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Toughness investigation on simulated weld HAZs of SQV-2A pressure vessel steel
Nuclear Engineering and Design 183 (1998) 9 – 20
Toughness investigation on simulated weld HAZs of SQV-2A pressure vessel steel Jinsun Liao 1,a, Kenji Ikeuchi a,*, Fukuhisa Matsuda b b
a Joining and Welding Research Institute, Osaka Uni6ersity, 11 -1 Mihogaoka, Ibaraki, Osaka 567, Japan Japan Power Engineering and Inspection corporation, Shin-Toranomon Building, 1 -5 -11 Akasaka, Minato-ku, Tokyo 107, Japan
Received 1 April 1997; received in revised form 12 May 1997; accepted 5 December 1997
Abstract In order to get detailed information about weld HAZs toughness of SQV-2A steel and determine the optimum welding and heat treatment parameters, the toughness of simulated CGHAZs (coarse grained heat affected zone) and CGHAZs (intercritically reheated CGHAZ) were systematically investigated. The influence of tempering thermal cycles on weld ICCGHAZs toughness was clarified. The effect of post weld heat treatments (PWHT) on weld CGHAZs toughness was also determined. The results showed that high toughness (absorbed energy \ 200 J) of weld HAZs could be achieved by selecting the optimum welding and PWHT parameters (cooling time Dt8/5: 6–40 s, PWHT: 893 K, 3.6–7.2 ks). Tempering thermal cycles with peak temperature of above 573 K could remarkably improve the toughness of deteriorated ICCGHAZs and reduce the hardness, when cooling time Dt8/5(2) of the reheating thermal cycle was 6 s, which implies that welding of SQV-2A without PWHT is possible, provided that low heat input welding is adopted and welding procedure is correctly arranged. Metallography and fractography revealed that M–A constituents in weld HAZs played an important role in controlling weld HAZ toughness. © 1998 Elsevier Science S.A. All rights reserved.
1. Introduction Pressure vessels are always fabricated from steels with excellent mechanical properties, of which toughness is one of the most important factors for ensuring the safety of the pressure vessels. For this reason, low alloy steel SQV-2A possessing superior low-temperature toughness and weldability is widely employed in building * Corresponding author. Tel.: + 81 6 8798644; fax: + 81 6 8798689. 1 Manufacturing Department, Sumiyoshi Factory Kurimoto Ltd., 8-45 Shibatani 2-chome, Suminoe-ku, Osaka 559, Japan.
pressure vessels. However, the superior toughness of SQV-2A base metal will be altered after experiencing thermal cycles imposed by welding process. Many studies have showed that the loss in toughness always happens in weld heat affected zone (HAZ) (Sato and Yamato, 1981; Tsuboi and Hirai, 1981). Several catastrophic service failures have been directly attributed to a lack of adequate notch toughness resulting from improper welding practice. To obtain a safe welded joint, therefore, it is necessary to know the effect of welding thermal cycles on weld HAZ toughness and determine the proper welding conditions.
0029-5493/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0029-5493(98)00151-4
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Table 1 Chemical composition of SQV–2A steel Material
SQV-2A
Chemical composition (mass%) C Si Mn
P
S
Ni
Mo
0.19
B0.01
B0.01
0.62
0.56
0.24
1.48
Not all regions of weld HAZ experience an equivalent decrease in toughness for a given set of welding conditions. The great damage is showed to occur in coarse grained heat affected zone (CGHAZ) which is the region immediately adjacent to the fusion zone where peak temperatures approach the melting point (Nippes et al., 1957; Nakao et al., 1985). Besides, during multi-pass welding, a part of the CGHAZ of a former weld pass can be intercritically reheated by subsequent thermal cycles to a/g two phase region when the peak temperature of the subsequent thermal cycle ranges between Ac1 and Ac3. This intercritically reheated CGHAZ (ICCGHAZ) has been reported to be another brittle zone in weld HAZ, and display even lower toughness than that of CGHAZ (Kim et al., 1991; Linton and King, 1992). The SQV-2A steel plate used for pressure vessels is very thick, so that the increase in weld heat input is desired in order to improve welding productivity. However, Savage and Owczarski (1966), Ohno et al. (1972) indicated that weld CGHAZ toughness decreases with increasing weld heat input for most low alloy steels. For SQV-2A steel, although it can be inferred that the weld CGHAZ toughness may also be deteriorated with increasing weld heat input like other low alloy steels, it is important to know the maximum heat input allowable for a satisfactory weld HAZ toughness, because it is useful in determining the limitation which must be placed on welding procedures to assure adequate weld HAZ toughness. For a thick steel plate, multi-pass welding is also unavoidable. The multi-pass welding leads to the formation of ICCGHAZs in welded joint. Davis and King (1993) found that ICCGHAZ toughness was related to the peak temperature of reheating thermal cycle, i.e. the second thermal cycle. When the peak temperature of the second
thermal cycle fell in a certain temperature range between Ac1 and Ac3, ICCGHAZs of some low alloy steels displayed a very low toughness. For SQV-2A steel, little information has been known about the ICCGHAZ toughness. To get a comprehensive understanding of whole weld HAZ toughness of SQV-2A steel, the characteristics of weld ICCGHAZs with various peak temperatures of reheating thermal cycle should also be determined. During multi-pass welding, the weld ICCGHAZs will be tempered by the following weld pass thermal cycles. The effect of tempering thermal cycles on the ICCGHAZs toughness has not been understood very well. Furthermore, the post weld heat treatment is usually performed to eliminate the residual stress after the welding of pressure vessels. It is also necessary to determine the effect of PWHT on the weld HAZ toughness of SQV-2A steel. In this paper, therefore, the toughness of weld CGHAZs with various heat inputs and that of ICCGHAZs with various peak temperatures of reheating thermal cycle were firstly investigated. Then the effects of tempering thermal cycles on weld ICCGHAZs toughness was clarified. The effect of PWHT on weld CGHAZs toughness was determined. On the basis of metallography and fractography, the factor controlling weld HAZs toughness was discussed. The purpose is to get detailed information about weld HAZ toughness of SQV-2A steel and provide a guidance for the correct choice of welding parameters. 2. Experimental details The chemical composition of SQV-2A pressure vessel steel is shown in Table 1. The as-received material was a quenched and tempered steel plate. The plate thickness (t) was 200 mm.
J. Liao et al. / Nuclear Engineering and Design 183 (1998) 9–20
Thermal/mechanical simulator Gleeble 1500 was used to simulate the weld CGHAZs, ICCGHAZs and tempering process. Samples (55× 10×10 mm) for the simulation were cut from the as-received steel plate at positions of 1/4t or 3/4t with the longest side of the samples parallel to the rolling direction. Thermal cycles used for simulation are schematically shown in Fig. 1. Single thermal cycle was used for the simulation of weld CGHAZs (Fig. 1a). The peak temperature Tp of the thermal cycle was 1623 K. Cooling time from 1073 to 773 K (Dt8/5) was varied from 6 to 1000 s to simulate weld CGHAZs with various weld heat inputs, because the cooling time Dt8/5 is determined by the heat input for a given welding process. Double thermal cycles were used for the simulation of weld ICCGHAZs (Fig. 1b). The first thermal cycle was
Fig. 1. Thermal cycles used for simulation of weld CGHAZs, ICCGHAZs and the tempering process.
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used to simulate a weld CGHAZ with a low heat input, and the second thermal cycle was used to simulate intercritically reheating thermal cycles. Peak temperature Tp2 of the second thermal cycle was varied from 873 to 1173 K covering the whole range of Ac1 –Ac3. For this steel, the Ac1 and Ac3 temperatures were determined under the Gleeble simulation condition to be about 945 and 1107 K, respectively. It should also be noted that microstructures of measured materials have influence on Ac1, and thus the Ac1 is different for as-received material and for the weld CGHAZ because of their different microstructures. Cooling time Dt8/5(2) of the second thermal cycle was chosen to be 6 and 40 s. Triple thermal cycles were used to simulate the ICCGHAZs and tempering thermal cycles (Fig. 1c). The first and second thermal cycles were used to simulate the ICCGHAZs with deteriorated toughness, and the third thermal cycle was used to simulate the tempering thermal cycles. For all thermal cycles, heating rate was 111 K s − 1, and holding time at peak temperatures was 6 s. PWHT was performed on the CGHAZs simulated with single thermal cycle. PWHT temperature was 893 K, and PWHT time was varied from 3.6 to 28.8 ks. Full size standard Charpy-V notch specimens were prepared from the samples subjected to weld HAZs simulation and PWHT. Charpy impact test was carried out at 293 K. The toughness was evaluated from absorbed energy 6E. Both scanning electron microscope (SEM) and transmission electron microscope (TEM) were used to characterize the microstructures of simulated weld HAZs. A two-stage electrolytic etching method was used to reveal the M–A constituent, by which the ferrite and carbide were etched preferentially in the 1st and 2nd stages, respectively. After two-stage electrolytic etching, the M–A constituent and carbide appear as light and dark phases in the ferrite matrix under SEM observation, respectively. In addition, fracture surface was examined with SEM. Hardness measurement was performed with a Vickers hardness tester.
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Fig. 2. Absorbed energy of weld CGHAZs as a function of cooling time Dt8/5.
3. Results and discussion
lem. Therefore, the heat treatments such as preheating and PWHT are necessary to eliminate the danger of cold cracking and improve the toughness further. Typical microstructures of weld CGHAZs with various cooling times Dt8/5 are shown in Fig. 4. For the cooling times Dt8/5 in the range of 6–8 s, the microstructure consisted mainly of martensite and lower bainite. Upper bainite was observed from 12 s of Dt8/5, and the M–A constituent was observed from 20 s of Dt8/5. When cooling time Dt8/5 was 40 s, microstructures consisted of lower bainite, upper bainite, M–A constituent and bainitic ferrite. With further increasing cooling time, the amount of lower bainite decreased, and those of upper bainite, M–A constituent and bainitic ferrite increased. When Dt8/5 was over 150 s, microstructures were almost upper bainite, M– A constituent and bainitic ferrite. Proeutectoid ferrite was not observed for Dt8/5 B 1000 s.
3.1. Weld CGHAZ toughness and microstructure The absorbed energies 6E of weld CGHAZs simulated with single thermal cycle are shown in Fig. 2, as a function of cooling time Dt8/5. With increasing cooling time Dt8/5, i.e. with increasing weld heat input, the absorbed energy decreased generally, although it increased slightly with increasing Dt8/5 from 6 to 8 s. When Dt8/5 was 20 s or longer, the remarkable decrease of absorbed energy was observed. The relation between hardness of weld CGHAZ and cooling time Dt8/5 is given in Fig. 3. Hardness decreased monotonously with increasing cooling time Dt8/5 from 6 to 1000 s. From the results above, it can be understood that when Dt8/5 was 12 s or shorter, the absorbed energies of weld CGHAZs were above 190 J, showing high toughness in weld CGHAZs. When Dt8/5 was 20 s or longer, however, the toughness deteriorated drastically. These results imply that the weld heat-input should be restricted to limit the cooling time D8/5 shorter than 12 s, in order that the weld CGHAZs have high toughness. However, when Dt8/5 was 12 s or shorter, the hardness was very high (higher than 400 HV), suggesting that cold cracking may become a prob-
3.2. Weld ICCGHAZ toughness and microstructure The absorbed energies of weld ICCGHAZs simulated with double thermal cycles are shown in Fig. 5, as a function of peak temperature Tp2 of the second thermal cycle. It could be seen that the
Fig. 3. Relation between hardness of CGHAZs and cooling time Dt8/5.
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Fig. 4. Typical microstructures of CGHAZs with various cooling times Dt8/5.
absorbed energies of the impact test were strongly influenced by both the peak temperature Tp2 and cooling time Dt8/5(2). When Tp2 was in the range from 873 to 923 K (below Ac1), the absorbed energy was much higher than that of weld CGHAZ simulated with single thermal cycle. This high toughness resulted from the tempering effect of the second thermal cycle on the martensitic structure of the CGHAZ. With increasing Tp2 from 923 to 973K (above Ac1), the absorbed energy decreased drastically for both cooling times Dt8/5(2) of 6 and 40 s. For the peak temperatures Tp2 \1003 K, the absorbed energies were closely related to the cooling time Dt8/5(2). For Dt8/5(2) of 6 s, the absorbed energy increased with further increasing Tp2 up to 1073K; but for Dt8/5(2) of 40 s, the absorbed energy decreased further with the increase of Tp2 from 923 to 1073 K. Only a slight increase in absorbed energy was observed for Tp2 higher than Ac3.
The relation between Vickers hardness of weld ICCGHAZs and the peak temperature Tp2 is given in Fig. 6. For Tp2 from 923 to 1003 K, hardness of weld ICCGHAZs was considerably lower than that of the CGHAZ simulated with single thermal cycle due to tempering effect. When Tp2 was above 1023 K, however, the hardness increased again. This phenomenon implied that when the Tp2 was \ 1023 K, most of prior structures were reaustenitized. Upon fast cooling, martensite formed; but upon slow cooling, structures consisting of bainite and martensite formed. The hardness of the martensitic structure or the complex structure of bainite and martensite was probably higher than that of a tempered martensitic structure, which might be the dominant structure for Tp2 B 1023 K. Microstructures of weld ICCGHAZs are revealed in Figs. 7 and 8 for Dt8/5(2) of 6 and 40 s, respectively.
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Fig. 5. Absorbed energy of ICCGHAZs as a function of peak temperature Tp2 of the second thermal cycle (Ac1 for as received material; (Ac1)% for CGHAZ of SQV-2A).
For both cooling times Dt8/5(2) of 6 and 40 s, the peak temperature Tp2 lower than (Ac1)’ was insufficient to cause reaustenitization. Carbides precipitated because of the tempering effect of the second thermal cycle on the martensitic structure of the CGHAZ, which resulted in superior toughness. In the case of Dt8/5(2) of 6 s, for peak temperatures Tp2 in the range of 923 – 1003 K, island-like
Fig. 7. SEM micrographs of ICCGHAZs (Dt8/5(2) =6 s).
Fig. 6. Relation between hardness of ICCGHAZs and Tp2 of the second thermal cycle.
M–A constituents were observed in the interior of the grains, and necklace-like M–A constituents at the grain boundary regions. For Tp2 from 1023 to 1123 K, reaustenitization occurred in large area. Upon cooling, martensite formed in this reaustenitized area. At the same time, untransformed areas were observed as dark areas in the interior of grains. In these untransformed areas, the prior structure was tempered by the second thermal cycle, thus carbides precipitated. When reheated over Ac3, prior structures were completely
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reaustenitized and transformed to martensite again after being cooled, as can be seen in Fig. 7. In the case of Dt8/5(2) of 40 s, M – A constituents also formed in the interior of the grain and at the grain boundary region for Tp2 from 923 to 1023 K. For peak temperatures Tp2 \1073 K, reaustenitization occurred in large area. Upon cooling, microstructures of upper bainite and lower bainite formed in the interior of the grain, and small M – A constituents at the grain boundary region because of the low cooling rate,
Fig. 9. Relation between Tp3 and absorbed energy of tempered ICCGHAZs.
as shown in Fig. 8.
3.3. Effect of tempering thermal cycle
Fig. 8. SEM micrographs of ICCGHAZs (Dt8/5(2) =40 s).
It can be known from Fig. 5 that when Tp2 is 973–1003K, the toughness deterioration occurs in ICCGHAZs irrespective of cooling time Dt8/5(2). When Tp2 is \ 1023 K, the toughness is also deteriorated in the ICCGHAZs for cooling time Dt8/5(2) of 40 s. In this section, investigation was taken to determine the effect of the third thermal cycle (i.e. tempering thermal cycle) on the toughness improvement in these deteriorated ICCGHAZs at peak temperatures Tp2 of 973 and 1073 K. Fig. 9 shows the relationship between peak temperature Tp3 of the tempering thermal cycle and the absorbed energy of the tempered ICCGHAZs. It can be seen that the changing tendency of the absorbed energy with Tp3 depends strongly on peak temperature Tp2 of the second thermal cycle. When Tp2 was at 973 K, the absorbed energies increased with increasing Tp3 of the third thermal cycle to 673 K. Further increasing Tp3 to 893 K resulted in the decrease of absorbed energy. When Tp2 was 1073 K, the absorbed energies increased monotonously with increasing Tp3. In addition, it can also be seen that the absorbed energies are strongly influenced by
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cooling time Dt8/5(2) of the second thermal cycle. When Dt8/5(2) was 6 s, the absorbed energies increased remarkably. When Dt8/5(2) was 40 s, however, the increase in absorbed energy was not as great as that for Dt8/5(2) of 6 s. The relation between peak temperature Tp3 and the hardness of the tempered ICCGHAZs is given in Fig. 10. The hardness of the ICCGHAZ was obviously reduced by tempering thermal cycles only when Tp2 was 1073 K and Dt8/5(2) was 6 s. When the peak temperature Tp3 of the tempering thermal cycles was \573 K, the hardness of all tempered ICCGHAZs was B350 HV. The above results demonstrate that when low heat input welding is adopted (for example, when Dt8/5(2) is 6 s), the deteriorated ICCGHAZs (e.g. at peak temperature Tp2 of 973 K) can be improved to have absorbed energy of above 200 J, provided that the Tp3 is controlled to be above 573 K. Although the absorbed energy decreases when Tp3 is increased to 893 K, the absorbed energy is still about 200 J, at the same level of the weld CGHAZ. However, when comparatively high heat input is adopted (for example, when Dt8/5(2) is 40 s), the toughness improvement by the tempering thermal cycle is not so apparent. Therefore, low heat input is necessary to the toughness improvement in ICCGHAZs by tempering thermal cycles.
Fig. 10. Relation between Tp3 and hardness of tempered ICCGHAZs.
Fig. 11. Effect of Tp3 on the decomposition of M – A constituents.
On the other hand, when low heat input is chosen (for example, when Dt8/5(2) is 6 s), the hardness of ICCGHAZs can be reduced below 350 HV, provided that Tp3 of the tempering thermal cycle is \ 573 K. Moreover, when Tp2 of the second thermal cycle is over Ac3, the high hardness (see Fig. 6) of martensitic structures can also be reduced by tempering thermal cycles to be B 350 HV, provided that the tempering temperature is high enough. That is to say, in the case of no PWHT, weld HAZs of high toughness (absorbed energy\ 200 J) and low hardness (B350 HV) can be achieved, when low heat input welding is adopted and welding procedure is correctly arranged so that the thermal cycles of following weld pass have tempering effects on the former weld pass.
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The microstructures of tempered ICCGHAZs (at Tp2 of 973 K) are given in Fig. 11. When Tp3 was 473–523 K, M – A constituents were decomposed partially. Increasing Tp3 to 573 – 673 K resulted in apparent decomposition of necklace-like M–A constituents. However, when Tp3 was further increased to 773 – 893 K, the decomposition of M–A constituent seems to be difficult. Small M–A constituents could be observed at the edge of original necklace-like M – A constituents. From Figs. 9 and 11, it can be found that the decomposition of M – A constituent corresponds to the increase of absorbed energy in ICCGHAZs.
3.4. Effect of PWHT Fig. 13. Effect of PWHT time on the hardness of CGHAZs.
The absorbed energies and hardness of weld CGHAZs after PWHT at 893 K are shown in Figs. 12 and 13, respectively, as a function of PWHT time. For cooling times Dt8/5 of 6, 40 and 150 s, the changing trends of absorbed energy with PWHT time were similar. The absorbed energies increased with PWHT time, and displayed the maximum value when PWHT time was 7.2 ks. Increasing PWHT time to 14.4 ks led to decrease in absorbed energy. The absorbed energies increased again with further increasing PWHT time to 28.8 ks. In addition, it could be seen that the absorbed energy of the CGHAZ was remarkably increased to be \ 300 J when cooling
time Dt8/5 was 6 s. When cooling time Dt8/5 was 40 s, the absorbed energy was also improved above 200 J after PWHT at 893 K for 3.6–7.2 ks. On the other hand, the hardness decreased markedly with increasing PWHT time to 7.2 ks for Dt8/5 of 6 s. For Dt8/5 of 40 and 150 s, decrease in hardness was not so great. When the PWHT time was 14.4 ks, a slight increase in hardness could be observed, comparing with the case when PWHT time was 7.2 ks. The changing trend of hardness implies that secondary hardening probably occurred when PWHT time was 14.4 ks, thus the toughness decreased slightly (see Fig. 12). From above results, it can be concluded that (i) when cooling times Dt8/5 are in the range of 6 to 40 s, weld CGHAZs toughness can be improved above 200 J, and hardness be reduced below 300 HV by PWHT, (ii) 893 K, 3.6–7.2 ks is the optimum PWHT condition from the view of toughness improvement.
3.5. Factor controlling weld HAZ toughness
Fig. 12. Effect of PWHT time on the absorbed energy of CGHAZs.
The toughness of weld CGHAZs decreased with increasing cooling time Dt8/5 (see Fig. 2). When Dt8/5 was 12 s or shorter, the absorbed energies of weld CGHAZs were above 190 J, showing high toughness in weld CGHAZs. When Dt8/5 was 20 s or longer, the toughness decreased drastically. This remarkable decrease in toughness
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was accompanied by the formation of M–A constituents in CGHAZs (see Fig. 4). The toughness of weld ICCGHAZs depended on not only peak temperature Tp2 but also cooling time Dt8/5(2) of the second thermal cycle. For Dt8/5(2) of 6 s, the toughness deterioration was observed when peak temperature Tp2 was in the range of 973 – 1023 K. For cooling time Dt8/5(2) of 40 s, the toughness deterioration was observed over a wider range of Tp2, i.e. 973 – 1173 K (see Fig. 5). The loss of toughness in ICCGHAZs was also accompanied by the formation of M–A constituents (see Figs. 7 and 8). Therefore, it can be concluded that the toughness deterioration in weld HAZs is mainly attributed to the formation of M – A constituents. The M–A constituent in weld HAZs is the main factor controlling weld HAZ toughness. By decomposing M – A constituents, toughness can be improved (see Figs. 9 and 11). To confirm the existence of M – A constituent in deteriorated weld HAZs, TEM microscopy was carried out on the thin foils prepared from the sample of toughness deteriorated CGHAZs and ICCGHAZs. Fig. 14 reveals TEM micrographs of M–A constituents in the CGHAZs and ICCGHAZs of SQV-2A steel. M – A constituents can be believed to exist in deteriorated weld HAZs of SQV-2A steel. The M – A constituent consists mainly of twin martensite. Carbon concentration in M – A constituents is very high and increases with cooling time Dt8/5 (Liao, 1995). With the increase of carbon content in M–A constituents, the hardness of M –A constituents increases (Okada, 1993). Generally, the hardness of M – A constituent is \ 600 HV, thus M–A constituents can likely become initiative sites of cleavage fracture. SEM observation on the surface of fractured specimens confirmed the cracking initiating role of M – A constituents. M– A constituents were observed at the initiative locations of the cleavage fracture, as can be seen in Fig. 15. The observation of microstructures near the fracture surface of Charpy impacted specimen subjected to the ICCGHAZ simulation also showed that microcracks initiated between the necklace-like M – A constituent and matrix, and
propagate into the interior of the grain, or along the grain boundary, as shown in Fig. 16. Therefore, it can be believed that the formation of M–A constituent is responsible for the toughness loss in weld HAZs of SQV-2A steel.
4. Summary and conclusion The toughness of weld CGHAZs with various heat inputs and that of ICCGHAZs with various peak temperatures of reheating thermal cycle were systematically investigated. The effect of tempering thermal cycles on the weld ICCGHAZs toughness was clarified. The effect of PWHT on weld CGHAZs toughness was determined. On the basis of metallography and fractography, the role of M–A constituents in controlling weld HAZs
Fig. 14. TEM micrographs of M – A constituents in (a) CGHAZs and (b) ICCGHAZs.
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1173 K in the case of Dt8/5(2) = 40 s. Tempering thermal cycles with peak temperature Tp3 of above 573 K can remarkably improve the toughness of deteriorated ICCGHAZs and reduce the hardness, when cooling time Dt8/5(2) is 6 s. These results imply that welding of SQV-2A steel without PWHT is possible, provided that low heat input welding is adopted and welding procedure is correctly arranged. (3) The formation of M–A constituent is responsible for the toughness deterioration in weld HAZs. (4) Based on the above results, the optimum welding and PWHT parameters are recommended as follows:
Fig. 15. SEM micrographs of initiation site of a cleavage fracture ((a) low magnification and (b) high magnification).
toughness was discussed. The results obtained in this study can be summarized as follows: (1) The toughness of weld CGHAZs of SQV2A steel decreases generally with increasing cooling time Dt8/5, namely with increasing weld heat input. When cooling times Dt8/5 are in the range of 6–12 s, the toughness of weld CGHAZs displays high values, but the hardness is too high (\400 HV). Suitable PWHT can significantly improve the toughness and reduce the hardness of weld CGHAZs. When Dt8/5 are in the range of 6–40 s, the absorbed energy can be increased above 200 J, and hardness be reduced below 300 HV after PWHT at 893 K for 3.6 – 7.2 ks. (2) Toughness deterioration occurs in ICCGHAZs, when peak temperature of the second thermal cycle Tp2 =973 – 1023 K in the case of cooling time Dt8/5(2) =6 s, and when Tp2 = 973–
Fig. 16. Microcracks near grain boundary of ICCGHAZs ((a) showing microcracks near grain boundaries and (b) showing cleavage fracture along grain boundary).
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1. When PWHT is carried out, (a) cooling time Dt8/5: 6– 40 s; (b) PWHT: 893 K, 3.6 –7.2 ks. 2. In the case of no PWHT, (a) cooling time Dt8/5: 6– 8 s; (b) peak temperature of the tempering thermal cycle should be controlled in the range 573 – 893 K so that deteriorated ICCGHAZs can be improved by the tempering thermal cycles. Above thermal cycles can result in weld HAZs of high absorbed energy (\ 200 J) and low hardness (B 350 HV).
References Davis, C.L., King, J.E., 1993. Effect of cooling rate on intercritically reheated microstructure and toughness in high strength low alloy steel. Mater. Sci. Technol. 9 (1), 8 – 15. Kim, B.C., Lee, S., Kim, N.J., Lee, D. Y., 1991. Microstructure and Local brittle zone phenomena in high-strength low alloy steel welds. Metall. Trans. A 22A (1), 139 – 149. Liao, J., 1995. Influence of M–A constituent on weld HAZ toughness low alloy steel SQV-2A for pressure vessels. PhD thesis, Tianjin University.
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Linton, V.M., King, J.E., 1992. The effect of heat input on the heat affected zone toughness of two offshore steels. Proc. 3rd Int. Conf. on Trend in Welding Research, Gatlinburg, TN, June 1 – 5 pp. 345 – 349. Nakao, Y., Oshige, H., Noi, S., 1985. Distribution of toughness in HAZ of multi-pass welded high strength steel. Q. J. Jpn. Weld. Soc. 3 – 4, 773 – 781 [in Japanese]. Nippes, E.F., Savage, W.F., Allio, R.J., 1957. Studies of the weld heat-affected zone of T-1 steel. Weld. J., 36 (12), 531s – 540s. Okada, H., 1993. Behavior of M – A constituent in large-heatinput weld HAZs of 780 – 980 MPa class HSLA steels, PhD thesis, Osaka University. Ohno, Y., Habu, R., Sekino, S., 1972. Thermal cycle test of high strength steel and steel plates for low temperature service. Tetsu To Hagane 58 (2), 306 – 316 [in Japanese]. Sato, M., Yamato, K., 1981. On the microstructure and toughness of HAZ in as rolled or normalized 50 and 60 Kg mm − 2 high strength steel. J. Jpn. Weld. Soc. 50 (1), 11–19 [in Japanese]. Savage W.F., Owczarski, W.A., 1966. The microstructure and notch impact behavior of a welded structure steel. Weld. J., 45 (2), 55s – 65s. Tsuboi, J., Hirai, Y., 1981. Microstructure and toughness of weld heat affected zone in quenched-tempered high strength steels. J. Jpn. Weld. Soc. 50 (1), 28 – 37 [in Japanese].
Toughness investigation on simulated weld HAZs of SQV-2A pressure vessel steel Jinsun Liao 1,a, Kenji Ikeuchi a,*, Fukuhisa Matsuda b b
a Joining and Welding Research Institute, Osaka Uni6ersity, 11 -1 Mihogaoka, Ibaraki, Osaka 567, Japan Japan Power Engineering and Inspection corporation, Shin-Toranomon Building, 1 -5 -11 Akasaka, Minato-ku, Tokyo 107, Japan
Received 1 April 1997; received in revised form 12 May 1997; accepted 5 December 1997
Abstract In order to get detailed information about weld HAZs toughness of SQV-2A steel and determine the optimum welding and heat treatment parameters, the toughness of simulated CGHAZs (coarse grained heat affected zone) and CGHAZs (intercritically reheated CGHAZ) were systematically investigated. The influence of tempering thermal cycles on weld ICCGHAZs toughness was clarified. The effect of post weld heat treatments (PWHT) on weld CGHAZs toughness was also determined. The results showed that high toughness (absorbed energy \ 200 J) of weld HAZs could be achieved by selecting the optimum welding and PWHT parameters (cooling time Dt8/5: 6–40 s, PWHT: 893 K, 3.6–7.2 ks). Tempering thermal cycles with peak temperature of above 573 K could remarkably improve the toughness of deteriorated ICCGHAZs and reduce the hardness, when cooling time Dt8/5(2) of the reheating thermal cycle was 6 s, which implies that welding of SQV-2A without PWHT is possible, provided that low heat input welding is adopted and welding procedure is correctly arranged. Metallography and fractography revealed that M–A constituents in weld HAZs played an important role in controlling weld HAZ toughness. © 1998 Elsevier Science S.A. All rights reserved.
1. Introduction Pressure vessels are always fabricated from steels with excellent mechanical properties, of which toughness is one of the most important factors for ensuring the safety of the pressure vessels. For this reason, low alloy steel SQV-2A possessing superior low-temperature toughness and weldability is widely employed in building * Corresponding author. Tel.: + 81 6 8798644; fax: + 81 6 8798689. 1 Manufacturing Department, Sumiyoshi Factory Kurimoto Ltd., 8-45 Shibatani 2-chome, Suminoe-ku, Osaka 559, Japan.
pressure vessels. However, the superior toughness of SQV-2A base metal will be altered after experiencing thermal cycles imposed by welding process. Many studies have showed that the loss in toughness always happens in weld heat affected zone (HAZ) (Sato and Yamato, 1981; Tsuboi and Hirai, 1981). Several catastrophic service failures have been directly attributed to a lack of adequate notch toughness resulting from improper welding practice. To obtain a safe welded joint, therefore, it is necessary to know the effect of welding thermal cycles on weld HAZ toughness and determine the proper welding conditions.
0029-5493/98/$19.00 © 1998 Elsevier Science S.A. All rights reserved. PII S0029-5493(98)00151-4
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Table 1 Chemical composition of SQV–2A steel Material
SQV-2A
Chemical composition (mass%) C Si Mn
P
S
Ni
Mo
0.19
B0.01
B0.01
0.62
0.56
0.24
1.48
Not all regions of weld HAZ experience an equivalent decrease in toughness for a given set of welding conditions. The great damage is showed to occur in coarse grained heat affected zone (CGHAZ) which is the region immediately adjacent to the fusion zone where peak temperatures approach the melting point (Nippes et al., 1957; Nakao et al., 1985). Besides, during multi-pass welding, a part of the CGHAZ of a former weld pass can be intercritically reheated by subsequent thermal cycles to a/g two phase region when the peak temperature of the subsequent thermal cycle ranges between Ac1 and Ac3. This intercritically reheated CGHAZ (ICCGHAZ) has been reported to be another brittle zone in weld HAZ, and display even lower toughness than that of CGHAZ (Kim et al., 1991; Linton and King, 1992). The SQV-2A steel plate used for pressure vessels is very thick, so that the increase in weld heat input is desired in order to improve welding productivity. However, Savage and Owczarski (1966), Ohno et al. (1972) indicated that weld CGHAZ toughness decreases with increasing weld heat input for most low alloy steels. For SQV-2A steel, although it can be inferred that the weld CGHAZ toughness may also be deteriorated with increasing weld heat input like other low alloy steels, it is important to know the maximum heat input allowable for a satisfactory weld HAZ toughness, because it is useful in determining the limitation which must be placed on welding procedures to assure adequate weld HAZ toughness. For a thick steel plate, multi-pass welding is also unavoidable. The multi-pass welding leads to the formation of ICCGHAZs in welded joint. Davis and King (1993) found that ICCGHAZ toughness was related to the peak temperature of reheating thermal cycle, i.e. the second thermal cycle. When the peak temperature of the second
thermal cycle fell in a certain temperature range between Ac1 and Ac3, ICCGHAZs of some low alloy steels displayed a very low toughness. For SQV-2A steel, little information has been known about the ICCGHAZ toughness. To get a comprehensive understanding of whole weld HAZ toughness of SQV-2A steel, the characteristics of weld ICCGHAZs with various peak temperatures of reheating thermal cycle should also be determined. During multi-pass welding, the weld ICCGHAZs will be tempered by the following weld pass thermal cycles. The effect of tempering thermal cycles on the ICCGHAZs toughness has not been understood very well. Furthermore, the post weld heat treatment is usually performed to eliminate the residual stress after the welding of pressure vessels. It is also necessary to determine the effect of PWHT on the weld HAZ toughness of SQV-2A steel. In this paper, therefore, the toughness of weld CGHAZs with various heat inputs and that of ICCGHAZs with various peak temperatures of reheating thermal cycle were firstly investigated. Then the effects of tempering thermal cycles on weld ICCGHAZs toughness was clarified. The effect of PWHT on weld CGHAZs toughness was determined. On the basis of metallography and fractography, the factor controlling weld HAZs toughness was discussed. The purpose is to get detailed information about weld HAZ toughness of SQV-2A steel and provide a guidance for the correct choice of welding parameters. 2. Experimental details The chemical composition of SQV-2A pressure vessel steel is shown in Table 1. The as-received material was a quenched and tempered steel plate. The plate thickness (t) was 200 mm.
J. Liao et al. / Nuclear Engineering and Design 183 (1998) 9–20
Thermal/mechanical simulator Gleeble 1500 was used to simulate the weld CGHAZs, ICCGHAZs and tempering process. Samples (55× 10×10 mm) for the simulation were cut from the as-received steel plate at positions of 1/4t or 3/4t with the longest side of the samples parallel to the rolling direction. Thermal cycles used for simulation are schematically shown in Fig. 1. Single thermal cycle was used for the simulation of weld CGHAZs (Fig. 1a). The peak temperature Tp of the thermal cycle was 1623 K. Cooling time from 1073 to 773 K (Dt8/5) was varied from 6 to 1000 s to simulate weld CGHAZs with various weld heat inputs, because the cooling time Dt8/5 is determined by the heat input for a given welding process. Double thermal cycles were used for the simulation of weld ICCGHAZs (Fig. 1b). The first thermal cycle was
Fig. 1. Thermal cycles used for simulation of weld CGHAZs, ICCGHAZs and the tempering process.
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used to simulate a weld CGHAZ with a low heat input, and the second thermal cycle was used to simulate intercritically reheating thermal cycles. Peak temperature Tp2 of the second thermal cycle was varied from 873 to 1173 K covering the whole range of Ac1 –Ac3. For this steel, the Ac1 and Ac3 temperatures were determined under the Gleeble simulation condition to be about 945 and 1107 K, respectively. It should also be noted that microstructures of measured materials have influence on Ac1, and thus the Ac1 is different for as-received material and for the weld CGHAZ because of their different microstructures. Cooling time Dt8/5(2) of the second thermal cycle was chosen to be 6 and 40 s. Triple thermal cycles were used to simulate the ICCGHAZs and tempering thermal cycles (Fig. 1c). The first and second thermal cycles were used to simulate the ICCGHAZs with deteriorated toughness, and the third thermal cycle was used to simulate the tempering thermal cycles. For all thermal cycles, heating rate was 111 K s − 1, and holding time at peak temperatures was 6 s. PWHT was performed on the CGHAZs simulated with single thermal cycle. PWHT temperature was 893 K, and PWHT time was varied from 3.6 to 28.8 ks. Full size standard Charpy-V notch specimens were prepared from the samples subjected to weld HAZs simulation and PWHT. Charpy impact test was carried out at 293 K. The toughness was evaluated from absorbed energy 6E. Both scanning electron microscope (SEM) and transmission electron microscope (TEM) were used to characterize the microstructures of simulated weld HAZs. A two-stage electrolytic etching method was used to reveal the M–A constituent, by which the ferrite and carbide were etched preferentially in the 1st and 2nd stages, respectively. After two-stage electrolytic etching, the M–A constituent and carbide appear as light and dark phases in the ferrite matrix under SEM observation, respectively. In addition, fracture surface was examined with SEM. Hardness measurement was performed with a Vickers hardness tester.
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Fig. 2. Absorbed energy of weld CGHAZs as a function of cooling time Dt8/5.
3. Results and discussion
lem. Therefore, the heat treatments such as preheating and PWHT are necessary to eliminate the danger of cold cracking and improve the toughness further. Typical microstructures of weld CGHAZs with various cooling times Dt8/5 are shown in Fig. 4. For the cooling times Dt8/5 in the range of 6–8 s, the microstructure consisted mainly of martensite and lower bainite. Upper bainite was observed from 12 s of Dt8/5, and the M–A constituent was observed from 20 s of Dt8/5. When cooling time Dt8/5 was 40 s, microstructures consisted of lower bainite, upper bainite, M–A constituent and bainitic ferrite. With further increasing cooling time, the amount of lower bainite decreased, and those of upper bainite, M–A constituent and bainitic ferrite increased. When Dt8/5 was over 150 s, microstructures were almost upper bainite, M– A constituent and bainitic ferrite. Proeutectoid ferrite was not observed for Dt8/5 B 1000 s.
3.1. Weld CGHAZ toughness and microstructure The absorbed energies 6E of weld CGHAZs simulated with single thermal cycle are shown in Fig. 2, as a function of cooling time Dt8/5. With increasing cooling time Dt8/5, i.e. with increasing weld heat input, the absorbed energy decreased generally, although it increased slightly with increasing Dt8/5 from 6 to 8 s. When Dt8/5 was 20 s or longer, the remarkable decrease of absorbed energy was observed. The relation between hardness of weld CGHAZ and cooling time Dt8/5 is given in Fig. 3. Hardness decreased monotonously with increasing cooling time Dt8/5 from 6 to 1000 s. From the results above, it can be understood that when Dt8/5 was 12 s or shorter, the absorbed energies of weld CGHAZs were above 190 J, showing high toughness in weld CGHAZs. When Dt8/5 was 20 s or longer, however, the toughness deteriorated drastically. These results imply that the weld heat-input should be restricted to limit the cooling time D8/5 shorter than 12 s, in order that the weld CGHAZs have high toughness. However, when Dt8/5 was 12 s or shorter, the hardness was very high (higher than 400 HV), suggesting that cold cracking may become a prob-
3.2. Weld ICCGHAZ toughness and microstructure The absorbed energies of weld ICCGHAZs simulated with double thermal cycles are shown in Fig. 5, as a function of peak temperature Tp2 of the second thermal cycle. It could be seen that the
Fig. 3. Relation between hardness of CGHAZs and cooling time Dt8/5.
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Fig. 4. Typical microstructures of CGHAZs with various cooling times Dt8/5.
absorbed energies of the impact test were strongly influenced by both the peak temperature Tp2 and cooling time Dt8/5(2). When Tp2 was in the range from 873 to 923 K (below Ac1), the absorbed energy was much higher than that of weld CGHAZ simulated with single thermal cycle. This high toughness resulted from the tempering effect of the second thermal cycle on the martensitic structure of the CGHAZ. With increasing Tp2 from 923 to 973K (above Ac1), the absorbed energy decreased drastically for both cooling times Dt8/5(2) of 6 and 40 s. For the peak temperatures Tp2 \1003 K, the absorbed energies were closely related to the cooling time Dt8/5(2). For Dt8/5(2) of 6 s, the absorbed energy increased with further increasing Tp2 up to 1073K; but for Dt8/5(2) of 40 s, the absorbed energy decreased further with the increase of Tp2 from 923 to 1073 K. Only a slight increase in absorbed energy was observed for Tp2 higher than Ac3.
The relation between Vickers hardness of weld ICCGHAZs and the peak temperature Tp2 is given in Fig. 6. For Tp2 from 923 to 1003 K, hardness of weld ICCGHAZs was considerably lower than that of the CGHAZ simulated with single thermal cycle due to tempering effect. When Tp2 was above 1023 K, however, the hardness increased again. This phenomenon implied that when the Tp2 was \ 1023 K, most of prior structures were reaustenitized. Upon fast cooling, martensite formed; but upon slow cooling, structures consisting of bainite and martensite formed. The hardness of the martensitic structure or the complex structure of bainite and martensite was probably higher than that of a tempered martensitic structure, which might be the dominant structure for Tp2 B 1023 K. Microstructures of weld ICCGHAZs are revealed in Figs. 7 and 8 for Dt8/5(2) of 6 and 40 s, respectively.
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Fig. 5. Absorbed energy of ICCGHAZs as a function of peak temperature Tp2 of the second thermal cycle (Ac1 for as received material; (Ac1)% for CGHAZ of SQV-2A).
For both cooling times Dt8/5(2) of 6 and 40 s, the peak temperature Tp2 lower than (Ac1)’ was insufficient to cause reaustenitization. Carbides precipitated because of the tempering effect of the second thermal cycle on the martensitic structure of the CGHAZ, which resulted in superior toughness. In the case of Dt8/5(2) of 6 s, for peak temperatures Tp2 in the range of 923 – 1003 K, island-like
Fig. 7. SEM micrographs of ICCGHAZs (Dt8/5(2) =6 s).
Fig. 6. Relation between hardness of ICCGHAZs and Tp2 of the second thermal cycle.
M–A constituents were observed in the interior of the grains, and necklace-like M–A constituents at the grain boundary regions. For Tp2 from 1023 to 1123 K, reaustenitization occurred in large area. Upon cooling, martensite formed in this reaustenitized area. At the same time, untransformed areas were observed as dark areas in the interior of grains. In these untransformed areas, the prior structure was tempered by the second thermal cycle, thus carbides precipitated. When reheated over Ac3, prior structures were completely
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reaustenitized and transformed to martensite again after being cooled, as can be seen in Fig. 7. In the case of Dt8/5(2) of 40 s, M – A constituents also formed in the interior of the grain and at the grain boundary region for Tp2 from 923 to 1023 K. For peak temperatures Tp2 \1073 K, reaustenitization occurred in large area. Upon cooling, microstructures of upper bainite and lower bainite formed in the interior of the grain, and small M – A constituents at the grain boundary region because of the low cooling rate,
Fig. 9. Relation between Tp3 and absorbed energy of tempered ICCGHAZs.
as shown in Fig. 8.
3.3. Effect of tempering thermal cycle
Fig. 8. SEM micrographs of ICCGHAZs (Dt8/5(2) =40 s).
It can be known from Fig. 5 that when Tp2 is 973–1003K, the toughness deterioration occurs in ICCGHAZs irrespective of cooling time Dt8/5(2). When Tp2 is \ 1023 K, the toughness is also deteriorated in the ICCGHAZs for cooling time Dt8/5(2) of 40 s. In this section, investigation was taken to determine the effect of the third thermal cycle (i.e. tempering thermal cycle) on the toughness improvement in these deteriorated ICCGHAZs at peak temperatures Tp2 of 973 and 1073 K. Fig. 9 shows the relationship between peak temperature Tp3 of the tempering thermal cycle and the absorbed energy of the tempered ICCGHAZs. It can be seen that the changing tendency of the absorbed energy with Tp3 depends strongly on peak temperature Tp2 of the second thermal cycle. When Tp2 was at 973 K, the absorbed energies increased with increasing Tp3 of the third thermal cycle to 673 K. Further increasing Tp3 to 893 K resulted in the decrease of absorbed energy. When Tp2 was 1073 K, the absorbed energies increased monotonously with increasing Tp3. In addition, it can also be seen that the absorbed energies are strongly influenced by
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cooling time Dt8/5(2) of the second thermal cycle. When Dt8/5(2) was 6 s, the absorbed energies increased remarkably. When Dt8/5(2) was 40 s, however, the increase in absorbed energy was not as great as that for Dt8/5(2) of 6 s. The relation between peak temperature Tp3 and the hardness of the tempered ICCGHAZs is given in Fig. 10. The hardness of the ICCGHAZ was obviously reduced by tempering thermal cycles only when Tp2 was 1073 K and Dt8/5(2) was 6 s. When the peak temperature Tp3 of the tempering thermal cycles was \573 K, the hardness of all tempered ICCGHAZs was B350 HV. The above results demonstrate that when low heat input welding is adopted (for example, when Dt8/5(2) is 6 s), the deteriorated ICCGHAZs (e.g. at peak temperature Tp2 of 973 K) can be improved to have absorbed energy of above 200 J, provided that the Tp3 is controlled to be above 573 K. Although the absorbed energy decreases when Tp3 is increased to 893 K, the absorbed energy is still about 200 J, at the same level of the weld CGHAZ. However, when comparatively high heat input is adopted (for example, when Dt8/5(2) is 40 s), the toughness improvement by the tempering thermal cycle is not so apparent. Therefore, low heat input is necessary to the toughness improvement in ICCGHAZs by tempering thermal cycles.
Fig. 10. Relation between Tp3 and hardness of tempered ICCGHAZs.
Fig. 11. Effect of Tp3 on the decomposition of M – A constituents.
On the other hand, when low heat input is chosen (for example, when Dt8/5(2) is 6 s), the hardness of ICCGHAZs can be reduced below 350 HV, provided that Tp3 of the tempering thermal cycle is \ 573 K. Moreover, when Tp2 of the second thermal cycle is over Ac3, the high hardness (see Fig. 6) of martensitic structures can also be reduced by tempering thermal cycles to be B 350 HV, provided that the tempering temperature is high enough. That is to say, in the case of no PWHT, weld HAZs of high toughness (absorbed energy\ 200 J) and low hardness (B350 HV) can be achieved, when low heat input welding is adopted and welding procedure is correctly arranged so that the thermal cycles of following weld pass have tempering effects on the former weld pass.
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The microstructures of tempered ICCGHAZs (at Tp2 of 973 K) are given in Fig. 11. When Tp3 was 473–523 K, M – A constituents were decomposed partially. Increasing Tp3 to 573 – 673 K resulted in apparent decomposition of necklace-like M–A constituents. However, when Tp3 was further increased to 773 – 893 K, the decomposition of M–A constituent seems to be difficult. Small M–A constituents could be observed at the edge of original necklace-like M – A constituents. From Figs. 9 and 11, it can be found that the decomposition of M – A constituent corresponds to the increase of absorbed energy in ICCGHAZs.
3.4. Effect of PWHT Fig. 13. Effect of PWHT time on the hardness of CGHAZs.
The absorbed energies and hardness of weld CGHAZs after PWHT at 893 K are shown in Figs. 12 and 13, respectively, as a function of PWHT time. For cooling times Dt8/5 of 6, 40 and 150 s, the changing trends of absorbed energy with PWHT time were similar. The absorbed energies increased with PWHT time, and displayed the maximum value when PWHT time was 7.2 ks. Increasing PWHT time to 14.4 ks led to decrease in absorbed energy. The absorbed energies increased again with further increasing PWHT time to 28.8 ks. In addition, it could be seen that the absorbed energy of the CGHAZ was remarkably increased to be \ 300 J when cooling
time Dt8/5 was 6 s. When cooling time Dt8/5 was 40 s, the absorbed energy was also improved above 200 J after PWHT at 893 K for 3.6–7.2 ks. On the other hand, the hardness decreased markedly with increasing PWHT time to 7.2 ks for Dt8/5 of 6 s. For Dt8/5 of 40 and 150 s, decrease in hardness was not so great. When the PWHT time was 14.4 ks, a slight increase in hardness could be observed, comparing with the case when PWHT time was 7.2 ks. The changing trend of hardness implies that secondary hardening probably occurred when PWHT time was 14.4 ks, thus the toughness decreased slightly (see Fig. 12). From above results, it can be concluded that (i) when cooling times Dt8/5 are in the range of 6 to 40 s, weld CGHAZs toughness can be improved above 200 J, and hardness be reduced below 300 HV by PWHT, (ii) 893 K, 3.6–7.2 ks is the optimum PWHT condition from the view of toughness improvement.
3.5. Factor controlling weld HAZ toughness
Fig. 12. Effect of PWHT time on the absorbed energy of CGHAZs.
The toughness of weld CGHAZs decreased with increasing cooling time Dt8/5 (see Fig. 2). When Dt8/5 was 12 s or shorter, the absorbed energies of weld CGHAZs were above 190 J, showing high toughness in weld CGHAZs. When Dt8/5 was 20 s or longer, the toughness decreased drastically. This remarkable decrease in toughness
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was accompanied by the formation of M–A constituents in CGHAZs (see Fig. 4). The toughness of weld ICCGHAZs depended on not only peak temperature Tp2 but also cooling time Dt8/5(2) of the second thermal cycle. For Dt8/5(2) of 6 s, the toughness deterioration was observed when peak temperature Tp2 was in the range of 973 – 1023 K. For cooling time Dt8/5(2) of 40 s, the toughness deterioration was observed over a wider range of Tp2, i.e. 973 – 1173 K (see Fig. 5). The loss of toughness in ICCGHAZs was also accompanied by the formation of M–A constituents (see Figs. 7 and 8). Therefore, it can be concluded that the toughness deterioration in weld HAZs is mainly attributed to the formation of M – A constituents. The M–A constituent in weld HAZs is the main factor controlling weld HAZ toughness. By decomposing M – A constituents, toughness can be improved (see Figs. 9 and 11). To confirm the existence of M – A constituent in deteriorated weld HAZs, TEM microscopy was carried out on the thin foils prepared from the sample of toughness deteriorated CGHAZs and ICCGHAZs. Fig. 14 reveals TEM micrographs of M–A constituents in the CGHAZs and ICCGHAZs of SQV-2A steel. M – A constituents can be believed to exist in deteriorated weld HAZs of SQV-2A steel. The M – A constituent consists mainly of twin martensite. Carbon concentration in M – A constituents is very high and increases with cooling time Dt8/5 (Liao, 1995). With the increase of carbon content in M–A constituents, the hardness of M –A constituents increases (Okada, 1993). Generally, the hardness of M – A constituent is \ 600 HV, thus M–A constituents can likely become initiative sites of cleavage fracture. SEM observation on the surface of fractured specimens confirmed the cracking initiating role of M – A constituents. M– A constituents were observed at the initiative locations of the cleavage fracture, as can be seen in Fig. 15. The observation of microstructures near the fracture surface of Charpy impacted specimen subjected to the ICCGHAZ simulation also showed that microcracks initiated between the necklace-like M – A constituent and matrix, and
propagate into the interior of the grain, or along the grain boundary, as shown in Fig. 16. Therefore, it can be believed that the formation of M–A constituent is responsible for the toughness loss in weld HAZs of SQV-2A steel.
4. Summary and conclusion The toughness of weld CGHAZs with various heat inputs and that of ICCGHAZs with various peak temperatures of reheating thermal cycle were systematically investigated. The effect of tempering thermal cycles on the weld ICCGHAZs toughness was clarified. The effect of PWHT on weld CGHAZs toughness was determined. On the basis of metallography and fractography, the role of M–A constituents in controlling weld HAZs
Fig. 14. TEM micrographs of M – A constituents in (a) CGHAZs and (b) ICCGHAZs.
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1173 K in the case of Dt8/5(2) = 40 s. Tempering thermal cycles with peak temperature Tp3 of above 573 K can remarkably improve the toughness of deteriorated ICCGHAZs and reduce the hardness, when cooling time Dt8/5(2) is 6 s. These results imply that welding of SQV-2A steel without PWHT is possible, provided that low heat input welding is adopted and welding procedure is correctly arranged. (3) The formation of M–A constituent is responsible for the toughness deterioration in weld HAZs. (4) Based on the above results, the optimum welding and PWHT parameters are recommended as follows:
Fig. 15. SEM micrographs of initiation site of a cleavage fracture ((a) low magnification and (b) high magnification).
toughness was discussed. The results obtained in this study can be summarized as follows: (1) The toughness of weld CGHAZs of SQV2A steel decreases generally with increasing cooling time Dt8/5, namely with increasing weld heat input. When cooling times Dt8/5 are in the range of 6–12 s, the toughness of weld CGHAZs displays high values, but the hardness is too high (\400 HV). Suitable PWHT can significantly improve the toughness and reduce the hardness of weld CGHAZs. When Dt8/5 are in the range of 6–40 s, the absorbed energy can be increased above 200 J, and hardness be reduced below 300 HV after PWHT at 893 K for 3.6 – 7.2 ks. (2) Toughness deterioration occurs in ICCGHAZs, when peak temperature of the second thermal cycle Tp2 =973 – 1023 K in the case of cooling time Dt8/5(2) =6 s, and when Tp2 = 973–
Fig. 16. Microcracks near grain boundary of ICCGHAZs ((a) showing microcracks near grain boundaries and (b) showing cleavage fracture along grain boundary).
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1. When PWHT is carried out, (a) cooling time Dt8/5: 6– 40 s; (b) PWHT: 893 K, 3.6 –7.2 ks. 2. In the case of no PWHT, (a) cooling time Dt8/5: 6– 8 s; (b) peak temperature of the tempering thermal cycle should be controlled in the range 573 – 893 K so that deteriorated ICCGHAZs can be improved by the tempering thermal cycles. Above thermal cycles can result in weld HAZs of high absorbed energy (\ 200 J) and low hardness (B 350 HV).
References Davis, C.L., King, J.E., 1993. Effect of cooling rate on intercritically reheated microstructure and toughness in high strength low alloy steel. Mater. Sci. Technol. 9 (1), 8 – 15. Kim, B.C., Lee, S., Kim, N.J., Lee, D. Y., 1991. Microstructure and Local brittle zone phenomena in high-strength low alloy steel welds. Metall. Trans. A 22A (1), 139 – 149. Liao, J., 1995. Influence of M–A constituent on weld HAZ toughness low alloy steel SQV-2A for pressure vessels. PhD thesis, Tianjin University.
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Linton, V.M., King, J.E., 1992. The effect of heat input on the heat affected zone toughness of two offshore steels. Proc. 3rd Int. Conf. on Trend in Welding Research, Gatlinburg, TN, June 1 – 5 pp. 345 – 349. Nakao, Y., Oshige, H., Noi, S., 1985. Distribution of toughness in HAZ of multi-pass welded high strength steel. Q. J. Jpn. Weld. Soc. 3 – 4, 773 – 781 [in Japanese]. Nippes, E.F., Savage, W.F., Allio, R.J., 1957. Studies of the weld heat-affected zone of T-1 steel. Weld. J., 36 (12), 531s – 540s. Okada, H., 1993. Behavior of M – A constituent in large-heatinput weld HAZs of 780 – 980 MPa class HSLA steels, PhD thesis, Osaka University. Ohno, Y., Habu, R., Sekino, S., 1972. Thermal cycle test of high strength steel and steel plates for low temperature service. Tetsu To Hagane 58 (2), 306 – 316 [in Japanese]. Sato, M., Yamato, K., 1981. On the microstructure and toughness of HAZ in as rolled or normalized 50 and 60 Kg mm − 2 high strength steel. J. Jpn. Weld. Soc. 50 (1), 11–19 [in Japanese]. Savage W.F., Owczarski, W.A., 1966. The microstructure and notch impact behavior of a welded structure steel. Weld. J., 45 (2), 55s – 65s. Tsuboi, J., Hirai, Y., 1981. Microstructure and toughness of weld heat affected zone in quenched-tempered high strength steels. J. Jpn. Weld. Soc. 50 (1), 28 – 37 [in Japanese].