Reversed Austenite Growth Behavior of a 13%Cr-5%Ni Stainless Steel during Intercritical Annealing

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1 ISIJ International, Vol. 56 (2016), ISIJ International, No. 1 Vol. 56 (2016), No. 1, pp Reversed Austenite Growth Behavior of a 13%Cr-5%Ni Stainless Steel during Intercritical Annealing Pengcheng SONG, 1) Wenbo LIU, 1) Chi ZHANG, 1) * Lu LIU 2) and Zhigang YANG 1) 1) Key LAMSBoratory of Advanced Materials of Ministry of Education, School of Materials Science and Engineering, Tsinghua University, Beijing, China. 2) State Nuclear Power Research Institute, Beijing, China. (Received on May 21, 2015; accepted on October 13, 2015) The austenite growth behavior during intercritical annealing of a martensitic stainless steel 13Cr-5Ni was investigated using X-ray diffraction (XRD), electron backscatter diffraction (EBSD), transmission electron microscopy (TEM) and Thermocalc-Dictra simulation. The samples were firstly heated to 650 C with the heating rate of 10 C/s, and then cooled to room temperature without holding for anytime. Experimental results show that reversed austenite tends to nucleate and grow with an acicular shape along lath boundaries (LAMSBs) and with a globular shape along grain boundaries (GBs). Alloy samples were then held at 650 C for different times, and both the EBSD and XRD results show that the amount of retained austenite at room temperature firstly increases with holding time, displaying a peak at about 20 minutes, and then starts to decrease. Together with the Thermocalc-Dictra simulation results and our previous results, this observation indicates that the austenite growth is accompanied with the diffusion of Ni from matrix to reversed austenite. KEY WORDS: martensitic stainless steel; reversed austenite growth; intercritical annealing. 1. Introduction Reaustenitization is generally one of the most effective industrially applied methods of grain refining in steel production. However, for some high alloy steels, 1 3) the average grain size can t be refined by reaustenitization because of the austenite memory, which can be described by the recovery of the shape and size of original austenite grains during the heat treatment process. 4) The nucleation and growth of reversed austenite during heating process are regarded to play critical roles in austenite memory. 5 7,9 13,17) In order to make clear the mechanisms of austenite memory, it s critical to investigate the formation mechanism of the reversed austenite. It is reported that austenite growth behavior at γ+α dual-phase region plays a major role in austenite memory phenomenon, and many works have been done to investigate this process. 5,8,11,13 16,18 21) It is commonly accepted that there are two kinds of reversed austenite formation mechanisms, namely acicular and globular mechanisms, respectively. Globular austenite tends to nucleate at original austenite grain boundaries or inside the lathes with substrate of precipitates, thus disintegrate the original austenite grain and lead to grain refinement 11,13,16,18) In the acicular mechanism, acicular austenite nucleates at lath boundaries and has a K-S orientation relationship with the matrix, and recovery of original austenite grain happens after phase transformation. 11,13,14,16,18) However, there is still a demand * Corresponding author: chizhang@tsinghua.edu.cn. DOI: for more comprehensive investigations for the formation mechanism of the reversed austenite, since possible factors that affect the nucleation and growth behaviors of reversed austenite are complicated, for instance, Ni/Cr ratio, 6) C content, 2) heating rate, 3) intercritical annealing temperature 5) and so on. A lot of work has been done to investigate the reversed austenite growth behavior in Fe Cr Ni alloys. 1 3,5,6) Ni diffusion has been proved to exhibit undeniable importance effects on the reversed austenite growth. With low Ni/ Cr ratio or low heating rate reversed austenite grows with Ni diffusion, while without Ni diffusion under high Ni/ Cr ratio and high heating rate. Quantitatively, the ratio is compared with 1 and the heating rate is compared with 10 k/s. 3,6) The C content also affects the diffusion process of Ni. With higher C content globular austenite tends to nucleate with Ni diffusion, while with lower C content acicular austenite tends to nucleate without Ni diffusion. 6) Heating rate has similar effects on Ni diffusion, i.e., higher heating rate corresponds to diffusionless process and lower heating rate gives rise to Ni diffusion. 3) Hence, it is important to understand the Ni diffusion during intercritical annealing. In our previous work, 5,15) effects of annealing temperature within dual-phase region on reversed austenite formation have been studied, showing that acicular austenite formed with the Ni diffusion and led to the recovery of original austenite grain, and the amount of retained austenite firstly increased and then decreased with the increase of annealing temperature. However, the effects of different annealing time on the reversed austenite formation at dual-phase region are still unclear. In the present work, more precise 2016 ISIJ 148

2 study on the reversed austenite formation during intercritical annealing has been done using electron backscatter diffraction (EBSD) and transmission electron microscopy (TEM). Dictra simulations about the Ni diffusion process in Fe-13Cr-5Ni ternary system have also been done at 650 C for different holding times. 2. Experiments and Dictra Simulations The composition of 13Cr-5Ni martensitic stainless steel used in this investigation is listed in Table 1. The A s temperature is 528 C, and the A f temperature is 830 C. 5) To homogenize alloy elements and other precipitates or phases, the alloy was homogenized at C for 5 h and sectioned into samples of size mm 3. To make clear the effects of intercritical annealing on retained austenite, the specimens were also isothermally held at 650 C from 0 min to min with the same heating rate of 10 C/s, and finally quenched into water. Optical microscopy (OM), TEM and EBSD were used to investigate the microstructure evolutions during intercritical annealing, and X-ray diffraction (XRD) was also used to detect the volume fraction of retained austenite. Composition of the phases was analyzed using the scanning transmission electronic microscopy (STEM) and energy-dispersive spectroscopy (EDS). Specimens for OM were mechanically polished and etched in chloroazotic acid. The specimens for TEM observation were first mechanically ground down to ~40 μm thickness followed by two-jet polishing in a solution of 10% percholoric acid and 90% alcohol at 30 C. Samples for EBSD were again mechanically polished and the strain caused by it was removed using electrochemical polishing. EBSD with the scanning width of 0.15 μm was used to study the microstructural evolution. X-ray diffraction tests were made by using Cu Kα radiation between 2θ value of 40 and 95. Thermocalc-Dictra simulation of Ni diffusion in 13Cr- 5Ni steel has also been carried out. In this simulation, annealing conditions are fitted as with a heating rate of 10 C/s to 650 C and holding for different times. The initial state of tempered martensite (BCC) structure is assumed as a bar with homogeneous Ni distribution, and reversed austenite formation is set to start from the right side. Databases that we used here are TCFE7 and Mob2 for thermodynamics and mobility respectively. no retained austenite exists after homogenization treatment followed by the deep quenching process in liquid nitrogen Austenite Formation during Heating Process For the purpose of studying the reversed austenite nucleation and growth behavior, homogenization treated specimens are heated to an intercritical temperature of 650 C with the heating rate of 10 C/s. The bright field image taken from the specimen, which is heated to 650 C without holding then quenched to room temperature, is shown in Fig. 2. Typical lath structure can be seen in Fig. 2, and the XRD result states that there are both of martensite and retained austenite, this is coincident with the EBSD results. Together with the composition of block 1 and block 2 that are shown in Table 2, from the Ni content of each block we know block 1 is retained austenite and block 2 is tempered martensite. And the retained austenite morphology at room temperature of the specimen in the bright field image is shown in Fig. 3. Blue region represents retained austenite, Fig. 1. Optical micrograph of initial martensite in 13Cr-5Ni. 3. Results and Discussions 3.1. Initial Martensitic Microstructure OM of initial structure after homogenization treatment (1 200 C for 5 h) is shown in Fig. 1. The original austenite grain boundaries and typical lath martensitic microstructure can be clearly seen. The average grain size of original austenite is about 180 μm, 5) and XRD results indicates that Fig. 2. The TEM morphology of martensite lathes after heating to 650 C, (a) bright field image, (b) SAED pattern, (c) index of (b). Table 1. Chemical composition (wt%) of the 13Cr-5Ni martensitic stainless steel. Alloy elements C Si Mn Ni Cr Mo V Cu P S Fe Composition Bal ISIJ

3 Table 2. Chemical composition (wt%) by STEM of block 1 (black) and block 2 (white) in Fig. 2(a). Alloy elements Si Ni Cr Fe Lath Lath Fig. 4. The EBSD morphologies after holding for different times at 650 C: (a) 20 min, (b) 60 min, (c) 600 min, (d) min. (Online version in color.) Fig. 3. The EBSD morphology after heating to 650 C, black arrows point to the globular austenite and the white arrows point to the acicular austenite. (Online version in color.) and red area stands for martensite phase. Grain boundaries (GBs) are indicated by the bold black curve, small angle boundaries (SABs) (2 θ<5 ) are the gray lines, and large angle martensite sub-block boundaries (LAMSBs) (5 θ) are represented by the black line, and from the work of Morito et al. 22) the angle of martensite sub-block is about 10. From our previous work, Liu 17) has pointed out that the preferential oritation are mainly 2 5 and around 45, here 2 5 corresponding to small angle boundaries 45 corresponding to random large angle boundaries. Comparing Fig. 3 to Figs. 4(a) and 4(b), that two kinds of retained austenite are found, acicular shaped retained austenite mainly distributed at lath boundaries as been pointed out by white arrows and globular shaped retained austenite mainly appears at original austenite grain boundaries and some of them are also in prior austenite grain as been pointed out by black arrows. Obviously, the amount of acicular shaped retained austenite is larger than that of globular shaped retained austenite. Most of the SABs (maybe defects) exist inside martensite lathes, and the (LAMSBs) are between variants. Although SABs, LAMSBs, GBs are all preferred sites for reversed austenite nucleation, in the present work most of the retained austenite exists at LAMSBs then GBs, few retained austenite are found at SABs Austenite Growth during Annealing Process In order to investigate the reversed austenite growth behavior during intercritical annealing process, the samples were held at 650 C for different times. EBSD maps of intercritically annealed alloy sample with different holding times are shown in Fig. 4. The amount of retained austenite decreased with holding time after the initial increase from 0 to 20 min. The average size and number density of retained austenite grains within the previous austenite grains are larger than that in the GBs. It can be inferred that the reversed austenite is preferred to nucleate and grow at lath boundaries or packet boundaries. After annealing at 650 C for 100 h (Fig. 4(d)), few austenite can be detected in the sample water-quenched to room temperature. Moreover, as you may notice LAMSBs are larger than those in Figs. 4(a), 4(b) and 4(c), this is because of the black line thickness is larger than those in Figs. 4(a), 4(b) and 4(c). The volume fraction of retained austenite at room temperature is also detected by XRD, and calculated with the direct comparison method: 23) RI γ α Vγ = + 1 RαI γ where V γ is the volume percent of retained austenite, R is a term that depends on interplanar spacing, composition of specimen and crystal structure. It can be calculated with first principles, and we get R γ = 1 from 24) for reasonable Rα simplicity, I is the integrated intensity. In our calculation we use integrated intensities of the peaks of (200) γ, (200) α, (220) γ and (211) α that are measured with the automated diffractometers. Multiple peaks intensities ratio minimizes the effect of preferred orientations, as the specimens are randomly oriented. Results in Fig. 5 show that volume percentage of retained austenite at room temperature first increases with holding time, displaying a peak at about 20 minutes, and then decreases at longer holding time. There is no retained austenite for each specimen at initial state. 15) The results are similar to our previous work where the volume ISIJ 150

4 Fig. 5. The volume percentage of retained austenite at room temperature vs holding time at 650 C. percent of retained austenite at room temperature depends on reverse treatment temperature. 5) Many works have been done about the mechanism of reversed austenite formation of Cr and Ni containing steels. 1 8) There are lots of factors that determine the reversed austenite formation behavior, such as composition (mainly the ratio of Ni/Cr and C content), 2,3,6) heating rate, 2,8) holding temperature 1,5) and holding time. It is accepted that proposed mechanisms of the reversed austenite growth can be divided into two kinds: diffusionless mechanism and diffusive mechanism. Reversed austenite formation behaviors in two kinds of Fe Cr Ni steels (16Cr-10Ni and 18Cr-9Ni) have been studied with the same experimental condition. 6) The one with higher Ni/Cr ratio (16Cr-10Ni) displayed diffusionless reversed austenite formation behavior, and the one with lower Ni/Cr ratio (18Cr-9Ni) displayed diffusive mechanism. This is because higher ratio of Ni/Cr leads to larger difference between free energy of austenite and martensite. Thus, lower austenitization temperature is needed for martensitic shear reversion. In the present work, the Ni/ Cr ratio of 13Cr-5Ni is even smaller than 18Cr-9Ni in, 6) but the reversed austenite formed with Ni diffusion. The reasons for this phenomenon can be attributed to the different heating rates in the present alloy (10 C/S) and in 18Cr-9Ni steel (300 C/S). 6) E. S. Park et al. 2) reported that in Fe Ni C alloys the mechanism for reversed austenite formation is decided by both of C content and heating rate. Low C content and high heating rate are diffusionless preferential factors, while high C content and low heating rate lead to diffusive reversed austenite formation. Our results agree with this statement. Effects of heating rate on reversed austenite formation have been further studied by 3) with the steel of Fe-13Cr-7Ni- 3Si. The results turned out to be that with the heating rates higher than 10 C/s reversed austenite tends to form with a diffusionless mechanism, while under the heating rates lower than 10 C/S diffusion process determines the formation of reversed austenite. In the present study, the heating rate is 10 C/S and reversed austenite formed diffusionally. Furthermore, Seo-Jae and his co-workers 7) reported that in annealing treatment of Fe-11Cr-9Ni-7Mn-0.02Si, acicular shaped reversed austenite form with a diffusionless mechanism at lath boundaries during continuous heating, and then globular austenite starts to form in tempered martensite in holding process. About the mechanism of the formation of acicular shaped reversed austenite, this work contradicts our previous studies, 5,15) but is the same at the point of formation of globular process. In our previous work of Fe-13Cr-5Ni steel, 5,15) during continuous heating from room temperature to austenite single phase region, acicular shaped retained austenite acts as pre-existing nuclei, and reversed austenite grew under a K-S orientation relationship with martensite lathes. Reversed austenite formation is accompanied with the Ni diffusion process. Acicular shaped retained martensite after solution treatment has been directly characterized with TEM in 1) in steel of Fe-15Cr-7Ni. Reversed austenite showed a plate-like shape and formation process is controlled by Ni diffusion. The amount of retained austenite at room temperature is affected by two factors, the volume fraction of reversed austenite at 650 C and the stability of reversed austenite. The stability of reversed austenite depends on the Ni content of reversed austenite and varies during the reverse treatment with variation in Ni concentration. At shorter holding time Ni atoms enrich in the reversed austenite which increases the stability of reversed austenite and the volume fraction of reversed austenite increases. However, for longer time the stability of reversed austenite starts to decrease due to the homogenization of Ni in reversed austenite and the decrease of the amount of retained austenite at room temperature. Therefore diffusion and redistribution of Ni between reversed austenite and matrix is responsible for the reversed austenite transformation and growth. 1,5,15) Dictra simulation with an attention to Ni diffusion during austenite reversion in the 13Cr-5Ni steel is shown in Fig. 6. The Ni concentration profiles between tempered martensite (BCC) and reversed austenite (FCC) formed with different holding times are exhibited in Fig. 6(a). Ni concentration becomes more uniform in both BCC and FCC phases with the increase of holding time. This means reverse transformation reaches the equilibrium at annealing temperature of 650 C. Moreover, the Ni concentration in reversed austenite is obviously higher than that in tempered martensite. However, as is displayed in Fig. 6(a), Ni diffusion distance in tempered martensite is much longer and slower to get to the equilibrium state. It means Ni distribution process between parent phase and reversed austenite is controlled by Ni diffusion in martensite phase. Position of phase interface has also been calculated as a functional of holding time, which is shown in Fig. 6(b). Length of the initial BCC bar is set as 100 grids, so the interface position moves from 100 to 76 (the length of BCC part, when transformation reach equilibrium state at 650 C) as the ordinate puts it. The horizontal ordinate demonstrates transformation time. Together with Fig. 6(a), it is easy to get the information that for each holding time, the position of phase interface and the position of Ni concentration profiles between two phases are almost the same. As has been stated above, the amount of retained austenite first increases, at 20 mins (holding time) reaches to the peak then decrease. Proposed mechanism for this phenomenon is the amount of retained austenite at room temperature depends on the ISIJ

5 Fig. 6. Thermocalc-Dictra simulation results of Ni distribution during reversed austenite formation. (a) content of Ni in BCC and FCC structures at different holding times, (b) position of BCC and FCC interface at each time.(online version in color.) volume fraction of reversed austenite at 650 C and the Ni content of reversed austenite, higher volume fraction and Ni content results in higher amount of retained austenite. From the DICTRA simulation data Fig. 6(b) we know that at 650 C it takes more than 1e+06s (277.8 hours) to get to the equilibrium volume fraction. In our experiments, the longest holding time, 100 h, is less than the time to get to the equilibrium state. Thus, in experiments with the increase of holding time the volume fraction of reversed austenite at 650 C increase linearly. However, from Fig. 6(a) we know that the Ni content of reversed austenite at high temperature is almoset the same at different holding times, which means the amount of retained austenite should be increased linearly with holding time. This is different with the conclusion in Fig. 5. Possible reason for this difference could be that in the DICTRA simulation Ni diffusion happens in Fe-13Cr-5Ni (wt %) tenary system, while in the experiments Ni diffusion happens in a more complicated circumstance. Still, we can get the statement that reversed austenite formation process is controlled by Ni diffusion. 5) 3.4. Proposed Mechanism of Reversed Austenite Growth Behavior Our experimental results and previous works 16,20,21,25) have demonstrated that mechanism of reversed austenite formation from non-equilibrium phase such as martensite can be divided into two kinds. The one is acicular shaped reversed austenite coherently nucleated at martensite lath boundaries, within martensite grains (acicular mechanism). The other is globular shaped reversed austenite randomly nucleated at martensite grain boundaries (globular mechanism). K-S orientation relationship between reversed austenite and martensite lathes, and diffusional mechanism of globular austenite have been reported by the previous work. 21,24) Proposed mechanism for reversed austenite formation in 13Cr-5Ni during intercritical annealing is schematically shown in Fig. 7. In former works, 5,16) acicular mechanism and globular mechanism are separately discussed within Fig. 7. Proposed mechanisms for (a) nucleation of reversed austenite, (b) growth of globular austenite, (c) growth of acicular austenite, in 13Cr-5Ni steel. different specimens. At beginning of reversed austenite formation in the present steel (Fig. 3), both of globular austenite nucleation and acicular austenite nucleation happen simultaneously in the same specimen, which have been shown schematically in Fig. 7(a). That is to say, GBs and LAMSBs are all favored nucleation sites for reversed austenite. In the following growth process, however, whether acicular mechanism or globular mechanism has a leading role depends on heat treatment conditions. As demonstrated above, acicular austenite formed coherently or maybe both coherently and diffusionally and globular austenite formed diffusionally, and it is reasonable to make the statement that with slow heating rate globular mechanism dominates the reversed austenite growth behavior 2,7) as has been shown in Fig. 7(b), while under fast heating rate acicular mechanism controls the reversed austenite growth behavior 2,7) as Fig. 7(c) puts it, and comparing with acicular austenite, the volume of globular austenite is much smaller and hasn t been demonstrated in this figure, vice versa for Fig. 7(b). In the present work, the acicular mechanism has been regarded as the dominant role during the reversed austenite 2016 ISIJ 152

6 nucleation and growth. Retained austenite is preferred to exist at lath boundaries with an acicular shape (Fig. 3), and during the intercritical annealing process it acts as reversed austenite nucleus, thus leading to a LAMSBs preferred acicular shaped reversed austenite nucleation process. It is reported that acicular austenite formation can lead to a surface relief effect from the EBSD situ observation. 5) Hence, it tends to establish a closer crystallographic relationship with martensite lathes. It is likely to infer that acicular austenite bonded to martensite lathes with an oriented manner to form coherent interface. Moreover, the oriented relation restricts acicular austenite could be mainly formed by one orientation, as shown schematically in Fig. 7(c). 4. Conclusion The reversed austenite transformation phenomenon in 13Cr-5Ni steel during intercritical annealing was studied. Following conclusions are drawn from the reported results and analysis: (1) Acicular-shaped retained austenite distributing along LAMSBs and globular-shaped retained austenite forming at GBs were observed after water qunching from 650 C, and the amount of retained austenite at room temperature first increases then decreases with the increase of holding time at 650 C. (2) Reversed austenite transformation is accompanied with the diffusion of Ni atoms from martensite to reversed austenite, and Ni content affects the stability of reversed austenite positively. Acknowledgments This work was supported by financial support from the National Natural Science Foundation of China (No ). REFERENCES 1) E. S. Park, D. K. Yoo, J. H. Sung, C. Y. Kang, J. H. Lee and J. H. Sung: Metal. Mater. Int., 10 (2004), ) C. A. Apple and G. Krauss: Acta Metall., 20 (1972), ) D. S. Leem, Y. D. Lee, J. H. Jun and C. S. Choi: Scr. Mater., 45 (2001), ) L. Liu, Z. G. Yang, C. Zhang and W. B. Liu: Mater. Sci. Eng. A, 527 (2010), ) K. Tomimura, S. Takaki and Y. Tokunaga: ISIJ Int., 31 (1991), ) S. J. Lee, Y. M. Park and Y. K. Lee: Mater. Sci. Eng. A, 515 (2009), 32. 7) V. D. Sadovsky: Structural Inheritance, Metallurgia, Moscow, (1973), 10. 8) A. S. Chaus, F. I. Rudnitskii and M. Murgas: Metal Sci. Heat Treat., 39 (1997), 53. 9) L. T. Zayats, D. O. Panov and M. G. Zakirova: Metal Sci. Heat Treat., 50 (2008), ) A. V. Supov: Metal Sci. Heat Treat., 38 (1996), ) T. Hara, N. Maruyama, Y. Shinohara, H. Asahi, G. Shigesato, M. Sugiyama and T. Koseki: ISIJ Int., 49 (2009), ) V. V. Sagaradze, V. E. Danilchenko, Ph. L Heritier and V. A. Shabashov: Mater. Sci. Eng. A, 337 (2002), ) Y. K. Lee, H. C. Shin, D. S. Leem, J. Y. Choi, W. Jin and C. S. Choi: Mater. Sci. Technol., 19 (2003), ) N. Nakada, T. Tsuchiyama, S. Takaki and S. Hashizume: ISIJ Int., 47 (2007), ) L. Liu, Z.-G. Yang and C. Zhang: J. Alloys Compd., 577 (2013), S ) S. T. Kimmins and D. J. Gooch: Met. Sci., 17 (1983), ) L. Liu: Tsinghua University, 2012, 12, PhD thesis. 18) H. Shirazi: Tohoku University, 2013, 53, PhD thesis. 19) S. S. D yachenko: Metal Sci. Heat Treat., 42 (2000), ) S. Matsuda and Y. Okamura: Trans. Iron Steel Inst. Jpn., 14 (1974), ) S. Matsuda and Y. Okamura: Trans. Iron Steel Inst. Jpn., 14 (1974), ) S. Morito, X. Huang, T. Furuhara, T. Maki and N. Hansen: Acta Mater., 54 (2006), ) Lambda Technologies: Diffraction Notes, No. 33, (2006), 1. 24) SAE : Retained Austenite and Its Measurement by X-Ray Diffraction, SAE Special Publication, 453, SAE, Warrendale, PA, ) S. Watanabe and T. Kunitake: Tetsu-to-Hagané, 61 (1975), ISIJ