Microstructural changes during superplastic deformation of Fe 24Cr 7Ni 3Mo 0.14N duplex stainless steel

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1 Materials Science and Engineering A266 (1999) Microstructural changes during superplastic deformation of Fe 24Cr 7Ni 3Mo 0.14N duplex stainless steel Young S. Han, Soon H. Hong * Department of Materials Science and Engineering, Korea Ad anced Institute of Science and Technology, Kusung-dong, Yusung-gu, Taejon, , South Korea Received 6 October 1998 Abstract The superplasticity of Fe 24Cr 7Ni 3Mo 0.14N duplex stainless steel after solution treatment at 1350 C followed by 90% cold rolling was investigated at 850 C with a strain rate ranging from 10 3 to 10 1 s 1. The microstructure of duplex stainless steel consists of matrix phase having low angle boundaries and phase as second phase particles before the deformation at 850 C. The constituent phases in the duplex stainless steel were found to be changed following through phase transformation during the deformation at 850 C. A maximum elongation of 750% was obtained at 850 C with strain rate of s 1. The low angle grain boundaries were changed into high angle grain boundaries by dynamic recrystallization of phase at an early stage of deformation. The dislocation density within matrix grains was low and a significant strain-induced grain growth was observed during the deformation. The misorientation angles between the neighboring grains increased with increasing strain; thus the low angle grain boundaries were transformed into high angle grain boundaries suitable for sliding by the dynamic recrystallization during the deformation. The grain boundary sliding assisted by dynamic recrystallization is considered a controlling mechanism for superplastic deformation at 850 C Elsevier Science S.A. All rights reserved. Keywords: Duplex stainless steel; Dynamic recrystallization; Grain boundary sliding; Phase; Phase; Superplasticity 1. Introduction The duplex stainless steels are defined as a family of stainless steels consisting of a two phase aggregated microstructure of -ferrite and -austenite. The duplex stainless steels have an attractive combination of high mechanical properties with excellent corrosion resistance, and are suitable for constructional and petrochemical applications [1 3]. Furthermore, the duplex stainless steels having a fine grained microstructure showing superplastic behavior since the grain growth is effectively suppressed at high temperature due to the two phase aggregated microstructure. The first investigation on the superplasticity of duplex stainless steels was reported by Hayden et al. [4,5]. They investigated the superplastic behavior of hot * Corresponding author. Tel.: ; fax: address: shhong@sorak.kaist.ac.kr (S.H. Hong) rolled duplex stainless steels with various Fe Cr Ni ternary compositions at temperatures near 1273K, and reported an elongation of 500% in a commercial 25Cr 6.5Ni 0.6Ti duplex stainless steel. Gibson et al. [6] reported that the superplasticity of 25Cr 6.5Ni 0.6Ti duplex stainless steel was enhanced by additional cold working of hot rolled sheets. Smith et al. [7] also reported elongation above 1000% in cold rolled 25Cr 6.5Ni 0.6Ti duplex stainless steel. Hayden and Smith [5 7] suggested that the grain boundary sliding is the dominant mechanism for superplasticity in duplex stainless steel. Maehara [8 11] recently reported an elongation above 1000% in a wide temperature range ( K) by controlled thermomechanical treatment in 25Cr 7Ni 3Mo 0.14N duplex stainless steel. Hayden and Gibson [4 6] obtained a fine grained duplex microstructure through recrystallization during thermomechanical treatment. However, Maehara [7 9] obtained a fine grained duplex microstructure through the precipitation of the second phase at test tempera /99/$ - see front matter 1999 Elsevier Science S.A. All rights reserved. PII: S (99)

2 Y.S. Han, S.H. Hong / Materials Science and Engineering A266 (1999) ture after solution treatment above 1250 C followed by cold rolling to 50% reduction. It is suggested [10,11] that the dynamic recrystallization of the softer phase in a duplex microstructure, which occurs continuously during deformation, could be the dominant mechanism for superplasticity at temperatures in the range C [10,11]. This conclusion emphasizes that the deformation mechanism for superplasticity in duplex stainless steel is not the grain boundary sliding which has generally been considered the dominant mechanism for superplastic deformation. On the other hand, Tsuzaki [12] suggested that grain boundary sliding is the dominant mechanism for superplasticity in duplex stainless steel, and that the role of dynamic recrystallization is to keep the grain size fine, suitable for grain boundary sliding. Duplex stainless steels exhibited an single phase just below the liquidus temperature, but the phase transformation of single phase into + duplex phases occurred below 1320 C. Since the volume fraction of phase increases with decreasing temperature below 1320 C, the precipitation of fine phase in matrix phase improved the hot ductility at temperatures between 950 and 1100 C. The phase is known to be formed through the eutectoid decomposition of phase into and phase at temperatures between 700 and 950 C. It was reported that the precipitation of phase played an important role in improving the hot ductility below 950 C [9]. However, most of the previous investigations on superplasticity in duplex stainless were focused on the + duplex microstructure near 1000 C, while the superplasticity of the + duplex microstructure below 950 C was not well understood. In this study, the superplasticity of Fe 24Cr 7Ni 3Mo 0.14N duplex stainless steel was investigated at 850 C. The primary purpose of this study is to investigate the microsturctural change during superplastic deformation and to understand the controlling mechanism for superplasticity in duplex stainless steel having + duplex microstructure. 2. Experimental procedures Fe 24Cr 7Ni 3Mo 0.14N duplex stainless was induction melted and cast into ingots of 7 kg. The composition of the melted alloy, analyzed by atomic absorption spectrometry, is shown in Table 1. The cast ingots were hot rolled into a 20-mm thick plate and Table 1 Chemical composition of melted duplex stainless steel (wt.%) Fe Cr Mo Bal Ni N 0.14 then solution treated at 1350 C for 30 min, followed by water quenching. The 20-mm thick hot rolled plates were cold rolled into 2-mm thick sheets by 90% reduction and the tensile specimens with thickness of 2 mm and gauge length of 5 mm were machined from the cold rolled plates. The tensile tests were performed with a constant initial strain rate ranging from 10 3 s 1 to 10 1 s 1 at 850 C. The specimens were held for 5 min at 850 C before the loading and were tested in an argon atmosphere to minimize surface oxidation. In order to examine the microstructural change in specimens during the high temperature tensile deformation, the tensile tests were interrupted at a fixed elongation and the specimens were rapidly cooled down to room temperature by flowing argon gas. X-ray diffraction analyses were performed using Cu K radiation in order to analyze the constituent phases at 850 C. The microstructures of the specimens were examined to analyze the volume fractions and grain size of constituent phases. Scanning electron micrographs were taken after electrolytic etching of the polished surface in a saturated oxalic acid solution. Thin foils for the transmission electron microscope were prepared by jet polishing at 50 C in a solution of 20% perchloric acid and 80% ethanol. Kikuchi patterns were obtained to measure the misorientation angle between adjacent grains. 3. Results and discussion Fig. 1(a) shows the variation of tensile elongation of duplex stainless steel with varying the initial strain rate at 850 C. A peak elongation of 750% was obtained at an initial strain rate of s 1. Fig. 1(b) shows the variation of peak stress of duplex stainless steel with varying initial strain rate at 850 C. The strain rate sensitivity was about 0.37, as shown in Fig. 1(b). The activation energy for superplastic deformation of duplex stainless steel was determined from the modulus compensated stress data with varying inverse of temperature from 850 C to 1000 C. The activation energy was 306 kj mol 1. The activation energy for superplastic deformation of the duplex stainless steel was similar to the activation energy for lattice diffusion of iron in and phase measured as 251 and 280 kj mol 1, respectively. The microstructures of duplex stainless steel consisted of a single phase just below the liquidus temperature, but the phase transformation of single phase into + duplex phase occurred below 1320 C. The phase is known to be formed through the eutectoid decomposition of phase into + phases at temperatures between 700 and 950 C [13,14]. The constituent phases of Fe 24Cr 7Ni 3Mo 0.14N du-

3 278 Y.S. Han, S.H. Hong / Materials Science and Engineering A266 (1999) phases after 5 min holding at 850 C just before the tensile test. This result showed that the phase transformation of duplex stainless steel following the sequence of was completed after 5 min holding at 850 C. When the duplex stainless steel was deformed up to 400% elongation at 850 C, the microstructure of also showed + duplex phases, shown in Fig. 2(b). The volume fraction of the phase is about 30% in both micrographs (Fig. 2(a) and (b)). Comparing the microstructures before and after deformation, a significant strain-induced grain growth was observed during deformation at 850 C. The grain size of the phase increases from 0.8 to 1.7 m Fig. 1. The variation of elongation and peak stress with varying initial strain rate of Fe 24Cr 7Ni 3Mo 0.14N duplex stainless steel at 850 C: (a) elongation; (b) peak stress. plex stainless steel changed following the sequence with increasing holding time at 850 C [15,16]. Fig. 2(a, b) shows the microstructures observed before and after the tensile deformation up to 400% at 850 C with an initial strain rate of s 1 in Fe 24Cr 7Ni 3Mo 0.14N duplex stainless steel. The duplex stainless steel as shown in Fig. 2(a) exhibited Fig. 2. The comparison of scanning electron micrographs of duplex stainless steel before and after tensile deformation of 400% at 850 C with an initial strain rate of s 1 : (a) undeformed, (b) deformed.

4 Y.S. Han, S.H. Hong / Materials Science and Engineering A266 (1999) Fig. 3. Microstructure and Kikuchi patterns of duplex stainless steel before deformation. (a) Transmission electron micrograph before deformation; the specimen was held for 5 min at 850 C. (b) Kikuchi pattern taken from the grain marked A. (c) Kikuchi pattern taken from the grain marked B. (d) Kikuchi pattern taken from the grain marked C. during deformation at 850 C. However, the matrix phase boundaries are unidentified in both SEM micrographs (Fig. 2(a) and (b)). The microstructures before deformation were observed by transmission electron microscopy and are shown in Fig. 3. Fig. 3(a) shows a transmission electron micrograph of the duplex stainless steel after 5 min holding at 850 C just before the tensile test. The grains of and phases are equiaxed with average grain sizes of about 1 m. Kikuchi patterns in Fig. 3(b d) were taken from the three grains marked A, B and C in Fig. 3(a). The misorientation angle between two neighboring grains marked A and B is 1.3 and the misorientation angle between two neighboring grains marked B and C is 0.8. This result clearly indicates that the / grain boundaries in the + duplex structure are the low angle boundaries before the deformation at 850 C. The schematic figures shown in Fig. 4 indicate the microstructural change in duplex stainless steel after solution treatment at 1350 C and 90% cold rolling. The duplex stainless steel showed a single phase with a grain size of about 100 m after solution treatment at 1350 C. The and phase were precipitated in matrix phase through the eutectoid decomposition of + as shown in Fig. 4(b). The constituent phases were finally transformed into + duplex structure as shown in Fig. 4(c). The grain size of phase was 0.4 m, which is much smaller than that of the phase measured at about 100 m. The phase that was

5 280 Y.S. Han, S.H. Hong / Materials Science and Engineering A266 (1999) precipitated in the grain should have the same orientation relationships with the matrix. Thus the / boundaries in the + duplex structure as shown in Fig. 4(c) are mainly the low angle boundaries before the tensile deformation. The grain boundaries between adjacent matrix grains should be high angle boundaries because grain boundary sliding is generally the predominant mode of deformation during superplastic flow [17 19]. Since the grain boundary sliding is difficult to proceed at low angle boundaries, exhibiting superplasticity in + duplex stainless steel having low angle grain boundaries is a special phenomenon. The transmission electron micrograph of duplex stainless steel shown in Fig. 5(a) was observed in an identical specimen deformed up to 20% elongation. Kikuchi patterns shown in Fig. 5(b d) were taken from Fig. 4. Schematic diagrams representing the microstructural changes in Fe 25Cr 7Ni 3Mo 0.14N duplex stainless steel during deformation at 850 C: (a) solution treated at 1350 C and cold rolled; (b) at an early stage in and phase precipitation occurred; (c) held 5 min at 850 C just before deformation.

6 Y.S. Han, S.H. Hong / Materials Science and Engineering A266 (1999) Fig. 5. Microstructure and Kikuchi patterns of duplex stainless steel after deformation up to 20%: (a) transmission electron micrograph after deformation up to 20%; (b) Kikuchi pattern taken from the grain marked A; (c) Kikuchi pattern taken from the grain marked B; (d) Kikuchi pattern taken from the grain marked C. the three grains marked A, B and C in Fig. 5(a). The misorientation angle between two neighboring grains marked A and B is 5.7 and between two neighboring grains marked A and C is Consequently, the misorientation between adjacent grains was increasing and some / boundaries changed their orientation from low angles to high angles with increasing strain up to 20%. Fig. 6(a) shows a transmission electron micrograph of duplex stainless steel after deformation up to 50% elongation. Dislocations are scarcely observed in the grains marked A and B as shown in Fig. 6(a, b). Kikuchi patterns in Fig. 6(b d) were taken from the three grains marked A, B and C in Fig. 6(a). The misorientation angle between two neighboring grains marked A and B is 11.0 and between two neighboring grains marked A and C is This result shows that the observed grain boundaries as shown in Fig. 6(a) are high angle boundaries. The misorientation angles between the neighboring grains increased with increasing the strain, thus the low angle grain boundaries were transformed into high angle grain boundaries suitable for sliding during the deformation. Fig. 7 shows the variation of measured misorientation angle of adjacent grains with increasing the strain. Most / boundaries are low angle grain boundaries before deformation and the large amount of low angle grain boundaries were changed into high angle grain boundaries with increasing strain up to 20%. Finally, all the low angle grain boundaries were changed into high angle grain boundaries with increasing strain up to

7 282 Y.S. Han, S.H. Hong / Materials Science and Engineering A266 (1999) %. The recrystallization of the matrix is completed in the early stage of deformation within 50%. It is concluded that the dynamic recrystallization transformed the microstructure into fine equiaxed duplex structure having high angle / boundaries suitable for the grain boundary sliding. The conventional recrystallization proceeds through nucleation and growth mechanism. During the conventional recrystallization, subgrains are formed by recovery and particular subgrains among a number of subgrains can grow as recrystallized grains or pre-existing high angle grain boundaries migrate into deformed matrix and dislocation free recrystallized grains are formed. However, in our results, the subgrain structure before the tensile deformation were changed into a fine recrystallized grain structure consisting of high angle boundaries in the early stage of the deformation. Furthermore, misorientations between adjacent grains continuously increased with increasing strain. The conventional features occurred during recrystallization were not observed in our results of microstructural analyses. Similar microstructural changes to our results have been reported in Al 6%Cu 0.4%Zr (Supral 100) [20,21] and Al 10Mg alloys [22] during superplastic deformation. These aluminum alloys were in recovery condition before the tensile deformation. The subgrain structure before the deformation was changed to a fine recrystallized grain structure consisting of high angle boundaries during deformation. Furthermore misorientations between adjacent grains continuously increased Fig. 6. Microstructure and Kikuchi patterns of duplex stainless steel after deformation up to 50%: (a) TEM micrograph after deformation up to 50%; (b) Kikuchi pattern taken from the grain marked A; (c) Kikuchi pattern taken from the grain marked B; (d) Kikuchi pattern taken from the grain marked C.

8 Y.S. Han, S.H. Hong / Materials Science and Engineering A266 (1999) phase increased rapidly with increasing time when deformed at 850 C, while the grain size of phase was kept almost constant with increasing time without deformation at 850 C. The strain-induced grain growth was pointed out as one of the typical features of the microstructural change during superplastic deformation controlled by the grain boundary sliding [30]. Especially, it had been reported that the strain-induced grain growth was obtained through the grain boundary slid- Fig. 7. The variation of misorientation angle between adjacent grains with varying tensile elongation. with an increase in strain. This type of recrystallization has been called continuous recrystallization, in contrast to conventional recrystallization [22]. It has been suggested [22] that continuous recrystallization is responsible for the development of the fine grained microstructure during the deformation. Fig. 8(a, b) show transmission electron micrographs of duplex stainless steel deformed up to elongation of 200% and 750%, respectively. The dislocation density was kept very low and the fine equiaxed grains were maintained up to a large elongation of 750%. This is indirect evidence demonstrating the occurrence of grain boundary sliding [23,24] These results also indicate that the deformation proceeded mainly by grain boundary sliding assisted by dynamic recrystallization [25 29]. Maehara [10] proposed that the dynamic recrystallization of the matrix, which occurs repeatedly during the deformation, is the mechanism of superplasticity of the duplex stainless steel in the + duplex microstructure at 850 C. He reported that the work hardened structure having a high dislocation density and recrystallized grains were observed in the deformed microstructure. However, the work hardened structure necessary for the recrystallization was not observed during the deformation up to 750% elongation in duplex stainless steel as shown from Figs Fig. 9 shows the change in average grain size of phase measured from scanning electron micrographs with varying the holding time with deformation and without deformation. Fig. 9 compares the static grain growth with the dynamic grain growth of phase in duplex stainless steel at 850 C. The grain size of Fig. 8. Transmission electron micrographs of duplex stainless steel after deformation at 850 C: (a) deformed 200% at 850 C; (b) deformed 750%.

9 284 Y.S. Han, S.H. Hong / Materials Science and Engineering A266 (1999) mechanism for superplastic deformation and the role of the dynamic recrystallization is to transform the low angle grain boundaries into high angle grain boundaries suitable for sliding. The misorientation angles between the neighboring grains increased with increasing strain; thus the low angle grain boundaries were transformed into high angle grain boundaries suitable for sliding by dynamic recrystallization during the deformation. The grain boundary sliding assisted by dynamic recrystallization is considered to be the controlling mechanism for superplastic deformation at 850 C. References Fig. 9. The variation of the grain size of phase in duplex stainless steel with time at 850 C with and without deformation under an initial strain rate of s 1. ing in microduplex structured materials during superplasic deformation [31]. Therefore, the extensive grain growth during the deformation suggests that the deformation proceeds mainly by grain boundary sliding. 4. Conclusions The superplastic deformation mechanism of Fe 24Cr 7Ni 3Mo 0.14N duplex stainless steel after solution treatment at 1350 C and followed by 90% cold rolling was investigated at 850 C and lead to the following conclusions. The constituent phases in the duplex stainless steel were found to be changed following a sequence of through phase transformation with increasing time at 850 C. The microstructure of duplex stainless steel consists of matrix phase having low angle boundaries and phase as second phase particles before the deformation at 850 C. The low angle grain boundaries between neighboring grains were continuously changed into high angle grain boundaries with increasing the strain by dynamic recrystallization of phase at an early stage of deformation. The dislocation density within matrix grains was low and a significant strain-induced grain growth was observed during the deformation. The grain boundary sliding is considered to be the controlling [1] M.A. Strecher, Met. Prog. 128 (1985) 31. [2] R.M. Davison, J.K. Redmond, Mater. Perform. 29 (1990) 57. [3] J.O. Nilsson, Mater. Sci. Technol. 8 (1992) 685. [4] H.W. Hayden, J.H. Brophy, Trans. ASM 61 (1967) 542. [5] H.W. Hayden, S. Floreen, P.G. Goodell, Met. Trans. 3 (1972) 833. [6] R.C. Gibson, H.W. Hayden, J.H. Brophy, Trans. ASM 61 (1968) 85. [7] C.I. Smith, B. Norgate, N. Ridley, Met. Sci. 10 (1976) 182. [8] Y. Maehara, Trans. ISIJ 25 (1985) 69. [9] Y. Maehara, Trans. ISIJ 27 (1987) 705. [10] Y. Maehara, Y. Ohmori, Met. Trans. A 18 (1987) 663. [11] Y. Maehara, Met. Trans. A 22 (1991) [12] K. Tsuzaki, H. Matsuyama, M. Nagao, T. Maki, Mater. Trans. JIM 31 (1990) 983. [13] Y. Maehara, N. Fujino, T. Kunitake, Trans. ISIJ 23 (1983) 240. [14] Y. Maehara, Y. Ohmori, T. Kunitake, Met. Technol. 10 (1983) 296. [15] Y.S. Han and S.H. Hong, Proceedings of Conference on Superplasticity and Superplastic Forming, Las Vegas, February 1995, 124th TMS Annual Meeting, p [16] Y.S. Han, S.H. Hong, Scripta Mater. 36 (1997) 577. [17] T.G. Nieh, J. Wadsworth, O.D. Sherby, Superplasticity in Metals and Ceramics, Cambridge University Press, Cambridge, 1997, p. 26. [18] H. Kokawa, T. Watanabe, S. Karashima, Philos. Mag. A 44 (1981) [19] J. Wadsworth, A.R. Pelton, Scripta Metall. 18 (1984) 387. [20] R. Grims, M.J. Stowell, B.M. Watts, Met. Technol. 3 (1976) 154. [21] R.H. Bricknell, J.W. Edingtion, Metall. Trans. A 10 (1979) [22] S.J. Hales, T.R. McNelly, Acta Metall. 36 (1988) [23] J. Pilling, N. Ridley, Superplasticity in Crystalline Solids, The Institute of Metals, 1989, p. 4. [24] O.D. Sherby, J. Wadsworth, Proc. Mater. Sci. 33 (1989) 169. [25] A.H. Chokshi, A.K. Mukherjee, T.G. Langdon, Mater. Sci. Eng. R 10 (1993) 237. [26] M.F. Ashby, R.A. Verrall, Acta Metall. 21 (1973) 149. [27] R.C. Gifkins, Metall. Trans. A 7 (1976) [28] J.H. Gittus, J. Eng. Mater. Technol. 99 (1977) 244. [29] J.R. Spingarn, W.D. Nix, Acta Metall. 26 (1978) [30] D.S. Wilkinson, J. Wadsworth, Acta Metall. 32 (1984) [31] K. Holm, J.D. Embury, G.R. Purdy, Acta Metall. 25 (1977)