Intergranular corrosion of high purity austenitic stainless steel containing silicon additions

Intergranular corrosion of high purity austenitic stainless steel containing silicon additions

Materials Science and Engineering, 51 (1981) 165 - 174 165 Intergranular Corrosion of High Purity Austenitic Stainless Steel Containing Silicon Addi...

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Materials Science and Engineering, 51 (1981) 165 - 174

165

Intergranular Corrosion of High Purity Austenitic Stainless Steel Containing Silicon Additions A. R. PERRIN and K. T. AUST

Department of Metallurgy and Materials Science, University of Toronto, Toronto, Ontario (Canada) (Received April 8, 1981; in revised form May 15, 1981)

SUMMARY

The intergranular corrosion o f high purity austenitic stainless steels was studied. Steels containing less than 500 ppm total impurities were doped with various amounts of silicon and solution treated before testing. Potentiodynamic anodic polarization curves and fixed potential measurements were used with optical and scanning electron microscopy to study the corrosion morphology in these alloys. Between 1.1 and 1.5 V with respect to a saturated calomel electrode (SCE) intergranular corrosion took place in the silicon-doped steels. Silicon concentrations in the range 100- 4800 ppm caused susceptibility to both intergranular and surface attack. The high purity base alloy (50 ppm) and high silicon content (4.15 wt.%) steels were immune to intergranular and surface attack. The transpassive current density (between 1.1 and 1.5 V (SCE)) was found to increase with increasing amounts of silicon.

1. INTRODUCTION

The intergranular corrosion of non-sensitized austenitic stainless steels was initially observed by DeLong [1] and later by Shirley and Truman [2] when these steels were exposed to unrefreshed Huey test solutions. Nonsensitized or solution-treated steels have no detectable carbide precipitation at the grain boundaries; thus sensitization or chromium depletion theories do not apply. Various corrosion tests have been used to show intergranular attack. The most popular has been the nitric acid-potassium dichromate solution [3, 4]. Austenitic stainless steels treated in this solution have been found to be susceptible to intergranular attack. It has been 0025-5416/81/0000-0000/$02.50

found that oxidizing ions such as Cr6+ depolarize the cathodic reactions and therefore raise the open-circuit potential of stainless steels immersed in nitric acids [ 5, 6]. Thus the anodic reaction rate is increased and the grain boundaries selectively dissolve. Schuller et al. [ 7] considered the effect of the polarization potential on the metallography of corrosion in sensitized and nonsensitized stainless steels. At very high potentials greater than 1150 mV with respect to a saturated calomel electrode (SCE), severe grain boundary grooving (funnel attack) was observed in the solution-treated steels but not in the sensitized steel. Osozawa et al. [8] found that intergranular corrosion took place in the transpassive region of anodic polarization curves (above 1200 mV (SCE)). Osozawa et aI. were studying sensitized stainless steels and found that the results could not be explained by using chromium depletion models where intergranular attack should not have occurred in the transpassive region. Several theories have been proposed to explain the localized anodic activity at the grain boundaries in these steels. Streicher [ 5] and Coriou et al. [6] suggested that the driving force for accelerated intergranular corrosion was provided by the strain energy associated with grain boundaries. However, this suggestion does not explain why high purity alloys are immune to attack [3, 9]. A solute segregation model has been proposed [ 10, 11] to account for the behaviour of these steels. According to this theory, the driving force for intergranular attack originates from the chemical composition differences between the grain boundaries and the matrix. The localized corrosion is associated with the presence of continuous grain boundary paths of solute-segregated regions. This theory accounts for the low corrosion rates of high © Elsevier Sequoia/Printed in The Netherlands

166 purity materials since these segregated regions would n o t be present. The effect of various elements on intergranular corrosion has been studied by Armijo [3]. It was shown that solution-treated high purity austenitic stainless steels d o p e d with phosphorus or silicon are susceptible to intergranular attack in the nitric acid-potassium dichromate test. Other elements such as carbon, nitrogen, oxygen, manganese and sulphur did n o t appear to p r o m o t e susceptibility. D i r e c t evidence relating intergranular corrosion to impurity segregation has been obtained by Joshi and Stein [12] from Auger spectroscopic analysis of austenitic stainless steel. Their results indicate a correlation between the grain boundary segregation of solute impurities, such as sulphur, phosphorus and silicon, and intergranular corrosion of t y p e 304 stainless steels in a strongly oxidizing solution. Auger electron spectroscopy represents a powerful tool; however, the cleavage m e t h o d of revealing grain boundary planes for study by Auger spectroscopy may tend to select only grain boundaries with a particular low cohesion, i.e. maximum segregation. Because of the ductility of austenitic stainless steels, particularly in the solution-treated condition, it is difficult to expose grain boundary surfaces for Auger analysis. It has been proposed that transpassive current densities should be sensitive to trace impurities at the grain boundaries [ 1 2 ] . If most of the corrosion attack takes place at the grain boundaries, then an impurity-free solution-treated austenitic stainless steel should have lower transpassive current densities than a steel that has impurities at the grain boundaries. The present work was carried o u t to determine whether anodic polarization curves may be a useful m e t h o d for studying segrega-

tion and intergranular corrosion in these solution-treated austenitic stainless steels. The importance of anodic polarization experiments in studies of the intergranular corrosion and segregation of sulphur in nickel [13] and carbon in ferritic F e - C r alloys [14] has recently been reported.

2. EXPERIMENTAL PROCEDURE Potentiodynamic polarization studies were carried o u t using alloys of high purity F e - C r - N i austenitic stainless steel (less than 500 ppm total impurities) d o p e d with controlled amounts of silicon. These alloys were prepared by the General Electric Company as described previously [ 3 ] . The alloy compositions based on analyses from the General Electric Vallecitos Atomic Laboratory are shown in Table 1. The samples were in the form of thin sheet, 2 cm long, 1 cm wide and 1 mm thick. The specimens were given a solution treatm e n t at 1250 °C for 24 h in vacuum and were then water quenched. All specimens were completely austenitic except for specimen 6 (Table 1) which was 50% ferritic. After heat treatment and before corrosion testing, the samples were electropolished to obtain a mirror-like finish in a solution of 78 ml perchloric acid (70%), 120 ml distilled water, 700 ml ethanol and 100 ml 2-butoxy ethanol at - - 2 0 °C. The equipment used to obtain polarization curves is depicted in Fig. 1. A Princeton Applied Research model 173 p o t e n t i o s t a t galvanostat and model 376 logarithmic current converter were used. The chart recorder was a Hewlett Packard 7044A x - y recorder. The electrode design used in this study to expose the thin sheet specimens to the corrosion solution is shown in Fig. 2. This design

TABLE 1 Chemical compositions of the high purity stainless steels with various silicon concentrations (balance, iron)

Specimen Cr (wt.%) Ni (wt.%) Si (wt.%)

C (wt.%)

Mn (wt.%) S (wt.%) P (wt.%) N (wt.%)

0 (wt.%)

1 2 3 4 5 6

0.0038 0.0046 0.0035 0.0031 0.0034 0.0044

0.0005 0.0005 0.0005 0.0005 0.0020 0.0005

0.0172 0.0179 0.0089 0.0059 0.0075 0.0045

15.5 15.0 15.2 15.1 16.0 15.2

13.1 13.5 13.6 13.4 12.6 13.5

0.001 0.01 0.18 0.28 0.48 4.15

0.008 0.0070 0.0070 0.0060 0.0020 0.0050

0.0015 0.0014 0.0013 0.0009 0.0017 0.0013

0.0011 0.0008 0.0008 0.0008 0.0016 0.0014

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reagent grade chemicals and doubly distilled water. A constant-potential scan rate of 0.6 V h -1 was applied during polarization and an atmosphere of high purity hydrogen was present in the polarization test cell. Fixedpotential polarization tests were also conducted in the transpassive region. Optical microscopy and scanning electron microscopy were used to study the corrosion morphology of the alloys after a full potentiodynamic polarization test up to 1.75 V (SCE) and after fixed-potential polarization.

3. R E S U L T S

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3.1. Anodic polarization studies Anodic polarization curves for the high purity steels are presented in Figs. 3 and 4. The presence of various amounts of silicon has no effect on the active and passive parts of the polarization curves. However, silicon had a considerable effect on the transpassive current density; the current density in this region increased with increasing silicon content. Figure 4 is a replot of the transpassive region of Fig. 3. The purer steels had the lowest current densities. The current density was found to increase with increasing silicon content except for the 2800 ppm Si specimen (its curve is n o t shown in Fig. 3) which behaved similarly to the 1800 ppm Si specimen. Above 1.6 V (SCE) the current density dependence on silicon content was reversed; greater current densities were obtained from 2.00 ........

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permitted easy removal of specimens, a constant exposed area and good electrical contact and did n o t allow any crevice corrosion. A detailed description of the polarization test cell, the constant-potential apparatus and the procedure for electrochemical measurements is given in ref. 15. The polarization tests were c o n d u c t e d in freshly prepared 1 N H2SO 4 at 30 °C using

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Fig. 5. Current density vs. time at a fixed potential (1.375 V (SCE)) for high purity silicon-doped austenitic stainless steels.

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the purer steels. Most of the attack above 1.6 V (SCE) was surface attack. The transpassive current densities of the high purity base alloy were slightly higher than those for the 100 p p m Si specimen below 1.6 V (SCE). After heat treatment a very thick oxide was present in the high purity base alloy (specimen 1). After the oxide had been removed, many surface defects were present which could n o t be removed by electropolishing. The higher transpassive current density of the base alloy (no silicon addition) was probably caused by surface damage that occurred during heat treatment. In order to determine whether the results shown in Figs. 3 and 4 obtained at a potential sweep rate of 0.6 V (SCE) h -z were equilibrium values, fixed-potential studies were performed. The specimens were held at 1.375 V (SCE) for 3 h and the current was measured. The results shown in Fig. 5 indicate that the current density increased rapidly for about 30 min and then almost completely levelled off. Again, higher silicon contents gave higher current densities. A comparison of current densities at 1.375 V (SCE) is shown in Fig. 6. The fixed potential current densities (the equilibrated values after 3 h) are slightly higher than the potentiodynamic values at this voltage; however, only the current density of the 1800 p p m Si specimen showed a significant difference from the potentiodynamic value obtained.

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3.2. Microscopic studies Photomicrographs of the stainless steels after a complete potentiodynamic polarization test up to 1.75 V (SCE) are shown in Fig. 7. Pitting is present but very little intergranular attack took place in specimen 1, the base alloy {Fig. 7(a)). The 100 ppm Si specimen {specimen 2) experienced a large a m o u n t of intergranular attack and some pitting, as shown in Fig. 7(b); however, continuous grain boundary grooves were n o t present. Twin boundaries in the high purity austenitic stainless steels were also attacked slightly, as was observed previously [3]. Figure 7(c) shows intergranular attack in the 4800 p p m Si specimen {specimen 5); the grain boundary grooves were continuous in this case. Considerable surface attack also occurred, as can

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be seen by the surface relief effect in Fig. 7(c). The steel containing 4.15 wt.% Si (specimen 6 ) experienced no intergranular attack and very little surface attack (Fig. 7(d)). The transpassive current densities of the stainless steels may be related to the metallography observed. Figures 3 - 5 show that the current density or corrosion rate increases with increasing amounts of silicon• Figure 7 indicates that up to 4800 ppm Si the metallography is consistent with the observed current densities. The intergranular and surface attack increase with increasing amounts of silicon. The greatest amount of intergranular attack took place in the 4800 ppm Si specimen.

However, the 4.15 wt.% Si specimen had very high transpassive current densities (e.g. Fig. 4) but yet did not suffer very much corrosion attack (Fig. 7(d)). The corroded microstructures observed during the fixed-potential studies of Fig. 5 are shown in Figs. 8 and 9. Intergranular corrosion is present in all the specimens shown here. The 100 ppm Si specimen suffered little surface attack and considerable intergranular attack (Figs. 8(a) and 9(a)). The same type of attack is present in the 1800 ppm Si specimen (Figs. 8(b) and 9(b)) except that some surface relief is also present. A large amount of intergranular attack and surface relief is

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evident in the 4800 ppm Si specimen (Figs. 8(c) and 9(c)). Surface relief at twins is very apparent. The m e t a l l o g r a p h y o f the corroded surfaces is consistent with Fig. 5.

4. DISCUSSION The addition of silicon had an effect o n the transpassive part of the anodic polarization curve; higher current densities were measured as the silicon c o n t e n t increased (Figs. 3 and 4). Secondary passivity was observed in all the specimens. Wilde and Armijo [16] observed secondary passivity in their study using similar

materials; however, a systematic variation in current density with silicon content was n o t observed (Fig. 10). In terms of measured parameters such as E(corr), I(crit) and E(primary passive), very little difference in the t w o studies is present except that E(corr) is slightly more negative in the work of Wilde and Armijo [16] than in the present results (an average of - - 0 . 3 8 0 V (SCE) compared with --0.304 V (SCE)). The greatest differences between the two studies are the passive and transpassive current densities. The passive current densities obtained in ref. 16 were between 2 and 3 p A cm -2 whereas the passive current density in Fig. 3 is

171

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(c) Fig. 9. Scanning electron microscopy micrographs of stainless steels after being held at 1.375 V (SCE) for 3 h: (a) specimen 2 (100 ppm Si); (b) specimen 3 (1800 ppm Si); (c) specimen 5 (4800 ppm Si). (Magnifications: (a) 1500X; (b), (c) 1650X.)

about 1 pA c m -2. Such differences may have been caused by the difference in surface preparation. The specimens of Wilde and Armijo [16] were polished with 600 grit silicon carbide paper and were then briefly electropolished and dried in absolute ethyl alcohol. The previous transpassive current densities [16] were also slightly higher in the region of secondary passivity. However, both specimens containing 4.15 wt.% Si behaved similarly and both had the highest transpassive current densities. Other differences in the two

studies include different solution heat treatments (1050 °C in ref. 16) and differences in the electrode design; a mechanical mount similar to the original design of France [ 17] was used in the work of ref. 16. The data in Fig. 6 may be compared with those of Armijo [3] in Fig. 11. Armijo's curve was obtained using the nitric acid-potassium dichromate weight loss test. Intergranular attack increased with increasing silicon content (up to 4800 ppm Si), as can be seen in both figures; however, the surface attack also

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ppm Si.

increased in Fig. 6. The nitric acid-potassium dichromate test relies primarily on grain dropping, whereas the polarization technique records both intergranular and surface attack. Figures 7(b) and 7(c) show the transition from a broken pitting-type intergranular corrosion to continuous grooving, This is also shown in Figs. 8 and 9. The surface attack was significant in specimens that contained 1800 ppm Si or more, The greatest amount of surface attack took place in the 4800 ppm Si specimen. It is evident from this study that silicon does not affect only intergranular attack. The results show that both the grain boundaries and the surfaces are affected. With increasing silicon, both intergranular and surface attack increase. However, at 4.15 wt.% Si the corrosion resistance of the material increases drastically. This is evident in Fig. 7(d) and is consistent with the nitric acid-potassium dichromate test results [3] where no intergranular corrosion took place when the

4.15 wt.% Si specimen was tested (Fig. 11). The current density measured for this steel was very high even though little attack took place {Figs. 3 and 4), Corrosion appeared to take place as uniform surface attack in the 4.15 wt.% Si specimen. These results support a protective film mechanism for the high silicon {4,15 wt.%) stainless steel. Desestret [ 18] concluded that silicon accelerates the anodic dissolution of grain boundaries but simultaneously forms a passivating film. Competition between the two effects causes the maximum in Fig, 11. At low concentrations (below 0.5 wt.% Si) the silicon promotes susceptibility to intergranular corrosion. This deleterious effect of silicon on corrosion may be associated with silicon segregation to grain boundaries. Such segregation has been found in low alloy steels [19] and in iron containing 2.5 wt.% Si [20] by Auger electron spectroscopy. At high concentrations the silicon decreases the

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susceptibility. This passivation effect of silicon may account for the region of secondary passivity and the decreasing current density with increasing silicon content above 1.6 V (SCE) (Fig. 4). The results demonstrate the complexity involved in the interpretation of polarization curves. For example, the transpassive current density may not be a direct indication of the severity of localized corrosion attack. In order to determine what is happening it is necessary to observe corroded microstructures at particular current densities and potentials. In addition, measurements of the current densities both at the surface and at the grain boundary, e.g. by the use of microelectrodes [21], may be a useful approach to this problem. 5. CONCLUSIONS

(1) High purity (less than 50 ppm Si) and high silicon (approximately 4 wt.%) solution-

treated austenitic stainless steels are immune to intergranular corrosion at transpassive potentials in 1 N HzSO 4. (2) In these high purity steels the silicon content range 100 - 4800 ppm causes intergranular and surface corrosion. Between 1 and 1.55 V (SCE) the transpassive current density increases with silicon content. (3) Silicon (up to about 0.5 wt.%) appears to accelerate the anodic dissolution of grain boundaries in stainless steel but provides resistance to localized corrosion at high concentrations (approximately 4 wt.%). ACKNOWLEDGMENTS

The authors thank the Natural Sciences and Engineering Research Council of Canada for financial support. One of the authors (A.R.P.) held a Stelco Fellowship in Metallurgy during the course of this work. We are indebted to Dr. D. Dautovich, Ontario Hydro Division, for many helpful

174

discussions during the course of this work. We would also like to thank Mr. D. Abdulla for his help during the design and machining of the specimen electrode and Mr. A. L. M. Perrin for his work in designing the linear ramp generator.

REFERENCES 1 W. B. DeLong, A S T M Spec. Tech. Publ. 93, 1949, p. 211. 2 H . T . Shirley and J. E. Truman, J. Iron Steel Inst., London, 171 ( 1 9 5 2 ) 3 5 4 . 3 J. S. Armijo, Corrosion, 24 (1968) 24. 4 K. T. Aust, J. S. Armijo, E. F. Koch and J. H. Westbrook, A S M Trans. Q., 61 (1968) 24. 5 M. A. Streicher, J. Electrochem. Soc., 106 (1959) 161. 6 H. Coriou, J. Hure and G. Plante, Electrochim. Acta, 5 (1961) 105. 7 H. J. Schuller, P. Schwaab and W. Schwenk, Arch. E~enhuttenwes., 33 (1962) 853. 8 K. Osozawa, K. Bohenkamp and H. J. Engell, Corros. Sci., 6 (1966) 421.

9 G. Chaudron, EURAEC Rep. 976, Q. Rep. 6, October - December 1963 (European Atomic Energy Commission). 10 K. T. Aust, J. S. Armijo and J. H. Westbrook, Trans. Am. Soc. Met., 59 (1966) 544. 11 K. T. Aust, Trans. AIME, 245 (1969) 2117. 12 A. Joshi and D. F. Stein, Corrosion, 28 (1972) 321. 13 H. Chaung, J. B. Lumsden and R. W. Staehle, Metall. Trans. A, 10 (1979) 1853. 14 D. Bouchet and L. Priester, J. Mater. Sci., 14 (1979) 2205. 15 A. R. Perrin, M.A.Sc. Thesis, University of Toronto, 1978. 16 B. E. Wilde and J. S. Armijo, Corrosion, 24 (1968) 393. 17 W. D. France, Jr., J. Electrochem. Soc., 114 (1967) 818. 18 A. Desestret, Thesis, University of Paris, 1964. 19 A. Joshi, P. W. Palmberg and D. F. Stein, Metall. Trans. A, 6 (1975) 2160. 20 T. Watanabe, T. Murakami and S. Karashima, Scr Metall., 12 (1978) 361. 21 J. A. Davis, Localized corrosion, in Proc. 3rd Conf. on Corrosion, Williamsburg, VA, 1971, National Association of Corrosion Engineers, Houston, TX, 1974, p. 168.