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Effect of Cu and B addition on tempering behavior in the weld CGHAZ of high strength low alloy plate steel
Materials Science and Engineering A 497 (2008) 153–159
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of Cu and B addition on tempering behavior in the weld CGHAZ of high strength low alloy plate steel Joonoh Moon a , Sanghoon Kim a , Jongho Lee a , Byoungchul Hwang b , Chang Gil Lee b , Changhee Lee a,∗ a b
Division of Material Sciences and Engineering, Hanyang University, 17 Haengdang-dong, Seongdong-ku, Seoul 133-791, Republic of Korea KIMS, 66 Sangnam-dong, Changwon City, Kyungnam 641-010, Republic of Korea
a r t i c l e
i n f o
Article history: Received 19 May 2008 Received in revised form 18 June 2008 Accepted 18 June 2008 Keywords: Tempering Heat affected zone Cu particle Precipitation Toughness
a b s t r a c t To understand the effect of Cu and B on mechanical properties of the weld coarse-grained heat-affected zone (CGHAZ) of high strength low alloy plate steels, six alloys including different Cu and B content were prepared. The CGHAZ and tempered conditions were simulated by a Gleeble simulator and mechanical properties were measured by Vickers hardness test and charpy impact test. Toughness of as-welded specimens was deteriorated with comparing to base steels and decreased with increasing Cu + B content. Meanwhile, Vickers hardness showed an opposite tendency, which is due to formation of martensite in the weld CGHAZ. For as-tempered conditions, Vickers hardness at first decreased with increasing tempering temperature and then was reinforced at around 550 ◦ C with an unexpected decrease of toughness, which was mainly due to the precipitation of Cu particles and carbides, such as (Ti, Nb)C and (Mo, Mn)2 C. © 2008 Published by Elsevier B.V.
1. Introduction It has been reported that Cu bearing steel with low C content is attractive as structural steel due to their proper balance of strength and toughness [1,2]. This excellent balance is ascribed to the metallurgical phenomena of the precipitation of fine Cu particles during aging [3–6]. Many researches have shown that BCC cluster of Cu particles is nucleated and transformed to stable -Cu particles after meta stable steps of 9R and 3R structure [3,7]. Meanwhile, in recent years, many studies have focused on the effect of boron addition in steels [8–10]. It is well known that B is the most effective element to improve the steel hardenability with only a small addition. That is, B is segregated in a prior austenite grain boundary at high temperature and then makes for a stable prior austenite grain boundary and delays the transformation to ferrite. Therefore, many efforts have been tried to design and control the Cu and B level and developed high strength low alloy steels having excellent mechanical properties [1,8]. Even though there are many interesting properties, when the steel is subjected to welding, microstructure and mechanical properties may be deteriorated by the formation of weld CGHAZ [9]. While many researches for various materials have
∗ Corresponding author. Tel.: +82 2 2220 0388; fax: +82 2 2293 4548. E-mail address: [email protected] (C. Lee). 0921-5093/$ – see front matter © 2008 Published by Elsevier B.V. doi:10.1016/j.msea.2008.06.031
shown that tempering as post weld heat treatment can contribute to the improvement of mechanical properties in the weld CGHAZ [11–13]. In the case of high strength low alloy steels containing Cu and B, only a little investigation on tempering effect in the weld CGHAZ has been dealt with previously despite great works during last several decades [4,9] and is still not clearly known as compared with base steels. In this study, to clear the effect of Cu and B addition on mechanical properties in the CGHAZ and as-tempered steels after welding, six alloys having different Cu and B content were prepared. The weld CGHAZ and subsequent tempering were simulated by a Gleeble simulator. Mechanical properties of base steels, the CGHAZ and as-tempered steels were measured by Vickers hardness test and charpy impact test and then mechanical properties variation will be discussed in regards to the precipitation of Cu particle and carbides.
2. Experimental Table 1 is the chemical composition of the experimental specimens. Cu content was systematically changed from 0 to 1.5 wt% with the same level of other elements and boron was added in steels D, E and F. It should be noted that experimental steels also have Ti and Nb, which contribute to the formation of Ti + Nb carbonitride particles. Fig. 1 shows the thermo-mechanical con-
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Table 1 Chemical compositions of the tested steels (in wt%) Alloy
C
Si
Mn
Ni
Cr
Mo
Ti
Nb
Cu
A
0.07
0.243
1.85
0.292
0.198
0.201
0.015
0.027
–
B (ppm) –
B
0.069
0.248
1.9
0.298
0.198
0.203
0.016
0.028
0.499
–
C
0.068
0.245
1.9
0.297
0.197
0.206
0.016
0.029
1.54
–
D
0.073
0.245
1.88
0.297
0.199
0.201
0.016
0.028
–
10
E
0.075
0.255
1.88
0.3
0.206
0.21
0.016
0.028
0.499
10
F
0.073
0.258
1.9
0.306
0.207
0.213
0.015
0.029
1.57
10
*The steels contain up to 0.034 wt% Al, 0.058 wt% V, <30 ppm N, <120 ppm P and <40 ppm S.
Fig. 1. Schematic schedule for a fabrication of HSLA plate steel.
trolled process (TMCP) for the fabrication of experimental steel plate. The plates were first reheated at 1150 ◦ C and subsequently cooled to 800 ◦ C with controlled rolling. At below 800 ◦ C, the plates were acceleratedly cooled by 20 ◦ C/s. The experimental schedule for cycles of the CGHAZ and subsequent tempering is given in Fig. 2. The CGHAZ of 1350 ◦ C peak temperature with 30 kJ/cm heat input was calculated by the Rosenthal equation [14] and simulated by the Gleeble simulator. Tempering was carried out during 30 min with tempering temperatures from 350 to 650 ◦ C. The tempering time of 30 min was determined for the Cu parti-
Fig. 2. Simulation cycle for the CGHAZ and tempering treatment.
cle precipitation. In the literature [6], it has been reported that aging time above 1000 s is required to precipitate the Cu particle at the temperature range of designed tempering schedule. Mechanical properties were measured for base steels, as-welded and as-tempered specimens by Vickers hardness test and charpy impact test. In order to confirm the microstructure and Cu particle, every specimen was first polished by conventional metallographic techniques and then thin-foiled specimens for TEM observations were prepared by electro-polishing in a mixed solution of 5% perchloric acid and 95% methanol. Meanwhile, TEM samples in order to observe carbides were prepared by conventional carbon replica technique. Heat-treated specimens were polished by using conventional metallographic techniques and then were etched in a mixed solution of methanol (80%) and perchloric acid (20%). A thin carbon film was deposited and extracted with Ni-grid in the same etchant at a voltage of 2 V. The thin-foiled specimens and collected carbon replica were observed by transmission electron microscopy.
3. Results and discussions 3.1. Microstructure and mechanical properties of base steel Fig. 3 shows the TEM micrographs of base steels and indicates that every steel consist of bainite and martensite lath. From TEM images, it can be noted that boron-free steels (steels A, B and C) consist of upper bainite and martensite mixture and steel containing boron (steels D, E and F) consist of full martensite microstructure with fine lath compared to boron-free steels. It is due to the increasing hardenability by boron addition. As mentioned above, B stabilizes the austenite grain boundary by segregation and prohibits the recrystallization of austenite at high temperature [8] and thus delays the austenite to ferrite transformation. Fig. 4 shows the Vickers hardness of base steels and indicates that hardness increases with increasing Cu and B content, which is mainly because of increasing solid solution hardening contribution and martensite formation by increasing hardenability. Meanwhile, Fig. 5 is the representative TEM micrograph showing the particles in the base steels and informs that base steel includes (Ti, Nb)C complex particle, while expected Cu particles have not been found in the base steel. This is due to the slow precipitation kinetics of Cu particle. That is, in the literature [6], precipitation of detectable Cu particle requires aging for larger than 1000 s at around 550 ◦ C. It has been widely accepted that stress applied by high pressure rolling during steel fabrication can accelerate the Cu particle precipitation kinetics [15] and but Cu particles have not been observed in the base steel used in this study despite of high stress level during fabrication.
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Fig. 3. TEM micrographs of base steels.
3.2. Microstructure and mechanical properties in the weld CGHAZ
Fig. 4. Vickers hardness of base steels.
Fig. 6 shows the optical micrographs in the weld CGHAZ of experimental steels and indicates that every specimen consists of full martensite microstructure with coarse prior austenite grain even if the steel does not contain Cu and B such as steel A. It is due to the austenite grain growth at high temperature and subsequent fast cooling during welding with low heat input of 30 kJ/cm. Thus, hardness in the CGHAZ increased as compared to base steels in Fig. 7(a). While toughness decreased after welding, this is also related to the formation of martensite by fast cooling. Fig. 8 shows the TEM micrograph of (Ti, Nb)C particle in the CGHAZ of D alloy. TEM observation indicates that the fraction of (Ti, Nb)C particle decreased largely as compared to base steel. It can be explained by the result of thermocalc calculation as shown in Fig. 9. In the thermo-calc calculation, this study considered all of the phases such as ferrite, austenite, (Ti, Nb)C, liquid, cementite and (Mo, Mn)2 C, together with all components. However, the calculation results of carbide are only plotted
Fig. 5. TEM micrographs of (Ti, Nb)C particle in the base metal of steel A and EDS analysis.
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J. Moon et al. / Materials Science and Engineering A 497 (2008) 153–159
Fig. 6. Optical micrographs in the weld CGHAZ of experimental steels, ×1000.
in Fig. 9. In this study, peak temperature in the CGHAZ is 1350 ◦ C which is higher than dissolution temperature of (Ti, Nb)C as shown in Fig. 9 and it means that whole (Ti, Nb)C particle can be dissolved during heating in the CGHAZ if an enough time was given at high temperature. However, thermal cycle of the CGHAZ in this study is very fast and thus some carbide can remain in the CGHAZ as shown in Fig. 8. 3.3. Precipitation and mechanical properties in as-tempered steels
Fig. 7. Mechanical properties of base steels and as-welded steels: (a) Vickers hardness, (b) charpy V-notch toughness.
To improve the mechanical properties of the CGHAZ, this study carried out the tempering with various temperatures after the CGHAZ simulation, as shown in Fig. 2. Fig. 10 shows the representative TEM images of tempered specimens (30 min at 550 ◦ C) and indicates that each tempered steel consists of tempered martensite microstructure with fine lath. Fig. 11 is the results of Vickers hardness test of as-tempered steels. Hardness initially decreased with increasing tempering temperature and then was reinforced at high tempering temperature above 400 ◦ C. This can be explained by the precipitation of Cu particle and carbide as follows. Previous works have shown that the hardness of tempered lath martensite can be improved by the precipitation of Cu particles and Mo2 C [1], where particles begin to precipitate at above 350 ◦ C [3]. Steels used in this study contain Mo and/or Cu as shown in Table 1. Therefore, reinforcement of tempered steels at high temperature in Fig. 11 is due to the precipitation of Cu particle or Mo2 C [3]. Actually, it has been found that Cu particles in steels C and F, which have 1.5 wt%Cu, precipitated after tempering. Fig. 12 is the Cu particle with 9R fringe pattern in tempered C alloy. Recent studies using high resolution electron microscopy have shown that Cu particles of twinned 9R structure with internal fringe are formed due to stacking fault by shear strain, before the transformation of -Cu particle of FCC structure. It has been found that 9R Cu particle has orientation relationship to the BCC matrix [7,16]. This study has considered that lattice condition of tempered martensite is nearly same with that of ferrite matrix by ejection of saturated carbon from martensite lattice during thermal aging. Thus the orientation relationship between 9R Cu particle and ferrite matrix in Fig. 11(b) has been analyzed as
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Fig. 8. TEM micrographs of (Ti, Nb)C in the CGHAZ of D alloy.
Fig. 9. Equilibrium volume fraction of carbides calculated by thermo-calc with increasing temperature.
follows: ¯ 9R //(1 1 0)Ferrite , (1 1 4)
[1 1¯ 0]9R //[1 1¯ 1]Ferrite
It is well matched with previous report [3,7]. Meanwhile, hardness of A and D alloy, which does not include Cu, was also improved
Fig. 11. Vickers hardness with increasing tempering temperature of experimental steels.
during tempering at 550 ◦ C. It is due to the precipitation of carbide [1], such as (Ti, Nb)C and (Mo, Mn)2 C, during tempering as shown in Fig. 9. Actually, many (Mo, Mn)2 C and (Ti, Nb)C were observed in tempered specimen of D alloy in Fig. 13. As a result, hardness of whole experimental specimens was improved by the precipitation
Fig. 10. Representative TEM micrographs in tempered specimens at 550 ◦ C: (a) C alloy and (b) F alloy.
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J. Moon et al. / Materials Science and Engineering A 497 (2008) 153–159
Fig. 12. Representative TEM micrographs of 9R Cu particle in tempered specimen in C alloy after 30 min aging at 550 ◦ C: (a) TEM bright field image, (b) FFT pattern analysis and (c) EDS analysis.
of Cu particle and/or carbide during tempering at high temperature in Fig. 11. In addition, hardness of Cu-added steels is higher than that of Cu-free steels and it is because that Cu added steels include the Cu particle as well as carbides. Fig. 14 indicates the variation of toughness of as-tempered steels with increasing temperature and shows the opposite tendency to the hardness profile in Fig. 11.
That is, toughness unexpectedly was not improved at high temperature despite of tempering treatment. Previous work has shown that small particles in thermally aged steel [17] make an embrittlement in tempered alloy. Therefore, toughness decrease in Fig. 14 is due to the embrittlement by precipitation of Cu particle and carbides.
Fig. 13. TEM micrographs of carbides in D alloy after tempering: (Ti, Nb)C and (Mo, Mn, Si)2 C in tempered specimen after 30 min aging at 550 ◦ C.
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(3) In tempered steels, Vickers hardness initially decreased with increasing tempering temperature and was reinforced at high temperature. TEM analysis showed that 9R Cu particles and carbides were precipitated during aging and then the reinforcement is mainly due to the precipitation of Cu particle and carbides during aging. While toughness tends to show opposite behavior and it was not improve at high temperature despite of tempering treatment. This result indicates that Cu addition is not useful choice to improve the toughness of the heat-affected zone. Acknowledgement This study was supported by a grant from the Fundamental R&D Program for the Core Technology of Materials funded by the Ministry of Commerce, Industry and Energy, Republic of Korea. References Fig. 14. Charpy V-notch toughness with an increasing tempering temperature of experimental steels.
4. Conclusion remarks (1) In the base steels, Cu particles and Mo2 C except (Ti, Nb)C were not found despite for high deformation by high pressure rolling during fabrication. This is related to the slow precipitation kinetics of them. In the case of B added steels, TEM results showed that the fine martensite microstructure was formed with an addition of boron, which plays a role in making prior austenite grain boundary to be stable and thus increases the hardenability. (2) Vickers hardness of base steels and the CGHAZ increased with increasing Cu and B content. This is due to the solid-solution hardening. While charpy V-notch toughness showed an opposite tendency, which is mainly due to the formation of the hard phase by increasing hardenability with Cu and B addition. Meanwhile, toughness in the CGHAZ is deteriorated as compared to that of the base steels.
[1] J.Y. Koo, M.J. Luton, N.V. Bangaru, R.A. Petkovic, Proc. of ISOPE-2003, 10. [2] K. Kishida, O. Akisue, Tetsu-to-Haganè 76 (1990) 759. [3] N. Maruyama, M. Sugiyama, T. Hara, H. Tamehiro, Mater. Trans. JIM 40 (1999) 268. [4] S.S. Ghasemi Banadkouki, D. Yu, D.P. Dunne, ISIJ Int. 36 (1996) 61. [5] J.B. Yang, M. Enomoto, ISIJ Int. 45 (2005) 1335. [6] J.B. Yang, T. Yamashita, N. Sano, M. Enomoto, Mater. Sci. Eng. A 487 (2008) 128. [7] P.J. Othen, M.L. Jenkins, G.D.W. Smith, Phil. Mag. A 70 (1994) 1. [8] J. Haga, N. Mizui, T. Nagamichi, A. Okamoto, ISIJ Int. 38 (1998) 580. [9] W. Chen, M.C. Chaturvedi, N.L. Richards, Metall. Mater. Trans. A 32 (2001) 931. [10] D.-H. Seo, N.-H. Heo, H.-C. Lee, Met. Mater. 3 (1997) 245. [11] T.-D. Park, K.-K. Baek, D.-S. Kim, Met. Mater. 3 (1997) 46. [12] M. Wang, Y. Wang, F. Sun, Mater. Sci. Eng. A 438–440 (2006) 1139. [13] G. Fu, F. Tian, H. Wang, J. Mater. Proc. Tech. 180 (2006) 216. [14] K. Easterling, Introduction to the Physical Metallurgy of Welding, Butterworths, 1983, p. 17. [15] B. Dutta, E.J. Palmiere, C.M. Sellars, Acta Mater. 49 (2001) 785. [16] B. Soylu, R.W.K. Honeycombe, Proceedings of the Conference on Phase Transformations, Cambridge, 1987, p. 135. [17] W.J. Phythian, C.A. English, J.T. Buswell, Proceedings of the 5th International Meeting on Environmental Degradation of Reactor Materials Water Reactors, Moterey, California, 1991.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Effect of Cu and B addition on tempering behavior in the weld CGHAZ of high strength low alloy plate steel Joonoh Moon a , Sanghoon Kim a , Jongho Lee a , Byoungchul Hwang b , Chang Gil Lee b , Changhee Lee a,∗ a b
Division of Material Sciences and Engineering, Hanyang University, 17 Haengdang-dong, Seongdong-ku, Seoul 133-791, Republic of Korea KIMS, 66 Sangnam-dong, Changwon City, Kyungnam 641-010, Republic of Korea
a r t i c l e
i n f o
Article history: Received 19 May 2008 Received in revised form 18 June 2008 Accepted 18 June 2008 Keywords: Tempering Heat affected zone Cu particle Precipitation Toughness
a b s t r a c t To understand the effect of Cu and B on mechanical properties of the weld coarse-grained heat-affected zone (CGHAZ) of high strength low alloy plate steels, six alloys including different Cu and B content were prepared. The CGHAZ and tempered conditions were simulated by a Gleeble simulator and mechanical properties were measured by Vickers hardness test and charpy impact test. Toughness of as-welded specimens was deteriorated with comparing to base steels and decreased with increasing Cu + B content. Meanwhile, Vickers hardness showed an opposite tendency, which is due to formation of martensite in the weld CGHAZ. For as-tempered conditions, Vickers hardness at first decreased with increasing tempering temperature and then was reinforced at around 550 ◦ C with an unexpected decrease of toughness, which was mainly due to the precipitation of Cu particles and carbides, such as (Ti, Nb)C and (Mo, Mn)2 C. © 2008 Published by Elsevier B.V.
1. Introduction It has been reported that Cu bearing steel with low C content is attractive as structural steel due to their proper balance of strength and toughness [1,2]. This excellent balance is ascribed to the metallurgical phenomena of the precipitation of fine Cu particles during aging [3–6]. Many researches have shown that BCC cluster of Cu particles is nucleated and transformed to stable -Cu particles after meta stable steps of 9R and 3R structure [3,7]. Meanwhile, in recent years, many studies have focused on the effect of boron addition in steels [8–10]. It is well known that B is the most effective element to improve the steel hardenability with only a small addition. That is, B is segregated in a prior austenite grain boundary at high temperature and then makes for a stable prior austenite grain boundary and delays the transformation to ferrite. Therefore, many efforts have been tried to design and control the Cu and B level and developed high strength low alloy steels having excellent mechanical properties [1,8]. Even though there are many interesting properties, when the steel is subjected to welding, microstructure and mechanical properties may be deteriorated by the formation of weld CGHAZ [9]. While many researches for various materials have
∗ Corresponding author. Tel.: +82 2 2220 0388; fax: +82 2 2293 4548. E-mail address: [email protected] (C. Lee). 0921-5093/$ – see front matter © 2008 Published by Elsevier B.V. doi:10.1016/j.msea.2008.06.031
shown that tempering as post weld heat treatment can contribute to the improvement of mechanical properties in the weld CGHAZ [11–13]. In the case of high strength low alloy steels containing Cu and B, only a little investigation on tempering effect in the weld CGHAZ has been dealt with previously despite great works during last several decades [4,9] and is still not clearly known as compared with base steels. In this study, to clear the effect of Cu and B addition on mechanical properties in the CGHAZ and as-tempered steels after welding, six alloys having different Cu and B content were prepared. The weld CGHAZ and subsequent tempering were simulated by a Gleeble simulator. Mechanical properties of base steels, the CGHAZ and as-tempered steels were measured by Vickers hardness test and charpy impact test and then mechanical properties variation will be discussed in regards to the precipitation of Cu particle and carbides.
2. Experimental Table 1 is the chemical composition of the experimental specimens. Cu content was systematically changed from 0 to 1.5 wt% with the same level of other elements and boron was added in steels D, E and F. It should be noted that experimental steels also have Ti and Nb, which contribute to the formation of Ti + Nb carbonitride particles. Fig. 1 shows the thermo-mechanical con-
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J. Moon et al. / Materials Science and Engineering A 497 (2008) 153–159
Table 1 Chemical compositions of the tested steels (in wt%) Alloy
C
Si
Mn
Ni
Cr
Mo
Ti
Nb
Cu
A
0.07
0.243
1.85
0.292
0.198
0.201
0.015
0.027
–
B (ppm) –
B
0.069
0.248
1.9
0.298
0.198
0.203
0.016
0.028
0.499
–
C
0.068
0.245
1.9
0.297
0.197
0.206
0.016
0.029
1.54
–
D
0.073
0.245
1.88
0.297
0.199
0.201
0.016
0.028
–
10
E
0.075
0.255
1.88
0.3
0.206
0.21
0.016
0.028
0.499
10
F
0.073
0.258
1.9
0.306
0.207
0.213
0.015
0.029
1.57
10
*The steels contain up to 0.034 wt% Al, 0.058 wt% V, <30 ppm N, <120 ppm P and <40 ppm S.
Fig. 1. Schematic schedule for a fabrication of HSLA plate steel.
trolled process (TMCP) for the fabrication of experimental steel plate. The plates were first reheated at 1150 ◦ C and subsequently cooled to 800 ◦ C with controlled rolling. At below 800 ◦ C, the plates were acceleratedly cooled by 20 ◦ C/s. The experimental schedule for cycles of the CGHAZ and subsequent tempering is given in Fig. 2. The CGHAZ of 1350 ◦ C peak temperature with 30 kJ/cm heat input was calculated by the Rosenthal equation [14] and simulated by the Gleeble simulator. Tempering was carried out during 30 min with tempering temperatures from 350 to 650 ◦ C. The tempering time of 30 min was determined for the Cu parti-
Fig. 2. Simulation cycle for the CGHAZ and tempering treatment.
cle precipitation. In the literature [6], it has been reported that aging time above 1000 s is required to precipitate the Cu particle at the temperature range of designed tempering schedule. Mechanical properties were measured for base steels, as-welded and as-tempered specimens by Vickers hardness test and charpy impact test. In order to confirm the microstructure and Cu particle, every specimen was first polished by conventional metallographic techniques and then thin-foiled specimens for TEM observations were prepared by electro-polishing in a mixed solution of 5% perchloric acid and 95% methanol. Meanwhile, TEM samples in order to observe carbides were prepared by conventional carbon replica technique. Heat-treated specimens were polished by using conventional metallographic techniques and then were etched in a mixed solution of methanol (80%) and perchloric acid (20%). A thin carbon film was deposited and extracted with Ni-grid in the same etchant at a voltage of 2 V. The thin-foiled specimens and collected carbon replica were observed by transmission electron microscopy.
3. Results and discussions 3.1. Microstructure and mechanical properties of base steel Fig. 3 shows the TEM micrographs of base steels and indicates that every steel consist of bainite and martensite lath. From TEM images, it can be noted that boron-free steels (steels A, B and C) consist of upper bainite and martensite mixture and steel containing boron (steels D, E and F) consist of full martensite microstructure with fine lath compared to boron-free steels. It is due to the increasing hardenability by boron addition. As mentioned above, B stabilizes the austenite grain boundary by segregation and prohibits the recrystallization of austenite at high temperature [8] and thus delays the austenite to ferrite transformation. Fig. 4 shows the Vickers hardness of base steels and indicates that hardness increases with increasing Cu and B content, which is mainly because of increasing solid solution hardening contribution and martensite formation by increasing hardenability. Meanwhile, Fig. 5 is the representative TEM micrograph showing the particles in the base steels and informs that base steel includes (Ti, Nb)C complex particle, while expected Cu particles have not been found in the base steel. This is due to the slow precipitation kinetics of Cu particle. That is, in the literature [6], precipitation of detectable Cu particle requires aging for larger than 1000 s at around 550 ◦ C. It has been widely accepted that stress applied by high pressure rolling during steel fabrication can accelerate the Cu particle precipitation kinetics [15] and but Cu particles have not been observed in the base steel used in this study despite of high stress level during fabrication.
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Fig. 3. TEM micrographs of base steels.
3.2. Microstructure and mechanical properties in the weld CGHAZ
Fig. 4. Vickers hardness of base steels.
Fig. 6 shows the optical micrographs in the weld CGHAZ of experimental steels and indicates that every specimen consists of full martensite microstructure with coarse prior austenite grain even if the steel does not contain Cu and B such as steel A. It is due to the austenite grain growth at high temperature and subsequent fast cooling during welding with low heat input of 30 kJ/cm. Thus, hardness in the CGHAZ increased as compared to base steels in Fig. 7(a). While toughness decreased after welding, this is also related to the formation of martensite by fast cooling. Fig. 8 shows the TEM micrograph of (Ti, Nb)C particle in the CGHAZ of D alloy. TEM observation indicates that the fraction of (Ti, Nb)C particle decreased largely as compared to base steel. It can be explained by the result of thermocalc calculation as shown in Fig. 9. In the thermo-calc calculation, this study considered all of the phases such as ferrite, austenite, (Ti, Nb)C, liquid, cementite and (Mo, Mn)2 C, together with all components. However, the calculation results of carbide are only plotted
Fig. 5. TEM micrographs of (Ti, Nb)C particle in the base metal of steel A and EDS analysis.
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J. Moon et al. / Materials Science and Engineering A 497 (2008) 153–159
Fig. 6. Optical micrographs in the weld CGHAZ of experimental steels, ×1000.
in Fig. 9. In this study, peak temperature in the CGHAZ is 1350 ◦ C which is higher than dissolution temperature of (Ti, Nb)C as shown in Fig. 9 and it means that whole (Ti, Nb)C particle can be dissolved during heating in the CGHAZ if an enough time was given at high temperature. However, thermal cycle of the CGHAZ in this study is very fast and thus some carbide can remain in the CGHAZ as shown in Fig. 8. 3.3. Precipitation and mechanical properties in as-tempered steels
Fig. 7. Mechanical properties of base steels and as-welded steels: (a) Vickers hardness, (b) charpy V-notch toughness.
To improve the mechanical properties of the CGHAZ, this study carried out the tempering with various temperatures after the CGHAZ simulation, as shown in Fig. 2. Fig. 10 shows the representative TEM images of tempered specimens (30 min at 550 ◦ C) and indicates that each tempered steel consists of tempered martensite microstructure with fine lath. Fig. 11 is the results of Vickers hardness test of as-tempered steels. Hardness initially decreased with increasing tempering temperature and then was reinforced at high tempering temperature above 400 ◦ C. This can be explained by the precipitation of Cu particle and carbide as follows. Previous works have shown that the hardness of tempered lath martensite can be improved by the precipitation of Cu particles and Mo2 C [1], where particles begin to precipitate at above 350 ◦ C [3]. Steels used in this study contain Mo and/or Cu as shown in Table 1. Therefore, reinforcement of tempered steels at high temperature in Fig. 11 is due to the precipitation of Cu particle or Mo2 C [3]. Actually, it has been found that Cu particles in steels C and F, which have 1.5 wt%Cu, precipitated after tempering. Fig. 12 is the Cu particle with 9R fringe pattern in tempered C alloy. Recent studies using high resolution electron microscopy have shown that Cu particles of twinned 9R structure with internal fringe are formed due to stacking fault by shear strain, before the transformation of -Cu particle of FCC structure. It has been found that 9R Cu particle has orientation relationship to the BCC matrix [7,16]. This study has considered that lattice condition of tempered martensite is nearly same with that of ferrite matrix by ejection of saturated carbon from martensite lattice during thermal aging. Thus the orientation relationship between 9R Cu particle and ferrite matrix in Fig. 11(b) has been analyzed as
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Fig. 8. TEM micrographs of (Ti, Nb)C in the CGHAZ of D alloy.
Fig. 9. Equilibrium volume fraction of carbides calculated by thermo-calc with increasing temperature.
follows: ¯ 9R //(1 1 0)Ferrite , (1 1 4)
[1 1¯ 0]9R //[1 1¯ 1]Ferrite
It is well matched with previous report [3,7]. Meanwhile, hardness of A and D alloy, which does not include Cu, was also improved
Fig. 11. Vickers hardness with increasing tempering temperature of experimental steels.
during tempering at 550 ◦ C. It is due to the precipitation of carbide [1], such as (Ti, Nb)C and (Mo, Mn)2 C, during tempering as shown in Fig. 9. Actually, many (Mo, Mn)2 C and (Ti, Nb)C were observed in tempered specimen of D alloy in Fig. 13. As a result, hardness of whole experimental specimens was improved by the precipitation
Fig. 10. Representative TEM micrographs in tempered specimens at 550 ◦ C: (a) C alloy and (b) F alloy.
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Fig. 12. Representative TEM micrographs of 9R Cu particle in tempered specimen in C alloy after 30 min aging at 550 ◦ C: (a) TEM bright field image, (b) FFT pattern analysis and (c) EDS analysis.
of Cu particle and/or carbide during tempering at high temperature in Fig. 11. In addition, hardness of Cu-added steels is higher than that of Cu-free steels and it is because that Cu added steels include the Cu particle as well as carbides. Fig. 14 indicates the variation of toughness of as-tempered steels with increasing temperature and shows the opposite tendency to the hardness profile in Fig. 11.
That is, toughness unexpectedly was not improved at high temperature despite of tempering treatment. Previous work has shown that small particles in thermally aged steel [17] make an embrittlement in tempered alloy. Therefore, toughness decrease in Fig. 14 is due to the embrittlement by precipitation of Cu particle and carbides.
Fig. 13. TEM micrographs of carbides in D alloy after tempering: (Ti, Nb)C and (Mo, Mn, Si)2 C in tempered specimen after 30 min aging at 550 ◦ C.
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(3) In tempered steels, Vickers hardness initially decreased with increasing tempering temperature and was reinforced at high temperature. TEM analysis showed that 9R Cu particles and carbides were precipitated during aging and then the reinforcement is mainly due to the precipitation of Cu particle and carbides during aging. While toughness tends to show opposite behavior and it was not improve at high temperature despite of tempering treatment. This result indicates that Cu addition is not useful choice to improve the toughness of the heat-affected zone. Acknowledgement This study was supported by a grant from the Fundamental R&D Program for the Core Technology of Materials funded by the Ministry of Commerce, Industry and Energy, Republic of Korea. References Fig. 14. Charpy V-notch toughness with an increasing tempering temperature of experimental steels.
4. Conclusion remarks (1) In the base steels, Cu particles and Mo2 C except (Ti, Nb)C were not found despite for high deformation by high pressure rolling during fabrication. This is related to the slow precipitation kinetics of them. In the case of B added steels, TEM results showed that the fine martensite microstructure was formed with an addition of boron, which plays a role in making prior austenite grain boundary to be stable and thus increases the hardenability. (2) Vickers hardness of base steels and the CGHAZ increased with increasing Cu and B content. This is due to the solid-solution hardening. While charpy V-notch toughness showed an opposite tendency, which is mainly due to the formation of the hard phase by increasing hardenability with Cu and B addition. Meanwhile, toughness in the CGHAZ is deteriorated as compared to that of the base steels.
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