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Effects of aging at 475 °C on corrosion properties of tungsten-containing duplex stainless steels
Corrosion Science 44 (2002) 2817–2830 www.elsevier.com/locate/corsci
Effects of aging at 475 °C on corrosion properties of tungsten-containing duplex stainless steels Chan-Jin Park, Hyuk-Sang Kwon
*
Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology, 373-1, Guseong-dong, Yuseong-gu, Taejon 305-701, South Korea Received 23 July 2001; accepted 21 March 2002
Abstract Effects of aging at 475 °C on the corrosion and mechanical properties of Fe–25Cr–7Ni– 0.25N–xMo–yW (x ¼ 0–3, y ¼ 0–6) duplex stainless steels were investigated by an anodic polarization test in HCl solution, a modified double-loop electrochemical potentiodynamic reactivation (DL-EPR) test, and an impact test. Corrosion resistance of the alloys was degraded with aging at 475 °C due to the depletion of Cr around a0 precipitates where numerous micropits were formed during the anodic polarization test. Especially for over-aged alloys, a second anodic current loop appeared in the passive region during the anodic polarization in 1 M HCl solution. The peak value of the second anodic current loop as well as the ratio of the maximum current in reactivation loop to that in anodic loop (ir =ia ) determined from the modified DL-EPR test were found to be an effective measure of the precipitation of a0 -phase during the aging. However, the degradation in corrosion resistance and impact toughness of the alloys during the aging was retarded with an increase in the W content of the designed DSS, suggesting that W in duplex stainless steels delays the precipitation rate of a0 -phase due to a slower diffusion rate of W compared with that of Mo in ferrite. Influences of aging on the galvanic corrosion behaviors between austenite and ferrite phases were discussed by atomic force microscopy observation. Ó 2002 Elsevier Science Ltd. All rights reserved. Keywords: Duplex stainless steels; 475 °C embrittlement; Aging; a0 -Phase; Tungsten; Localized corrosion; EPR test (double-loop electrochemical potentiokinetic reactivation); Galvanic corrosion
*
Corresponding author. Tel.: +82-42-869-3326; fax: +82-42-869-3310. E-mail address: [email protected] (H.-S. Kwon).
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1. Introduction Austenitic–ferritic duplex stainless steel (DSS) is very attractive as a structural material in the fields of energy/environmental systems where both high mechanical strength and excellent resistance to localized and stress corrosion are required. Generally, these alloys have two to three times higher yield strength and exhibit greater resistance to localized and stress corrosion than type 300-series austenitic stainless steels at a comparable cost [1–3]. It is well known that DSS undergoes ‘475 °C embrittlement’ due to a decomposition of ferrite phase into a Cr-rich a0 -phase and Fe-rich phase a0 -phase when exposed to temperatures of 300–550 °C. This phase separation can be related to a miscibility gap in the Fe–Cr phase diagram [4]. Depending on the temperature of aging and the chemical composition of alloys, the nature of this decomposition is either spinodal separation or nucleation and growth. At lower temperatures range around 300 °C, only phase splitting due to spinodal decomposition can take place, but above about 300 °C, the precipitation of Cr-rich a0 -phase is a cause of 475 °C embrittlement. The Cr-rich a0 precipitates cause DSS to be embrittled by lowering the mobility of dislocation and by creating microvoid near them in ferrite matrix, and further, they deteriorate corrosion resistance of DSS by forming the Cr-depleted regions around them [4–6]. Since the a0 precipitates have very fine size of about few nm and same crystal structure as the ferrite phase, it is difficult to analyze quantitatively the influence of a0 precipitates on various properties by observing them with an optical or an electron microscopy. Therefore, recently a few researchers have tried to evaluate the degree of 475 °C embrittlement of the alloys in terms of the change in electrochemical properties induced by the precipitation of a0 -phase [7,8]. According to the previous reports [9,10], the partial substitution of Mo by W in 25%Cr DSSs retarded the nucleation and growth rate of sigma (r)-phase during high temperature aging, thereby delaying degradation of corrosion and mechanical properties of the alloys. The retardation of r-phase precipitation by W resulted from its inherently slower diffusion rate compared with that of Mo in DSS. Both Mo and W are major alloying elements constituting the r-phase in W-containing DSSs [10]. Since the precipitation of a0 -phase occurs by diffusion controlled process, the precipitation rate of a0 -phase may be delayed by partial substitution of Mo by W, thereby retarding the 475 °C embrittlement of DSS. The objective of this study is to examine the effects of aging at 475 °C and W on the precipitation kinetics of a0 -phase and the resultant degradation in corrosion resistance of Fe–25Cr–7Ni–0.25N–xMo–yW (x ¼ 0–3, y ¼ 0–6) alloys, and also to develop a new electrochemical method to evaluate the degree of precipitation of a0 phase in DSS.
2. Experimental Experimental alloys were designed so as to have the same value (42) of the pitting resistance equivalent (PRE) by varying the contents of Mo and W for Fe–25Cr–7Ni–
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Table 1 Chemical compositions (wt.%) of designed DSSs Alloy designation
Cr
Ni
Mo
W
N
Fe
3Mo 2Mo–2W 3W–1.5Mo 6W
24.60 25.01 24.82 24.34
6.60 6.81 6.79 6.54
3.12 2.15 1.60 –
– 2.10 3.25 6.30
0.25–0.30 0.25–0.30 0.25–0.30 0.25–0.30
Balance Balance Balance Balance
xMo–yW–0.25N (x þ 1=2y ¼ 3) alloys. The PRE of DSSs, a criterion for pitting resistance, is defined by the experimental equation (1) [11–13], and calculated by chemical compositions (wt.%) of major alloying elements such as Cr, Mo, W, and N that affect significantly the resistance to pitting corrosion of stainless steels. According to the equation, W enhances the resistance to pitting corrosion of DSS half as much as does Mo [14]: PRE ¼ %Cr þ 3:3ð%Mo þ 1=2%WÞ þ 30%N
ð1Þ
The designed alloys were high-purity heats melted in a laboratory scale vacuum induction furnace and cast in 25 kg ingot. Chemical compositions of the alloys were presented in Table 1. In this study, each alloy was named as 3Mo, 2Mo–2W, 3W– 1.5Mo and 6W after its major compositions. The alloys were prepared in the forms of hot-rolled sheet 4 mm thick and plate 12 mm thick. All the alloys were solution annealed for 2 h at 1050 °C, and then aged respectively for 1, 10, 100 and 300 h at 475 °C. The specimens were ground with silicon carbide paper up to 2000 mesh finish before electrochemical tests. Some specimens were polished to 1 lm finish before testing for metallographic observations. Anodic polarization tests were performed at a scan rate of 0.5 mV s1 to investigate the effects of the aging and W addition on the localized corrosion of the alloys in deaerated 4 M NaCl solution at 80 °C, and also in 1 M HCl solution at 30 °C. The electrochemical cell for the tests consisted of a 1 l-multineck flask, as specified in ASTM G5, with a platinum counter-electrode and a saturated calomel electrode (SCE) positioned in a salt bridge with a high-silica tip. In order to measure the degree of depletion of Cr near the a0 precipitates of the designed DSSs with aging at 475 °C, a double-loop electrochemical potentiokinetic reactivation (DL-EPR) test was conducted on the aged alloys in 0.5 M H2 SO4 þ 0:001 M TA (thioacetamid), a solution proposed by Shultze et al. [15]. The degree of Cr-depletion in the alloys was evaluated by measuring the ratio of reactivation peak current (ir ) to activation peak current (ia ) when the potential was applied at a scan rate of 1 mV s1 from 500 to 200 mVSCE , and then reversely to 500 mVSCE . Galvanic corrosion behavior between austenitic and ferritic phases in the designed DSSs was examined by observing the boundary region of the two phases with an atomic force microscopy (AFM) after immersing the aged alloys for 1 h in 1 M HCl solution at 30 °C. The dissolution depth of ferrite grain relative to that of neighboring austenite grain near the phase boundary was considered as a measure of embrittlement as reported in the previous study by Yi and Shoji [8]. The degradation
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of mechanical properties of the designed alloys induced by aging at 475 °C was measured by Charpy V-notch impact tests performed at room temperature. The surface of the specimen subjected to an anodic polarization test was observed with a scanning electron microscopy (SEM). The a0 precipitates formed in DSS during aging were analyzed by a transmission electron microscopy (TEM). Thin foil specimens for TEM observation were prepared electrochemically in 5% perchloric acid plus 95% acetic acid. 3. Results and discussion 3.1. Precipitation of a0 -phase TEM bright field image of the 3Mo alloy aged for 300 h at 475 °C is shown in Fig. 1, where a modulated contrast observed appears to be associated with a spinodal decomposition of ferrite phase to a- and a0 -phase that has occurred during the aging. It was found from analysis of the diffraction pattern in Fig. 1 that a0 -phase has body , which is almost centered cubic (bcc) structure and its lattice constant is about 2.90 A same as that of ferrite matrix. In addition, other precipitates such as carbides were not seen in grain or at grain boundaries. Therefore, it is considered that the variation in mechanical and corrosion properties of DSS with aging was due predominantly to the precipitation of the a0 -phase. 3.2. Effect of tungsten (W) on pitting potential Fig. 2 shows the effects of W content on the pitting potential of the designed alloys in 4 M NaCl solution at 80 °C. The pitting potential of each alloy was taken as an
Fig. 1. TEM bright field image and selected area diffraction pattern for [0 0 1] zone axis in 3Mo alloy aged for 300 h at 475 °C.
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Fig. 2. Pitting potential of designed alloys as a function of W content in deaerated 4 M NaCl solution at 80 °C.
average value determined from five polarization tests. The 3W–1.5Mo alloy exhibited the highest pitting potential among the designed alloys. The results shown in Fig. 2 demonstrated that the resistance to localized corrosion of the designed alloys was maximum when alloyed by 3%W and 1.5%Mo, even if the PRE values of the alloys were equal. Further, the results suggested that there is a synergic effect of Mo and W on enhancing the resistance to pitting corrosion of 25%Cr DSSs, which was most effective when Mo and W is alloyed at the ratio of 1 to 2. Both Mo and W belong to the same family of VIB in a periodic table, and hence will exhibit similar chemical behaviors. These elements are added to stainless steels to improve the resistance to localized corrosion such as pitting and stress corrosion. Okamoto [14] suggested that the contribution of W to the pitting corrosion resistance in DSS is half of that of Mo. Thus, the PRE formula already presented in Eq. (1) has been used as an index of pitting corrosion resistance of DSS. However, Kim and Kwon [10] recently reported that the resistance to localized corrosion of DSS was improved with a partial substitution of Mo by W in alloying composition even although all the designed alloys have the same PRE value of 42. This reveals that there exists a synergic effect of Mo and W on the improvement in the resistance to localized corrosion of DSS. The above results shown in Fig. 2 confirm again that the resistance to pitting of 25%Cr DSSs can be most improved by alloying Mo and W at a specific ratio such as 1 to 2. 3.3. Effect of aging on pitting corrosion resistance In order to investigate the effects of aging at 475 °C on the resistance to pitting corrosion of DSS, the pitting potential of the alloys were determined form anodic
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Fig. 3. Effect of aging on the pitting potentials of the designed alloys determined from anodic polarization tests in 4 M NaCl solution at 80 °C.
polarization curves measured in 4 M NaCl solution at 80 °C, and presented in Fig. 3. While the pitting potentials of the 3Mo and the 2Mo–2W alloy decreased with aging, those of the 3W–1.5Mo and 6W alloys remained relatively high values above 600 mVSCE in spite of aging. Moreover, the degradation in the resistance to pitting corrosion of the alloys with aging was delayed with the ratio of W to Mo being increased. The decrease in the pitting potential of the alloys with aging is due to the precipitation of the a0 -phase and the resultant depletion of Cr around a0 precipitates. On the other hand, the delay in the degradation of the resistance to pitting corrosion of the alloys with an increase of W content appears to be due to the retardation in precipitation rate of a0 -phase by W. Lizlovs and Bond [16] reported that the pitting potential of 18Cr–2Mo Ti stabilized ferritic stainless steels decreased with aging at 475 °C in NaCl solution, and the main cause of that was attributed to the precipitation of the Cr-rich a0 -phase in matrix. 3.4. Effect of aging on the anodic polarization response of DSS in acidic solution Fig. 4 shows effects of aging on the anodic polarization responses of the designed alloys in 1 M HCl solution at 30 °C. It is evident that the critical anodic current density (ic ) and the passive current density (ip ) for each alloy increased with aging as a result of a0 precipitation. Especially, second anodic current loops with the peak current at about 300 mV were appeared in otherwise passive region for all the alloys when aged for 300 h. The second anodic current loop appears to be associated with the formation of Cr-depleted region around the a0 precipitates. In the range of potential where the second anodic current loop is formed, active elements such as Fe and Cr may be dissolved preferentially in the Cr-depleted regions around a0 precipitates with noble alloying element of Ni being enriched on the surface of the alloy, which cause the alloy to be repassivated again with an increase in applied potential,
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Fig. 4. Anodic polarization responses of (a) 3Mo, (b) 2Mo–2W, (c) 3W–1.5Mo, and (d) 6W alloy in 1 M HCl solution at 30 °C.
thereby forming the anodic current loop. When comparing the second anodic current peak for all the alloys, the peak current value decreased with an increase of W content in the designed alloys; the highest peak value for 3Mo alloy, and the lowest peak value for 6W alloy. This again confirms that the precipitation of a0 -phase was retarded effectively by increasing the W content in the designed alloys, thereby producing a less Cr-depleted region around a0 precipitates for alloys containing higher W content. In order to examine why the second anodic current loop appeared only for alloys aged for 300 h, the 3W–1.5Mo specimens with or without aging were anodically polarized up to 600 mVSCE in 1 M HCl solution at 30 °C, and then removed from the environment to observe the surface morphology of the specimens. SEM surface morphologies of the specimens are shown in Fig. 5. For the specimen solution annealed, a uniform corrosion occurred in both ferrite and austenite phases as shown in
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Fig. 5. SEM surface morphologies of 3W–1.5Mo alloy after anodic polarization tests in 1 M HCl solution at 30 °C; (a) solution annealed and (b) aged for 300 h.
Fig. 5(a). In contrast, a severe localized corrosion with numerous micropits occurred within grain as well as along grain boundary in the ferrite phase for the specimen aged for 300 h as shown in Fig. 5(b). This is due to the fact that the precipitation of a0 -phase and the resultant Cr-depletion occurred within the grain and along the grain boundary of the ferrite phase. In summary, the second anodic current loop is formed by the preferential dissolution of Cr-depleted regions associated with the a0 precipitation, and the decrease in the peak current of second anodic current loop with an increase in the W content of the designed alloys resulted from the retardation of the a0 precipitation by W. 3.5. Double-loop electrochemical potentiokinetic reactivation test The degree of Cr-depletion around a0 precipitates formed during aging was examined by (DL-EPR) tests conducted in 0.5 M H2 SO4 þ 0:001 M TA solution at 60 °C, and the results for 3W–1.5Mo alloy are shown in Fig. 6. The reactivation current peak (ir ) was not appeared in the solution annealed sample, whereas a large reactivation current peak was obtained in the sample aged for 300 h as shown in Fig. 6. Two current peaks observed at 324 and 362 mVSCE for the solution annealed sample during the reactivation potential scanning resulted respectively from hydrogen reduction reaction and metal dissolution reaction. The degree of Cr-depletion with aging for 3Mo and 3W–1.5Mo alloys are represented by the ratio of reactivation peak current (ir ) to activation peak current (ia ) that were determined from the double-loop EPR tests as shown in Fig. 7. The ir /ia ratio for 3Mo alloy was much higher than that for 3W–1.5Mo alloy when equivalently aged, which demonstrated that the degree of Cr-depletion in 3Mo alloy is greater than that in 3Mo alloy due to the faster precipitation rate of a0 -phase in the 3Mo alloy. Fig. 8 shows SEM micrographs on the surfaces of the solution annealed or the aged 3W–1.5Mo alloy that have been subject to DL-EPR test. For the aged sample, numerous micropits were distributed uniformly in ferrite grain, whereas the
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Fig. 6. DL-EPR curves for 3W–1.5Mo alloy in 0.5 M H2 SO4 þ 0:001 M TA solution at 60 °C; (a) solution annealed and (b) aged for 100 h.
Fig. 7. Effect of aging on the reactivation/activation current peak ratio (ir /ia ) of 3Mo and 3W–1.5Mo alloys determined from DL-EPR tests.
solution annealed sample showed a smooth surface only with several micropits formed due probably to non-metallic inclusions. Evidently, the numerous micropits formed in ferrite grains of the aged sample confirm that the reactivation current loop observed in the EPR test resulted from Crdepletion due to the precipitation of a0 -phase during the aging at 475 °C. 3.6. Effect of aging on impact toughness Fig. 9 shows the results of Charpy-impact test conducted at room temperature for 3Mo and 3W–1.5Mo alloys that were aged at 475 °C for 1, 10, 100 and 300 h. The impact toughness of the alloys was decreased with aging due to an embrittlement
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Fig. 8. SEM surface morphologies of 3W–1.5Mo alloy after DL-EPR tests; (a) solution annealed and (b) aged for 100 h.
Fig. 9. Effect of aging on the Charpy-impact toughness of 3Mo and 3W–1.5Mo alloys at room temperature.
induced by the precipitation of a0 -phase. Nevertheless, 3W–1.5Mo alloy exhibited relatively higher toughness values compared with those of the 3Mo alloy under the same aged condition, which also corresponds to the results of the electrochemical tests shown in Figs. 4 and 7. However, the impact toughness results in Fig. 9 shows slightly different tendency from the ir /ia ratio in Fig. 7, where after aging for 300 h there is small difference in ir /ia ratio between 3Mo and 3W–1.5Mo alloy but there is clear difference in impact toughness between two. The main causes of the degradation in corrosion resistance and in mechanical properties of DSS with aging at 475 °C are not exactly the same. Strictly speaking, the former is due to the depletion of the matrix of Cr after the precipitation of a0 -phases and the latter is due to the hindrance of a0 -phase to the dislocation movement. The decrease in impact energy of
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the alloys shown in Fig. 9 directly reflects the degree of precipitation of a0 -phase, while the ratio of reactivation to activation current (ir /ia ) in DL-EPR test in Fig. 7 shows the degree of the depletion of Cr with the precipitation of a0 -phase. It is a general phenomenon that the ir /ia ratio increases slowly or even decreases after long periods of aging because Cr in Cr-depleted areas can be partly replenished by the diffusion of Cr from surrounding matrix with an over-aging, even though the volume fraction of precipitate is continuously increased. Thus, it is considered that the ir /ia ratio of the 3Mo alloy increased slowly after aging for 100 h, even though the impact toughness that is directly associated with the volume fraction of a0 -phase decreases abruptly with aging. 3.7. Galvanic corrosion behavior between austenite and ferrite in DSS A galvanic cell is formed when two dissimilar metals are connected electrically while both are immersed in a solution. DSS consisted of austenite and ferrite, therefore, produces numerous galvanic cells between the two phases when exposed to an acidic solution. Generally, ferrite is more active than austenite in DSS due to lower Ni content in ferrite compared with that in austenite, and hence is dissolved preferentially in solution with austenite being galvanically protected. Both solution annealed sample and samples aged respectively for 100 and 300 h at 475 °C were immersed for 1 h in 1 M HCl solution at 30 °C, and then difference in dissolution depth at ferrite grain relative to neighboring austenite grain for each sample was measured by an AFM. The results are shown in Table 2. The dissolution depth of ferrite grain relative to austenite grain was found to decrease with aging. It is noteworthy that the dissolution depth of ferrite grain for the aged sample was found to be much smaller than that for the solution annealed sample. This suggests that galvanic corrosion behavior of the aged alloys is somewhat different from that of the solution annealed alloy. In addition, the dissolution depth of ferrite grain for the aged 3W–1.5Mo alloy was found to be greater than that for the aged 3Mo alloy. AFM images of the solution annealed and the aged samples for the 3W–1.5Mo alloy are presented in Fig. 10. The entire ferrite regions corroded uniformly in the alloy aged for 300 h while most of dissolution occurred into the ferrite region at austenite/ferrite grain boundaries in the solution annealed alloy. These large differences in dissolution morphology between the solution annealed and the aged samples are associated with microstructural differences between the two samples as schematically shown in Fig. 11. Galvanic cells are formed between neighboring austenite Table 2 Dissolution depth of ferrite grain relative to that of neighboring austenite grain for both 3Mo and 3W– ) 1.5Mo alloys with aging at 475 °C (unit: A Aging time (h) 3Mo 3W–1.5Mo
0
100
300
6783 5506
320 802
186 436
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Fig. 10. AFM images of 3W–1.5Mo alloy after immersion for 1 h in 1 M HCl solution; (a) solution annealed and (b) aged for 100 h.
Fig. 11. Schematic of galvanic corrosion behaviors of (a) solution annealed and (b) aged sample.
grain and ferrite grain in the solution annealed sample, and the ferrite region nearer the grain boundary may be dissolved deeper than the rest of the ferrite grain when convection and migration of ions are not so vigorous. In contrast to this, numerous microgalvanic cells of about few nm formed between Cr-rich a0 precipitate and neighboring Fe-rich phase in ferrite matrix of the aged alloy lead to uniform corrosion in ferrite phase as shown in Fig. 11. The deeper dissolution depth of ferritic phase relative to that of austenitic phase for the aged 1.5Mo–3W alloy compared with that of the aged 3Mo alloy due primarily to the fact that the precipitation of a0 is more delayed or retarded in the 1.5Mo–3W alloy, producing less microcells between a0 - and a-phases than the 3Mo alloy.
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3.8. Role of W to retard the precipitation of a0 -phase Results of this work showed that the impact toughness and the resistance to localized corrosion of DSS were significantly degraded by the precipitation of a0 -phase with aging at 475 °C. It is well known that a0 -phase is formed by a spinodal decomposition process in ferrite phase during aging. The spinodal decomposition is generally known as a transformation process where there is no barrier to nucleation. When aged at temperature ranges where the spinodal decomposition can occur, the ferrite phase in DSS is separated into Cr-rich and Fe-rich phases by a small compositional fluctuation. This process is accompanied with an up-hill diffusion to the opposite direction of compositional gradient [17], and proceeds until reaching an equilibrium composition. A question is raised why the precipitation rate of a0 -phase is retarded with increasing the %W/%Mo ratio in 25%Cr based DSS during aging at 475 °C. Generally, both Mo and W in DSS are known to promote 475 °C embrittlement [18]. However, it was reported that the diffusion rate of W at 475 °C was 10–100 times slower than that of Mo in iron or ferrous alloys [19–21]. Thus, it appears that W diffuses more slowly than does Mo in the spinodal decomposition of ferritic phase during aging, and further it also obstructs the diffusion of other elements during the aging, thereby delaying the precipitation rate of a0 -phase.
4. Conclusion 1. The corrosion and mechanical properties of the designed DSS were degraded with aging at 475 °C due to the precipitation of a0 -phase and its resultant Cr-depletion around the precipitates. However, the degradation rate of the alloys was retarded with an increase in the W content in the designed alloys. 2. The degradation in the corrosion resistance of the DSS aged at 475 °C, confirmed by the increase in both the critical current density and the passivation current density in the anodic polarization curve, results from the Cr-depletion around a0 precipitates. Especially, for the over-aged alloys, a second anodic current loop appeared in the passive range during anodic polarization in 1 M HCl solution due to the Cr-depletion by precipitation of Cr-rich a0 -phase. 3. The degree of Cr-depletion with aging for DSS can be evaluated by the ratio of ir / ia determined from the modified DL-EPR test. 4. Both the peak value of the second anodic current loop determined from anodic polarization curve in 1 M HCl solution and the ratio of (ir /ia ) determined from the modified DL-EPR tests were found to be an effective measure of the precipitation of a0 -phase during the aging, and correlated with the degree of embrittlement of DSS by the aging at 475 °C. 5. The degradation in corrosion resistance and impact toughness of the alloys with aging was retarded with an increase in the W content of the designed DSS, suggesting that W in DSSs delays the precipitation rate of a0 -phase due to a slower diffusion rate of W compared with that of Mo in ferritic phase.
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6. Galvanic corrosion behavior of the aged alloy exhibited somewhat different from that of the solution annealed alloy. For the solution annealed, ferrite nearer the a=c grain boundary was dissolved deeper by galvanic corrosion between a and c, whereas the ferrite phase of the aged alloy was almost uniformly dissolved due to formation of numerous microgalvanic cells between Cr-rich a0 precipitate and surrounding Fe-rich a-phase.
References [1] S. Bernhardsson, Proceedings of the Conference on Duplex Stainless Steels 91, 1991, p. 185. [2] S. Karlsson, Proceedings of the Conference on Stainless Steels 84, 1985, p. 438. [3] R. Kiesseling, S. Bernhardsson, Proceedings of the Conference on Duplex Stainless Steels 91, 1991, p. 605. [4] R.N. Wright, in: R.A. Rula (Ed.), Toughness of Ferritic Stainless Steels, ASTM STP 706, 1979, p. 2. [5] N. Sridhar, J. Kolts, L.H. Flashe, J. Met. March (1983) 31. [6] M.J. Marcinkowski, R.M. Fisher, A. Szirmae, Trans. AIME 230 (1964) 676. [7] X.C. Jiang, T. Yoshimura, Y. Ishikawa, J. Electrochem. Soc. 139 (1992) 1001. [8] Y.S. Yi, T. Shoji, J. Nucl. Mater. 231 (1996) 20. [9] B.W. Oh, J.I. Kim, S.W. Jeong, H.-S. Kwon, S.H. Hong, Y.G. Kim, Proceedings of 1st European Stainless Steel Conference, vol. 3, 1993, p. 59. [10] J.S. Kim, H.-S. Kwon, Corrosion 55 (1999) 512. [11] M. Barteri, M.G. Mecozzi, I. Nembrini, Proceedings of the Conference on Duplex Stainless Steels 94, paper no. 60, 1994. [12] J.M. Nicholls, Proceedings of the Conference on Duplex Stainless Steels 94, paper no. KIII, 1994. [13] L.F. Garfias-Mesias, J.M. Sykes, C.D. Tuck, Corros. Sci. 38 (1996) 1319. [14] H. Okamoto, Proceedings Conference on Application of Stainless Steels 92, vol. 1, 1992, p. 1251. [15] S. Shultze, J. Golner, K. Eick, P. Veit, I. Garz, Proceedings of the Conference on Duplex Stainless Steels 97, 1997, p. 639. [16] E.A. Lizlovs, A.P. Bond, J. Electrochem. Soc. May (1975) 589. [17] D.A. Porter, K.E. Easterling, in: Phase transformations in metals and alloys, Chapman & Hall, New York, 1991, p. 308. [18] J. Charlse, Proceedings of the Conference on Duplex Stainless Steels 91, 1991, p. 3. [19] V.T. Borisov, V.M. Golikov, G.V. Sherbedinskiy, Phys. Met. Metallogr. 22 (1996) 175. [20] H.M. Lee, S.M. Allen, M. Grujicic, Met. Trans. 22A (1991) 2869. [21] J.E. Shackelford, W. Alexander, J.S. Park, Materials Science and Engineering Handbook, second ed., CRC Press, Boca Raton, FL, 1994, p. 219.
Effects of aging at 475 °C on corrosion properties of tungsten-containing duplex stainless steels Chan-Jin Park, Hyuk-Sang Kwon
*
Department of Materials Science and Engineering, Korea Advanced Institute of Science and Technology, 373-1, Guseong-dong, Yuseong-gu, Taejon 305-701, South Korea Received 23 July 2001; accepted 21 March 2002
Abstract Effects of aging at 475 °C on the corrosion and mechanical properties of Fe–25Cr–7Ni– 0.25N–xMo–yW (x ¼ 0–3, y ¼ 0–6) duplex stainless steels were investigated by an anodic polarization test in HCl solution, a modified double-loop electrochemical potentiodynamic reactivation (DL-EPR) test, and an impact test. Corrosion resistance of the alloys was degraded with aging at 475 °C due to the depletion of Cr around a0 precipitates where numerous micropits were formed during the anodic polarization test. Especially for over-aged alloys, a second anodic current loop appeared in the passive region during the anodic polarization in 1 M HCl solution. The peak value of the second anodic current loop as well as the ratio of the maximum current in reactivation loop to that in anodic loop (ir =ia ) determined from the modified DL-EPR test were found to be an effective measure of the precipitation of a0 -phase during the aging. However, the degradation in corrosion resistance and impact toughness of the alloys during the aging was retarded with an increase in the W content of the designed DSS, suggesting that W in duplex stainless steels delays the precipitation rate of a0 -phase due to a slower diffusion rate of W compared with that of Mo in ferrite. Influences of aging on the galvanic corrosion behaviors between austenite and ferrite phases were discussed by atomic force microscopy observation. Ó 2002 Elsevier Science Ltd. All rights reserved. Keywords: Duplex stainless steels; 475 °C embrittlement; Aging; a0 -Phase; Tungsten; Localized corrosion; EPR test (double-loop electrochemical potentiokinetic reactivation); Galvanic corrosion
*
Corresponding author. Tel.: +82-42-869-3326; fax: +82-42-869-3310. E-mail address: [email protected] (H.-S. Kwon).
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1. Introduction Austenitic–ferritic duplex stainless steel (DSS) is very attractive as a structural material in the fields of energy/environmental systems where both high mechanical strength and excellent resistance to localized and stress corrosion are required. Generally, these alloys have two to three times higher yield strength and exhibit greater resistance to localized and stress corrosion than type 300-series austenitic stainless steels at a comparable cost [1–3]. It is well known that DSS undergoes ‘475 °C embrittlement’ due to a decomposition of ferrite phase into a Cr-rich a0 -phase and Fe-rich phase a0 -phase when exposed to temperatures of 300–550 °C. This phase separation can be related to a miscibility gap in the Fe–Cr phase diagram [4]. Depending on the temperature of aging and the chemical composition of alloys, the nature of this decomposition is either spinodal separation or nucleation and growth. At lower temperatures range around 300 °C, only phase splitting due to spinodal decomposition can take place, but above about 300 °C, the precipitation of Cr-rich a0 -phase is a cause of 475 °C embrittlement. The Cr-rich a0 precipitates cause DSS to be embrittled by lowering the mobility of dislocation and by creating microvoid near them in ferrite matrix, and further, they deteriorate corrosion resistance of DSS by forming the Cr-depleted regions around them [4–6]. Since the a0 precipitates have very fine size of about few nm and same crystal structure as the ferrite phase, it is difficult to analyze quantitatively the influence of a0 precipitates on various properties by observing them with an optical or an electron microscopy. Therefore, recently a few researchers have tried to evaluate the degree of 475 °C embrittlement of the alloys in terms of the change in electrochemical properties induced by the precipitation of a0 -phase [7,8]. According to the previous reports [9,10], the partial substitution of Mo by W in 25%Cr DSSs retarded the nucleation and growth rate of sigma (r)-phase during high temperature aging, thereby delaying degradation of corrosion and mechanical properties of the alloys. The retardation of r-phase precipitation by W resulted from its inherently slower diffusion rate compared with that of Mo in DSS. Both Mo and W are major alloying elements constituting the r-phase in W-containing DSSs [10]. Since the precipitation of a0 -phase occurs by diffusion controlled process, the precipitation rate of a0 -phase may be delayed by partial substitution of Mo by W, thereby retarding the 475 °C embrittlement of DSS. The objective of this study is to examine the effects of aging at 475 °C and W on the precipitation kinetics of a0 -phase and the resultant degradation in corrosion resistance of Fe–25Cr–7Ni–0.25N–xMo–yW (x ¼ 0–3, y ¼ 0–6) alloys, and also to develop a new electrochemical method to evaluate the degree of precipitation of a0 phase in DSS.
2. Experimental Experimental alloys were designed so as to have the same value (42) of the pitting resistance equivalent (PRE) by varying the contents of Mo and W for Fe–25Cr–7Ni–
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Table 1 Chemical compositions (wt.%) of designed DSSs Alloy designation
Cr
Ni
Mo
W
N
Fe
3Mo 2Mo–2W 3W–1.5Mo 6W
24.60 25.01 24.82 24.34
6.60 6.81 6.79 6.54
3.12 2.15 1.60 –
– 2.10 3.25 6.30
0.25–0.30 0.25–0.30 0.25–0.30 0.25–0.30
Balance Balance Balance Balance
xMo–yW–0.25N (x þ 1=2y ¼ 3) alloys. The PRE of DSSs, a criterion for pitting resistance, is defined by the experimental equation (1) [11–13], and calculated by chemical compositions (wt.%) of major alloying elements such as Cr, Mo, W, and N that affect significantly the resistance to pitting corrosion of stainless steels. According to the equation, W enhances the resistance to pitting corrosion of DSS half as much as does Mo [14]: PRE ¼ %Cr þ 3:3ð%Mo þ 1=2%WÞ þ 30%N
ð1Þ
The designed alloys were high-purity heats melted in a laboratory scale vacuum induction furnace and cast in 25 kg ingot. Chemical compositions of the alloys were presented in Table 1. In this study, each alloy was named as 3Mo, 2Mo–2W, 3W– 1.5Mo and 6W after its major compositions. The alloys were prepared in the forms of hot-rolled sheet 4 mm thick and plate 12 mm thick. All the alloys were solution annealed for 2 h at 1050 °C, and then aged respectively for 1, 10, 100 and 300 h at 475 °C. The specimens were ground with silicon carbide paper up to 2000 mesh finish before electrochemical tests. Some specimens were polished to 1 lm finish before testing for metallographic observations. Anodic polarization tests were performed at a scan rate of 0.5 mV s1 to investigate the effects of the aging and W addition on the localized corrosion of the alloys in deaerated 4 M NaCl solution at 80 °C, and also in 1 M HCl solution at 30 °C. The electrochemical cell for the tests consisted of a 1 l-multineck flask, as specified in ASTM G5, with a platinum counter-electrode and a saturated calomel electrode (SCE) positioned in a salt bridge with a high-silica tip. In order to measure the degree of depletion of Cr near the a0 precipitates of the designed DSSs with aging at 475 °C, a double-loop electrochemical potentiokinetic reactivation (DL-EPR) test was conducted on the aged alloys in 0.5 M H2 SO4 þ 0:001 M TA (thioacetamid), a solution proposed by Shultze et al. [15]. The degree of Cr-depletion in the alloys was evaluated by measuring the ratio of reactivation peak current (ir ) to activation peak current (ia ) when the potential was applied at a scan rate of 1 mV s1 from 500 to 200 mVSCE , and then reversely to 500 mVSCE . Galvanic corrosion behavior between austenitic and ferritic phases in the designed DSSs was examined by observing the boundary region of the two phases with an atomic force microscopy (AFM) after immersing the aged alloys for 1 h in 1 M HCl solution at 30 °C. The dissolution depth of ferrite grain relative to that of neighboring austenite grain near the phase boundary was considered as a measure of embrittlement as reported in the previous study by Yi and Shoji [8]. The degradation
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of mechanical properties of the designed alloys induced by aging at 475 °C was measured by Charpy V-notch impact tests performed at room temperature. The surface of the specimen subjected to an anodic polarization test was observed with a scanning electron microscopy (SEM). The a0 precipitates formed in DSS during aging were analyzed by a transmission electron microscopy (TEM). Thin foil specimens for TEM observation were prepared electrochemically in 5% perchloric acid plus 95% acetic acid. 3. Results and discussion 3.1. Precipitation of a0 -phase TEM bright field image of the 3Mo alloy aged for 300 h at 475 °C is shown in Fig. 1, where a modulated contrast observed appears to be associated with a spinodal decomposition of ferrite phase to a- and a0 -phase that has occurred during the aging. It was found from analysis of the diffraction pattern in Fig. 1 that a0 -phase has body , which is almost centered cubic (bcc) structure and its lattice constant is about 2.90 A same as that of ferrite matrix. In addition, other precipitates such as carbides were not seen in grain or at grain boundaries. Therefore, it is considered that the variation in mechanical and corrosion properties of DSS with aging was due predominantly to the precipitation of the a0 -phase. 3.2. Effect of tungsten (W) on pitting potential Fig. 2 shows the effects of W content on the pitting potential of the designed alloys in 4 M NaCl solution at 80 °C. The pitting potential of each alloy was taken as an
Fig. 1. TEM bright field image and selected area diffraction pattern for [0 0 1] zone axis in 3Mo alloy aged for 300 h at 475 °C.
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Fig. 2. Pitting potential of designed alloys as a function of W content in deaerated 4 M NaCl solution at 80 °C.
average value determined from five polarization tests. The 3W–1.5Mo alloy exhibited the highest pitting potential among the designed alloys. The results shown in Fig. 2 demonstrated that the resistance to localized corrosion of the designed alloys was maximum when alloyed by 3%W and 1.5%Mo, even if the PRE values of the alloys were equal. Further, the results suggested that there is a synergic effect of Mo and W on enhancing the resistance to pitting corrosion of 25%Cr DSSs, which was most effective when Mo and W is alloyed at the ratio of 1 to 2. Both Mo and W belong to the same family of VIB in a periodic table, and hence will exhibit similar chemical behaviors. These elements are added to stainless steels to improve the resistance to localized corrosion such as pitting and stress corrosion. Okamoto [14] suggested that the contribution of W to the pitting corrosion resistance in DSS is half of that of Mo. Thus, the PRE formula already presented in Eq. (1) has been used as an index of pitting corrosion resistance of DSS. However, Kim and Kwon [10] recently reported that the resistance to localized corrosion of DSS was improved with a partial substitution of Mo by W in alloying composition even although all the designed alloys have the same PRE value of 42. This reveals that there exists a synergic effect of Mo and W on the improvement in the resistance to localized corrosion of DSS. The above results shown in Fig. 2 confirm again that the resistance to pitting of 25%Cr DSSs can be most improved by alloying Mo and W at a specific ratio such as 1 to 2. 3.3. Effect of aging on pitting corrosion resistance In order to investigate the effects of aging at 475 °C on the resistance to pitting corrosion of DSS, the pitting potential of the alloys were determined form anodic
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Fig. 3. Effect of aging on the pitting potentials of the designed alloys determined from anodic polarization tests in 4 M NaCl solution at 80 °C.
polarization curves measured in 4 M NaCl solution at 80 °C, and presented in Fig. 3. While the pitting potentials of the 3Mo and the 2Mo–2W alloy decreased with aging, those of the 3W–1.5Mo and 6W alloys remained relatively high values above 600 mVSCE in spite of aging. Moreover, the degradation in the resistance to pitting corrosion of the alloys with aging was delayed with the ratio of W to Mo being increased. The decrease in the pitting potential of the alloys with aging is due to the precipitation of the a0 -phase and the resultant depletion of Cr around a0 precipitates. On the other hand, the delay in the degradation of the resistance to pitting corrosion of the alloys with an increase of W content appears to be due to the retardation in precipitation rate of a0 -phase by W. Lizlovs and Bond [16] reported that the pitting potential of 18Cr–2Mo Ti stabilized ferritic stainless steels decreased with aging at 475 °C in NaCl solution, and the main cause of that was attributed to the precipitation of the Cr-rich a0 -phase in matrix. 3.4. Effect of aging on the anodic polarization response of DSS in acidic solution Fig. 4 shows effects of aging on the anodic polarization responses of the designed alloys in 1 M HCl solution at 30 °C. It is evident that the critical anodic current density (ic ) and the passive current density (ip ) for each alloy increased with aging as a result of a0 precipitation. Especially, second anodic current loops with the peak current at about 300 mV were appeared in otherwise passive region for all the alloys when aged for 300 h. The second anodic current loop appears to be associated with the formation of Cr-depleted region around the a0 precipitates. In the range of potential where the second anodic current loop is formed, active elements such as Fe and Cr may be dissolved preferentially in the Cr-depleted regions around a0 precipitates with noble alloying element of Ni being enriched on the surface of the alloy, which cause the alloy to be repassivated again with an increase in applied potential,
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Fig. 4. Anodic polarization responses of (a) 3Mo, (b) 2Mo–2W, (c) 3W–1.5Mo, and (d) 6W alloy in 1 M HCl solution at 30 °C.
thereby forming the anodic current loop. When comparing the second anodic current peak for all the alloys, the peak current value decreased with an increase of W content in the designed alloys; the highest peak value for 3Mo alloy, and the lowest peak value for 6W alloy. This again confirms that the precipitation of a0 -phase was retarded effectively by increasing the W content in the designed alloys, thereby producing a less Cr-depleted region around a0 precipitates for alloys containing higher W content. In order to examine why the second anodic current loop appeared only for alloys aged for 300 h, the 3W–1.5Mo specimens with or without aging were anodically polarized up to 600 mVSCE in 1 M HCl solution at 30 °C, and then removed from the environment to observe the surface morphology of the specimens. SEM surface morphologies of the specimens are shown in Fig. 5. For the specimen solution annealed, a uniform corrosion occurred in both ferrite and austenite phases as shown in
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Fig. 5. SEM surface morphologies of 3W–1.5Mo alloy after anodic polarization tests in 1 M HCl solution at 30 °C; (a) solution annealed and (b) aged for 300 h.
Fig. 5(a). In contrast, a severe localized corrosion with numerous micropits occurred within grain as well as along grain boundary in the ferrite phase for the specimen aged for 300 h as shown in Fig. 5(b). This is due to the fact that the precipitation of a0 -phase and the resultant Cr-depletion occurred within the grain and along the grain boundary of the ferrite phase. In summary, the second anodic current loop is formed by the preferential dissolution of Cr-depleted regions associated with the a0 precipitation, and the decrease in the peak current of second anodic current loop with an increase in the W content of the designed alloys resulted from the retardation of the a0 precipitation by W. 3.5. Double-loop electrochemical potentiokinetic reactivation test The degree of Cr-depletion around a0 precipitates formed during aging was examined by (DL-EPR) tests conducted in 0.5 M H2 SO4 þ 0:001 M TA solution at 60 °C, and the results for 3W–1.5Mo alloy are shown in Fig. 6. The reactivation current peak (ir ) was not appeared in the solution annealed sample, whereas a large reactivation current peak was obtained in the sample aged for 300 h as shown in Fig. 6. Two current peaks observed at 324 and 362 mVSCE for the solution annealed sample during the reactivation potential scanning resulted respectively from hydrogen reduction reaction and metal dissolution reaction. The degree of Cr-depletion with aging for 3Mo and 3W–1.5Mo alloys are represented by the ratio of reactivation peak current (ir ) to activation peak current (ia ) that were determined from the double-loop EPR tests as shown in Fig. 7. The ir /ia ratio for 3Mo alloy was much higher than that for 3W–1.5Mo alloy when equivalently aged, which demonstrated that the degree of Cr-depletion in 3Mo alloy is greater than that in 3Mo alloy due to the faster precipitation rate of a0 -phase in the 3Mo alloy. Fig. 8 shows SEM micrographs on the surfaces of the solution annealed or the aged 3W–1.5Mo alloy that have been subject to DL-EPR test. For the aged sample, numerous micropits were distributed uniformly in ferrite grain, whereas the
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Fig. 6. DL-EPR curves for 3W–1.5Mo alloy in 0.5 M H2 SO4 þ 0:001 M TA solution at 60 °C; (a) solution annealed and (b) aged for 100 h.
Fig. 7. Effect of aging on the reactivation/activation current peak ratio (ir /ia ) of 3Mo and 3W–1.5Mo alloys determined from DL-EPR tests.
solution annealed sample showed a smooth surface only with several micropits formed due probably to non-metallic inclusions. Evidently, the numerous micropits formed in ferrite grains of the aged sample confirm that the reactivation current loop observed in the EPR test resulted from Crdepletion due to the precipitation of a0 -phase during the aging at 475 °C. 3.6. Effect of aging on impact toughness Fig. 9 shows the results of Charpy-impact test conducted at room temperature for 3Mo and 3W–1.5Mo alloys that were aged at 475 °C for 1, 10, 100 and 300 h. The impact toughness of the alloys was decreased with aging due to an embrittlement
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Fig. 8. SEM surface morphologies of 3W–1.5Mo alloy after DL-EPR tests; (a) solution annealed and (b) aged for 100 h.
Fig. 9. Effect of aging on the Charpy-impact toughness of 3Mo and 3W–1.5Mo alloys at room temperature.
induced by the precipitation of a0 -phase. Nevertheless, 3W–1.5Mo alloy exhibited relatively higher toughness values compared with those of the 3Mo alloy under the same aged condition, which also corresponds to the results of the electrochemical tests shown in Figs. 4 and 7. However, the impact toughness results in Fig. 9 shows slightly different tendency from the ir /ia ratio in Fig. 7, where after aging for 300 h there is small difference in ir /ia ratio between 3Mo and 3W–1.5Mo alloy but there is clear difference in impact toughness between two. The main causes of the degradation in corrosion resistance and in mechanical properties of DSS with aging at 475 °C are not exactly the same. Strictly speaking, the former is due to the depletion of the matrix of Cr after the precipitation of a0 -phases and the latter is due to the hindrance of a0 -phase to the dislocation movement. The decrease in impact energy of
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the alloys shown in Fig. 9 directly reflects the degree of precipitation of a0 -phase, while the ratio of reactivation to activation current (ir /ia ) in DL-EPR test in Fig. 7 shows the degree of the depletion of Cr with the precipitation of a0 -phase. It is a general phenomenon that the ir /ia ratio increases slowly or even decreases after long periods of aging because Cr in Cr-depleted areas can be partly replenished by the diffusion of Cr from surrounding matrix with an over-aging, even though the volume fraction of precipitate is continuously increased. Thus, it is considered that the ir /ia ratio of the 3Mo alloy increased slowly after aging for 100 h, even though the impact toughness that is directly associated with the volume fraction of a0 -phase decreases abruptly with aging. 3.7. Galvanic corrosion behavior between austenite and ferrite in DSS A galvanic cell is formed when two dissimilar metals are connected electrically while both are immersed in a solution. DSS consisted of austenite and ferrite, therefore, produces numerous galvanic cells between the two phases when exposed to an acidic solution. Generally, ferrite is more active than austenite in DSS due to lower Ni content in ferrite compared with that in austenite, and hence is dissolved preferentially in solution with austenite being galvanically protected. Both solution annealed sample and samples aged respectively for 100 and 300 h at 475 °C were immersed for 1 h in 1 M HCl solution at 30 °C, and then difference in dissolution depth at ferrite grain relative to neighboring austenite grain for each sample was measured by an AFM. The results are shown in Table 2. The dissolution depth of ferrite grain relative to austenite grain was found to decrease with aging. It is noteworthy that the dissolution depth of ferrite grain for the aged sample was found to be much smaller than that for the solution annealed sample. This suggests that galvanic corrosion behavior of the aged alloys is somewhat different from that of the solution annealed alloy. In addition, the dissolution depth of ferrite grain for the aged 3W–1.5Mo alloy was found to be greater than that for the aged 3Mo alloy. AFM images of the solution annealed and the aged samples for the 3W–1.5Mo alloy are presented in Fig. 10. The entire ferrite regions corroded uniformly in the alloy aged for 300 h while most of dissolution occurred into the ferrite region at austenite/ferrite grain boundaries in the solution annealed alloy. These large differences in dissolution morphology between the solution annealed and the aged samples are associated with microstructural differences between the two samples as schematically shown in Fig. 11. Galvanic cells are formed between neighboring austenite Table 2 Dissolution depth of ferrite grain relative to that of neighboring austenite grain for both 3Mo and 3W– ) 1.5Mo alloys with aging at 475 °C (unit: A Aging time (h) 3Mo 3W–1.5Mo
0
100
300
6783 5506
320 802
186 436
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Fig. 10. AFM images of 3W–1.5Mo alloy after immersion for 1 h in 1 M HCl solution; (a) solution annealed and (b) aged for 100 h.
Fig. 11. Schematic of galvanic corrosion behaviors of (a) solution annealed and (b) aged sample.
grain and ferrite grain in the solution annealed sample, and the ferrite region nearer the grain boundary may be dissolved deeper than the rest of the ferrite grain when convection and migration of ions are not so vigorous. In contrast to this, numerous microgalvanic cells of about few nm formed between Cr-rich a0 precipitate and neighboring Fe-rich phase in ferrite matrix of the aged alloy lead to uniform corrosion in ferrite phase as shown in Fig. 11. The deeper dissolution depth of ferritic phase relative to that of austenitic phase for the aged 1.5Mo–3W alloy compared with that of the aged 3Mo alloy due primarily to the fact that the precipitation of a0 is more delayed or retarded in the 1.5Mo–3W alloy, producing less microcells between a0 - and a-phases than the 3Mo alloy.
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3.8. Role of W to retard the precipitation of a0 -phase Results of this work showed that the impact toughness and the resistance to localized corrosion of DSS were significantly degraded by the precipitation of a0 -phase with aging at 475 °C. It is well known that a0 -phase is formed by a spinodal decomposition process in ferrite phase during aging. The spinodal decomposition is generally known as a transformation process where there is no barrier to nucleation. When aged at temperature ranges where the spinodal decomposition can occur, the ferrite phase in DSS is separated into Cr-rich and Fe-rich phases by a small compositional fluctuation. This process is accompanied with an up-hill diffusion to the opposite direction of compositional gradient [17], and proceeds until reaching an equilibrium composition. A question is raised why the precipitation rate of a0 -phase is retarded with increasing the %W/%Mo ratio in 25%Cr based DSS during aging at 475 °C. Generally, both Mo and W in DSS are known to promote 475 °C embrittlement [18]. However, it was reported that the diffusion rate of W at 475 °C was 10–100 times slower than that of Mo in iron or ferrous alloys [19–21]. Thus, it appears that W diffuses more slowly than does Mo in the spinodal decomposition of ferritic phase during aging, and further it also obstructs the diffusion of other elements during the aging, thereby delaying the precipitation rate of a0 -phase.
4. Conclusion 1. The corrosion and mechanical properties of the designed DSS were degraded with aging at 475 °C due to the precipitation of a0 -phase and its resultant Cr-depletion around the precipitates. However, the degradation rate of the alloys was retarded with an increase in the W content in the designed alloys. 2. The degradation in the corrosion resistance of the DSS aged at 475 °C, confirmed by the increase in both the critical current density and the passivation current density in the anodic polarization curve, results from the Cr-depletion around a0 precipitates. Especially, for the over-aged alloys, a second anodic current loop appeared in the passive range during anodic polarization in 1 M HCl solution due to the Cr-depletion by precipitation of Cr-rich a0 -phase. 3. The degree of Cr-depletion with aging for DSS can be evaluated by the ratio of ir / ia determined from the modified DL-EPR test. 4. Both the peak value of the second anodic current loop determined from anodic polarization curve in 1 M HCl solution and the ratio of (ir /ia ) determined from the modified DL-EPR tests were found to be an effective measure of the precipitation of a0 -phase during the aging, and correlated with the degree of embrittlement of DSS by the aging at 475 °C. 5. The degradation in corrosion resistance and impact toughness of the alloys with aging was retarded with an increase in the W content of the designed DSS, suggesting that W in DSSs delays the precipitation rate of a0 -phase due to a slower diffusion rate of W compared with that of Mo in ferritic phase.
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6. Galvanic corrosion behavior of the aged alloy exhibited somewhat different from that of the solution annealed alloy. For the solution annealed, ferrite nearer the a=c grain boundary was dissolved deeper by galvanic corrosion between a and c, whereas the ferrite phase of the aged alloy was almost uniformly dissolved due to formation of numerous microgalvanic cells between Cr-rich a0 precipitate and surrounding Fe-rich a-phase.
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