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Influence of annealing temperature on deformation behavior of 329LA lean duplex stainless steel
Materials Science & Engineering A 679 (2017) 531–537
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
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Influence of annealing temperature on deformation behavior of 329LA lean duplex stainless steel ⁎
Nithi Saenarjhana, Jee-Hyun Kanga, Soo Chan Leeb, Sung-Joon Kima, a b
Graduate Institute of Ferrous Technology (GIFT), Pohang University of Science and Technology (POSTECH), Pohang 37673, Republic of Korea Technical Research Laboratories, POSCO, Pohang 37859, Republic of Korea
A R T I C L E I N F O
A BS T RAC T
Keywords: Duplex stainless steel Transformation induced plasticity Mechanical properties Electron microscopy
Tensile behavior of a 329LA lean duplex stainless steel at room temperature is investigated after annealing at different temperatures. The austenite composition and thereby, its stability depend on the annealing temperature. For recrystallized alloys, austenite stability increases with the annealing temperature. Therefore, martensitic transformation is effectively suppressed during cooling and deformation, which results in low tensile strength and strain hardening rate, when the steel is annealed at higher temperature. Accordingly, transformation-induced plasticity is more pronounced in alloys annealed at lower temperature. During deformation, α′-martensite forms with a blocky morphology with the absence of the deformation bands indexed as an hcp phase. Thus, it is suggested that austenite directly transforms to α′-martensite during deformation rather than via the deformation bands due to a relatively high driving force for α′-martensite transformation in the present alloy.
1. Introduction Duplex stainless steels (DSSs) exhibit excellent combination of corrosion resistance and mechanical properties, which has enabled their development as structural components in chemical and pulp industries as well as on/offshore applications. In order to achieve the requirements for the application, conventional DSSs contain considerable amounts of expensive alloying elements such as Cr, Ni, and Mo [1]. Recently, efforts have been made to develop lean DSSs with lower amount of the expensive alloying elements. In this regard, Ni and Mo were replaced with Mn and N [1]. Mn plays an important role in both stabilizing austenite and increasing the solubility of interstitial elements such as C and N. N also acts as a strong austenite stabilizer and has beneficial effects on the strength and pitting corrosion resistance [2–5]. DSSs contain two phases, ferrite and austenite, the equilibrium fractions of which depend on annealing temperature. Accordingly, there is a possibility to alter the partitioning of alloying elements into the phases by annealing at different temperatures. Changing the composition of austenite affects its stability, and therefore, determines if it transforms into martensite during quenching as well as deformation. The austenite stability can be represented by the thermodynamic driving force, which is calculated as the difference in Gibbs free energies of fcc and bcc structures. The driving force can be further
⁎
related with martensite start temperature [6] and stacking fault energy [7]. Especially, stacking fault energy is often connected to the deformation mechanisms. It is often reported that austenite in lean DSSs experiences deformation-induced martensitic transformation (DIMT) [8–10]. DIMT plays an important role in mechanical properties by enhancing tensile strength and ductility, which is known as transformation-induced plasticity (TRIP). TRIP results in an excellent combination of strength and elongation over 1 GPa-60% in a DSS [11]. Further increase in stacking fault energy by adjusting the chemical composition [12] or deformation conditions [13–15] can lead to different deformation mechanisms such as twinning and pure dislocation glide rather than DIMT. Hence, the present work is conducted to study the influence of different annealing temperatures on the mechanical behavior and deformation mechanism of a 329LA lean DSS. Tensile tests were conducted along with microstructure analysis, and deformation mechanism is discussed with the assistance of thermodynamic calculations. 2. Experimental procedures The chemical composition of a 329LA lean DSS used in this study is Fe-0.17N-0.02C-20.5Cr-2.0Ni-1.8Mn-0.7Cu-0.6Mo-0.5Si (wt%). A 4.0 mm thick hot-rolled plate was cold- rolled to 1.0 mm in thickness.
Corresponding author. E-mail address: [email protected] (S.-J. Kim).
http://dx.doi.org/10.1016/j.msea.2016.10.062 Received 10 August 2016; Received in revised form 9 October 2016; Accepted 12 October 2016 Available online 21 October 2016 0921-5093/ © 2016 Elsevier B.V. All rights reserved.
Materials Science & Engineering A 679 (2017) 531–537
N. Saenarjhan et al.
structures, which are ferrite and martensite. On the other hand, the annealing temperatures higher than 900 °C, clearly separate the fcc structure from the bcc structure. Moreover, some bcc structure is present inside austenite grains which implies that α′-martensite has formed during water quenching due to low austenite stability. Note that there is a possibility that some α′-martensite has formed during sample preparation. Therefore, the martensite fraction was not obtained from the phase maps but by a ferrite detector. The measured percentages of magnetic phases after annealing are shown in Fig. 2a with respect to the annealing temperature. Calculated equilibrium percentages of ferrite at different temperatures are shown together for comparison. The result shows that while the equilibrium and measured values are similar for the annealing temperature above 1100 °C, the differences between these values are significant at lower temperatures. The differences indicate that some austenite transforms into α′-martensite during quenching which is consistent with the EBSD results. In order to estimate the percentage of α′-martensite, these differences were multiplied with a correction factor, 1.71 to consider the different magnetic property between ferrite and α′-martensite [16]. Then, the remaining portion other than ferrite and α′-martensite is regarded as austenite, and the estimated phase percentages in the asannealed alloys are shown in Fig. 2b. Note that such estimation is not applicable to unrecrystallized specimens since the equilibrium austenite fraction has not been reached during annealing. Since the fraction of the measured magnetic phases and that of equilibrium ferrite were similar for the alloy annealed at 1200 °C, it seems that no α′-martensite has formed during quenching, and the ferrite percentage was taken as 62%. From Fig. 2, it is concluded that the austenite percentage decreases with respect to annealing temperature over 900 °C; however, the resulting austenite percentage after quenching is approximately 40% for all alloys because less austenite transforms into α′-martensite during quenching.
The plate was machined into subsize tensile specimens with gauge length and width, 33 mm and 10 mm, respectively. The specimens were then subjected to annealing at different temperatures from 800 to 1200 °C for 10 min with the temperature interval of 50 °C, which was followed by quenching in water. Equilibrium percentages and compositions of austenite at different annealing temperatures were calculated with ThermoCalc Version 4.1 with TCFE7 database. Based on the obtained compositions of austenite, molar Gibbs free energies for fcc, hcp, and bcc are calculated to be Gmfcc , Gmhcp , and Gmbcc . Then, the driving forces of the transformations from austenite to an hcp phase and α′-martensite are obtained with F fcc → hcp = −(Gmhcp − Gmfcc ) and F fcc → bcc = −(Gmbcc − Gmfcc ), respectively. Tensile tests were carried out at room temperature with a constant crosshead speed of 1 mm min−1. Interrupted tensile tests were performed to study microstructure evolution. The as-annealed microstructures were analyzed by electron backscattered diffraction (EBSD) with a field emission scanning electron microscope, JEOL JSM7100F. Moreover, N concentrations in austenite were obtained with electron probe microanalysis (EPMA) by JEOL JXA-8530F. In order to identify different phases, X-ray diffraction (XRD) was analyzed by a Bruker D8 Advance with a Cu source. Volume percentages of magnetic phases were measured by a Fisher FMP30 ferrite detector. The deformed structure of austenite was investigated by a transmission electron microscope (TEM), JEOL JEM 2100. Specimens for EBSD were final-polished with colloidal silica after grinding with SiC papers up to #2000. TEM specimens were prepared by grinding and polishing a thin foil down to < 100 µm thick, which were followed by jet polishing at room temperature with a solution of 12% perchloric acid in acetic acid.
3. Results and discussion 3.1. As-annealed microstructure
3.2. Tensile properties Phase maps of undeformed microstructures after annealing at 800, 950, 1000, and 1200 °C obtained by EBSD are shown in Fig. 1. It is clear that both bcc and fcc structures become coarser with increasing annealing temperature especially from 950 °C. For annealing temperature up to 900 °C, most of the microstructure consists of the bcc
The engineering stress-strain curves at room temperature are shown in Fig. 3. It is observed that the tensile behavior of 329LA is sensitive to the annealing temperature. Yield and tensile strengths decrease gradually as the annealing temperature increases, while
Fig. 1. Electron backscattered diffraction (EBSD) phase maps of as-annealed microstructures after annealing at (a) 800 °C, (b) 950 °C, (c) 1000 °C and (d) 1200 °C. Note that the scale of (d) is different from others to reveal large microstructures.
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True stress or strain hardening rate (MPa)
Engineering stress (MPa)
Fig. 2. (a) Percentages of measured magnetic phases after annealing and those of equilibrium ferrite with respect to annealing temperature. (b) Estimated phase percentages in quenched alloys after annealing at 950–1200 °C.
800 600 400 o
800 C o 850 C o 900 C
200 0
0
o
950 C o 1000 C o 1050 C
o
1100 C o 1150 C o 1200 C
10 20 30 40 50 Engineering strain (%)
5000 4000
o
1000 C o 1100 C o 1200 C
3000 2000 1000 0
60
o
950 C o 1050 C o 1150 C
0.0
0.1
0.2 0.3 True strain
0.4
0.5
Fig. 4. True stress and strain hardening rate with respect to true strain after annealing at 950–1200 °C.
Fig. 3. Engineering stress-strain curves after annealing at different temperatures.
and maximum points of SHR. Moreover, the peak value of SHR is reduced. Therefore, when the annealing temperature increases, DIMT not only starts and ends later, but also takes place less actively. As a result, TRIP effect is reduced with respect to the annealing temperature and the tensile strength decreases while ductility increases accordingly (Fig. 3).
elongation is enhanced. Improvement in ductility over 50% can be achieved when the annealing temperature surpasses 1000 °C. In addition, all curves show a sigmoidal shape which indicates the TRIP due to DIMT as will be shown later. Based on the as-annealed microstructures (Fig. 1), the tensile curves can be categorized into two groups: (i) the curves of unrecrystallized microstructures obtained after annealing at 800–900 °C, and (ii) the curves of fully recrystallized microstructures after annealing at 950–1200 °C. Since the microstructures are complicated in the first group with unrecrystallized martensite and austenite which are yet to reach equilibrium, the main focus will be on the second group. The true stress-strain curves as well as corresponding strain hardening rates (SHRs) are plotted in Fig. 4. The yield strength tends to decrease with respect to the annealing temperature; this is attributed to the combined effects of change in both grain size (Fig. 1) and the phase fraction (Fig. 2b). Since martensite is a hard phase and ferrite is a soft phase, the yield strength is reduced with the decrease in martensite fraction and the increase in ferrite fraction as the alloy is annealed at higher temperature. The SHRs of all alloys decrease to minimum points and then increase with straining. The minimum points are caused by the inflection points in tensile curves (Fig. 3), which are known to be the activation points of DIMT [17,18]. DIMT gradually introduces martensite into microstructure and effectively enhances SHR. When the SHR reaches its maximum, the martensite fraction saturates and dynamic recovery of dislocations reduces SHR subsequently. It is obvious that increasing the annealing temperature delays the minimum
3.3. Microstructure evolution Fig. 5 shows the percentage of α′-martensite which formed during deformation for the specimens annealed at different temperatures. The values were obtained by taking the measured percentage difference of magnetic phases between undeformed and failed specimens and multiplying it by a correction factor, 1.71 [16]. Additionally, Fig. 5 includes the values which are normalized by the percentage of austenite in the as-annealed alloys. The trends of normalized and unnormalized percentages are similar due to similar austenite fraction in as-annealed samples regardless of the annealing temperature (Fig. 2b). Fig. 5 suggests that α′-martensite formed during deformation in all specimens except that annealed at 1200 °C. The percentage of deformationinduced martensite (DIM) increases as annealing temperature decreases. It proves that DIMT plays an important role in strain hardening; the higher peak in SHR is available with annealing at lower temperature (Fig. 4), since a higher degree of DIMT is achieved (Fig. 5). The microstructure evolution of the alloy annealed at 1000 °C is investigated in detail with interrupted tensile testing. XRD spectra in 533
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DIM percntage (%)
70
grains as shown in Fig. 7b. Although the deformation bands resemble deformation twins, they are not twins since twin diffraction spots do not appear. No hexagonally indexed deformation band is found, which is consistent with the XRD results (Fig. 6a). Moreover, deformationinduced α′-martensite is not observed, which is attributed to the low percentage less than 2% (Fig. 6b). At the strain of 30% (Fig. 8a), α′martensite often appears in austenite with Nishiyama-Wasserman (NW) ( < 110 > γ// < 001 > α′) orientation relationship. For specimen deformed up to failure (Fig. 8b), a large number of dislocations and stacking faults were observed and both of N-W orientation relationship and Kurdjumov-Sachs (K-S) orientation relationship of < 110 > γ // < 111 > α′ were attained between austenite and α′-martensite. The TEM observation of 1000 °C annealed specimen reveals the morphologies of α′-martensite to be big and bulky rather than needlelike and sharp (Fig. 8a). The morphology of α′-martensite varies depending on the transformation sequence, while the shape of hexagonally indexed deformation band is mostly reported in the form of a needle or sharp band [9–12,16,20–22]. When DIMT follows austenite (γ)→hexagonally indexed deformation band→α′, a small α′-martensite is usually revealed at the intersection of the hexagonally indexed deformation bands [13,16,20,22]. On the other hand, α′-martensite is blocky if the direct transformation, γ→α′, takes place [22,25]. Hence, it is suggested that DIMT occurs mainly via γ→α′ in the present alloy based on the observation of the blocky α′-martensite. This is consistent with the absence of hcp peaks in Fig. 6a. The microstructure evolution during deformation for the alloy annealed at 1200 °C is shown in Fig. 9. Before deformation, an austenite grain contains considerable densities of stacking faults and sharp thin dislocation bands; however, no quenched α′-martensite is found (Fig. 9a). At the strain of 30%, α′-martensite is observed on a rare occasion, and mostly only large densities of dislocations and stacking faults are observed without any extra diffraction spots (Fig. 9b). Throughout the deformation until failure, only dislocated austenite was mostly observed and no clear evidence of deformationinduced martensite was found (Fig. 9c). It is important to note that a high density of sharp thin bands which are observed at every stage do not show any twin spots. According to Welsch et al. [19], these bands are called as dynamic slip bands, which can improve strain hardening. Thus, the SHR plateau in the specimen annealed at 1200 °C at the true strain of 0.18–0.35 (Fig. 4) may be the result of the dynamic slip band formation in the present alloy.
Unnormalized Normalized
60 50 40 30 20 10 0 800
900
1000
1100
1200
Fig. 5. Measured percentages of α′-martensite formed during deformation in failed specimens with respect to annealing temperature. The red markers represent the values normalized with the initial austenite percentages.
Fig. 6a clearly show that the fraction of austenite decreases while that of a bcc structure increases according to the strain. The result is consistent with the magnetic phase measurement presented in Fig. 6b. However, additional diffraction peaks did not appear in the present alloy although an hcp phase is sometimes detected during deformation [13,16,20,22]. Since it has been suggested that the deformation bands which are clusterings of stacking faults can appear as the hcp structure in XRD spectra and EBSD phase maps [23,24], such structure is called as a hexagonally indexed deformation band in the present work. Accordingly, the absence of the hcp peaks in Fig. 6a suggests that the formation of the deformation bands is not a pre-requisite for transforming austenite into α′-martensite. Moreover, DIMT is significantly enhanced between the engineering strains of 10% and 30%. These correspond to the true strains equal to 0.095 and 0.262, respectively, where the SHR reaches its peak from the minimum point (Fig. 4). Thus, it demonstrates that DIMT effectively improves strain hardening. Figs. 7 and 8 show the transmission electron micrographs before and after deformation taken from the specimen annealed at 1000 °C. Before deformation, the dislocation densities are low in both ferrite and austenite and annealing twins in austenite are observed. Moreover, α′martensite formed during quenching can be differentiated from ferrite by the high dislocation density (Fig. 7a). After 10% strain, sharp straight deformation bands and stacking faults appear within austenite
Fig. 6. (a) X-ray diffraction spectra and (b) measured percentages of α′-martensite of the specimens annealed at 1000 °C, which are deformed to the engineering strains of 0%, 10%, 30%, and 51%, which correspond to the true strains of 0, 0.095, 0.262, and 0.412, respectively.
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Fig. 7. Transmission electron micrographs and diffraction patterns of austenitic regions for specimens annealed at 1000 °C, when the alloy is deformed to the engineering strains of (a) 0 and (b) 10%. The second image in (a) illustrates the orientation relationship between austenite and quenched martensite. All diffraction patterns are obtained by aligning the zone axis (ZA) parallel to [110]γ.
Fig. 8. Transmission electron micrographs and diffraction patterns of austenitic regions for specimens annealed at 1000 °C, when the alloy is deformed to the engineering strains (a) of 30% and (b) up to failure. The second images in (a) and (b) are dark field images taken from the reflections of α′-martensite which has N-W and K-S orientation relationship with austenite, respectively. All the diffraction patterns are obtained by aligning the zone axis (ZA) parallel to [110]γ.
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Fig. 9. Transmission electron micrographs and diffraction patterns of austenitic regions for specimens annealed at 1200 °C, which are deformed to the engineering strains (a) of 0, (b) of 30%, and (c) up to failure. All diffraction patterns are obtained by aligning the zone axis (ZA) parallel to [110]γ.
Fig. 10. Concentrations of (a) N, (b) Cr, Ni, and Mn in austenite. Calculated and experimental data are shown with filled and unfilled markers, respectively. Table 1 Equilibrium compositions of austenite at different annealing temperatures. Fe is balanced. Temperature (°C)
Cr (wt%)
Ni (wt%)
Mn (wt%)
C (wt%)
N (wt%)
Cu (wt%)
Mo (wt%)
950 1000 1050 1100 1150 1200
19.40 19.30 19.16 19.00 18.90 18.83
2.27 2.30 2.34 2.39 2.45 2.50
1.92 1.93 1.95 1.97 1.99 2.00
0.026 0.027 0.028 0.031 0.034 0.039
0.23 0.23 0.25 0.27 0.30 0.33
0.76 0.78 0.80 0.80 0.82 0.82
0.51 0.51 0.49 0.48 0.47 0.45
to the annealing temperature are obvious. Since the equilibrium percentage of ferrite increases (Fig. 2a), the concentrations of austenite forming elements such as C, N, Ni, Mn, and Cu are enriched in austenite according to the annealing temperature. Based on the estimated austenite compositions, the driving force for martensitic transformation (Ffcc→hcp, Ffcc→bcc) are calculated and plotted with respect to the annealing temperature in Fig. 11. It is obvious that both Ffcc→hcp and Ffcc→bcc decrease with respect to the temperature, which proves more stable austenite is produced accordingly. This explains the reason why more α′-martensite forms during quenching (Fig. 2b) and DIMT starts more rapidly (Fig. 4) in the alloy annealed at a lower temperature. Moreover, austenite stability seems to determine the growth kinetics and the fraction of DIM. As a result, the peak of the strain hardening rate is the lowest, and the peak is reached most slowly for the alloy annealed at 1200 °C, where the highest stability of austenite is achieved.
3.4. Austenite stability The experimental results clearly show that austenite stability relies on annealing temperature; thus, annealing at different temperatures not only alters the microstructures after quenching (Fig. 2b) but also affects the degree of DIMT (Fig. 5), which determines the tensile behavior (Fig. 4). Therefore, in order to study the stability of austenite produced by different annealing, the compositions of austenite were obtained with EPMA and the concentrations of the main alloying elements, N, Cr, Ni, and Mn are shown in Fig. 10. Moreover, thermodynamic calculation was carried out and the results are shown together in the same figure while the exact values are given in Table 1. It is clear that the effect of the annealing temperature is pronounced in the concentrations of N and Cr in comparison with Ni and Mn. Although the experimental data are slightly deviated from the calculated values, the trends of increasing N and decreasing Cr with respect
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Fig. 11. Calculated driving forces for the transformations from austenite (a) to the hcp phase (Ffcc→hcp), and (b) to α′-martensite (Ffcc→bcc).
Furthermore, it seems that Ffcc→hcp and Ffcc→bcc can provide a hint for a DIMT sequence. While Ffcc→bcc is always positive, Ffcc→ hcp becomes negative with the annealing above 1000 °C, which suggests that transformation from austenite to the hcp phase is unfavorable. In addition, Zhang and Hu have carried out a similar study with a DSS containing 0.14N-0.02C-20Cr-2.1Ni-5.1Mn-0.2Cu0.3Mo-0.4Si (wt%) which has a higher concentration of Mn than the present alloy [21]. They have claimed that DIMT follows γ→hexagonally indexed deformation band→α′, since the deformation bands actively appeared at the early stage of deformation, and later, α′martensite formed at the intersection of the deformation bands. Ffcc→ hcp and Ffcc→bcc of their alloy have been calculated and shown together with those of the present alloy in Fig. 10. Moreover, the higher Mn in Zhang and Hu's alloy results in a negligible effect on Ffcc→hcp while it significantly reduces Ffcc→bcc. Hence, it is plausible that α′martensite forms directly from austenite because Ffcc→bcc is sufficiently high for the present alloy. On the other hand, it was energetically favorable for α′-martensite to form via hexagonally indexed deformation band rather than from austenite for the Zhang and Hu's alloy. Therefore, the DIMT sequence seems to be determined by the value of Ffcc→hcp and Ffcc→bcc as well as their relative difference.
hardening rate was lower. 5. During the deformation, blocky α′-martensite forms directly from austenite (γ→α′), not via hexagonally indexed deformation band. This is because the driving force for α′-martensite transformation is relatively higher for the present alloy in comparison with the alloy which experiences γ→hexagonally indexed deformation band→α′. Acknowledgment The authors cordially appreciate POSCO (Pohang, South Korea) for supplying the materials. References [1] J. Charles, P. Chemelle, in: Proceedings of the 8th Duplex Stainless Steels, Baunne, France, 2010, pp. 29-82. [2] R. Jargelius-Pettersson, Corros. Sci. 41 (1999) 1639–1664. [3] M. Byrnes, M. Grujicic, W. Owen, Acta Met. 35 (1987) 1853–1862. [4] M. Fujisawa, Y. Kato, T. Ujiro, CAMP-ISIJ 22 (2009) 1163–1164. [5] J. Simmons, Mater. Sci. Eng. A 207 (1996) 159–169. [6] G. Ghosh, G. Olson, Acta Metall. Mater. 42 (1994) 3361–3370. [7] G.B. Olson, M. Cohen, Metall. Trans. A 7 (1976) 1897–1904. [8] Q. Ran, Y. Xu, J. Li, J. Wan, X. Xiao, H. Yu, L. Jiang, Mater. Des. 56 (2014) 959–965. [9] J. Choi, J. Ji, S. Hwang, K.-T. Park, Mater. Sci. Eng. A 535 (2012) 32–39. [10] J. Choi, J. Ji, S. Hwang, K.-T. Park, Mater. Sci. Eng. A 534 (2012) 673–680. [11] C. Herrera, D. Ponge, D. Raabe, Acta Mater. 59 (2011) 4653–4664. [12] J. Choi, J. Ji, S. Hwang, K.-T. Park, Metall. Mater. Trans. A 44 (2013) 597–601. [13] Z. Tang, R. Misra, M. Ma, N. Zan, Z. Wu, H. Ding, Mater. Sci. Eng. A 624 (2015) 186–192. [14] L. Mosecker, D. Pierce, A. Schwedt, M. Beighmohamadi, J. Mayer, W. Bleck, J. Wittig, Mater. Sci. Eng. A 642 (2015) 71–83. [15] L. Kru¨ger, S. Wolf, S. Martin, U. Martin, A. Jahn, A. Weiß, P. Scheller, Steel Res. Inter. 82 (9) (2011) 1087–1093. [16] J. Choi, J. Ji, S. Hwang, K.-T. Park, Mater. Sci. Eng. A 528 (2011) 6012–6019. [17] G. Huang, D. Matlock, G. Krauss, Metall. Trans. A 20 (1989) 1239–1246. [18] G. Olson, M. Azrin, Metall. Trans. A 9 (1978) 713–721. [19] E. Welsch, D. Ponge, S.M. Hafez Haghighat, S. Sandlӧbes, P. Choi, M. Herbig, S. Zaefferer, D. Raabe, Acta Mater. 116 (2016) 188–199. [20] T.-H. Lee, E. Shin, C.-S. Oh, H.-Y. Ha, S.-J. Kim, Acta Mater. 58 (2010) 3173–3186. [21] W. Zhang, J. Hu, Mater. Charact. 79 (2013) 37–42. [22] S. Martin, S. Wolf, U. Martin, L. Kru¨ger, D. Rafaja, Metall. Mater. Trans. A 47 (2016) 49–58. [23] S. Martin, C. Ullrich, D. Šimek, U. Martin, D. Rafaja, J. Appl. Cryst. 44 (2011) 779–787. [24] A. Weidner, S. Martin, V. Klemm, U. Martin, H. Biermann, Scr. Mater. 64 (2011) 513–516. [25] C.-H. Seo, K. Kwon, K. Choi, K.-H. Kim, J. Kwak, S. Lee, N. Kim, Scr. Mater. 66 (2012) 519–522.
4. Conclusions The effect of annealing temperature on deformation behavior of a 329LA lean duplex stainless steel was investigated. 1. As the annealing temperature increases, the percentage of ferrite increases and austenite is enriched with austenite stabilizing elements. Accordingly, the stability of austenite upon quenching and deformation is significantly improved. 2. In the alloy with less stable austenite, more α′-martensite forms during quenching, which enhances the strength and strain hardening rate. 3. Deformation-induced martensitic transformation (DIMT) is the main deformation mechanism at room temperature regardless of the annealing temperature between 950 and 1150 °C, while specimen annealed at 1200 °C shows almost no DIMT. 4. The stability of austenite decelerates the nucleation and growth of deformation-induced martensite. Thus, the strain hardening rate starts to increase and reaches the peak more slowly for the alloy annealed at higher temperature; moreover, its maximum strain
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Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea
crossmark
Influence of annealing temperature on deformation behavior of 329LA lean duplex stainless steel ⁎
Nithi Saenarjhana, Jee-Hyun Kanga, Soo Chan Leeb, Sung-Joon Kima, a b
Graduate Institute of Ferrous Technology (GIFT), Pohang University of Science and Technology (POSTECH), Pohang 37673, Republic of Korea Technical Research Laboratories, POSCO, Pohang 37859, Republic of Korea
A R T I C L E I N F O
A BS T RAC T
Keywords: Duplex stainless steel Transformation induced plasticity Mechanical properties Electron microscopy
Tensile behavior of a 329LA lean duplex stainless steel at room temperature is investigated after annealing at different temperatures. The austenite composition and thereby, its stability depend on the annealing temperature. For recrystallized alloys, austenite stability increases with the annealing temperature. Therefore, martensitic transformation is effectively suppressed during cooling and deformation, which results in low tensile strength and strain hardening rate, when the steel is annealed at higher temperature. Accordingly, transformation-induced plasticity is more pronounced in alloys annealed at lower temperature. During deformation, α′-martensite forms with a blocky morphology with the absence of the deformation bands indexed as an hcp phase. Thus, it is suggested that austenite directly transforms to α′-martensite during deformation rather than via the deformation bands due to a relatively high driving force for α′-martensite transformation in the present alloy.
1. Introduction Duplex stainless steels (DSSs) exhibit excellent combination of corrosion resistance and mechanical properties, which has enabled their development as structural components in chemical and pulp industries as well as on/offshore applications. In order to achieve the requirements for the application, conventional DSSs contain considerable amounts of expensive alloying elements such as Cr, Ni, and Mo [1]. Recently, efforts have been made to develop lean DSSs with lower amount of the expensive alloying elements. In this regard, Ni and Mo were replaced with Mn and N [1]. Mn plays an important role in both stabilizing austenite and increasing the solubility of interstitial elements such as C and N. N also acts as a strong austenite stabilizer and has beneficial effects on the strength and pitting corrosion resistance [2–5]. DSSs contain two phases, ferrite and austenite, the equilibrium fractions of which depend on annealing temperature. Accordingly, there is a possibility to alter the partitioning of alloying elements into the phases by annealing at different temperatures. Changing the composition of austenite affects its stability, and therefore, determines if it transforms into martensite during quenching as well as deformation. The austenite stability can be represented by the thermodynamic driving force, which is calculated as the difference in Gibbs free energies of fcc and bcc structures. The driving force can be further
⁎
related with martensite start temperature [6] and stacking fault energy [7]. Especially, stacking fault energy is often connected to the deformation mechanisms. It is often reported that austenite in lean DSSs experiences deformation-induced martensitic transformation (DIMT) [8–10]. DIMT plays an important role in mechanical properties by enhancing tensile strength and ductility, which is known as transformation-induced plasticity (TRIP). TRIP results in an excellent combination of strength and elongation over 1 GPa-60% in a DSS [11]. Further increase in stacking fault energy by adjusting the chemical composition [12] or deformation conditions [13–15] can lead to different deformation mechanisms such as twinning and pure dislocation glide rather than DIMT. Hence, the present work is conducted to study the influence of different annealing temperatures on the mechanical behavior and deformation mechanism of a 329LA lean DSS. Tensile tests were conducted along with microstructure analysis, and deformation mechanism is discussed with the assistance of thermodynamic calculations. 2. Experimental procedures The chemical composition of a 329LA lean DSS used in this study is Fe-0.17N-0.02C-20.5Cr-2.0Ni-1.8Mn-0.7Cu-0.6Mo-0.5Si (wt%). A 4.0 mm thick hot-rolled plate was cold- rolled to 1.0 mm in thickness.
Corresponding author. E-mail address: [email protected] (S.-J. Kim).
http://dx.doi.org/10.1016/j.msea.2016.10.062 Received 10 August 2016; Received in revised form 9 October 2016; Accepted 12 October 2016 Available online 21 October 2016 0921-5093/ © 2016 Elsevier B.V. All rights reserved.
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structures, which are ferrite and martensite. On the other hand, the annealing temperatures higher than 900 °C, clearly separate the fcc structure from the bcc structure. Moreover, some bcc structure is present inside austenite grains which implies that α′-martensite has formed during water quenching due to low austenite stability. Note that there is a possibility that some α′-martensite has formed during sample preparation. Therefore, the martensite fraction was not obtained from the phase maps but by a ferrite detector. The measured percentages of magnetic phases after annealing are shown in Fig. 2a with respect to the annealing temperature. Calculated equilibrium percentages of ferrite at different temperatures are shown together for comparison. The result shows that while the equilibrium and measured values are similar for the annealing temperature above 1100 °C, the differences between these values are significant at lower temperatures. The differences indicate that some austenite transforms into α′-martensite during quenching which is consistent with the EBSD results. In order to estimate the percentage of α′-martensite, these differences were multiplied with a correction factor, 1.71 to consider the different magnetic property between ferrite and α′-martensite [16]. Then, the remaining portion other than ferrite and α′-martensite is regarded as austenite, and the estimated phase percentages in the asannealed alloys are shown in Fig. 2b. Note that such estimation is not applicable to unrecrystallized specimens since the equilibrium austenite fraction has not been reached during annealing. Since the fraction of the measured magnetic phases and that of equilibrium ferrite were similar for the alloy annealed at 1200 °C, it seems that no α′-martensite has formed during quenching, and the ferrite percentage was taken as 62%. From Fig. 2, it is concluded that the austenite percentage decreases with respect to annealing temperature over 900 °C; however, the resulting austenite percentage after quenching is approximately 40% for all alloys because less austenite transforms into α′-martensite during quenching.
The plate was machined into subsize tensile specimens with gauge length and width, 33 mm and 10 mm, respectively. The specimens were then subjected to annealing at different temperatures from 800 to 1200 °C for 10 min with the temperature interval of 50 °C, which was followed by quenching in water. Equilibrium percentages and compositions of austenite at different annealing temperatures were calculated with ThermoCalc Version 4.1 with TCFE7 database. Based on the obtained compositions of austenite, molar Gibbs free energies for fcc, hcp, and bcc are calculated to be Gmfcc , Gmhcp , and Gmbcc . Then, the driving forces of the transformations from austenite to an hcp phase and α′-martensite are obtained with F fcc → hcp = −(Gmhcp − Gmfcc ) and F fcc → bcc = −(Gmbcc − Gmfcc ), respectively. Tensile tests were carried out at room temperature with a constant crosshead speed of 1 mm min−1. Interrupted tensile tests were performed to study microstructure evolution. The as-annealed microstructures were analyzed by electron backscattered diffraction (EBSD) with a field emission scanning electron microscope, JEOL JSM7100F. Moreover, N concentrations in austenite were obtained with electron probe microanalysis (EPMA) by JEOL JXA-8530F. In order to identify different phases, X-ray diffraction (XRD) was analyzed by a Bruker D8 Advance with a Cu source. Volume percentages of magnetic phases were measured by a Fisher FMP30 ferrite detector. The deformed structure of austenite was investigated by a transmission electron microscope (TEM), JEOL JEM 2100. Specimens for EBSD were final-polished with colloidal silica after grinding with SiC papers up to #2000. TEM specimens were prepared by grinding and polishing a thin foil down to < 100 µm thick, which were followed by jet polishing at room temperature with a solution of 12% perchloric acid in acetic acid.
3. Results and discussion 3.1. As-annealed microstructure
3.2. Tensile properties Phase maps of undeformed microstructures after annealing at 800, 950, 1000, and 1200 °C obtained by EBSD are shown in Fig. 1. It is clear that both bcc and fcc structures become coarser with increasing annealing temperature especially from 950 °C. For annealing temperature up to 900 °C, most of the microstructure consists of the bcc
The engineering stress-strain curves at room temperature are shown in Fig. 3. It is observed that the tensile behavior of 329LA is sensitive to the annealing temperature. Yield and tensile strengths decrease gradually as the annealing temperature increases, while
Fig. 1. Electron backscattered diffraction (EBSD) phase maps of as-annealed microstructures after annealing at (a) 800 °C, (b) 950 °C, (c) 1000 °C and (d) 1200 °C. Note that the scale of (d) is different from others to reveal large microstructures.
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True stress or strain hardening rate (MPa)
Engineering stress (MPa)
Fig. 2. (a) Percentages of measured magnetic phases after annealing and those of equilibrium ferrite with respect to annealing temperature. (b) Estimated phase percentages in quenched alloys after annealing at 950–1200 °C.
800 600 400 o
800 C o 850 C o 900 C
200 0
0
o
950 C o 1000 C o 1050 C
o
1100 C o 1150 C o 1200 C
10 20 30 40 50 Engineering strain (%)
5000 4000
o
1000 C o 1100 C o 1200 C
3000 2000 1000 0
60
o
950 C o 1050 C o 1150 C
0.0
0.1
0.2 0.3 True strain
0.4
0.5
Fig. 4. True stress and strain hardening rate with respect to true strain after annealing at 950–1200 °C.
Fig. 3. Engineering stress-strain curves after annealing at different temperatures.
and maximum points of SHR. Moreover, the peak value of SHR is reduced. Therefore, when the annealing temperature increases, DIMT not only starts and ends later, but also takes place less actively. As a result, TRIP effect is reduced with respect to the annealing temperature and the tensile strength decreases while ductility increases accordingly (Fig. 3).
elongation is enhanced. Improvement in ductility over 50% can be achieved when the annealing temperature surpasses 1000 °C. In addition, all curves show a sigmoidal shape which indicates the TRIP due to DIMT as will be shown later. Based on the as-annealed microstructures (Fig. 1), the tensile curves can be categorized into two groups: (i) the curves of unrecrystallized microstructures obtained after annealing at 800–900 °C, and (ii) the curves of fully recrystallized microstructures after annealing at 950–1200 °C. Since the microstructures are complicated in the first group with unrecrystallized martensite and austenite which are yet to reach equilibrium, the main focus will be on the second group. The true stress-strain curves as well as corresponding strain hardening rates (SHRs) are plotted in Fig. 4. The yield strength tends to decrease with respect to the annealing temperature; this is attributed to the combined effects of change in both grain size (Fig. 1) and the phase fraction (Fig. 2b). Since martensite is a hard phase and ferrite is a soft phase, the yield strength is reduced with the decrease in martensite fraction and the increase in ferrite fraction as the alloy is annealed at higher temperature. The SHRs of all alloys decrease to minimum points and then increase with straining. The minimum points are caused by the inflection points in tensile curves (Fig. 3), which are known to be the activation points of DIMT [17,18]. DIMT gradually introduces martensite into microstructure and effectively enhances SHR. When the SHR reaches its maximum, the martensite fraction saturates and dynamic recovery of dislocations reduces SHR subsequently. It is obvious that increasing the annealing temperature delays the minimum
3.3. Microstructure evolution Fig. 5 shows the percentage of α′-martensite which formed during deformation for the specimens annealed at different temperatures. The values were obtained by taking the measured percentage difference of magnetic phases between undeformed and failed specimens and multiplying it by a correction factor, 1.71 [16]. Additionally, Fig. 5 includes the values which are normalized by the percentage of austenite in the as-annealed alloys. The trends of normalized and unnormalized percentages are similar due to similar austenite fraction in as-annealed samples regardless of the annealing temperature (Fig. 2b). Fig. 5 suggests that α′-martensite formed during deformation in all specimens except that annealed at 1200 °C. The percentage of deformationinduced martensite (DIM) increases as annealing temperature decreases. It proves that DIMT plays an important role in strain hardening; the higher peak in SHR is available with annealing at lower temperature (Fig. 4), since a higher degree of DIMT is achieved (Fig. 5). The microstructure evolution of the alloy annealed at 1000 °C is investigated in detail with interrupted tensile testing. XRD spectra in 533
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DIM percntage (%)
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grains as shown in Fig. 7b. Although the deformation bands resemble deformation twins, they are not twins since twin diffraction spots do not appear. No hexagonally indexed deformation band is found, which is consistent with the XRD results (Fig. 6a). Moreover, deformationinduced α′-martensite is not observed, which is attributed to the low percentage less than 2% (Fig. 6b). At the strain of 30% (Fig. 8a), α′martensite often appears in austenite with Nishiyama-Wasserman (NW) ( < 110 > γ// < 001 > α′) orientation relationship. For specimen deformed up to failure (Fig. 8b), a large number of dislocations and stacking faults were observed and both of N-W orientation relationship and Kurdjumov-Sachs (K-S) orientation relationship of < 110 > γ // < 111 > α′ were attained between austenite and α′-martensite. The TEM observation of 1000 °C annealed specimen reveals the morphologies of α′-martensite to be big and bulky rather than needlelike and sharp (Fig. 8a). The morphology of α′-martensite varies depending on the transformation sequence, while the shape of hexagonally indexed deformation band is mostly reported in the form of a needle or sharp band [9–12,16,20–22]. When DIMT follows austenite (γ)→hexagonally indexed deformation band→α′, a small α′-martensite is usually revealed at the intersection of the hexagonally indexed deformation bands [13,16,20,22]. On the other hand, α′-martensite is blocky if the direct transformation, γ→α′, takes place [22,25]. Hence, it is suggested that DIMT occurs mainly via γ→α′ in the present alloy based on the observation of the blocky α′-martensite. This is consistent with the absence of hcp peaks in Fig. 6a. The microstructure evolution during deformation for the alloy annealed at 1200 °C is shown in Fig. 9. Before deformation, an austenite grain contains considerable densities of stacking faults and sharp thin dislocation bands; however, no quenched α′-martensite is found (Fig. 9a). At the strain of 30%, α′-martensite is observed on a rare occasion, and mostly only large densities of dislocations and stacking faults are observed without any extra diffraction spots (Fig. 9b). Throughout the deformation until failure, only dislocated austenite was mostly observed and no clear evidence of deformationinduced martensite was found (Fig. 9c). It is important to note that a high density of sharp thin bands which are observed at every stage do not show any twin spots. According to Welsch et al. [19], these bands are called as dynamic slip bands, which can improve strain hardening. Thus, the SHR plateau in the specimen annealed at 1200 °C at the true strain of 0.18–0.35 (Fig. 4) may be the result of the dynamic slip band formation in the present alloy.
Unnormalized Normalized
60 50 40 30 20 10 0 800
900
1000
1100
1200
Fig. 5. Measured percentages of α′-martensite formed during deformation in failed specimens with respect to annealing temperature. The red markers represent the values normalized with the initial austenite percentages.
Fig. 6a clearly show that the fraction of austenite decreases while that of a bcc structure increases according to the strain. The result is consistent with the magnetic phase measurement presented in Fig. 6b. However, additional diffraction peaks did not appear in the present alloy although an hcp phase is sometimes detected during deformation [13,16,20,22]. Since it has been suggested that the deformation bands which are clusterings of stacking faults can appear as the hcp structure in XRD spectra and EBSD phase maps [23,24], such structure is called as a hexagonally indexed deformation band in the present work. Accordingly, the absence of the hcp peaks in Fig. 6a suggests that the formation of the deformation bands is not a pre-requisite for transforming austenite into α′-martensite. Moreover, DIMT is significantly enhanced between the engineering strains of 10% and 30%. These correspond to the true strains equal to 0.095 and 0.262, respectively, where the SHR reaches its peak from the minimum point (Fig. 4). Thus, it demonstrates that DIMT effectively improves strain hardening. Figs. 7 and 8 show the transmission electron micrographs before and after deformation taken from the specimen annealed at 1000 °C. Before deformation, the dislocation densities are low in both ferrite and austenite and annealing twins in austenite are observed. Moreover, α′martensite formed during quenching can be differentiated from ferrite by the high dislocation density (Fig. 7a). After 10% strain, sharp straight deformation bands and stacking faults appear within austenite
Fig. 6. (a) X-ray diffraction spectra and (b) measured percentages of α′-martensite of the specimens annealed at 1000 °C, which are deformed to the engineering strains of 0%, 10%, 30%, and 51%, which correspond to the true strains of 0, 0.095, 0.262, and 0.412, respectively.
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Fig. 7. Transmission electron micrographs and diffraction patterns of austenitic regions for specimens annealed at 1000 °C, when the alloy is deformed to the engineering strains of (a) 0 and (b) 10%. The second image in (a) illustrates the orientation relationship between austenite and quenched martensite. All diffraction patterns are obtained by aligning the zone axis (ZA) parallel to [110]γ.
Fig. 8. Transmission electron micrographs and diffraction patterns of austenitic regions for specimens annealed at 1000 °C, when the alloy is deformed to the engineering strains (a) of 30% and (b) up to failure. The second images in (a) and (b) are dark field images taken from the reflections of α′-martensite which has N-W and K-S orientation relationship with austenite, respectively. All the diffraction patterns are obtained by aligning the zone axis (ZA) parallel to [110]γ.
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Fig. 9. Transmission electron micrographs and diffraction patterns of austenitic regions for specimens annealed at 1200 °C, which are deformed to the engineering strains (a) of 0, (b) of 30%, and (c) up to failure. All diffraction patterns are obtained by aligning the zone axis (ZA) parallel to [110]γ.
Fig. 10. Concentrations of (a) N, (b) Cr, Ni, and Mn in austenite. Calculated and experimental data are shown with filled and unfilled markers, respectively. Table 1 Equilibrium compositions of austenite at different annealing temperatures. Fe is balanced. Temperature (°C)
Cr (wt%)
Ni (wt%)
Mn (wt%)
C (wt%)
N (wt%)
Cu (wt%)
Mo (wt%)
950 1000 1050 1100 1150 1200
19.40 19.30 19.16 19.00 18.90 18.83
2.27 2.30 2.34 2.39 2.45 2.50
1.92 1.93 1.95 1.97 1.99 2.00
0.026 0.027 0.028 0.031 0.034 0.039
0.23 0.23 0.25 0.27 0.30 0.33
0.76 0.78 0.80 0.80 0.82 0.82
0.51 0.51 0.49 0.48 0.47 0.45
to the annealing temperature are obvious. Since the equilibrium percentage of ferrite increases (Fig. 2a), the concentrations of austenite forming elements such as C, N, Ni, Mn, and Cu are enriched in austenite according to the annealing temperature. Based on the estimated austenite compositions, the driving force for martensitic transformation (Ffcc→hcp, Ffcc→bcc) are calculated and plotted with respect to the annealing temperature in Fig. 11. It is obvious that both Ffcc→hcp and Ffcc→bcc decrease with respect to the temperature, which proves more stable austenite is produced accordingly. This explains the reason why more α′-martensite forms during quenching (Fig. 2b) and DIMT starts more rapidly (Fig. 4) in the alloy annealed at a lower temperature. Moreover, austenite stability seems to determine the growth kinetics and the fraction of DIM. As a result, the peak of the strain hardening rate is the lowest, and the peak is reached most slowly for the alloy annealed at 1200 °C, where the highest stability of austenite is achieved.
3.4. Austenite stability The experimental results clearly show that austenite stability relies on annealing temperature; thus, annealing at different temperatures not only alters the microstructures after quenching (Fig. 2b) but also affects the degree of DIMT (Fig. 5), which determines the tensile behavior (Fig. 4). Therefore, in order to study the stability of austenite produced by different annealing, the compositions of austenite were obtained with EPMA and the concentrations of the main alloying elements, N, Cr, Ni, and Mn are shown in Fig. 10. Moreover, thermodynamic calculation was carried out and the results are shown together in the same figure while the exact values are given in Table 1. It is clear that the effect of the annealing temperature is pronounced in the concentrations of N and Cr in comparison with Ni and Mn. Although the experimental data are slightly deviated from the calculated values, the trends of increasing N and decreasing Cr with respect
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Fig. 11. Calculated driving forces for the transformations from austenite (a) to the hcp phase (Ffcc→hcp), and (b) to α′-martensite (Ffcc→bcc).
Furthermore, it seems that Ffcc→hcp and Ffcc→bcc can provide a hint for a DIMT sequence. While Ffcc→bcc is always positive, Ffcc→ hcp becomes negative with the annealing above 1000 °C, which suggests that transformation from austenite to the hcp phase is unfavorable. In addition, Zhang and Hu have carried out a similar study with a DSS containing 0.14N-0.02C-20Cr-2.1Ni-5.1Mn-0.2Cu0.3Mo-0.4Si (wt%) which has a higher concentration of Mn than the present alloy [21]. They have claimed that DIMT follows γ→hexagonally indexed deformation band→α′, since the deformation bands actively appeared at the early stage of deformation, and later, α′martensite formed at the intersection of the deformation bands. Ffcc→ hcp and Ffcc→bcc of their alloy have been calculated and shown together with those of the present alloy in Fig. 10. Moreover, the higher Mn in Zhang and Hu's alloy results in a negligible effect on Ffcc→hcp while it significantly reduces Ffcc→bcc. Hence, it is plausible that α′martensite forms directly from austenite because Ffcc→bcc is sufficiently high for the present alloy. On the other hand, it was energetically favorable for α′-martensite to form via hexagonally indexed deformation band rather than from austenite for the Zhang and Hu's alloy. Therefore, the DIMT sequence seems to be determined by the value of Ffcc→hcp and Ffcc→bcc as well as their relative difference.
hardening rate was lower. 5. During the deformation, blocky α′-martensite forms directly from austenite (γ→α′), not via hexagonally indexed deformation band. This is because the driving force for α′-martensite transformation is relatively higher for the present alloy in comparison with the alloy which experiences γ→hexagonally indexed deformation band→α′. Acknowledgment The authors cordially appreciate POSCO (Pohang, South Korea) for supplying the materials. References [1] J. Charles, P. Chemelle, in: Proceedings of the 8th Duplex Stainless Steels, Baunne, France, 2010, pp. 29-82. [2] R. Jargelius-Pettersson, Corros. Sci. 41 (1999) 1639–1664. [3] M. Byrnes, M. Grujicic, W. Owen, Acta Met. 35 (1987) 1853–1862. [4] M. Fujisawa, Y. Kato, T. Ujiro, CAMP-ISIJ 22 (2009) 1163–1164. [5] J. Simmons, Mater. Sci. Eng. A 207 (1996) 159–169. [6] G. Ghosh, G. Olson, Acta Metall. Mater. 42 (1994) 3361–3370. [7] G.B. Olson, M. Cohen, Metall. Trans. A 7 (1976) 1897–1904. [8] Q. Ran, Y. Xu, J. Li, J. Wan, X. Xiao, H. Yu, L. Jiang, Mater. Des. 56 (2014) 959–965. [9] J. Choi, J. Ji, S. Hwang, K.-T. Park, Mater. Sci. Eng. A 535 (2012) 32–39. [10] J. Choi, J. Ji, S. Hwang, K.-T. Park, Mater. Sci. Eng. A 534 (2012) 673–680. [11] C. Herrera, D. Ponge, D. Raabe, Acta Mater. 59 (2011) 4653–4664. [12] J. Choi, J. Ji, S. Hwang, K.-T. Park, Metall. Mater. Trans. A 44 (2013) 597–601. [13] Z. Tang, R. Misra, M. Ma, N. Zan, Z. Wu, H. Ding, Mater. Sci. Eng. A 624 (2015) 186–192. [14] L. Mosecker, D. Pierce, A. Schwedt, M. Beighmohamadi, J. Mayer, W. Bleck, J. Wittig, Mater. Sci. Eng. A 642 (2015) 71–83. [15] L. Kru¨ger, S. Wolf, S. Martin, U. Martin, A. Jahn, A. Weiß, P. Scheller, Steel Res. Inter. 82 (9) (2011) 1087–1093. [16] J. Choi, J. Ji, S. Hwang, K.-T. Park, Mater. Sci. Eng. A 528 (2011) 6012–6019. [17] G. Huang, D. Matlock, G. Krauss, Metall. Trans. A 20 (1989) 1239–1246. [18] G. Olson, M. Azrin, Metall. Trans. A 9 (1978) 713–721. [19] E. Welsch, D. Ponge, S.M. Hafez Haghighat, S. Sandlӧbes, P. Choi, M. Herbig, S. Zaefferer, D. Raabe, Acta Mater. 116 (2016) 188–199. [20] T.-H. Lee, E. Shin, C.-S. Oh, H.-Y. Ha, S.-J. Kim, Acta Mater. 58 (2010) 3173–3186. [21] W. Zhang, J. Hu, Mater. Charact. 79 (2013) 37–42. [22] S. Martin, S. Wolf, U. Martin, L. Kru¨ger, D. Rafaja, Metall. Mater. Trans. A 47 (2016) 49–58. [23] S. Martin, C. Ullrich, D. Šimek, U. Martin, D. Rafaja, J. Appl. Cryst. 44 (2011) 779–787. [24] A. Weidner, S. Martin, V. Klemm, U. Martin, H. Biermann, Scr. Mater. 64 (2011) 513–516. [25] C.-H. Seo, K. Kwon, K. Choi, K.-H. Kim, J. Kwak, S. Lee, N. Kim, Scr. Mater. 66 (2012) 519–522.
4. Conclusions The effect of annealing temperature on deformation behavior of a 329LA lean duplex stainless steel was investigated. 1. As the annealing temperature increases, the percentage of ferrite increases and austenite is enriched with austenite stabilizing elements. Accordingly, the stability of austenite upon quenching and deformation is significantly improved. 2. In the alloy with less stable austenite, more α′-martensite forms during quenching, which enhances the strength and strain hardening rate. 3. Deformation-induced martensitic transformation (DIMT) is the main deformation mechanism at room temperature regardless of the annealing temperature between 950 and 1150 °C, while specimen annealed at 1200 °C shows almost no DIMT. 4. The stability of austenite decelerates the nucleation and growth of deformation-induced martensite. Thus, the strain hardening rate starts to increase and reaches the peak more slowly for the alloy annealed at higher temperature; moreover, its maximum strain
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