Passive films at the nanoscale

Electrochimica Acta 84 (2012) 129–138

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Review

Passive films at the nanoscale Vincent Maurice ∗ , Philippe Marcus ∗ CNRS - Chimie ParisTech (UMR 7045), Laboratoire de Physico-Chimie des Surfaces (LPCS), Ecole Nationale Supérieure de Chimie de Paris (ENSCP), 11 rue Pierre et Marie Curie, F-75005 Paris, France

a r t i c l e

i n f o

Article history: Received 14 December 2011 Received in revised form 30 March 2012 Accepted 31 March 2012 Available online 6 April 2012 Keywords: Corrosion Nanoscale Passivation Metals Alloys

a b s t r a c t The nanometer scale chemical and structural aspects of ultrathin oxide passive films providing selfprotection against corrosion to metals and alloys in aqueous environments are reviewed. Data on the nucleation and growth of 2D anodic oxide films, details on the atomic structure and nanostructure of 3D passive films, the preferential role of surface step edges in dissolution in the passive state and the preferential role of grain boundaries of the passive films in passivity breakdown are presented. Future perspectives are discussed, and exemplified by new data obtained on the relationship between the nanostructure of oxide passive films and their local electronic properties. Atomistic corrosion modeling by ab initio density functional theory (DFT) is illustrated by the example of interactions of chloride ions with hydroxylated oxide surfaces, including the role of surface step edges. Data obtained on well-defined substrate surfaces with surface analytical techniques are emphasized. © 2012 Elsevier Ltd. All rights reserved.

Contents 1. 2.

3. 4.

Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Passive films at the nanoscale . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.1. Nucleation and 2D growth of passive films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.2. 3D growth and structure of passive films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.3. Nanostructure and relation with passivity breakdown . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 2.4. Electronic properties of grain boundaries in passive films . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Atomistic modeling of passivity breakdown . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . Conclusion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . References . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

1. Introduction Self-protection against corrosion of metals and alloys in aqueous environments is provided by the growth of oxide passive films. This surface property of metal and alloy surfaces is essential to sustainable development in numerous applications and industries where metallic components are used. Oxide passive films most commonly do not exceed a few nanometers in thickness at room temperature, and are hydroxylated, well adherent, and effectively isolate the substrate from the corrosive environment. They are however sensitive to local breakdown eventually leading, in the presence of aggressive species (e.g. chlorides), to accelerated dissolution of the metal (or alloy) substrate at localized sites (e.g. pitting). Understanding

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the mechanisms and controlling the properties of self-protection against corrosion by surface oxide films of nanometric thickness requires a detailed knowledge of these systems at the nanometer and atomic scales, that can be obtained by experimental analytical studies and modeling studies. In the first part of this article, we review the recent insight gained at the nanoscale on the growth and structure of oxide passive films using high resolution scanning nanoprobes and then illustrate the perspectives of future studies with such techniques. In the second part we present the recent progresses in atomistic modeling of passivity breakdown and discuss the perspectives of development of such studies. 2. Passive films at the nanoscale

∗ Corresponding authors. E-mail addresses: [email protected] (V. Maurice), [email protected] (P. Marcus). 0013-4686/$ – see front matter © 2012 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.electacta.2012.03.158

Oxide passive films grown on metal and alloy surfaces have been recently reviewed [1–7]. Their structure has been studied mostly with scanning tunneling microscopy (STM) and also with

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atomic force microscopy (AFM) for Ni [8–17], Cr [18,19], Fe [20–25], Cu [26–31], Co [32,33], and stainless steel [34–38] and nickelbased [39] alloys. In situ grazing incidence X-ray diffraction (GIXD) experiments have confirmed and complemented the data for Ni [14,40] and Fe [41,42]. In many cases the passive films consist, at the nanoscale, of crystalline oxide grains joined by grain boundaries. Non-ordered areas can form between the crystalline grains, particularly during the growth process, i.e. under non-stationary conditions. These inter-granular sites of the passive film play a key role in the breakdown of passivity and the initiation of localized corrosion [43–45]. Data obtained over the last decade on copper and nickel singlecrystal surfaces are selected below to illustrate the nucleation and 2D growth mechanism, the structure of 3D oxide passive films, and the role of their nanostructure on passivity breakdown. New data obtained on the local electronic properties of passive films are then presented to illustrate the perspectives of investigation of the relationship between the nanostructure of passive films and their local properties.

2.1. Nucleation and 2D growth of passive films The formation of a 2D adlayer similar to a surface oxide can result from the adsorption of OH− in the potential range preceding 3D anodic oxidation. The structure of this 2D adlayer has been shown to be a precursor for the growth of 3D metal oxide occurring at higher potential, by in situ electrochemical STM (ECSTM) on copper [26–29] and silver [46,47] in alkaline aqueous

solutions. Fig. 1 illustrates this process for a Cu(1 1 1) single-crystal surface in 0.1 M NaOH on which 3D Cu2 O(1 1 1) growth is observed at U > −0.2 V/SHE. The sequence of STM images shown in Fig. 1a was obtained after a positive potential step to −0.6 V/SHE, i.e. at the onset of adsorption of OH− anions according to the following reaction: −

Me + OH− → Me − OHads + (1 − )e−

(1)

The initially bare and atomically smooth terraces of the surface (marked m. in Fig. 1a) become progressively covered by darker appearing 2D islands (marked ad.) that grow laterally with time and coalesce to cover completely the terraces. Initial growth of the adsorbed layer occurs preferentially at the step edges, confirming the preferential reactivity of these pre-existing defect sites of the surface, a general finding of Surface Science and Interfacial Electrochemistry studies of solid–gas and solid–liquid interfaces. Fig. 1a also reveals the displacement of the step edges and the formation of monoatomic protruding adislands (marked i.) at the end of the growth process. These two features are indicative of the reconstruction of the topmost Cu plane induced by the adsorption of the OH− anions. The reconstruction causes the ejection of Cu atoms from the topmost plane. The ejected atoms diffuse on the surface and aggregate at step edges. This causes the observed step flow. In the final stages of the adsorption process where most of the surface is already covered by the 2D adlayer, the ejected atoms have a reduced mobility on the OH-covered terraces and aggregate to form the observed adislands. Some fraction of the ejected Cu atoms may also dissolve in the electrolyte during this process.

Fig. 1. (a) Sequence of ECSTM images showing the growth of the OHads adlayer on Cu(1 1 1) in 0.1 M NaOH(aq) at −0.6 V/SHE. The adlayer islands and the bare substrate are marked ad. and m., respectively. (b) ECSTM image and model of the ordered structure of OHads formed on Cu(1 1 1) in 0.1 M NaOH(aq) in the potential region below oxidation. The large and small cells mark the lattice of OHads species and of the reconstructed CuR plane, respectively. Adapted from Refs. [27,28].

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Fig. 1b shows an atomically resolved image of the adlayer and the model of the ordered surface according to a Cu(1 1 1)/CuR /OHads topmost plane sequence. The reconstructed Cu topmost metal plane (CuR ) is directly observed. A hexagonal lattice with a parameter of 0.6 ± 0.02 nm is measured. Each unit cell contains one minimum and four maxima of intensity. The period between the maxima (assigned to Cu atoms and forming a sub-lattice) is ∼0.3 nm, which is larger than the period in the bulk-like Cu(1 1 1) plane (0.256 nm) and confirms the reconstruction of the topmost Cu plane into a lower density CuR plane. A coverage of ∼0.2 OH per Cu(1 1 1) atom is deduced from the density of the intensity minima, in excellent agreement with the coverage of 0.19 obtained from the electrochemical charge transfer measurements assuming  = 0 in Eq. (1). In addition, the STM data show that the position of the OHads groups corresponds to three-fold hollow sites of the reconstructed CuR plane on which they form a (2 × 2) superstructure. This structure of the CuR /OHads topmost plane mimics that of the Cu and O sub-lattices in (1 1 1)-oriented Cu2 O (presented below), thus forming a surface (hydr)oxide adlayer that can be viewed as a structural 2D precursor for the 3D growth of the Cu2 O(1 1 1) anodic oxide. OH-induced reconstruction of the topmost metal plane has been also observed for Cu(0 0 1) in the potential range below 3D oxidation [29] and for Ag(1 1 1) [46,47]. 2.2. 3D growth and structure of passive films The influence of the overpotential, i.e. the driving force for oxide formation, on the growth, crystallization and structure of the passive film formed on Cu(1 1 1) in the potential range of Cu(I) oxidation is illustrated in Fig. 2 [26,28]. At low oversaturation (U = −0.25 V/SHE), poorly crystallized and one monolayer thick islands covering partially the substrate are formed after preferential nucleation at step edges (Fig. 2a). They are separated by islands of the ordered adlayer described above. At higher oversaturation (U = −0.2 V/SHE), well crystallized and several monolayer thick (3D) films are formed (Fig. 2b). The equivalent thickness of the oxide layer was ∼0.5 and 7 equivalent monolayers (ML, one ML corresponds to one (1 1 1)-oriented O2− –Cu+ –O2− slab) after growth at −0.25 and −0.2 V/SHE, respectively, as determined from subsequent measurements of the charge density transfer during electrochemical reduction. The lattice of the oxide layer, observed by STM (insert in Fig. 2b), is hexagonal with a parameter of ∼0.3 nm, consistent with the Cu sub-lattice in the (1 1 1)-oriented cuprite. The oxide grows in parallel (or anti-parallel) epitaxy defined by the following relationship: Cu2 O(1 1 1) [1 1¯ 0] || Cu(1 1 1) [1 1¯ 0] or [1¯ 1 0]. On Cu(0 0 1) (Fig. 2c), the oxide layer, also a few ML thick, has a square symmetry and the same periodicity of 0.3 nm, consistent with the Cu-sub-lattice of (0 0 1)-oriented Cu2 O. The rotation of 45◦ of the close-packed [1 1¯ 0] direction of the oxide lattice with respect to the close-packed direction of the Cu(0 0 1) lattice gives an epitaxial relationship noted as Cu2 O(0 0 1) [1 1¯ 0] || Cu(0 0 1) [1 0 0]. On both substrates, the crystalline Cu(I) oxide layers have a nanostructured and faceted surface (Fig. 2). The surface faceting indicates a tilt of a few degrees of the orientation of the oxide lattices with respect to the Cu lattice as shown in Fig. 3. The tilt is thought to result, at least in part, from the relaxation of the stress at the metal/oxide interface resulting from the large mismatch between the two lattices (17%). The height of the surface steps of the oxide layers corresponds to 1 ML of cuprite, indicating an identical chemical termination of the Cu2 O(1 1 1) and Cu2 O(0 0 1) oxide terraces. The oxide layer surface is most likely hydroxylated in the aqueous solution and DFT modeling for the Cu2 O(1 1 1) surface has shown that OH adsorption on the oxide lattice stabilizes the unreconstructed structure of the Cu sub-lattice observed in situ by STM [48].

Fig. 2. ECSTM images of the Cu(I) oxide grown on Cu(1 1 1) at −0.25 V/SHE (a) and −0.20 V/SHE (b), and on Cu(0 0 1) at −0.11 V/SHE (c) in 0.1 M NaOH(aq). At low oversaturation (a), non-crystalline 2D oxide islands (ox.) separated by the adsorbed OH layer (ad) cover partially the substrate. At higher oversaturation (b, c), a 3D crystalline oxide layer fully covers the substrate. Its atomic lattice, shown in the inset, corresponds to that of Cu2 O(1 1 1) (b) and Cu2 O(0 0 1) (c) on Cu(1 1 1) and Cu(0 0 1), respectively. (a, b) Adapted with permission from Ref. [26]. Copyright 2001 American Chemical Society. (c) Reprinted from Ref. [29].

In the potential range of Cu(II) oxidation, crystalline Cu(I)/Cu(II) duplex passive films having a total thickness of 4–5 nm are formed in 0.1 M NaOH [31]. The composition has been confirmed by ex situ XPS and ISS combined with adequate transfer systems from the electrochemical cell avoiding oxygen contamination [49–52] and

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Fig. 3. Model (section view) of the 7 ML thick Cu(I) oxide grown in tilted epitaxy on Cu(1 1 1). The stacking sequence of the O2− and Cu+ planes in the oxide is illustrated by the gray lines. The oxide surface is terminated by a monolayer of hydroxyl/hydroxide groups. The faceted surface corresponds to a tilt of ∼5◦ between the two lattices. Interfacial misfit dislocations and misorientation dislocations are illustrated by the T and |— symbols, respectively. White and black symbols correspond to dislocations of the oxide and metal lattices, respectively. Adapted with permission from Ref. [26]. Copyright 2001 American Chemical Society.

by in situ techniques, mostly Raman spectroscopy [53–55]. On both Cu(1 1 1) and Cu(0 0 1), a terrace and step topography of the passivated surfaces is observed with terraces extending up to 20 nm in width and with a step height of ∼0.25 nm corresponding to the thickness of one equivalent monolayer of CuO(0 0 1). Accordingly, the atomic lattices observed are consistent with a bulk-terminated CuO(0 0 1) surface. The epitaxial relationships for the duplex layers are then: CuO(0 0 1)[1¯ 1 0]||Cu2 O(1 1 1)[1 1¯ 0]||Cu(1 1 1)[1 1¯ 0] or [1¯ 1 0]

(2)

CuO(0 0 1)[1¯ 1 0]||Cu2 O(0 0 1)[1 1¯ 0]||Cu(0 0 1)[1 0 0]

(3)

which corresponds in both cases to the parallel alignment of the closed packed directions of the CuO and Cu2 O lattices. The common (0 0 1) orientation of the CuO outer layers of the duplex films on the two substrates is assigned to their surface hydroxylation at the passive film/electrolyte interface, that would be necessary to stabilize the bulk-like termination of CuO(0 0 1) which is polar and unstable when anhydrous. This assignment is supported by the step height measurements that indicate an identical chemical termination of all terraces. It is thought that the surface is terminated by an OH− (or OH) layer on the topmost plane of the Cu2+ sub-lattice, then O2− and Cu2+ planes alternate towards the bulk of the oxide layer. For Ni(1 1 1) in 1 mM NaOH at U ≥ −0.13 V/SHE, i.e. beyond the top of the anodic peak corresponding to the growth of the passivating oxide, the 3D growth of the passive film is observed (Fig. 4a) [16]. As on copper, it is characterized by the formation of a faceted topography indicative of the tilt between the oxide lattice of the passive film and the lattice of the substrate. The hexagonal lattice at −0.13 and 0.22 V/SHE has a parameter of 0.32 ± 0.01 nm, slightly larger than that measured at −0.28 V/SHE when a 2D passivating oxide is formed. The parameter is in good agreement with the value of 0.317 nm expected for an unstrained 3D layer assigned to ␤-Ni(OH)2 (0 0 0 1). This is consistent with a crystalline 3D outer hydroxide layer of the passive film formed in alkaline electrolytes, as opposed to the 2D outer hydroxide layer existing on the inner oxide layer in acid electrolytes [56–61].

Fig. 4. ECSTM images of Ni(1 1 1) surfaces passivated (a) in 1 mM NaOH(aq) (pH ∼ 11) at U = −0.13 V/SHE and (b) in 0.05 M H2 SO4 + 0.095 M NaOH(aq.) (pH ∼ 2.9) at U = +0.95 V/SHE. The inserts show the high resolution images revealing the atomic lattices. (a) Adapted from Ref. [16] with kind permission from Springer. (b) Reproduced from Ref. [12] by permission of the Electrochemical Society.

Fig. 4b illustrates the typical topography of a Ni(1 1 1) singlecrystal surface polarized in the middle of the passive domain in a sulfuric acid solution (0.05 M H2 SO4 + 0.095 M NaOH, pH ∼ 2.9) [12]. The passivated surface is also faceted, exhibiting terraces and steps. The presence of terraces at the surface of the crystallized passive film is indicative of a slightly tilted epitaxy between the NiO lattice forming the barrier oxide layer, and the Ni(1 1 1) substrate terraces. It has been proposed that this tilt partly relaxes the interfacial stress associated with the mismatch of 16% between the two lattices. This tilt has been confirmed by GIXD measurements performed in situ on Ni(1 1 1) [40], showing good agreement between local information obtained by STM and average information obtained by GIXD. A similar surface faceting is observed for the Cu(I) oxide layer grown on copper (as described above) for which the lattice misfit between Cu2 O and Cu lattices is similar [26,28]. The lattice measured on the terraces is hexagonal with a parameter of 0.3 ± 0.02 nm assigned to NiO(1 1 1) based on STM data [8,9,11,12,14], which was also confirmed by GIXD measurements [40]. It must be pointed out that the (1 1 1) surface of NiO which has the NaCl structure is normally polar and unstable if bulk-like terminated. It is however the surface which is obtained by passivation. The reason for this is that the surface is stabilized by adsorption of a monolayer of hydroxyl groups, as confirmed by DFT modeling (see below). UHV studies of thermal oxide ultrathin films also confirm

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that the (1 × 1) unreconstructed NiO(1 1 1) surface is stabilized by hydroxylation [62,63]. These data show that the direction of growth of the oxide film is governed, at least in part, by the minimization of the oxide surface energy by the hydroxyl/hydroxide groups. The presence of water is thus a major factor for the structural aspects of the growth mechanism. The presence of steps at the surface of the passive film plays a key role in the dissolution mechanism in the passive state as revealed for passivated Ni(1 1 1) by ECSTM sequence imaging [43,64]. The passive film dissolves at the edges of the facets produced by the tilted epitaxy between the NiO(1 1 1) and Ni(1 1 1) lattices. The resulting 2D step flow process is dependent on the step orientation: 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 process leads to the stabilization of the facets with edges oriented along the close-packed directions of the oxide lattice, and produces steps that are oriented along the {1 0 0} planes, the most stable orientations of the NiO structure. A model for such a surface is presented in Fig. 9 discussed further on. 2.3. Nanostructure and relation with passivity breakdown The lateral dimensions of the crystalline grains forming the passive films on nickel are relatively well documented. 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 30–230 nm for 3D films in stationary conditions of passivity [9,12,14,16,17,43]. A large dispersion could be found on the same sample, suggesting a varying degree of advancement of the coalescence of the oxide grains during the nucleation and growth mechanism. A lower average value of ∼8 nm for the NiO(1 1 1) single-crystal domain size was obtained from GIXD data [14,40]. This difference is assigned to the fact that STM and AFM measurements are unable to resolve the multiple twin or grain boundaries that can exist in the inner part of the passive film if they do not markedly affect the surface topography. For the passive film grown on iron, the lateral grain size was determined to be 5 nm and 8 nm from in situ GIXD data obtained on passivated Fe(1 1 0) and Fe(1 0 0) surfaces, respectively [41,42]. A value of 5 nm was obtained from ex situ STM data on passivated sputter deposited pure iron films [24]. Intergranular sites, i.e. grain boundaries for a well crystallized passive film, play a major role in passivity breakdown, acting as preferential nanostructural defects where nanopits are observed to nucleate both without and with aggressive anions (Cl− ) present in the electrolyte [43–45,65,66]. This is illustrated by Fig. 5 for passivated Ni(1 1 1) surfaces [43]. After passivation at 0.55 V/SHE in a chloride-free sulfuric acid solution (pH 2.9) (Fig. 5a), the surface is completely covered by the nanogranular passive film. It is the faceted surface of the grains that is revealed by the typical higher magnification STM images shown in Fig. 4. Between the grains, depressions with a depth varying from 0.4 to 1.4 nm are measured. Their formation was assigned to the competitive dissolution occurring on the non yet passivated (or less protected) sites that are formed in the transient process of growth of the oxide film, prior to steady state passivation of the surface. Increasing the potential in the passive range by successive potential steps up to 1.05 V/SHE locally corrodes the pre-passivated surface. In the absence of chlorides, the formation of local depressions of nanometer dimensions (20–30 nm at the surface) is observed with a density of (3 ± 2) × 1010 cm−2 (some are marked in Fig. 5b). The depth of these depressions ranges from 2.2 to 3.8 nm, larger than after pre-passivation at 0.55 V/SHE and than the thickness of the passive film formed in these conditions (<2 nm), indicating the locally enhanced corrosion of the substrate in these sites. This implies that the surface, pre-passivated at 0.55 V/SHE, has been locally depassivated, with a local enhancement of the

Fig. 5. AFM images (z range = 3 nm) of the Ni(1 1 1) surface after pre-passivation for 30 min in 0.05 M H2 SO4 + 0.095 M NaOH (pH 2.9) at +0.55 VSHE (a), and subsequent increase of the potential stepwise (steps of 0.1 V) every 30 min up to +1.05 VSHE in the absence (b) or presence (c) of chloride (0.05 M NaCl). Reprinted from Ref. [44].

corrosion of the substrate, and subsequently repassivated, thus showing that nanopits are formed in the passive state in the absence of chloride. The largest dimension of the nanopits measured by AFM corresponds to a charge transient of ∼8 × 10−14 C per nanopit. Assuming a duration of 100 ms, the current transient would be ∼0.8 pA. This is of the order of magnitude of ∼1 pA that can be detected only with advanced transient measurement techniques,

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Fig. 6. Model of a metal/oxide/electrolyte system with a surface oxide barrier layer consisting of nanograins separated by grain boundaries and schemes of the potential drop for the normal defect-free barrier layer and an inter-granular boundary. Adapted from Ref. [44].

provided the electrode area is reduced to micrometer dimensions [67]. The presence of chlorides promotes the growth of the nanopits as illustrated in Fig. 5c. The depressions observed between the grains have, for the most part, the same dimensions as those described above. However, significantly larger depressions are also observed (some are marked). Their lateral dimension ranges between 40 and 50 nm at the surface and their depth between 5 and 6 nm. Their density is (2 ± 1) × 109 cm−2 . Fig. 5c also shows an area, marked T, corresponding to an extended terrace of the substrate that exhibits much less local attacks than the rest of the imaged surface where narrower terraces are observed due to a higher density of substrate steps. This shows the role of the defects of the substrate (monoatomic steps) in passivity breakdown causing preferential alignment of the metastable pits along the substrate terrace edges. A model has been recently proposed to take into account the effect of oxide grains boundaries on the mechanisms of passivity breakdown and localized corrosion [44,45]. The model postulates local redistribution of the potential drop in the defect sites that are less resistive to ion migration and thus take over less potential drop (Fig. 6). With a given total potential difference between the substrate metal and the electrolyte, a larger drop will be located at the oxide/electrolyte and/or the metal/oxide interface which accelerates the electrochemical reactions at these defect sites. According to the interface at which the potential drop is predominant, passivity breakdown can occur by enhanced dissolution (oxide thinning) at the oxide/electrolyte interface or by accelerated interfacial voiding at the metal/oxide interface followed by collapse or rupture of the passive film. A combination of the two mechanisms is possible. In all cases, the rate of the interfacial process is increased at the inter-granular defective sites of the passive film. 2.4. Electronic properties of grain boundaries in passive films As shown above, grain boundaries in nanometric passive films are preferential reactive sites prone to initiate localized corrosion

phenomena, e.g. pitting. Investigating the relationships between the nanostructure of passive film and the local properties appears then a pertinent perspective in order to develop the understanding of the mechanisms of passivity breakdown and pit initiation. As an example, we present here new data on the local electronic properties of passive films obtained by scanning tunneling spectroscopy (STS). Only a brief account of these data is given here as the results will be reported and discussed in details separately.The STS technique allows, in combination with STM, to record the tunneling current I versus the applied bias voltage Vt between surface and tip at pre-selected locations on the imaged surface. The normalized differential conductance (dI:dVt )/(I/Vt ) extracted from the I–Vt curves is proportional to the surface density of electronic states [68,69]. The electronic structure at the surface is thus probed locally with a resolution similar to imaging. The application of this method is exemplified here for Ni(1 1 1) single-crystal surfaces passivated in acidic sulfuric acid solution (pH 2.3). Spectroscopic I–Vt measurements (TS) are most often restricted to a few hundreds mV range in an ECSTM cell [70,71]. This is because with the potential of the substrate fixed by the electrochemical process under study and with the potential of the tip adjusted to minimize the electrochemical tip current and optimize the STM measurement, the tunnel voltage can no longer be adjusted in a wide range as in UHV conditions. However, recent developments allowing in situ single-point spectroscopic measurements (ECTS) over a larger range of tunnel voltage of nearly 3 V have been applied to oxide films on Fe [72–75] and Cu [76]. As an example, Fig. 7 shows, in the upper panel, tunneling spectra obtained by scanning the tip potential (UT ) between −800 and +1500 mV/SCC at 5 different sample potential (US ) values corresponding to various oxidation states of the Fe electrode and plotted versus UT [75]. SSC is used by the authors to refer to a miniaturized home built Ag/AgCl reference electrode. The lower panel shows the differential tunneling conductance dIT /dUT obtained by numerical differentiation of the curves in the upper panel. The upper panel shows linear spectra for iron polarized at −800 and −650 mV/SCC. This results from a high and nearly constant density of surface states (or high surface conductance) of the metallic Fe(0) electrode. At −400 mV, an hydrated Fe(II) layer, a few nanometer thick and with unknown electronic properties, is formed on the electrode surface in the specific conditions used in this study. A double exponential behavior is measured revealing a region of low current due to a decrease of the DOS. At 500 mV, the oxide layer is oxidized to Fe(III) with well-characterized n-type semiconductive properties. The band gap in the DOS increases in width, increasing the separation between the exponential current branches. Its width is consistent with values reported for passivated iron (Eg = 1.6–1.9 eV). The positive current branch obtained for (UT − US ) > 0 results from hole injection from the tip to the valence band (VB). At 500 mV, it is maintained by the slight depletion of the Fe(III) oxide film (upper inset marked #). At 900 mV, the positive current branch becomes extinguished due to stronger depletion conditions and evidences a low density of charge carriers (upper inset marked ##). Both at 500 and 900 mV, the negative current branches result from electron injection from the tip to the conduction band (CB) of the oxide. Such spectroscopic data show the capability of ECTS to obtain in situ the electronic properties of the solid/electrolyte interface and to measure the availability of charge carriers at the electrode surface. Series of tunneling spectra can be acquired stepwise over a large range of polarization potential and can be used to plot so-called conductograms (map of the conductance (i.e. the charge exchange properties) of the electrode as a function of potential), and to study the modifications resulting from variation of the pH of the electrolyte or from the presence of chlorides. However, up to now, such tunneling spectra have only been obtained in singlepoint mode, i.e. without combined imaging of the surface. Besides

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Fig. 7. ECTS spectra of Fe polarized at US = −800, −650, −400, +500 and +900 mV in a borate buffer aqueous solution. The upper panel shows the tip current versus tip potential curves. Solid lines and gray bands show the average curves and standard deviations, respectively. The lower panel shows the conductance curves. See text for details. Adapted with permission from Ref. [75]. Copyright 2006 American Chemical Society.

the oxide films studied with ECTS were in some cases anodically deposited and not genuine ultra-thin passive films grown by solid state reaction at the substrate surface. The results reported here are STS data (i.e. TS spectra combined with STM imaging) obtained on genuinely passivated nickel surfaces. They were obtained in ex situ conditions of measurements, i.e. in the absence of polarization of the pre-passivated surface. Such STS measurements are less appropriate than ECTS to investigate the effect of the polarization conditions on the electronic properties of the surface. However, STS is at present the unique technique enabling to investigate the relationship between electronic properties and nanostructure of a passivated surface. Fig. 8 shows the nanometer scale STM/STS data for Ni(1 1 1) passivated at 0.815 V/SHE in a sulfuric acid solution (0.05 M H2 SO4 + 0.095 M NaOH, pH ∼ 2.3). The topography (Fig. 8a) is characterized by the presence of grains, grain clusters and intergranular sites. The individual grains and the grain clusters range from 3 to 5 nm and from 12 to 24 nm in width, respectively. The apparent depth measured in the inter-granular sites ranges from 0.5 to 1.8 nm, which is consistent with a fully covering film since it does not exceed the equivalent thickness of 1.7–2.2 nm of the passive film in these conditions of passivation [12,14,60]. Fig. 8b shows the normalized differential conductance curves plotted versus Vt , the surface-to-tip bias. The I–Vt curves were assigned to grains or grains boundaries of the passive film based on line profile analysis of their locations on the topographic images recorded

Fig. 8. STM/STS data obtained on Ni(1 1 1) surfaces passivated in 0.05 M H2 SO4 + 0.095 M NaOH(aq.) (pH ∼ 2.3) at U = +0.815 V/SHE. (a) Topographic image, X = Y = 100 nm, Vt = 1.1 V, I = 0.5 nA, z range = 1.6 nm. The 25 locations pre-selected for spectroscopic I(Vt ) measurements are marked with letters from A to Y. Arrows point to grain aggregates. (b) Normalized differential conductance curves obtained from the averages of 251 and 55 I–Vt curves obtained on grain and grain boundaries of the passive film, respectively. The energy band model of the passivated Ni(1 1 1) surface deduced from the STS data obtained at the grain and grain boundaries of the passive film is shown. The arrows pointing downwards (upwards) indicate the possible electronic transfer from (to) the passivated surface. See text for details.

simultaneously. 251 and 55 measurements obtained on grain and grain boundaries, respectively, were then averaged before being numerically differentiated and normalized. The resulting (dI/dVt )/(I/Vt ) curves (Fig. 8b) reveal the distinctive local density of states obtained on the grains and grain boundaries of the passive film, schematically illustrated on the energy band model.For the occupied states measured at negative bias, the increase of the density measured on the grains for Vt ≤ −0.54 ± 0.03 V with respect to the Fermi level (at Vt = 0 V) is in good agreement with UPS data that position the top of the valence band 0.5 V below the Fermi level for NiO [77]. This marks the position EV for the grains of the passive film below which the occupied states of the passive film act as electron donors as illustrated by the arrows pointing from the energy band model. Since the width of the electronic gap of

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the bulk passive film is about 3.5 V [78–80], the position of the valence band edge is consistent with a p-type character of the film. At the grain boundaries, the slight shift to Vt ≤ −0.66 ± 0.03 V of the position where the increase of the density is measured suggests that the p character of the passive film is slightly attenuated on these sites. In the range −0.54 ± 0.03 < Vt < 0.20 ± 0.03 V on the grains and −0.64 ± 0.03 < Vt < 0.51 ± 0.03 V at the grain boundaries, the density of states is minimum which is assigned to the electronic gap of bulk passive film. However the measured density is not nil since direct tunneling can occur from/to the substrate at negative/positive bias as illustrated by arrows pointing downwards/upwards in the model. This direct mechanism is possible because the passive film does not exceed 2 nm in thickness in the selected conditions of growth. The density of unoccupied states measured at positive bias shows peaks measured on the grains for 0.5 < Vt < 1.5 V. These peaks suggest the presence of surface states (illustrated by ovals in the model) since the width of the electronic gap of the bulk passive film is about 3.5 V, and thus extends up to Vt ∼ 3 V. These surface states can promote indirect electronic transfer to the empty states of the substrate as illustrated by the arrows in the model. A remarkable feature of these STS data is to show that, at the grain boundaries, the density of these states is markedly reduced compared to the grains in the range 0.5 < Vt < 1.2 V before reaching the same level than on the grains above 1.2 V. This marked decrease of the density of states observed between 0.2 and 1.2 V suggests the absence of the surface states at the grain boundaries of the passive film, and thus no promotion of electronic transfer in this energy window in these sites. The origin of this difference and its implication for the properties of corrosion protection of passive films will be discussed when reporting in details these new data.Local tunneling spectroscopy allows us to investigate the local electronic properties of passive films, and the above data provide direct evidence that the grain boundaries of the surface oxide layer have different properties than the oxide grains. Future studies should reveal the relation between local electronic properties and surface electrochemical reactivity.

3. Atomistic modeling of passivity breakdown The atomistic modeling of corrosion phenomena in the context of the metal/liquid interfaces has been developed over the recent years mostly using ab initio density functional theory (DFT) [48,81–95]. Ab initio DFT methods can provide useful atomic and sub-atomic scale information on hydroxylated oxide surfaces of passive films and their interaction with corrosive species. However, their application is limited by the complexity of the systems that must include three phases, metal(alloy)/oxide/electrolyte, their interfaces, electric field and temperature for a realistic description. Besides, and as shown above, the description of the oxide layer must take into account its orientation, the presence of surface defects and that of nanostructural defects that are key actors of the reactivity. Periodic DFT calculations have been applied to study cation mobility in a Fe(II)–Fe(III) oxide [85], thermodynamic stability of Fe–Cr spinel oxides [92], surface hydroxylation on NiO(1 1 1) [81] and Cu2 O(1 1 1) [48], and Cl− interaction with hydroxylated NiO(1 1 1) [93–95]. As an example, we present below the interaction of Cl− with a hydroxylated NiO(1 1 1) surface for which the influence of steps has been recently studied [94]. The reactions of adsorption and sub-surface insertion of Cl atoms were first investigated on non-defective hydroxylated NiO(1 1 1) terraces characteristic of the terraces at the surface of the passive film on nickel [93]. These two reactions were selected as they are characteristic of the adsorption-induced local thinning mechanism of passivity breakdown and of the initial step of the penetration-induced voiding mechanism, respectively [96,97].

Adsorption was modeled by substituting the surface OH groups by Cl atoms at coverages of 25%, 50%, 75% and 100%. Sub-surface insertion was modeled by exchanging one adsorbed Cl of the topmost anionic layer with one O atom of the first inner anionic layer of the oxide. All structures mapping the configuration space were optimized and the adsorption and insertion energies were calculated. The adsorbed Cl surface structures tend to form a O–Ni–OH/Cl atomic trilayer characteristic of the building trilayers encountered in the layered crystalline structure of the nickel hydroxychloride (Ni(OH)Cl) bulk compound. Strong relaxation is observed at low coverage (25%), splitting the mixed OH/Cl topmost anionic layer. It decreases with increasing Cl coverage and vanishes at surface saturation. Increasing the Cl surface coverage increases the repulsive interactions in the topmost layer constrained by the lattice parameter of the oxide. The adsorption energy increases by ∼1.1 eV for a Cl coverage increasing from 25% to 100%, going from exothermic at low coverage to endothermic at high coverage. Cl adsorption is thus energetically favorable at low coverage on the undefective hydroxylated NiO(1 1 1) surface but not at high surface coverage. It does not seem to promote dissolution. The sub-surface insertion of a Cl atom into the first anionic plane of the oxide leads to surface relaxations that depend on the Cl surface coverage. At 25%, little relaxation is observed in the topmost Ni–OH bilayer but bonding is lost between the Ni atoms and O atoms of the underlying oxide, giving a highly unstable surface structure likely to dissolve if insertion can be forced and suggesting a hybrid (i.e. combining adsorption and sub-surface penetration) mechanism of local thinning of the passive film. At 50% and 75%, relaxations are observed in the surface O/Cl–Ni–OH/Cl trilayer but the bonding with the underlying oxide increases with coverage. At 100%, sub-surface insertion leads to reconstruction characterized by a strong inter-layer mixing occuring in response to the increase of the electrostatic repulsion in the topmost layer. The reconstruction allows to maintain bonding with the underlying oxide and to stabilize the saturated adsorbed structure. A (5 3 3)-oriented NiO periodic model including monoatomic (0 1 0) step edges and (1 1 1) terraces was built to model the effect of steps at the surface of the passive film (Fig. 9). Likewise Cl adsorption was modeled by substituting the surface OH groups by Cl atoms at 25%, 50%, 75% and 100% coverages and sub-surface insertion was modeled by exchanging one adsorbed Cl of the topmost anionic layer with one O atom of the first inner anionic layer of the oxide [94]. A noticeable effect of the presence of the step edges on the surface reconstruction of the adsorbed Cl structures is observed. Substructures of Ni(OH)2 , Ni(OH)Cl or Ni(Cl)2 composition are formed and detached from the step edges, as shown in Fig. 9, suggesting a major role of the step edges in the Cl adsorption-induced thinning mechanism of the oxide film. 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. Due to repulsive interactions in the topmost layer constrained by the lattice parameter of the oxide, the adsorption energy increases by 0.4 eV for a Cl coverage increasing from 25% to 100%, but less markedly than on the undefective surface, and remains exothermic. Unlike the adsorbed structures, the reconstruction of the subsurface inserted structures does not lead to the formation of substructures of Ni(OH)Cl or Ni(OH)2 type detaching from the step edges, showing that, after sub-surface insertion of the Cl atoms, film dissolution is not promoted. At surface saturation, the subsurface inserted structures become more stable than the adsorbed structures, indicating a possible bifurcation from the Cl adsorptioninduced oxide thinning mechanism to the penetration-induced mechanism of passivity breakdown. This requires the saturation in adsorbed Cl of the step edges and their immediate vicinity but not

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lattice. The step edges of the passive film surface are preferential sites of dissolution, causing a step flow dissolution mechanism in the passive state. Numerous grain boundaries are present resulting from nucleation and growth of a crystalline passive film. They act as preferential nanostructural defects in passivity breakdown and localized corrosion initiation. A major future perspective is the investigation of the relationship between the nanostructure of the ultrathin passive layers and their local properties. The example of an STS study of the local electronic properties of the passive film grown on single-crystal nickel surfaces has been presented. The results, which are very promising, show that grain boundaries display a markedly different density of states compared to the oxide grains, which implies different surface reactivity. Such studies applied to other metals and alloys will be insightful for a better understanding of the mechanisms of corrosion and corrosion protection at the nanoscale. Atomistic modeling by ab initio methods has already proved its usefulness in the field of corrosion, in providing further understanding of localized corrosion initiation mechanisms. Future work in this area will have to include explicitly the metal substrate, the aqueous medium and the electrode potential. References

Fig. 9. Side view models of the hydroxylated NiO(5 3 3) surface showing the (0 1 0) step edges, the (1 1 1) terraces and Cl-adsorbed structures before and after optimization at coverages of 25%, 50%, 75% and 100%. Adapted from Ref. [94].

necessarily of the extended defect-free terraces. Such data provide atomic scale evidence of a Cl-induced dissolution of step edges on passivated metal surfaces and of a possible mechanism for the local penetration of Cl atoms into the oxide lattice.The above data clearly show the interest of atomistic modelling, which is not only to confirm the interpretation of experimental data, but to allow us to validate or rebut hypotheses that cannot be tested experimentally. 4. Conclusion Data obtained at the nanoscale on the growth, structure and local properties of ultrathin oxide passive films providing corrosion protection of metal substrates have been reviewed. The selected examples on Cu and Ni surfaces show that in the potential range preceding 3D anodic oxide growth, OH adsorption induces surface reconstruction of the topmost metal plane to adopt the structural arrangement found in 3D passivating oxides, thus forming 2D surface oxides that play the role of structural precursors in the growth of passive films. Passive films on Cu and Ni are crystalline. The substrate structure influences the crystallographic orientation in which the oxide grows. The oxide surface is hydroxylated. Step edges, faceting the surface of the passivating oxide grains, result from a few degree tilt of the oxide lattice with respect of the metal

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