Marcus P. Corrosion mechanisms in theory and practice

Подождите немного. Документ загружается.

of much thicker salt films in larger pits of several tens of μm provides further

stabilization but is not a necessary condition for localized corrosion. The presence

of nonaggressive anions in the bulk electrolyte also caused their accumulation but did

not suppress the additional presence of aggressive anions that is absolutely necessary

for a stable pit growth. The electropolishing conditions in sulfate-containing solution

caused large amounts of these anions corresponding to several hundred nm of

FeSO

within the hemispherical pits but did not avoid the additional and necessary

accumulation of ~ 5 nm FeCl

[66].

Repassivation of Corrosion Pits

The kinetics of repassivation of small pits in an early stage of their development

seem to be related to the transport of aggressive anions, such as the chloride

accumulated locally during the intense dissolution process from the pit to the bulk

[4,30]. If this transport is the rate-determining step, one expects the repassivation

time to increase with the depth of a corrosion pit. If we simply apply the Einstein-

Smoluchowski relation for the transport time t

out of a pit of radius r [Eq. (30)]

and if the radius r is given by the local current density i

and the lifetime t

of the

pit by Eq. (31), we obtain for the repassivation time t

[Eq. (32)]

Mechanisms of Pitting Corrosion 279

The t

increases with the square of the lifetime t

of a pit. The constants k

have

been estimated from D = 5 × 10

–6

/s, zF/V

= 2.72 × 10

A s/cm

for Fe and

2.93 × 10

A s/cm

for Ni to k

= 1.35 × 10

–4

and k

= 1.16 × 10

–4

for Fe

and Ni, respectively. To check the i

and t

dependence of t

, potentiostatic pulse

measurements have been performed. The potential was pulsed for some 10 ms up

to 1 s to values above the pitting potential E

and back to E ≤ E

for increasing

time Δt. When the potential was finally pulsed back to the potential in the pitting

region, the current density started to increase further at the same level reached at

the end of the previous pitting period if Δt was short enough. When Δt was too

long, at least some of the pits repassivated and the subsequent pitting period had

to start again with the nucleation of new pits and a smaller value of the average

current density of the electrode. The interval Δt of a fast deviation from a continuous

increase of the current density for a set of pulse experiments is taken as the

repassivation time t

. A double logarithmic plot of t

and t

shows a slope of

approximately 2 as required by Eq. (32). Also, t

changes proportionally with i

The estimated values agree sufficiently well with the experimental data of Figure 23.

One may therefore conclude that the rate-determining step for the repassivation

is the transport of accumulated aggressive anions out of small pits. This result

coincides well with the explanation that localized corrosion is stabilized at the initial

stage by the accumulation of aggressive anions that prevent the formation of a

passive layer at the active pit surface.

Factors Stabilizing Pit Growth

The factors that stabilize the growth of a pit depend on the stage of its development.

Ohmic drops cannot serve as an explanation for small pits with micrometer

dimensions. The ohmic drops are <1 mV for strongly acidic, well-conducting

electrolytes such as 0.5 M H

and a pit radius of 1 μm. This is far too small to

shift the potential from its potentiostatically fixed value below the passivation

potential by more than 1 V (Fig. 20). However, in later stages for pits of several

millimeters radius the ohmic drop might become high enough to satisfy this

condition. The extremely high local current densities are a further proof of the

absence of large ohmic drops. The local potential at the pit surface is in the vicinity

of the potentiostatically fixed value in order to permit these high dissolution currents.

Ohmic drops have already been corrected in Figure 15, presenting these high current

densities for Ni.

280 Strehblow

Figure 23 Repassivation time t

vs. pitting time t

for Fe (A= 0.13 cm

) and Ni (A= 0.2 cm

)

at different potentials for pitting as indicated, t

leading to deviations of continuous current

increase for one potential sequence. () Fe 0.01 M Cl

–

; (c) Fe 0.01 M Cl

–

, (z) Ni 0.1 M

–

Δ Fe 0.01 M Cl

–

leading to continuous increase of current for many subsequent

potential cycles, phathalate buffer, pH 5.0 (From Ref. 4.)

pH shifts also cannot explain the stable pit growth for metals such as Fe, Ni,

and steels that are also passive in strongly acidic electrolytes. For these systems

passivity is a kinetically determined phenomenon that is even in disagreement

with the Pourbaix diagram. But there are metals with passivation potentials in

close agreement with their Pourbaix diagrams that form passivating oxide layers

which dissolve at a high rate in strongly acidic electrolytes, such as Al and Cu. For

these systems, local destabilization of passivity may occur on acidification of the

pit electrolyte with pitting in neutral or weakly acidic bulk electrolytes. However,

even for these metals there still remains the question why halides are necessary for

pit nucleation and pit growth. The pH dependence of the Flade potential of –0.059

V/pH as obtained for metals such as Fe and Ni cannot serve as a stabilizing factor.

The pH shift required to obtain active conditions at the pit surface is much too

large in comparison with the applied electrode potentials.

The requirement for aggressive anions and their influence on the repassivation

kinetics demonstrate that their accumulation stabilizes localized corrosion in an

early stage. The autocatalytically enhanced dissolution of the passive layer by

these anions prevents any repassivation once a pit has formed. It should be mentioned

again that pure Cr is one of the few metals that is not subject to pitting by chloride.

In agreement with this finding, Cr(III) salts dissolve very slowly in water. Fast

transfer of the complexed ions to the electrolyte is necessary for the attack on the

passive layer, which provides a good explanation for the pitting resistance of this

metal compared with Fe and Ni.

For sufficiently positive potentials the intense dissolution process leads to

severe accumulation of corrosion products and finally, after a transition time, to

the formation of a salt film, which is a further stabilizing factor for a pit. It may

act as a reservoir to maintain a high concentration of the aggressive anions. These

explanations are supported by the logarithmic dependence of the critical pitting

and inhibition potential on the concentration of aggressive (A) or inhibiting (I)

anions according to Eqs. (1) and (2).

These equations with large b factors have been found experimentally for

Fe and Ni and various aggressive and inhibiting anions [2,67]. They have been

deduced on the basis of a thermodynamic theory of adsorption of the anions at

the metal surface of a pit in competition with each other and with the formation

of the passive layer [37,67]. This explanation agrees well with the concept of

repassivation prevented by accumulation of aggressive anions. According to this

approach, a potential-assisted Langmuir type of adsorption isotherm for the

aggressive anions is assumed at the parts of the bare metal surface where the

passive layer has been removed by any of the mechanisms that have been discussed.

For simplicity, homogeneous adsorption at the metal surface is assumed with no

lateral interaction between the adsorbed ions. The accumulation of corrosion

products andconsequently of the aggressive anions due to the extremely high

dissolution rates supports their adsorption and increases their surface coverage

Θ, which also increases with the electrode potential E. Stable pit growth will be

achieved when Θ exceeds a critical value Θ

crit

. It should be large enough to

prevent the formation of a continuous oxide film, i.e., the repassivation of the pit

surface. In the light of this concept, a given composition of the bulk electrolyte

requires a minimum electrode potentials E

to exceed Θ

crit

and to achieve

stable pit growth. Detailed evaluation of this adsorption equilibrium yields

Mechanisms of Pitting Corrosion 281

where ΔG°

ad,A

is the standard Gibbs free energy (free enthalpy) of adsorption of

the aggressive anions A, Δφ

is the potential drop at the standard hydrogen electrode,

and γ is the electrosorption valence, which takes into account the closest approach

of the anions to the metal surface because of their own dimensions.

If inhibiting anions I are present as well, they compete with the aggressive

anions A for the adsorption sites according to the same type of adsorption isotherm

as given in Eq. (33) with their specific ΔG°

values and bulk concentrations.

Combining the equations for the inhibiting and aggressive anions leads to the critical

conditions of inhibition of localized corrosion. The detailed discussion [2] yields

for the inhibition potential E

282 Strehblow

where ΔG°

is the difference in Gibbs free energies of species A and I, Θ

A,ad

and

I,ad

are the surface coverages of both ions, and Δγ is the difference in their

adsorption valences. The negative Δγ value causes a positive shift of E

with

increasing [A] and decreasing [I], which is an upper potential limit of the pitting

range. The small values of Δγ corresponding to small differences for the approach

of ions A and I to the electrode surface cause large b factors for Eq. (34), which is

found experimentally for nitrate and perchlorate [2,58].

Pits may also be stabilized by other complicating effects. In some cases the

pit mouth is covered by precipitated corrosion products such as hydroxides or

even a metal cover when the dissolution occurs irregularly (Fig. 14). In these

situations an ohmic drop can lead to a potential shift of the pit surface below the

passivation potential. The same will hold for large and especially deep and narrow

pits. Crevice corrosion may be seen as a special case of localized corrosion. In

these cases ohmic drops act as stabilizing factors for pitting. Last but not least, the

presence a gas bubble within a pit will cause a situation similar to that of a crevice

with related ohmic drops. Hydrogen bubbles have been observed by Pickering and

Frankenthal [68] within pits and served as an explanation for stable pitting. In this

situation it is still not understood how the potential becomes sufficiently negative

within a small pit despite a potentiostatically fixed positive value to cause

hydrogen evolution. Apparently gas bubbles may be an important stabilizing factor

for some special systems and later stages of pit growth.

A further factor in the stability of a growing pit is alloy components that

prevent or reduce pitting. They usually accumulate within the passive layer due to

their slow dissolution kinetics. The special situation of Cr has already been discussed

in detail. Fe-Si alloys with more than 20 at % Si are another example. A sufficiently

large Si content causes the formation of a continuous SiO

layer especially in acidic

electrolytes [69], which prevents localized corrosion. The dissolution of SiO

is not

enhanced by halides, which apparently is a consequence of the missing complexation

of Si(IV) so that the passive layer remains protective even in the presence of a large

chloride concentration. If a pit starts, a large bulk concentration of Si will lead to Si

accumulation at the pit surface, which then may form a protective Si(IV)-rich film.

This situation will lead to repassivation of the corrosion pit in an early stage as

observed experimentally [70].

CONCLUSION

The discussion of various experimental results has shown that different effects may

serve as stabilizing factors for localized corrosion. It depends on the stage of

development of a corrosion pit and the environmental conditions which one is the

most effective. Therefore one has to refer to the special situation under discussion.

Often the controversial points of view of authors depend on the different situations

that are examined and discussed. It was in the aim of this contribution to sort out

the different factors and to describe the conditions in which they are effective. In this

sense, localized corrosion is a very complicated process involving different effects.

The same statement holds for the nucleation of corrosion pits. Here three main

mechanisms have been discussed. Again, it depends on the experimental or

environmental conditions which one is the most effective and dominates the

nucleation step.

Although there are many common features in localized corrosion of the

different systems, one should also take into account the specific chemistry of a system.

In this sense the tendency of halides to complex with metal cations is very important

and seems to be the key for an understanding of the stabilization of a corrosion pit

by prevention of the repassivation of a defect site within the passive layer.

Enhancement of the transfer of metal cations from the oxide to the electrolyte by

halides, especially the strongly complexing fluoride, holds for many metals. The

special situation of Cr, which is resistant to localized corrosion, may be understood

in terms of the slow dissolution kinetics of the Cr(III) salts. One should also

consider the different passivation mechanisms. Metals such as Al and Cu are not

passive in strongly acidic electrolytes. In these cases localized acidification by the

hydrolysis of corrosion products may well serve as a stabilizing factor for pitting.

Other metals such as Fe, Ni, and steels are passive even in strongly acidic electrolytes,

in disagreement with the predictions of the Pourbaix diagrams. Local acidification

cannot be a stabilizing effect for pitting in these cases. In summary, all explanations

of stable pit growth should include chemical arguments to explain why halides are

so essential for localized corrosion. Their complexing and catalytic properties for

the dissolution of the passive layers should not be ignored.

Besides the possible differences and the special influence of the chemical

properties of the various systems, many common features are found for localized

corrosion. Electrochemical and surface analytical methods for the study of passive

layers and their breakdown have been very successful in the past, and closer insight

into the effective mechanisms has been achieved. STM and SFM (Scanning Force

Microscopy) will give valuable information about the nucleation and very early

stages of growth of corrosion pits. The application of these methods for fundamental

studies of pitting corrosion has just started. Further investigations using these

methods and extension to other metals will broaden the experience and knowledge

and will give better insight into the mechanisms of this technologically important

type of corrosion.

Mechanisms of Pitting Corrosion 283

REFERENCES

1. P. Schmutz and G. S. Frankel, J. Electrochem. Soc. 145:2285 (1998).

2. H.-H. Strehblow and B. Titze, Corros. Sci. 17:461 (1977).

3. H. Prinz and H.-H. Strehblow, Corros. Sci. 40:1671 (1998).

4. H.-H. Strehblow, Werkst. Koros. 27:792 (1976).

5. T. P. Hoar, D. C. Mears, and G. P. Rothwell, Corros. Sci. 5:279 (1965).

6. K. J. Vetter and H.-H. Strehblow, Ber. Bunsenges. phys. Chem. 74:1024 (1970).

7. N. Sato, Electrochim. Acta 16:1683 (1971).

8. N. Sato, K. Kudo, and T. Noda, Electrochim. Acta 16:1909 (1971).

9. Ya. J. Kolotyrkin, Corrosion 19:261t (1964).

10. T. P. Hoar and W. R. Jacob, Nature 216:1299 (1967).

11. A. G. Revez and J. Kruger, Passivity of Metals (R.P. Frankenthal and J. Kruger, eds.),

The Electrochemical Society, Princeton, NJ, 1978, p. 137.

12. J. Kruger, G. G. Long, M. Kuriyama, and A. J. Goldman, Passivity of Metals and

Semiconductors (M. Froment, ed.), Elsevier Science Publishers, Amsterdam, 1983,

p. 163.

13. C. Y. Chao, L. F. Lin, and D. D. Macdonald, J. Electrochem. Soc. 128:1191 (1981).

14. C. Y. Chao, L. F. Lin, and D. D. Macdonald, J. Electrochem. Soc. 128:1194 (1981).

15. J. Augustynski, in Ref. 11, p. 285.

16. S. Meitra and E. D. Verink, in Ref. 11, p. 309.

17. T. E. Pou, O. J. Murphy, J. O. M. Bockris, L. L. Tongson, and M. Monkowski,

Proceedings, 9th International Congress on Metallic Corrosion, Toronto, 1984, Vol.

2, p. 141.

18. G. C. Wood, J. A. Richardson, M. F. Abd Rabbo, L. P. Mapa, and W. H. Sutton, in

Ref. 11, p. 973.

19. S. Szlarska-Smialowska, H. Viefhaus, and M. Janik-Czachor, Corros. Sci. 16:644

(1976).

20. M. Janik-Czachor and K. Kaszczyszyn, Werkst. Korros 33:500 (1982).

21. B. P. Löchel and H.-H. Strehblow, Electrochim. Acta 28:565 (1983).

22. B. P. Löchel and H.-H. Strehblow, J. Electrochem. Soc. 131:713 (1984).

23. W. P. Yang, D. Costa, and P. Marcus, Oxide Films on Metals and Alloys (B.

MacDougall, R. S. Alwitt, T. A. Ramanarayanan, eds.) Proceedings Vol. 92–22, The

Electrochemical Society, Princeton, NJ, 1992, p. 516.

24. S. Mischler, A. Vogel, H. J. Mathieu, and D. Landolt, Corros. Sci. 32:925 (1991).

25. P. Marcus and J. M. Herbelin, Corros. Sci. 34:1123 (1993).

26. H.-H. Strehblow, Proceedings Corrosion 91, NACE, Houston, 1991, p. 76/1.

27. H.-H. Strehblow, Proceedings Corrosion Prevention, Australasian Corrosion

Association, Brisbane, Australia, November 1997.

28. C. L. Me Bee and J. Kruger, Localized Corrosion, NACE, Houston, 1974, p. 252.

29. W. Paatsch, Ber. Bunsenges. Phys. Chem. 77:895 (1973).

30. H.-H. Strehblow, Werkst. Korros. 35:437 (1984).

31. K. J. Vetter and H.-H. Strehblow, in Ref. 28, p. 240.

32. S. Haupt and H.-H. Strehblow, Langmuir 3:873 (1987).

33. B. P. Löchel and H.-H. Strehblow, Werkst. Korros. 31:353 (1980).

34. U. Bertocci, Advances in Localized Corrosion,

NACE Proceedings of 2nd International

Conference on Localized Corrosion, Orlando, FL (H. S. Isaacs, U. Bertocci, J. Kruger,

and S. Smialowska, eds.), NACE, Houston, 1987. p. 127.

35. J. Steward and D. E. Williams, in Ref. 34, p. 131.

36. S. P. Mattin and G. T. Burstein, Progress in the Understanding and Prevention of

Corrosion (J. M. Costa and A. D. Mercer, eds.), Institute of Materials, 1993, p. 1109.

284 Strehblow

37. H.-H. Strehblow, dissertation, Freie Universität, Berlin, 1971.

38. K. J. Vetter and F. Gorn, Electrochim. Acta 18:321 (1973).

39. S Benamar, thesis, Université Pierre et Marie Curie, Paris, 1998.

40. I. Milosev and H.-H. Strehblow, Biomed. Mater. Res., submitted.

41. K. E. Heusler and L. Fischer, Werkst. Korros. 27:551 (1976).

42. H.-H. Strehblow, B. Titze, and B. P. Löchel, Corros. Sci. 19:1047 (1979).

43. H.-H. Strehblow and B. P. Löchel, Passivity of Metals and Semiconductors (M.

Froment, ed.), Elsevier Science Publishers, New York, 1983, p. 379.

44. B. P. Löchel and H.-H. Strehblow, J. Electrochem. Soc. 131:522 (1984).

45. Stability Constants of Metal Ion Complexes, Suppl. 1, Special Publ. 25, The

Chemical Society, London, 1971, p. 171.

46. Ref. 45, p. 167.

47. Ref. 45. p. 217.

48. W. Khalil, S. Haupt, and H.-H. Strehblow, Werkst. Korros. 36:16 (1985).

49. F. A. Cotton and G. Wilkinson, Anorganische Chemie, Verlag Chemie, Weinheim,

1980, p. 920.

50. Handbook of Chemistry and Physics, CRC Press, Boca Raton, FL, 1978, p. B105.

51. V. Maurice, V. Inard, and P. Marcus, Proceedings of the Symposium Critical Factors

in Localized Corrosion III, 194th meeting of the Electrochemical Society, Boston,

November 2–6, 1998 (R. G. Kelley, P. M. Natishan, G. S. Frankel, and R. C. Newman,

eds.), The Electrochemical Society Proceedings Series, PV 98-17, Pennington, NJ

(1998), pp. 552–562.

52. J. Kunze, diploma thesis, Heinrich-Heine-Universität, Düsseldorf, 1998, p. 64.

53. H.-H. Strehblow and J. Wenners, Z. Phys. Chem NF, 98:199 (1975).

54. H. Kaesche, Die Korrosion der Metalle, Springer, Berlin 1979, p. 268

55. H.-H. Strehblow and M. B. Ives, Corros. Sci.l6:317 (1976).

56. H.-H. Strehblow and J. Wenners, Electrochim. Acta 22:421 (1977).

57. Th. R. Beck and R. C. Alkire, J. Electrochem. Soc. 126:1662 (1979).

58. H.-H. Strehblow, Habilitationsschrift, Freie Universität, Berlin, 1977, p. 53.

59. G. S. Frankel, J. R. Scully, and V. Jahnes, J. Electrochem. Soc. 143:1834 (1996).

60. A. Seghal, D. Lu, and G. S. Frankel, J. Electrochem. Soc. 145:2834 (1998).

61. G. R. Engelhardt, A. D. Davydov, and T. B. Zhukova, Sov. Electrochem. 26:888

(1990).

62. J. Newman, D. N. Hansen, and K. J. Vetter, Electrochim. Acta 22:829 (1974).

63. G. Engelhardt and H.-H. Strehblow, J. Electroanal. Chem. 365:7 (1994).

64. P. Heimgartner and H. Böhni, Corrosion 41:715 (1985).

65. F. H. Hunkeler and H. Böhni, Werkst. Korros. 34:593 (1983).

66. H.-H. Strehblow, K. J. Vetter, and A. Willgallis, Ber. Bunsen Ges. Phys. Chem.

75:823 (1971).

67. K. J. Vetter and H.-H. Strehblow. Ber. Bunsenges. Phys. Chem. 74:449 (1970).

68. H. W. Pickering and R. P. Frankenthal, in Ref. 28, p. 261.

69. C. Schmidt and H.-H. Strehblow, J. Electrochem. Soc. 145:834 (1998).

70. U. Wolff, dissertation, Technische Universität Dresden, 1999.

Mechanisms of Pitting Corrosion 285

287

Sulfur-Assisted Corrosion Mechanisms

and the Role of Alloyed Elements

Philippe Marcus

Laboratoire de Physico-Chimie des Surfaces, CNRS, Université Pierre et Marie Curie,

Ecole Nationale Supérieure de Chimie de Paris, Paris, France

INTRODUCTION

Sulfur is an element which is found in the bulk of most materials (in saturated solid

solution and often in inclusions) as well as in many environments (in the form of,

e.g., HS

–

, H

S, SO

, S

–

The role of sulfide inclusions in corrosion has been recognized in early works.

The fact that sulfur-containing species are detrimental to the resistance of metals and

alloys to localized corrosion has been established for a long time but the mechanisms

have remained unclear until recently. In the area of passivity breakdown, where

substantial research effort has been expended for several years, the effects of chloride

ions have been investigated much more than the effects of sulfur. The aim of this

chapter is to review the fundamental aspects of the mechanisms of S-induced

corrosion, with special emphasis on the role played by adsorbed (or chemisorbed)

sulfur.

In the first sections of this chapter, the basic effects of adsorbed sulfur on

anodic dissolution, on passivation, and on the breakdown of passive films are

presented. In a subsequent section the effects of alloyed elements (essentially Cr

and Mo) are presented and the way in which they can counteract the detrimental

influence of sulfur is emphasized. In the next section, implications of the

mechanisms of sulfur-induced corrosion for different areas of practical importance

are given, with connections to the related chapters of this book. In the last section

of this chapter, the thermodynamic predictions of the conditions of adsorption of

sulfur on metal surfaces in water are given.

FUNDAMENTAL ASPECTS OF SULFUR-INDUCED CORROSION

The basic concepts and the associated mechanisms underlying the effects of adsorbed

sulfur on the corrosion of metals and alloys are presented here. They have been

derived to a large extent from studies performed on well-defined metal and alloy

surfaces. In many cases, oriented single crystals have been used to control the

structure of the surface on which the corrosion reactions were investigated.

Controlled amounts of sulfur in the range of a monolayer were produced prior to the

electrochemical measurements by dosing the surface with gaseous H

S. Surface

science techniques, including Auger electron spectroscopy (AES), electron

spectro-scopy for chemical analysis (ESCA), also called X-ray photoelectron

spectroscopy (XPS), low-energy electron diffraction (LEED), and a more specific

radiochemical technique utilizing the

S radioisotope, were combined to analyze

the surface prior to and after the electrochemical tests, in order to relate the

electrochemical and corrosion behavior with the chemical composition and the

structure of the surface. This approach has allowed us to determine the mechanisms

of the sulfur-assisted corrosion reactions on an atomic scale.

Activation (or Acceleration) of Anodic Dissolution by

Adsorbed Sulfur

The first studies were conducted on Ni [1–3] and Ni-Fe [4] single crystals with

well-controlled coverages of radioactive sulfur. They clearly demonstrated that adsorbed

sulfur accelerates the anodic dissolution of the metal and the alloys. This effect

is seen in Figure 1, where the polarization diagrams for a nickel-iron single-

crystal alloy [Ni-25Fe(100)(at %)] with and without adsorbed sulfur are shown.

Two major findings are apparent from the measurements: (a) a significant increase

288 Marcus

Figure 1 The influence of adsorbed sulfur on the i-E diagram for Ni-25Fe (at %) single

crystal [(100) face] in 0.05 M H

. θ is the coverage of the surface by adsorbed sulfur,

measured with the radiochemical technique (

S radiotracer).

in anodic current density is observed when the surface is covered by a monolayer

of adsorbed sulfur, showing that the dissolution is accelerated by the presence of

sulfur, and (b) the surface sulfur is highly stable, since it remains adsorbed during

dissolution of the metal. This result was established by the measurements of the

coverage, denoted by θ in Figure 1. This gives evidence of the catalytic nature of

the accelerating effect of adsorbed sulfur on the dissolution. In this way, a micro-

scopic amount of adsorbed sulfur can stimulate the dissolution of a significant

amount of material (about 1 μm of alloy is dissolved in the active region of Figure

1 under the catalytic action of a monoatomic layer of adsorbed sulfur). The cat-

alytic effect of a monolayer of adsorbed sulfur on anodic dissolution of metals and

alloys is summarized in the schematic i-E diagram of Figure 2 with, in the inset,

potential shifts measured with adsorbed sulfur on different metals and alloys. Since

the first direct demonstration of the accelerating effect of dissolution by adsorbed

sulfur, which was for nickel and nickel-iron alloys, the more general nature of this

effect has been revealed by similar experiments on nickel alloy 600 [5,6] and on

austenitic stainless steels [7]. The acceleration factor R (ratio of the current densi-

ties at the maximum of the active peak with and without adsorbed sulfur) for nickel-

base and iron-base alloys are reported in Table 1. Values in the range 5 to 12 have

been observed. This striking effect has been attributed to a weakening of the

metal-metal bonds induced by adsorbed sulfur, which leads to a lowering of the

Sulfur-Assisted Corrosion Mechanisms 289

Figure 2 Schematic i-E plots showing the effect of adsorbed sulfur on anodic dissolution.

Inset: Potential shift (ΔE

= 1

) induced by a complete monolayer of adsorbed sulfur on

different metals and alloys (ΔE

= 1

is measured at half-maximum of the active peak of the

sulfur-free surface).