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Investigations of Structure and Stability of Passive Films by Surface Analytical Techniques
Investigations of Structure and Stability of Passive Films by Surface Analytical Techniques P Marcus and V Maurice, PSL Research University, CNRS - Chimie ParisTech, Institut de Recherche de Chimie Paris/Physical Chemistry of Surfaces Group, Paris, France © 2018 Elsevier Inc. All rights reserved.
Introduction Structure of Passive Films Dissolution of Passive Films Nanostructure of Passive Films Passivity Breakdown and Initiation of Localized Corrosion Conclusion and Outlook Further Reading
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Nomenclature AFM DFT EC-STM fcc GIXD ML SHE STM STS ToF-SIMS UHV XPS
Atomic force microscopy Density functional theory Electrochemical scanning tunneling microscopy Face centered cubic Grazing incidence X-ray diffraction Monolayer Saturated hydrogen electrode Scanning tunneling microscopy Scanning tunneling spectroscopy Time-of-flight secondary ion mass spectrometry Ultra high vacuum X-ray photoelectron spectroscopy
Introduction The formation of passive films on metals and alloys is a central issue in corrosion science and engineering since it is the best of all means for protection of metallic materials against corrosion, and a key for their use in our environment. Passivity originates from the growth of an anodic oxide film on the surface of metal and alloy substrates in aqueous environment. The oxide film is either thermodynamically stable, or it dissolves very slowly, so that it protects the metal against corrosion. The oxide passive films most commonly do not exceed a few nanometers in thickness, are hydroxylated, well adherent, and effectively isolate the substrate from the corrosive environment. They are self-healing and may repair after breakdown. Surface analytical methods combined with electrochemical methods have been widely used to characterize the passivity of numerous metals and alloys and determine the thickness, composition, and electronic properties of passive films. Insight in the atomic structure of passive films and their morphology at the nanometer scale can be obtained using singlecrystal surfaces, prepared atomically flat, exposed to well-defined corrosive aqueous solutions, and investigated in situ under electrochemical control or ex situ after well-controlled electrochemical treatment and transfer of the sample to the surface analysis instrument. Scanning tunneling microscopy and atomic force microscopy, implemented (EC-STM, EC-AFM) or not (STM, AFM) with electrochemical control, as well as in situ grazing incidence X-ray diffraction (GIXD) with synchrotron light source have been applied on pure metal substrates like Cu, Ni, Fe, Cr, and Co, and on alloy substrates like stainless steel and nickel-based alloys. Passive films are polycrystalline with grains of nanometric lateral dimensions due to a high density of sites for oxide nucleation. The microstructure is textured by the preferential crystallographic orientation adopted by the oxide grains. Nonordered areas can form between the crystalline grains, particularly during the growth process, that is, under nonstationary conditions. These intergranular sites of the passive film play a key role in the breakdown of passivity and the initiation of localized corrosion. For Cr-rich passive films, the development of crystallinity is conditioned by electrode potential and aging under polarization. Structural characterization of passive films is usually complemented by surface chemical analysis using photoelectron spectroscopy (XPS) and ion spectrometry (ToF-SIMS).
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Structure of Passive Films Hydroxide ions from the liquid water adsorb on the surface in the potential range preceding 3D anodic oxide growth. On Cu(111), this process occurs at 0.6 V/SHE in 0.1 M NaOH (Fig. 1A). It induces surface reconstruction and formation of a 2D layer of adsorbed hydroxide (OHads). At low oversaturation potential for anodic oxide formation ( 0.25 V/SHE), the metal substrate is partially covered by poorly crystallized and one monolayer thick oxide islands formed after preferential nucleation at step edges and separated by the hydroxide adlayer. At higher oversaturation ( 0.2 V/SHE), the metal substrate becomes fully covered by well crystallized and several monolayer thick (111)-oriented Cu(I) oxide films (Fig. 1B). The equivalent thickness of the oxide layer deduced from charge transfer measurements during subsequent cathodic reduction is 0.5 and 7 equivalent monolayers (ML) after growth at 0.25 and 0.2 V/SHE, respectively. One ML corresponds to one (111)-oriented O2 -Cuþ-O2 slab of the cuprite structure (Fig. 1C). The 2D hydroxide adlayer is a structural precursor for the 3D passive film, adopting the same building block. The oxide layer has a hexagonal lattice with a parameter of 0.3 nm (inset in Fig. 1B), consistent with the Cu sublattice in the (111)-oriented cuprite. The fcc oxide film grows in parallel (or antiparallel) epitaxy on the fcc substrate: Cu2O(111)[1 1 0] || Cu(111)[1 1 0] or [1 10]. On Cu(001), the Cu2O passive film is also a few ML but is (001)-oriented with a 45-degree rotation between the close-packed directions of the oxide and substrate lattices: Cu2O(001)[1 1 0] || Cu(001)[100].
Fig. 1 Passivation of copper as observed on Cu(111) in 0.1 M NaOH(aq). (A) Electrochemical (voltammetry) data recorded between the hydrogen and oxygen evolution limits with marked Cu(0) active region including OH adsorption and passive Cu(I) and Cu(I)/Cu(II) regions. (B) Cu(I) oxide passive film grown as observed by EC-STM at 0.20 V/SHE. The atomic lattice (inset) observed in situ corresponds to the Cu sublattice in Cu2O(111). (C) The side view model illustrating the stacking sequence of the O2 and Cuþ planes in the oxide and the surface termination by a monolayer of hydroxyl/hydroxide groups. (A) Adapted from Kunze, J.; Maurice, V.; Klein, L.H.; Strehblow, H.-H.; Marcus, P. In Situ STM Study of the Effect of Chlorides on the Initial Stages of Anodic Oxidation of Cu (111) in Alkaline Solutions. Electrochim. Acta 2003, 48, 1157–1167. Copyright 2003, Elsevier. (B,C) Reprinted with permission from Kunze, J.; Maurice, V.; Klein, L.H.; Strehblow, H.-H.; Marcus, P. In Situ Scanning Tunneling Microscopy Study of the Anodic Oxidation of Cu (111) in 0.1 M NaoH. J. Phys. Chem. B 2001, 105, 4263–4269. Copyright 2001, American Chemical Society.
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Owing to a tilt of a few degrees of the orientation of the oxide lattice with respect to that of the metal lattice, crystalline passive film expose a faceted surface with periodic step edges as evidenced on copper by EC-STM (Fig. 1). The tilt results in part from the relaxation of the stress at the metal/oxide interface resulting from the large mismatch between the two lattices (17% on Cu). The height of the surface steps of the Cu(I) oxide layer corresponds to 1 ML of cuprite, indicating an identical chemical termination of the Cu2O oxide terraces. The oxide layer surface is hydroxylated at the interface with liquid water, as confirmed by density functional theory (DFT) modeling for the Cu2O(111) surface showing that OH adsorption on the oxide lattice stabilizes the unreconstructed structure of the Cu sublattice observed in situ by EC-STM. The structure of the duplex Cu(I)/Cu(II) passive films formed in the potential range of Cu(II) anodic oxidation is also crystalline with surface hydroxylation stabilizing the otherwise polar (001)-oriented CuO structure of the Cu(II) outer layer of the passive film. The passive film on nickel is also crystalline, as observed by EC-STM on Ni(111) in acid solution (Fig. 2). Like on copper, the passive film surface is faceted, exhibiting regular terraces and step edges due to the slightly tilted epitaxy that allows relaxation
Fig. 2 Passivation of nickel as observed on Ni(111) in 0.05 M H2SO4 þ 0.095 M NaOH(aq.) (pH 2.9). (A) Electrochemical polarization curve with the active, passive, and transpassive regions marked. (B) Ni(II) oxide passive film as observed by EC-STM at 0.95 V/SHE. The inset shows the high resolution image recorded in situ, revealing the passive film atomic lattice. (C) The side view model of the hydroxylated NiO surface with (010) step edges and (111) terraces as optimized by DFT simulation. (B) Adapted with permission from Maurice, V.; Klein, L.H.; Marcus, P. Atomic-Scale Investigation of the Localized Corrosion of Passivated Nickel Surfaces. Surf. Interf. Anal. 2002, 34, 139–143. Copyright 2002, Wiley. (C) Reprinted from Bouzoubaa, A.; Diawara, B.; Maurice, V.; Minot, C.; Marcus, P. Ab Initio Modelling of Localized Corrosion: Study of the Role of Surface Steps in the Interaction of Chlorides With Passivated Nickel Surfaces. Corros. Sci. 2009, 51, 2174–2182. Copyright 2009, Elsevier.
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of the interfacial stress associated with the lattice mismatch (16%) between the NiO(111) lattice forming the barrier oxide layer and the Ni(111) lattice of the substrate. The (111) orientation of the NiO lattice and its tilt of about 3 degree with respect to the Ni(111) orientation have been confirmed by in situ GIXD measurements. The fcc oxide grows in antiparallel epitaxy on the fcc substrate: NiO(111)[1 1 0] || Ni(111)[1 10]. The atomic lattice measured by EC-STM on the terraces is hexagonal with a parameter of 0.3 0.02 nm consistent with the O (and Ni) sublattice of NiO(111). The NiO lattice obtained by passivation is (111)-oriented despite the fact that this surface is polar and thus unstable if bulk-like terminated. DFT modeling confirmed that the reason for this preferential (111) orientation of the terraces is that the surface is stabilized by adsorption of a monolayer of hydroxyl groups. Fig. 2C shows a model for such a faceted and hydroxylated surface as optimized by DFT. The direction of growth of the oxide film is thus governed by the minimization of the interface energy by a tilted growth and by the minimization of the oxide surface energy by the hydroxyl/hydroxide group termination. On chromium and stainless alloys, including stainless steels, the formation of Cr(III)-enriched oxide passive layers provides excellent corrosion resistance. Potential and aging under polarization are critical factors for the development of crystalline passive films. Electrode potential and aging-dependent crystallization observed by STM are paralleled by dehydroxylation measured by X-ray photoelectron spectroscopy (XPS). On chromium passivated in acid solution at low potential, the passive film is highly hydrated and the Cr(III) oxide inner layer is only partially developed. The passive film consists of small Cr(III) nanocrystals buried in a disordered Cr(III) hydroxide matrix. At high potential, the inner part of the passive film is dehydrated and consists mostly of larger crystals of Cr(III) oxide. A faceted topography extending over several tens of nanometers has been observed by EC-STM. The nanocrystals have a lattice consistent with the O sublattice in a-Cr2O3(0001) (corundum structure). The basal plane of the oxide is parallel to bcc Cr(110) with the closed-packed rows of the O sublattice nearly parallel to those of the metal and a mismatch of 12%. On stainless steels passivated in acid solution, the crystallinity of the passive films decreases with increasing Cr content of the alloy for short polarization times ( 2 h). Structural changes also occur during aging under anodic polarization. The major structural modification is an increase of the crystallinity of the film and the coalescence of Cr(III) oxide nanocrystals in the inner oxide as observed on Fe-22Cr and Fe-18Cr-13Ni alloys studied over time periods of up to 65 h. Like on chromium, the measured atomic lattice is consistent with the O sublattice of the (0001)-oriented a-Cr2O3 corundum structure. The rate of crystallization is more rapid on the austenitic stainless steel than on the ferritic one.
Dissolution of Passive Films The steps at the surface of the passive film play a key role for dissolution in the passive state, and thus for the stability of passivated surfaces. The passive film dissolves at the edges of the facets produced by the tilted epitaxy between the oxide and metal lattices (Fig. 2), which produces a receding step flow. The process is similar to that of active dissolution of oxide-free metal surfaces at moderate potential. The receding 2D step flow is dependent on the step orientation since the step edges oriented along the closed-packed directions of the oxide lattice dissolve much less rapidly, due to the higher coordination of their atoms. This structure-dependent process leads to the stabilization of the facets with edges oriented along the close-packed directions of the oxide lattice. On passivated nickel, it produces steps that are oriented along the {100} planes, the most stable orientations of the NiO structure. This receding step flow dissolution mechanism is observed independently of the presence or absence of chloride anions in the electrolyte. Dissolution of the passive film, catalyzed by chlorides complexing metal cations at the surface of the passive film, is a key aspect of the adsorption mechanism of passivity breakdown and initiation of localized corrosion. DFT modeling has been applied to hydroxylated and stepped NiO(111) surfaces in order to investigate the interaction with Cl and its effect on the dissolution of the oxide. A (533)-oriented NiO periodic model was built-up to simulate the (111) termination of the passive film terraces with monoatomic (010) step edges (Fig. 3). Cl adsorption structures were modeled by substituting the surface OH groups by Cl atoms at 25%, 50%, 75% and 100% coverages. After DFT optimization, substructures of Ni(OH)2, Ni(OH)Cl, or Ni(Cl)2 composition
Fig. 3 DFT modeling of the hydroxylated NiO(533) surface including a (010) step edge and a (111) terrace. Side view models of Cl-adsorbed structures at 25% coverage are shown before and after optimization. Adapted from Bouzoubaa, A.; Diawara, B.; Maurice, V.; Minot, C.; Marcus, P. Ab Initio Modelling of Localized Corrosion: Study of the Role of Surface Steps in the Interaction of Chlorides With Passivated Nickel Surfaces. Corros. Sci. 2009, 51, 2174–2182. Copyright 2009, Elsevier.
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were observed to form and detach from the step edges of the adsorbed structures, confirming the major role of the step edges in the dissolution of the oxide. The calculated energies of detachment of the substructures revealed that the Cl-containing substructures are easier to detach, showing that dissolution at the step edges can be promoted by the adsorption of Cl and confirming a Cl adsorption-induced oxide thinning as a possible breakdown mechanism of the oxide passive film.
Nanostructure of Passive Films Even when grown on single-crystal substrates, passive films have a polycrystalline microstructure with grains of nanometric lateral dimensions. This is a result of a high density of sites for oxide nucleation in their initial stages of formation. The lateral dimensions of the crystalline grains forming the passive films are well documented by STM, AFM, and GIXD data for nickel. Values determined from the morphology observed by STM and AFM have been reported to range from 2 nm in the initial stages of 2D growth to 30 to 230 nm for 3D films in stationary conditions of passivity. A large dispersion could be found on the same sample, suggesting a varying degree of advancement of the coalescence of the oxide grains during growth after nucleation. A lower average value of 8 nm for the NiO(111) single-crystal domain size was derived from GIXD data obtained on passivated Ni(111) surfaces. This difference can be assigned to the fact that STM and AFM measurements are unable to resolve the multiple twin or subgrain boundaries that may exist in the passive film if they do not markedly emerge at the topmost surface of the film. On iron, the lateral grain size was determined to be 5 nm and 6–8 nm from the in situ GIXD data obtained on passivated Fe(110) and Fe(100) surfaces, respectively. On alloys such as stainless steels, not only the structure exhibits nanoscale defects such as intergranular sites but the most recent investigations suggest that the Cr enrichment in the passive film, which is a key parameter for the corrosion resistance, is not homogeneous at the nanometer scale, and varies between the oxide nanograins themselves and also depends on the coordination sites (steps vs. terraces) of the substrate as a result of the growth mechanism of the passive film. Combined surface chemical analysis by ToF-SIMS and XPS of a FeCrNiMo(100) single-crystal alloy surface passivated in acid solution shows that a duplex hydroxylated oxide matrix, 1.8-1.9 nm thick, is formed with a strong partition between Cr(III) and Fe(II)-Fe(III) in the inner and outer layers, respectively (Fig. 4). Cr(III) is increasingly enriched upon passivation and aging, from 57% to 72% Cr(III) and from 41% to 23% Fe(II)–Fe(III) as obtained from the XPS analysis, due to preferential iron oxide dissolution. Ni, only present as oxide traces in the film, is enriched in the alloy underneath. Mo, mostly present as Mo(IV) in the Cr-rich inner layer prior to anodic polarization, becomes increasingly enriched (from 2% up to 4%–5%) and concentrated in the Fe-rich outer layer of the passive film, mostly as Mo(VI). Metallic Mo is not significantly enriched below the passive film produced from the native oxide covered surface. STM analysis reveals a typical terrace and step topography of the alloy surface with the covering oxide film presenting a nanogranular morphology (lateral grain size of 11.5 2.6 nm) (Fig. 4C). Depressions are also evidenced that result from transient dissolution that compete with the transformation of the initial native oxide film into the passive oxide film upon the passivation treatment. Since the oxide becomes further enriched in chromium after passivation, dissolution is a marker for the least Cr-enriched areas of the passivated surface. Dissolution is preferentially located on the terraces and not at the step edges of the substrate as observed in the active state in the absence of a protecting oxide film. This difference arises from the competing passivation by Cr enrichment of the oxide film, favored at step edges and thus promoting the protection of these substrate sites. As a result, dissolution is re-located on the substrate sites where the protection provided by the oxide film is less effective because of a lower local Cr enrichment on the oxide film. These data lead to the conclusion that the homogeneity of the Cr enrichment in the passive film is a key issue of the local corrosion resistance.
Passivity Breakdown and Initiation of Localized Corrosion Passive films are sensitive to local breakdown, eventually leading, in the presence of aggressive species (e.g., chlorides), to accelerated dissolution of the metallic substrate at localized sites (e.g., pitting), whereas the rest of the surface remains well protected. Local passivity breakdown is the first stage in the process leading to localized corrosion by pitting. Pits can be preferential sites for crack initiation. The grain boundaries separating the oxide grains play a key role in the breakdown of passivity and the initiation of localized corrosion as revealed by STM and AFM studies performed on single-crystal substrate surfaces. These nanostructure defects act as preferential sites for passivity breakdown and pit initiation both without and with aggressive anions (Cl) present in the electrolyte, as illustrated in Fig. 5 for passivated Ni(111) surfaces. After prepassivation in a chloride-free sulfuric acid solution in the lower part of the passive domain (0.55 V/SHE), the surface is completely covered by the polycrystalline passive film (Fig. 5A). This prepassivated surface corrodes locally with formation of nanopits when the potential is increased (up to 1.05 V/SHE) in the passive range. In the absence of chlorides, the nanopits (some are circled in Fig. 5B) observed between the grains have lateral dimensions of 20–30 nm at the surface. Their measured depth (2.2–3.8 nm) is larger than the variation of the surface level at grain boundaries after prepassivation (0.4–1.4 nm), and also larger than the thickness of the passive film formed in these conditions (< 2 nm). This implies that the prepassivated surface has been locally depassivated at grain boundaries of the passive
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Fig. 4 Passivation of stainless steel in 0.05 M H2SO4(aq) as observed on Fe-l7Cr-l4.5Ni-2.3Mo(100). (A) Depth profiles of oxidized chromium as measured by ToF-SIMS for surfaces covered by the native oxide film (Nat.) and after passivation at 0.5 V/SHE (Pass.) for 2 and 20 h. The Cr enrichment of the inner part of the surface oxide is enhanced after passivation. (B) Cr2p3/2 core level as measured by XPS after passivation at 0.5 V/SHE (2 h). The oxide film contains Cr(III) oxide and hydroxide species. (C) Cation fraction in the surface oxide film as obtained by XPS. (D) Nanoscale morphology observed by STM after passivation at 0.5 V/SHE (2 h). The oxide passive film has a granular morphology (some grains are pointed) and covers the substrate terraces and step edges. Substrate terraces display depressions (some are circled) evidencing local protection failure caused by competing transient dissolution during passivation. Substrate step edges (marked by dashed lines) are more corrosion resistant owing to preferential local Cr enrichment of the passive film. (A,B,D) Adapted with permission from Maurice, V.; Peng, H.; Klein, L.H.; Seyeux, A.; Zanna, S.; Marcus, P.; Effects of Molybdenum on the Composition and Nanoscale Morphology of Passivated Austenitic Stainless Steel Surfaces. Faraday Discuss. 2015, 180, 151–170. Copyright 2015, Royal Society of Chemistry.
film, with a local and transient enhancement of the corrosion of the substrate, and subsequently repassivated. Thus nanopits are formed (and repassivated) after breakdown at the less resistive grain boundary sites of the passive film even in the absence of chlorides. The presence of chlorides promotes the growth of some of the nanopits (Fig. 5C). The nanopits observed between the grains have, for the most part, the same dimensions as those formed in the chloride-free electrolyte. However, significantly larger ones are also observed (some are marked) with lateral dimension of 40–50 nm at the surface and a depth of 5–6 nm. Their density ( 2 109 cm 2) is about one order of magnitude lower than that of the smaller nanopits, showing that only a fraction of the less resistive sites of the passive film is impacted by the effect of chlorides during the depassivation/repassivation process. The STM image in Fig. 5D confirms passivity breakdown at the grain boundaries of the passive film, here a triple joint between oxide grains. Local electronic properties probed by scanning tunneling spectroscopy (STS) on passivated nickel showed a markedly different surface density of states explained by the presence of oxygen vacancies at grain boundaries of the oxide passive film. Anionic transport and chloride entry can be expected to be locally promoted in these sites.
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Fig. 5 Nanoscale pit initiation as observed by AFM (A–C) and STM (D) on Ni(111). Surface after prepassivation for 30 min in Cl-free 0.05 M H2SO4 þ 0.095 M NaOH (pH 2.9) at 0.55 V/SHE (A), and subsequent increase of the potential up to 1.05 V/SHE in the absence (B) or presence (C) of chloride (0.05 M NaCl). Surface prepassivated in Cl-free 0.05 M H2SO4 þ 0.095 M NaOH (pH 2.9) at 0.95 V/SHE and subsequently exposed to chlorides (0.05 M NaCl) (D). (A–C) Reprinted from Marcus, P.; Strehblow, H.-H.; Maurice, V. Localized Corrosion (Pitting): A Model of Passivity Breakdown Including the Role of the Oxide Layer Nanostructure. Corros. Sci. 2008, 50, 2698–2704. Copyright 2008, Elsevier. (D) Adapted with permission from Maurice, V.; Inard, V.; Marcus, P. STM Investigation of the Localized Corrosion of Passivated Ni(111) Single-Crystal Surfaces in Critical Factors. In: Natishan, P.M.; Kelly, R.G.; Frankel, G.S.; Newman, R.C. (eds.) Localized Corrosion III. The Electrochemical Society Proceedings Series. Pennington, NJ: 1999; PV 98–17: pp. 552–562. Copyright 1999, The Electrochemical Society.
Conclusion and Outlook Surface analytical methods applied to metal and alloy surfaces exposed to corroding liquid environments in well-controlled conditions enable to investigate the structure and stability of passive films providing self-protection against corrosion and the instability at the origin of localized corrosion. This surface science approach of passivity has established that oxide passive films are in most cases polycrystalline with grains of nanometer lateral dimensions and textured. The grains expose a facetted surface owing to a few degree tilt of the oxide lattice with respect of the metal lattice accommodating the mismatch between the oxide and the metal lattices. The oxide surface is stabilized by hydroxylation. Hydroxide ions adsorption precedes 3D oxide growth and produces 2D overlayer structures of adsorbed hydroxyl groups acting as structural precursors for the growth of 3D passive films. Aging under polarization is critical to the crystallization of chromium-rich passive films. On stainless steel, the Cr enrichment of the passive films may be inhomogeneous at the nanometer scale, influencing the local resistance to localized corrosion initiation. The step edges of the passive film surface are preferential sites of dissolution, causing a step flow dissolution mechanism in the passive state, promoted by the formation of metal hydroxychloride complexes or metal chlorides, as simulated by DFT. At the nanometer scale, grain boundaries in polycrystalline oxide passive films on metals act as preferential nanostructural defects for passivity breakdown and localized corrosion initiation. Nonuniform composition of the passive film creates weak points in the passive film on Crcontaining alloys. Despite its success, this approach has been developed so far by a small number of research groups in the corrosion science community, mostly because of the difficulty to prepare model solid/liquid interfaces well controlled at atomic or nanometric scale, particularly for highly reactive metals like Ti, Al, Mg, and their alloys. Combining UHV facilities for surface preparation and characterization with in situ electrochemical testing of the corrosion properties and characterization of the corrosioninduced alterations will enable further progress. A prerequisite to such a combined surface science and electrochemistry approach
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is to ensure no or minimum or no alteration of the surface during return transfer between UHV and electrolyte. The design of model interfaces relevant for the studied corrosion process and the availability of equipments combining UHV surface characterization techniques (microscopy and spectroscopy) with electrochemistry is a key to the success of this approach.
See also: Kinetics of Oxide Growth of Passive Films on Transition Metals; Passivation of Steel and Stainless Steel in Alkaline Media Simulating Concrete; Passivity of Metals and the Kelvin Probe Technique.
Further Reading Bouzoubaa, A.; Diawara, B.; Maurice, V.; Minot, C.; Marcus, P. Ab Initio Modelling of Localized Corrosion: Study of the Role of Surface Steps in the Interaction of Chlorides With Passivated Nickel Surfaces. Corros. Sci. 2009, 51, 2174–2182. Davenport, A. J.; Oblonsky, L. J.; Ryan, M. P.; Toney, M. F. The Structure of the Passive Film That Forms on Iron in Aqueous Environments. J. Electrochem. Soc. 2000, 147, 2162–2173. Frankel, G. S. Pitting Corrosion of Metals a Review of the Critical Factors. J. Electrochem. Soc. 1998, 145, 2186–2198. Macdonald, D. D. Passivitydthe Key to Our Metals-Based Civilization. Pure Appl. Chem. 1999, 71, 951–978. Machet, A.; Galtayries, A.; Zanna, S.; Klein, L. H.; Maurice, V.; Jolivet, P.; Foucault, M.; Combrade, P.; Scott, P.; Marcus, P. XPS and STM Study of the Growth and Structure of Passive Films in High Temperature Water on a Nickel-Base Alloy. Electrochim. Acta 2004, 49, 3957–3964. Marcus, P.; Strehblow, H.-H.; Maurice, V. Localized Corrosion (Pitting): A Model of Passivity Breakdown Including the Role of the Oxide Layer Nanostructure. Corrosion Sci. 2008, 50, 2698–2704. Marcus, P., Ed. Corrosion Mechanisms in Theory and Practice, 3rd ed.; CRC Press, Taylor and Francis: Boca Raton, FL, 2011. Marcus, P., Mansfeld, F., Eds.; Analytical Methods in Corrosion Science and Engineering, CRC Press, Taylor and Francis: Boca Raton, FL, 2006. Marcus, P.; Maurice, V. Oxide Passive Films and Corrosion Protection. In Oxide Ultrathin Films. Science and Technology; Pacchioni, G., Valeri, S., Eds.; Wiley-VCH Verlag GmbH & Co. KGaA: Weinheim, 2012; pp 119–144. Massoud, T.; Wiame, F.; Klein, L. H.; Seyeux, A.; Marcus, P. Local Electronic Properties of the Passive Film on Nickel Studied by Scanning Tunneling Spectroscopy. J. Electrochem. Soc. 2012, 159, C351–C356. Maurice, V.; Marcus, P. Passive Films at the Nanoscale. Electrochim. Acta 2012, 84, 129–138. Olsson, C. O. A.; Landolt, D. Passive Films on Stainless SteelsdChemistry, Structure and Growth. Electrochim. Acta 2003, 48, 1093–1104. Scherer, J.; Ocko, B. M.; Magnussen, O. M. Structure, Dissolution, and Passivation of Ni(111) Electrodes in Sulfuric Acid Solution: An In Situ STM, X-Ray Scattering, and Electrochemical Study. Electrochim. Acta 2003, 48, 1169–1191. Soltis, J. Passivity Breakdown, Pit Initiation and Propagation of Pits in Metallic Materials–Review. Corrosion Sci. 2015, 90, 5–22. Strehblow, H.-H.; Maurice, V.; Marcus, P. Passivity of Metals. In Corrosion Mechanisms in Theory and Practice; Marcus, P., Ed.; 3rd Ed.; CRC Press, Taylor and Francis: Boca Raton, FL, 2011; pp 235–326. Strehblow, H.-H.; Marcus, P. Mechanisms of Pitting Corrosion. In Corrosion Mechanisms in Theory and Practice; Marcus, P., Ed.; 3rd Ed.; CRC Press, Taylor and Francis: Boca Raton, FL, 2011; pp 349–393. Strehblow, H.-H. Passivity of Metals Studied by Surface Analytical Methods, a Review. Electrochim. Acta 2016, 212, 630–648. Szklarska-Smialowska, Z. ZS-Smialowska. Pitting and Crevice Corrosion, NACE International: Houston, TX, 2005. Taylor, C. D.; Marcus, P. Molecular Modeling of Corrosion ProcessesdScientific Development and Engineering Applications, John Wiley & Sons: Hoboken, 2015. Zavadil, K. R.; Ohlhausen, J. A.; Kotula, P. G. Nanoscale Void Nucleation and Growth in the Passive Oxide on Aluminum as a Prepitting Process. J. Electrochem. Soc. 2006, 153, B296–B303.
Introduction Structure of Passive Films Dissolution of Passive Films Nanostructure of Passive Films Passivity Breakdown and Initiation of Localized Corrosion Conclusion and Outlook Further Reading
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Nomenclature AFM DFT EC-STM fcc GIXD ML SHE STM STS ToF-SIMS UHV XPS
Atomic force microscopy Density functional theory Electrochemical scanning tunneling microscopy Face centered cubic Grazing incidence X-ray diffraction Monolayer Saturated hydrogen electrode Scanning tunneling microscopy Scanning tunneling spectroscopy Time-of-flight secondary ion mass spectrometry Ultra high vacuum X-ray photoelectron spectroscopy
Introduction The formation of passive films on metals and alloys is a central issue in corrosion science and engineering since it is the best of all means for protection of metallic materials against corrosion, and a key for their use in our environment. Passivity originates from the growth of an anodic oxide film on the surface of metal and alloy substrates in aqueous environment. The oxide film is either thermodynamically stable, or it dissolves very slowly, so that it protects the metal against corrosion. The oxide passive films most commonly do not exceed a few nanometers in thickness, are hydroxylated, well adherent, and effectively isolate the substrate from the corrosive environment. They are self-healing and may repair after breakdown. Surface analytical methods combined with electrochemical methods have been widely used to characterize the passivity of numerous metals and alloys and determine the thickness, composition, and electronic properties of passive films. Insight in the atomic structure of passive films and their morphology at the nanometer scale can be obtained using singlecrystal surfaces, prepared atomically flat, exposed to well-defined corrosive aqueous solutions, and investigated in situ under electrochemical control or ex situ after well-controlled electrochemical treatment and transfer of the sample to the surface analysis instrument. Scanning tunneling microscopy and atomic force microscopy, implemented (EC-STM, EC-AFM) or not (STM, AFM) with electrochemical control, as well as in situ grazing incidence X-ray diffraction (GIXD) with synchrotron light source have been applied on pure metal substrates like Cu, Ni, Fe, Cr, and Co, and on alloy substrates like stainless steel and nickel-based alloys. Passive films are polycrystalline with grains of nanometric lateral dimensions due to a high density of sites for oxide nucleation. The microstructure is textured by the preferential crystallographic orientation adopted by the oxide grains. Nonordered areas can form between the crystalline grains, particularly during the growth process, that is, under nonstationary conditions. These intergranular sites of the passive film play a key role in the breakdown of passivity and the initiation of localized corrosion. For Cr-rich passive films, the development of crystallinity is conditioned by electrode potential and aging under polarization. Structural characterization of passive films is usually complemented by surface chemical analysis using photoelectron spectroscopy (XPS) and ion spectrometry (ToF-SIMS).
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Structure of Passive Films Hydroxide ions from the liquid water adsorb on the surface in the potential range preceding 3D anodic oxide growth. On Cu(111), this process occurs at 0.6 V/SHE in 0.1 M NaOH (Fig. 1A). It induces surface reconstruction and formation of a 2D layer of adsorbed hydroxide (OHads). At low oversaturation potential for anodic oxide formation ( 0.25 V/SHE), the metal substrate is partially covered by poorly crystallized and one monolayer thick oxide islands formed after preferential nucleation at step edges and separated by the hydroxide adlayer. At higher oversaturation ( 0.2 V/SHE), the metal substrate becomes fully covered by well crystallized and several monolayer thick (111)-oriented Cu(I) oxide films (Fig. 1B). The equivalent thickness of the oxide layer deduced from charge transfer measurements during subsequent cathodic reduction is 0.5 and 7 equivalent monolayers (ML) after growth at 0.25 and 0.2 V/SHE, respectively. One ML corresponds to one (111)-oriented O2 -Cuþ-O2 slab of the cuprite structure (Fig. 1C). The 2D hydroxide adlayer is a structural precursor for the 3D passive film, adopting the same building block. The oxide layer has a hexagonal lattice with a parameter of 0.3 nm (inset in Fig. 1B), consistent with the Cu sublattice in the (111)-oriented cuprite. The fcc oxide film grows in parallel (or antiparallel) epitaxy on the fcc substrate: Cu2O(111)[1 1 0] || Cu(111)[1 1 0] or [1 10]. On Cu(001), the Cu2O passive film is also a few ML but is (001)-oriented with a 45-degree rotation between the close-packed directions of the oxide and substrate lattices: Cu2O(001)[1 1 0] || Cu(001)[100].
Fig. 1 Passivation of copper as observed on Cu(111) in 0.1 M NaOH(aq). (A) Electrochemical (voltammetry) data recorded between the hydrogen and oxygen evolution limits with marked Cu(0) active region including OH adsorption and passive Cu(I) and Cu(I)/Cu(II) regions. (B) Cu(I) oxide passive film grown as observed by EC-STM at 0.20 V/SHE. The atomic lattice (inset) observed in situ corresponds to the Cu sublattice in Cu2O(111). (C) The side view model illustrating the stacking sequence of the O2 and Cuþ planes in the oxide and the surface termination by a monolayer of hydroxyl/hydroxide groups. (A) Adapted from Kunze, J.; Maurice, V.; Klein, L.H.; Strehblow, H.-H.; Marcus, P. In Situ STM Study of the Effect of Chlorides on the Initial Stages of Anodic Oxidation of Cu (111) in Alkaline Solutions. Electrochim. Acta 2003, 48, 1157–1167. Copyright 2003, Elsevier. (B,C) Reprinted with permission from Kunze, J.; Maurice, V.; Klein, L.H.; Strehblow, H.-H.; Marcus, P. In Situ Scanning Tunneling Microscopy Study of the Anodic Oxidation of Cu (111) in 0.1 M NaoH. J. Phys. Chem. B 2001, 105, 4263–4269. Copyright 2001, American Chemical Society.
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Owing to a tilt of a few degrees of the orientation of the oxide lattice with respect to that of the metal lattice, crystalline passive film expose a faceted surface with periodic step edges as evidenced on copper by EC-STM (Fig. 1). The tilt results in part from the relaxation of the stress at the metal/oxide interface resulting from the large mismatch between the two lattices (17% on Cu). The height of the surface steps of the Cu(I) oxide layer corresponds to 1 ML of cuprite, indicating an identical chemical termination of the Cu2O oxide terraces. The oxide layer surface is hydroxylated at the interface with liquid water, as confirmed by density functional theory (DFT) modeling for the Cu2O(111) surface showing that OH adsorption on the oxide lattice stabilizes the unreconstructed structure of the Cu sublattice observed in situ by EC-STM. The structure of the duplex Cu(I)/Cu(II) passive films formed in the potential range of Cu(II) anodic oxidation is also crystalline with surface hydroxylation stabilizing the otherwise polar (001)-oriented CuO structure of the Cu(II) outer layer of the passive film. The passive film on nickel is also crystalline, as observed by EC-STM on Ni(111) in acid solution (Fig. 2). Like on copper, the passive film surface is faceted, exhibiting regular terraces and step edges due to the slightly tilted epitaxy that allows relaxation
Fig. 2 Passivation of nickel as observed on Ni(111) in 0.05 M H2SO4 þ 0.095 M NaOH(aq.) (pH 2.9). (A) Electrochemical polarization curve with the active, passive, and transpassive regions marked. (B) Ni(II) oxide passive film as observed by EC-STM at 0.95 V/SHE. The inset shows the high resolution image recorded in situ, revealing the passive film atomic lattice. (C) The side view model of the hydroxylated NiO surface with (010) step edges and (111) terraces as optimized by DFT simulation. (B) Adapted with permission from Maurice, V.; Klein, L.H.; Marcus, P. Atomic-Scale Investigation of the Localized Corrosion of Passivated Nickel Surfaces. Surf. Interf. Anal. 2002, 34, 139–143. Copyright 2002, Wiley. (C) Reprinted from Bouzoubaa, A.; Diawara, B.; Maurice, V.; Minot, C.; Marcus, P. Ab Initio Modelling of Localized Corrosion: Study of the Role of Surface Steps in the Interaction of Chlorides With Passivated Nickel Surfaces. Corros. Sci. 2009, 51, 2174–2182. Copyright 2009, Elsevier.
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of the interfacial stress associated with the lattice mismatch (16%) between the NiO(111) lattice forming the barrier oxide layer and the Ni(111) lattice of the substrate. The (111) orientation of the NiO lattice and its tilt of about 3 degree with respect to the Ni(111) orientation have been confirmed by in situ GIXD measurements. The fcc oxide grows in antiparallel epitaxy on the fcc substrate: NiO(111)[1 1 0] || Ni(111)[1 10]. The atomic lattice measured by EC-STM on the terraces is hexagonal with a parameter of 0.3 0.02 nm consistent with the O (and Ni) sublattice of NiO(111). The NiO lattice obtained by passivation is (111)-oriented despite the fact that this surface is polar and thus unstable if bulk-like terminated. DFT modeling confirmed that the reason for this preferential (111) orientation of the terraces is that the surface is stabilized by adsorption of a monolayer of hydroxyl groups. Fig. 2C shows a model for such a faceted and hydroxylated surface as optimized by DFT. The direction of growth of the oxide film is thus governed by the minimization of the interface energy by a tilted growth and by the minimization of the oxide surface energy by the hydroxyl/hydroxide group termination. On chromium and stainless alloys, including stainless steels, the formation of Cr(III)-enriched oxide passive layers provides excellent corrosion resistance. Potential and aging under polarization are critical factors for the development of crystalline passive films. Electrode potential and aging-dependent crystallization observed by STM are paralleled by dehydroxylation measured by X-ray photoelectron spectroscopy (XPS). On chromium passivated in acid solution at low potential, the passive film is highly hydrated and the Cr(III) oxide inner layer is only partially developed. The passive film consists of small Cr(III) nanocrystals buried in a disordered Cr(III) hydroxide matrix. At high potential, the inner part of the passive film is dehydrated and consists mostly of larger crystals of Cr(III) oxide. A faceted topography extending over several tens of nanometers has been observed by EC-STM. The nanocrystals have a lattice consistent with the O sublattice in a-Cr2O3(0001) (corundum structure). The basal plane of the oxide is parallel to bcc Cr(110) with the closed-packed rows of the O sublattice nearly parallel to those of the metal and a mismatch of 12%. On stainless steels passivated in acid solution, the crystallinity of the passive films decreases with increasing Cr content of the alloy for short polarization times ( 2 h). Structural changes also occur during aging under anodic polarization. The major structural modification is an increase of the crystallinity of the film and the coalescence of Cr(III) oxide nanocrystals in the inner oxide as observed on Fe-22Cr and Fe-18Cr-13Ni alloys studied over time periods of up to 65 h. Like on chromium, the measured atomic lattice is consistent with the O sublattice of the (0001)-oriented a-Cr2O3 corundum structure. The rate of crystallization is more rapid on the austenitic stainless steel than on the ferritic one.
Dissolution of Passive Films The steps at the surface of the passive film play a key role for dissolution in the passive state, and thus for the stability of passivated surfaces. The passive film dissolves at the edges of the facets produced by the tilted epitaxy between the oxide and metal lattices (Fig. 2), which produces a receding step flow. The process is similar to that of active dissolution of oxide-free metal surfaces at moderate potential. The receding 2D step flow is dependent on the step orientation since the step edges oriented along the closed-packed directions of the oxide lattice dissolve much less rapidly, due to the higher coordination of their atoms. This structure-dependent process leads to the stabilization of the facets with edges oriented along the close-packed directions of the oxide lattice. On passivated nickel, it produces steps that are oriented along the {100} planes, the most stable orientations of the NiO structure. This receding step flow dissolution mechanism is observed independently of the presence or absence of chloride anions in the electrolyte. Dissolution of the passive film, catalyzed by chlorides complexing metal cations at the surface of the passive film, is a key aspect of the adsorption mechanism of passivity breakdown and initiation of localized corrosion. DFT modeling has been applied to hydroxylated and stepped NiO(111) surfaces in order to investigate the interaction with Cl and its effect on the dissolution of the oxide. A (533)-oriented NiO periodic model was built-up to simulate the (111) termination of the passive film terraces with monoatomic (010) step edges (Fig. 3). Cl adsorption structures were modeled by substituting the surface OH groups by Cl atoms at 25%, 50%, 75% and 100% coverages. After DFT optimization, substructures of Ni(OH)2, Ni(OH)Cl, or Ni(Cl)2 composition
Fig. 3 DFT modeling of the hydroxylated NiO(533) surface including a (010) step edge and a (111) terrace. Side view models of Cl-adsorbed structures at 25% coverage are shown before and after optimization. Adapted from Bouzoubaa, A.; Diawara, B.; Maurice, V.; Minot, C.; Marcus, P. Ab Initio Modelling of Localized Corrosion: Study of the Role of Surface Steps in the Interaction of Chlorides With Passivated Nickel Surfaces. Corros. Sci. 2009, 51, 2174–2182. Copyright 2009, Elsevier.
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were observed to form and detach from the step edges of the adsorbed structures, confirming the major role of the step edges in the dissolution of the oxide. The calculated energies of detachment of the substructures revealed that the Cl-containing substructures are easier to detach, showing that dissolution at the step edges can be promoted by the adsorption of Cl and confirming a Cl adsorption-induced oxide thinning as a possible breakdown mechanism of the oxide passive film.
Nanostructure of Passive Films Even when grown on single-crystal substrates, passive films have a polycrystalline microstructure with grains of nanometric lateral dimensions. This is a result of a high density of sites for oxide nucleation in their initial stages of formation. The lateral dimensions of the crystalline grains forming the passive films are well documented by STM, AFM, and GIXD data for nickel. Values determined from the morphology observed by STM and AFM have been reported to range from 2 nm in the initial stages of 2D growth to 30 to 230 nm for 3D films in stationary conditions of passivity. A large dispersion could be found on the same sample, suggesting a varying degree of advancement of the coalescence of the oxide grains during growth after nucleation. A lower average value of 8 nm for the NiO(111) single-crystal domain size was derived from GIXD data obtained on passivated Ni(111) surfaces. This difference can be assigned to the fact that STM and AFM measurements are unable to resolve the multiple twin or subgrain boundaries that may exist in the passive film if they do not markedly emerge at the topmost surface of the film. On iron, the lateral grain size was determined to be 5 nm and 6–8 nm from the in situ GIXD data obtained on passivated Fe(110) and Fe(100) surfaces, respectively. On alloys such as stainless steels, not only the structure exhibits nanoscale defects such as intergranular sites but the most recent investigations suggest that the Cr enrichment in the passive film, which is a key parameter for the corrosion resistance, is not homogeneous at the nanometer scale, and varies between the oxide nanograins themselves and also depends on the coordination sites (steps vs. terraces) of the substrate as a result of the growth mechanism of the passive film. Combined surface chemical analysis by ToF-SIMS and XPS of a FeCrNiMo(100) single-crystal alloy surface passivated in acid solution shows that a duplex hydroxylated oxide matrix, 1.8-1.9 nm thick, is formed with a strong partition between Cr(III) and Fe(II)-Fe(III) in the inner and outer layers, respectively (Fig. 4). Cr(III) is increasingly enriched upon passivation and aging, from 57% to 72% Cr(III) and from 41% to 23% Fe(II)–Fe(III) as obtained from the XPS analysis, due to preferential iron oxide dissolution. Ni, only present as oxide traces in the film, is enriched in the alloy underneath. Mo, mostly present as Mo(IV) in the Cr-rich inner layer prior to anodic polarization, becomes increasingly enriched (from 2% up to 4%–5%) and concentrated in the Fe-rich outer layer of the passive film, mostly as Mo(VI). Metallic Mo is not significantly enriched below the passive film produced from the native oxide covered surface. STM analysis reveals a typical terrace and step topography of the alloy surface with the covering oxide film presenting a nanogranular morphology (lateral grain size of 11.5 2.6 nm) (Fig. 4C). Depressions are also evidenced that result from transient dissolution that compete with the transformation of the initial native oxide film into the passive oxide film upon the passivation treatment. Since the oxide becomes further enriched in chromium after passivation, dissolution is a marker for the least Cr-enriched areas of the passivated surface. Dissolution is preferentially located on the terraces and not at the step edges of the substrate as observed in the active state in the absence of a protecting oxide film. This difference arises from the competing passivation by Cr enrichment of the oxide film, favored at step edges and thus promoting the protection of these substrate sites. As a result, dissolution is re-located on the substrate sites where the protection provided by the oxide film is less effective because of a lower local Cr enrichment on the oxide film. These data lead to the conclusion that the homogeneity of the Cr enrichment in the passive film is a key issue of the local corrosion resistance.
Passivity Breakdown and Initiation of Localized Corrosion Passive films are sensitive to local breakdown, eventually leading, in the presence of aggressive species (e.g., chlorides), to accelerated dissolution of the metallic substrate at localized sites (e.g., pitting), whereas the rest of the surface remains well protected. Local passivity breakdown is the first stage in the process leading to localized corrosion by pitting. Pits can be preferential sites for crack initiation. The grain boundaries separating the oxide grains play a key role in the breakdown of passivity and the initiation of localized corrosion as revealed by STM and AFM studies performed on single-crystal substrate surfaces. These nanostructure defects act as preferential sites for passivity breakdown and pit initiation both without and with aggressive anions (Cl) present in the electrolyte, as illustrated in Fig. 5 for passivated Ni(111) surfaces. After prepassivation in a chloride-free sulfuric acid solution in the lower part of the passive domain (0.55 V/SHE), the surface is completely covered by the polycrystalline passive film (Fig. 5A). This prepassivated surface corrodes locally with formation of nanopits when the potential is increased (up to 1.05 V/SHE) in the passive range. In the absence of chlorides, the nanopits (some are circled in Fig. 5B) observed between the grains have lateral dimensions of 20–30 nm at the surface. Their measured depth (2.2–3.8 nm) is larger than the variation of the surface level at grain boundaries after prepassivation (0.4–1.4 nm), and also larger than the thickness of the passive film formed in these conditions (< 2 nm). This implies that the prepassivated surface has been locally depassivated at grain boundaries of the passive
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Fig. 4 Passivation of stainless steel in 0.05 M H2SO4(aq) as observed on Fe-l7Cr-l4.5Ni-2.3Mo(100). (A) Depth profiles of oxidized chromium as measured by ToF-SIMS for surfaces covered by the native oxide film (Nat.) and after passivation at 0.5 V/SHE (Pass.) for 2 and 20 h. The Cr enrichment of the inner part of the surface oxide is enhanced after passivation. (B) Cr2p3/2 core level as measured by XPS after passivation at 0.5 V/SHE (2 h). The oxide film contains Cr(III) oxide and hydroxide species. (C) Cation fraction in the surface oxide film as obtained by XPS. (D) Nanoscale morphology observed by STM after passivation at 0.5 V/SHE (2 h). The oxide passive film has a granular morphology (some grains are pointed) and covers the substrate terraces and step edges. Substrate terraces display depressions (some are circled) evidencing local protection failure caused by competing transient dissolution during passivation. Substrate step edges (marked by dashed lines) are more corrosion resistant owing to preferential local Cr enrichment of the passive film. (A,B,D) Adapted with permission from Maurice, V.; Peng, H.; Klein, L.H.; Seyeux, A.; Zanna, S.; Marcus, P.; Effects of Molybdenum on the Composition and Nanoscale Morphology of Passivated Austenitic Stainless Steel Surfaces. Faraday Discuss. 2015, 180, 151–170. Copyright 2015, Royal Society of Chemistry.
film, with a local and transient enhancement of the corrosion of the substrate, and subsequently repassivated. Thus nanopits are formed (and repassivated) after breakdown at the less resistive grain boundary sites of the passive film even in the absence of chlorides. The presence of chlorides promotes the growth of some of the nanopits (Fig. 5C). The nanopits observed between the grains have, for the most part, the same dimensions as those formed in the chloride-free electrolyte. However, significantly larger ones are also observed (some are marked) with lateral dimension of 40–50 nm at the surface and a depth of 5–6 nm. Their density ( 2 109 cm 2) is about one order of magnitude lower than that of the smaller nanopits, showing that only a fraction of the less resistive sites of the passive film is impacted by the effect of chlorides during the depassivation/repassivation process. The STM image in Fig. 5D confirms passivity breakdown at the grain boundaries of the passive film, here a triple joint between oxide grains. Local electronic properties probed by scanning tunneling spectroscopy (STS) on passivated nickel showed a markedly different surface density of states explained by the presence of oxygen vacancies at grain boundaries of the oxide passive film. Anionic transport and chloride entry can be expected to be locally promoted in these sites.
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Fig. 5 Nanoscale pit initiation as observed by AFM (A–C) and STM (D) on Ni(111). Surface after prepassivation for 30 min in Cl-free 0.05 M H2SO4 þ 0.095 M NaOH (pH 2.9) at 0.55 V/SHE (A), and subsequent increase of the potential up to 1.05 V/SHE in the absence (B) or presence (C) of chloride (0.05 M NaCl). Surface prepassivated in Cl-free 0.05 M H2SO4 þ 0.095 M NaOH (pH 2.9) at 0.95 V/SHE and subsequently exposed to chlorides (0.05 M NaCl) (D). (A–C) Reprinted from Marcus, P.; Strehblow, H.-H.; Maurice, V. Localized Corrosion (Pitting): A Model of Passivity Breakdown Including the Role of the Oxide Layer Nanostructure. Corros. Sci. 2008, 50, 2698–2704. Copyright 2008, Elsevier. (D) Adapted with permission from Maurice, V.; Inard, V.; Marcus, P. STM Investigation of the Localized Corrosion of Passivated Ni(111) Single-Crystal Surfaces in Critical Factors. In: Natishan, P.M.; Kelly, R.G.; Frankel, G.S.; Newman, R.C. (eds.) Localized Corrosion III. The Electrochemical Society Proceedings Series. Pennington, NJ: 1999; PV 98–17: pp. 552–562. Copyright 1999, The Electrochemical Society.
Conclusion and Outlook Surface analytical methods applied to metal and alloy surfaces exposed to corroding liquid environments in well-controlled conditions enable to investigate the structure and stability of passive films providing self-protection against corrosion and the instability at the origin of localized corrosion. This surface science approach of passivity has established that oxide passive films are in most cases polycrystalline with grains of nanometer lateral dimensions and textured. The grains expose a facetted surface owing to a few degree tilt of the oxide lattice with respect of the metal lattice accommodating the mismatch between the oxide and the metal lattices. The oxide surface is stabilized by hydroxylation. Hydroxide ions adsorption precedes 3D oxide growth and produces 2D overlayer structures of adsorbed hydroxyl groups acting as structural precursors for the growth of 3D passive films. Aging under polarization is critical to the crystallization of chromium-rich passive films. On stainless steel, the Cr enrichment of the passive films may be inhomogeneous at the nanometer scale, influencing the local resistance to localized corrosion initiation. The step edges of the passive film surface are preferential sites of dissolution, causing a step flow dissolution mechanism in the passive state, promoted by the formation of metal hydroxychloride complexes or metal chlorides, as simulated by DFT. At the nanometer scale, grain boundaries in polycrystalline oxide passive films on metals act as preferential nanostructural defects for passivity breakdown and localized corrosion initiation. Nonuniform composition of the passive film creates weak points in the passive film on Crcontaining alloys. Despite its success, this approach has been developed so far by a small number of research groups in the corrosion science community, mostly because of the difficulty to prepare model solid/liquid interfaces well controlled at atomic or nanometric scale, particularly for highly reactive metals like Ti, Al, Mg, and their alloys. Combining UHV facilities for surface preparation and characterization with in situ electrochemical testing of the corrosion properties and characterization of the corrosioninduced alterations will enable further progress. A prerequisite to such a combined surface science and electrochemistry approach
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is to ensure no or minimum or no alteration of the surface during return transfer between UHV and electrolyte. The design of model interfaces relevant for the studied corrosion process and the availability of equipments combining UHV surface characterization techniques (microscopy and spectroscopy) with electrochemistry is a key to the success of this approach.
See also: Kinetics of Oxide Growth of Passive Films on Transition Metals; Passivation of Steel and Stainless Steel in Alkaline Media Simulating Concrete; Passivity of Metals and the Kelvin Probe Technique.
Further Reading Bouzoubaa, A.; Diawara, B.; Maurice, V.; Minot, C.; Marcus, P. Ab Initio Modelling of Localized Corrosion: Study of the Role of Surface Steps in the Interaction of Chlorides With Passivated Nickel Surfaces. Corros. Sci. 2009, 51, 2174–2182. Davenport, A. J.; Oblonsky, L. J.; Ryan, M. P.; Toney, M. F. The Structure of the Passive Film That Forms on Iron in Aqueous Environments. J. Electrochem. Soc. 2000, 147, 2162–2173. Frankel, G. S. Pitting Corrosion of Metals a Review of the Critical Factors. J. Electrochem. Soc. 1998, 145, 2186–2198. Macdonald, D. D. Passivitydthe Key to Our Metals-Based Civilization. Pure Appl. Chem. 1999, 71, 951–978. Machet, A.; Galtayries, A.; Zanna, S.; Klein, L. H.; Maurice, V.; Jolivet, P.; Foucault, M.; Combrade, P.; Scott, P.; Marcus, P. XPS and STM Study of the Growth and Structure of Passive Films in High Temperature Water on a Nickel-Base Alloy. Electrochim. Acta 2004, 49, 3957–3964. Marcus, P.; Strehblow, H.-H.; Maurice, V. Localized Corrosion (Pitting): A Model of Passivity Breakdown Including the Role of the Oxide Layer Nanostructure. Corrosion Sci. 2008, 50, 2698–2704. Marcus, P., Ed. Corrosion Mechanisms in Theory and Practice, 3rd ed.; CRC Press, Taylor and Francis: Boca Raton, FL, 2011. Marcus, P., Mansfeld, F., Eds.; Analytical Methods in Corrosion Science and Engineering, CRC Press, Taylor and Francis: Boca Raton, FL, 2006. Marcus, P.; Maurice, V. Oxide Passive Films and Corrosion Protection. In Oxide Ultrathin Films. Science and Technology; Pacchioni, G., Valeri, S., Eds.; Wiley-VCH Verlag GmbH & Co. KGaA: Weinheim, 2012; pp 119–144. Massoud, T.; Wiame, F.; Klein, L. H.; Seyeux, A.; Marcus, P. Local Electronic Properties of the Passive Film on Nickel Studied by Scanning Tunneling Spectroscopy. J. Electrochem. Soc. 2012, 159, C351–C356. Maurice, V.; Marcus, P. Passive Films at the Nanoscale. Electrochim. Acta 2012, 84, 129–138. Olsson, C. O. A.; Landolt, D. Passive Films on Stainless SteelsdChemistry, Structure and Growth. Electrochim. Acta 2003, 48, 1093–1104. Scherer, J.; Ocko, B. M.; Magnussen, O. M. Structure, Dissolution, and Passivation of Ni(111) Electrodes in Sulfuric Acid Solution: An In Situ STM, X-Ray Scattering, and Electrochemical Study. Electrochim. Acta 2003, 48, 1169–1191. Soltis, J. Passivity Breakdown, Pit Initiation and Propagation of Pits in Metallic Materials–Review. Corrosion Sci. 2015, 90, 5–22. Strehblow, H.-H.; Maurice, V.; Marcus, P. Passivity of Metals. In Corrosion Mechanisms in Theory and Practice; Marcus, P., Ed.; 3rd Ed.; CRC Press, Taylor and Francis: Boca Raton, FL, 2011; pp 235–326. Strehblow, H.-H.; Marcus, P. Mechanisms of Pitting Corrosion. In Corrosion Mechanisms in Theory and Practice; Marcus, P., Ed.; 3rd Ed.; CRC Press, Taylor and Francis: Boca Raton, FL, 2011; pp 349–393. Strehblow, H.-H. Passivity of Metals Studied by Surface Analytical Methods, a Review. Electrochim. Acta 2016, 212, 630–648. Szklarska-Smialowska, Z. ZS-Smialowska. Pitting and Crevice Corrosion, NACE International: Houston, TX, 2005. Taylor, C. D.; Marcus, P. Molecular Modeling of Corrosion ProcessesdScientific Development and Engineering Applications, John Wiley & Sons: Hoboken, 2015. Zavadil, K. R.; Ohlhausen, J. A.; Kotula, P. G. Nanoscale Void Nucleation and Growth in the Passive Oxide on Aluminum as a Prepitting Process. J. Electrochem. Soc. 2006, 153, B296–B303.