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Effects of Nb and C additions on the microstructure and tensile properties of lightweight ferritic Fe–8Al–5Mn alloy
Available online at www.sciencedirect.com
ScienceDirect Scripta Materialia 89 (2014) 37–40 www.elsevier.com/locate/scriptamat
Effects of Nb and C additions on the microstructure and tensile properties of lightweight ferritic Fe–8Al–5Mn alloy A. Zargaran,a H.S. Kim,a J.H. Kwakb and Nack J. Kima,⇑ a
Graduate Institute of Ferrous Technology, POSTECH, Pohang 790-784, Republic of Korea b Technical Research Lab., POSCO, Gwangyang, Republic of Korea Received 30 April 2014; revised 23 June 2014; accepted 23 June 2014 Available online 8 July 2014
Additions of Nb and C result in the formation of NbC and j-carbide in Fe–8Al–5Mn alloy. Both NbC and j-carbide particles inhibit recrystallization during hot rolling with the resultant formation of pancake-shaped grain structure in Fe–8Al–5Mn–0.1Nb– 0.1C alloy, but j-carbide particles also promote recrystallization via particle-stimulated nucleation mechanism during subsequent warm/cold rolling and annealing. As a result, Fe–8Al–5Mn–0.1Nb–0.1C alloy has a much finer grain size and a better combination of tensile properties than Fe–8Al–5Mn alloy. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Lightweight steels; Ferrite; j-Carbides; Recrystallization; Grain refinement
It is well known that the addition of Al to steels considerably decreases their densities [1,2], making the so-called Al-containing lightweight steels potential structural materials particularly for automotive applications that require high specific strength (strength to weight ratio). There are basically three variants of lightweight steels: ferrite-based [1,3–13], austenite-based [14–16] and ferrite–austenite duplex [11,17] steels. The solute contents in the ferrite-based steels are generally lower than those in the austenite-based and duplex steels, which require high solute contents, particularly of Mn, to stabilize austenite as the matrix phase. However, the ferrite-based steels have a poorer combination of strength and ductility than the austenite-based and duplex steels [2]. Although Al can increase the strength of the lightweight steels considerably (46 MPa per 1 wt.% Al) by solid-solution hardening [5,10], the maximum amount of Al is usually limited to 8 wt.% in ferrite-based lightweight steels since above this level there is a large decrease in ductility due to ordering reactions [1,3] as well as the reduced mobility of dislocations due to the large solid-solution hardening effect of Al [5]. Although
⇑ Corresponding
author at: Graduate Institute of Ferrous Technology, POSTECH, Pohang 790-784, Republic of Korea; e-mail: [email protected]
C is frequently utilized as an important alloying element in austenite-based and duplex steels [2,14–16], addition of C has usually been avoided in ferrite-based lightweight steels due to the deleterious effect of j-carbide, (Fe,Mn)3AlC, particles forming at grain boundaries on ductility [3–5]. However, recent studies have shown that j-carbide can be utilized as a strengthening phase in the high C (0.3–1.1 wt.%) containing ferrite-based lightweight steels without drastically degrading ductility when its morphology is optimized through appropriate thermomechanical treatments (TMTs) [7,8,11,12]. Microalloying additions of Nb, V and Ti have also been attempted to increase the strength of the ferrite-based lightweight steels [1,3,4], but their effects on tensile properties vary depending on investigations and moreover the related mechanism is not clear. Interestingly, despite the important role of grain size in the mechanical properties of structural materials, there have been only a few studies on the grain refinement of the ferrite-based lightweight steels [3,4]. In fact, most of the ferrite-based lightweight steels have rather coarse grain size ranging from 40 to 90 lm in the final products [3,4,6,11–13]. Unlike other conventional steels showing austenite-to-ferrite transformation during cooling after hot rolling, the ferrite-based lightweight steels are usually subjected to hot rolling in the ferrite region due to their expanded ferrite phase field and accordingly
http://dx.doi.org/10.1016/j.scriptamat.2014.06.018 1359-6462/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
38
A. Zargaran et al. / Scripta Materialia 89 (2014) 37–40
cannot utilize the grain refinement by phase transformation during cooling. Therefore, one plausible approach to refine the grain size is the utilization of warm/cold rolling followed by annealing, which induces the recrystallization of deformed ferrite grains into finer ferrite grains. Nevertheless, the grain sizes of the final products are still coarse, ranging from 40 to 90 lm as mentioned above, indicating that an additional approach needs to be taken to further refine the grain size besides the conventional approach of rolling followed by annealing. In the present paper, we report our work on the grain refinement and the resultant tensile properties of the ferrite-based lightweight steel by the microalloying additions of Nb and C, which have been actively utilized in conventional highstrength low-alloy (HSLA) steels [18]. Two alloys with nominal compositions of Fe–8Al– 5Mn and Fe–8Al–5Mn–0.1Nb–0.1C (wt.%) were prepared by vacuum induction melting. The compositions of the alloys based on wet chemical analysis are shown in Table 1. After homogenization at 1200 °C for 1 h, they were hot rolled at 1100 °C with a finish rolling temperature of 1000 °C to produce plates 20 mm thick. These hot-rolled plates were then warm rolled at 650 °C to 3 mm thick sheets (85% reduction). After an intermediate annealing at 850 °C for 15 min, the sheets were cold rolled to 1 mm in thickness (67% reduction) followed by final annealing at 750 °C for 1 h. All specimens were air cooled after rolling or annealing. Microstructure was analyzed by optical microscopy (OM), scanning Table 1. Analyzed chemical compositions of the alloys (in wt.%). Alloy
Al
Mn
Nb
C
Fe–8Al–5Mn Fe–8Al–5Mn–0.1Nb–0.1C
7.9 8.3
4.5 4.9
– 0.095
0.017 0.094
Figure 1. Microstructure of the alloys in as-hot-rolled condition: OM micrographs of (a) Fe–8Al–5Mn and (b) Fe–8Al–5Mn–0.1Nb–0.1C, (c) SEM micrograph of j-carbide within ferrite matrix of Fe–8Al– 5Mn–0.1Nb–0.1C, (d) TEM micrograph of NbC and j-carbide along grain boundary of Fe–8Al–5Mn–0.1Nb–0.1C, (e) schematics of diffraction patterns of j-carbide (dotted/black) and ferrite (solid/ red), and (f) schematics of diffraction patterns of NbC (dotted/black) and ferrite (solid/red). ND: normal direction, RD: rolling direction. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
electron microscopy (SEM) at an acceleration voltage of 15 kV, electron backscatter diffraction (EBSD) with a step size of 0.5 lm, and transmission electron microscopy (TEM) at an acceleration voltage of 200 kV. The average grain size was measured for the grains having minimum misorientation angles of 15°. In addition, an in situ heating EBSD experiment was conducted in order to investigate the recrystallization behavior. A step size of 0.05 lm was used to examine the microstructure before heating and this was increased to 0.1 lm after heating to reduce the scanning time due fast recrystallization kinetics. Tensile tests were conducted using the specimens with a gauge length of 12.6 mm, a gauge width of 5 mm, and a gauge thickness of 1 mm at the strain rate of 10 3 s 1. Tensile loading direction was parallel to the rolling direction (RD). Figure 1 shows the micrographs of the alloys in as-hotrolled condition. Very coarse grains are elongated along RD in both alloys, but the size of the grains through the thickness of sheet is much smaller in the Fe–8Al–5Mn– 0.1Nb–0.1C alloy than in the Fe–8Al–5Mn alloy (160 vs. 850 lm). Although both alloys show an inhomogeneous grain structure with some partially recrystallized grains, it can be seen that Fe–8Al–5Mn–0.1Nb–0.1C alloy has much more elongated pancake-shaped grain structure than Fe–8Al–5Mn alloy, indicating that the recrystallization is prohibited more in Fe–8Al–5Mn–0.1Nb–0.1C alloy than in Fe–8Al–5Mn alloy. Moreover, partially recrystallized grains in Fe–8Al–5Mn alloy are coarser than those in Fe–8Al–5Mn–0.1Nb–0.1C alloy. While Fe-8Al-5Mn alloy shows a fully ferritic structure after hot rolling, Fe-8Al-5Mn-0.1Nb-0.1C alloy shows the formation of j-carbide particles within ferrite matrix as well as along grain boundaries as shown in Figure 1b–d. While the j-carbide particles within ferrite matrix exist as needle-shape particles as shown in Figure 1c, the j-carbide particles along grain boundaries are polygonal in shape and are connected to each other, forming a continuous network as shown in Figure 1d. They both have a Nishiyama–Wasserman (N–W) orientation relationship with ferrite, (0–11)Fe // (1–11) j-carbide and [100] Fe // [110] j-carbide, as shown in Figure 1d and e. In addition, NbC particles have been observed along the grain boundaries, having a Baker–Nutting (B–N) orientation relationship with ferrite, (002)Fe // (002)NbC and [100] Fe // [110] NbC (Fig. 1d–f). After cold rolling followed by final annealing at 750 °C for 1 h, the microstructure becomes much more homogeneous as compared to that of as-hot-rolled alloys as shown in Figure 2. Figure 2 also shows that the average grain size of Fe–8Al–5Mn–0.1Nb–0.1C alloy (10 lm) is considerably finer than that of Fe–8Al–5Mn alloy (32 lm). Room-temperature tensile stress–strain curves of the investigated alloys are shown in Figure 3a and their tensile properties are summarized in Table 2. It can be seen that the Fe–8Al–5Mn–0.1Nb–0.1C alloy has higher yield and tensile strengths than Fe–8Al–5Mn alloy. A yield point phenomenon occurs in Fe–8Al–5Mn–0.1Nb–0.1C alloy. Although there are rather large differences in yield and tensile strengths, Fe–8Al–5Mn–0.1Nb–0.1C alloy has only slightly smaller elongation than Fe–8Al–5Mn alloy. As shown in Figure 3b, compared to other representative
A. Zargaran et al. / Scripta Materialia 89 (2014) 37–40
Figure 2. Inverse pole figure (IPF) maps of (a) Fe–8Al–5Mn and (b) Fe–8Al–5Mn–0.1Nb–0.1C after final annealing.
Figure 3. (a) Engineering tensile stress–strain curves of the alloys and (b) yield strength and total elongation of the ferritic lightweight steels.
ferritic lightweight steels, the present alloys show a better combination of strength and ductility. The present study shows that additions of Nb and C have significant effects on the microstructure and tensile properties of Fe–8Al–5Mn alloy. As shown above (Fig. 1b–d), Fe–8Al–5Mn–0.1Nb–0.1C alloy contains NbC and j-carbide particles, which are undoubtedly related to its resultant grain size and tensile properties. Analyses of the microstructural evolution during various stages of TMTs show the respective roles of NbC and j-carbide in grain refinement. As can be seen in Figure 1d, NbC particles are present at grain boundaries along with j-carbide particles in the as-hot-rolled condition. Calculation of the solubilities of Nb and C shows that NbC could form at below 1200 °C in ferrite [19] and therefore the pancake-shaped grain structure of Fe–8Al–0.1Nb– 0.1C alloy can be thought to be due to the effect of NbC on inhibiting the recrystallization of ferrite during hot rolling, similar to its role in inhibiting the recrystallization of austenite in microalloyed steels during rolling [18]. The presence of j-carbide particles along grain boundaries in the as-hot-rolled condition also suggests their possible role in inhibiting the recrystallization of ferrite during rolling. Since there is no reliable database for the thermodynamic calculation of phase diagrams for the Fe–Al–Mn–Nb–C alloy system, it is not quite clear at what temperature j-carbide forms during cooling. It is possible for j-carbide to form during cooling after rolling considering its continuous network along grain boundaries. However, if j-carbide particles can form during hot rolling at between 1100 and 1000 °C as reported by Rigaud et al. [20], it is quite possible for j-carbide particles to play the similar role to that
39
of NbC particles such that they would inhibit the recrystallization of ferrite during rolling. In fact, Baligidad also observed the formation of pancake-shaped ferrite grains in hot-rolled Fe–8.8Al–0.1C alloy in as-hot-rolled condition [9], although neither the detailed analyses on microstructural development nor the effect of subsequent warm/cold rolling and annealing were conducted. Although it is known that the formation of j-carbide particles along grain boundaries often adversely affects the rollability during subsequent TMTs [13], their presence can be favorably utilized to refine the grain size. To understand the microstructural evolution of Fe–8Al– 5Mn–0.1Nb–0.1C alloy, microstructural observation has been conducted on specimens subjected to subsequent TMTs, i.e. warm/cold rolling and annealing, as shown in Figure 4. Figure 4 also shows that the coarse needle-shaped j-carbide particles (Fig. 1c) in the ashot-rolled condition become elongated particles that form a banded structure along RD after warm rolling to 3 mm (85% reduction) at 650 °C (Fig. 4a). It can also be seen that some of j-carbide particles are cut and fractured by deformation. During annealing of the warm-rolled specimen at 850 °C for 15 min, the banded j-carbide particles become separated, forming discrete particles, similar to the spheroidization of pearlite lamellae observed in warm-rolled and annealed 5140 steel [21]. Recrystallization also occurs, forming a much finer grain structure (grain size: 25 lm) than that of as-hotrolled specimen (Fig. 4b). The distribution of j-carbide particles becomes quite uniform after final cold rolling to 1 mm (67% reduction) and annealing at 750 °C for 1 h (Fig. 4c) and there is a significant grain refinement as shown in Figures 2b and 4c. On the other hand, the grain sizes of the Fe–8Al–5Mn alloy after warm rolling and annealing (Fig. 4d) and final cold rolling and annealing (Fig. 2a) are 55 lm and 32 lm, respectively, which are much coarser than those of the Fe–8Al–5Mn–0.1Nb–0.1C alloy at the same processing conditions. It suggests that the presence of j-carbide particles in Fe–8Al–5Mn–0.1Nb–0.1C alloy plays a major role in refining its grain size by promoting recrystallization during subsequent TMTs after hot rolling. It is well known that recrystallization of deformed alloys can occur at grain boundaries, deformation bands and second-phase particles [22]. As shown in Figure 1, the Fe–8Al–5Mn–0.1Nb–0.1C alloy shows a much more elongated pancake-shaped grain structure than Fe-8Al5Mn alloy in as-hot-rolled condition, suggesting that there would be more deformation bands in the former than in the latter. These deformation bands could act as nucleation sites for recrystallization during warm rolling; however, the warm-rolled microstructure (Fig. 4a) does not show the occurrence of recrystallization, indicating that the amount of strain stored in as-hot-rolled condition is not sufficient to induce recrystallization at the warm-rolling temperature of 650 °C possibly due
Table 2. Room-temperature tensile properties of the alloys after final annealing. Alloy
YS (MPa)
UTS (MPa)
Uniform elongation (%)
Total elongation (%)
Fe–8Al–5Mn Fe–8Al–5Mn–0.1Nb–0.1C
484.0 599.6
592.2 694.5
16.9 16.5
35.9 30.7
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A. Zargaran et al. / Scripta Materialia 89 (2014) 37–40
suggests that PSN is the dominant grain refinement mechanism for Fe–8Al–5Mn–0.1Nb–0.1C alloy since it has much finer grain size than the Fe–8Al–5Mn alloy, which does not contain any second-phase particles. In summary, the effects of 0.1Nb and 0.1C additions on the microstructure and tensile properties of Fe–8Al–5Mn ferritic lightweight steels have been characterized. Our results show that the microstructure of Fe–8Al–5Mn– 0.1Nb–0.1C alloy contains NbC and j-carbide particles within ferrite matrix and along grain boundaries. The microstructure of the Fe–8Al–5Mn–0.1Nb–0.1C alloy is much more elongated along RD than that of Fe–8Al– 5Mn alloy due to the effect of both NbC and j-carbide particles on inhibiting the recrystallization of ferrite during hot rolling. The Fe–8Al–5Mn–0.1Nb–0.1C alloy has a much finer grain size than the Fe–8Al–5Mn alloy after final annealing due to the enhanced recrystallization by PSN mechanism associated with j-carbide particles in the former and accordingly the former shows a better combination of tensile properties than the latter. Authors are grateful for the generous support of POSCO. Figure 4. Microstructure of the alloys after TMTs: (a) Fe–8Al–5Mn– 0.1Nb–0.1C after warm rolling (inset: magnified image of black rectangular area), (b) Fe–8Al–5Mn–0.1Nb-0.1C after annealing of warm-rolled sheet, (c) Fe–8Al–5Mn–0.1Nb–0.1C after final annealing of cold-rolled sheet, (d) Fe–8Al–5Mn after annealing of warm-rolled sheet, and IPF maps of cold-rolled Fe–8Al–5Mn–0.1Nb–0.1C (e) before heating, (f) after heating at 700 °C for 5 min showing the nucleation of new ferrite grains at j-carbide particles (arrowed), and (g) after heating at 700 °C for 15 min.
to the strain relief during cooling after hot rolling. It also indicates that the amount of strain received during warm rolling at 650 °C is not sufficient for dynamic recrystallization to occur. However, the warm-rolled specimens show the occurrence of static recrystallization after annealing at 850 °C, which is higher than the warm-rolling temperature (650 °C). On the other hand, the presence of j-carbide particles in the Fe–8Al–5Mn– 0.1Nb–0.1C alloy would play an important role in grain refinement by the operation of particle-stimulated nucleation (PSN) mechanism [22]. An in situ heating EBSD study of the cold-rolled Fe–8Al–5Mn–0.1Nb– 0.1C alloy shows that recrystallization readily occurs along j-carbide particles with the formation of new grains having different orientations from that of the matrix (Fig 4e–g). j-carbide particles can also inhibit the growth of newly formed grains by pinning grain boundaries as can be seen in Figure 4c, which shows the location of j-carbide particles mostly along grain boundaries. On the other hand, NbC particles would play a minor role in inhibiting grain growth considering their much smaller volume fraction than that of j-carbide particles (0.2% vs. 2.3%). As mentioned previously, the grain sizes of the Fe–8Al–5Mn–0.1Nb– 0.1C alloy in the warm-rolled and annealed condition and the cold-rolled and annealed condition are 25 and 10 lm, respectively, which are much finer that those of the Fe–8Al–5Mn alloy under the same TMT conditions (55 lm in the warm-rolled and annealed condition and 32 lm in the cold-rolled and annealed condition). It
[1] G. Frommeyer, E.J. Drewes, B. Engl, Rev. Metall. 97 (2000) 1245. [2] H. Kim, D.-W. Suh, N.J. Kim, Sci. Technol. Adv. Mater. 14 (2013) 014205. [3] U. Bru¨x, G. Frommeyer, J. Jimenez, Steel Res. Int. 73 (2002) 543. [4] R. Rana, C. Liu, R.K. Ray, Scripta Mater. 68 (2013) 354. [5] J. Herrmann, G. Inden, G. Sauthoff, Acta Mater. 51 (2003) 2847. [6] S. Vayas, S. Viswanathan, V.K. Sikka, Scripta Metall. Mater. 27 (1992) 185. [7] R.G. Baligidad, U. Prakash, A. Radhakrishna, Mater. Sci. Eng., A 269 (1999) 120. [8] R.G. Baligidad, A. Radhakrishna, Mater. Sci. Eng., A 308 (2001) 136. [9] R.G. Baligidad, J. Mater. Sci. 39 (2004) 6623. [10] H. Ishii, K. Ohkubo, S. Miura, T. Mohri, Mater. Trans. 44 (2003) 1679. [11] S.S. Sohn, B.J. Lee, S. Lee, N.J. Kim, J.H. Kwak, Acta Mater. 61 (2013) 5050. [12] S.Y. Han, S.Y. Shin, B.J. Lee, S. Lee, N.J. Kim, J.H. Kwak, Metall. Mater. Trans. 44A (2013) 235. [13] S.Y. Han, S.Y. Shin, S. Lee, N.J. Kim, J.-H. Kwak, K.-G. Chin, Metall. Mater. Trans. 42A (2011) 138. [14] G. Fromeyer, U. Bru¨x, Steel Res. Int. 77 (2006) 627. [15] K. Choi, C.-H. Seo, H. Lee, S.K. Kim, J.H. Kwak, K.G. Chin, K.-T. Park, N.J. Kim, Scripta Mater. 63 (2010) 1028. [16] I. Gutierrez-Urrutia, D. Raabe, Acta Mater. 60 (2012) 5791. [17] C.-H. Seo, K.H. Kwon, K. Choi, K.-H. Kim, J.H. Kwak, S. Lee, N.J. Kim, Scripta Mater. 66 (2012) 519. [18] M.G. Akben, I. Weiss, J.J. Jonas, Acta Metall. 29 (1981) 111. [19] T. Klymko, M.H.F. Sluiter, J. Mater. Sci. 47 (2012) 7601. [20] V. Rigaud, D. Daloz, J. Drillet, A. Perlade, P. Maugis, G. Lesoult, ISIJ Int. 47 (2007) 898. [21] J. Arruabarrena, B. Lopez, J.M. Rodriguez-Ibabe, Metall. Mater. Trans. 45A (2014) 1470. [22] F.J. Humphreys, M. Hatherley, Recrystallization and Related Annealing Phenomena, second ed., Elsevier, Oxford, 1995.
ScienceDirect Scripta Materialia 89 (2014) 37–40 www.elsevier.com/locate/scriptamat
Effects of Nb and C additions on the microstructure and tensile properties of lightweight ferritic Fe–8Al–5Mn alloy A. Zargaran,a H.S. Kim,a J.H. Kwakb and Nack J. Kima,⇑ a
Graduate Institute of Ferrous Technology, POSTECH, Pohang 790-784, Republic of Korea b Technical Research Lab., POSCO, Gwangyang, Republic of Korea Received 30 April 2014; revised 23 June 2014; accepted 23 June 2014 Available online 8 July 2014
Additions of Nb and C result in the formation of NbC and j-carbide in Fe–8Al–5Mn alloy. Both NbC and j-carbide particles inhibit recrystallization during hot rolling with the resultant formation of pancake-shaped grain structure in Fe–8Al–5Mn–0.1Nb– 0.1C alloy, but j-carbide particles also promote recrystallization via particle-stimulated nucleation mechanism during subsequent warm/cold rolling and annealing. As a result, Fe–8Al–5Mn–0.1Nb–0.1C alloy has a much finer grain size and a better combination of tensile properties than Fe–8Al–5Mn alloy. Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Lightweight steels; Ferrite; j-Carbides; Recrystallization; Grain refinement
It is well known that the addition of Al to steels considerably decreases their densities [1,2], making the so-called Al-containing lightweight steels potential structural materials particularly for automotive applications that require high specific strength (strength to weight ratio). There are basically three variants of lightweight steels: ferrite-based [1,3–13], austenite-based [14–16] and ferrite–austenite duplex [11,17] steels. The solute contents in the ferrite-based steels are generally lower than those in the austenite-based and duplex steels, which require high solute contents, particularly of Mn, to stabilize austenite as the matrix phase. However, the ferrite-based steels have a poorer combination of strength and ductility than the austenite-based and duplex steels [2]. Although Al can increase the strength of the lightweight steels considerably (46 MPa per 1 wt.% Al) by solid-solution hardening [5,10], the maximum amount of Al is usually limited to 8 wt.% in ferrite-based lightweight steels since above this level there is a large decrease in ductility due to ordering reactions [1,3] as well as the reduced mobility of dislocations due to the large solid-solution hardening effect of Al [5]. Although
⇑ Corresponding
author at: Graduate Institute of Ferrous Technology, POSTECH, Pohang 790-784, Republic of Korea; e-mail: [email protected]
C is frequently utilized as an important alloying element in austenite-based and duplex steels [2,14–16], addition of C has usually been avoided in ferrite-based lightweight steels due to the deleterious effect of j-carbide, (Fe,Mn)3AlC, particles forming at grain boundaries on ductility [3–5]. However, recent studies have shown that j-carbide can be utilized as a strengthening phase in the high C (0.3–1.1 wt.%) containing ferrite-based lightweight steels without drastically degrading ductility when its morphology is optimized through appropriate thermomechanical treatments (TMTs) [7,8,11,12]. Microalloying additions of Nb, V and Ti have also been attempted to increase the strength of the ferrite-based lightweight steels [1,3,4], but their effects on tensile properties vary depending on investigations and moreover the related mechanism is not clear. Interestingly, despite the important role of grain size in the mechanical properties of structural materials, there have been only a few studies on the grain refinement of the ferrite-based lightweight steels [3,4]. In fact, most of the ferrite-based lightweight steels have rather coarse grain size ranging from 40 to 90 lm in the final products [3,4,6,11–13]. Unlike other conventional steels showing austenite-to-ferrite transformation during cooling after hot rolling, the ferrite-based lightweight steels are usually subjected to hot rolling in the ferrite region due to their expanded ferrite phase field and accordingly
http://dx.doi.org/10.1016/j.scriptamat.2014.06.018 1359-6462/Ó 2014 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
38
A. Zargaran et al. / Scripta Materialia 89 (2014) 37–40
cannot utilize the grain refinement by phase transformation during cooling. Therefore, one plausible approach to refine the grain size is the utilization of warm/cold rolling followed by annealing, which induces the recrystallization of deformed ferrite grains into finer ferrite grains. Nevertheless, the grain sizes of the final products are still coarse, ranging from 40 to 90 lm as mentioned above, indicating that an additional approach needs to be taken to further refine the grain size besides the conventional approach of rolling followed by annealing. In the present paper, we report our work on the grain refinement and the resultant tensile properties of the ferrite-based lightweight steel by the microalloying additions of Nb and C, which have been actively utilized in conventional highstrength low-alloy (HSLA) steels [18]. Two alloys with nominal compositions of Fe–8Al– 5Mn and Fe–8Al–5Mn–0.1Nb–0.1C (wt.%) were prepared by vacuum induction melting. The compositions of the alloys based on wet chemical analysis are shown in Table 1. After homogenization at 1200 °C for 1 h, they were hot rolled at 1100 °C with a finish rolling temperature of 1000 °C to produce plates 20 mm thick. These hot-rolled plates were then warm rolled at 650 °C to 3 mm thick sheets (85% reduction). After an intermediate annealing at 850 °C for 15 min, the sheets were cold rolled to 1 mm in thickness (67% reduction) followed by final annealing at 750 °C for 1 h. All specimens were air cooled after rolling or annealing. Microstructure was analyzed by optical microscopy (OM), scanning Table 1. Analyzed chemical compositions of the alloys (in wt.%). Alloy
Al
Mn
Nb
C
Fe–8Al–5Mn Fe–8Al–5Mn–0.1Nb–0.1C
7.9 8.3
4.5 4.9
– 0.095
0.017 0.094
Figure 1. Microstructure of the alloys in as-hot-rolled condition: OM micrographs of (a) Fe–8Al–5Mn and (b) Fe–8Al–5Mn–0.1Nb–0.1C, (c) SEM micrograph of j-carbide within ferrite matrix of Fe–8Al– 5Mn–0.1Nb–0.1C, (d) TEM micrograph of NbC and j-carbide along grain boundary of Fe–8Al–5Mn–0.1Nb–0.1C, (e) schematics of diffraction patterns of j-carbide (dotted/black) and ferrite (solid/ red), and (f) schematics of diffraction patterns of NbC (dotted/black) and ferrite (solid/red). ND: normal direction, RD: rolling direction. (For interpretation of the references to color in this figure legend, the reader is referred to the web version of this article.)
electron microscopy (SEM) at an acceleration voltage of 15 kV, electron backscatter diffraction (EBSD) with a step size of 0.5 lm, and transmission electron microscopy (TEM) at an acceleration voltage of 200 kV. The average grain size was measured for the grains having minimum misorientation angles of 15°. In addition, an in situ heating EBSD experiment was conducted in order to investigate the recrystallization behavior. A step size of 0.05 lm was used to examine the microstructure before heating and this was increased to 0.1 lm after heating to reduce the scanning time due fast recrystallization kinetics. Tensile tests were conducted using the specimens with a gauge length of 12.6 mm, a gauge width of 5 mm, and a gauge thickness of 1 mm at the strain rate of 10 3 s 1. Tensile loading direction was parallel to the rolling direction (RD). Figure 1 shows the micrographs of the alloys in as-hotrolled condition. Very coarse grains are elongated along RD in both alloys, but the size of the grains through the thickness of sheet is much smaller in the Fe–8Al–5Mn– 0.1Nb–0.1C alloy than in the Fe–8Al–5Mn alloy (160 vs. 850 lm). Although both alloys show an inhomogeneous grain structure with some partially recrystallized grains, it can be seen that Fe–8Al–5Mn–0.1Nb–0.1C alloy has much more elongated pancake-shaped grain structure than Fe–8Al–5Mn alloy, indicating that the recrystallization is prohibited more in Fe–8Al–5Mn–0.1Nb–0.1C alloy than in Fe–8Al–5Mn alloy. Moreover, partially recrystallized grains in Fe–8Al–5Mn alloy are coarser than those in Fe–8Al–5Mn–0.1Nb–0.1C alloy. While Fe-8Al-5Mn alloy shows a fully ferritic structure after hot rolling, Fe-8Al-5Mn-0.1Nb-0.1C alloy shows the formation of j-carbide particles within ferrite matrix as well as along grain boundaries as shown in Figure 1b–d. While the j-carbide particles within ferrite matrix exist as needle-shape particles as shown in Figure 1c, the j-carbide particles along grain boundaries are polygonal in shape and are connected to each other, forming a continuous network as shown in Figure 1d. They both have a Nishiyama–Wasserman (N–W) orientation relationship with ferrite, (0–11)Fe // (1–11) j-carbide and [100] Fe // [110] j-carbide, as shown in Figure 1d and e. In addition, NbC particles have been observed along the grain boundaries, having a Baker–Nutting (B–N) orientation relationship with ferrite, (002)Fe // (002)NbC and [100] Fe // [110] NbC (Fig. 1d–f). After cold rolling followed by final annealing at 750 °C for 1 h, the microstructure becomes much more homogeneous as compared to that of as-hot-rolled alloys as shown in Figure 2. Figure 2 also shows that the average grain size of Fe–8Al–5Mn–0.1Nb–0.1C alloy (10 lm) is considerably finer than that of Fe–8Al–5Mn alloy (32 lm). Room-temperature tensile stress–strain curves of the investigated alloys are shown in Figure 3a and their tensile properties are summarized in Table 2. It can be seen that the Fe–8Al–5Mn–0.1Nb–0.1C alloy has higher yield and tensile strengths than Fe–8Al–5Mn alloy. A yield point phenomenon occurs in Fe–8Al–5Mn–0.1Nb–0.1C alloy. Although there are rather large differences in yield and tensile strengths, Fe–8Al–5Mn–0.1Nb–0.1C alloy has only slightly smaller elongation than Fe–8Al–5Mn alloy. As shown in Figure 3b, compared to other representative
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Figure 2. Inverse pole figure (IPF) maps of (a) Fe–8Al–5Mn and (b) Fe–8Al–5Mn–0.1Nb–0.1C after final annealing.
Figure 3. (a) Engineering tensile stress–strain curves of the alloys and (b) yield strength and total elongation of the ferritic lightweight steels.
ferritic lightweight steels, the present alloys show a better combination of strength and ductility. The present study shows that additions of Nb and C have significant effects on the microstructure and tensile properties of Fe–8Al–5Mn alloy. As shown above (Fig. 1b–d), Fe–8Al–5Mn–0.1Nb–0.1C alloy contains NbC and j-carbide particles, which are undoubtedly related to its resultant grain size and tensile properties. Analyses of the microstructural evolution during various stages of TMTs show the respective roles of NbC and j-carbide in grain refinement. As can be seen in Figure 1d, NbC particles are present at grain boundaries along with j-carbide particles in the as-hot-rolled condition. Calculation of the solubilities of Nb and C shows that NbC could form at below 1200 °C in ferrite [19] and therefore the pancake-shaped grain structure of Fe–8Al–0.1Nb– 0.1C alloy can be thought to be due to the effect of NbC on inhibiting the recrystallization of ferrite during hot rolling, similar to its role in inhibiting the recrystallization of austenite in microalloyed steels during rolling [18]. The presence of j-carbide particles along grain boundaries in the as-hot-rolled condition also suggests their possible role in inhibiting the recrystallization of ferrite during rolling. Since there is no reliable database for the thermodynamic calculation of phase diagrams for the Fe–Al–Mn–Nb–C alloy system, it is not quite clear at what temperature j-carbide forms during cooling. It is possible for j-carbide to form during cooling after rolling considering its continuous network along grain boundaries. However, if j-carbide particles can form during hot rolling at between 1100 and 1000 °C as reported by Rigaud et al. [20], it is quite possible for j-carbide particles to play the similar role to that
39
of NbC particles such that they would inhibit the recrystallization of ferrite during rolling. In fact, Baligidad also observed the formation of pancake-shaped ferrite grains in hot-rolled Fe–8.8Al–0.1C alloy in as-hot-rolled condition [9], although neither the detailed analyses on microstructural development nor the effect of subsequent warm/cold rolling and annealing were conducted. Although it is known that the formation of j-carbide particles along grain boundaries often adversely affects the rollability during subsequent TMTs [13], their presence can be favorably utilized to refine the grain size. To understand the microstructural evolution of Fe–8Al– 5Mn–0.1Nb–0.1C alloy, microstructural observation has been conducted on specimens subjected to subsequent TMTs, i.e. warm/cold rolling and annealing, as shown in Figure 4. Figure 4 also shows that the coarse needle-shaped j-carbide particles (Fig. 1c) in the ashot-rolled condition become elongated particles that form a banded structure along RD after warm rolling to 3 mm (85% reduction) at 650 °C (Fig. 4a). It can also be seen that some of j-carbide particles are cut and fractured by deformation. During annealing of the warm-rolled specimen at 850 °C for 15 min, the banded j-carbide particles become separated, forming discrete particles, similar to the spheroidization of pearlite lamellae observed in warm-rolled and annealed 5140 steel [21]. Recrystallization also occurs, forming a much finer grain structure (grain size: 25 lm) than that of as-hotrolled specimen (Fig. 4b). The distribution of j-carbide particles becomes quite uniform after final cold rolling to 1 mm (67% reduction) and annealing at 750 °C for 1 h (Fig. 4c) and there is a significant grain refinement as shown in Figures 2b and 4c. On the other hand, the grain sizes of the Fe–8Al–5Mn alloy after warm rolling and annealing (Fig. 4d) and final cold rolling and annealing (Fig. 2a) are 55 lm and 32 lm, respectively, which are much coarser than those of the Fe–8Al–5Mn–0.1Nb–0.1C alloy at the same processing conditions. It suggests that the presence of j-carbide particles in Fe–8Al–5Mn–0.1Nb–0.1C alloy plays a major role in refining its grain size by promoting recrystallization during subsequent TMTs after hot rolling. It is well known that recrystallization of deformed alloys can occur at grain boundaries, deformation bands and second-phase particles [22]. As shown in Figure 1, the Fe–8Al–5Mn–0.1Nb–0.1C alloy shows a much more elongated pancake-shaped grain structure than Fe-8Al5Mn alloy in as-hot-rolled condition, suggesting that there would be more deformation bands in the former than in the latter. These deformation bands could act as nucleation sites for recrystallization during warm rolling; however, the warm-rolled microstructure (Fig. 4a) does not show the occurrence of recrystallization, indicating that the amount of strain stored in as-hot-rolled condition is not sufficient to induce recrystallization at the warm-rolling temperature of 650 °C possibly due
Table 2. Room-temperature tensile properties of the alloys after final annealing. Alloy
YS (MPa)
UTS (MPa)
Uniform elongation (%)
Total elongation (%)
Fe–8Al–5Mn Fe–8Al–5Mn–0.1Nb–0.1C
484.0 599.6
592.2 694.5
16.9 16.5
35.9 30.7
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suggests that PSN is the dominant grain refinement mechanism for Fe–8Al–5Mn–0.1Nb–0.1C alloy since it has much finer grain size than the Fe–8Al–5Mn alloy, which does not contain any second-phase particles. In summary, the effects of 0.1Nb and 0.1C additions on the microstructure and tensile properties of Fe–8Al–5Mn ferritic lightweight steels have been characterized. Our results show that the microstructure of Fe–8Al–5Mn– 0.1Nb–0.1C alloy contains NbC and j-carbide particles within ferrite matrix and along grain boundaries. The microstructure of the Fe–8Al–5Mn–0.1Nb–0.1C alloy is much more elongated along RD than that of Fe–8Al– 5Mn alloy due to the effect of both NbC and j-carbide particles on inhibiting the recrystallization of ferrite during hot rolling. The Fe–8Al–5Mn–0.1Nb–0.1C alloy has a much finer grain size than the Fe–8Al–5Mn alloy after final annealing due to the enhanced recrystallization by PSN mechanism associated with j-carbide particles in the former and accordingly the former shows a better combination of tensile properties than the latter. Authors are grateful for the generous support of POSCO. Figure 4. Microstructure of the alloys after TMTs: (a) Fe–8Al–5Mn– 0.1Nb–0.1C after warm rolling (inset: magnified image of black rectangular area), (b) Fe–8Al–5Mn–0.1Nb-0.1C after annealing of warm-rolled sheet, (c) Fe–8Al–5Mn–0.1Nb–0.1C after final annealing of cold-rolled sheet, (d) Fe–8Al–5Mn after annealing of warm-rolled sheet, and IPF maps of cold-rolled Fe–8Al–5Mn–0.1Nb–0.1C (e) before heating, (f) after heating at 700 °C for 5 min showing the nucleation of new ferrite grains at j-carbide particles (arrowed), and (g) after heating at 700 °C for 15 min.
to the strain relief during cooling after hot rolling. It also indicates that the amount of strain received during warm rolling at 650 °C is not sufficient for dynamic recrystallization to occur. However, the warm-rolled specimens show the occurrence of static recrystallization after annealing at 850 °C, which is higher than the warm-rolling temperature (650 °C). On the other hand, the presence of j-carbide particles in the Fe–8Al–5Mn– 0.1Nb–0.1C alloy would play an important role in grain refinement by the operation of particle-stimulated nucleation (PSN) mechanism [22]. An in situ heating EBSD study of the cold-rolled Fe–8Al–5Mn–0.1Nb– 0.1C alloy shows that recrystallization readily occurs along j-carbide particles with the formation of new grains having different orientations from that of the matrix (Fig 4e–g). j-carbide particles can also inhibit the growth of newly formed grains by pinning grain boundaries as can be seen in Figure 4c, which shows the location of j-carbide particles mostly along grain boundaries. On the other hand, NbC particles would play a minor role in inhibiting grain growth considering their much smaller volume fraction than that of j-carbide particles (0.2% vs. 2.3%). As mentioned previously, the grain sizes of the Fe–8Al–5Mn–0.1Nb– 0.1C alloy in the warm-rolled and annealed condition and the cold-rolled and annealed condition are 25 and 10 lm, respectively, which are much finer that those of the Fe–8Al–5Mn alloy under the same TMT conditions (55 lm in the warm-rolled and annealed condition and 32 lm in the cold-rolled and annealed condition). It
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