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Twinning and Tripping in 10% Mn steels
Materials Science & Engineering A 591 (2014) 90–96
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Twinning and Tripping in 10% Mn steels Huseyin Aydin a,n, In-Ho Jung a, Elhachmi Essadiqi b, Stephen Yue a a b
McGill University, Department of Mining and Materials Engineering, Montreal, QC, Canada Université Internationale de Rabat, UIR, Aerospace Engineering School, Rabat, Technopolis Shore Bypass Rabat-Salé, Morocco
art ic l e i nf o
a b s t r a c t
Article history: Received 14 August 2013 Received in revised form 22 October 2013 Accepted 29 October 2013 Available online 5 November 2013
In the present work, a medium Mn, Fe–Mn–C–Al–Si alloy was subjected to different heat treatment conditions and subsequent deformation to understand the effect of these processes on austenite. It was found that, after intercritical annealing, the microstructure was ferrite plus austenite duplex phase (FADP) regardless of cooling rate to room temperature. When cold rolled, the retained austenite of the FADP structures exhibited both twinning and strain induced transformation (SIT) to martensite. A detailed characterization of the co-existence of twinning and SITing after cold rolling is presented. & 2013 Elsevier B.V. All rights reserved.
Keywords: Twinning Strain induced transformation (SIT) Stacking fault energy (SFE) Retained austenite (RA)
1. Introduction Over the last few decades, growing demands for weight saving and safety requirement have motivated new concepts of automotive steels to achieve improved mechanical properties in comparison with the existing Advanced High Strength Steels (AHSS). Among various developments, medium Mn content steels are considered to be potential candidates to achieve the performance targets of the so-called third generation AHSS [1]. These steels are currently one of the materials being increasingly developed to harness the promise of second generation steels without high alloy costs and difficult processing issues associated with these latter steels. Indeed, several literature studies proved that the medium Mn steels containing 5–10 wt% Mn have enhanced mechanical properties compared to the first generation AHSS due to the occurrence of deformation induced martensitic transformation [2–5]. It is also stated that the partitioning of Mn and C in the austenite during intercritical annealing are the two main contributions for the austenite stability, and mechanical stabilization of the austenite does not contribute to the austenite stability due to the very low dislocation density of the austenite grains [3–5]. Hence, these steels typically achieve very high strengths, but they have very limited strain hardening and, as a consequence, lower elongations compared to second generation AHSS. This paper describes a novel ferrite plus retained austenite microstructure based on the above alloying concepts. The behavior
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of retained austenite during room temperature deformation of three heat treated variants was investigated, with a focus on twinning and strain induced transformation characteristics. 2. Materials and experimental methods 2.1. Compositions The steel used throughout this work was supplied by CANMETMTL (Hamilton, Ontario, Canada). Castings were done in an induction furnace. The composition (wt%) of the examined steel is shown in Table 1. This is one composition selected from a series of alloys which were designed on the basis of attaining a certain level of stacking fault energy (SFE) in the metastable austenite to promote twinning as opposed to transformation to martensite [2]. The thermodynamic modeling approach was adopted to calculate the magnitude of stacking fault energy and most of the data used in this study was taken from the literature and Fsteel database of FactSage (version 6.2) computational thermodynamic software [6]. 2.2. Processing The cast ingot was sectioned into plates having a thickness of 30 mm, and was homogenized for 3 h at 1200 1C, immediately hot rolled to 6 mm at temperatures between 1100 and 750 1C and then water quenched. The as-hot rolled plates were intercritically annealed at elevated temperatures and then either air (AC), furnace (FC) cooled or water quenched (WQ) prior to mechanical testing and microstructural characterization. Fig. 1 shows the detailed heat treatment process of the steel samples. The
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Table 1 The bulk chemical composition of Fe10Mn steel. Alloy
C (%)
Mn (%)
Si (%)
Al (%)
Mo (%)
Fe
Fe10Mn
0.2
10.02
3.17
3.19
0.06
Bal.
Fig. 2. The phase diagram of Fe10Mn steel alloy (α: ferrite, γ: austenite).
3. Results and discussions Fig. 1. Heat cycle of the intercritical annealing process.
3.1. Microstructure
annealing temperatures were chosen to observe the transformation behavior of austenite at different temperatures and cooling rates (i.e. temperature conditions at highest FCC and/or C content etc.). The main objectives of the heat treatments were firstly to maximize the amount of retained austenite and then, to control the stability of retained austenite. Finally, cold rolling was also conducted on the annealed samples to observe the deformation behavior of the austenite.
2.3. Characterization The microstructural examinations were initially carried out on a Nikon L150 optical microscope, Bruker D8 X-ray difractometer and Philips XL-30 field emission scanning electron microscope (FE-SEM). All samples were mechanically polished down to 0.05 μm with colloidal silica and then chemically etched using a solution of 2% Nital followed by 10% aqueous sodium metabisulfite (Na2S2O5). The etchants used in this study have been widely used for AHSS steels to differentiate ferrite, austenite and martensite. So it is indeed required to differentiate them before analyses [7]. Furthermore, since there is no carbon in the etching solutions it was ignored during the Electron Probe Micro-Analyze (EPMA). The volume percentage of the phases were calculated with the help of Clemex Captive image analyze software and an XRD intensity correlation method [8]. Transmission electron microscopy (TEM) studies were conducted on Philips CM200 transmission electron microscope (TEM). The thin foil preparation involved mechanical thinning and jet polishing at 50 V in a solution containing 340 ml butanol, 600 ml methanol, and 60 ml perchloric acid at 20 1C. The BAHR DIL 805 quench dilatometer was used to evaluate the transformation behavior and to confirm the martensite start (Ms) temperature of steel composition at fast cooling rates. Finally, mechanical tests were performed using an MTS hydraulic machine by using 100 kN load cell. Tensile test samples were machined according to ASTM E-8 sub-size standard with a strain rate of 0.1 mm/s.
According to the FactSage predicted phase diagram (Fig. 2), the equilibrium structure at room temperature of this composition has no austenite. However, this composition and the above heat treatments generated a microstructure of ferrite plus austenite duplex (FADP) microstructure. Fig. 3 shows the morphology of the FADP structure after the two different heat treatment conditions. The first generation AHSS, relied on bainite to help retain (metastable) austenite, thus leading to an equiaxed ferrite plus bainite structure, with the retained austenite largely adopting morphologies associated with the bainite. In this structure, no bainite is required to retain the austenite, and the structure is comprised of equiaxed austenite and ferrite grains. Since the predicted equilibrium microstructure at room temperature has no austenite, the austenite in this structure is assumed to be unstable. As shown in Fig. 3, the grain sizes of the ferrite and austenite are more or less comparable to each other, with the ferrite appearing to be a little coarser than the austenite grains. The microstructures also showed equiaxed austenitic grains with a limited number of annealing twins and very small amount of martensite in the austenitic phases, but apparently with more annealing twins after air cooling. One of the most interesting points of these microstructures is that there is no significant effect of cooling rate; the retained austenite levels are very similar (about 50 vol%) even after furnace cooling, indicating that the kinetics of the transformation to ferrite are relatively slow. Particularly striking is the absence of martensite after water quenching, indicating the martensite start (Ms) temperature is well below room temperature. In order to estimate the Ms of the austenite formed at the intercritical temperature, the retained austenite composition in the WQ condition (i.e., as-quenched) was measured by EPMAWDS from the microstructure in Fig. 4 and shown in Table 2. The partitioning of C and Mn to austenite can be seen in Table 2. Unfortunately, there is no empirical equation in the literature to predict Ms from the austenite composition that has been created for medium Mn steels. Moreover, all such empirical equations have been developed by experiments that begin with 100% austenite, whereas the transformation that is of interest here is from the two phase ferrite plus austenite condition. Nevertheless, the above composition was used to determine the Ms temperatures
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RA
RA
F F
F
RA
Fig. 3. The optical images of Fe10Mn after two different heat treatment conditions; (a) air cooled, (b) water quenched and (c) furnace cooled. RA and F are retained austenite and ferrite, respectively.
F
Spectrum RA
Fig. 4. The BSE image and WDS analysis of Fe10Mn after WQ.
Fig. 5. Dilatometer analysis of Fe10Mn at gas nitrogen cooling.
Table 2 The retained austenite compositions measured by EPMA.
very high cooling rates (Fig. 5) but no sign of any transformation was observed.
Alloy
C (%)
Mn (%)
Si (%)
Al (%)
Mo (%)
Fe
Fe10Mn
1.10
11.2
2.95
2.97
0.09
Bal.
according to the following equation [9]: M s ¼ 539 423C 30:4Mn 7:5Si þ 30Al
ð1Þ
The Ms temperature is predicted to be 199.8 1C, therefore the austenite is expected not to transform martensite during quenching to room temperature. The transformation behavior of austenite to martensite was also confirmed with dilatometer experiments at
3.2. Effect of cold rolling on the microstructure Having noted the above insignificant effect of cooling rate on the amount of austenite it might be argued that this austenite is stable and not unstable as predicted by FactSage. Indeed, in contrast to the FactSage predictions, the ThermoCalc studies of De-Cooman et al. show that the pseudo-binary Fe–C phase diagrams of 10 Mn% steels have single phase austenitic area at elevated temperatures and stable austenite can be seen at room temperature conditions [10].
H. Aydin et al. / Materials Science & Engineering A 591 (2014) 90–96
However, after cold rolling, all the microstructures exhibited significant reduction in retained austenite, as measured by XRD (Table 3). Since it can be assumed that this was due to strain induced transformation (SIT) to martensite, this indicates that the austenite is unstable (i.e. retained). In the case of the differences of the austenite prediction, there is no obvious kinetic issue that could explain transformation of austenite during room temperature deformation if it is not unstable. Furthermore, cold rolling not only transformed approximately half of the retained austenite to strain induced martensitic, but also generated twinning, as is clearly observed both optically (Fig. 6) and in the TEM examination (Fig. 7). The only other reference to the simultaneous observation of twinning and SITing in retained austenite that the authors have found is by Timokhina et al. [11] for a conventional TRIP steel microstructure (i.e. ferrite, bainite and retained austenite); however, this was in the as-hot rolled condition – no other papers have reported twinning and SITing of retained austenite after cold deformation. Some austenite grains exhibited a mixture of SIT and twinning, whilst other grains revealed either twinning only or SIT only. It is clearly of interest to understand exactly why the retained austenite in the FADP steel can undergo both twinning and SIT during deformation, whereas the retained austenite in TRIP steels undergoes only strain induced transformation and the retained austenite in TWIP steels undergoes only twinning. However, in all Table 3 The amount of retained austenite after cold rolling. Austenite volume fraction (%)
Air cooled Water quenched Furnace cooled
Heat treated
Cold rolled
Δtransformed
47 53 47.5
23 32 28.4
24 21 19.1
93
previous studies, the microstructures that undergo twinning are 100% austenite which is energetically stable, e.g., TWIP and Hadfield steels [12–15], so there is no chance for strain induced transformation to martensite. Therefore the relevant comparison of FADP steels is with TRIP steels rather than TWIP steels. In this regard, the retained austenite grains in FADP are much larger and the level of retained austenite is much higher. Moreover, there is only one morphological type of retained austenite in FADP, i.e. equiaxed, whereas there are at least three types of retained austenites in TRIP steels ((i) equiaxed small grains, (ii) interlath layers in the bainite and (iii) layers between bainite packets and equiaxed ferrite grains) [16,17]. It is well known that decreasing the retained austenite size in TRIP steels increases the resistance against SIT. It may be that size is also a factor in twinning. Equally important is to determine when twinning and SIT take place during deformation, i.e., the evolution of the microstructure with strain. In fact, there are almost certainly three deformation mechanisms operating: (i) slip, (ii) SIT, and (iii) twinning. In terms of modeling the events, if these mechanisms are treated purely as ways to relieve the applied stress, then the obvious approach is to use the concept of critical resolved shear stress via the Schmidt factor. This approach is well known for slip and twinning; the concept of critical stress is also used in SIT [18,19], but the Schmidt factor is not considered. Another way to consider the problem is by a classic Gibbs energy approach; this is commonly used to explain TRIP [17,20] but not for twinning. With this approach, twinned retained austenite and strain induced martensite could be considered as being two phases in equilibrium. Fig. 8 shows a schematic diagram of the effect of composition on the free energies of these two phases. If the diffusion of solutes is allowed, the two phase assemblage is delineated by the common tangent between these curves, which leads to the two co-existing phases having different compositions; this does not occur in this case because both SIT and twinning are diffusionless ‘transformations’, and the compositions
Twins
F Twins M
F
RA M
M RA
F
Twins Fig. 6. The optical images of Fe10Mn after cold rolling (a) air cooled, (b) water quenched and (c) furnace cooled samples. RA, M and F are retained austenite, martensite and ferrite, respectively.
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Fig. 7. Deformation twins and SIT in Fe10Mn (a) bright field image and (b) weak-beam dark field TEM image of only twins and (c) bright field image and (d) dark field image of both twins (oriented along [011] zone axis (FCC)) and SIT (oriented along [111] (BCT)).
Fig. 9. Engineering stress–strain curves.
Fig. 8. Gibbs energy–composition diagram for the austenite and strain induced martensite phases at a given temperature.
of twinned and SIT phases are therefore the same. Therefore, in order to satisfy the co-existence of twinned austenite (γtwin) and SIT phase (α1), the Gibbs energies of the two phases should be the same or, in practice, very similar for the existing composition of the retained austenite. This scenario maybe feasible as twinning and tripping are mechanisms at high (about 20%) and low (2%) Mn concentrations, respectively. EPMA analysis [2] has shown that the Mn in the retained austenite of FADP is 11.2%, which is more or less
Table 4 The mechanical properties of Fe10Mn after tensile tests. Fe10Mn
UTS (MPa)
% εtotal
Rp0.2 (MPa)
WQ þ Tensile AC þ Tensile FCþ Tensile
868 903 776
31 23 25
637 694 567
H. Aydin et al. / Materials Science & Engineering A 591 (2014) 90–96
95
Fig. 10. Effect of (a) retained austenite fraction on strain and (b) SIT transformation on stress.
exactly between the high and low Mn levels. However, it would be very fortuitous to have attained the exact compositional ‘fit’ in this experiment where the free energies of strain induced martensite and tinned retained austenite exactly coincide. Therefore, this ‘classical’ ΔG curve approach does not satisfactorily explain the coexistence of twinning and tripping. A ‘hybrid’ of the CRSS and free energy approaches may be required to satisfactorily analyze this simultaneous twinning and SITing phenomenon. Since it is accepted that there is a strong effect of Mn and C on the SFE and transformation kinetics during deformation, one other reason why both SIT and twinning are seen in the retained austenite may be due to an inhomogeneous carbon distribution within the retained austenite. This could lead an incomplete transformation of retained austenite and these regions of austenite tend to transform to martensite or twinning at lower strains [11]. 3.3. Mechanical properties Engineering stress–strain curves are displayed in Fig. 9 for the two annealed samples and Table 4 summarizes their mechanical properties. Firstly, the WQ condition leads the highest total elongation of more than 30%, which is significant enhancement for medium Mn steel compositions, and it was balanced with a relatively high UTS value of 868 MPa. Annealing with AC led to slightly higher UTS more than 900 MPa but lower ductility values of 23%. In case of FC the total elongation was compatible with AC but at a much lower strength value because of the softer and coarser grain morphology. As shown in Fig. 10, increasing amount of transformation increases strength, which is basically due to increasing strain induced martensite. However, increasing transformation does not lead to increasing ductility; this demonstrates that increasing SIT increases ductility only if it postpones strain instability. In this case, it appears that SIT may take place well before strain instability, thus contributing little to ductility. In fact, increasing ductility correlates better with the increasing retained austenite, which might be due to FCC being fundamentally more ductile than ferrite or martensite. It is also interesting to note that the ductility of the asquenched specimens is slightly higher than the air cooled ones for a given RA vol%. This may be related to the fact that in the WQ specimen, the intercritical austenite was formed at higher temperature, thus leading to a lower carbon content [21], as confirmed by detailed EPMA studies in Reference [2]. This lower C content might make retained austenite softer, thereby increasing the ductility. The effect of twinning has not been considered but it is expected enhance the mechanical properties. Unfortunately there
is not enough information to quantify the separate contributions of twinning and SIT in the microstructure.
4. Summary i. Fe10Mn generated a ferrite and retained austenite duplex structures (FADP microstructures) after an intercritical anneal, regardless of whether the specimen was quenched or air cooled. ii. The microstructure comprised of significant amount of retained austenite (about 50 vol%) and an equiaxed structure. iii. Cold rolling significantly reduced the retained austenite by strain induced transformation to martensite, and simultaneously twinned some of the retained austenite. iv. Despite the similarity between the microstructures, the mechanical properties of all these steels are slightly different between each other. The results clearly showed that the ductility of these 10% Mn steels increases with the retained austenite fraction of the microstructure, and the strength increases with the increase of transformed strain induced martensitic volume fraction.
Acknowledgments This work was carried out in the framework of automotive R&D program of the AUTO-21 in Canada. The authors also acknowledge financial support from the CANMET-MTL. References [1] D.K. Matlock, J.G. Speer, in: A. Haldar, S. Suwas, D. Bhattacharjee (Eds.), Proceedings of the International Conference on Microstructure and Texture in Steels, Springer, London, 2009, pp. 185–205. [2] H. Aydin, E. Essadiqi, I.-H. Jung, S. Yue, Mater. Sci. Eng. A 564 (2013) 501–508. [3] B. Cooman, P. Gibbs, S. Lee, D. Matlock, Metall. Mater. Trans. A 44 (2013) 2563–2572. [4] S. Lee, S.-J. Lee, B.C. De Cooman, Scr. Mater. 65 (2011) 225–228. [5] S.-J. Lee, S. Lee, B.C. De Cooman, Scr. Mater. 64 (2011) 649–652. [6] C.W. Bale, E. Bélisle, P. Chartrand, S.A. Decterov, G. Eriksson, K. Hack, I.H. Jung, Y.B. Kang, J. Melançon, A.D. Pelton, C. Robelin, S. Petersen, Calphad 33 (2009) 295–311. [7] A.K. De, J.G. Speer, D.K. Matlock, Adv. Mater. Process. 161 (2003) 27–31. [8] B.D. Cullity, Elements of X-Ray Diffraction, Addison-Wesley Pub. Co., Reading, Mass, 1978. [9] J. Mahieu, B.C. Cooman, J. Maki, Metall. Mater. Trans. A 33 (2002) 2573–2580. [10] S.L.B.C. De Cooman, in: K.D. Hickey (Ed.), Proceedings of the International Symposium on New Developments in Advanced High-Strength Sheet Steels, Ronald E, Ashburn, Vail, Colo.,USA, 2013, pp. 95–103.
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[11] I.B. Timokhina, P.D. Hodgson, H. Beladi, E.V. Pereloma, Metall. Ital. 101 (2009) 43–48. [12] P. Adler, G. Olson, W. Owen, Metall. Mater. Trans. A 17 (1986) 1725–1737. [13] B.C. De Cooman, Kwang-geun Chin, Jinkyung Kim, in: M. Chiaberge (Ed.), New Trends and Developments in Automotive System Engineering, 2011. [14] O. Bouaziz, Scr. Mater. 66 (2012) 982–985. [15] O. Bouaziz, S. Allain, C.P. Scott, P. Cugy, D. Barbier, Curr. Opin. Solid State Mater. Sci. 15 (2011) 141–168. [16] D.Q. Bai, A. Di Chiro, S. Yue, Mater. Sci. Forum 284–192 (1998) 253–260.
[17] H.K.D.H. Bhadeshia, R.W.K. Honeycombe, Steels: Microstructure and Properties, Elsevier; Butterworth-Heinemann, Amsterdam; Boston, 2006. [18] A. García-Junceda, F.G. Caballero, C. Capdevila, C. García de Andrés, Scr. Mater. 57 (2007) 89–92. [19] A.R. Marder, Metall. Trans. A 12 (1981) 1569–1579. [20] J. Speer, D.K. Matlock, B.C. De Cooman, J.G. Schroth, Acta Mater. 51 (2003) 2611–2622. [21] B.S. Seong, E.J. Shin, Y.S. Han, C.H. Lee, Y.J. Kim, S.J. Kim, Phys. B: Condens. Matter 350 (2004) E467–E469.
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
Twinning and Tripping in 10% Mn steels Huseyin Aydin a,n, In-Ho Jung a, Elhachmi Essadiqi b, Stephen Yue a a b
McGill University, Department of Mining and Materials Engineering, Montreal, QC, Canada Université Internationale de Rabat, UIR, Aerospace Engineering School, Rabat, Technopolis Shore Bypass Rabat-Salé, Morocco
art ic l e i nf o
a b s t r a c t
Article history: Received 14 August 2013 Received in revised form 22 October 2013 Accepted 29 October 2013 Available online 5 November 2013
In the present work, a medium Mn, Fe–Mn–C–Al–Si alloy was subjected to different heat treatment conditions and subsequent deformation to understand the effect of these processes on austenite. It was found that, after intercritical annealing, the microstructure was ferrite plus austenite duplex phase (FADP) regardless of cooling rate to room temperature. When cold rolled, the retained austenite of the FADP structures exhibited both twinning and strain induced transformation (SIT) to martensite. A detailed characterization of the co-existence of twinning and SITing after cold rolling is presented. & 2013 Elsevier B.V. All rights reserved.
Keywords: Twinning Strain induced transformation (SIT) Stacking fault energy (SFE) Retained austenite (RA)
1. Introduction Over the last few decades, growing demands for weight saving and safety requirement have motivated new concepts of automotive steels to achieve improved mechanical properties in comparison with the existing Advanced High Strength Steels (AHSS). Among various developments, medium Mn content steels are considered to be potential candidates to achieve the performance targets of the so-called third generation AHSS [1]. These steels are currently one of the materials being increasingly developed to harness the promise of second generation steels without high alloy costs and difficult processing issues associated with these latter steels. Indeed, several literature studies proved that the medium Mn steels containing 5–10 wt% Mn have enhanced mechanical properties compared to the first generation AHSS due to the occurrence of deformation induced martensitic transformation [2–5]. It is also stated that the partitioning of Mn and C in the austenite during intercritical annealing are the two main contributions for the austenite stability, and mechanical stabilization of the austenite does not contribute to the austenite stability due to the very low dislocation density of the austenite grains [3–5]. Hence, these steels typically achieve very high strengths, but they have very limited strain hardening and, as a consequence, lower elongations compared to second generation AHSS. This paper describes a novel ferrite plus retained austenite microstructure based on the above alloying concepts. The behavior
n
Corresponding author. Tel.: þ 1 514 773 1559; fax: þ 1 514 398 4492. E-mail address: [email protected] (H. Aydin).
0921-5093/$ - see front matter & 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.msea.2013.10.088
of retained austenite during room temperature deformation of three heat treated variants was investigated, with a focus on twinning and strain induced transformation characteristics. 2. Materials and experimental methods 2.1. Compositions The steel used throughout this work was supplied by CANMETMTL (Hamilton, Ontario, Canada). Castings were done in an induction furnace. The composition (wt%) of the examined steel is shown in Table 1. This is one composition selected from a series of alloys which were designed on the basis of attaining a certain level of stacking fault energy (SFE) in the metastable austenite to promote twinning as opposed to transformation to martensite [2]. The thermodynamic modeling approach was adopted to calculate the magnitude of stacking fault energy and most of the data used in this study was taken from the literature and Fsteel database of FactSage (version 6.2) computational thermodynamic software [6]. 2.2. Processing The cast ingot was sectioned into plates having a thickness of 30 mm, and was homogenized for 3 h at 1200 1C, immediately hot rolled to 6 mm at temperatures between 1100 and 750 1C and then water quenched. The as-hot rolled plates were intercritically annealed at elevated temperatures and then either air (AC), furnace (FC) cooled or water quenched (WQ) prior to mechanical testing and microstructural characterization. Fig. 1 shows the detailed heat treatment process of the steel samples. The
H. Aydin et al. / Materials Science & Engineering A 591 (2014) 90–96
91
Table 1 The bulk chemical composition of Fe10Mn steel. Alloy
C (%)
Mn (%)
Si (%)
Al (%)
Mo (%)
Fe
Fe10Mn
0.2
10.02
3.17
3.19
0.06
Bal.
Fig. 2. The phase diagram of Fe10Mn steel alloy (α: ferrite, γ: austenite).
3. Results and discussions Fig. 1. Heat cycle of the intercritical annealing process.
3.1. Microstructure
annealing temperatures were chosen to observe the transformation behavior of austenite at different temperatures and cooling rates (i.e. temperature conditions at highest FCC and/or C content etc.). The main objectives of the heat treatments were firstly to maximize the amount of retained austenite and then, to control the stability of retained austenite. Finally, cold rolling was also conducted on the annealed samples to observe the deformation behavior of the austenite.
2.3. Characterization The microstructural examinations were initially carried out on a Nikon L150 optical microscope, Bruker D8 X-ray difractometer and Philips XL-30 field emission scanning electron microscope (FE-SEM). All samples were mechanically polished down to 0.05 μm with colloidal silica and then chemically etched using a solution of 2% Nital followed by 10% aqueous sodium metabisulfite (Na2S2O5). The etchants used in this study have been widely used for AHSS steels to differentiate ferrite, austenite and martensite. So it is indeed required to differentiate them before analyses [7]. Furthermore, since there is no carbon in the etching solutions it was ignored during the Electron Probe Micro-Analyze (EPMA). The volume percentage of the phases were calculated with the help of Clemex Captive image analyze software and an XRD intensity correlation method [8]. Transmission electron microscopy (TEM) studies were conducted on Philips CM200 transmission electron microscope (TEM). The thin foil preparation involved mechanical thinning and jet polishing at 50 V in a solution containing 340 ml butanol, 600 ml methanol, and 60 ml perchloric acid at 20 1C. The BAHR DIL 805 quench dilatometer was used to evaluate the transformation behavior and to confirm the martensite start (Ms) temperature of steel composition at fast cooling rates. Finally, mechanical tests were performed using an MTS hydraulic machine by using 100 kN load cell. Tensile test samples were machined according to ASTM E-8 sub-size standard with a strain rate of 0.1 mm/s.
According to the FactSage predicted phase diagram (Fig. 2), the equilibrium structure at room temperature of this composition has no austenite. However, this composition and the above heat treatments generated a microstructure of ferrite plus austenite duplex (FADP) microstructure. Fig. 3 shows the morphology of the FADP structure after the two different heat treatment conditions. The first generation AHSS, relied on bainite to help retain (metastable) austenite, thus leading to an equiaxed ferrite plus bainite structure, with the retained austenite largely adopting morphologies associated with the bainite. In this structure, no bainite is required to retain the austenite, and the structure is comprised of equiaxed austenite and ferrite grains. Since the predicted equilibrium microstructure at room temperature has no austenite, the austenite in this structure is assumed to be unstable. As shown in Fig. 3, the grain sizes of the ferrite and austenite are more or less comparable to each other, with the ferrite appearing to be a little coarser than the austenite grains. The microstructures also showed equiaxed austenitic grains with a limited number of annealing twins and very small amount of martensite in the austenitic phases, but apparently with more annealing twins after air cooling. One of the most interesting points of these microstructures is that there is no significant effect of cooling rate; the retained austenite levels are very similar (about 50 vol%) even after furnace cooling, indicating that the kinetics of the transformation to ferrite are relatively slow. Particularly striking is the absence of martensite after water quenching, indicating the martensite start (Ms) temperature is well below room temperature. In order to estimate the Ms of the austenite formed at the intercritical temperature, the retained austenite composition in the WQ condition (i.e., as-quenched) was measured by EPMAWDS from the microstructure in Fig. 4 and shown in Table 2. The partitioning of C and Mn to austenite can be seen in Table 2. Unfortunately, there is no empirical equation in the literature to predict Ms from the austenite composition that has been created for medium Mn steels. Moreover, all such empirical equations have been developed by experiments that begin with 100% austenite, whereas the transformation that is of interest here is from the two phase ferrite plus austenite condition. Nevertheless, the above composition was used to determine the Ms temperatures
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H. Aydin et al. / Materials Science & Engineering A 591 (2014) 90–96
RA
RA
F F
F
RA
Fig. 3. The optical images of Fe10Mn after two different heat treatment conditions; (a) air cooled, (b) water quenched and (c) furnace cooled. RA and F are retained austenite and ferrite, respectively.
F
Spectrum RA
Fig. 4. The BSE image and WDS analysis of Fe10Mn after WQ.
Fig. 5. Dilatometer analysis of Fe10Mn at gas nitrogen cooling.
Table 2 The retained austenite compositions measured by EPMA.
very high cooling rates (Fig. 5) but no sign of any transformation was observed.
Alloy
C (%)
Mn (%)
Si (%)
Al (%)
Mo (%)
Fe
Fe10Mn
1.10
11.2
2.95
2.97
0.09
Bal.
according to the following equation [9]: M s ¼ 539 423C 30:4Mn 7:5Si þ 30Al
ð1Þ
The Ms temperature is predicted to be 199.8 1C, therefore the austenite is expected not to transform martensite during quenching to room temperature. The transformation behavior of austenite to martensite was also confirmed with dilatometer experiments at
3.2. Effect of cold rolling on the microstructure Having noted the above insignificant effect of cooling rate on the amount of austenite it might be argued that this austenite is stable and not unstable as predicted by FactSage. Indeed, in contrast to the FactSage predictions, the ThermoCalc studies of De-Cooman et al. show that the pseudo-binary Fe–C phase diagrams of 10 Mn% steels have single phase austenitic area at elevated temperatures and stable austenite can be seen at room temperature conditions [10].
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However, after cold rolling, all the microstructures exhibited significant reduction in retained austenite, as measured by XRD (Table 3). Since it can be assumed that this was due to strain induced transformation (SIT) to martensite, this indicates that the austenite is unstable (i.e. retained). In the case of the differences of the austenite prediction, there is no obvious kinetic issue that could explain transformation of austenite during room temperature deformation if it is not unstable. Furthermore, cold rolling not only transformed approximately half of the retained austenite to strain induced martensitic, but also generated twinning, as is clearly observed both optically (Fig. 6) and in the TEM examination (Fig. 7). The only other reference to the simultaneous observation of twinning and SITing in retained austenite that the authors have found is by Timokhina et al. [11] for a conventional TRIP steel microstructure (i.e. ferrite, bainite and retained austenite); however, this was in the as-hot rolled condition – no other papers have reported twinning and SITing of retained austenite after cold deformation. Some austenite grains exhibited a mixture of SIT and twinning, whilst other grains revealed either twinning only or SIT only. It is clearly of interest to understand exactly why the retained austenite in the FADP steel can undergo both twinning and SIT during deformation, whereas the retained austenite in TRIP steels undergoes only strain induced transformation and the retained austenite in TWIP steels undergoes only twinning. However, in all Table 3 The amount of retained austenite after cold rolling. Austenite volume fraction (%)
Air cooled Water quenched Furnace cooled
Heat treated
Cold rolled
Δtransformed
47 53 47.5
23 32 28.4
24 21 19.1
93
previous studies, the microstructures that undergo twinning are 100% austenite which is energetically stable, e.g., TWIP and Hadfield steels [12–15], so there is no chance for strain induced transformation to martensite. Therefore the relevant comparison of FADP steels is with TRIP steels rather than TWIP steels. In this regard, the retained austenite grains in FADP are much larger and the level of retained austenite is much higher. Moreover, there is only one morphological type of retained austenite in FADP, i.e. equiaxed, whereas there are at least three types of retained austenites in TRIP steels ((i) equiaxed small grains, (ii) interlath layers in the bainite and (iii) layers between bainite packets and equiaxed ferrite grains) [16,17]. It is well known that decreasing the retained austenite size in TRIP steels increases the resistance against SIT. It may be that size is also a factor in twinning. Equally important is to determine when twinning and SIT take place during deformation, i.e., the evolution of the microstructure with strain. In fact, there are almost certainly three deformation mechanisms operating: (i) slip, (ii) SIT, and (iii) twinning. In terms of modeling the events, if these mechanisms are treated purely as ways to relieve the applied stress, then the obvious approach is to use the concept of critical resolved shear stress via the Schmidt factor. This approach is well known for slip and twinning; the concept of critical stress is also used in SIT [18,19], but the Schmidt factor is not considered. Another way to consider the problem is by a classic Gibbs energy approach; this is commonly used to explain TRIP [17,20] but not for twinning. With this approach, twinned retained austenite and strain induced martensite could be considered as being two phases in equilibrium. Fig. 8 shows a schematic diagram of the effect of composition on the free energies of these two phases. If the diffusion of solutes is allowed, the two phase assemblage is delineated by the common tangent between these curves, which leads to the two co-existing phases having different compositions; this does not occur in this case because both SIT and twinning are diffusionless ‘transformations’, and the compositions
Twins
F Twins M
F
RA M
M RA
F
Twins Fig. 6. The optical images of Fe10Mn after cold rolling (a) air cooled, (b) water quenched and (c) furnace cooled samples. RA, M and F are retained austenite, martensite and ferrite, respectively.
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Fig. 7. Deformation twins and SIT in Fe10Mn (a) bright field image and (b) weak-beam dark field TEM image of only twins and (c) bright field image and (d) dark field image of both twins (oriented along [011] zone axis (FCC)) and SIT (oriented along [111] (BCT)).
Fig. 9. Engineering stress–strain curves.
Fig. 8. Gibbs energy–composition diagram for the austenite and strain induced martensite phases at a given temperature.
of twinned and SIT phases are therefore the same. Therefore, in order to satisfy the co-existence of twinned austenite (γtwin) and SIT phase (α1), the Gibbs energies of the two phases should be the same or, in practice, very similar for the existing composition of the retained austenite. This scenario maybe feasible as twinning and tripping are mechanisms at high (about 20%) and low (2%) Mn concentrations, respectively. EPMA analysis [2] has shown that the Mn in the retained austenite of FADP is 11.2%, which is more or less
Table 4 The mechanical properties of Fe10Mn after tensile tests. Fe10Mn
UTS (MPa)
% εtotal
Rp0.2 (MPa)
WQ þ Tensile AC þ Tensile FCþ Tensile
868 903 776
31 23 25
637 694 567
H. Aydin et al. / Materials Science & Engineering A 591 (2014) 90–96
95
Fig. 10. Effect of (a) retained austenite fraction on strain and (b) SIT transformation on stress.
exactly between the high and low Mn levels. However, it would be very fortuitous to have attained the exact compositional ‘fit’ in this experiment where the free energies of strain induced martensite and tinned retained austenite exactly coincide. Therefore, this ‘classical’ ΔG curve approach does not satisfactorily explain the coexistence of twinning and tripping. A ‘hybrid’ of the CRSS and free energy approaches may be required to satisfactorily analyze this simultaneous twinning and SITing phenomenon. Since it is accepted that there is a strong effect of Mn and C on the SFE and transformation kinetics during deformation, one other reason why both SIT and twinning are seen in the retained austenite may be due to an inhomogeneous carbon distribution within the retained austenite. This could lead an incomplete transformation of retained austenite and these regions of austenite tend to transform to martensite or twinning at lower strains [11]. 3.3. Mechanical properties Engineering stress–strain curves are displayed in Fig. 9 for the two annealed samples and Table 4 summarizes their mechanical properties. Firstly, the WQ condition leads the highest total elongation of more than 30%, which is significant enhancement for medium Mn steel compositions, and it was balanced with a relatively high UTS value of 868 MPa. Annealing with AC led to slightly higher UTS more than 900 MPa but lower ductility values of 23%. In case of FC the total elongation was compatible with AC but at a much lower strength value because of the softer and coarser grain morphology. As shown in Fig. 10, increasing amount of transformation increases strength, which is basically due to increasing strain induced martensite. However, increasing transformation does not lead to increasing ductility; this demonstrates that increasing SIT increases ductility only if it postpones strain instability. In this case, it appears that SIT may take place well before strain instability, thus contributing little to ductility. In fact, increasing ductility correlates better with the increasing retained austenite, which might be due to FCC being fundamentally more ductile than ferrite or martensite. It is also interesting to note that the ductility of the asquenched specimens is slightly higher than the air cooled ones for a given RA vol%. This may be related to the fact that in the WQ specimen, the intercritical austenite was formed at higher temperature, thus leading to a lower carbon content [21], as confirmed by detailed EPMA studies in Reference [2]. This lower C content might make retained austenite softer, thereby increasing the ductility. The effect of twinning has not been considered but it is expected enhance the mechanical properties. Unfortunately there
is not enough information to quantify the separate contributions of twinning and SIT in the microstructure.
4. Summary i. Fe10Mn generated a ferrite and retained austenite duplex structures (FADP microstructures) after an intercritical anneal, regardless of whether the specimen was quenched or air cooled. ii. The microstructure comprised of significant amount of retained austenite (about 50 vol%) and an equiaxed structure. iii. Cold rolling significantly reduced the retained austenite by strain induced transformation to martensite, and simultaneously twinned some of the retained austenite. iv. Despite the similarity between the microstructures, the mechanical properties of all these steels are slightly different between each other. The results clearly showed that the ductility of these 10% Mn steels increases with the retained austenite fraction of the microstructure, and the strength increases with the increase of transformed strain induced martensitic volume fraction.
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