Marcus P. Corrosion mechanisms in theory and practice

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activation energy barrier for the passage of a metal atom from the surface to the

solution. This mechanism is shown schematically in Figure 3 [8]. The existence of a

dipole M

δ+

– S

δ–

associated with the presence of sulfur can also promote the

anodic dissolution.

The influence of sulfur can be strongly localized if the sulfur is not homogeneously

distributed over the surface but adsorbed in specific sites, such as surface defects,

where sulfur atoms are more tightly bonded.

Direct observation of the dissolution of sulfur-modified Ni(100) [9] using in situ

scanning tunneling microscopy (STM) is in agreement with the previously shown

effects of adsorbed S on dissolution (including the enhanced dissolution and the

stability of the S layer). It also provides structural data at the atomic scale, indicating

that the dissolution of nickel atoms is taking place at step edges, with a displacement

of adsorbed S atoms from the edge of the upper terrace to the adjacent lower terrace,

as shown in Figure 4.

On Cu, it was also shown that the anodic dissolution, in a mildly alkaline

borate buffer solution, is catalyzed by a submonolayer of sulfur and that S remains

adsorbed on the copper surface [10].

The anodic dissolution of iron in acidic medium is also accelerated by H

S [11].

290 Marcus

Table 1 Ratio (R) of the Current Densities at the Maximum of the Active Peaks with and

without Adsorbed Sulfur

Coverages of the surfaces by preadsorbed sulfur are (in 10

– 9

g/cm

): 42 on Ni(100), 47 on Ni(111), 44

on Ni(110), 40 on Ni-25Fe (100), 42 on Ni-25Fe(111), 41 on Ni-25Fe (110), 40 on Fe-17Cr-13Ni(100),

and 40 on Ni-15Cr-10Fe.

Figure 3 Weakening of the metal-metal bond by adsorbed sulfur and acceleration of the

dissolution of the metal.

The major conclusion of the results reviewed in this section is that a

monoatomic layer of sulfur can promote the dissolution of macroscopic amounts

of material. This gives evidence of a direct link between the atomic-level interactions

of adsorbed sulfur atoms with metal surfaces and their macroscopic manifestations in

corrosion.

Blocking or Retarding Effect of Adsorbed Sulfur on Passivation

Effect on the Growth of the Passive Film

Another important effect of adsorbed sulfur that was observed first on Ni [1–3]

and later on Ni-Fe alloys [4] is the poisoning of passivation. This effect takes place

above a critical coverage of the surface by sulfur that was found to be 0.7–0.8

monolayer. This effect of sulfur is observed in the i-E diagram shown previously in

Figure 1, where the passivation potential of the surface covered by a complete

monolayer of sulfur (θ = 1) is shifted (by ~100mV) to a more anodic value with

respect to the sulfur-free surface. In using the

S radiotracer, it was demonstrated

that the desorption (or electro-oxidation) of ~20% of the full monolayer is

necessary to allow the passive film to be formed. The retarding effect of sulfur on

the growth of the passive film is caused by blocking the sites of adsorption of

hydroxyl ions, which are the precursors in the formation of the passive layer. X-ray

photoelectron spectroscopy (XPS) measurements have revealed that in the

presence of adsorbed sulfur the adsorption of OH groups is not totally inhibited,

but the OH groups on the surface are more diluted and thus the disproportionation

reaction of adjacent OH groups is prevented. A schematic representation of this

mechanism is given in Figure 5. In this way the oxide film, which would normally

passivate the surface, is not formed.

Effect on the Passivation Kinetics

Measurements of the passivation kinetics by means of potential steps applied to

surfaces without or with adsorbed sulfur have revealed that the time to complete

passivation increases when the surface is covered by adsorbed sulfur. This effect is

Sulfur-Assisted Corrosion Mechanisms 291

Figure 4 Model of the anodic dissolution process of sulfur-covered Ni(100) electrodes,

showing dissolution of two Ni atoms from the step edge and displacement of an S atom from

the upper terrace to a fourfold hollow site on the lower terrace (small circles denote Ni atoms,

large circles denote S atoms, the gray color indicates atoms of the upper terrace). (Adapted

from Ref. 9.)

represented in the i-t diagram of Figure 6. A rapid decrease of the current

immediately after the potential step is observed for the clean surface, whereas for

the sulfur-covered surface the current density decreases slowly, indicative of a slow

process of lateral growth of the film on top of the remaining fraction of a monolayer

of adsorbed sulfur. This is consistent with the view that a complete monolayer of

sulfur inhibits the nucleation of the oxide, as observed on nickel, whereas a fraction

of a monolayer does not inhibit the nucleation but may diminish the density of

nucleation sites for the oxide and make the lateral growth of the film more difficult.

These two effects can slow down the passivation kinetics.

Such kinetic effects have been observed on sulfur-contaminated nickel [8], on

nickel alloy 600 [6], and on copper [12].

292 Marcus

Figure 5 The mechanisms of the blocking effect of adsorbed sulfur on passivation.

Figure 6 The effect of adsorbed sulfur on the passivation kinetics of nickel [Ni(111), 0.05 M

, 540 mV/SHE].

Effect on the Structure and Properties of the Passive Film

Below the critical sulfur coverage of 0.7–0.8 monolayer on Ni, the passive film grows

on top of the remaining adsorbed sulfur. However, the structure and the properties

of the film formed on the sulfur-contaminated surface are modified. A study [2] by

reflection high-energy electron diffraction (RHEED) of the structure of the passive

films formed on nickel single crystals [Ni(111)] has shown that, in the absence of

sulfur, a crystalline film epitaxial with the substrate is formed (a fact that has been

confirmed by both ex situ and in situ atomic resolution imaging of the film by

scanning tunneling microscopy [13–15]), whereas in the presence of sulfur the

epitaxial growth is disrupted and a more defective polycrystalline film is obtained.

Accordingly, the current density in the passive state is about four times larger for a

film formed with sulfur at the metal-oxide interface. XPS measurements of the

binding energy of sulfur (S 2p) provide evidence that it remains at the interface

because the binding energy remains close to that of sulfur bonded to the metallic

surface (~ 162 eV).

Anodic Segregation of Sulfur

The concept of anodic segregation was first proposed after the discovery that

sulfur present as an impurity in the bulk of Ni and of Ni-25Fe alloys is enriched on

the surface during anodic dissolution [2,4,16]. This phenomenon was also

observed on Fe-50Ni [17] and on alloy 600 [5,18]. The experiments were performed

on Ni and Ni-Fe alloys that had been doped with sulfur prior to the electrochemical

treatments. The introduction of radioactive S into the bulk was performed by

exposing the samples to an H

S-H

gas mixture (with

S-labeled H

S) at high

temperature (1000 to 1200°C) under thermodynamic conditions in which sulfur is

in solid solution in the metal or the alloy. Rapid quenching to room temperature

after the high-temperature treatment avoids sulfide precipitation. The i-E curves

recorded on a sulfur-free (i.e., annealed in pure hydrogen) Ni electrode and on Ni

containing 50 ppm sulfur are shown in Figure 7a. The sulfur-free Ni sample

exhibits the normal active-passive transition observed in 0.05 M H

. The

sample with bulk sulfur exhibits a completely different behavior: there is no active-

passive transition, and the current density increases with increasing potential in the

whole range of potentials corresponding to the passive state of sulfur-free Ni. The

results of the measurements of the sulfur concentration on the surface, using

radioactive sulfur, are shown in Figure 7a. They demonstrate that sulfur segregates

on the surface during dissolution of the electrode with bulk sulfur (the measured

values of the surface sulfur concentration are denoted S

segregated

in Fig. 7a). The

values denoted S

exposed

in Figure 7a represent the integrated amount of sulfur that

became available on the surface due to the dissolution of the metal matrix.

Comparison of the data for S

segregated

and S

exposed

revealed that all the sulfur present

in the bulk remains on the surface during dissolution, up to a surface concentration of

40 × 10

–9

g/cm

, which corresponds to a complete monolayer of sulfur on Ni(100).

Above this surface coverage, nickel sulfide precipitates and then a thin layer of nickel

sulfide is formed. The amount of sulfur in this layer is ~40 × 10

–8

g/cm

, which corresponds

to ~25 Å of Ni

. Sulfur in excess of this amount is dissolved. The formation, before

reaching the passivation potential, of a complete monolayer of sulfur and the subsequent

Sulfur-Assisted Corrosion Mechanisms 293

growth of a sulfide layer preclude the formation of the passive oxide film. The

sulfide layer is not protective and thus high corrosion rates are obtained in the range

of potentials in which pure Ni is normally passivated. This phenomenon was called

anodic segregation [3,4]. The anodic segregation rate depends on the sulfur content

of the material and on the rate of anodic dissolution. Because of differences in

dissolution rates for single-crystal surfaces of different crystallographic orientation,

the sensitivity to sulfur increases in the order (111) ≤ (100) < (110), which is the order

of increasing dissolution rates observed for Ni in 0.05 M H

[2]. The mechanism

of anodic segregation is represented in Figure 7b.

294 Marcus

Figure 7 (a) Surface enrichment of sulfur during anodic dissolution: i-E curves and radio-

tracer measurements of the sulfur concentration on the surface of Ni(100) containing 50 ppm

of sulfur (0.05 M H

, 1 V h

–1

). (b) The mechanism of anodic segregation.

Passivity Breakdown Caused by Interfacial Sulfur

After the formation of the passive film on a bare metal surface, slow dissolution of

the metal cations continues, involving dissolution at the passive film surface and

transport of ions through the oxide (see Chap. 6). The question then arises as to where

the bulk impurities, e.g., sulfur, go during this process. This question was addressed

in a detailed investigation of Ni and Ni-Fe alloys that were doped with radioactive

sulfur in order to trace the path taken by the sulfur atoms during the slow dissolution

in the passive state [19]. The radiochemical results gave direct evidence that sulfur

present in the bulk accumulates at the metal–passive film interface. The enrichment

rate was shown to be proportional to the sulfur content in the metal and to the

dissolution current density. Above a critical concentration of sulfur at the interface

(metal-oxide), breakdown of the passive film was observed. This critical concentration

of interfacial sulfur was measured with radioactive sulfur and found to be close to one

monolayer of sulfur (i.e. ~ 40 × 10

–9

g/cm

). The loss of adherence (or the decohesion)

at the interface for this critical coverage has been attributed to aweakening of the

bonding of the oxide to the substrate caused by sulfur. The nucleation of the sulfide

that is expected to take place when the sulfur concentration exceeds a

complete monolayer may be responsible for the observed local breakdown of the film

and pitting. The defects that are likely to exist at the interface could serve as specific

sites for the nucleation of the sulfide, once a quasi-complete monolayer of sulfur has

been accumulated at the interface. A schematic representation of the proposed

mechanism of the sulfur-induced breakdown of the passive film is shown in Figure 8.

It is to be noted that the three following factors are in favor of sulfur remaining

at the metal–passive film interface: (a) the sulfur-metal chemical bond is very strong

(see Chap. 2), (b) the solubility of S in nickel oxide (which constitutes the inner part

of the passive film on Ni and Ni-Fe alloys) is very low, and (c) the electric field across

the passive film, which assists the passage of cations from the metal–passive film

interface to the passive film–solution interface, should impede the transport of sulfur,

which would be negatively charged.

The Joint Action of Sulfur and Chloride Ions

The passivity breakdown and pitting caused by Cl

–

ions has been the subject of an

enormous amount of work. This area has been reviewed in Chapter 8.

Sulfur-Assisted Corrosion Mechanisms 295

Figure 8 Mechanism of the breakdown of the passive film induced by enrichment of sulfur

at the metal–passive film interface.

In the preceding parts of this chapter we have seen that S alone can enhance the

dissolution, block or retard the growth of the passive film, and cause passivity

breakdown and pitting by enrichment at the metal-film interface. All these effects can

evidently also take place in the presence of Cl

–

. A major difference between the

mechanisms of action of Cl

–

and S is that S does not seem to interact directly with

the oxide surface as strongly as Cl

–

The effect of sulfur species and chloride ions has been studied in some detail in

the case of corrosion of stainless steels in solutions containing thiosulfate (S

–

)

and chloride (Cl

–

) ions [20–27].

It has been demonstrated [27] that S

–

is not reduced on the surface of the

passive film formed in neutral solution on an Fe-17Cr alloy, whereas the reduction to

adsorbed sulfur or sulfide does take place on the nonpassivated alloy surface. In this

study, the combined action of chloride and sulfur was also investigated. The

conclusion was that with 30 ppm thiosulfate added to the Cl

–

-containing solution

(0.02 M) at pH 7, the initial breakdown of the passive film is caused by Cl

–

and the

thiosulfates do not play a major role in this step, whereas the presence of thiosulfate

has a drastic effect on the subsequent step, the occurrence of stable pits. With Cl

–

, and

without S

–

, unstable pits (i.e., pits that are repassivated) are formed over a wide

range of potentials below the classical pitting potential corresponding to the formation

of stable pits, whereas with the thiosulfates, the reaction on the bare metal surface

produces reduced sulfur (adsorbed sulfur or sulfide), which precludes repassivation

and causes a marked increase of stable pits at low potential. This effect is shown in

Figure 9. The formation of adsorbed sulfur by reaction of S

–

with the bare

metal surface was demonstrated by XPS measurements of the binding energy of the

sulfur core level electrons (S2p), which was found to be 169eV for adsorbed thiosulfate

on the passive film surface and 162eV after reduction of the thiosulfate on the

Fe-17% Cr alloy surface. In previous works on the effects of addition of S

–

corrosion of stainless steels [21,25], such a mechanism had been hypothesized, but

direct experimental evidence was lacking. A synergistic effect of Cl

–

and S

–

pitting corrosion was also observed on 310 stainless steel [28]. Type 304L stainless

steel suffers stress corrosion cracking in chloride solutions with 10

–3

to 10

–1

M S

–

(at 353 K). It was shown that the cracks initiated at pits [29] and the authors

concluded that S

–

affects the initiation process of pitting and cracking.

To summarize, a likely mechanism for the combined action of chlorides and sulfur

is the local breakdown of the film by Cl

–

, which then permits all the effects of S shown

above on metal and alloy surfaces to take place. A major consequence is the stabilization

of otherwise unstable (or metastable) pits. This mechanism is consistent with the lower

pitting potential and/or shorter incubation time observed experimentally.

Sulfur species in aqueous solution may also cause directly the breakdown of the

film if the sulfide is thermodynamically more stable than the oxide, a case that would

require a high concentration of the sulfur species in solution to provide a reaction of

the type M

+ zH

S↔M

+ yH

O+2(z–y)H

+2(z–y)e

–

Sulfide inclusions have been known for many years to be detrimental to the corro-

sion resistance of steels and extensive studies have been performed on

stainless steels ([30–33] and references in Chapter 10). Preferential adsorption of Cl

–

sulfide inclusions has been suggested to be an important step in the corrosion mechanism.

The precise mechanism of the joint action of Cl

–

and the sulfide is not fully elucidated.

296 Marcus

COUNTERACTING THE SULFUR EFFECTS WITH ALLOYED

ELEMENTS

The Role of Molybdenum

Numerous electrochemical and surface analytical studies have been performed to

understand the effect of alloyed molybdenum on the corrosion resistance of stainless

steels (see, e.g., the section on the role of Mo in Mo-containing austenitic stainless

steels in Chapter 7).

Sulfur-Assisted Corrosion Mechanisms 297

Figure 9 The joint action of S

–

and Cl

–

on passivity breakdown and localized corrosion

of an Fe-17Cr alloy. (a) 0.02 M Cl

–

; (b) 0.02 M Cl

–

+ 30 ppm S

–

. S

–

is reduced on the

alloy surface, as evidenced by the difference in the binding energy of S 2p for S

–

(169 eV)

and S

ads

(162 eV).

In order to reach a better understanding of the role of molybdenum in the

presence of sulfur, studies were undertaken on simple systems, i.e., binary and ternary

single-crystal alloys on which the surface concentrations of sulfur and of Mo could be

precisely measured by radiochemical (

S) and spectroscopic (XPS) techniques. The

experiments performed on nickel-molybdenum alloys [Ni-2Mo and Ni-6Mo(at %)] pro-

vided the first direct evidence of a specific surface interaction between molybdenum and

adsorbed sulfur leading to the removal of sulfur from the surface and thereby the atten-

uation or the disappearance of the detrimental effects of adsorbed sulfur [34,35].

The most conclusive experiments were done in (a) preadsorbing, in a gas

mixture of H

S and H

, a known amount of sulfur on the well-defined surface of an

Ni-2 at % Mo alloy [the (100) crystallographic orientation of this fee alloy was used]

and (b) measuring the variation of the sulfur coverage during anodic dissolution in the

active state (at 320mV/SHE) in 0.05 M H

. The first and essential observation was

that the sulfur coverage decreases, whereas it had been found to remainconstant on pure

Ni (discussed earlier in this chapter). The precise measurement of the desorption kinet-

ics is shown in Figure 10 in terms of sulfur coverage (weight per cm

and number of

sulfur atoms per metal atom on the surface) versus the amount of dissolved material

(expressed in number of dissolved atom layers). The initial coverage corresponds to

saturation of the surface by adsorbed sulfur [42 ng/cm

, one S atom for two M atoms

on the surface of the (100) face]. Because sulfur was initially present only on the

surface, the results demonstrate that a surface reaction between Mo and sulfur is the

cause of the loss of the surface sulfur. This is a dynamic process in which the dissolu-

tion of the alloy brings Mo to the surface, which bonds to adsorbed sulfur and is then

dissolved with it. A theoretical model was proposed for the mechanism of the effect of

Mo. In this model, one S atom adsorbed on the surface reacts with k Mo atoms of the

surface plane and the cluster that is formed (in which O or OH may participate) is

dissolved. It was shown that the sulfur coverage on the surface after dissolution of

(n+ 1) atom planes of the alloy, denoted θ

(n + 1), has the following expression [34]:

(n + l) = θ

(n) – [(θ

(n))

)

]

where C

is the surface concentration of Mo, which was assumed to be identical to

the bulk content of Mo, and k is the number, defined above, of Mo atoms required

to form with sulfur a cluster that is dissolved. This equation was used to fit the

experimental variation of the sulfur coverage. The results for k = 1, 2, 3 are computed

in Figure 10. The curve with k = 2 provides a very good fit to the experimental curve.

Similar findings were obtained with Ni-6% Mo (100). These results strongly support

the proposed mechanism in which, on a (100) surface, two Mo atoms bond to and

remove one adsorbed S atom. It is to be noted that in this mechanism the passive

film is not directly involved, but of course the major consequence of the removal

of sulfur by Mo is that the passive film may be formed on an otherwise blocked

surface. The same model has been used successfully to interpret in a quantitative

manner the fact that on other Ni-Mo alloys the surface enrichment of sulfur by

anodic segregation of bulk sulfur was very limited [36] compared with what had been

reported previously for Ni [2]. On Ni-6%Mo (100) with 32 ppm sulfur in the bulk,

the coverage by sulfur measured after different times of anodic polarization in the

active region (in 0.05 M H

) was always less than a complete monolayer. The

theoretical amount of sulfur, calculated using an equation similar to that given

298 Marcus

above for the Mo-induced desorption of adsorbed S, but in which a term accounting

for the supply of sulfur by the dissolution of the alloy was added, gave a good fit to

the experimental values when the value of k was taken equal to 2. This confirmed

the validity of the mechanism in which two Mo atoms bond to one sulfur atom and

Sulfur-Assisted Corrosion Mechanisms 299

Figure 10 The loss of surface sulfur induced by molybdenum (experimental and

theoretical curves).