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Stress corrosion cracking of noble metals and their alloys in solutions containing cations of the noble metal: Review of observations relevant to competing models of SCC
Corrosion Science 50 (2008) 1807–1810
Contents lists available at ScienceDirect
Corrosion Science journal homepage: www.elsevier.com/locate/corsci
Letter
Stress corrosion cracking of noble metals and their alloys in solutions containing cations of the noble metal: Review of observations relevant to competing models of SCC Roger C. Newman * University of Toronto, Department of Chemical Engineering and Applied Chemistry, 200 College Street, Toronto, Ontario, Canada M5S 3E5
a r t i c l e
i n f o
Available online 19 June 2008
Keywords: Stress corrosion cracking Copper alloy Noble metal Surface mobility
a b s t r a c t The stress corrosion cracking of Ag, Au and Cu-base alloys, and of pure Ag and Cu, has been studied by J.R. Galvele and others. These authors used solutions that contained the cation of the more-noble metal, so that the tested specimen was at or close to its equilibrium potential in the given solution. The opportunity is taken to review the history of this far-from-new observation and some of its implications. The role of the exchange current density in such cracking is discussed. Observations of Sieradzki and Torchio are used to suggest that in alloys such as brass, SCC is favoured by low, not high, surface mobility, in line with the film-induced cleavage model, which requires very fine nanoporosity at the crack tip – such a favourable condition can only be achieved if dealloying is either very fast or occurs under conditions of low surface mobility. Observations of very slow intergranular SCC in pure metals under dynamic loading are interesting, but not really suggestive of mechanistic continuity with the dramatic mixed-mode cracking that occurs under static loading in brass or AuAg alloys. Torchio’s observations on brass U-bends in CuSO4 solutions of various pH and Cu2+ concentrations are particularly hard to interpret using the surface mobility model. Ó 2008 Elsevier Ltd. All rights reserved.
Studies of SCC in solutions containing noble metal cations, by the Galvele group In 1994, Montoto et al. [1] reported stress corrosion cracking (SCC) in slow strain rate tests of polycrystalline Ag–30Cd in solutions containing AgNO3.1 A number of other studies were done by the same group on the Ag–Cd system, on a-brass in solutions containing Cu(NO3)2, and on ‘22 carat’ gold in solutions with Au3+ [2–7]. Generally a wire of the more-noble metal was used as reference electrode and the tests were done potentiostatically at 0 V, i.e. at the nominal equilibrium potential – disregarding complexities that may arise because of the formation of Cu I species. Recently pure Ag and Cu were shown to crack, though very slowly, under similar conditions [6]; in this case the tests were done at open circuit but the tensile sample was coupled to a longer piece of the same metal. On a log(crack velocity) scale, it was claimed that there was kinetic and therefore mechanistic continuity between the behaviour of alloys and pure metals, at least for these severe testing conditions. The nature of the crack path, especially at high contents of noble metal and in pure metals, is not always clear from the published discussion and micrographs.
* Tel.: +1 416 946 0604; fax: +1 416 978 8605. E-mail address: [email protected] 1 All alloy compositions in this paper are in atomic percent. 0010-938X/$ - see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2008.06.011
All the above observations were interpreted using the surface mobility model of SCC. For solutions containing the noble metal cation [1–7] this is based on an increase in surface diffusivity and crack velocity, caused by the exchange current density. The crack tip captures surface vacancies, thus growing the crack. The only role of dealloying, where this is possible, is to increase the number of surface vacancies – nanoporosity plays no role at all. Plots of log crack velocity (average, from slow strain rate tests) against log cation concentration were used to support this interpretation. Different slopes were found in Cu-base and Ag-base systems. The crack velocity and the exchange current density were claimed to be in exact proportion.2 Unfortunately, the only external literature cited by Montoto et al. was early work that had used ammoniacal Cu II solutions [8,9]. This is quite a different situation: in ammonia solution Cu II is an oxidant, because Cu I (aq) is strongly stabilized by complexation. For Ag in AgNO3 solution there should be a true equilibrium, at least for the unstrained surface (this is not so clear for Cu in Cu(NO3)2 solution, as discussed later).
2 The argument used for the copper system [2] was quite convoluted, involving Cu I intermediates and published kinetic data obtained in a different solution.
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R.C. Newman / Corrosion Science 50 (2008) 1807–1810
Early precedents for the cation-equilibration effect on dealloying and SCC The occurrence of SCC under conditions of equilibrium with respect to the more-noble metal was studied quite intensively in the 1980s, and observations similar to those of Giordano et al. [2] were made in the 1970s. In 1987, Sieradzki et al. [10] published the results of a study of SCC and bare-surface electrochemistry of CuZn and CuAl monocrystals, using ammoniacal Cu I (not Cu II) solutions, with all testing performed at the equilibrium potential of Cu. For the particular strain rate and other conditions used, both dealloying and transgranular SCC ceased below 20 at.% Zn and 16 at.% Al. This was claimed to reflect a causal connection between dealloying and SCC in these alloy systems, via film-induced cleavage (intergranular SCC may persist below these limits – certainly it persists in similar solutions with Cu II [11], where it replaces transgranular SCC below 20% Zn). The parting limits were discussed in the context of the site percolation threshold for the fcc lattice, which is 19.9%. The specific role of Cu Iaq was not discussed in detail in the 1987 paper by Sieradzki et al., but around that time the idea was developed that aqueous cations of the more-noble alloy component can promote dealloying at low contents (20–50%) of less-noble metal by exchange with the surface [12]. Single-shot cleavage of precorroded brass foils was displayed in the ‘equilibrium’ condition [13]. In another paper, responding to criticisms that there may have been a slight oxidizing effect of trace Cu II cations, SCC propagation was demonstrated even at considerable (up to 40 mV) cathodic polarization with respect to the Cu equilibrium potential [14]. All this work was inspired by a paper from Bertocci et al. [15], who were the first to show SCC of brass in ammoniacal Cu I solution, and considered this a key observation suggesting a role of dealloying in SCC. Prior to the work of Sieradzki et al. [10], there were well-documented cases of SCC of brass in CuSO4 solutions [20–22]. The solutions were not de-aerated, but that would not invalidate the results in the present context – Giordano et al. [2] reported anodic current densities in the range 1–10 mA/cm2 for their tests conducted at 0 V vs. Cu, which is much greater than the oxygen limiting current density in aerated solution. So if they had done their tests in aerated solution with no potential control, the potential would still have been lower than the one they used. They stated that their samples became brown during testing at 0 V (Cu), and ascribed the high currents to dezincification – this may be true in part, but it is also possible they were forming Cu2O as discussed later.
expectations of the surface mobility model. We can even add ammonium ions to an acidic chloride solution and cause a transition from macro-dealloying with no SCC to nanoscale dealloying with the possibility of SCC [16]. There is not really a contradiction between saying that Ag or Cu or Au ions increase surface mobility and saying that we need to maintain low surface mobility to have SCC of brass, or Ag–30Cd, or Au–30Ag. As recent simulations have shown [17], ‘surface mobility’ is not quite an adequate description of the property that is altered by introducing the exchange of noble metal atoms at the surface. Rather, the dissolution of less-noble metal atoms from critical sites is facilitated by the exchange of the more-noble metal atoms. It was shown that the usual 55% parting limit occurs because below this concentration of less-noble metal, dealloying cannot occur without dissolving atoms from sites with more than 9 nearest-neighbours. By exchanging the more-noble atoms, we enable such sites, at least fleetingly, to reduce their coordination, allowing their dissolution and consequent porosity formation. This bypasses the usual parting limit and gives us a 20% rather than 55% dealloying threshold. There does not have to be much lateral mobility of the more-noble metal, and indeed too much mobility will tend to suppress SCC because it grows the porosity and thus suppresses cleavage nucleation. Several publications [13,18,19] have shown that the ability of dealloyed layers to inject a substrate crack is reversible, and that this correlates with coarsening of porosity by surface diffusion. Addition of a substance (pyridine) that suppressed surface mobility prolonged the ability to inject a crack [19]. The film-induced cleavage model would not necessarily predict a linear relation between crack velocity and exchange current density, but clearly it predicts a positive correlation between them: the higher the exchange current density, the more frequently paths for dissolution of less-noble metal are created in competition with smoothening of the surface by surface diffusion. This is a complex issue, because tests done at equilibrium with different concentrations of noble metal cation are being done at different absolute potentials, and absolute potential may affect SCC velocity in its own right. The fact that AgCd alloys show a slope d(log v)/d(log i0) of about 0.5, while CuZn alloys show a slope of about 0.25, could simply reflect the fact that d(E0)/d(log [Mn+]) is 60 mV and 30 mV, respectively! In their paper dealing with dealloying of brass in chloride, Newman et al. [12] addressed this issue by varying both [Cu Iaq] and [Cl ] to obtain a range of absolute potentials. They concluded that whilst the absolute potential did have some effect, the main influence came from [Cu Iaq] per se. This, at least, is in agreement with the picture suggested by the Galvele group.
Does low or high surface mobility favour SCC? What is meant by ‘surface mobility’ in the context of the exchange current density?
SCC of pure metals, and of alloys with more than 80% of noble metal
Within the framework of the surface mobility model, it is not at all clear why ammonia should cause such easy and rapid SCC in brass. Too many observations seem to be in direct contradiction with the model. On the other hand, the film-induced cleavage model appears qualitatively consistent with experimental data, in that it considers that there can be too much surface mobility: SCC is favoured by very fine nanoporosity such as that found in ammonia (or, alternatively, such rapid dealloying kinetics that there is no time for coarsening of the porosity). It was proposed [12] that the beneficial effect of alloyed arsenic on macro-dezincification in chloride was due to a reduction in the surface mobility of Cu, and indeed arsenical brass shows much easier SCC than ordinary brass in tests conducted at the Cu/Cu Iaq/CuCl ternary equilibrium condition in cuprous chloride solutions [16]. So the behaviour of brass as we vary its surface mobility is exactly contrary to the
Whilst the observations of Farina et al. [6] on SCC of pure Cu and Ag are novel in that they were conducted exactly at the equilibrium potential (at least for Ag, and assuming that the solutions were de-aerated; this is not mentioned), SCC of pure Cu and Ag was known already under conditions of slow anodic dissolution [23,24]. Much earlier, Pugh had discovered intergranular SCC of pure Cu in non-tarnishing ammoniacal Cu II solutions, which amounts to the same thing [11]. Surface mobility may be responsible for SCC of pure metals in the equilibrated state, but one can think of a host of other explanations. Ag and Cu have high exchange current densities in the AgNO3 and Cu(NO3)2 solutions where Farina et al. [6] reported SCC. Under dynamic straining conditions, one can expect a significant depression of the local equilibrium potential at the crack tip, due to the elevated population of low-coordination metal atoms.
R.C. Newman / Corrosion Science 50 (2008) 1807–1810
If one were to postulate an anodic dissolution mechanism of SCC, a crack velocity of ca. 10 10–10 9 m/s [6] requires an anodic current density at the crack tip of only ca. 0.1–1 mA/cm2. It is not at all clear that anodic dissolution can be ruled out as a mechanism of SCC in such circumstances, even without consideration of grainboundary impurity segregation. Transgranular SCC would be more convincing, in that it is hard to envisage this being caused by anodic dissolution, but although Cu SCC appears to be mixed-mode [6], we face the difficulty that testing pure Cu in Cu(NO3)2 solution is not really an ‘equilibrium’ experiment. According to Sieradzki and Kim [23] SCC of pure Cu is associated with etch-pitting and porosity. The experiments of Farina et al. [6] used a large area of pure unstrained Cu or Ag, immersed and coupled to the slow strain rate specimen. The authors were trying, without using a potentiostat, to maintain the usual metal/metal-ion equilibrium potential. It is unfortunate that the strained and unstrained metals were not coupled through a zero-resistance ammeter (or that the test was not done potentiostatically at 0 V), as this would have given useful additional information relevant to the possibility of an anodic dissolution mechanism of SCC. Again, one has to distinguish carefully between intergranular and transgranular cracking when considering the SCC mechanism in pure metals or in alloys very rich in the more-noble metal, such as Cu–10Zn or Ag–10Cd. Sieradzki et al. [10], under their particular conditions of strain rate etc., found no SCC at all below 20% Zn in brass monocrystals: the correlation with dealloying was precise in this system and in Cu–Al. It is possible, then, that there was an unreported fracture-mode transition between 10 and 20% Zn in the work of Giordano et al. [2], as would be suggested by the work of Pugh using ammoniacal solutions [11]. On the other hand, it is also possible that the strain rate used by Sieradzki et al. (7 10 6 cm s 1 extension rate on a jeweller’s saw notch) was too high to see the slow cracking reported by Giordano et al. for Cu–10Zn (1 10 10–4 10 10 m/s). The slowest crack velocity reported by Sieradzki et al. was 5 10 8 m/s (of course crack velocities are generally higher in the ammonia than the nitrate system).
Contradictions between the observations of Torchio and the surface mobility mechanism In 2007, Farina et al. [7] cited a previous study of SCC by Torchio [22] in near-equilibrium conditions – for brass in naturally aerated, aqueous CuSO4 at various pH values. But in fact, Torchio did not think the work dealt with equilibrium at all, referring at length to the importance of Cu2O formation. This encouraged us to look a little more closely at the behaviour of pure copper in CuSO4 solutions of various pH. Indeed, for de-aerated solutions without deliberate acidification, the potential of Cu is several tens of mV higher than the Cu2+/Cu equilibrium value [25], and there is visible formation of Cu2O. Stirring with copper turnings will almost eliminate the blue colour of the solution under such conditions, but not at a low pH such as 2.0. In other words, copper corrodes to Cu2O in unacidified, de-aerated CuSO4 solutions – and probably in the Cu(NO3)2 solutions used by Giordano et al. [2]. Torchio [22] found dramatic effects of pH on the SCC behaviour of brass in CuSO4 solutions – Fig. 1. Importantly, he used U-bends and constant-load tensile testing, and thus avoided the distorting effect of slow strain rate testing at low levels of susceptibility. At low pH where the conditions would be close to equilibrium for Cu2+/Cu (bearing in mind that oxygen was present and that Cu I also has a certain low equilibrium solubility), slow intergranular cracking of cold-worked U-bends was found. But as the pH was increased (and the surface mobility was decreased – by formation of Cu2O), dramatic and rapid transgranular SCC was found for pH
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Fig. 1. Occurrence of SCC in 240-h U-bend tests of cold-worked admiralty brass, using stagnant 0.1 M Na2SO4 plus H2SO4 plus CuSO4 solutions of the indicated [initial] pH and Cu2+ concentrations, by Torchio [22]. For 10 1 M initial concentration of CuSO4, the crack depths in microns after 240 h are indicated. ‘F’ signifies that the 600 lm-thick U-bend samples failed completely.
values greater than about 3, in line with the observations of Pinchback et al. [20] and Giordano et al. [2] in solutions without deliberate acidification. This is the pH threshold where Cu2O starts to form in the relevant range of [Cu2+] [26]. For constant-load specimens tested at a stress of 81 MPa, the slow intergranular SCC persisted but the dramatic transgranular SCC did not occur, reflecting the low value of stress. Torchio gave a complicated rationalization, involving the presence or absence of stirring or solution refreshment, but we cannot see how that alters the basic picture presented in Fig. 1. Once again, the literature tends to contradict the predictions of the surface mobility model: SCC of brass, especially under static load, is favoured by conditions of low surface mobility – in this case nearneutral conditions allowing the formation of Cu2O, rather than acidic pH where the Cu is unoxidized and free to diffuse on the surface. So low surface mobility can arise in at least three ways: presence of ammonia, alloying with arsenic, or formation of Cu2O. At least one of these factors is required for film-induced cleavage to occur. In such a context, it hardly matters, from a mechanistic point of view, whether or not pure metals or Cu–10Zn alloys can be made to crack – slowly and intergranularly – under extreme conditions. This must be another kind of mechanism, and surface mobility may be as good an explanation as any – we prefer some kind of anodic dissolution. Acknowledgements This research was funded by an NSERC Senior Industrial Research Chair in Nuclear Engineering awarded to R.C. Newman in association with UNENE, the University Network of Excellence in Nuclear Engineering. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]
M.L. Montoto, G.S. Duffo, J.R. Galvele, Corros. Sci. 36 (1994) 1805–1808. C.M. Giordano, G.S. Duffo, J.R. Galvele, Corros. Sci. 39 (1997) 915–1923. J.R. Galvele, G.S. Duffo, Corros. Sci. 39 (1997) 605–608. M.L. Montoto, G.S. Duffo, J.R. Galvele, Corros. Sci. 43 (2001) 755–764. G.S. Duffo, S.B. Farina, J.R. Galvele, Corros. Sci. 46 (2004) 1–4. S.B. Farina, G.S. Duffo, J.R. Galvele, Corros. Sci. 47 (2005) 239–245. S.B. Farina, G.S. Duffo, J.R. Galvele, Corros. Sci. 49 (2007) 687–1695. E.N. Pugh, A.R.C. Westwood, Philos. Mag. 13 (1966) 167–183. H.H. Uhlig, K. Gupta, W. Liang, J. Electrochem. Soc. 122 (1975) 343–350. K. Sieradzki, J.S. Kim, A.T. Cole, R.C. Newman, J. Electrochem. Soc. 134 (1987) 1635–1639.
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[11] E.N. Pugh, Fundamental Aspects of Stress Corrosion Cracking, NACE, Houston, 1969, pp. 119–158. [12] R.C. Newman, T. Shahrabi, K. Sieradzki, Corros. Sci. 28 (1988) 873–886. [13] R.C. Newman, T. Shahrabi, K. Sieradzki, Scripta Metall. 23 (1989) 71–74. [14] T. Shahrabi, R.C. Newman, K. Sieradzki, J. Electrochem. Soc. 140 (1993) 348– 352. [15] U. Bertocci, F.I. Thomas, E.N. Pugh, Corrosion 40 (1984) 439–440. [16] T. Shahrabi, R.C. Newman, Dealloying and stress-corrosion cracking of copper alloys in Cu I solutions, Mater. Sci. Forum 44–45 (1989) 169–176. [17] D.M. Artymowicz, R.C. Newman, J. Erlebacher, Insights into the parting limit for de-alloying from revision of atomistic considerations, in: ECS Transactions Volume, Critical Factors in Localized Corrosion, a Symposium in Honor of Hugh S. Isaacs, in press.
[18] R.G. Kelly, T. Shahrabi, A.J. Frost, R.C. Newman, Metall. Trans. A 22A (1991) 191–197. [19] M. Saito, G.S. Smith, R.C. Newman, Corros. Sci. 35 (1993) 411. [20] T.R. Pinchback, S.P. Clough, L.A. Heldt, Corrosion 32 (1976) 369–371. [21] W.K. Blanchard, D.A. Koss, L.A. Heldt, Metall. Trans. A 15A (1984) 1281– 1286. [22] S. Torchio, Corros. Sci. 26 (1986) 133–151. [23] K. Sieradzki, J.S. Kim, Acta Metall. 40 (1992) 625–635. [24] S.G. Corcoran, K. Sieradzki, Scripta Metall. 26 (1992) 633–637. [25] W. Zhang, R.C. Newman, Unpublished Research, University of Toronto. [26] M. Pourbaix, Atlas of Electrochemical Equilibria in Aqueous Solutions, NACE, Houston, 1974.
Contents lists available at ScienceDirect
Corrosion Science journal homepage: www.elsevier.com/locate/corsci
Letter
Stress corrosion cracking of noble metals and their alloys in solutions containing cations of the noble metal: Review of observations relevant to competing models of SCC Roger C. Newman * University of Toronto, Department of Chemical Engineering and Applied Chemistry, 200 College Street, Toronto, Ontario, Canada M5S 3E5
a r t i c l e
i n f o
Available online 19 June 2008
Keywords: Stress corrosion cracking Copper alloy Noble metal Surface mobility
a b s t r a c t The stress corrosion cracking of Ag, Au and Cu-base alloys, and of pure Ag and Cu, has been studied by J.R. Galvele and others. These authors used solutions that contained the cation of the more-noble metal, so that the tested specimen was at or close to its equilibrium potential in the given solution. The opportunity is taken to review the history of this far-from-new observation and some of its implications. The role of the exchange current density in such cracking is discussed. Observations of Sieradzki and Torchio are used to suggest that in alloys such as brass, SCC is favoured by low, not high, surface mobility, in line with the film-induced cleavage model, which requires very fine nanoporosity at the crack tip – such a favourable condition can only be achieved if dealloying is either very fast or occurs under conditions of low surface mobility. Observations of very slow intergranular SCC in pure metals under dynamic loading are interesting, but not really suggestive of mechanistic continuity with the dramatic mixed-mode cracking that occurs under static loading in brass or AuAg alloys. Torchio’s observations on brass U-bends in CuSO4 solutions of various pH and Cu2+ concentrations are particularly hard to interpret using the surface mobility model. Ó 2008 Elsevier Ltd. All rights reserved.
Studies of SCC in solutions containing noble metal cations, by the Galvele group In 1994, Montoto et al. [1] reported stress corrosion cracking (SCC) in slow strain rate tests of polycrystalline Ag–30Cd in solutions containing AgNO3.1 A number of other studies were done by the same group on the Ag–Cd system, on a-brass in solutions containing Cu(NO3)2, and on ‘22 carat’ gold in solutions with Au3+ [2–7]. Generally a wire of the more-noble metal was used as reference electrode and the tests were done potentiostatically at 0 V, i.e. at the nominal equilibrium potential – disregarding complexities that may arise because of the formation of Cu I species. Recently pure Ag and Cu were shown to crack, though very slowly, under similar conditions [6]; in this case the tests were done at open circuit but the tensile sample was coupled to a longer piece of the same metal. On a log(crack velocity) scale, it was claimed that there was kinetic and therefore mechanistic continuity between the behaviour of alloys and pure metals, at least for these severe testing conditions. The nature of the crack path, especially at high contents of noble metal and in pure metals, is not always clear from the published discussion and micrographs.
* Tel.: +1 416 946 0604; fax: +1 416 978 8605. E-mail address: [email protected] 1 All alloy compositions in this paper are in atomic percent. 0010-938X/$ - see front matter Ó 2008 Elsevier Ltd. All rights reserved. doi:10.1016/j.corsci.2008.06.011
All the above observations were interpreted using the surface mobility model of SCC. For solutions containing the noble metal cation [1–7] this is based on an increase in surface diffusivity and crack velocity, caused by the exchange current density. The crack tip captures surface vacancies, thus growing the crack. The only role of dealloying, where this is possible, is to increase the number of surface vacancies – nanoporosity plays no role at all. Plots of log crack velocity (average, from slow strain rate tests) against log cation concentration were used to support this interpretation. Different slopes were found in Cu-base and Ag-base systems. The crack velocity and the exchange current density were claimed to be in exact proportion.2 Unfortunately, the only external literature cited by Montoto et al. was early work that had used ammoniacal Cu II solutions [8,9]. This is quite a different situation: in ammonia solution Cu II is an oxidant, because Cu I (aq) is strongly stabilized by complexation. For Ag in AgNO3 solution there should be a true equilibrium, at least for the unstrained surface (this is not so clear for Cu in Cu(NO3)2 solution, as discussed later).
2 The argument used for the copper system [2] was quite convoluted, involving Cu I intermediates and published kinetic data obtained in a different solution.
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R.C. Newman / Corrosion Science 50 (2008) 1807–1810
Early precedents for the cation-equilibration effect on dealloying and SCC The occurrence of SCC under conditions of equilibrium with respect to the more-noble metal was studied quite intensively in the 1980s, and observations similar to those of Giordano et al. [2] were made in the 1970s. In 1987, Sieradzki et al. [10] published the results of a study of SCC and bare-surface electrochemistry of CuZn and CuAl monocrystals, using ammoniacal Cu I (not Cu II) solutions, with all testing performed at the equilibrium potential of Cu. For the particular strain rate and other conditions used, both dealloying and transgranular SCC ceased below 20 at.% Zn and 16 at.% Al. This was claimed to reflect a causal connection between dealloying and SCC in these alloy systems, via film-induced cleavage (intergranular SCC may persist below these limits – certainly it persists in similar solutions with Cu II [11], where it replaces transgranular SCC below 20% Zn). The parting limits were discussed in the context of the site percolation threshold for the fcc lattice, which is 19.9%. The specific role of Cu Iaq was not discussed in detail in the 1987 paper by Sieradzki et al., but around that time the idea was developed that aqueous cations of the more-noble alloy component can promote dealloying at low contents (20–50%) of less-noble metal by exchange with the surface [12]. Single-shot cleavage of precorroded brass foils was displayed in the ‘equilibrium’ condition [13]. In another paper, responding to criticisms that there may have been a slight oxidizing effect of trace Cu II cations, SCC propagation was demonstrated even at considerable (up to 40 mV) cathodic polarization with respect to the Cu equilibrium potential [14]. All this work was inspired by a paper from Bertocci et al. [15], who were the first to show SCC of brass in ammoniacal Cu I solution, and considered this a key observation suggesting a role of dealloying in SCC. Prior to the work of Sieradzki et al. [10], there were well-documented cases of SCC of brass in CuSO4 solutions [20–22]. The solutions were not de-aerated, but that would not invalidate the results in the present context – Giordano et al. [2] reported anodic current densities in the range 1–10 mA/cm2 for their tests conducted at 0 V vs. Cu, which is much greater than the oxygen limiting current density in aerated solution. So if they had done their tests in aerated solution with no potential control, the potential would still have been lower than the one they used. They stated that their samples became brown during testing at 0 V (Cu), and ascribed the high currents to dezincification – this may be true in part, but it is also possible they were forming Cu2O as discussed later.
expectations of the surface mobility model. We can even add ammonium ions to an acidic chloride solution and cause a transition from macro-dealloying with no SCC to nanoscale dealloying with the possibility of SCC [16]. There is not really a contradiction between saying that Ag or Cu or Au ions increase surface mobility and saying that we need to maintain low surface mobility to have SCC of brass, or Ag–30Cd, or Au–30Ag. As recent simulations have shown [17], ‘surface mobility’ is not quite an adequate description of the property that is altered by introducing the exchange of noble metal atoms at the surface. Rather, the dissolution of less-noble metal atoms from critical sites is facilitated by the exchange of the more-noble metal atoms. It was shown that the usual 55% parting limit occurs because below this concentration of less-noble metal, dealloying cannot occur without dissolving atoms from sites with more than 9 nearest-neighbours. By exchanging the more-noble atoms, we enable such sites, at least fleetingly, to reduce their coordination, allowing their dissolution and consequent porosity formation. This bypasses the usual parting limit and gives us a 20% rather than 55% dealloying threshold. There does not have to be much lateral mobility of the more-noble metal, and indeed too much mobility will tend to suppress SCC because it grows the porosity and thus suppresses cleavage nucleation. Several publications [13,18,19] have shown that the ability of dealloyed layers to inject a substrate crack is reversible, and that this correlates with coarsening of porosity by surface diffusion. Addition of a substance (pyridine) that suppressed surface mobility prolonged the ability to inject a crack [19]. The film-induced cleavage model would not necessarily predict a linear relation between crack velocity and exchange current density, but clearly it predicts a positive correlation between them: the higher the exchange current density, the more frequently paths for dissolution of less-noble metal are created in competition with smoothening of the surface by surface diffusion. This is a complex issue, because tests done at equilibrium with different concentrations of noble metal cation are being done at different absolute potentials, and absolute potential may affect SCC velocity in its own right. The fact that AgCd alloys show a slope d(log v)/d(log i0) of about 0.5, while CuZn alloys show a slope of about 0.25, could simply reflect the fact that d(E0)/d(log [Mn+]) is 60 mV and 30 mV, respectively! In their paper dealing with dealloying of brass in chloride, Newman et al. [12] addressed this issue by varying both [Cu Iaq] and [Cl ] to obtain a range of absolute potentials. They concluded that whilst the absolute potential did have some effect, the main influence came from [Cu Iaq] per se. This, at least, is in agreement with the picture suggested by the Galvele group.
Does low or high surface mobility favour SCC? What is meant by ‘surface mobility’ in the context of the exchange current density?
SCC of pure metals, and of alloys with more than 80% of noble metal
Within the framework of the surface mobility model, it is not at all clear why ammonia should cause such easy and rapid SCC in brass. Too many observations seem to be in direct contradiction with the model. On the other hand, the film-induced cleavage model appears qualitatively consistent with experimental data, in that it considers that there can be too much surface mobility: SCC is favoured by very fine nanoporosity such as that found in ammonia (or, alternatively, such rapid dealloying kinetics that there is no time for coarsening of the porosity). It was proposed [12] that the beneficial effect of alloyed arsenic on macro-dezincification in chloride was due to a reduction in the surface mobility of Cu, and indeed arsenical brass shows much easier SCC than ordinary brass in tests conducted at the Cu/Cu Iaq/CuCl ternary equilibrium condition in cuprous chloride solutions [16]. So the behaviour of brass as we vary its surface mobility is exactly contrary to the
Whilst the observations of Farina et al. [6] on SCC of pure Cu and Ag are novel in that they were conducted exactly at the equilibrium potential (at least for Ag, and assuming that the solutions were de-aerated; this is not mentioned), SCC of pure Cu and Ag was known already under conditions of slow anodic dissolution [23,24]. Much earlier, Pugh had discovered intergranular SCC of pure Cu in non-tarnishing ammoniacal Cu II solutions, which amounts to the same thing [11]. Surface mobility may be responsible for SCC of pure metals in the equilibrated state, but one can think of a host of other explanations. Ag and Cu have high exchange current densities in the AgNO3 and Cu(NO3)2 solutions where Farina et al. [6] reported SCC. Under dynamic straining conditions, one can expect a significant depression of the local equilibrium potential at the crack tip, due to the elevated population of low-coordination metal atoms.
R.C. Newman / Corrosion Science 50 (2008) 1807–1810
If one were to postulate an anodic dissolution mechanism of SCC, a crack velocity of ca. 10 10–10 9 m/s [6] requires an anodic current density at the crack tip of only ca. 0.1–1 mA/cm2. It is not at all clear that anodic dissolution can be ruled out as a mechanism of SCC in such circumstances, even without consideration of grainboundary impurity segregation. Transgranular SCC would be more convincing, in that it is hard to envisage this being caused by anodic dissolution, but although Cu SCC appears to be mixed-mode [6], we face the difficulty that testing pure Cu in Cu(NO3)2 solution is not really an ‘equilibrium’ experiment. According to Sieradzki and Kim [23] SCC of pure Cu is associated with etch-pitting and porosity. The experiments of Farina et al. [6] used a large area of pure unstrained Cu or Ag, immersed and coupled to the slow strain rate specimen. The authors were trying, without using a potentiostat, to maintain the usual metal/metal-ion equilibrium potential. It is unfortunate that the strained and unstrained metals were not coupled through a zero-resistance ammeter (or that the test was not done potentiostatically at 0 V), as this would have given useful additional information relevant to the possibility of an anodic dissolution mechanism of SCC. Again, one has to distinguish carefully between intergranular and transgranular cracking when considering the SCC mechanism in pure metals or in alloys very rich in the more-noble metal, such as Cu–10Zn or Ag–10Cd. Sieradzki et al. [10], under their particular conditions of strain rate etc., found no SCC at all below 20% Zn in brass monocrystals: the correlation with dealloying was precise in this system and in Cu–Al. It is possible, then, that there was an unreported fracture-mode transition between 10 and 20% Zn in the work of Giordano et al. [2], as would be suggested by the work of Pugh using ammoniacal solutions [11]. On the other hand, it is also possible that the strain rate used by Sieradzki et al. (7 10 6 cm s 1 extension rate on a jeweller’s saw notch) was too high to see the slow cracking reported by Giordano et al. for Cu–10Zn (1 10 10–4 10 10 m/s). The slowest crack velocity reported by Sieradzki et al. was 5 10 8 m/s (of course crack velocities are generally higher in the ammonia than the nitrate system).
Contradictions between the observations of Torchio and the surface mobility mechanism In 2007, Farina et al. [7] cited a previous study of SCC by Torchio [22] in near-equilibrium conditions – for brass in naturally aerated, aqueous CuSO4 at various pH values. But in fact, Torchio did not think the work dealt with equilibrium at all, referring at length to the importance of Cu2O formation. This encouraged us to look a little more closely at the behaviour of pure copper in CuSO4 solutions of various pH. Indeed, for de-aerated solutions without deliberate acidification, the potential of Cu is several tens of mV higher than the Cu2+/Cu equilibrium value [25], and there is visible formation of Cu2O. Stirring with copper turnings will almost eliminate the blue colour of the solution under such conditions, but not at a low pH such as 2.0. In other words, copper corrodes to Cu2O in unacidified, de-aerated CuSO4 solutions – and probably in the Cu(NO3)2 solutions used by Giordano et al. [2]. Torchio [22] found dramatic effects of pH on the SCC behaviour of brass in CuSO4 solutions – Fig. 1. Importantly, he used U-bends and constant-load tensile testing, and thus avoided the distorting effect of slow strain rate testing at low levels of susceptibility. At low pH where the conditions would be close to equilibrium for Cu2+/Cu (bearing in mind that oxygen was present and that Cu I also has a certain low equilibrium solubility), slow intergranular cracking of cold-worked U-bends was found. But as the pH was increased (and the surface mobility was decreased – by formation of Cu2O), dramatic and rapid transgranular SCC was found for pH
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Fig. 1. Occurrence of SCC in 240-h U-bend tests of cold-worked admiralty brass, using stagnant 0.1 M Na2SO4 plus H2SO4 plus CuSO4 solutions of the indicated [initial] pH and Cu2+ concentrations, by Torchio [22]. For 10 1 M initial concentration of CuSO4, the crack depths in microns after 240 h are indicated. ‘F’ signifies that the 600 lm-thick U-bend samples failed completely.
values greater than about 3, in line with the observations of Pinchback et al. [20] and Giordano et al. [2] in solutions without deliberate acidification. This is the pH threshold where Cu2O starts to form in the relevant range of [Cu2+] [26]. For constant-load specimens tested at a stress of 81 MPa, the slow intergranular SCC persisted but the dramatic transgranular SCC did not occur, reflecting the low value of stress. Torchio gave a complicated rationalization, involving the presence or absence of stirring or solution refreshment, but we cannot see how that alters the basic picture presented in Fig. 1. Once again, the literature tends to contradict the predictions of the surface mobility model: SCC of brass, especially under static load, is favoured by conditions of low surface mobility – in this case nearneutral conditions allowing the formation of Cu2O, rather than acidic pH where the Cu is unoxidized and free to diffuse on the surface. So low surface mobility can arise in at least three ways: presence of ammonia, alloying with arsenic, or formation of Cu2O. At least one of these factors is required for film-induced cleavage to occur. In such a context, it hardly matters, from a mechanistic point of view, whether or not pure metals or Cu–10Zn alloys can be made to crack – slowly and intergranularly – under extreme conditions. This must be another kind of mechanism, and surface mobility may be as good an explanation as any – we prefer some kind of anodic dissolution. Acknowledgements This research was funded by an NSERC Senior Industrial Research Chair in Nuclear Engineering awarded to R.C. Newman in association with UNENE, the University Network of Excellence in Nuclear Engineering. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10]
M.L. Montoto, G.S. Duffo, J.R. Galvele, Corros. Sci. 36 (1994) 1805–1808. C.M. Giordano, G.S. Duffo, J.R. Galvele, Corros. Sci. 39 (1997) 915–1923. J.R. Galvele, G.S. Duffo, Corros. Sci. 39 (1997) 605–608. M.L. Montoto, G.S. Duffo, J.R. Galvele, Corros. Sci. 43 (2001) 755–764. G.S. Duffo, S.B. Farina, J.R. Galvele, Corros. Sci. 46 (2004) 1–4. S.B. Farina, G.S. Duffo, J.R. Galvele, Corros. Sci. 47 (2005) 239–245. S.B. Farina, G.S. Duffo, J.R. Galvele, Corros. Sci. 49 (2007) 687–1695. E.N. Pugh, A.R.C. Westwood, Philos. Mag. 13 (1966) 167–183. H.H. Uhlig, K. Gupta, W. Liang, J. Electrochem. Soc. 122 (1975) 343–350. K. Sieradzki, J.S. Kim, A.T. Cole, R.C. Newman, J. Electrochem. Soc. 134 (1987) 1635–1639.
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[11] E.N. Pugh, Fundamental Aspects of Stress Corrosion Cracking, NACE, Houston, 1969, pp. 119–158. [12] R.C. Newman, T. Shahrabi, K. Sieradzki, Corros. Sci. 28 (1988) 873–886. [13] R.C. Newman, T. Shahrabi, K. Sieradzki, Scripta Metall. 23 (1989) 71–74. [14] T. Shahrabi, R.C. Newman, K. Sieradzki, J. Electrochem. Soc. 140 (1993) 348– 352. [15] U. Bertocci, F.I. Thomas, E.N. Pugh, Corrosion 40 (1984) 439–440. [16] T. Shahrabi, R.C. Newman, Dealloying and stress-corrosion cracking of copper alloys in Cu I solutions, Mater. Sci. Forum 44–45 (1989) 169–176. [17] D.M. Artymowicz, R.C. Newman, J. Erlebacher, Insights into the parting limit for de-alloying from revision of atomistic considerations, in: ECS Transactions Volume, Critical Factors in Localized Corrosion, a Symposium in Honor of Hugh S. Isaacs, in press.
[18] R.G. Kelly, T. Shahrabi, A.J. Frost, R.C. Newman, Metall. Trans. A 22A (1991) 191–197. [19] M. Saito, G.S. Smith, R.C. Newman, Corros. Sci. 35 (1993) 411. [20] T.R. Pinchback, S.P. Clough, L.A. Heldt, Corrosion 32 (1976) 369–371. [21] W.K. Blanchard, D.A. Koss, L.A. Heldt, Metall. Trans. A 15A (1984) 1281– 1286. [22] S. Torchio, Corros. Sci. 26 (1986) 133–151. [23] K. Sieradzki, J.S. Kim, Acta Metall. 40 (1992) 625–635. [24] S.G. Corcoran, K. Sieradzki, Scripta Metall. 26 (1992) 633–637. [25] W. Zhang, R.C. Newman, Unpublished Research, University of Toronto. [26] M. Pourbaix, Atlas of Electrochemical Equilibria in Aqueous Solutions, NACE, Houston, 1974.