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Microstructure and stress corrosion cracking in simulated heat-affected zones of duplex stainless steels
Corrosion Science 44 (2002) 2841–2856 www.elsevier.com/locate/corsci
Microstructure and stress corrosion cracking in simulated heat-affected zones of duplex stainless steels Horng-Yih Liou
a,b,*
, Rong-Iuan Hsieh b, Wen-Ta Tsai
a
a
b
Department of Materials Science and Engineering, National Cheng Kung University, Tainan 701, Taiwan, ROC Steel and Aluminum Research and Development Department, China Steel Corporation, Kaohsiung 812, Taiwan, ROC Received 12 September 2001; accepted 28 February 2002
Abstract The effects of nitrogen content and the cooling rate on the reformation of austenite in the Gleeble simulated heat-affected zone (HAZ) of 2205 duplex stainless steels (DSSs) were investigated. The variation of stress corrosion cracking (SCC) behavior in the HAZ of 40 wt% CaCl2 solution at 100 °C was also studied. Grain boundary austenite (GBA), Widmanstatten austenite (WA), intergranular austenite (IGA) and partially transformed austenite (PTA) were present in the HAZ. The types and amounts of these reformed austenites varied with the cooling rate and nitrogen content in the DSS. U-bend tests revealed that pitting corrosion and selective dissolution might assist the crack initiation, while the types and amounts of reformed austenite in the HAZ affected the mode of crack propagation. The presence of GBA was found to promote the occurrence of intergranular stress corrosion cracking. WA, IGA and PTA were found to exhibit a beneficial effect on SCC resistance by deviating the crack propagation path. Ó 2002 Elsevier Science Ltd. All rights reserved. Keywords: Duplex stainless steel; Heat-affected zone; Reformed austenite; Stress corrosion cracking
* Corresponding author. Address: Steel and Aluminum Research and Development Department, China Steel Corporation, Kaohsiung 812, Taiwan, ROC. Fax: +886-7-805-1093. E-mail address: [email protected] (H.-Y. Liou).
0010-938X/02/$ - see front matter Ó 2002 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 0 - 9 3 8 X ( 0 2 ) 0 0 0 6 8 - 9
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1. Introduction Although commercial austenitic stainless steels (SSs), such as 304 and 316 SSs, have superior general corrosion resistance and low temperature impact toughness, their pitting, crevice corrosion and stress corrosion cracking (SCC) resistances are poor in chloride-containing environments. This restricts the application of these SSs to environments that are not severely corrosive [1]. The duplex stainless steels (DSSs) containing both ferrite and austenite phases have better localized corrosion resistance than single-phase austenitic SSs in chloride-containing solution. Moreover, the price of DSS is competitive to those conventional austenitic, ferritic SSs and nickelbase alloys. Therefore, the applications of DSSs as structural materials in various industrial sectors, such as chemical, petrochemical, pulp and paper, power generation, desalination, and oil and gas, have been steadily increasing [2–5]. A DSS is an Fe–Cr–Ni alloy that has a two-phase structure, with both phases being above 30% in volume fraction [5]. The good mechanical property and corrosion resistance are dependent upon the proper austenite–ferrite balance, which is approximately 1:1, and high Cr content together with the high Mo and N contents [6]. However, during the fabrication process, such as welding, the austenite–ferrite ratio may be changed together with the precipitation of intermetallic compounds (e.g. r phase, v phase, Cr23 C6 , Cr2 N, c2 , etc.) [7]. For example, the phase (austenite– ferrite) ratio either in the weld-fused zone or in the heat-affected zone (HAZ) tends to deviate from 1:1, due to the weld thermal cycle and different alloy compositions. Different ratios of Creq to Nieq (Creq /Nieq ), peak temperature, heat input, preheat temperature, layer temperature and cooling rate will cause the formations of ferrite and austenite with different grain sizes and volume ratios. These factors also affect the formations of chromium nitride (Cr2 N) and other precipitates [8,9]. In general, the weld region of DSS exhibits higher ferrite content, coarser grain and more extensive Cr2 N precipitation than the base metal. All of these factors have the tendency to reduce both the corrosion resistance and mechanical property (mainly in ductility and toughness) of the weldment [8,9]. Nitrogen is one of the most important elements to improve the corrosion resistance and mechanical properties of the HAZ. It can retard the precipitation of the intermetallic compound by raising the d-ferrite (d-ferrite is also called a-ferrite in many places and hereinafter a-ferrite or a-phase is unified to use) to austenite transformation temperature (A4 temperature) and assisting the reformation of austenite phase in the HAZ [10–12]. Meanwhile, nitrogen has a stronger solid solution strengthening capability and lowers the stacking fault energy, which consequently increases the strength of DSS [13,14]. It can also improve the passivation behavior, facilitate the repassivation process of the base metal within the pit, increase the pitting potential and promote the surface adsorption of the NHþ 4 [15–17]. Therefore, the addition of nitrogen into DSS has become popular because of these multiple beneficial effects. The effects of nitrogen and cooling cycle during welding on the SCC of DSS have only been qualitatively treated in the literature [18,19]. The purpose of this investigation, thus, is to systematically evaluate the effects of the nitrogen content and the
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cooling cycle after welding simulation on the changes of microstructures and their effects on the SCC resistance of DSS in chloride solution.
2. Experimental procedures 2.1. Materials preparation 2205 DSSs, with three different compositions were melted in a 250 kg vacuum furnace and then cast as 160 160 mm2 square ingots. The chemical compositions, listed in Table 1, were varied in nitrogen content based on 22.3Cr–5.7Ni–3.2Mo– 0.9Mn–0.5Si–0.25Cu–0.013C–xN. The ingots were then reheated at 1250 °C for 1.5 h and hot rolled, using a laboratory hot rolling mill, into 13 mm thick plates. Prior to weld simulation, these steels were subjected to a solution treatment at 1100 °C for 10 min and then quenched in water. 2.2. Weld simulation The specimens, 10:5 10:5 80 mm3 , used for the welding HAZ simulation with a Gleeble 1500 thermo-mechanical simulator were cut from the 13 mm thick plates with their longitudinal direction parallel to the rolling direction. The thermal history of the simulation was based on the temperature/time function derived by Lindblom et al. [20]. The ambient temperature was controlled at 25 °C. The peak temperature was chosen as 1350 °C and was held for 1 s before cooling. The welding guide line is given by the cooling time between 1200 and 800 °C, Dt12=8 , which is the typical temperature range within which austenite reformation and precipitation occur for DSSs [20]. However, a relationship exists between Dt12=8 and Dt8=5 (cooling time between 800 and 500 °C), giving the following expression [20]. 1 2
1
2 Dt8=5 ð500 T0 Þ ð800 T0 Þ ¼ ; 1 1 Dt12=8 ð800 T0 Þ2 ð1200 T0 Þ2
where T0 ¼ starting temperature (°C). Therefore, it is preferable to apply Dt8=5 in this discussion since it is more easily and accurately measured [20]. In this investigation, the cooling times from 800 to 500 °C (Dt8=5 ) were controlled at 5, 20, 60 and 100 s, respectively. Fig. 1 shows an Table 1 Chemical compositions of the experimental SSs used (wt%) C
Si
Mn
P
S
Ni
Cr
Mo
N
Cu
Fe
0.013 0.013 0.013
0.48 0.50 0.48
0.89 0.90 0.89
0.020 0.021 0.020
0.0038 0.0042 0.0038
5.60 5.69 5.60
22.3 22.3 22.3
3.22 3.21 3.22
0.096 0.135 0.165
0.25 0.25 0.25
Bal. Bal. Bal.
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Fig. 1. Thermal cycles of Gleeble simulation.
example of the thermal cycle for the Gleeble simulation, with the relationship between Dt8=5 and the heat input derived from the following equation [20]: Q=d ¼ k ðDt8=5 Þ1=2 ; where Q is the net heat input (J/mm), d is the thickness (mm), k is the thermal coefficient ðJ=mm2 s1=2 Þ ; 25:52 (for DSS), and Dt8=5 is the cooling time from 800 to 500 °C (s). All Gleeble simulation parameters and calculated heat input are listed in Table 2. 2.3. Metallography The microstructures of the simulated HAZ were examined by using optical microscopy (OM). The etching solution used for the OM examination was the boiling Murukami reagent (30 g K3 Fe(CN)6 þ 30 g KOH þ 100 ml H2 O), which made the austenite white, ferrite tan, and carbide black. Quantitative measurement of the austenite phase content was carried out by an image analyzer plus point counting method, with the examined area being about 4 mm2 . In addition, the amount of Cr2 N in the simulated HAZ was determined by analyzing the chromium Table 2 Parameters and heat inputs for Gleeble simulation Dt12=8 (s)
Dt8=5 (s)
Tmax (°C)
Q (kJ/mm)
2 8 21 35
5 20 60 100
1350 1350 1350 1350
0.57 1.14 1.98 2.55
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content of precipitates. Electrolytic extraction and analysis of precipitates from the simulated HAZ was performed by potentiostatic electrolysis in a nonaqueous electrolyte solution [21]. Detailed experimental procedures can be seen in a previous work [22]. 2.4. Stress corrosion cracking test The U-bend specimen was used to evaluate the sensitivity to SCC of the simulated HAZ in DSSs with different nitrogen contents. The test solution was 40 wt% calcium chloride (CaCl2 ) at 100 °C. The time-to-failure was recorded as a measure for the resistance to SCC. Post-fracture observation was carried out by either an OM or a scanning electron microscope (SEM).
3. Results and discussion 3.1. Effects of nitrogen content and cooling rate on the microstructure Fig. 2 shows the effect of the cooling time (Dt8=5 ) on the microstructural change in the simulated HAZ of a 0.096 wt% N DSS. When the peak temperature was held at 1350 °C for 1 s, the c phase was completely dissolved and only the a phase was stable. With a Dt8=5 of 5 s, the HAZ primarily consisted of the ferrite phase with a small amount of austenite formed at the grain boundary, Fig. 2a. The a grain coarsened to about 400 lm when the cooling time was held for 5 s from 800 to 500 °C. As the Dt8=5 was increased to 20 s, the formation of grain boundary austenite (GBA) became noticeable (Fig. 2b). In addition to GBA, Widmanstatten austenite (WA) and intragranular austenite (IGA) were observed when the Dt8=5 was increased to 60 s (Fig. 2c). At a much slower cooling rate (Dt8=5 ¼ 100 s), the amounts of WA and IGA increased and the GBA was almost fully replaced by WA (Fig. 2d). Since the precipitation of austenite is a diffusion-controlled nucleation and growth process [23,24], the reformation of austenite is controlled by a para-equilibrium transformation mechanism in which the diffusion of the interstitial elements (carbon and nitrogen) is the controlling process [11]. Leone and Kerr [25] pointed out that the diffusion distance of nitrogen during one weld thermal cycle (Dt8=5 ¼ 15 s) was 50–100 lm. This is much longer than those of chromium and nickel, being 0.9 and 4.4 lm respectively. Prolonging the cooling time will, thus, produce the diffusion of c-stabilization elements, such as N and Ni, and ensure more austenite been transformed from ferrite [26]. For a high nitrogen containing (0.165 wt% N) DSS, the temperature for full dissolution of austenite was demonstrated to be about 1360 °C by thermodynamic calculation in a previous work [27]. Consequently, only part of the original c phase in the solution annealed DSS was dissolved by holding the temperature at 1350 °C for 1 s. The partially transformed austenite (PTA) was found to inhibit the grain growth of the a phase [26], which led to a finer ferritic grain size than that of the DSS with 0.096 wt% N. The effect of cooling rate (or Dt8=5 ) on the microstructure change
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Fig. 2. The microstructure of the HAZ of 0.096 wt% N DSSs after different cooling time treatments, Dt8=5 ¼ 5 s (a), 20 s (b), 60 s (c) and 100 s (d).
in HAZ for the high nitrogen containing 0.165 wt% N DSS is demonstrated in Fig. 3. Comparing Fig. 2 with Fig. 3, the austenite content increases markedly and the ferrite grain size becomes finer as the nitrogen content is increased from 0.096 to 0.165 wt%. GBA, WA, and IGA were found to precipitate when the temperature was lowered to 800–500 °C. The amount of reformed austenite (GBA, WA and IGA) increased with increasing Dt8=5 as revealed in Fig. 3. The effect of cooling time on the change in austenite content for different levels of nitrogen containing DSSs is shown in Fig. 4. The austenite content in the solution annealed (1100 °C/10 min) DSS increased from 24% to 42% as the nitrogen content was raised from 0.096 to 0.165 wt%. In the simulated HAZ, the austenite varied not only with the nitrogen content but also with the cooling rate. For instance, the austenite contents in the simulated HAZ increased from 2% to 13% for the DSS with 0.096% N, and from 16% to 35% for that with 0.165% N, as Dt8=5 was increased from 5 to 100 s. With a slow cooling rate (Dt8=5 ¼ 100 s), the c content in the simulated HAZ was close to that of the solution annealed DSS, especially in the higher nitrogen bearing DSS. It has been found that the ferrite solvus temperature can be raised by increasing the nitrogen content in the DSS [11,12]. Consequently, the reformation of austenite in DSS after welding becomes more feasible.
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Fig. 3. The microstructure of the HAZ of 0.165 wt% N DSSs after different cooling time treatments, Dt8=5 ¼ 5 s (a), 20 s (b), 60 s (c) and 100 s (d).
A previous investigation revealed that Cr2 N is the main precipitate formed in the simulated HAZ of DSS [22]. The amount of Cr2 N formed was found to increase with decreasing nitrogen content in the DSS. The results also showed that the precipitation of Cr2 N was inhibited as the cooling rate in the simulated HAZ was lowered. This is because the solubility of nitrogen in the ferrite phase is rather low, around 0.05% [28], and supersaturation of nitrogen will occur in the ferrite if the austenite forms too slowly during cooling. Since the reformation of austenite is prevalent at low cooling rates (or high Dt8=5 ), the precipitation of Cr2 N will become less favourable. The higher the austenite, the lower the nitrogen in the form of Cr2 N [22]. 3.2. Effects of nitrogen content and cooling rate on the SCC The effect of N content and cooling time on the SCC sensitivity in 40 wt% CaCl2 solution of solution annealed and simulated HAZs of DSSs is shown in Fig. 5. According to Bernhardsson [29], 40 wt% calcium chloride (CaCl2 ) solution at 100 °C was chosen to be the SCC test solution due to its good reproduction and correspondence to predict the actual application. In this investigation, no matter what the
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Fig. 4. Effects of cooling time on the austenite content in the solution annealed matrix and simulated HAZs of DSSs with various nitrogen contents of DSS.
Fig. 5. Effect of cooling time on the time-to-failure in 40 wt% CaCl2 solution (100 °C) of solution annealed matrix and simulated HAZs with various nitrogen contents of DSSs.
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variation in nitrogen content was, the solution annealed DSSs had similar SCC resistance. The average time-to-failure was more than 450 h. This result demonstrated that in CaCl2 solution the SCC susceptibility of DSSs was almost independent of nitrogen content in the range of 0.096–0.165%. But, the sensitivity of SCC in simulated HAZs compared with the solution-annealed matrix of DSSs was dramatically increased. The time-to-failure of HAZ at high cooling rate (Dt8=5 ¼ 5 s) of DSS with 0.096% N was only 18 h. However, the resistance to SCC gradually improved along with the increase in either the cooling time or the nitrogen content. The higher the nitrogen content or the cooling time, the better the resistance to SCC in HAZs of DSSs. For DSS with 0.165% N and at Dt8=5 of 100 s, the time to failure in HAZs increased to about 365 h. At such a low cooling rate, the SCC resistance of the HAZ was raised almost to the same level as the solution annealed DSS. The susceptibility of DSS to SCC depends on many factors such as: alloying elements [30,31], microstructures [2,31–33], applied stresses [34], environments [29,35,36], etc. From the U-bend tests in the 40% CaCl2 solution at 100 °C, the timeto-failure of the HAZ as a function of austenite content is demonstrated in Fig. 6. Clearly, the time-to-failure increased with increasing austenite content which varied with the cooling rate (Dt8=5 ) as revealed in Fig. 4. As mentioned in the previous section, the quantities of Cr2 N decreased with increasing austenite content in the HAZ, and it seemed that the SCC susceptibility of the HAZ increased with increasing amount of Cr2 N precipitation. The change in microstructure in the HAZ
Fig. 6. Effect of austenite content on the time-to-failure in 40 wt% CaCl2 solution (100 °C) of solution annealed matrix and simulated HAZs with various nitrogen contents of DSSs.
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Fig. 7. SEM micrograph showing SCC crack resulting from cross-link of pit, U-bend test for an HAZ of DSS in 40 wt% CaCl2 solution.
has a significant effect on the SCC behavior as will be discussed later, in addition to the effects resulting from austenite content and Cr2 N precipitate. The fracture morphology of SCC in DSSs was further examined. Fig. 7 shows an example of the SEM micrograph of the surface crack of an HAZ after the U-bend test in 40 wt% CaCl2 solution (100 °C). As shown in this figure, pits were formed on the surface. Micro-fissures were found to be induced in the vicinity of pits and gradually cross-linked one from the other. Clearly, pitting-assisted stress corrosion cracking manipulated the fracture at the HAZ of DSS. As mentioned in a previous work [22], Cr2 N precipitates acted as preferential nucleation sites for pitting corrosion. The pitting-assisted SCC as found in Fig. 7 could also be attributed to the presence of Cr2 N in the HAZ of DSS. The occurrence of pitting corrosion was particularly important in the crack initiation during the process of SCC. Fig. 8 shows the near surface cross-section micrographs of the failed U-bend specimens for the solution annealed and the HAZ of a DSS with 0.165 wt% N after testing in the 40 wt% CaCl2 solution at 100 °C. The results revealed that the ferrite phase either in the solution annealed matrix or in the HAZ corroded preferentially in this testing environment. Selective dissolution of ferrite phase could result in surface crack initiation and subsequently induced SCC of DSS, either solution annealed or welded. Selective dissolution depends on the distribution of alloying elements between the a and c phases, the amount of precipitates formed, and the testing solutions used [37,38]. Sridhar and Kolts [37] proposed that corrosion occurred mostly in the austenite for the high nitrogen containing DSS (Ferralium alloy 255, 0.17% N) in both phosphoric acid and sulfuric acid solutions. Whereas, in hydrochloric acid and other chloride environments, corrosion occurred preferentially in the ferrite phase [37]. However, Fourie and Robinson [38] observed that the selective dissolution behavior of 2205 DSS was a function of the chloride concentration in 1 M H2 SO4 solution at 60 °C. At high chloride concentrations (more than 0.2 M NaCl) selective
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Fig. 8. The OM micrograph showing the selective dissolution of ferrite phase in the (a) solution annealed matrix and (b) HAZ (Dt8=5 ¼ 60 s) of 0.165 wt% N DSS after SCC testing in 40 wt% CaCl2 solution.
dissolution of ferrite occurred while selective dissolution of austenite was observed in the absence of chloride. At low chloride concentrations (less than 0.1 M NaCl) both austenite and ferrite phases were dissolved. In this study, preferential corrosion of ferrite phase was probably due to the high concentrations of chloride (40 wt% CaCl2 solution), despite the absence of H2 SO4 . Fig. 9 depicts the cross-section crack propagation path of a secondary crack in a simulated HAZ (Dt8=5 ¼ 5 s) of DSS with 0.096 wt% N in 40 wt% CaCl2 solution at 100 °C. It is interesting to note that the crack propagated in a transgranular manner in the early stage, which was followed by intergranular stress corrosion cracking (IGSCC) in the latter stage. In other words, crack initiation was not confined to the grain boundary even though the crack might propagate intergranularly. The SEM fractograph showing the combined feature of transgranular crack initiation and intergranular crack propagation in the HAZ of of 0.096 wt% N DSS is depicted in Fig. 10. With Dt8=5 ¼ 5 s, GBA was formed in the interfaces between ferrite grains in the HAZ of 0.096 wt% N DSS. The intergranular mode of SCC was thought to be related to the presence of GBA. EDS examination of the fracture surface could be employed to identify the presence of austenite in the grain boundary or at the
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Fig. 9. The cross-sectioned crack propagation path in the HAZ (Dt8=5 ¼ 5 s) of 0.096 wt% N DSS after SCC testing in the 40 wt% CaCl2 solution.
interface of ferrite grains. As depicted in Fig. 10b, the clean and smooth ferrite grain facets were clearly revealed. On top of some of the ferrite grain facets, a thin layer of GBA with an appearance like orange peel was found to exist. This thin layer of GBA could be very brittle and cause the occurrence of IGSCC. For the DSS with 0.096 wt% N, a significant change in the fracture surface morphology of the simulated HAZ was observed upon varying Dt8=5 . At Dt8=5 of 20 s, transgranular stress corrosion cracking (TGSCC) mixed with IGSCC was observed (Fig. 11a). As the cooling time was increased to 100 s, exclusive TGSCC was observed (Fig. 11b). For DSSs with nitrogen contents of 0.135 and 0.165 wt%, the fracture surface morphologies of the simulated HAZs were also affected by the cooling rate. At Dt8=5 of 5 s, mixed TGSCC/IGSCC mode was observed but the percentage of IGSCC in the HAZ of DSS with 0.165 wt% N was less than that of DSS with 0.135 wt% N. As the cooling time increased beyond 20 s, only TGSCC was observed in the HAZs of both the DSSs. Fig. 11c gives an example showing TGSCC in the HAZ, with Dt8=5 ¼ 60 s, of a DSS containing 0.135 wt% of N. As illustrated in Figs. 2 and 3, the microstructure of the simulated HAZ depends on the nitrogen content and the cooling rate. GBA was present in the HAZ with low nitrogen content and cooled at a high cooling rate. The occurrence of IGSCC in that of 0.096 wt% N DSS was believed to be associated with the presence of GBA. The cause of the deleterious effect of GBA is not known and needs further investigation. As the reformed austenite increased in grain size and volume fraction, the resistance to IGSCC also improved. The presence of WA, IGA and PTA seemed not to promote IGSCC in the HAZ of DSS in 40% CaCl2 solution at 100 °C. These types of reformed austenite, on the contrary, exhibited a beneficial effect in delaying crack propagation by deviating its growth path as demonstrated in Fig. 12 for the HAZ, with Dt8=5 ¼ 5 s, of a DSS of 0.165 wt% N. As shown in this micrograph, the crack propagated transgranularly in the ferrite phase and changed path when it encountered WA or IGA.
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Fig. 10. (a) The SEM micrograph of the HAZ (Dt8=5 ¼ 5 s) of 0.096 wt% N DSS after SCC testing in 40 wt% CaCl2 solution, (b) enlargement of (a).
4. Conclusions 1. Grain boundary austenite, WA, IGA and partially transformed austenite could be formed in the HAZ of DSSs. The types and amounts of these reformed austenites in the HAZ depended on the nitrogen content in the DSS and the cooling rate employed. 2. The SCC resistance in the HAZ was inferior to that of the solution annealed DSS. However, the susceptibility to SCC of the HAZ in 40 wt% CaCl2 solution at 100 °C decreased as the amount of reformed austenite increased. The SCC resistance of the HAZ was also found to increase as the nitrogen content in the DSS increased and the cooling rate decreased.
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Fig. 11. The SEM micrograph of the HAZ of DSS after SCC testing in 40 wt% CaCl2 solution, (a) N ¼ 0:096 wt%, Dt8=5 ¼ 20 s, (b) N ¼ 0:096 wt%, Dt8=5 ¼ 100 s, (c) N ¼ 0:135 wt%, Dt8=5 ¼ 60 s.
3. Pitting corrosion and selective dissolution of ferrite phases could assist crack initiation and induce SCC in the HAZ of DSSs. 4. IGSCC could be found in the HAZ when GBA was present. The transformation from IGSCC to TGSCC occurred as GBA was gradually replaced by WA or IGA.
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Fig. 12. The cross-sectioned crack propagation path in the HAZ (Dt8=5 ¼ 5 s) of 0.165 wt% N DSS after SCC testing in the 40 wt% CaCl2 solution.
5. The presence of WA, IGA and PTA exhibited beneficial effect on SCC resistance by deviating the crack propagation path.
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S.L. Chou, M.J. Tsai, W.T. Tsai, J.T. Lee, Materials Chemistry and Physics 51 (1997) 97. Y.S. Lim, J.S. Kim, S.J. Ahn, H.S. Kwon, Y. Katada, Corrosion Science 43 (2001) 53. C.A. Olsson, Corrosion Science 37 (3) (1995) 467. T. Kuroda, F. Matsuda, in: Proceedings of Duplex Stainless Steels Confference, Glasgow, Scotland, November 13–16, 1994, Paper 35. T.G. Gooch, in: Proceedings of Duplex Stainless Steels Conference, Beaune, Bourgogne, France, October 28–30, 1991, vol. 1, p. 325. B.E.S. Lindblom, B. Lundquist, N.E. Hannerz, Scandinavian Journal of Metallurgy 20 (1991) 305. M. Saeki, F. Kurosawa, M. Matsuo, Transaction ISIJ 26 (1986) 1017. H.Y. Liou, R.I. Hsieh, W.T. Tsai, Materials Chemistry and Physics 74 (2002) 33. R. Mundt, H. Hoffmeister, in: G.L. Dunlop (Ed.), Stainless Steels‘84, The Institute of Metals, 1984, p. 315. S. Hertzman, W. Robert, M. Lindenmo, in: Stainless Steels’86, The Hague, 1986, Paper 30, p. 257. G. Leone, H.W. Kerr, Welding Journal 61 (1982) 13s. S. Atamert, J.E. King, Materials Science and Technology 8 (1992) 896. R.I. Hsieh, H.Y. Liou, Y.T. Pan, China Steel Technical Report 13 (1999) 14. J. Charles, in: J. Charles, S. Bernhardsson (Eds.), Duplex Stainless Steels’91, vol. 1, Les Editions de Physique, France, 1991, p. 3. S. Bernhardsson, Sandvik Steel Report, S-51-32-ENG, 1984, p. 1. K.Y. Kim, P.Q. Zhang, T.H. Ha, Y.H. Lee, Corrosion 54 (11) (1998) 910. A. Laitinen, H. Hanninen, Corrosion 52 (4) (1996) 295. S. Mark Wilhelm, R.D. Kane, Corrosion 40 (8) (1984) 431. T. Magnin, J.M. Lardon, Materials Science and Engineering A 104 (1988) 21. A. Miyasaka, T. Kanamaru, H. Ogawa, Corrosion 52 (8) (1996) 592. K. Van Gelder, J.G. Erlings, J.W.M. Damen, A. Visser, Corrosion Science 27 (10/11) (1987) 1271. R.F.A. Jargelius, R. Blom, S. Hertzman, J. Linder, in: J. Charles, S. Bernhardsson (Eds.), Duplex Stainless Steel’91, Les Editions de Physique, 1992, p. 211. N. Sridhar, J. Kolts, Corrosion 43 (11) (1987) 646. J.W. Fourie, F.P.A. Robinson, in: Proceedings of International Conference on Stainless Steels, ISIJ, Chiba, 1991, p. 111.
Microstructure and stress corrosion cracking in simulated heat-affected zones of duplex stainless steels Horng-Yih Liou
a,b,*
, Rong-Iuan Hsieh b, Wen-Ta Tsai
a
a
b
Department of Materials Science and Engineering, National Cheng Kung University, Tainan 701, Taiwan, ROC Steel and Aluminum Research and Development Department, China Steel Corporation, Kaohsiung 812, Taiwan, ROC Received 12 September 2001; accepted 28 February 2002
Abstract The effects of nitrogen content and the cooling rate on the reformation of austenite in the Gleeble simulated heat-affected zone (HAZ) of 2205 duplex stainless steels (DSSs) were investigated. The variation of stress corrosion cracking (SCC) behavior in the HAZ of 40 wt% CaCl2 solution at 100 °C was also studied. Grain boundary austenite (GBA), Widmanstatten austenite (WA), intergranular austenite (IGA) and partially transformed austenite (PTA) were present in the HAZ. The types and amounts of these reformed austenites varied with the cooling rate and nitrogen content in the DSS. U-bend tests revealed that pitting corrosion and selective dissolution might assist the crack initiation, while the types and amounts of reformed austenite in the HAZ affected the mode of crack propagation. The presence of GBA was found to promote the occurrence of intergranular stress corrosion cracking. WA, IGA and PTA were found to exhibit a beneficial effect on SCC resistance by deviating the crack propagation path. Ó 2002 Elsevier Science Ltd. All rights reserved. Keywords: Duplex stainless steel; Heat-affected zone; Reformed austenite; Stress corrosion cracking
* Corresponding author. Address: Steel and Aluminum Research and Development Department, China Steel Corporation, Kaohsiung 812, Taiwan, ROC. Fax: +886-7-805-1093. E-mail address: [email protected] (H.-Y. Liou).
0010-938X/02/$ - see front matter Ó 2002 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 0 - 9 3 8 X ( 0 2 ) 0 0 0 6 8 - 9
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1. Introduction Although commercial austenitic stainless steels (SSs), such as 304 and 316 SSs, have superior general corrosion resistance and low temperature impact toughness, their pitting, crevice corrosion and stress corrosion cracking (SCC) resistances are poor in chloride-containing environments. This restricts the application of these SSs to environments that are not severely corrosive [1]. The duplex stainless steels (DSSs) containing both ferrite and austenite phases have better localized corrosion resistance than single-phase austenitic SSs in chloride-containing solution. Moreover, the price of DSS is competitive to those conventional austenitic, ferritic SSs and nickelbase alloys. Therefore, the applications of DSSs as structural materials in various industrial sectors, such as chemical, petrochemical, pulp and paper, power generation, desalination, and oil and gas, have been steadily increasing [2–5]. A DSS is an Fe–Cr–Ni alloy that has a two-phase structure, with both phases being above 30% in volume fraction [5]. The good mechanical property and corrosion resistance are dependent upon the proper austenite–ferrite balance, which is approximately 1:1, and high Cr content together with the high Mo and N contents [6]. However, during the fabrication process, such as welding, the austenite–ferrite ratio may be changed together with the precipitation of intermetallic compounds (e.g. r phase, v phase, Cr23 C6 , Cr2 N, c2 , etc.) [7]. For example, the phase (austenite– ferrite) ratio either in the weld-fused zone or in the heat-affected zone (HAZ) tends to deviate from 1:1, due to the weld thermal cycle and different alloy compositions. Different ratios of Creq to Nieq (Creq /Nieq ), peak temperature, heat input, preheat temperature, layer temperature and cooling rate will cause the formations of ferrite and austenite with different grain sizes and volume ratios. These factors also affect the formations of chromium nitride (Cr2 N) and other precipitates [8,9]. In general, the weld region of DSS exhibits higher ferrite content, coarser grain and more extensive Cr2 N precipitation than the base metal. All of these factors have the tendency to reduce both the corrosion resistance and mechanical property (mainly in ductility and toughness) of the weldment [8,9]. Nitrogen is one of the most important elements to improve the corrosion resistance and mechanical properties of the HAZ. It can retard the precipitation of the intermetallic compound by raising the d-ferrite (d-ferrite is also called a-ferrite in many places and hereinafter a-ferrite or a-phase is unified to use) to austenite transformation temperature (A4 temperature) and assisting the reformation of austenite phase in the HAZ [10–12]. Meanwhile, nitrogen has a stronger solid solution strengthening capability and lowers the stacking fault energy, which consequently increases the strength of DSS [13,14]. It can also improve the passivation behavior, facilitate the repassivation process of the base metal within the pit, increase the pitting potential and promote the surface adsorption of the NHþ 4 [15–17]. Therefore, the addition of nitrogen into DSS has become popular because of these multiple beneficial effects. The effects of nitrogen and cooling cycle during welding on the SCC of DSS have only been qualitatively treated in the literature [18,19]. The purpose of this investigation, thus, is to systematically evaluate the effects of the nitrogen content and the
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cooling cycle after welding simulation on the changes of microstructures and their effects on the SCC resistance of DSS in chloride solution.
2. Experimental procedures 2.1. Materials preparation 2205 DSSs, with three different compositions were melted in a 250 kg vacuum furnace and then cast as 160 160 mm2 square ingots. The chemical compositions, listed in Table 1, were varied in nitrogen content based on 22.3Cr–5.7Ni–3.2Mo– 0.9Mn–0.5Si–0.25Cu–0.013C–xN. The ingots were then reheated at 1250 °C for 1.5 h and hot rolled, using a laboratory hot rolling mill, into 13 mm thick plates. Prior to weld simulation, these steels were subjected to a solution treatment at 1100 °C for 10 min and then quenched in water. 2.2. Weld simulation The specimens, 10:5 10:5 80 mm3 , used for the welding HAZ simulation with a Gleeble 1500 thermo-mechanical simulator were cut from the 13 mm thick plates with their longitudinal direction parallel to the rolling direction. The thermal history of the simulation was based on the temperature/time function derived by Lindblom et al. [20]. The ambient temperature was controlled at 25 °C. The peak temperature was chosen as 1350 °C and was held for 1 s before cooling. The welding guide line is given by the cooling time between 1200 and 800 °C, Dt12=8 , which is the typical temperature range within which austenite reformation and precipitation occur for DSSs [20]. However, a relationship exists between Dt12=8 and Dt8=5 (cooling time between 800 and 500 °C), giving the following expression [20]. 1 2
1
2 Dt8=5 ð500 T0 Þ ð800 T0 Þ ¼ ; 1 1 Dt12=8 ð800 T0 Þ2 ð1200 T0 Þ2
where T0 ¼ starting temperature (°C). Therefore, it is preferable to apply Dt8=5 in this discussion since it is more easily and accurately measured [20]. In this investigation, the cooling times from 800 to 500 °C (Dt8=5 ) were controlled at 5, 20, 60 and 100 s, respectively. Fig. 1 shows an Table 1 Chemical compositions of the experimental SSs used (wt%) C
Si
Mn
P
S
Ni
Cr
Mo
N
Cu
Fe
0.013 0.013 0.013
0.48 0.50 0.48
0.89 0.90 0.89
0.020 0.021 0.020
0.0038 0.0042 0.0038
5.60 5.69 5.60
22.3 22.3 22.3
3.22 3.21 3.22
0.096 0.135 0.165
0.25 0.25 0.25
Bal. Bal. Bal.
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Fig. 1. Thermal cycles of Gleeble simulation.
example of the thermal cycle for the Gleeble simulation, with the relationship between Dt8=5 and the heat input derived from the following equation [20]: Q=d ¼ k ðDt8=5 Þ1=2 ; where Q is the net heat input (J/mm), d is the thickness (mm), k is the thermal coefficient ðJ=mm2 s1=2 Þ ; 25:52 (for DSS), and Dt8=5 is the cooling time from 800 to 500 °C (s). All Gleeble simulation parameters and calculated heat input are listed in Table 2. 2.3. Metallography The microstructures of the simulated HAZ were examined by using optical microscopy (OM). The etching solution used for the OM examination was the boiling Murukami reagent (30 g K3 Fe(CN)6 þ 30 g KOH þ 100 ml H2 O), which made the austenite white, ferrite tan, and carbide black. Quantitative measurement of the austenite phase content was carried out by an image analyzer plus point counting method, with the examined area being about 4 mm2 . In addition, the amount of Cr2 N in the simulated HAZ was determined by analyzing the chromium Table 2 Parameters and heat inputs for Gleeble simulation Dt12=8 (s)
Dt8=5 (s)
Tmax (°C)
Q (kJ/mm)
2 8 21 35
5 20 60 100
1350 1350 1350 1350
0.57 1.14 1.98 2.55
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content of precipitates. Electrolytic extraction and analysis of precipitates from the simulated HAZ was performed by potentiostatic electrolysis in a nonaqueous electrolyte solution [21]. Detailed experimental procedures can be seen in a previous work [22]. 2.4. Stress corrosion cracking test The U-bend specimen was used to evaluate the sensitivity to SCC of the simulated HAZ in DSSs with different nitrogen contents. The test solution was 40 wt% calcium chloride (CaCl2 ) at 100 °C. The time-to-failure was recorded as a measure for the resistance to SCC. Post-fracture observation was carried out by either an OM or a scanning electron microscope (SEM).
3. Results and discussion 3.1. Effects of nitrogen content and cooling rate on the microstructure Fig. 2 shows the effect of the cooling time (Dt8=5 ) on the microstructural change in the simulated HAZ of a 0.096 wt% N DSS. When the peak temperature was held at 1350 °C for 1 s, the c phase was completely dissolved and only the a phase was stable. With a Dt8=5 of 5 s, the HAZ primarily consisted of the ferrite phase with a small amount of austenite formed at the grain boundary, Fig. 2a. The a grain coarsened to about 400 lm when the cooling time was held for 5 s from 800 to 500 °C. As the Dt8=5 was increased to 20 s, the formation of grain boundary austenite (GBA) became noticeable (Fig. 2b). In addition to GBA, Widmanstatten austenite (WA) and intragranular austenite (IGA) were observed when the Dt8=5 was increased to 60 s (Fig. 2c). At a much slower cooling rate (Dt8=5 ¼ 100 s), the amounts of WA and IGA increased and the GBA was almost fully replaced by WA (Fig. 2d). Since the precipitation of austenite is a diffusion-controlled nucleation and growth process [23,24], the reformation of austenite is controlled by a para-equilibrium transformation mechanism in which the diffusion of the interstitial elements (carbon and nitrogen) is the controlling process [11]. Leone and Kerr [25] pointed out that the diffusion distance of nitrogen during one weld thermal cycle (Dt8=5 ¼ 15 s) was 50–100 lm. This is much longer than those of chromium and nickel, being 0.9 and 4.4 lm respectively. Prolonging the cooling time will, thus, produce the diffusion of c-stabilization elements, such as N and Ni, and ensure more austenite been transformed from ferrite [26]. For a high nitrogen containing (0.165 wt% N) DSS, the temperature for full dissolution of austenite was demonstrated to be about 1360 °C by thermodynamic calculation in a previous work [27]. Consequently, only part of the original c phase in the solution annealed DSS was dissolved by holding the temperature at 1350 °C for 1 s. The partially transformed austenite (PTA) was found to inhibit the grain growth of the a phase [26], which led to a finer ferritic grain size than that of the DSS with 0.096 wt% N. The effect of cooling rate (or Dt8=5 ) on the microstructure change
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Fig. 2. The microstructure of the HAZ of 0.096 wt% N DSSs after different cooling time treatments, Dt8=5 ¼ 5 s (a), 20 s (b), 60 s (c) and 100 s (d).
in HAZ for the high nitrogen containing 0.165 wt% N DSS is demonstrated in Fig. 3. Comparing Fig. 2 with Fig. 3, the austenite content increases markedly and the ferrite grain size becomes finer as the nitrogen content is increased from 0.096 to 0.165 wt%. GBA, WA, and IGA were found to precipitate when the temperature was lowered to 800–500 °C. The amount of reformed austenite (GBA, WA and IGA) increased with increasing Dt8=5 as revealed in Fig. 3. The effect of cooling time on the change in austenite content for different levels of nitrogen containing DSSs is shown in Fig. 4. The austenite content in the solution annealed (1100 °C/10 min) DSS increased from 24% to 42% as the nitrogen content was raised from 0.096 to 0.165 wt%. In the simulated HAZ, the austenite varied not only with the nitrogen content but also with the cooling rate. For instance, the austenite contents in the simulated HAZ increased from 2% to 13% for the DSS with 0.096% N, and from 16% to 35% for that with 0.165% N, as Dt8=5 was increased from 5 to 100 s. With a slow cooling rate (Dt8=5 ¼ 100 s), the c content in the simulated HAZ was close to that of the solution annealed DSS, especially in the higher nitrogen bearing DSS. It has been found that the ferrite solvus temperature can be raised by increasing the nitrogen content in the DSS [11,12]. Consequently, the reformation of austenite in DSS after welding becomes more feasible.
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Fig. 3. The microstructure of the HAZ of 0.165 wt% N DSSs after different cooling time treatments, Dt8=5 ¼ 5 s (a), 20 s (b), 60 s (c) and 100 s (d).
A previous investigation revealed that Cr2 N is the main precipitate formed in the simulated HAZ of DSS [22]. The amount of Cr2 N formed was found to increase with decreasing nitrogen content in the DSS. The results also showed that the precipitation of Cr2 N was inhibited as the cooling rate in the simulated HAZ was lowered. This is because the solubility of nitrogen in the ferrite phase is rather low, around 0.05% [28], and supersaturation of nitrogen will occur in the ferrite if the austenite forms too slowly during cooling. Since the reformation of austenite is prevalent at low cooling rates (or high Dt8=5 ), the precipitation of Cr2 N will become less favourable. The higher the austenite, the lower the nitrogen in the form of Cr2 N [22]. 3.2. Effects of nitrogen content and cooling rate on the SCC The effect of N content and cooling time on the SCC sensitivity in 40 wt% CaCl2 solution of solution annealed and simulated HAZs of DSSs is shown in Fig. 5. According to Bernhardsson [29], 40 wt% calcium chloride (CaCl2 ) solution at 100 °C was chosen to be the SCC test solution due to its good reproduction and correspondence to predict the actual application. In this investigation, no matter what the
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Fig. 4. Effects of cooling time on the austenite content in the solution annealed matrix and simulated HAZs of DSSs with various nitrogen contents of DSS.
Fig. 5. Effect of cooling time on the time-to-failure in 40 wt% CaCl2 solution (100 °C) of solution annealed matrix and simulated HAZs with various nitrogen contents of DSSs.
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variation in nitrogen content was, the solution annealed DSSs had similar SCC resistance. The average time-to-failure was more than 450 h. This result demonstrated that in CaCl2 solution the SCC susceptibility of DSSs was almost independent of nitrogen content in the range of 0.096–0.165%. But, the sensitivity of SCC in simulated HAZs compared with the solution-annealed matrix of DSSs was dramatically increased. The time-to-failure of HAZ at high cooling rate (Dt8=5 ¼ 5 s) of DSS with 0.096% N was only 18 h. However, the resistance to SCC gradually improved along with the increase in either the cooling time or the nitrogen content. The higher the nitrogen content or the cooling time, the better the resistance to SCC in HAZs of DSSs. For DSS with 0.165% N and at Dt8=5 of 100 s, the time to failure in HAZs increased to about 365 h. At such a low cooling rate, the SCC resistance of the HAZ was raised almost to the same level as the solution annealed DSS. The susceptibility of DSS to SCC depends on many factors such as: alloying elements [30,31], microstructures [2,31–33], applied stresses [34], environments [29,35,36], etc. From the U-bend tests in the 40% CaCl2 solution at 100 °C, the timeto-failure of the HAZ as a function of austenite content is demonstrated in Fig. 6. Clearly, the time-to-failure increased with increasing austenite content which varied with the cooling rate (Dt8=5 ) as revealed in Fig. 4. As mentioned in the previous section, the quantities of Cr2 N decreased with increasing austenite content in the HAZ, and it seemed that the SCC susceptibility of the HAZ increased with increasing amount of Cr2 N precipitation. The change in microstructure in the HAZ
Fig. 6. Effect of austenite content on the time-to-failure in 40 wt% CaCl2 solution (100 °C) of solution annealed matrix and simulated HAZs with various nitrogen contents of DSSs.
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Fig. 7. SEM micrograph showing SCC crack resulting from cross-link of pit, U-bend test for an HAZ of DSS in 40 wt% CaCl2 solution.
has a significant effect on the SCC behavior as will be discussed later, in addition to the effects resulting from austenite content and Cr2 N precipitate. The fracture morphology of SCC in DSSs was further examined. Fig. 7 shows an example of the SEM micrograph of the surface crack of an HAZ after the U-bend test in 40 wt% CaCl2 solution (100 °C). As shown in this figure, pits were formed on the surface. Micro-fissures were found to be induced in the vicinity of pits and gradually cross-linked one from the other. Clearly, pitting-assisted stress corrosion cracking manipulated the fracture at the HAZ of DSS. As mentioned in a previous work [22], Cr2 N precipitates acted as preferential nucleation sites for pitting corrosion. The pitting-assisted SCC as found in Fig. 7 could also be attributed to the presence of Cr2 N in the HAZ of DSS. The occurrence of pitting corrosion was particularly important in the crack initiation during the process of SCC. Fig. 8 shows the near surface cross-section micrographs of the failed U-bend specimens for the solution annealed and the HAZ of a DSS with 0.165 wt% N after testing in the 40 wt% CaCl2 solution at 100 °C. The results revealed that the ferrite phase either in the solution annealed matrix or in the HAZ corroded preferentially in this testing environment. Selective dissolution of ferrite phase could result in surface crack initiation and subsequently induced SCC of DSS, either solution annealed or welded. Selective dissolution depends on the distribution of alloying elements between the a and c phases, the amount of precipitates formed, and the testing solutions used [37,38]. Sridhar and Kolts [37] proposed that corrosion occurred mostly in the austenite for the high nitrogen containing DSS (Ferralium alloy 255, 0.17% N) in both phosphoric acid and sulfuric acid solutions. Whereas, in hydrochloric acid and other chloride environments, corrosion occurred preferentially in the ferrite phase [37]. However, Fourie and Robinson [38] observed that the selective dissolution behavior of 2205 DSS was a function of the chloride concentration in 1 M H2 SO4 solution at 60 °C. At high chloride concentrations (more than 0.2 M NaCl) selective
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Fig. 8. The OM micrograph showing the selective dissolution of ferrite phase in the (a) solution annealed matrix and (b) HAZ (Dt8=5 ¼ 60 s) of 0.165 wt% N DSS after SCC testing in 40 wt% CaCl2 solution.
dissolution of ferrite occurred while selective dissolution of austenite was observed in the absence of chloride. At low chloride concentrations (less than 0.1 M NaCl) both austenite and ferrite phases were dissolved. In this study, preferential corrosion of ferrite phase was probably due to the high concentrations of chloride (40 wt% CaCl2 solution), despite the absence of H2 SO4 . Fig. 9 depicts the cross-section crack propagation path of a secondary crack in a simulated HAZ (Dt8=5 ¼ 5 s) of DSS with 0.096 wt% N in 40 wt% CaCl2 solution at 100 °C. It is interesting to note that the crack propagated in a transgranular manner in the early stage, which was followed by intergranular stress corrosion cracking (IGSCC) in the latter stage. In other words, crack initiation was not confined to the grain boundary even though the crack might propagate intergranularly. The SEM fractograph showing the combined feature of transgranular crack initiation and intergranular crack propagation in the HAZ of of 0.096 wt% N DSS is depicted in Fig. 10. With Dt8=5 ¼ 5 s, GBA was formed in the interfaces between ferrite grains in the HAZ of 0.096 wt% N DSS. The intergranular mode of SCC was thought to be related to the presence of GBA. EDS examination of the fracture surface could be employed to identify the presence of austenite in the grain boundary or at the
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Fig. 9. The cross-sectioned crack propagation path in the HAZ (Dt8=5 ¼ 5 s) of 0.096 wt% N DSS after SCC testing in the 40 wt% CaCl2 solution.
interface of ferrite grains. As depicted in Fig. 10b, the clean and smooth ferrite grain facets were clearly revealed. On top of some of the ferrite grain facets, a thin layer of GBA with an appearance like orange peel was found to exist. This thin layer of GBA could be very brittle and cause the occurrence of IGSCC. For the DSS with 0.096 wt% N, a significant change in the fracture surface morphology of the simulated HAZ was observed upon varying Dt8=5 . At Dt8=5 of 20 s, transgranular stress corrosion cracking (TGSCC) mixed with IGSCC was observed (Fig. 11a). As the cooling time was increased to 100 s, exclusive TGSCC was observed (Fig. 11b). For DSSs with nitrogen contents of 0.135 and 0.165 wt%, the fracture surface morphologies of the simulated HAZs were also affected by the cooling rate. At Dt8=5 of 5 s, mixed TGSCC/IGSCC mode was observed but the percentage of IGSCC in the HAZ of DSS with 0.165 wt% N was less than that of DSS with 0.135 wt% N. As the cooling time increased beyond 20 s, only TGSCC was observed in the HAZs of both the DSSs. Fig. 11c gives an example showing TGSCC in the HAZ, with Dt8=5 ¼ 60 s, of a DSS containing 0.135 wt% of N. As illustrated in Figs. 2 and 3, the microstructure of the simulated HAZ depends on the nitrogen content and the cooling rate. GBA was present in the HAZ with low nitrogen content and cooled at a high cooling rate. The occurrence of IGSCC in that of 0.096 wt% N DSS was believed to be associated with the presence of GBA. The cause of the deleterious effect of GBA is not known and needs further investigation. As the reformed austenite increased in grain size and volume fraction, the resistance to IGSCC also improved. The presence of WA, IGA and PTA seemed not to promote IGSCC in the HAZ of DSS in 40% CaCl2 solution at 100 °C. These types of reformed austenite, on the contrary, exhibited a beneficial effect in delaying crack propagation by deviating its growth path as demonstrated in Fig. 12 for the HAZ, with Dt8=5 ¼ 5 s, of a DSS of 0.165 wt% N. As shown in this micrograph, the crack propagated transgranularly in the ferrite phase and changed path when it encountered WA or IGA.
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Fig. 10. (a) The SEM micrograph of the HAZ (Dt8=5 ¼ 5 s) of 0.096 wt% N DSS after SCC testing in 40 wt% CaCl2 solution, (b) enlargement of (a).
4. Conclusions 1. Grain boundary austenite, WA, IGA and partially transformed austenite could be formed in the HAZ of DSSs. The types and amounts of these reformed austenites in the HAZ depended on the nitrogen content in the DSS and the cooling rate employed. 2. The SCC resistance in the HAZ was inferior to that of the solution annealed DSS. However, the susceptibility to SCC of the HAZ in 40 wt% CaCl2 solution at 100 °C decreased as the amount of reformed austenite increased. The SCC resistance of the HAZ was also found to increase as the nitrogen content in the DSS increased and the cooling rate decreased.
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Fig. 11. The SEM micrograph of the HAZ of DSS after SCC testing in 40 wt% CaCl2 solution, (a) N ¼ 0:096 wt%, Dt8=5 ¼ 20 s, (b) N ¼ 0:096 wt%, Dt8=5 ¼ 100 s, (c) N ¼ 0:135 wt%, Dt8=5 ¼ 60 s.
3. Pitting corrosion and selective dissolution of ferrite phases could assist crack initiation and induce SCC in the HAZ of DSSs. 4. IGSCC could be found in the HAZ when GBA was present. The transformation from IGSCC to TGSCC occurred as GBA was gradually replaced by WA or IGA.
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Fig. 12. The cross-sectioned crack propagation path in the HAZ (Dt8=5 ¼ 5 s) of 0.165 wt% N DSS after SCC testing in the 40 wt% CaCl2 solution.
5. The presence of WA, IGA and PTA exhibited beneficial effect on SCC resistance by deviating the crack propagation path.
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