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Effects of Cu on the precipitation of intermetallic compounds and the intergranular corrosion of hyper duplex stainless steels
Corrosion Science 66 (2013) 217–224
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Effects of Cu on the precipitation of intermetallic compounds and the intergranular corrosion of hyper duplex stainless steels Soon-Hyeok Jeon a, Soon-Tae Kim a, In-Sung Lee a, Ji-Soo Kim b, Kwang-Tae Kim b, Yong-Soo Park a,⇑ a b
Department of Material Science and Engineering, Yonsei University, 134 Shinchon-dong, Seodaemun-gu, Seoul 120-749, Republic of Korea Stainless Steel Research Group, POSCO Technical Research Laboratories Goedong-dong, Nam-Gu, Pohang, Gyeongbuk 790-785, Republic of Korea
a r t i c l e
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Article history: Received 27 June 2012 Accepted 18 September 2012 Available online 27 September 2012 Keywords: A. Copper A. Stainless B. Polarization B. SEM C. Intergranular corrosion
a b s t r a c t To elucidate the effects of Cu on the precipitation of intermetallic compounds and the intergranular corrosion of the hyper duplex stainless steels, a double loop potentiokinetic reactivation test, a scanning electron microscope analysis and thermodynamic calculation were conducted. The addition of Cu to the base alloy reduces the total amount of intermetallic compounds. Particularly, Cu addition to the base alloy results in pronouncedly suppressing the amount of sigma phase whereas it slightly increases chi phases. The Cu added alloy reduces the degree of sensitization due to the delayed precipitation of intermetallic compounds, compared with that of the base alloy. Ó 2012 Elsevier Ltd. All rights reserved.
1. Introduction Duplex stainless steel (DSS) is the stainless steel that has microstructure where both ferrite (a) and austenite (c) phases are present in approximately equal volume fraction. In general, it is well known that standard stainless steel such as UNS S32205 is defined as DSS with a pitting resistance equivalent number (PREN = Cr + 3.3 (Mo + 0.5 W) + 16 N [1,2]) of 35 and super duplex stainless steels (SDSSs) such as UNS S32750, UNS S32760 and UNS S32550 are defined as DSSs with PREN of 40–45. SDSSs have been increasingly used for various applications such as power plants, desalination facilities and chemical plants due to high resistance to pitting and crevice corrosion, excellent mechanical properties and a relatively low cost, compared with other higher performance materials such as super austenite stainless steel (SASS) [3–5]. However, in heat exchanger application, corrosion resistance of SDSS is insufficient for higher temperature service or for a long service life, and where materials, with even higher corrosion resistance, are needed. Hyper duplex stainless steel (HDSS) such as UNS S32707 is defined as a highly alloyed duplex stainless steel with a PREN in excess of 45 [1]. HDSSs were developed to meet industry demands for higher operating temperatures and longer run times and replace the costly SASS that has high levels of Ni and Mo contents [6,7]. The existing researches were progressed for DSS and SDSS [8–12]. However, few studies have been focused
⇑ Corresponding author. Tel.: +82 2 2123 5833; fax: +82 2 362 9199. E-mail address: [email protected] (Y.-S. Park). 0010-938X/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.corsci.2012.09.023
on HDSS, having higher Cr, Mo and W than DSS and SDSS, to substitute the 6% Mo austenite stainless steels (ASSs) in more severe environment. The intermetallic phases such as sigma (r) phase and chi (v) phase are the intermetallic phases that forms at high temperature 600–950 °C [13,14]. These intermetallic phases lower the fracture toughness and the corrosion resistance significantly [15–17]. To avoid this embrittlement and corrosion, DSSs are subjected to solution heat treatment followed by water quenching in the final stage of production or fabrication. The PREN 40 grade SDSS (e.g. UNS S32750) contains larger amounts of Cr and Mo than the PREN 35 grade DSS (e.g. UNS S32205), and as a result, has a tendency to precipitate the intermetallic phases. According to time–temperature-transformation (TTT) curve [18,19], the PREN 40 grade SDSS precipitates 1% r phase after aging for 100 s at 850 °C whereas the PREN 35 grade DSS precipitates 1% r phase after aging for 2000 s at 850 °C. That is, the precipitation rate of the r phase in the PREN 40 grade is much faster than that in the PREN 35 grade. If the furnace to quench time and quench rate is not sufficient, these brittle r phases in the center of large sections of the PREN 40 grade SDSS can be easily precipitated. In reality, these brittle r phases in the PREN 40 grade SDSS with heavier section have been observed to be forming during the cooling process after the continuous casting, during the slow cooling process of hot rolling, in the center of a casting, and in the heat affected zone after welding. The effect of addition of various elements on the precipitation of secondary phases has been investigated. Park et al. [20] assumed that the addition of rare earth metals suppress the precipitation
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Table 1 Chemical composition of the experimental alloys (wt.%).
⁄
Alloy designation
C
Cr
Ni
Mo
W
Si
Mn
Cu
S
N
Fe
PREN
WBASE
0.020
27.29
7.06
2.58
3.39
0.22
1.46
–
0.0037
0.33
Bal.
46.7
WCU
0.031
27.35
6.60
2.58
3.42
0.30
1.43
1.32
0.0040
0.33
Bal.
46.8
PREN = Cr + 3.3 (Mo + 0.5 W) + 16 N.
Fig. 1. Optical micrograph of the experimental alloys after solution heat-treated at 1090 °C for 30 min: (a) the alloy-WBASE and (b) the alloy-WCU.
Fig. 2. BSE images of the experimental alloys at 850 °C for 10 and 30 min: (a) the alloy-WBASE and (b) the alloy-WCU.
of the r and v phases in SDSS, because the REM (rare earth metals) atoms occupy the vacancies, effectively limiting the diffusion of Cr, Mo and W and REM oxides/oxy-sulfides which acts as the obstacles seemed to enhance the retardation effect. Huang et al. [21] showed that the higher the N content, the lower r phase content in DSSs. However, Si addition retards d phase decomposition into the c phase when SUS 309L stainless steels are solution treated at 1050 °C and also enhances the d-ferrite decomposition rate and accelerates r phase precipitation [22].
Ogawa et al. [23] showed that W in the HAZ of DSS is effective in delaying r phase precipitation due to the enhancement the tendency to form the v phase with a higher nucleation efficiency and lower growth rate than the r phase. Kim et al. [24] showed that a partial substitution of W for Mo in DSS retarded the formation of r phase in the alloys. Park et al. [25] reported that a partial substitution of W for Mo in ferritic stainless steels (FSSs) not only increased resistance to pitting corrosion of the alloys by a synergic contribution of Mo and W but also reduced the amount of
S.-H. Jeon et al. / Corrosion Science 66 (2013) 217–224 Table 2 Chemical composition of the intermetallic compounds formed in experimental alloys after aging at 850 °C (wt.%). Phase
Matrix Chi (v) Sigma (r)
WBASE
WCU
Fe
Cr
Mo
W
Fe
Cr
Mo
W
59.1 42.6 55.2
27.3 25.9 29.8
2.6 11.1 4.8
3.4 13.9 4.9
58.0 44.4 54.5
27.4 26.2 29.7
2.6 9.8 4.7
3.4 13.6 4.9
precipitation of r phase in the alloys. Smuk et al. [26] suggested that Cu addition to the alloy delays the precipitation of r phase during slow cooling due to fact that Cu particles pin moving r phase boundaries. However, the mechanism of the effects of Cu addition on the precipitation of intermetallic compounds and the associated intergranular corrosion are not well documented. In consideration of the workability and the productivity, it is certainly worthwhile for the large product which is difficult to avoid formation of intermetallic compounds during cooling if the retardation of intermetallic compounds formation can be
219
improved through optimizing alloy design by the addition of Cu and the Mo substitution by W to HDSS. In this study, to elucidate the effects of Cu addition on the precipitation of intermetallic compounds and intergranular corrosion of the HDSS, double loop electrochemical potentiokinetic reactivation (DL-EPR) test, a scanning electron microscope-energy dispersive X-ray spectroscope (SEM–EDS) analysis, an electron probe micro-analyzer (EPMA) and thermodynamic calculation using Thermo-Calc software were conducted.
2. Experimental procedures 2.1. Material and heat treatment Ingots weighing 50 kg with dimensions 150 by 150 by 300 mm (width by length by height) were manufactured using a high frequency vacuum induction furnace. After these ingots were hot rolled in the range of 1050–1250 °C, plates of 6 mm thickness were manufactured. The specimens were cut into dimensions of 15 by 15 by 6 mm (width by length by thickness) and solution
Fig. 3. Effects of Cu addition and aging time at 850 °C on the area fraction of intermetallic compounds precipitated in the experimental alloys: (a) r phase, (b) v phase and (c) total amount of r and v phases.
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heat-treated in air for 5 min per 1 mm thickness at 1090 °C and then quenched in water. The specimens were then isothermally aged at 850 °C for 10 and 30 min. Chemical analyses were performed by ARL 3460 optical emission spectrometer (OES). The nitrogen concentration was analyzed using a LECO N/O (TC-300) analyzer. Chemical compositions of the experimental alloys are presented in Table 1. 2.2. Double loop electrochemical potentiokinetic test The double loop electrochemical potentiokinetic tests are conducted following the recommendations of Majidi and Streicher [27]. DL-EPR test was originally used to evaluate the degree of sensitization of stainless steels [28]. DL-EPR test was applied to selectively attack the matrix around fine precipitates formed during aging, without an attack on the precipitates themselves. DL-EPR test was carried out using an EG&G PAR 263A potentiostat model. The standard test solution (0.5 M H2SO4 + 0.01 M KSCN) was used for ASSs. On the other hand, a more aggressive solution (2 M H2SO4 + 0.01 M KSCN + 0.5 M NaCl) was used for the more corrosion resistant DSSs [29]. DL-EPR test was conducted in 2 M H2SO4 + 0.01 M KSCN + 0.5 M NaCl at 30 °C. The degree of sensitization 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 0.06 V/min from 600 to 200 mVSCE, and then reversely to 600 mVSCE, respectively.
2.3. Micro-structural characterization To observe the microstructures of the alloys, they were ground to 2000 grit using SiC abrasive papers, polished surface using diamond paste. The sample was ultrasonic cleaned in acetone and distilled water to remove any impurities from the polished surface of the sample. The intermetallic compounds were observed using a JEOL JSM 700 SEM in a JEOL JSM 840A BSE. In addition, the chemical compositions of the intermetallic compounds were analyzed by a OXFORD instruments INCA X-art (51-ADD0069) EDS attached to a SEM. The line profiles of the Cr, Mo and W in the intermetallic compounds were measured using a SHIMADZU EPMA-1600 EPMA.
2.4. Thermodynamic equilibrium calculation To predict the effects of Cu on the activity of Cr, Mo and W in the alloy, thermodynamic calculation using Thermo-Calc software was conducted. The commercial thermodynamic computing software, Thermo-Calc was used to perform the thermodynamic calculation and it should be stressed that such calculation give the equilibrium state of system with the TCFE5-TCS Steels/Fe-alloys database [30]. The POLY and POST modules in Thermo-Calc software were used to perform the calculation. The POLY module can calculate various complex heterogeneous equilibrium states. And the POST module makes it possible to plot many kinds of phase diagrams and property diagrams [31].
Fig. 4. EPMA line profile of the intermetallic compounds in alloy-WBASE at 850 °C for 10 min.
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3.1. Microstructural analysis Fig. 1 shows the microstructures of the experimental alloys which were solution heat treated at 1090 °C for 30 min. The c phase can be found as island phase on the background of a phase which looks relatively dark. No intermetallic compounds were observed in either alloy. Besides, affected by hot rolling, the specimen has elongated texture parallel to the rolling direction. Fig. 2 shows the BSE images of experimental alloys aged for 10 and 30 min at 850 °C. After aging for 10 min at 850 °C, the precipitates formed continuously networks along grain boundaries and also randomly appeared within grains. These precipitates were identified as sigma and chi phases by SEM-EDS analyses. In presented in Table 2, the chemical composition of r and v phase is analyzed using SEM-EDS. A few v phases precipitated continuously along phase boundaries between c phases and a phases, and r phase appeared randomly within a grains in the both alloys according to the eutectoid reaction. After aging for 30 min at 850 °C, the total amount of intermetallic compounds at grain boundaries and within grains increased in the both alloys. Although the intermetallic compounds became further grown in the both alloys, the precipitation rate and type of intermetallic compounds in the base alloy is different to the Cu added alloy. Particularly, in the case of alloy-WBASE, numerous r phases formed widely within a grains and v phases precipitated continuously along phase boundaries between c phases and a phases. However, in the case of alloy-WCU, the r phases do not grow well at grain boundaries and within grains, but a number of v phases became further grown along phase boundaries between c phases and a phases and a phases as well as within a grains. Fig. 3 shows the effects of Cu addition and aging time at 850 °C on the area fraction of r phase, v phase and total intermetallic compounds in the experimental alloys by image analyzer. With an increase in aging time, the area fraction of r phase precipitated in the alloy-WBASE is much higher than that of alloy-WCU (Fig. 3(a)). However, the area fraction of v phase precipitated in the alloy-WCU is higher than that of alloy-WBASE (Fig. 3(b)). Fig. 3(c) shows that the area fraction of total intermetallic compounds (r + v) precipitated in alloy-WBASE is much higher than that of the alloy-WCU. In summary, as aging time at 850 °C increases, the addition of Cu to the base alloy reduces the total amount of intermetallic compounds. Particularly, the addition of Cu to the base alloy results in pronouncedly suppressing the amount of r phase whereas it slightly increases v phases. This result suggests that Cu addition to the base alloy effectively retards the precipitation of the intermetallic compounds.
reactivation current peak of both alloy-WBASE and alloy-WCU is more severely. The degree of Cr, Mo and W depletion with aging for alloyWBASE and alloy-WCU are represented by the ratio of reactivation peak current (ir) to activation peak current (ia) that were obtained from the DL-EPR tests as shown in Fig. 6. The ratio of ir/ia for alloyWBASE was much higher than that for alloy-WCU when equivalently aged, which demonstrated that the degree of Cr, Mo and W depletion in alloy-WBASE is greater than that in alloy-WCU. The degree of Cr, Mo and W depletion in alloy-WBASE is greater than that in alloy-WCU, which demonstrated that the amount of intermetallic compounds precipitation of the alloy-WBASE is more than that of the alloy-WCU. 3.3. Mechanism on the effect of Cu addition on the precipitation of intermetallic compounds Cr and Mo are the main elements which increase the susceptibility to r phase precipitation. The precipitation index (PI) is a formula for measure the tendency of the duplex stainless steels to r precipitation [2].
PI ¼ Cr þ 3:3 Mo
300 200
10 min.
0 -100 -200 -300
ia
ir
-400 -500 -600 -8 10
-7
10
-6
10
-5
10
-4
10
-3
10
-2
10
-1
10
2
Current Density(A /cm ) 300
Potential (mV vs. SCE)
Fig. 4 shows the BSE images and the EPMA line analyses of the intermetallic compounds of the alloy-WBASE. The EPMA line analyses of the intermetallic compounds show that the formation of r phases and v phases lead to Cr, Mo and W depleted areas at the interfaces between a phases and intermetallic phases. The degree of Cr, Mo and W depletion around intermetallic compounds formed during aging was examined by DL-EPR tests conducted in 2 M H2SO4 + 0.01 M KSCN + 0.5 M NaCl solution at 30 °C. Fig. 5 shows the results of DL-EPR test for the alloy-WBASE and alloy-WCU. A small reactivation current peak was appeared in the alloy-WCU, whereas a large reactivation current peak was obtained in the alloy-WBASE aged for 10 min at 850 °C as shown in Fig. 5(a). With increasing aging time, the difference of
(a)
WBASE WCU
100
200
3.2. Effect of Cu addition on the sensitization
ð1Þ
Alloy-WCU presents slightly higher concentrations of Cr than alloy-WBASE and the concentrations of Mo in the both alloys are same. By Eq. (1), the PI value of alloy-WCU is slightly higher than that of alloy-WBASE. Hence, it is not reasonable for the experimental results to suggest that the total amount of intermetallic
Potential (mV vs. SCE)
3. Results and discussion
(b)
WBASE WCU
30 min.
100 0 -100 -200 -300 -400 -500 -600 -8 10
-7
10
-6
10
-5
10
-4
10
-3
10
-2
10
-1
10
2
Current Density(A /cm ) Fig. 5. Effects of Cu addition and aging time at 850 °C on the DL-EPR behavior for the experimental alloys in 2 M H2SO4 + 0.01 M KSCN + 0.5 M NaCl solution at 30 °C: (a) aged for 10 min and (b) aged for 30 min.
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be a strong stabilizer for the v phase. Park et al. [20] reported that the v phase has higher affinity for W than for Mo and the great affinity of v phase for W can affect the reduction of the formation of r phase in the ferritic stainless steels. Cao et al. [32] also suggested that the v phase had a lower nucleation barrier than the r phase, but only stable r phase survived after long aging. As noted in the previous metallographic observations, the v phase precipitates earlier than the r phase, and v phase appears to be meta-stable and act as a transition phase in the FSSs. The effect of Cu addition on the activity of Cr, Mo and W can be calculated as a function of Cu content as shown in Fig. 7. As shown in Fig. 7, the activity of Cr and W linearly increase with an increase of Cu content. However, the activity of Mo linearly decreases with an increase of Cu content. Since W is a strong stabilizer for the v phase, the addition of Cu to the base alloy increases the precipitation rate of v phase with a higher nucleation efficiency and lower growth rate than the r phase due to the increase of activity of W. The preferential precipitation of v phase seems to be closely associated with the retardation of the precipitation r phase. Since the formation of r phase in the both alloys requires high concentrations of Mo and W, the preferential precipitation of v phase in the alloy during the initial period of aging can inhibit the nucleation and growth of r phase by depleting W and Mo adjacent to the v precipitates. Thus, the formation of r phase would be influenced significantly by the compositional variation of v phase. In the initial stage of v phase formation, r phase hardly forms, because the W or Mo required for r phase formation is insufficient around the v phase with its very high W or Mo content. The addition of Cu to the base alloy results in pronouncedly suppressing the amount of r phase due to the number of precipitated v phase. Although the activity of Cr and W linearly increase with an increase of Cu content, the activity of Mo linearly decreases with an
0.30 WBASE WCU
0.25
WBASE
Ir / Ia
0.20 0.15
WCU
0.10 0.05 0.00 -0.05 -5
0
5
10
15
20
25
30
35
o
Aging time at 850 C(minutes) Fig. 6. Effects of Cu addition and aging time at 850 °C on the ratio (ir/ia) of a reactivation current density to activation current density determined from the DL-EPR test of the experimental alloys.
compounds precipitation of alloy-WBASE is more than that of alloy-WCU. The effect of Cu addition was more determinant factor to precipitation of intermetallic compounds than the difference of PI value between alloy-WBASE and alloy-WCU. In presented in Table 2, the chemical composition of r and v phase is analyzed using SEM-EDS. The r phases contained slightly higher Cr, Mo and W compared to the matrix of the both alloys. The v phases contained the lower Cr and much higher Mo and W contents than the r phases. Especially, the v phase showed much higher W contents, 13.8 compared to that, 4.8% in the r phase while the difference of Mo contents between the r and v phase was relatively small. Accordingly, during aging, W is likely to segregate preferentially in v phase rather than in r phase. W should
0.0520
(a)
Activity of molybdenum
Activity of chromium
0.480 0.475 0.470 0.465 0.460 0.455
(b)
0.0515
0.0510
0.0505
0.0500
0.450 0.0
0.3
0.6
0.9
1.2
0.0
1.5
0.3
Cu (wt.%)
0.6
0.9
1.2
1.5
Cu (wt.%)
0.0610
Activity of tungsten
(c) 0.0605
0.0600
0.0595
0.0590
0.0585 0.0
0.3
0.6
0.9
1.2
1.5
Cu (wt.%) Fig. 7. Effects of Cu addition on the activity of the alloying elements in the alloy-WBASE as calculated using Thermo-Calc: (a) Cr, (b) Mo and (c) W.
S.-H. Jeon et al. / Corrosion Science 66 (2013) 217–224
223
Fig. 8. Schematic of the effect of Cu addition on the precipitation of intermetallic compounds of the experimental alloys.
increase of Cu content. Mo is also the main element stabilizing thermodynamically the r and v phases like Cr and W. Since Mo is one of the key elements in deleterious intermetallic compounds, the decrease of activity of Mo has an effect on the precipitation of intermetallic compounds. Since Mo is a strong stabilizer for the r phase, the addition of Cu to the base alloy results in pronouncedly suppressing the amount of r phase due to the decrease of activity of Mo. In summary, the effect of Cu addition on the precipitation of intermetallic compounds of the experimental alloys is schematically presented in Fig. 8. With an increase in aging time, the total amount of intermetallic compounds precipitated in the both alloys increases. Although the intermetallic compounds became further grown in the both alloys, the precipitation rate and type of intermetallic compounds in the base alloy is different to the Cu added alloy. Particularly, in the case of base alloy, numerous r phases formed widely within a grains and v phases precipitated continuously along phase boundaries between c phases and a phases. However, in the case of Cu added alloy, the r phases do not grow well at grain boundaries and within grains, but a number of v phases became further grown along phase boundaries between c phases and a phases and a phases as well as within a grains. The addition of Cu to the base alloy results in pronouncedly suppressing the amount of r phase whereas it slightly increases v phases. The addition of Cu to the base alloy reduces the total amount of intermetallic compounds. The Cu added alloy reduces the degree of sensitization due to the delayed precipitation of intermetallic compounds, compared with that of the base alloy.
4. Conclusions (1) The activity of Cr and W linearly increase with an increase of Cu content. However, the activity of Mo linearly decreases with an increase of Cu content. The addition of Cu to the base alloy increases the precipitation rate of v phase due to the increase of activity of W and retards the precipitation of r phase due to the decrease of activity of Mo. The addition of Cu to the base alloy results in pronouncedly suppressing the amount of r phase whereas it slightly increases v phases. Therefore, the addition of Cu to the base alloy reduces the total amount of intermetallic compounds. (2) The addition of Cu to the base alloy results in pronouncedly suppressing the amount of r phase due to the number of precipitated v phase. Since the formation of r phase in the both alloys requires high concentrations of Mo and W, the preferential precipitation of v phase in the alloys during the initial period of aging can inhibit the nucleation and growth of r phase by depleting W and Mo adjacent to the v precipitates. (3) DL-EPR test was originally used to evaluate the degree of sensitization of stainless steels. The degree of sensitization with aging for experimental alloys is represented by the ratio of ir to ia The ratio of ir/ia for alloy-WBASE was much higher than that for alloy-WCU when equivalently aged. These results indicated that the addition of Cu to the base alloy reduces the degree of sensitization due to the delayed precipitation of intermetallic compounds.
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Effects of Cu on the precipitation of intermetallic compounds and the intergranular corrosion of hyper duplex stainless steels Soon-Hyeok Jeon a, Soon-Tae Kim a, In-Sung Lee a, Ji-Soo Kim b, Kwang-Tae Kim b, Yong-Soo Park a,⇑ a b
Department of Material Science and Engineering, Yonsei University, 134 Shinchon-dong, Seodaemun-gu, Seoul 120-749, Republic of Korea Stainless Steel Research Group, POSCO Technical Research Laboratories Goedong-dong, Nam-Gu, Pohang, Gyeongbuk 790-785, Republic of Korea
a r t i c l e
i n f o
Article history: Received 27 June 2012 Accepted 18 September 2012 Available online 27 September 2012 Keywords: A. Copper A. Stainless B. Polarization B. SEM C. Intergranular corrosion
a b s t r a c t To elucidate the effects of Cu on the precipitation of intermetallic compounds and the intergranular corrosion of the hyper duplex stainless steels, a double loop potentiokinetic reactivation test, a scanning electron microscope analysis and thermodynamic calculation were conducted. The addition of Cu to the base alloy reduces the total amount of intermetallic compounds. Particularly, Cu addition to the base alloy results in pronouncedly suppressing the amount of sigma phase whereas it slightly increases chi phases. The Cu added alloy reduces the degree of sensitization due to the delayed precipitation of intermetallic compounds, compared with that of the base alloy. Ó 2012 Elsevier Ltd. All rights reserved.
1. Introduction Duplex stainless steel (DSS) is the stainless steel that has microstructure where both ferrite (a) and austenite (c) phases are present in approximately equal volume fraction. In general, it is well known that standard stainless steel such as UNS S32205 is defined as DSS with a pitting resistance equivalent number (PREN = Cr + 3.3 (Mo + 0.5 W) + 16 N [1,2]) of 35 and super duplex stainless steels (SDSSs) such as UNS S32750, UNS S32760 and UNS S32550 are defined as DSSs with PREN of 40–45. SDSSs have been increasingly used for various applications such as power plants, desalination facilities and chemical plants due to high resistance to pitting and crevice corrosion, excellent mechanical properties and a relatively low cost, compared with other higher performance materials such as super austenite stainless steel (SASS) [3–5]. However, in heat exchanger application, corrosion resistance of SDSS is insufficient for higher temperature service or for a long service life, and where materials, with even higher corrosion resistance, are needed. Hyper duplex stainless steel (HDSS) such as UNS S32707 is defined as a highly alloyed duplex stainless steel with a PREN in excess of 45 [1]. HDSSs were developed to meet industry demands for higher operating temperatures and longer run times and replace the costly SASS that has high levels of Ni and Mo contents [6,7]. The existing researches were progressed for DSS and SDSS [8–12]. However, few studies have been focused
⇑ Corresponding author. Tel.: +82 2 2123 5833; fax: +82 2 362 9199. E-mail address: [email protected] (Y.-S. Park). 0010-938X/$ - see front matter Ó 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.corsci.2012.09.023
on HDSS, having higher Cr, Mo and W than DSS and SDSS, to substitute the 6% Mo austenite stainless steels (ASSs) in more severe environment. The intermetallic phases such as sigma (r) phase and chi (v) phase are the intermetallic phases that forms at high temperature 600–950 °C [13,14]. These intermetallic phases lower the fracture toughness and the corrosion resistance significantly [15–17]. To avoid this embrittlement and corrosion, DSSs are subjected to solution heat treatment followed by water quenching in the final stage of production or fabrication. The PREN 40 grade SDSS (e.g. UNS S32750) contains larger amounts of Cr and Mo than the PREN 35 grade DSS (e.g. UNS S32205), and as a result, has a tendency to precipitate the intermetallic phases. According to time–temperature-transformation (TTT) curve [18,19], the PREN 40 grade SDSS precipitates 1% r phase after aging for 100 s at 850 °C whereas the PREN 35 grade DSS precipitates 1% r phase after aging for 2000 s at 850 °C. That is, the precipitation rate of the r phase in the PREN 40 grade is much faster than that in the PREN 35 grade. If the furnace to quench time and quench rate is not sufficient, these brittle r phases in the center of large sections of the PREN 40 grade SDSS can be easily precipitated. In reality, these brittle r phases in the PREN 40 grade SDSS with heavier section have been observed to be forming during the cooling process after the continuous casting, during the slow cooling process of hot rolling, in the center of a casting, and in the heat affected zone after welding. The effect of addition of various elements on the precipitation of secondary phases has been investigated. Park et al. [20] assumed that the addition of rare earth metals suppress the precipitation
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Table 1 Chemical composition of the experimental alloys (wt.%).
⁄
Alloy designation
C
Cr
Ni
Mo
W
Si
Mn
Cu
S
N
Fe
PREN
WBASE
0.020
27.29
7.06
2.58
3.39
0.22
1.46
–
0.0037
0.33
Bal.
46.7
WCU
0.031
27.35
6.60
2.58
3.42
0.30
1.43
1.32
0.0040
0.33
Bal.
46.8
PREN = Cr + 3.3 (Mo + 0.5 W) + 16 N.
Fig. 1. Optical micrograph of the experimental alloys after solution heat-treated at 1090 °C for 30 min: (a) the alloy-WBASE and (b) the alloy-WCU.
Fig. 2. BSE images of the experimental alloys at 850 °C for 10 and 30 min: (a) the alloy-WBASE and (b) the alloy-WCU.
of the r and v phases in SDSS, because the REM (rare earth metals) atoms occupy the vacancies, effectively limiting the diffusion of Cr, Mo and W and REM oxides/oxy-sulfides which acts as the obstacles seemed to enhance the retardation effect. Huang et al. [21] showed that the higher the N content, the lower r phase content in DSSs. However, Si addition retards d phase decomposition into the c phase when SUS 309L stainless steels are solution treated at 1050 °C and also enhances the d-ferrite decomposition rate and accelerates r phase precipitation [22].
Ogawa et al. [23] showed that W in the HAZ of DSS is effective in delaying r phase precipitation due to the enhancement the tendency to form the v phase with a higher nucleation efficiency and lower growth rate than the r phase. Kim et al. [24] showed that a partial substitution of W for Mo in DSS retarded the formation of r phase in the alloys. Park et al. [25] reported that a partial substitution of W for Mo in ferritic stainless steels (FSSs) not only increased resistance to pitting corrosion of the alloys by a synergic contribution of Mo and W but also reduced the amount of
S.-H. Jeon et al. / Corrosion Science 66 (2013) 217–224 Table 2 Chemical composition of the intermetallic compounds formed in experimental alloys after aging at 850 °C (wt.%). Phase
Matrix Chi (v) Sigma (r)
WBASE
WCU
Fe
Cr
Mo
W
Fe
Cr
Mo
W
59.1 42.6 55.2
27.3 25.9 29.8
2.6 11.1 4.8
3.4 13.9 4.9
58.0 44.4 54.5
27.4 26.2 29.7
2.6 9.8 4.7
3.4 13.6 4.9
precipitation of r phase in the alloys. Smuk et al. [26] suggested that Cu addition to the alloy delays the precipitation of r phase during slow cooling due to fact that Cu particles pin moving r phase boundaries. However, the mechanism of the effects of Cu addition on the precipitation of intermetallic compounds and the associated intergranular corrosion are not well documented. In consideration of the workability and the productivity, it is certainly worthwhile for the large product which is difficult to avoid formation of intermetallic compounds during cooling if the retardation of intermetallic compounds formation can be
219
improved through optimizing alloy design by the addition of Cu and the Mo substitution by W to HDSS. In this study, to elucidate the effects of Cu addition on the precipitation of intermetallic compounds and intergranular corrosion of the HDSS, double loop electrochemical potentiokinetic reactivation (DL-EPR) test, a scanning electron microscope-energy dispersive X-ray spectroscope (SEM–EDS) analysis, an electron probe micro-analyzer (EPMA) and thermodynamic calculation using Thermo-Calc software were conducted.
2. Experimental procedures 2.1. Material and heat treatment Ingots weighing 50 kg with dimensions 150 by 150 by 300 mm (width by length by height) were manufactured using a high frequency vacuum induction furnace. After these ingots were hot rolled in the range of 1050–1250 °C, plates of 6 mm thickness were manufactured. The specimens were cut into dimensions of 15 by 15 by 6 mm (width by length by thickness) and solution
Fig. 3. Effects of Cu addition and aging time at 850 °C on the area fraction of intermetallic compounds precipitated in the experimental alloys: (a) r phase, (b) v phase and (c) total amount of r and v phases.
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heat-treated in air for 5 min per 1 mm thickness at 1090 °C and then quenched in water. The specimens were then isothermally aged at 850 °C for 10 and 30 min. Chemical analyses were performed by ARL 3460 optical emission spectrometer (OES). The nitrogen concentration was analyzed using a LECO N/O (TC-300) analyzer. Chemical compositions of the experimental alloys are presented in Table 1. 2.2. Double loop electrochemical potentiokinetic test The double loop electrochemical potentiokinetic tests are conducted following the recommendations of Majidi and Streicher [27]. DL-EPR test was originally used to evaluate the degree of sensitization of stainless steels [28]. DL-EPR test was applied to selectively attack the matrix around fine precipitates formed during aging, without an attack on the precipitates themselves. DL-EPR test was carried out using an EG&G PAR 263A potentiostat model. The standard test solution (0.5 M H2SO4 + 0.01 M KSCN) was used for ASSs. On the other hand, a more aggressive solution (2 M H2SO4 + 0.01 M KSCN + 0.5 M NaCl) was used for the more corrosion resistant DSSs [29]. DL-EPR test was conducted in 2 M H2SO4 + 0.01 M KSCN + 0.5 M NaCl at 30 °C. The degree of sensitization 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 0.06 V/min from 600 to 200 mVSCE, and then reversely to 600 mVSCE, respectively.
2.3. Micro-structural characterization To observe the microstructures of the alloys, they were ground to 2000 grit using SiC abrasive papers, polished surface using diamond paste. The sample was ultrasonic cleaned in acetone and distilled water to remove any impurities from the polished surface of the sample. The intermetallic compounds were observed using a JEOL JSM 700 SEM in a JEOL JSM 840A BSE. In addition, the chemical compositions of the intermetallic compounds were analyzed by a OXFORD instruments INCA X-art (51-ADD0069) EDS attached to a SEM. The line profiles of the Cr, Mo and W in the intermetallic compounds were measured using a SHIMADZU EPMA-1600 EPMA.
2.4. Thermodynamic equilibrium calculation To predict the effects of Cu on the activity of Cr, Mo and W in the alloy, thermodynamic calculation using Thermo-Calc software was conducted. The commercial thermodynamic computing software, Thermo-Calc was used to perform the thermodynamic calculation and it should be stressed that such calculation give the equilibrium state of system with the TCFE5-TCS Steels/Fe-alloys database [30]. The POLY and POST modules in Thermo-Calc software were used to perform the calculation. The POLY module can calculate various complex heterogeneous equilibrium states. And the POST module makes it possible to plot many kinds of phase diagrams and property diagrams [31].
Fig. 4. EPMA line profile of the intermetallic compounds in alloy-WBASE at 850 °C for 10 min.
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3.1. Microstructural analysis Fig. 1 shows the microstructures of the experimental alloys which were solution heat treated at 1090 °C for 30 min. The c phase can be found as island phase on the background of a phase which looks relatively dark. No intermetallic compounds were observed in either alloy. Besides, affected by hot rolling, the specimen has elongated texture parallel to the rolling direction. Fig. 2 shows the BSE images of experimental alloys aged for 10 and 30 min at 850 °C. After aging for 10 min at 850 °C, the precipitates formed continuously networks along grain boundaries and also randomly appeared within grains. These precipitates were identified as sigma and chi phases by SEM-EDS analyses. In presented in Table 2, the chemical composition of r and v phase is analyzed using SEM-EDS. A few v phases precipitated continuously along phase boundaries between c phases and a phases, and r phase appeared randomly within a grains in the both alloys according to the eutectoid reaction. After aging for 30 min at 850 °C, the total amount of intermetallic compounds at grain boundaries and within grains increased in the both alloys. Although the intermetallic compounds became further grown in the both alloys, the precipitation rate and type of intermetallic compounds in the base alloy is different to the Cu added alloy. Particularly, in the case of alloy-WBASE, numerous r phases formed widely within a grains and v phases precipitated continuously along phase boundaries between c phases and a phases. However, in the case of alloy-WCU, the r phases do not grow well at grain boundaries and within grains, but a number of v phases became further grown along phase boundaries between c phases and a phases and a phases as well as within a grains. Fig. 3 shows the effects of Cu addition and aging time at 850 °C on the area fraction of r phase, v phase and total intermetallic compounds in the experimental alloys by image analyzer. With an increase in aging time, the area fraction of r phase precipitated in the alloy-WBASE is much higher than that of alloy-WCU (Fig. 3(a)). However, the area fraction of v phase precipitated in the alloy-WCU is higher than that of alloy-WBASE (Fig. 3(b)). Fig. 3(c) shows that the area fraction of total intermetallic compounds (r + v) precipitated in alloy-WBASE is much higher than that of the alloy-WCU. In summary, as aging time at 850 °C increases, the addition of Cu to the base alloy reduces the total amount of intermetallic compounds. Particularly, the addition of Cu to the base alloy results in pronouncedly suppressing the amount of r phase whereas it slightly increases v phases. This result suggests that Cu addition to the base alloy effectively retards the precipitation of the intermetallic compounds.
reactivation current peak of both alloy-WBASE and alloy-WCU is more severely. The degree of Cr, Mo and W depletion with aging for alloyWBASE and alloy-WCU are represented by the ratio of reactivation peak current (ir) to activation peak current (ia) that were obtained from the DL-EPR tests as shown in Fig. 6. The ratio of ir/ia for alloyWBASE was much higher than that for alloy-WCU when equivalently aged, which demonstrated that the degree of Cr, Mo and W depletion in alloy-WBASE is greater than that in alloy-WCU. The degree of Cr, Mo and W depletion in alloy-WBASE is greater than that in alloy-WCU, which demonstrated that the amount of intermetallic compounds precipitation of the alloy-WBASE is more than that of the alloy-WCU. 3.3. Mechanism on the effect of Cu addition on the precipitation of intermetallic compounds Cr and Mo are the main elements which increase the susceptibility to r phase precipitation. The precipitation index (PI) is a formula for measure the tendency of the duplex stainless steels to r precipitation [2].
PI ¼ Cr þ 3:3 Mo
300 200
10 min.
0 -100 -200 -300
ia
ir
-400 -500 -600 -8 10
-7
10
-6
10
-5
10
-4
10
-3
10
-2
10
-1
10
2
Current Density(A /cm ) 300
Potential (mV vs. SCE)
Fig. 4 shows the BSE images and the EPMA line analyses of the intermetallic compounds of the alloy-WBASE. The EPMA line analyses of the intermetallic compounds show that the formation of r phases and v phases lead to Cr, Mo and W depleted areas at the interfaces between a phases and intermetallic phases. The degree of Cr, Mo and W depletion around intermetallic compounds formed during aging was examined by DL-EPR tests conducted in 2 M H2SO4 + 0.01 M KSCN + 0.5 M NaCl solution at 30 °C. Fig. 5 shows the results of DL-EPR test for the alloy-WBASE and alloy-WCU. A small reactivation current peak was appeared in the alloy-WCU, whereas a large reactivation current peak was obtained in the alloy-WBASE aged for 10 min at 850 °C as shown in Fig. 5(a). With increasing aging time, the difference of
(a)
WBASE WCU
100
200
3.2. Effect of Cu addition on the sensitization
ð1Þ
Alloy-WCU presents slightly higher concentrations of Cr than alloy-WBASE and the concentrations of Mo in the both alloys are same. By Eq. (1), the PI value of alloy-WCU is slightly higher than that of alloy-WBASE. Hence, it is not reasonable for the experimental results to suggest that the total amount of intermetallic
Potential (mV vs. SCE)
3. Results and discussion
(b)
WBASE WCU
30 min.
100 0 -100 -200 -300 -400 -500 -600 -8 10
-7
10
-6
10
-5
10
-4
10
-3
10
-2
10
-1
10
2
Current Density(A /cm ) Fig. 5. Effects of Cu addition and aging time at 850 °C on the DL-EPR behavior for the experimental alloys in 2 M H2SO4 + 0.01 M KSCN + 0.5 M NaCl solution at 30 °C: (a) aged for 10 min and (b) aged for 30 min.
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be a strong stabilizer for the v phase. Park et al. [20] reported that the v phase has higher affinity for W than for Mo and the great affinity of v phase for W can affect the reduction of the formation of r phase in the ferritic stainless steels. Cao et al. [32] also suggested that the v phase had a lower nucleation barrier than the r phase, but only stable r phase survived after long aging. As noted in the previous metallographic observations, the v phase precipitates earlier than the r phase, and v phase appears to be meta-stable and act as a transition phase in the FSSs. The effect of Cu addition on the activity of Cr, Mo and W can be calculated as a function of Cu content as shown in Fig. 7. As shown in Fig. 7, the activity of Cr and W linearly increase with an increase of Cu content. However, the activity of Mo linearly decreases with an increase of Cu content. Since W is a strong stabilizer for the v phase, the addition of Cu to the base alloy increases the precipitation rate of v phase with a higher nucleation efficiency and lower growth rate than the r phase due to the increase of activity of W. The preferential precipitation of v phase seems to be closely associated with the retardation of the precipitation r phase. Since the formation of r phase in the both alloys requires high concentrations of Mo and W, the preferential precipitation of v phase in the alloy during the initial period of aging can inhibit the nucleation and growth of r phase by depleting W and Mo adjacent to the v precipitates. Thus, the formation of r phase would be influenced significantly by the compositional variation of v phase. In the initial stage of v phase formation, r phase hardly forms, because the W or Mo required for r phase formation is insufficient around the v phase with its very high W or Mo content. The addition of Cu to the base alloy results in pronouncedly suppressing the amount of r phase due to the number of precipitated v phase. Although the activity of Cr and W linearly increase with an increase of Cu content, the activity of Mo linearly decreases with an
0.30 WBASE WCU
0.25
WBASE
Ir / Ia
0.20 0.15
WCU
0.10 0.05 0.00 -0.05 -5
0
5
10
15
20
25
30
35
o
Aging time at 850 C(minutes) Fig. 6. Effects of Cu addition and aging time at 850 °C on the ratio (ir/ia) of a reactivation current density to activation current density determined from the DL-EPR test of the experimental alloys.
compounds precipitation of alloy-WBASE is more than that of alloy-WCU. The effect of Cu addition was more determinant factor to precipitation of intermetallic compounds than the difference of PI value between alloy-WBASE and alloy-WCU. In presented in Table 2, the chemical composition of r and v phase is analyzed using SEM-EDS. The r phases contained slightly higher Cr, Mo and W compared to the matrix of the both alloys. The v phases contained the lower Cr and much higher Mo and W contents than the r phases. Especially, the v phase showed much higher W contents, 13.8 compared to that, 4.8% in the r phase while the difference of Mo contents between the r and v phase was relatively small. Accordingly, during aging, W is likely to segregate preferentially in v phase rather than in r phase. W should
0.0520
(a)
Activity of molybdenum
Activity of chromium
0.480 0.475 0.470 0.465 0.460 0.455
(b)
0.0515
0.0510
0.0505
0.0500
0.450 0.0
0.3
0.6
0.9
1.2
0.0
1.5
0.3
Cu (wt.%)
0.6
0.9
1.2
1.5
Cu (wt.%)
0.0610
Activity of tungsten
(c) 0.0605
0.0600
0.0595
0.0590
0.0585 0.0
0.3
0.6
0.9
1.2
1.5
Cu (wt.%) Fig. 7. Effects of Cu addition on the activity of the alloying elements in the alloy-WBASE as calculated using Thermo-Calc: (a) Cr, (b) Mo and (c) W.
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Fig. 8. Schematic of the effect of Cu addition on the precipitation of intermetallic compounds of the experimental alloys.
increase of Cu content. Mo is also the main element stabilizing thermodynamically the r and v phases like Cr and W. Since Mo is one of the key elements in deleterious intermetallic compounds, the decrease of activity of Mo has an effect on the precipitation of intermetallic compounds. Since Mo is a strong stabilizer for the r phase, the addition of Cu to the base alloy results in pronouncedly suppressing the amount of r phase due to the decrease of activity of Mo. In summary, the effect of Cu addition on the precipitation of intermetallic compounds of the experimental alloys is schematically presented in Fig. 8. With an increase in aging time, the total amount of intermetallic compounds precipitated in the both alloys increases. Although the intermetallic compounds became further grown in the both alloys, the precipitation rate and type of intermetallic compounds in the base alloy is different to the Cu added alloy. Particularly, in the case of base alloy, numerous r phases formed widely within a grains and v phases precipitated continuously along phase boundaries between c phases and a phases. However, in the case of Cu added alloy, the r phases do not grow well at grain boundaries and within grains, but a number of v phases became further grown along phase boundaries between c phases and a phases and a phases as well as within a grains. The addition of Cu to the base alloy results in pronouncedly suppressing the amount of r phase whereas it slightly increases v phases. The addition of Cu to the base alloy reduces the total amount of intermetallic compounds. The Cu added alloy reduces the degree of sensitization due to the delayed precipitation of intermetallic compounds, compared with that of the base alloy.
4. Conclusions (1) The activity of Cr and W linearly increase with an increase of Cu content. However, the activity of Mo linearly decreases with an increase of Cu content. The addition of Cu to the base alloy increases the precipitation rate of v phase due to the increase of activity of W and retards the precipitation of r phase due to the decrease of activity of Mo. The addition of Cu to the base alloy results in pronouncedly suppressing the amount of r phase whereas it slightly increases v phases. Therefore, the addition of Cu to the base alloy reduces the total amount of intermetallic compounds. (2) The addition of Cu to the base alloy results in pronouncedly suppressing the amount of r phase due to the number of precipitated v phase. Since the formation of r phase in the both alloys requires high concentrations of Mo and W, the preferential precipitation of v phase in the alloys during the initial period of aging can inhibit the nucleation and growth of r phase by depleting W and Mo adjacent to the v precipitates. (3) DL-EPR test was originally used to evaluate the degree of sensitization of stainless steels. The degree of sensitization with aging for experimental alloys is represented by the ratio of ir to ia The ratio of ir/ia for alloy-WBASE was much higher than that for alloy-WCU when equivalently aged. These results indicated that the addition of Cu to the base alloy reduces the degree of sensitization due to the delayed precipitation of intermetallic compounds.
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