Continuous Cooling Bainite Transformation Characteristics of a Low Carbon Microalloyed Steel under the Simulated Welding Thermal Cycle Process

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1 Available online at SciVerse ScienceDirect J. Mater. Sci. Technol., 2013, 29(5), 446e450 Continuous Cooling Bainite Transformation Characteristics of a Low Carbon Microalloyed Steel under the Simulated Welding Thermal Cycle Process Xiangwei Kong 1)*, Chunlin Qiu 2) 1) School of Mechanical Engineering and Automation, Northeastern University, Shenyang , China 2) State Key Laboratory of Rolling Technology and Automation, Northeastern University, Shenyang , China [Manuscript received November 25, 2012, in revised form February 15, 2013, Available online 26 March 2013] Continuous cooling transformation of a low carbon microalloyed steel was investigated after it was subjected to the simulation welding thermal cycle process and the interrupted cooling test. Microstructure observation was performed by optical microscopy and transmission electron microscopy. On the basis of the dilatometric data and microstructure observation, the continuous cooling transformation (CCT) diagram was determined, which showed that the main microstructure changes from a mixture of lath martensite and bainitic ferrite to full granular bainite with the increase in the cooling time t 8/5 from 10 to 600 s, accompanied with a decrease in the microhardness. The interrupted cooling test confirmed that the bainitic ferrite can form attached to grain boundaries at the beginning of transformation even if the final microstructure contains a mixture of granular bainite and bainitic ferrite. KEY WORDS: Low carbon microalloyed steel; Bainitic ferrite; Continuous cooling transformation; Welding 1. Introduction In recent decades, the hot rolled low carbon microalloyed steels have been widely used in building, bridge, pipeline and offshore structures because of their excellent balance of high strength and good toughness [1,2]. However, these excellent mechanical properties can be disturbed by the welding thermal cycles characterized by rapid heating and uneven cooling rate with high peak temperature, especially for large heat input welding process [3e5]. Therefore, how to predict the microstructure evolution in the heat affected zone (HAZ) of the weldment is a necessity to optimize the mechanical properties and to ensure the safety of the weldable steels. Normally, continuous cooling transformation (CCT) diagram is regarded as the effective method to predict microstructural evolution during the welding thermal cycle because the composition, cooling rate, and austenite grain size of the material are related to the phase transformation temperature and resultant microstructure [6]. Although quite a number of CCT diagrams have been provided for the traditional steels according to the previous literature [6e8], continued development of the chemical composition design of the * Corresponding author. Prof., Ph.D.; Tel.: þ ; address: xwkong@me.neu.edu.cn (X. Kong) /$ e see front matter Copyright Ó 2013, The editorial office of Journal of Materials Science & Technology. Published by Elsevier Limited. All rights reserved. modern steels requires an understanding of their transformation behavior in detail under the welding thermal cycle conditions. Furthermore, because the cooling rate is very uneven at the whole weld cooling stage [9], the microstructure evolution may exhibit different features at the different transformation stages. For example, the product phase may be controlled by a shear mechanism at the beginning stage of the transformation due to the relatively high cooling rate. On the contrary, the carbon diffusion rate may be mainly responsible for the transformation rate as the cooling rate decreases at the finish stage of the transformation. Therefore, understanding of the change in microstructure morphology at different stages of the transformation is also very important to completely exploit the transformation behavior during the welding. However, rare research focused on the microstructure morphology at different stages of the transformation [10]. In this study, the simulation welding thermal cycle technique was employed to conduct the continuous cooling transformation test of the coarse grained HAZ (CGHAZ) of a modern low carbon microalloyed steel. The corresponding CCT diagram was determined according to the dilatometric data and microstructure observation. Meanwhile, the interrupted cooling process by water quench was used to investigate in detail the change of microstructure morphology with the increase in volume fraction transformation. 2. Experimental The low carbon bainitic steel for the study was a 20-mm thick plate produced by thermo mechanical controlled process

2 X. Kong and C. Qiu: J. Mater. Sci. Technol., 2013, 29(5), 446e Table 1 Chemical compositions and mechanical properties of the as-rolled low carbon bainitic steel Chemical compositions (wt%) C Si Mn Nb V Ti B Cr, Mo, Ni P cm * Mechanical properties R t0.5 (MPa) R m (MPa) A (%) Toughness (J) ( 20 C) Note: *P cm ¼ C þ Si Mn þ Cu þ Cr þ þ Ni þ Mo 15 þ V 10 þ 5B. (TMCP). Its chemical composition and mechanical properties are shown in Table 1. The carbon content and welding crack susceptibility index P cm are 0.053% and 0.183%, respectively, in order to improve the steel resistance to the welding cold crack in the HAZ. The specimens cut from the steel plate were machined into long cylindrical shape with the dimension of V6 50 mm. The welding thermal cycle simulation was carried out on a thermomechanical simulator to study microstructural evolution when the specimens were subjected to different heat inputs welding thermal cycle process. The welding thermal cycle was controlled by a 2-D Rykalin mathematical model and the detailed process parameters were as follows: the specimens were rapidly heated at a rate of 130 C/s to the peak temperature of 1350 C and held for 2 s. The cooling rates were determined by the cooling times from 800 to 500 C(t 8/5 ), the range of which were chosen from 10 to 600 s in order to simulate different welding heat inputs, and the temperature was controlled by the Rykalin model until it fell to 250 C at which the phase changes have completed. The typical thermal cycle temperature curves measured are shown in Fig. 1. For the water quenching tests, two welding thermal cycle processes with the cooling time of 30 and 120 s were selected. When the temperature controlled by the Rykalin model reached the setting temperature, i.e. 440, 470, 500 C, the temperature control program was interrupted immediately and the specimens were cooled down through the manual water quenching method, as indicated with arrows in Fig. 1. After the simulation welding thermal cycle tests, all metallographic specimens were obtained by mechanical polishing and etching with 3% nital etchant, and then observed using a Leica optical microscope. The microhardness was determined by utilizing an FM 700 hardness testing machine employing a load of 49 N In order to identify the refinement structure of some water quenched specimens, the thin-foil specimens were observed by transmission electron microscopy (FEI-Tecnai G 2 F20 TEM, Netherlands). 3. Results and Discussion 3.1. Microstructure morphology and CCT diagram Fig. 2 shows the microstructural variation with the increase of the cooling time after the specimens were subjected to the simulation welding process with the peak temperature of 1350 C. The main microstructure is bainitic ferrite as the cooling time t 8/5 is 30 s (Fig. 2(a)). It is obvious that the prior austenite grain boundaries are present in the final microstructure because the bainitic ferrite forms by a shear mechanism reported by Bhadeshia [11], who considered that this kind of microstructure always nucleates along the prior austenite grain boundary. The typical morphology of bainitic ferrite can be also found in the previous study [12]. Additionally, it is evident that the abnormal growth of the prior austenite grains occurs at the high peak temperature and the average grain size is about 87 mm measured by linear intercept method. As the cooling time increases to 120 s, the microstructure is characterized by a mixture of bainitic ferrite and granular bainite (see Fig. 2(b)). With the further increase in the cooling time, the morphology of bainitic ferrite disappears gradually and the microstructure is predominantly granular bainite even though the cooling time increases to 300 s and 600 s (Fig. 2(c) and (d)). Meanwhile, the longer the cooling time, the larger the average austenite grain size forms, mainly due to the extended duration time under higher temperature condition [4]. On the other hand, the amount of the second phase, martensite-austenite (MA) constituent, increases and its morphology changes from dot shape to massive shape with increasing the cooling time. On the basis of the dilatometric data which was measured by using a dilatometer, the transformation start/finish temperature can be determined. In combination with the microstructure observation as mentioned above, the CCT diagram of the CGHAZ can be obtained, as shown in Fig. 3. At a heating rate of 130 C/s the A c3 temperature measured is about 868 C. The fully martensite transformation start temperature (M s ) can be calculated based on the chemical composition using the empirical formula shown in Eq. (1) [13]. M s ¼ x C 30x Mn 11x Si 12x Cr 18x Ni 7x Mo (1) Fig. 1 Typical thermal cycle temperature curves (the dash arrow lines signified the manual water quenched temperature schematically). It can be found that the transformation start temperatures are very close to the M s dot-lines when the cooling time is 10 s, and the final microstructure contains higher fraction of lath martensite. As the cooling time increases, the transformation start temperature rises gradually but the transformation finish temperature shows a

3 448 X. Kong and C. Qiu: J. Mater. Sci. Technol., 2013, 29(5), 446e450 Fig. 2 Optical images showing the microstructure variation with the increase of cooling time t 8/5 : (a) 30 s, (b) 120 s, (c) 300 s, (d) 600 s (BF: bainitic ferrite, GB: granular bainite, MA: martensite-austenite constituent). trend of first increase and then decrease, which is mainly responsible for the results that the main microstructure changes from a mixture of lath martensite and bainitic ferrite to fully granular bainite. On the other hand, the average microhardness decreases with the increase in cooling time. The maximum microhardness of a mixture of lath martensite and bainitic ferrite is much less than 350 HV which is regarded as the critical value of the HAZ for avoiding hydrogen cracking [1]. Thus, this result probably suggests that the experimental steel has an excellent ability of resistance to welding cold crack Microstructure evolution during interrupted cooling process The partially transformed microstructure can be effectively obtained by water quenching method. Fig. 4 shows the Fig. 3 Continuous cooling transformation diagram of the simulated CGHAZ welding thermal cycle (LM: lath martensite). microstructure evolution with the cooling time of 30 and 120 s at different quenching temperatures. When the cooling time is 30 s, very few amount of bainitic ferrites with lath morphology exist at the quenching temperature of 500 C (Fig. 4(a)), due to the quenching temperature just corresponding to the transformation start temperature as shown in the CCT diagram. When the quenching temperature decreases to 470 C, some bainitic ferrites with fine lath width form attached to the grain boundaries (Fig. 4(b)) and the volume fraction transformed is only about 21% measured by using the point count method based on several random micrographs. The amount of bainitic ferrite increases and the bainitic ferrite lath becomes wider when the quenching temperature further decreases to 440 C (Fig. 4(c)). The change in ferrite lath width is partly attributed to carbon segregation and different phase transformation temperatures [11]. Meanwhile, it is evident that the bainitic ferrite sheaves with different growth directions have a hard impingement event inside the grain [12] and some secondary bainite sheaves seem to nucleate on the primary bainite sheaves, which is usually called sympathetic nucleation [14]. By contrast, the morphologies of microstructures with the cooling time of 120 s are presented in Figs. 4(d)e(f) when the quenching temperature changes from 530 to 470 C. The final microstructure contains a mixture of bainitic ferrite and granular ferrite (Fig. 4(f)). In general, the formation of bainitic ferrite is later than that of granular ferrite because of the higher transformation driving force needed for the formation of bainitic ferrite [15]. The present result represented in Fig. 2 also supports this explanation because the bainitic ferrite almost disappears when the cooling time t 8/5 prolongs to 300 s or more. However, it can be found that the morphology of bainitic ferrite lath nucleates along the grain boundary at the beginning of transformation (see Fig. 4(d)). Meanwhile, when the quenching temperature

4 X. Kong and C. Qiu: J. Mater. Sci. Technol., 2013, 29(5), 446e Fig. 4 Optical micrographs showing microstructural evolution with the cooling process at the different cooling times of (aec) 30 s and (def) 120 s and different water quenching temperatures: (a) 500 C, (b) 470 C, (c) 440 C, (d) 530 C, (e) 500 C, (f) 470 C (M: martensite). decreases to 500 C, the amount of bainitic ferrite increases largely while the amount of granular bainite is still very little (see Fig. 4(e)). This strongly indicates that the bainitic ferrite seems to be transformed first at the mixture of product phase, which may be mainly attributed to the grain boundaries with high-energy to promote the nucleation of bainitic ferrite[14] Refined microstructure Fig. 5 shows the refined microstructure of the specimen with the cooling time of 120 s and the quenching temperature of 470 C. It can be found that the quenching temperature of 470 C is close to the transformation finish temperature based on the CCT diagram and Fig. 5 TEM images showing the refinement morphologies of microstructure with the cooling time of 120 s and the water quenching temperature of 470 C.

5 450 X. Kong and C. Qiu: J. Mater. Sci. Technol., 2013, 29(5), 446e450 optical microstructure in Fig. 4(f). Fig. 5(a) shows that massive MA constituent with irregular boundaries precipitates on the ferrite matrix and micro twinning structure occurs in the MA constituents (signified with arrow in Fig. 5(a)). According to the previous study [16], the MA constituent formed during the welding thermal cycle has a variable structure in nature, such as high dislocation/ twinning structure and full austenite of the fcc crystal structure. Fig. 5(b) shows that bainitic ferrite forms in a small region and the thin film MA constituent occurs at the lath boundaries. Additionally, the length of bainitic ferrite seems to be limited by grain boundaries and the massive MA constituents with high carbon content are presented at the front of growth direction of bainitic ferrite. The width of bainitic ferrite laths is distinct when they transformed in the different temperature ranges as compared Fig. 5(b) with Fig. 5(c). However, all of these bainitic ferrite laths contain high density dislocation, as shown in Fig. 5(d). These high density dislocations mainly come from three factors [17,18] : the first one is that the dislocation debris introduced into the parent phase can be inherited by the bainitic ferrite; the second is from the transformation dislocation; the third factor is about the thermal stress resulted from the uneven welding thermal cycle. Meanwhile, the bainitic ferrite containing the relatively high dislocation density is usually regarded as an important factor to support the idea that it forms by a shear mechanism [17]. And high density dislocation produced in the welding microstructure is also considered as an effective way to relax the welding residual internal stress [18]. 4. Conclusions (1) The main microstructure of the low carbon microalloyed steel changes from a mixture of lath martensite and bainitic ferrite to full granular ferrite with the increase in the cooling time t 8/5 when the specimens are subjected to the simulated CGHAZ welding thermal cycle process. The maximum microhardness is about 295 HV, indicating that this steel has an excellent ability of resistance to welding cold crack. (2) The interrupted cooling test shows that the bainitic ferrite can nucleate firstly along the grain boundary even if the final microstructure contains a mixture of granular bainite and bainitic ferrite. And the bainitic ferrite lath always contains high dislocation density, which is mainly resulted from the transformation mechanism and welding internal stress. Acknowledgments The authors gratefully acknowledge the financial support of Shenyang Key Laboratory of Construction Project (Grant No. F ) and Science Foundation for the Excellent Youth Scholars of Ministry of Education of China (Grant No ). REFERENCES [1] W.B. Morrison, Mater. Sci. Technol. 25 (2009) 1066e1073. [2] F. Xiao, M. Zhao, Y. Shan, B. Liao, K. Yang, J. Mater. Sci. Technol. 20 (2004) 779e781. [3] D.M. Viano, N.U. Ahmed, G.O. Schumann, Sci. Technol. Weld. Join. 5 (2000) 26e34. [4] L. Lan, C. Qiu, D. Zhao, X. Gao, L. Du, Mater. Sci. Eng. A 529 (2011) 192e200. [5] D.P. Fairchild, D.G. Howden, W.A.T. Clark, Metall. Mater. Trans. A 31 (2000) 641e652. [6] P.L. Harrison, R.A. Farrar, Int. Mater. Rev. 34 (1989) 35e51. [7] R.W. Fonda, G. Spanos, Metall. Mater. Trans. A 31 (2000) 2145e [8] Y.Q. Zhang, H.Q. Zhang, W.M. Liu, H. Hou, Mater. Sci. Eng. A 499 (2009) 182e186. [9] K. Poorhaydari, B.M. Patchett, D.G. Ivey, Weld. J. 84 (2005) 149se155s. [10] A. Lambert-Perlade, A.F. Gourgues, A. Pineau, Acta. Mater. 52 (2004) 2337e2348. [11] H.K.D.H. Bhadeshia, Mater. Sci. Eng. A 273e275 (1999) 58e 66. [12] L.Y. Lan, C.L. Qiu, D.W. Zhao, X.H. Gao, L.X. Du, Mater. Sci. Technol. 27 (2011) 1657e1663. [13] S.M.C. Van Bohemen, Mater. Sci. Technol. 28 (2012) 487e495. [14] H.I. Aaronson, W.T. Reynolds Jr., G.R. Purdy, Metall. Mater. Trans. A 37 (2006) 1731e1745. [15] K. Wu, Z. Li, A.M. Guo, X. He, L. Zhang, A. Fang, L. Cheng, ISIJ Int. 46 (2006) 161e165. [16] L. Lan, C. Qiu, D. Zhao, X. Gao, L. Xiu, J. Mater. Sci. 47 (2012) 4732e4742. [17] H.K.D.H. Bhadeshia, J.W. Christian, Metall. Trans. A 21 (1990) 767e797. [18] K. Poorhaydari, B.M. Patchett, D.G. Ivey, Mater. Sci. Eng. A 435e 436 (2006) 371e382.