Effect of Carbon and Manganese on the Quenching and Partitioning Response of CMnSi Steels

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1 , pp Effect of Carbon and Manganese on the Quenching and Partitioning Response of CMnSi Steels Emmanuel De MOOR, 1) John Gordon SPEER, 1) David Kidder MATLOCK, 1) Jai-Hyun KWAK 2) and Seung-Bok LEE 2) 1) Advanced Steel Processing and Products Research Center, George S. Ansell Department of Metallurgical and Materials Engineering, Colorado School of Mines, 1500 Illinois Street, Golden, CO USA. 2) POSCO, 700 Gumbo, Gwangyang, Jeonnam, South Korea. (Received on June 14, 2010; accepted on August 30, 2010) CMnSi steel grades with carbon contents ranging from 0.2 to 0.3 wt% and manganese contents of 3 and 5 wt% were Quenched and Partitioned (Q&P). Tensile properties were assessed and retained austenite fractions measured. Intercritically annealed and fully austenitized conditions were studied. The best combinations of tensile strength and total elongation obtained in the 0.2C 3Mn 1.6Si grade after intercritical annealing were associated with strength levels in the MPa range and total elongations ranging from 14 to 20%. Optimum properties were obtained in the 0.3C 3Mn 1.6Si steel after full austenitization with tensile strength levels ranging from to MPa and total elongations ranging from 11 to 18%. The 0.2C 3Mn 1.6Si fully austenitized samples also exhibited remarkable strength/ductility combinations albeit at lower strength levels of MPa UTS with 9 15% total elongation indicating the effectiveness of the manganese addition to develop novel property combinations. KEY WORDS: third generation AHSS; Quenching and Partitioning; martensite; retained austenite. 1. Introduction Quenching and Partitioning (Q&P) is receiving increased attention as a potential processing route to develop socalled third generation advanced high strength sheet steel (AHSS) properties. 1 14) The proposed heat treatment aims at stabilizing retained austenite in martensitic microstructures through decarburization of the martensitic matrix and carbon enrichment of the austenite. The Q&P design identifies an optimum quench temperature, QT, corresponding to the greatest fraction of austenite that can be stabilized with the carbon available in the martensite formed during quenching to the QT. 1) Carbide formation should be avoided as it may act as a carbon sink. Thus, TRIP steel type compositions may be suitable for Q&P processing, and have been studied by several authors. 2 4,6) The addition of molybdenum to a CMnSi TRIP steel composition has been shown to result in increased austenite stabilization whereas the partial replacement of Si by Al resulted in lower fractions. 6) The present work investigates the effect of carbon and manganese modifications on the Q&P response of a typical CMnSi TRIP steel composition. Two levels of carbon (0.2 and 0.3 wt%) and manganese (3 and 5 wt%) were considered. 2. Experimental Procedure The chemical compositions of the investigated laboratory-prepared steel grades are given in Table 1. The material was received as cold rolled sheet with a thickness of 1 mm. Longitudinal tensile specimens were machined according to the ASTM E8 geometry. Specimens were heat treated using salt pots, following the heat treating matrix given in Table 2. A fixed reheating time of 2 min was used and the employed reheating temperatures are based on Thermo-Calc thermodynamic calculations shown in Fig. 1 to obtain an intercritical microstructure consisting of 50 vol% intercritical ferrite or a fully austenitized microstructure. The amount of intercritical ferrite was measured using light optical microscopy (LOM) on intercritically annealed and water quenched samples to verify the Thermo-Calc predictions. The results are shown in Table 3. Reasonable agreement was obtained between the predicted and measured volume fractions of intercritical ferrite. To ensure full austenitization, the reheating temperatures were approximately 50 C above the predicted A e3 Table 1. Chemical composition, in wt%, of laboratory prepared steel grades. 137

2 Table 2. Heat Treating Matrix. Table 3. Volume fraction of intercritical ferrite (% a int ) obtained by LOM phase area point counting of annealed and water quenched samples. temperature and LOM was conducted to verify the absence of intercritical ferrite. The quench temperatures were calculated according to the approach proposed by Speer et al. 1) assuming full carbon depletion of the martensite and homogenization throughout the austenite in the absence of cementite or transition carbide formation. Fully stable austenite is assumed at the quench temperature and during partitioning. The results are shown in Fig. 2(a) for intercritical annealing and in Fig. 2(b) for full austenitization. Quench temperatures 20 C above or below the calculated optimum QT, and also 250 C were included in some cases. Quenching was done in Durferrit AS 140 heat treating salt for quench temperatures above 160 C. Quenching to lower temperatures was done using organic Paratherm NF heat transfer fluid. Initial heat treatments were conducted for each QT with tensile specimens having a spot welded thermocouple, connected to a data recorder in order to assess the quench time required to reach the desired quench temperature. It was found that the desired QT was obtained after a quenching time of 10 s for quenching in salt, whereas 18 s were required for quenching in the oil bath to the lower quench temperatures. The heat lost during transfer from the quench bath to the partitioning bath, amounting to 20 C in some cases, was also taken into account. Partitioning was done for 10, 30, and 100 s at 400 C and for 10 s at 450 C. Following partitioning, the samples were water quenched to room temperature. Tensile testing was done on an electro-mechanical tensile machine at a constant strain rate of /s, with a 2 inch (50.8 mm) 50% extensometer. Two samples were tested for each heat treating condition. Measured yield strengths were based on the 0.2% offset method, uniform strains were determined as the engineering strain at the peak load used for UTS calculation, and total strains to failure were obtained from the extensometer output at final fracture. All samples were observed to fail within the specified extensometer gage length. The retained austenite content was measured for each microstructure using X-ray diffraction (XRD). The samples were lightly ground to remove surface oxides followed by chemical thinning for 3 min using a solution of 50 parts water, 50 parts 30% hydrogen peroxide, 1 part hydrofluoric acid. The XRD analysis was performed on a Phillips X-pert diffractometer operating at 45 kv and 40 ma, using an Fig. 1. Thermo-Calc predicted equilibrium phase distributions for the three alloys. X celerator detector, filtered copper radiation, and a 1 degree slit. Quantification of the austenite content employed four austenite peaks and four ferrite/martensite peaks: {111}, {200}, {220}, {311} and {110}, {200}, {211}, {220}, respectively. A 2-theta scan was run from 40 to 105 degrees. Each sample was run twice. Data were analyzed using Profit software to produce integrated intensity values. The volume fractions of retained austenite were calculated according to the SAE method. 15) 3. Results and Discussion 3.1. Intercritically Annealed Steels The tensile properties obtained in the intercritically annealed and Q&P processed microstructures are given in Table 4 for the 0.2C 3Mn 1.6Si composition. The retained austenite fractions and carbon contents measured by XRD are also given in the table. The same properties are given for the 0.3C 3Mn 1.6Si and 0.3C 5Mn 1.6Si steels in Table 5 and Table 6, respectively. Tensile strength levels of MPa are combined with total elongations ranging from 14 to 20% in the 0.2C 3Mn 1.6Si grade. The 0.3C 3Mn 1.6Si grade exhibits tensile strength levels of MPa with total elongations of 14 18% whereas the properties in the 0.3C 5Mn 1.6Si grade range from to MPa tensile strength and 7 to 15% total elongation. A low yield to tensile ratio is observed in the 0.2C 3Mn 1.6Si grade and the ratio increases with increased C and Mn alloying. Lower retained austenite fractions are measured than were predicted assuming full carbon depletion of the martensite as shown in Fig. 2. High austenite carbon contents are observed indicating carbon enrichment. Representative SEM micrographs of the microstructure of a 0.2C 3Mn 1.6Si intercritically annealed Q&P sample, with a QT of 185 C, PT of 400 C and Pt of 30 s are given in Fig. 3. Austenite pools and laths are ob- 138

3 Fig. 2. Calculated retained austenite fractions as a function of quench temperature assuming full carbon partitioning for a) intercritical annealing and b) full austenitization. Table 4. Yield strength (YS), ultimate tensile strength (UTS), uniform elongation (UE), total elongation (TE), retained austenite fraction (fg ret ) and carbon content obtained for the intercritically annealed 0.2C 3Mn 1.6Si grade. Table 5. Yield strength (YS), ultimate tensile strength (UTS), uniform elongation (UE), total elongation (TE), retained austenite fraction (fg ret ) and carbon content obtained for the intercritically annealed 0.3C 3Mn 1.6Si grade. Table 6. Yield strength (YS), ultimate tensile strength (UTS), uniform elongation (UE), total elongation (TE), retained austenite fraction (fg ret ) and carbon content obtained for the intercritically annealed 0.3C 5Mn 1.6Si grade. served. Spherical particles are present in the intercritical ferrite which are believed to be cementite that was not dissolved during intercritical annealing. Figure 4 shows an SEM micrograph of an intercritically annealed and water quenched 0.3C 5Mn 1.6Si sample where the remnants of the predecessor pearlitic microstructure are clearly ob- 139

4 ISIJ International, Vol. 51 (2011), No. 1 low retained austenite volume fractions measured by XRD, and further characterization of the austenite fractions by other techniques may be helpful. The cementite observed in Fig. 3, believed to be inherited from the prior microstructure due to insufficient dissolution during intercritical annealing, lowers the carbon available in the intercritical austenite, reducing the potential for austenite retention by Q&P. A clear correlation between the measured volume fractions of retained austenite and the strain hardening behavior is not apparent, and thus the role of retained austenite in fine microstructures is perhaps not fully understood. Further work is needed to understand the mechanical behavior of these microstructures. Stress strain and instantaneous strain hardening as a function of true strain curves are given in Fig. 6 for the intercritically annealed 0.3C 3Mn 1.6Si steel for a QT of 140 C. Clearly higher yield strengths and lower strain hardening are obtained as compared to the 0.2C 3Mn 1.6Si grades. Fairly low fractions of retained austenite were again measured. Figure 7 shows stress strain curves and instantaneous strain hardening as a function of true strain for the 0.3C 5Mn 1.6Si grades for a fixed QT of 70 C. High yield strengths and limited strain hardening are obtained. The stress strain curves and strain hardening behavior observed in Fig. 6 and Fig. 7 is similar to the behavior of tempered Dual Phase steels.4) 3.2. Fully Austenitized Steels Tensile properties, austenite fractions and carbon contents obtained after full austenitization and Q&P processing are given in Tables 7 through 9. Tensile strength levels of MPa and total elongations of 9 15% are obtained in the 0.2C 3Mn 1.6Si grade, and MPa and 2 17% in the 0.3C 3Mn 1.6Si grade. The 0.3C 5Mn 1.6Si grades exhibit low ductility in most cases. Retained austenite fractions between 4 and 9 vol% are obtained in the 0.2C 3Mn 1.6Si grade whereas fractions ranging from 1 to 15 vol% are obtained in the 0.3C 3Mn 1.6Si grade. A peak in austenite fraction is observed with partitioning time in most cases in the latter grade. It can be noted that much more austenite is stabilized following full austenitization as compared to intercritical annealing in this grade. A pronounced effect of QT is observed for the austenite fraction stabilized in the 0.3C 5Mn 1.6Si grade with the highest tested QT resulting in the highest volume fraction of about 14 vol%. Representative microstructures are presented in Figs. 8 through 10. Large spherical precipitates are absent which suggests that the austenitizing treatment was successful in dissolving the cementite of the prior microstructure. Thick-film-like austenite is observed in the microstructure of the 0.2C 3Mn 1.6Si grade quenched to 210 C and partitioned at 400 C for 120 s as indicated by the arrows in Fig. 8. Some fine precipitates are observed in the martensitic regions. Similar observations are made for the 0.3C 3Mn 1.6Si Q&P grade (QT 200 C, PT 400 C and Pt 30 s) shown in Fig. 9. Flat featureless regions are observed in the micrograph of the 0.3C 5Mn 1.6Si Q&P grade (QT 160 C, PT 400 C and Pt 30 s) shown in Fig. 10, presumably indicative of fresh martensite where carbon depletion through tempering or partitioning did not take place to a significant extent. Fig. 3. SEM images of a 0.2C 3Mn 1.6Si sample reheated to 725 C for 120 s, quenched to 185 C and partitioned at 400 C for 30 s followed by final water quenching. 2% nital etched. Image (a) shows undissolved carbides and image (b) shows retained austenite regions (highlighted by arrows) adjacent to martensite. Fig. 4. SEM image of a 0.3C 5Mn 1.6Si sample reheated to 660 C for 120 s and water quenched. 2% nital etched. served. Undissolved cementite was present in the intercritical microstructures of all grades, for the annealing conditions employed in this work. It is clear from Table 4 by the large difference between UTS and YS that the intercritically annealed Q&P processed 0.2C 3Mn 1.6Si samples exhibit significant strain hardening as also observed by the stress strain curves and instantaneous n-values4) plotted as a function of true strain in Fig. 5 for a QT of 165 C and various partitioning conditions. The tensile strength levels decrease with increasing partitioning time and temperature, whereas ductility increases. The n-values decrease with strain after passing through a maximum. This strain hardening behavior may be somewhat unexpected especially given the fairly 140

5 Fig. 5. (a) Stress strain curves and (b) plots of instantaneous strain hardening as a function of true strain for the intercritically annealed 0.2C 3Mn 1.6Si grade for a QT of 165 C for the indicated partitioning conditions temperature and time. The retained austenite volume fractions are also given. Fig. 6. (a) Stress strain curves and (b) plots of instantaneous strain hardening as a function of true strain for the intercritically annealed 0.3C 3Mn 1.6Si grade for a QT of 140 C for different partitioning conditions. Fig. 7. Stress strain curves for the 0.3C 5Mn 1.6Si alloy for a reheating temperature of 660 C and QT of 70 C for different partitioning conditions. Remarkable tensile strength levels are found in combination with significant ductility levels in the 0.2C 3Mn 1.6Si and 0.3C 3Mn 1.6Si Q&P steels resulting from pronounced strain hardening as shown in Fig. 11. A significant effect of partitioning time and temperature on strength and ductility is observed. Low ductility was exhibited by the 0.3C 5Mn 1.6Si steel with numerous samples failing before yielding. The low ductility may be related to the untempered martensitic regions in the microstructure as shown in Fig. 10. Given the low retained austenite fractions obtained, the quench temperature and partitioning conditions may not have been optimized for this steel Comparison with Other Studies The tensile properties obtained in this work are summarized in Fig. 12, plotting total elongation versus tensile strength. It is clear that high strength levels are combined with significant ductility in the 0.2C 3Mn 1.6Si intercritically annealed Q&P heat treated grade and in the 0.2C 3Mn 1.6Si and 0.3C 3Mn 1.6Si grades after full austenitization. The properties are compared with other proposed third generation AHSS processing routes in Fig. 13. The different approaches have been reviewed in 16) and include ultrafine Dual Phase grades obtained through special hot deformation practices, 17,18) modified TRIP processing through either alloying 19 21) or austempering modifications, 3,22,23) bainite where judicious alloying is employed to reduce the bainite transformation to lower temperatures 24 29) in order to obtain ultrafine microstructures, and TWIP/TRIP alloys with lower manganese levels than typically observed for second generation austenitic alloys ) Previously reported properties resulting from Q&P processing of a variety of compositions are also given. 2 4,6 8) Multiple sample geometries and sizes were employed in these studies and, in order to facilitate comparison, the literature data were corrected for sample geometry and size according to the ISO 2566/1-1984(E) standard 34) to an ASTM E8 sample geometry. It is clear from Fig. 13 that exceptional properties have been developed through Q&P processing in the present work, when compared to other promising approaches. Some of these properties are within the desired range for future 3rd generation AHSS steels. Tensile 141

6 Table 7. Yield strength (YS), ultimate tensile strength (UTS), uniform elongation (UE), total elongation (TE), retained austenite fraction (fg ret ) and carbon content obtained for the fully austenitized 0.2C 3Mn 1.6Si grade. Table 8. Yield strength (YS), ultimate tensile strength (UTS), uniform elongation (UE), total elongation (TE), retained austenite fraction (fg ret ) and carbon content obtained for the fully austenitized 0.3C 3Mn 1.6Si grade. Table 9. Yield strength (YS), ultimate tensile strength (UTS), uniform elongation (UE), total elongation (TE), retained austenite fraction (fg ret ) and carbon content obtained for the fully austenitized 0.3C 5Mn 1.6Si grade. strength levels above 2 GPa with appreciably high ductility have also been developed through Q&P heat treating of a 0.41C 1.30Mn 1.27Si 1.01Ni 0.56Cr alloy. 7) 4. Conclusions Three grades with different carbon and manganese contents were heat treated through quenching and partitioning. Tensile properties were assessed and retained austenite fractions and carbon contents were determined by XRD analysis. High strength in combination with significant ductility was exhibited in particular by the 0.2C 3Mn 1.6Si and 0.3C 3Mn 1.6Si grades after full austenitization. Novel property combinations were obtained in the present research. The results also point out the need to understand Fig. 8. SEM image of a 0.2C 3Mn 1.6Si sample reheated to 840 C for 120 s, quenched to 210 C and partitioned at 400 C for 30 s followed by final water quenching. Retained austenite regions are indicated by the arrows. 2% nital etched. 142

7 ISIJ International, Vol. 51 (2011), No. 1 Fig. 10. SEM image of a 0.3C 5Mn 1.6Si sample reheated to 780 C for 120 s, quenched to 160 C and partitioned at 400 C for 30 s followed by final water quenching. 2% nital etched. Fig. 9. SEM image of a 0.3C 3Mn 1.6Si sample reheated to 820 C for 120 s, quenched to 200 C and partitioned at 400 C for 30 s followed by final water quenching. 2% nital etched. Fig. 11. Stress strain curves and plots of instantaneous strain hardening as a function of true strain for the a) fully austenitized 0.2C 3Mn 1.6Si grade for a QT of 250 C and b) fully austenitized 0.3C 3Mn 1.6Si grade for a quench temperature of 200 C and different partitioning conditions. Fig. 12. Tensile properties obtained for the three alloys following a) intercritical annealing and b) full austenitization represented on total elongation versus tensile strength diagrams. 143

8 Fig. 13. Situation of properties obtained in current research with respect to other proposed third generation AHSS development routes. 16) The data were collected from literature and corrected according to ISO 2566/1-1984(E) 34) to the ASTM E8 geometry. better the structure/property relationships in Q&P steels; in particular the contribution of small retained austenite fractions to strain hardening. Acknowledgements POSCO is gratefully acknowledged for providing material and financial support. The sponsors of the Advanced Steel Processing and Products Research Center are acknowledged for their continued support to overall operations of the center. REFERENCES 1) J. G. Speer, D. K. Matlock, B. C. De Cooman and J. G. Schroth: Acta Mater., 51 (2003), ) A. M. Streicher, J. G. Speer, D. K. Matlock and B. C. De Cooman: Proc. of the Int. Conf. on Advanced High Strength Sheet Steels for Automotive Applications, AIST Warrendale, PA, (2004), 51. 3) H. J. Jun and N. Fonstein: Proc. of the Int. Conf. on New Developments in Advanced High Strength Sheet Steels, AIST Warrendale, PA, (2008), ) E. De Moor, S. Lacroix, A. J. Clarke, J. Penning and J. G. Speer: Metall. Trans. A, 39 (2008), ) M. J. Santofimia, T. Nguyen-Minh, L. Zhao, D. N. Hanlon, T. A. Kop and J. Sietsma: Proc. of the Int. Conf. on New Developments in Advanced High Strength Sheet Steels, AIST Warrendale, PA, (2008), ) E. De Moor, J. G. Speer, D. K. Matlock, C. Föjer and J. Penning: Proc. of Materials Science and Technology (MS&T), AIST Warrendale, PA, (2009), ) H. Y. Li, X. W. Lu, W. J. Li and X. J. Jin: Metall. Trans. A, 41 (2010), ) L. Wang and W. Feng: SAE Technical Paper No , Society of Automotive Engineers (SAE) Int., Warrendale, PA, (2010). 9) M. J. Santofimia, J. G. Speer, A. J. Clarke, L. Zhao and J. Sietsma: Acta Mater., 57 (2009), ) M. J. Santofimia, L. Zhao and J. Sietsma: Scr. Mater., 59 (2008), ) A. J. Clarke, J. G. Speer, D. K. Matlock, F. C. Rizzo, D. V. Edmonds and M. J. Santofimia: Scr. Mater., 61 (2009), ) D. H. Kim, J. G. Speer, H. S. Kim and B. C. De Cooman: Metall. Trans. A, 40 (2009), ) G. A. Thomas, J. G. Speer and D. K. Matlock: Proc. of the Int. Conf. on New Developments in Advanced High Strength Sheet Steels, AIST Warrendale, PA, (2008), ) D. H. Kim, J. G. Speer, H. S. Kim and B. C. De Cooman: Proc. of Materials Science and Technology (MS&T), AIST, Warrendale, PA, (2009), ) SP-453, Society of Automotive Engineers (SAE) Int., Warrendale, PA, (1980), ) E. De Moor, P. J. Gibbs, J. G. Speer and D. K. Matlock: Iron Steel Technol., 7 (2010), ) M. Militzer, S. Sarkar, K. Mukherjee, H. Azizi-Alizamini and W. J. Poole: Proc. of New Developments on Metallurgy and Applications of High Strength Steels, TMS, Warrendale, PA, (2008), ) M. Calcagnotto, D. Ponge and D. Raabe: ISIJ Int., 48 (2008), ) O. Matsumura, Y. Sakuma and H. Takechi: Trans. Iron Steel Inst. Jpn., 27, (1987), ) O. Matsumura, Y. Sakuma, Y. Ishii and J. Zhao: ISIJ Int., 32 (1992), ) D. Krizan, B. C. De Cooman and J. Antonissen: Proc. of the Int. Conf. on Advanced High Strength Sheet Steels for Automotive Applications, AIST, Warrendale, PA, (2004), ) K. Sugimoto, T. Iida, J. Sakaguchi and T. Kashima, ISIJ Int., 40 (2000), ) S. Cobo, C. Colin and S. Alain: Proc. of New Developments on Metallurgy and Applications of High Strength Steels, TMS, Warrendale, PA, (2008), ) H. K. D. H. Bhadeshia and D. V. Edmonds: Metall. Sci., 17 (1983), ) H. K. D. H. Bhadeshia and D. V. Edmonds: Metall. Sci., 17 (1983), ) V. T. T. Miihkinen and D. V. Edmonds: Mater. Sci. Technol., 3 (1987), ) V. T. T. Miihkinen and D. V. Edmonds: Mater. Sci. Technol., 3 (1987), ) F. G. Caballero, M. J. Santofimia, C. García-Mateo, J. Chao and C. García de Andrés: Mater. Des., 30 (2009), ) C. Garcia-Mateo and F. G. Caballero: ISIJ Int., 45 (2005), ) G. Frommeyer, U. Brüx and P. Neumann: ISIJ Int., 43 (2003), ) Y. N. Dastur and W. C. Leslie: Metall. Trans. A, 12 (1981), ) M. J. Merwin: Proc. of Materials Science and Technology (MS&T), (2007), ) M. J. Merwin: Mater. Sci. Forum, (2007), ) International Organization for Standardization: Steel Conversion of elongation values Part 1: Carbon and low alloy steels, International Standard ISO 2566/1-1984(E),