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Electrochemical passivation of iron alloys and the film characterisation by XPS
Corrosion Science 42 (2000) 1709±1724 www.elsevier.com/locate/corsci
Electrochemical passivation of iron alloys and the ®lm characterisation by XPS Santanu Bera, S. Rangarajan, S.V. Narasimhan* Water and Steam Chemistry Laboratory, Bhabha Atomic Res. Ctr., BARC Facilities, IGCAR Campus, Kalpakkam 603 102, Tamil Nadu, India Received 4 April 1997; accepted 20 January 2000
Abstract The initial stages of oxide growth and composition of passive ®lms developed electrochemically on iron alloys viz. AISI 304, 316 and incoloy 800 were studied by X-ray photoelectron spectroscopy. The results showed an enrichment of chromium in the passive ®lms and the thickness of the ®lms were of the same order on all the alloys. The hydroxide species were found to dominate at the oxide solution interface. The air oxidised ®lm showed the presence of both chromium and iron in the oxidised state unlike the electrochemically developed passive ®lm where chromium was found to be the dominant oxidised species. 7 2000 Elsevier Science Ltd. All rights reserved. Keywords: Iron alloys; X-ray photoelectron spectroscopy; Passive ®lms
1. Introduction Passivation of stainless steel in sulphuric acid is a widely studied subject in corrosion science. The major part of the study includes the understanding of the surface microstructure, chemical structure, ®lm thickness and near surface depth distribution of the alloy components using techniques like electrochemical polarization, impedance and elipsometric measurements [1±3]. But none of these techniques yields a precise and direct chemical analysis of the passive ®lms. * Corresponding author. Tel.: +91-4114-40397; fax: +91-4114-40360. 0010-938X/00/$ - see front matter 7 2000 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 0 - 9 3 8 X ( 0 0 ) 0 0 0 2 8 - 7
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Realising the thickness of the passive ®lm is of the order of a few nanometer, several surface analytical tools like X-ray photoelectron spectroscopy (XPS), Auger electron spectroscopy (AES) and secondary ion mass spectroscopy (SIMS) are being employed for such analysis to derive the required information [4,5]. The growth kinetics of passivation in its initial stages was explained through the formation of iron and oxygen ions at the solution oxide interface and entrance of oxygen ions into the lattice by the `place exchange' mechanism [6]. This process is followed by the diusion mechanism for nucleation or further thickening of the ®lm. The mechanism of nucleation and growth of oxide ®lm on bare surface of stainless steel was explained by Burstein and Marshall [7] by analysing the current transients due to rapid mechanical removal of the oxide ®lm. Formation of the two distinct oxidised monolayers was observed on Ag in alkaline solution [8] by measuring current transients. During passivation, surface enrichment of alloy components, grain boundary phenomenon, monolayer formation and chemical states of the ®lm constituents are understood directly by surface sensitive techniques. The XPS measurements had shown that the surface of passive ®lms was partially hydrated [9]. The growth of passive ®lms at dierent potentials had been studied by SIMS [10]. The absence of (OH)ÿ ions within the passive ®lms on iron had been con®rmed by Mitchell et al. [11,12]. The presence of exchangable OHÿ ions was seen for Fe±Cr alloy passivation using SIMS. AES was employed mostly to get the quantitative chemical composition and depth pro®ling of passive ®lms. Enrichment of Mo was observed in Cr±Mo and Fe±Cr±Mo alloy by AES [13,14] and XPS [15]. Analytical approaches are also developed for analysing the complex nature of the passive ®lms. From the loss tail height of Fe 2p photoelectric peak, the in-depth layer information on iron alloy had been studied by Castle et al. [16]. A detailed numerical analysis of Auger depth pro®le technique was done on oxide ®lms on stainless steel by Olsson [17]. Quantitative AES depth pro®ling of passive ®lm formed on stainless steel was performed by Lorang et al., using sequential layer sputtering model [18]. There is a general agreement in literature about the passivation of stainless steel in sulphuric acid, e.g. (i) formation of Cr3+ compound dominantly and enrichment of chromium in the ®lm and (ii) thickness of the passive ®lm in nm range. Nevertheless, several important aspects of the ®lm structure and the initial stages of the ®lm growth or the kinetics of the ®lm formation are not fully established. In this paper, detailed XPS analysis of electrochemically developed passive ®lms at its initial stages of growth on iron alloys are presented. 2. Experimental methods The coupons of stainless steel 304 (composition in wt.%: C, 0.03; Si, <0.75; Mn, 1.5; P, <0.04; S, 0.03; Cr, 18; Ni, 8; Fe, bal.), stainless steel 316 (composition in wt.%: C, 0.04; Si, <0.75; Mn, 2.0; P, <0.04; S, 0.03; Mo, 2; Cr, 17.5; Ni, 11; Fe, bal.) and incoloy 800 (composition in wt.%: C, 0.03; Si, <0.75; Mn, 1.5; P,
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<0.04; Nb, 0.15; S, 0.03; Cr, 21; Ni, 32; Fe, bal.) were produced by casting hot and cold rolled to 2.5 mm thickness. The coupons were ®rst wet ground on 600±1000 grit silicon carbide papers followed by polishing with 2±3 mm diamond powder. The polished samples were rinsed in acetone and ultrasonically cleaned in methanol. The specimens were washed by deionised water and then dried in air. The dry coupons were taken for XPS analysis to see the oxide ®lm formed naturally by air. Potentiodynamic polarization curves for the three alloys were obtained in 0.1 M sulphuric acid and a suitable passive potential was selected for each of the alloys. Initially the surface oxide layer was removed by applying a negative potential to the sample. After 5 min, the coupons were passivated at the corresponding anodic potential for 40 min in the deaerated solution. A calomel electrode was chosen as the reference electrode and a passivating potential of 506 mV for AISI 304 and 316 and 306 mV for incoloy 800 was applied in order to grow the oxide ®lms. The passive potentials were chosen from the passive region of the polarization curves and were in the same range used for passivation of other alloy steels [19]. The sulphuric acid left after passivation was analysed by inductively coupled plasma mass spectroscopy (ICPMS) to estimate the metal ion dissolved in the process of passivation. The samples, after passivation were taken out, rinsed with deionised water, dried and loaded immediately in the fast-entry chamber (vacuum: 10ÿ7 mbar) of XPS system. The samples were kept in the chamber for several hours for drying prior to shifting to the analysis chamber (vacuum: 10ÿ10 mbar). The XPS analysis was carried out using a VG make ESCALAB MK200X equipped with dual X-ray source. Al Ka was used as the exciting source for photoelectrons, which were collected by hemispherical analyser kept at 0.4 eV resolution capability. The calibration of the binding energy was done by observing Au 4f7/2 photoelectron peak from Au/Si system at 84.0 eV [20]. The binding energy of Cls peak from contaminated C was observed at 284.6 eV and was used as the reference for peak binding energy calculation and peak identi®cation. The depth pro®ling was done by in-situ Ar ion sputtering gun, working at 4 keV beam energy and 1 mA beam current. The data processing and deconvolution of photoelectron peaks were done by `ECLIPSE V1.7T' packages supplied with the instrument. 3. Results and discussions 3.1. Nature of polarization curves The potentiodynamic polarization curves for the three alloys in 0.1 M sulphuric acid are shown in Fig. 1. It can be observed from the ®gure that AISI 316 is nobler and the corrosion rate is approximately 100 times smaller than AISI 304. However, there is not much dierence in the width of passive potential region between these two alloys. The extra nobilities and higher corrosion resistance of AISI 316 could be due to the presence of Mo. Though the corrosion potential of
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incloloy 800 was more positive than AISI 316, its corrosion rate was found to be higher in sulphuric acid medium possibly due to less protective nature of the surface ®lm compared to that on stainless steel. The passive current densities are higher in case of AISI 304 and incoloy 800 (compared to AISI 316) due to the absence of Mo which is a good passivator. The presence of Mo normally makes the passive ®lm more stable and protective [19]. Fig. 2 gives the variation of current with time for the three alloys. It is seen that the current reaches a steady state value within 15 min. The area under the curve represents the total charge passed through the electrodes during 40 min of
Fig. 1. Polarization curves for alloys polarised in 0.1 M H2SO4 solution. The section of the curves with open symbols represent the cathodic branches of the polarization curves.
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passivation. It is noticed from the curves that the rate of achieving the steady state increases in order AISI 304 < AISI 316 < incoloy 800. 3.2. Analysis of the oxide ®lms by XPS The chemical states of the alloy components in the passive ®lms were identi®ed
Fig. 2. Variation of current with exposure time for dierent alloys.
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by XPS. The photoelectrons corresponding to Cr 2p3/2, Fe 2p3/2, Ols and Ni 2p3/2 emanating from the surface and at dierent depths were recorded. It was observed that in the three alloys chromium was present in Cr0 and Cr3+ oxidation states. The Cr 2p3/2 peak registered for ®lms developed on AISI 304 for 40 and 60 min of passivation are shown in Fig. 3 as curves a and b, respectively. Cr 2p3/2 peak positions were observed at 574.1 and 577.2 eV for Cr0 and Cr3+ respectively. In case of the ®lm passivated for 60 min the Cr0 signal was not observed. Normally, in XPS, at a K.E. of approximately 1000 eV, 60% of the signal comes from mean free path depth l and 95% of the signal comes from 3l [21]. As the ®lms on the coupons are very thin compared to the analytical depth, the signals from substrate below the ®lm are expected. The thickness of the ®lm grown in 40 min was not sucient enough to attenuate the Cr0 signal from the substrate completely as indicated in Fig. 3 (peak at 574.1 eV) but in 60 min
Fig. 3. XPS of Cr 2p3/2 and Fe 2p3/2 for passivated (at +506 mV, SCE) AISI 304: for (a) 40 min and (b) 60 min. The peaks at 574.1 eV (marked arrow) and 577.2 eV represent the Cr0 and oxidised Cr, respectively. The peaks at 706.9 and 710.2 eV indicate the presence of Fe0 and oxidised iron.
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passivation, the non-appearance of Cr0 signal meant that signal from the substrate was completely attenuated. Though the Cr0 signal vanished with time of passivation the Fe 2p3/2 signal did not vanish much with time as shown in Fig. 3. The peak at 706.9 eV corresponds to the Fe0 and a broad peak centering at 710.2 eV corresponds to the Fe3+ or Fe2+ species [1]. The mean free path of photoelectrons l is proportional to the square root of kinetic energy. The calculation, [22] show almost the same mean free path l considering the kinetic energies of Fe 2p3/2 and Cr 2p3/2 photoelectrons and hence should give the same analytical depth of information. Thus, the presence of Fe0 and the absence of Cr0 peak indicated that the Fe0 signal was coming from the surface ®lm. Hence, in the ®lm developed for 40 min, the intense Fe0 signal was coming from the substrate and the ®lm. In case of AISI 316 and incoloy 800 a lot of similarities to AISI 304 were observed in the passive ®lms. Chromium in these alloys had undergone dominating chemical modi®cation to Cr3+. Iron was mostly non-oxidised, but a small broad peak around 710.2 eV indicated the formation of Fe3+ or Fe2+ oxidation state. In AISI 316 Mo showed chemical modi®cation with peak positions of Mo 3d5/2 at 227.5 eV and at around 232.5 eV owing to the formation of Mo+6 chemical species [1]. In all the three cases it was observed that iron (be it in Fe2+ or Fe3+) was sputtered out very quickly indicating that Fe2+/3+ was present at the oxide solution interface only. However, the chromium (Cr3+) state was not sputtered out so quickly like iron (Fe2+/3+), due to the fact that Cr3+ was lodged in the oxide matrix throughout. The variation in the ratio of Cr3+ to Cr0 with ®lm depth, obtained by deconvoluting Cr 2p3/2 photoelectric peaks at dierent sputtering time is shown in Fig. 4, where the curves sharply bend downwards. This indicated that non-reacted Cr and Cr3+ were not present in the same atomic layer, rather they formed sharp interface in the ®lm. Thus, the thickness of the ®lm was logically governed by the depth upto which Cr3+ was present. The ®lm was sputtered at the rate of around 2 AÊ/min and within 10 min the Cr3+ was removed almost completely indicating the thickness was of the order of 20 AÊ for 40 min of passivation. Fig. 5 shows the enrichment of Cr with respect to Fe. The photoelectric peak area ratio of chromium to iron was observed to vary with ®lm depth. The Cr/Fe ratio was observed to be 0.15, 0.17 and 0.22 for AISI 304, 316 and incoloy 800 before passivation and increased in all the ®lm surfaces after passivation. The enrichment of chromium on the ®lm surface might have happened because of preferential dissolution of Ni into the solution [15]. Thus, chromium enrichment was the major process and chromium oxidation was the major chemical modi®cation even at the very initial stages of the growth. In all these ®lms the Cr3+ 2p3/2 peak was observed to have larger FWHM (full width at half maxima) than Cr3+ in Cr2O3 (ca. 3 eV, [23] for XPS of oxides). Thus, the presence of chromium hydroxide was inferred. The deconvolution of Cr 2p3/2 peak into Cr3+ in hydroxide and Cr3+ in oxide (separated by around 0.5 eV [24]) for the three alloy systems were carried out to obtain the chromium-oxide and chromium-hydroxide contributions. The deconvoluted peak area ratio
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Fig. 4. Plots of oxidised Cr to unoxidised Cr atomic concentration ratio (Cr3+/Cr0) in the passive ®lm with depth.
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Fig. 5. Plots of Cr 2p3/2 to Fe 2p3/2 photoelectron counts (intensity) ratio in the passive ®lm with depth. (The photoelectron count is proportional to the atomic concentration of the element.)
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(chromium-hydroxide:chromium-oxide) obtained at dierent depths are plotted with ®lm depth as shown in Fig. 6. It was observed that Cr3+ for hydroxide dominantly stayed at the ®lm solution interface and Cr3+ for oxide grew below the hydroxide. The kinetics of growth was dominantly governed by the dissolution of the ions into the solution. The ionic species of the alloy components were formed during passivation and were dissolved into the solution according to the selectivity factor [15]. The selectivity factor is de®ned as the S(Mi) = (Mi% in solution)/(Mi% in alloy), where Mi% represents the composition ratio of any element to sum of all components. In AISI 304, the analysis of the solution by ICPMS for iron showed S(Fe) =1.0 and indicated the uniform dissolution of iron. S(Ni) was 1.2 which indicated the selective dissolution of Ni during passivation. S(Cr) was 0.8 and hence dissolution of chromium was less than the amount present in the alloy. These results agrees well with the previous work by castle et al. [15]. Similarly for AISI 316 S(Fe) = 1.0, S(Ni) = 1.2 and S(Cr) = 0.65 and for incoloy S(Fe) = 1.0, S(Ni) = 1.2 and S(Cr) = 0.62 was observed. The amount of Ni at dierent depth of the ®lm was estimated by XPS. In Fig. 7 the variation of absolute counts of Ni 2p3/2 is shown to increase with time. This complemented the results obtained by ICPMS that Ni was preferentially dissolved in the solution during passivation, hence decreasing the counts at the surface. Thus, understanding the nature of the passive ®lm described so far, it was predicted that initially the OHÿ group was associated with chromium of the alloy dominantly to form the chromium hydroxide at the ®lm solution interface and gradually that might have formed oxide layer below the hydroxide layer. As Ni was preferentially dissolved into the solution, either Ni did not form any compound species or probably was below the detectable limit. Iron dissolved as Fe2+ into the solution [15] and its extent of dissolution was almost the same as the percentage in the alloy. Hence iron present in oxide solution interface was in trace quantity and was mostly in Fe3+ or Fe2+ chemical state. As the selectivity factor of Cr < 1, i.e. it dissolves less than its amount present in the alloy, Cr3+ becomes the dominating component in the ®lm. From the above results, it could be concluded that as passivation proceeded, chromium was oxidised to (+3) state and segregated in the form of a localised oxide coating. Ni and the major portion of iron were released in solution. The amount of iron that was left behind, remained in unoxidised form. Since unoxidised iron can not coexist in the lattice of chromium oxide it has to be adjacent to the chromium-oxide grain. Thus, the model explains the appearance of intense Fe0 signal in 40 or 60 min passive ®lm. As the analytical depth is greater than the ®lm thickness (developed for 40 min) a part of the signal was coming from the substrate. From substrate Cr0 signal area the proportional Fe signal area should be around 4.8 times the area of Cr0 signal. But in the present case overall Fe0 signal was around 10 times the Cr0 signal. So it is logical to presume that the Fe0 signal was coming from the ®lm also. The photoelectron peaks for Cr 2p, Fe 2p, and O ls are compared for the three passive ®lms. The Cr 2p did not show much change due to the change in the alloy
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Fig. 6. Plots of chromium-hydroxide to chromium-oxide atomic concentration ratio in the passive ®lms with depth.
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Fig. 7. Depth pro®le of Ni 2p3/2 photoelectron counts (intensity) in the passive ®lms for the three alloys. Near the surface of the ®lm Ni was observed to be less because of the selective dissolution of Ni.
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composition and in all cases shoulder of Cr0 was observed along with the Cr3+ species as discussed earlier. In Fig. 8, Fe 2p3/2 signal acquired from the passive ®lms developed on the alloys are compared. The broad peak resulted from the Fe2+ and Fe3+ species was observed around 710.5 eV for AISI 304 and AISI 316, but shifted to 711.2 eV in case of incoloy which may be due to the formation of Fe±OH like species in the passive ®lm on incoloy. The FWHM of O ls was observed to be more than 3 eV in all cases due to the presence of O in dierent chemical environment in the ®lms. O ls peak was observed to be asymmetric, which was deconvoluted into two peaks at 531.7 and 530.3 eV. The peak at 531.7 eV con®rmed the presence of hydroxide in the ®lm. The deconvolution of these peaks showed the presence of M±O and M±OH groups in the ®lms, which complemented the results obtained from Cr 2p3/2 peak analysis. Sulphur 2p line was noted for the three cases and in each case it showed two peaks at 169.5 and 162.5 eV. The peak at 169.5 eV may be contributed by the sulphate group adsorbed on the ®lm which was con®rmed by a subtle tailing of O ls peak towards higher binding energy and was removed very easily from the ®lm. But the peak at 162.5 eV vanished only after several layers of sputtering. However, this peak was shifted around 1.5 eV below the normal S at 164.04 eV. Hence, H±S group formation was also predicted. In Table 1, the peak energies of Fe 2p3/2, Cr 2p3/2 and O ls photoelectrons from the ®lm and naturally developed oxide surface are given. It is clear form that table that the chemical states of iron and chromium for the naturally developed oxide ®lm is dierent from the electrochemically grown oxide ®lms. It was observed that both chromium enrichment and the chemical composition could distinguish the two processes. In passivated ®lm Cr3+ was the dominating oxidation process, but Table 1 XPS peak positions for various components in the oxide ®lmsa
AISI 304 Naturally grown oxide Passivated at +506 mV (SCE) AISI 316 Naturally grown oxide Passivated at +506 mV (SCE) Incoloy 800 Naturally grown oxide Passivated at +306 mV (SCE) a
Fe 2p3/2 (eV)
Cr 2p3/2 (eV)
O ls (eV)
709.9 706.9 710.2 706.7
576.6
530.5
577.2 574.1
532.0 530.5
576.5
532.0 530.2 531.7 530.3
710.2 706.9 710.4 706.8 711.1 706.9 711.1 706.9
Bold represents the dominating species.
576.9 574.1 576.8 577.5 574.5
532.1 530.5 530.8 532.2
Cr/O (area)
0.56
0.74
0.84
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Fig. 8. Fe 2p3/2 peak for the three alloys after passivation for 40 min: (a) AISI 316 passivated at +506 mV (SCE), (b) AISI 304 passivated at +506 mV (SCE) and (c) incoloy 800 passivated at +306 mV (SCE).
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in the air grown oxide both iron and chromium were oxidised. In comparison of the three alloys, Fe/O ratio for AISI 304 and AISI 316 were almost same but Cr/ O ratio increased suggesting that AISI 316 is a better alloy for passivation. The change in the Cr/O content may be dominantly due to the change in the chromium content in the analyzing area. In case of incoloy the change may be logical because of more chromium content in the alloy itself, but in AISI 316 the chromium increase might have occurred due to the presence of Mo in the alloy. Because of the dissolution of Mo, the dissolution of chromium was decreased consequently increasing the ratio. 4. Conclusions 1. The phenomenon of chromium enrichment was observed in all the three alloys during initial stage of passivation and the Cr3+ oxidation state was found to be the dominating oxidation process. 2. The oxygen containing species from solution initially attacks chromium of the alloy in contact with the solution and forms hydroxides at the ®lm solution interface and successively oxide below the interface probably by place exchange mechanism. The nucleation of the oxide ®lm is occurred by the segregation of Cr3+ species leaving unoxidised iron and Ni in the neighbourhood of the Cr2O3. 3. Normally grown oxide ®lm showed both chromium and iron in the oxidised state, where as the passive ®lm showed only chromium was oxidised dominantly.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17]
P. Marcus, I. Olefjord, Corros. Sci. 28 (1988) 589. H.W. Hoppe, S. Haupt, H.H. Strehblow, Surf. Interface Anal. 21 (1994) 514. N.E. Hakiki, S. Boudin, B. Rondot, M. DA Cunha Belo, Corros. Sci. 37 (1995) 1809. G. Okammoto, T. Shibata, Corros. Sci. 10 (1970) 371. M.J. Graham, Corros. Sci. 37 (1995) 1377. N. Sato, M. Cohen, J. Electrochem. Soc. 111 (1964) 512. G.T. Burstein, P.I. Marshall, Corros. Sci. 23 (1983) 125. G.T. Burstein, R.C. Newman, Electrochem. Acta 25 (1980) 1009. K. Asami, K. Hosimoto, S. Shimodaria, Corros. Sci. 18 (1978) 151. J.A. Bardwell, B. MacDougall, M.J. Graham, J. Electrochem. Soc. 135 (1988) 413. D.F. Mitchell, G.I. Proule, M.J. Graham, Appl. Surf. Sci. 21 (1985) 199. D.F. Mitchell, M.J. Graham, J. Electrochem. Soc. 133 (1986) 936. R. Goetz, D. Landot, Electrochim. Acta 27 (1982) 1061. R. Goetz, D. Landot, Electrochim. Acta 29 (1984) 667. J.E. Castle, J.H. Qiu, J. Electrochem. Soc. 137 (1990) 2033. J.E. Castle, Ruoru Ke, J.F. Natts, Corros. Sci. 30 (1990) 771. C.-O.A. Olsson, Surf. Interface Anal. 21 (1994) 846.
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[18] G. Lorang, J.L. Xu, J.-P. Langeron, Surf. Interface Anal. 19 (1992) 60. [19] J.E. Castle, J.H. Qiu, Corros. Sci. 29 (1989) 591. [20] D.K. Sarkar, S. Bera, S. Dhara, K.G.M. Nair, S.V. Narasimhan, S. Chowdhury, Appl. Surf. Sci. 120 (1997) 159. [21] C.J. Powell, Surf. Sci. 44 (1974) 29. [22] D.B. Briggs, M.P. Seah, Practical Surface Analysis by Auger and XPS, New York, 1987, p. 187. [23] K. Asami, K. Hashimoto, Corros. Sci. 17 (1977) 559. [24] I. Grohmann, E. Kemnitz, A. Lippitz, W.E.S. Unger, Surf. Interface Anal. 23 (1995) 887.
Electrochemical passivation of iron alloys and the ®lm characterisation by XPS Santanu Bera, S. Rangarajan, S.V. Narasimhan* Water and Steam Chemistry Laboratory, Bhabha Atomic Res. Ctr., BARC Facilities, IGCAR Campus, Kalpakkam 603 102, Tamil Nadu, India Received 4 April 1997; accepted 20 January 2000
Abstract The initial stages of oxide growth and composition of passive ®lms developed electrochemically on iron alloys viz. AISI 304, 316 and incoloy 800 were studied by X-ray photoelectron spectroscopy. The results showed an enrichment of chromium in the passive ®lms and the thickness of the ®lms were of the same order on all the alloys. The hydroxide species were found to dominate at the oxide solution interface. The air oxidised ®lm showed the presence of both chromium and iron in the oxidised state unlike the electrochemically developed passive ®lm where chromium was found to be the dominant oxidised species. 7 2000 Elsevier Science Ltd. All rights reserved. Keywords: Iron alloys; X-ray photoelectron spectroscopy; Passive ®lms
1. Introduction Passivation of stainless steel in sulphuric acid is a widely studied subject in corrosion science. The major part of the study includes the understanding of the surface microstructure, chemical structure, ®lm thickness and near surface depth distribution of the alloy components using techniques like electrochemical polarization, impedance and elipsometric measurements [1±3]. But none of these techniques yields a precise and direct chemical analysis of the passive ®lms. * Corresponding author. Tel.: +91-4114-40397; fax: +91-4114-40360. 0010-938X/00/$ - see front matter 7 2000 Elsevier Science Ltd. All rights reserved. PII: S 0 0 1 0 - 9 3 8 X ( 0 0 ) 0 0 0 2 8 - 7
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Realising the thickness of the passive ®lm is of the order of a few nanometer, several surface analytical tools like X-ray photoelectron spectroscopy (XPS), Auger electron spectroscopy (AES) and secondary ion mass spectroscopy (SIMS) are being employed for such analysis to derive the required information [4,5]. The growth kinetics of passivation in its initial stages was explained through the formation of iron and oxygen ions at the solution oxide interface and entrance of oxygen ions into the lattice by the `place exchange' mechanism [6]. This process is followed by the diusion mechanism for nucleation or further thickening of the ®lm. The mechanism of nucleation and growth of oxide ®lm on bare surface of stainless steel was explained by Burstein and Marshall [7] by analysing the current transients due to rapid mechanical removal of the oxide ®lm. Formation of the two distinct oxidised monolayers was observed on Ag in alkaline solution [8] by measuring current transients. During passivation, surface enrichment of alloy components, grain boundary phenomenon, monolayer formation and chemical states of the ®lm constituents are understood directly by surface sensitive techniques. The XPS measurements had shown that the surface of passive ®lms was partially hydrated [9]. The growth of passive ®lms at dierent potentials had been studied by SIMS [10]. The absence of (OH)ÿ ions within the passive ®lms on iron had been con®rmed by Mitchell et al. [11,12]. The presence of exchangable OHÿ ions was seen for Fe±Cr alloy passivation using SIMS. AES was employed mostly to get the quantitative chemical composition and depth pro®ling of passive ®lms. Enrichment of Mo was observed in Cr±Mo and Fe±Cr±Mo alloy by AES [13,14] and XPS [15]. Analytical approaches are also developed for analysing the complex nature of the passive ®lms. From the loss tail height of Fe 2p photoelectric peak, the in-depth layer information on iron alloy had been studied by Castle et al. [16]. A detailed numerical analysis of Auger depth pro®le technique was done on oxide ®lms on stainless steel by Olsson [17]. Quantitative AES depth pro®ling of passive ®lm formed on stainless steel was performed by Lorang et al., using sequential layer sputtering model [18]. There is a general agreement in literature about the passivation of stainless steel in sulphuric acid, e.g. (i) formation of Cr3+ compound dominantly and enrichment of chromium in the ®lm and (ii) thickness of the passive ®lm in nm range. Nevertheless, several important aspects of the ®lm structure and the initial stages of the ®lm growth or the kinetics of the ®lm formation are not fully established. In this paper, detailed XPS analysis of electrochemically developed passive ®lms at its initial stages of growth on iron alloys are presented. 2. Experimental methods The coupons of stainless steel 304 (composition in wt.%: C, 0.03; Si, <0.75; Mn, 1.5; P, <0.04; S, 0.03; Cr, 18; Ni, 8; Fe, bal.), stainless steel 316 (composition in wt.%: C, 0.04; Si, <0.75; Mn, 2.0; P, <0.04; S, 0.03; Mo, 2; Cr, 17.5; Ni, 11; Fe, bal.) and incoloy 800 (composition in wt.%: C, 0.03; Si, <0.75; Mn, 1.5; P,
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<0.04; Nb, 0.15; S, 0.03; Cr, 21; Ni, 32; Fe, bal.) were produced by casting hot and cold rolled to 2.5 mm thickness. The coupons were ®rst wet ground on 600±1000 grit silicon carbide papers followed by polishing with 2±3 mm diamond powder. The polished samples were rinsed in acetone and ultrasonically cleaned in methanol. The specimens were washed by deionised water and then dried in air. The dry coupons were taken for XPS analysis to see the oxide ®lm formed naturally by air. Potentiodynamic polarization curves for the three alloys were obtained in 0.1 M sulphuric acid and a suitable passive potential was selected for each of the alloys. Initially the surface oxide layer was removed by applying a negative potential to the sample. After 5 min, the coupons were passivated at the corresponding anodic potential for 40 min in the deaerated solution. A calomel electrode was chosen as the reference electrode and a passivating potential of 506 mV for AISI 304 and 316 and 306 mV for incoloy 800 was applied in order to grow the oxide ®lms. The passive potentials were chosen from the passive region of the polarization curves and were in the same range used for passivation of other alloy steels [19]. The sulphuric acid left after passivation was analysed by inductively coupled plasma mass spectroscopy (ICPMS) to estimate the metal ion dissolved in the process of passivation. The samples, after passivation were taken out, rinsed with deionised water, dried and loaded immediately in the fast-entry chamber (vacuum: 10ÿ7 mbar) of XPS system. The samples were kept in the chamber for several hours for drying prior to shifting to the analysis chamber (vacuum: 10ÿ10 mbar). The XPS analysis was carried out using a VG make ESCALAB MK200X equipped with dual X-ray source. Al Ka was used as the exciting source for photoelectrons, which were collected by hemispherical analyser kept at 0.4 eV resolution capability. The calibration of the binding energy was done by observing Au 4f7/2 photoelectron peak from Au/Si system at 84.0 eV [20]. The binding energy of Cls peak from contaminated C was observed at 284.6 eV and was used as the reference for peak binding energy calculation and peak identi®cation. The depth pro®ling was done by in-situ Ar ion sputtering gun, working at 4 keV beam energy and 1 mA beam current. The data processing and deconvolution of photoelectron peaks were done by `ECLIPSE V1.7T' packages supplied with the instrument. 3. Results and discussions 3.1. Nature of polarization curves The potentiodynamic polarization curves for the three alloys in 0.1 M sulphuric acid are shown in Fig. 1. It can be observed from the ®gure that AISI 316 is nobler and the corrosion rate is approximately 100 times smaller than AISI 304. However, there is not much dierence in the width of passive potential region between these two alloys. The extra nobilities and higher corrosion resistance of AISI 316 could be due to the presence of Mo. Though the corrosion potential of
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incloloy 800 was more positive than AISI 316, its corrosion rate was found to be higher in sulphuric acid medium possibly due to less protective nature of the surface ®lm compared to that on stainless steel. The passive current densities are higher in case of AISI 304 and incoloy 800 (compared to AISI 316) due to the absence of Mo which is a good passivator. The presence of Mo normally makes the passive ®lm more stable and protective [19]. Fig. 2 gives the variation of current with time for the three alloys. It is seen that the current reaches a steady state value within 15 min. The area under the curve represents the total charge passed through the electrodes during 40 min of
Fig. 1. Polarization curves for alloys polarised in 0.1 M H2SO4 solution. The section of the curves with open symbols represent the cathodic branches of the polarization curves.
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passivation. It is noticed from the curves that the rate of achieving the steady state increases in order AISI 304 < AISI 316 < incoloy 800. 3.2. Analysis of the oxide ®lms by XPS The chemical states of the alloy components in the passive ®lms were identi®ed
Fig. 2. Variation of current with exposure time for dierent alloys.
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by XPS. The photoelectrons corresponding to Cr 2p3/2, Fe 2p3/2, Ols and Ni 2p3/2 emanating from the surface and at dierent depths were recorded. It was observed that in the three alloys chromium was present in Cr0 and Cr3+ oxidation states. The Cr 2p3/2 peak registered for ®lms developed on AISI 304 for 40 and 60 min of passivation are shown in Fig. 3 as curves a and b, respectively. Cr 2p3/2 peak positions were observed at 574.1 and 577.2 eV for Cr0 and Cr3+ respectively. In case of the ®lm passivated for 60 min the Cr0 signal was not observed. Normally, in XPS, at a K.E. of approximately 1000 eV, 60% of the signal comes from mean free path depth l and 95% of the signal comes from 3l [21]. As the ®lms on the coupons are very thin compared to the analytical depth, the signals from substrate below the ®lm are expected. The thickness of the ®lm grown in 40 min was not sucient enough to attenuate the Cr0 signal from the substrate completely as indicated in Fig. 3 (peak at 574.1 eV) but in 60 min
Fig. 3. XPS of Cr 2p3/2 and Fe 2p3/2 for passivated (at +506 mV, SCE) AISI 304: for (a) 40 min and (b) 60 min. The peaks at 574.1 eV (marked arrow) and 577.2 eV represent the Cr0 and oxidised Cr, respectively. The peaks at 706.9 and 710.2 eV indicate the presence of Fe0 and oxidised iron.
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passivation, the non-appearance of Cr0 signal meant that signal from the substrate was completely attenuated. Though the Cr0 signal vanished with time of passivation the Fe 2p3/2 signal did not vanish much with time as shown in Fig. 3. The peak at 706.9 eV corresponds to the Fe0 and a broad peak centering at 710.2 eV corresponds to the Fe3+ or Fe2+ species [1]. The mean free path of photoelectrons l is proportional to the square root of kinetic energy. The calculation, [22] show almost the same mean free path l considering the kinetic energies of Fe 2p3/2 and Cr 2p3/2 photoelectrons and hence should give the same analytical depth of information. Thus, the presence of Fe0 and the absence of Cr0 peak indicated that the Fe0 signal was coming from the surface ®lm. Hence, in the ®lm developed for 40 min, the intense Fe0 signal was coming from the substrate and the ®lm. In case of AISI 316 and incoloy 800 a lot of similarities to AISI 304 were observed in the passive ®lms. Chromium in these alloys had undergone dominating chemical modi®cation to Cr3+. Iron was mostly non-oxidised, but a small broad peak around 710.2 eV indicated the formation of Fe3+ or Fe2+ oxidation state. In AISI 316 Mo showed chemical modi®cation with peak positions of Mo 3d5/2 at 227.5 eV and at around 232.5 eV owing to the formation of Mo+6 chemical species [1]. In all the three cases it was observed that iron (be it in Fe2+ or Fe3+) was sputtered out very quickly indicating that Fe2+/3+ was present at the oxide solution interface only. However, the chromium (Cr3+) state was not sputtered out so quickly like iron (Fe2+/3+), due to the fact that Cr3+ was lodged in the oxide matrix throughout. The variation in the ratio of Cr3+ to Cr0 with ®lm depth, obtained by deconvoluting Cr 2p3/2 photoelectric peaks at dierent sputtering time is shown in Fig. 4, where the curves sharply bend downwards. This indicated that non-reacted Cr and Cr3+ were not present in the same atomic layer, rather they formed sharp interface in the ®lm. Thus, the thickness of the ®lm was logically governed by the depth upto which Cr3+ was present. The ®lm was sputtered at the rate of around 2 AÊ/min and within 10 min the Cr3+ was removed almost completely indicating the thickness was of the order of 20 AÊ for 40 min of passivation. Fig. 5 shows the enrichment of Cr with respect to Fe. The photoelectric peak area ratio of chromium to iron was observed to vary with ®lm depth. The Cr/Fe ratio was observed to be 0.15, 0.17 and 0.22 for AISI 304, 316 and incoloy 800 before passivation and increased in all the ®lm surfaces after passivation. The enrichment of chromium on the ®lm surface might have happened because of preferential dissolution of Ni into the solution [15]. Thus, chromium enrichment was the major process and chromium oxidation was the major chemical modi®cation even at the very initial stages of the growth. In all these ®lms the Cr3+ 2p3/2 peak was observed to have larger FWHM (full width at half maxima) than Cr3+ in Cr2O3 (ca. 3 eV, [23] for XPS of oxides). Thus, the presence of chromium hydroxide was inferred. The deconvolution of Cr 2p3/2 peak into Cr3+ in hydroxide and Cr3+ in oxide (separated by around 0.5 eV [24]) for the three alloy systems were carried out to obtain the chromium-oxide and chromium-hydroxide contributions. The deconvoluted peak area ratio
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Fig. 4. Plots of oxidised Cr to unoxidised Cr atomic concentration ratio (Cr3+/Cr0) in the passive ®lm with depth.
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Fig. 5. Plots of Cr 2p3/2 to Fe 2p3/2 photoelectron counts (intensity) ratio in the passive ®lm with depth. (The photoelectron count is proportional to the atomic concentration of the element.)
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(chromium-hydroxide:chromium-oxide) obtained at dierent depths are plotted with ®lm depth as shown in Fig. 6. It was observed that Cr3+ for hydroxide dominantly stayed at the ®lm solution interface and Cr3+ for oxide grew below the hydroxide. The kinetics of growth was dominantly governed by the dissolution of the ions into the solution. The ionic species of the alloy components were formed during passivation and were dissolved into the solution according to the selectivity factor [15]. The selectivity factor is de®ned as the S(Mi) = (Mi% in solution)/(Mi% in alloy), where Mi% represents the composition ratio of any element to sum of all components. In AISI 304, the analysis of the solution by ICPMS for iron showed S(Fe) =1.0 and indicated the uniform dissolution of iron. S(Ni) was 1.2 which indicated the selective dissolution of Ni during passivation. S(Cr) was 0.8 and hence dissolution of chromium was less than the amount present in the alloy. These results agrees well with the previous work by castle et al. [15]. Similarly for AISI 316 S(Fe) = 1.0, S(Ni) = 1.2 and S(Cr) = 0.65 and for incoloy S(Fe) = 1.0, S(Ni) = 1.2 and S(Cr) = 0.62 was observed. The amount of Ni at dierent depth of the ®lm was estimated by XPS. In Fig. 7 the variation of absolute counts of Ni 2p3/2 is shown to increase with time. This complemented the results obtained by ICPMS that Ni was preferentially dissolved in the solution during passivation, hence decreasing the counts at the surface. Thus, understanding the nature of the passive ®lm described so far, it was predicted that initially the OHÿ group was associated with chromium of the alloy dominantly to form the chromium hydroxide at the ®lm solution interface and gradually that might have formed oxide layer below the hydroxide layer. As Ni was preferentially dissolved into the solution, either Ni did not form any compound species or probably was below the detectable limit. Iron dissolved as Fe2+ into the solution [15] and its extent of dissolution was almost the same as the percentage in the alloy. Hence iron present in oxide solution interface was in trace quantity and was mostly in Fe3+ or Fe2+ chemical state. As the selectivity factor of Cr < 1, i.e. it dissolves less than its amount present in the alloy, Cr3+ becomes the dominating component in the ®lm. From the above results, it could be concluded that as passivation proceeded, chromium was oxidised to (+3) state and segregated in the form of a localised oxide coating. Ni and the major portion of iron were released in solution. The amount of iron that was left behind, remained in unoxidised form. Since unoxidised iron can not coexist in the lattice of chromium oxide it has to be adjacent to the chromium-oxide grain. Thus, the model explains the appearance of intense Fe0 signal in 40 or 60 min passive ®lm. As the analytical depth is greater than the ®lm thickness (developed for 40 min) a part of the signal was coming from the substrate. From substrate Cr0 signal area the proportional Fe signal area should be around 4.8 times the area of Cr0 signal. But in the present case overall Fe0 signal was around 10 times the Cr0 signal. So it is logical to presume that the Fe0 signal was coming from the ®lm also. The photoelectron peaks for Cr 2p, Fe 2p, and O ls are compared for the three passive ®lms. The Cr 2p did not show much change due to the change in the alloy
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Fig. 6. Plots of chromium-hydroxide to chromium-oxide atomic concentration ratio in the passive ®lms with depth.
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Fig. 7. Depth pro®le of Ni 2p3/2 photoelectron counts (intensity) in the passive ®lms for the three alloys. Near the surface of the ®lm Ni was observed to be less because of the selective dissolution of Ni.
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composition and in all cases shoulder of Cr0 was observed along with the Cr3+ species as discussed earlier. In Fig. 8, Fe 2p3/2 signal acquired from the passive ®lms developed on the alloys are compared. The broad peak resulted from the Fe2+ and Fe3+ species was observed around 710.5 eV for AISI 304 and AISI 316, but shifted to 711.2 eV in case of incoloy which may be due to the formation of Fe±OH like species in the passive ®lm on incoloy. The FWHM of O ls was observed to be more than 3 eV in all cases due to the presence of O in dierent chemical environment in the ®lms. O ls peak was observed to be asymmetric, which was deconvoluted into two peaks at 531.7 and 530.3 eV. The peak at 531.7 eV con®rmed the presence of hydroxide in the ®lm. The deconvolution of these peaks showed the presence of M±O and M±OH groups in the ®lms, which complemented the results obtained from Cr 2p3/2 peak analysis. Sulphur 2p line was noted for the three cases and in each case it showed two peaks at 169.5 and 162.5 eV. The peak at 169.5 eV may be contributed by the sulphate group adsorbed on the ®lm which was con®rmed by a subtle tailing of O ls peak towards higher binding energy and was removed very easily from the ®lm. But the peak at 162.5 eV vanished only after several layers of sputtering. However, this peak was shifted around 1.5 eV below the normal S at 164.04 eV. Hence, H±S group formation was also predicted. In Table 1, the peak energies of Fe 2p3/2, Cr 2p3/2 and O ls photoelectrons from the ®lm and naturally developed oxide surface are given. It is clear form that table that the chemical states of iron and chromium for the naturally developed oxide ®lm is dierent from the electrochemically grown oxide ®lms. It was observed that both chromium enrichment and the chemical composition could distinguish the two processes. In passivated ®lm Cr3+ was the dominating oxidation process, but Table 1 XPS peak positions for various components in the oxide ®lmsa
AISI 304 Naturally grown oxide Passivated at +506 mV (SCE) AISI 316 Naturally grown oxide Passivated at +506 mV (SCE) Incoloy 800 Naturally grown oxide Passivated at +306 mV (SCE) a
Fe 2p3/2 (eV)
Cr 2p3/2 (eV)
O ls (eV)
709.9 706.9 710.2 706.7
576.6
530.5
577.2 574.1
532.0 530.5
576.5
532.0 530.2 531.7 530.3
710.2 706.9 710.4 706.8 711.1 706.9 711.1 706.9
Bold represents the dominating species.
576.9 574.1 576.8 577.5 574.5
532.1 530.5 530.8 532.2
Cr/O (area)
0.56
0.74
0.84
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Fig. 8. Fe 2p3/2 peak for the three alloys after passivation for 40 min: (a) AISI 316 passivated at +506 mV (SCE), (b) AISI 304 passivated at +506 mV (SCE) and (c) incoloy 800 passivated at +306 mV (SCE).
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in the air grown oxide both iron and chromium were oxidised. In comparison of the three alloys, Fe/O ratio for AISI 304 and AISI 316 were almost same but Cr/ O ratio increased suggesting that AISI 316 is a better alloy for passivation. The change in the Cr/O content may be dominantly due to the change in the chromium content in the analyzing area. In case of incoloy the change may be logical because of more chromium content in the alloy itself, but in AISI 316 the chromium increase might have occurred due to the presence of Mo in the alloy. Because of the dissolution of Mo, the dissolution of chromium was decreased consequently increasing the ratio. 4. Conclusions 1. The phenomenon of chromium enrichment was observed in all the three alloys during initial stage of passivation and the Cr3+ oxidation state was found to be the dominating oxidation process. 2. The oxygen containing species from solution initially attacks chromium of the alloy in contact with the solution and forms hydroxides at the ®lm solution interface and successively oxide below the interface probably by place exchange mechanism. The nucleation of the oxide ®lm is occurred by the segregation of Cr3+ species leaving unoxidised iron and Ni in the neighbourhood of the Cr2O3. 3. Normally grown oxide ®lm showed both chromium and iron in the oxidised state, where as the passive ®lm showed only chromium was oxidised dominantly.
References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17]
P. Marcus, I. Olefjord, Corros. Sci. 28 (1988) 589. H.W. Hoppe, S. Haupt, H.H. Strehblow, Surf. Interface Anal. 21 (1994) 514. N.E. Hakiki, S. Boudin, B. Rondot, M. DA Cunha Belo, Corros. Sci. 37 (1995) 1809. G. Okammoto, T. Shibata, Corros. Sci. 10 (1970) 371. M.J. Graham, Corros. Sci. 37 (1995) 1377. N. Sato, M. Cohen, J. Electrochem. Soc. 111 (1964) 512. G.T. Burstein, P.I. Marshall, Corros. Sci. 23 (1983) 125. G.T. Burstein, R.C. Newman, Electrochem. Acta 25 (1980) 1009. K. Asami, K. Hosimoto, S. Shimodaria, Corros. Sci. 18 (1978) 151. J.A. Bardwell, B. MacDougall, M.J. Graham, J. Electrochem. Soc. 135 (1988) 413. D.F. Mitchell, G.I. Proule, M.J. Graham, Appl. Surf. Sci. 21 (1985) 199. D.F. Mitchell, M.J. Graham, J. Electrochem. Soc. 133 (1986) 936. R. Goetz, D. Landot, Electrochim. Acta 27 (1982) 1061. R. Goetz, D. Landot, Electrochim. Acta 29 (1984) 667. J.E. Castle, J.H. Qiu, J. Electrochem. Soc. 137 (1990) 2033. J.E. Castle, Ruoru Ke, J.F. Natts, Corros. Sci. 30 (1990) 771. C.-O.A. Olsson, Surf. Interface Anal. 21 (1994) 846.
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S. Bera et al. / Corrosion Science 42 (2000) 1709±1724
[18] G. Lorang, J.L. Xu, J.-P. Langeron, Surf. Interface Anal. 19 (1992) 60. [19] J.E. Castle, J.H. Qiu, Corros. Sci. 29 (1989) 591. [20] D.K. Sarkar, S. Bera, S. Dhara, K.G.M. Nair, S.V. Narasimhan, S. Chowdhury, Appl. Surf. Sci. 120 (1997) 159. [21] C.J. Powell, Surf. Sci. 44 (1974) 29. [22] D.B. Briggs, M.P. Seah, Practical Surface Analysis by Auger and XPS, New York, 1987, p. 187. [23] K. Asami, K. Hashimoto, Corros. Sci. 17 (1977) 559. [24] I. Grohmann, E. Kemnitz, A. Lippitz, W.E.S. Unger, Surf. Interface Anal. 23 (1995) 887.