DRX at 800 1C, the tensile ductility continues to decrease. 5. By increasing the temperature to 1000 1C, the dynamically recrystallized new grains are well developed in the microstructure thereby improving the ductility via isolation of the micro-cracks.
Acknowledgment Authors would like to acknowledge the financial support of University of Tehran for this research under the grant number of 8107001/1/01.
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