Electrochemical formation of highly pitting resistant passive films on a biomedical grade 316LVM stainless steel surface

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Materials Science and Engineering C 28 (2008) 94 – 106 www.elsevier.com/locate/msec

Electrochemical formation of highly pitting resistant passive films on a biomedical grade 316LVM stainless steel surface Abdullah Shahryari a , Sasha Omanovic a,⁎, Jerzy A. Szpunar b a b

Department of Chemical Engineering, McGill University, Montreal, Quebec, Canada H3A 2B2 Department of Mining and Materials, McGill University, Montreal, Quebec, Canada H3A 2B2

Received 13 July 2007; received in revised form 31 August 2007; accepted 12 September 2007 Available online 20 September 2007

Abstract The results discussed in the paper demonstrate that a significant improvement in pitting corrosion resistance of a biomedical grade 316LVM stainless steel can be achieved by electrochemically forming highly-protective passive oxide films on the material's surface, under cyclic potentiodynamic polarization conditions. The film formed in a sodium nitrate electrolyte is completely resistant to pitting corrosion in simulating physiological solutions even at high temperatures (60 °C), and after sterilization. The high pitting resistance of the electrochemically-formed films was explained on the basis of their semiconducting properties. Namely, the enrichment of the outer part of the electrochemically formed passive film with Cr(VI)-species results in a decrease in the density of oxygen vacancies, which act as pitting initiation sites, and their ‘replacement’ by metal vacancies formed by the electrochemical oxidation of Cr(III) to Cr(VI). In this configuration, the outer Cr(VI)-rich oxide layer behaves as cation selective, which results in the increased pitting corrosion resistance of the film. The simple electrochemical passivation technique discussed in the paper can be efficiently used to form highly pitting resistant passive films on 316LVM-built medical implant devices of any geometry. © 2007 Elsevier B.V. All rights reserved. Keywords: Biomaterials; Stainless steel; Pitting corrosion; Passive films; Electrochemical potentiodynamic passivation; Semiconducting properties

1. Introduction Stainless steels (SSs) are widely used as biomaterials and materials of construction. In biomedical applications they are used as coronary and pulmonary stents, hip prosthesis, screws, external fixations, etc. This is mainly due to their good resistance to general corrosion. Generally, the first requirement of any material serving in a biological system is that it should be inert and not cause any undesirable reaction with its surrounding. When SS is placed inside a tissue, the interaction between the implant and the tissue determines the degree of its biocompatibility. Corrosion, as an electrochemical process, commences on the surface of SS implants and subsequently affects the body's response. This is undesirably followed by release of ions such as chromium and nickel in the surrounding ⁎ Corresponding author. Department of Chemical Engineering, McGill University, 3610 University Street, Montreal, QC, Canada H3A 2B2. Tel.: +1 514 398 4273; fax: +1 514 398 6678. E-mail address: [email protected] (S. Omanovic). 0928-4931/$ - see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.msec.2007.09.002

tissue which can negatively affect the response of the host body to the biomaterial. Okazaki et al. [1] studied the degradation rate of a number of commonly used biomaterials in different biological solutions. In the case of SS316LVM, they reported a relatively high release rate of the alloying elements, mainly chromium and iron, into the solution, reaching a concentration of 2 μg cm− 2 after a seven-day exposure. The interaction of various elements, originated from dissolution of SS, with the human body has attracted much attention and has been extensively studied [2–4]. The literature has proven that any deviation from the natural concentration of the ions such as chromium, iron and nickel inside the body may severely prevent the body from functioning properly, and can ultimately lead to severe health-related consequences. Therefore, corrosion of the material is the first issue to be considered when it is designed to function in a human body. Although, the resistance of SS passive films to general corrosion is relatively high, the films are highly susceptible to localized forms of corrosion. Pitting corrosion is one of the most severe types of localized attack on SSs, which can limit their

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application in bio-systems. Pitting corrosion can adversely affect both biocompatibility and mechanical strength of the implant. For instance, it has been shown that pitting corrosion initiated on the surface of a coronary stent strut can lead to a complete mechanical failure of the stent [5]. Shih et al. [6] have investigated the influence of the surface passive oxide film properties on the pitting resistance of SS stents. They showed by in vivo experiments that the pitting resistance of the implanted stents in the dog's abdominal aorta is directly dependent on the stent material's passive film properties. Further, the thrombogenic activity of the SS surface is also believed to be dependent on the biomaterial's passive oxide film physico-chemical properties and its pitting resistance [7–9]. Investigation of the thrombogenicity of variously treated SS cardiovascular stents has shown that the minimum thrombosis is detected on the implanted stent exhibiting the highest resistance to pitting corrosion in in vivo experiments [7–9]. Further, SS fracture fixation parts such as screw heads have been frequently reported to fail due to pitting corrosion [10,11]. Also, localized forms of corrosion on hip prostheses are one of the commonly reported defects on the retrieved parts [12]. One of the other major problems related to pitting corrosion of orthopedic implants is the elevated friction originated from the pitting corrosion products acting as debris at the joints. Release of corrosion products into the tissue surrounding the implant can elicit different types of reactions in the host tissue. The increase of the local and overall concentration of certain species associated with corrosion of the implants has been cited in many clinical repots [13–15]. In the case of SSs, Williams et al. [16] showed that at the screw-plate junctions of the bone fixation components, signs of inflammations along with the corrosion products, ironcontaining granules and micro-plates containing chromium compounds, have been found. The resistance of SS to corrosion is owed to the presence of a thin oxide film on its surface, known as passive film. The physico-chemical properties of this passive film control the material's corrosion behavior, its interaction with the body, and thus the degree of the material's biocompatibility. A number of different properties of the passive film, such as chemical composition, electronic properties, thickness, etc. determine the capability of the SS to resist pitting. It is a common agreement that the passive film on a biomedical grade SS, e.g. 316LVM, has a duplex structure, being composed of Cr-oxide and Feoxide sub-layers. However, the passive/protective nature of the passive film is generally attributed to the presence of Cr-oxide. Consequently, it is believed that an increased content of chromium in the passive film results in an increased resistance of the SS to pitting corrosion [17–20]. However, Shahryari et al. [21] have shown that the enrichment of electrochemically formed passive films with Cr(III) does not yield an improvement in their pitting resistance relative to the naturally grown passive film, but rather their the enrichment with Cr(VI) species. It has also been frequently reported that molybdenum plays a positive role in pitting corrosion [22,23]. Merello et al. [24] claimed that presence of Mo in SS in the amount higher than 1 wt.% results in a considerable improvement in their pitting resistance of the material. On the other hand, the presence of

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some inclusions on the SS surface, such as sulfide inclusions, represents a major problem due to their acceleration of pitting corrosion [25–27]. Ryan et al. [26] suggested that the depletion of chromium around the MnS inclusions triggers the initial stage of pitting corrosion. This has been questioned by Meng et al. [28] and Schmuki et al. [29], who suggested that no chromiumdepleted zone was detected around any of the MnS inclusions. While it is known that low-sulphur steels (e.g. a biomedical grade 316LVM) have improved corrosion resistance, some researchers believe that pitting corrosion in these steels is also always associated with MnS inclusions [26]. Significant efforts have been made to develop methods for modification of the SS surface and/or its passive films in order to improve the material's pitting corrosion resistance, and thus its biocompatibility. These methods have been focused mainly on the removal of surface inclusions [17,30], on the modification of chemical properties and element distribution in the passive film [31–33], or on the increase in the Cr/Fe ratio in the film [17–20]. Some improvement in pitting resistance has been achieved by these and some other methods [17,19,34,35]. Nitric acid has been used as one of the most popular chemical passivation reagents for surface treatment of surgical SS implants [36]. The literature has emphasized a beneficial effect of nitric acid on chromium enrichment in the modified passive layer [17,19,30]. Coating of 316 L SS with hydroxyapatite has also been used in order to increase the material's biocompatibility in terms of general and pitting corrosion resistance [37]. Shih et al. [6,18] have tested the properties of a 316 L surface modified using various passivation methods, and have concluded that the method that produces an amorphous surface oxide layer gives the highest increase in the material's biocompatibility. Laser surface modification has been studied by many researchers as a method of improving the pitting resistance of the SS surface [33,38–41]. According to the literature, the general idea on the influence of laser surface treatment is that this method produces a homogeneous surface having a fine-grained structure, and also dissolves or confines the carbide particles. Electrochemical polishing of 316 L SS slotted tube coronary stents has been used as a surface pretreatment method in order to increase the material's biocompatibility [42]. Nitriding, a surface treatment method commonly used to increase the wear resistance of SSs, has also been used as a SS surface modification method [43,44]. Ion implantation of the 304 SS surface by molybdenum, yttrium, titanium and nitrogen has also been applied to increase the material's corrosion resistance [45–47]. The goal of the research presented in this paper has been to investigate the suitability of a simple electrochemical technique, cyclic voltammetry, for the formation of a highly pitting resistant passive film on a biomedical grade 316LVM SS surface. The influence of a range of experimental parameters (temperature, chloride concentration, number of polarization scans, etc.) on the resulting pitting behavior of the material is discussed. It will be shown that the formed passive film offers a significantly higher pitting corrosion resistance than the naturally grown passive film, thus offering an increased biocompatibility of the material. The latter has also been confirmed by preliminary

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experiments (not shown) in which a 95% increase in surface density of pre-osteoblast cells (MC3T3) has been achieved on the electrochemically modified 316LVM surface. The improvement in pitting corrosion resistance of the modified surface has been explained on the basis of chemical composition and semiconducting properties of the film. 2. Experimental All the electrochemical measurements were carried out in a single-compartment electrochemical cell. The counter electrode (CE) was a platinum wire, and a saturated calomel electrode (SCE) was used as the reference electrode (RE). Stainless steel SS316LVM was used as the working electrode (WE). Its chemical composition is shown in Table 1. The WE was prepared by embedding SS316LVM pieces into epoxy resin, exposing a two-dimensional rectangular cross-section of 0.36 cm2 to the electrolyte. Several different electrolytes were used in this research. For passivation of 316LVM, 0.1 M NaNO3 (NN) and 0.1 M phosphate buffer pH 7.4 (PB) solutions were used as passivating electrolytes. The latter was also used in capacitance measurements. Three different corrosion testing electrolytes were used: pure saline solution (0.16 M NaCl), phosphate buffer saline (PBS), and Hank's solution. All the solutions were prepared using deionized water having resistivity of 18.2 MΩ cm. All the electrochemical testing experiments were performed in oxygen-free electrolytes which was achieved by continuously purging the electrolyte with argon, starting 30 min prior to the measurement and continuing during the measurement. Prior to each experiment, the working electrode surface was wet polished using 600 grit abrasive sandpaper followed by sonication in deionized water to remove any polishing residues. The specimens were next degreased with ethanol and immersed in the cell containing an appropriate electrolyte. This was followed either by leaving the WE at open-circuit potential (OCP) for 1 h before the subsequent measurements (this surface is named here as “unmodified”), or by electrochemically passivating the 316LVM surface prior to the OCP measurement in either 0.1 M NaNO3 or 0.1 M phosphate buffer pH 7.4 (these samples are named here as “ME-NN” or “ME-PB”, respectively). The electrochemical passivation was done by polarizing the WE between two potential limits specified in the paper for each measurement and at a sweep rate of 100 mV s− 1. Prior to each passivation experiment, in order to reduce the surface oxides formed spontaneously during the 316LVM surface polishing and degreasing, the WE was cathodically polarized at − 1.0 V for 5 min. The first step in characterization of the surface corrosion resistance was the electrode stabilization at OCP in a corrosion

testing solution for a period of 1 h. Next, pitting experiments were done by anodically polarizing the working electrode from 50 mV negative of the OCP to the potential at which a current density of 1 mA cm− 2 was reached. This was immediately followed by a reverse cathodic polarization bias to the starting potential. Chronoamperometric measurements were performed following the electrode stabilization at OCP in a corrosion testing solution by potentiostatically polarizing the WE and measuring the resulting current. To investigate the semiconducting properties of naturally grown and electrochemically formed passive films, capacitance measurements were carried out in the PB solution under a potentiostatic control and cathodic (negative) bias, starting from 0.9 V down to − 0.9 V. The applied ac amplitude was ± 10 mV. X-ray photoelectron spectroscopy (XPS) measurements were made using a VG instrument Escalab 220i XL equipped with an argon ion gun. The X-ray unmonochromatic source was Al (1486.6 eV). The ion etching beam was used at 3 keV with a magnification of ten, which provided an etching area of 1.5 × 1.5 mm2. The etching rate was 2 nm min− 1 and the pressure during the etching was kept at 10− 8 mbar. The reference energies used for calibration of the binding energies were the Ag3d5/2 signal at 367.9 eV and the Cu2p3/2 signal at 932.7 eV. The analyzer was fixed at normal position (90°) to the surface. A survey spectrum was first recorded to identify all elements present on the sample surface, followed by recording high resolution spectra. The spectra were fitted using linear background subtraction and a combination of Gaussian and Lorentzian line shapes with addition of an asymmetry factor for the metal peaks. In order to derive a quantitative analysis, 2p spectra of the selected elements were recorded. Next, the calculated area under the deconvoluted peaks after background subtraction was correlated to the atomic concentration of the corresponding element using the related correction factors. 3. Results and discussion 3.1. Cyclic polarization Fig. 1 shows the cyclic voltammogram (CV) of a freshly polished 316LVM SS electrode recorded in a potential region between −0.8 V and 0.9 V, in which the solid curve represents the 1st polarization sweep and the dashed curve represents the 200th polarization sweep. The shape of the CV changes greatly with the number of polarization sweeps. The CV recorded in the first sweep (solid line) displays an anodic hump (A1) at ca.− 0.5 V and a broad current shoulder (A2) cantered at ca. 0.5 V. In the reverse cathodic polarization for the 1st sweep, the voltammogram shows two reduction peaks, C1 and C2. The relation of the shoulders/peaks to specific redox reactions has

Table 1 Chemical composition of AISI 316LVM stainless steel (wt.%) Fe

Cr

Ni

C

Mo

Mn

S

Si

P

Cu

Sn

Co

N

O

Nb

Bal

16.57

10.34

0.016

2.13

1.54

0.001

0.54

0.024

0.28

0.009

0.09

0.03

34 ppm

0.01

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Fig. 1. Cyclic voltammograms of the 316LVM surface recorded in 0.1 M NaNO3. The solid curve shows the 1st polarization sweep and the dashed curve shows the 200th polarization sweep. Scan rate = 100 mV s− 1. Inset: Variation of the anodic-to-cathodic total charge ratio with the number of sweeps. Temperature = 22 °C.

already been well explained in the literature [48–51]. Briefly, the anodic peak (A1) is related to the formation of a Fe(II)/Fe (III) layer (i.e. Fe2O3) on the pre-existing Cr(III)-oxide surface layer, followed by the oxidation of Cr(III)-oxide to soluble Cr (VI)-species (A2), most likely Cr2O72−. The cathodic peak (C2) is related to the reduction of Cr(VI)-species back to Cr(III)-oxide, and (C1) to the reductive decomposition of an oxide layer composed of Fe2O3 back to Fe(II)-species. On the freshly polished 316LVM surface (1st sweep, solid curve), anodically formed Cr(VI)-species are mainly lost in solution by diffusion through a very thin and non-compact preformed Fe2O3 passive film, with subsequent dissolution in the aqueous phase. This explains the absence of passive transition associated with shoulder A2 in the anodic scan, and poor definition (small change) of the cathodic peak C2 (solid curve) in the cathodic scan. On the other hand, the 200th sweep in Fig. 1 (dashed line) shows that the charge associated with the Cr(III)to-Cr(VI) transition in the potential region of A2 and C2 is negligible compared to that recorded in the first sweep. This demonstrates that Cr(VI) species formed in the 200th anodic sweep remain ‘arrested’ in the growing surface passive film. It also indicates that the passive film formed during the prolonged cyclization of the electrode is more compact and thicker than the film formed in the 1st sweep, thus effectively preventing the dissolution of Cr(VI) species into the solution. As it will be shown later in the text (Fig. 10), the major product of chromium oxidation during the electrode cyclization is Cr(III), formed by oxidation of metallic chromium, with a small contribution of Cr (VI) species accumulated in the outer part of the passive film. In order to get a better quantitative insight into changes during the cyclization of the 316LVM electrode, the ratio of the total anodic-to-cathodic charge is presented as an inset to Fig. 1. The curve shows that the anodic reactions related to shoulders A1 and A2 appearing in the 1st sweep are quite irreversible (QA / QC = 11.09) and that only ca. 9% of the charge related to these anodic processes is used to form species that do not dissolve into the solution but form the surface passive oxide

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film. However, by cyclization of the surface, the reversibility of the anodic processes rapidly increases, and after ca. 20 sweeps it approaches almost 100% (QA / QC = 1.05, after the 200th sweep). This indicates that the cyclic polarization of the 316LVM surface under the given conditions significantly contributes to the enhancement of the passive properties of the surface oxide film, preventing the dissolution of the underlying metals into the solution even at high anodic potentials. Hence, this could be potentially used as an efficient and convenient method for the formation of passive films on the 316LVM surface that offer better corrosion resistance than naturally grown passive films. In order to verify this hypothesis, pitting corrosion polarization measurements have been made on both the unmodified surface (on which the passive film was grown naturally) and on surfaces modified under various passivation conditions. 3.2. Pitting polarization Measurements of pitting resistance of modified and unmodified surfaces were performed in various solutions: pure saline, PBS, and Hank's. Fig. 2 shows the anodic polarization curves of the 316LVM surface modified in two different electrolytes (sodium nitrate, ME-NN, and phosphate buffer, ME-PB) and also the response of the unmodified surface. The unmodified surface represents the 316LVM surface on which the passive film was grown naturally during the stabilization of the electrode at OCP. All the curves were recorded in PBS. The pitting resistance of the unmodified surface is in agreement with the literature [26,52–54] with the onset of pitting at ca. 0.4 V. On the other hand, the polarization curve recorded on the ME-PB surface shows a considerable improvement in pitting resistance. The passive region extends to high anodic potentials, and the breakdown of the passive film commences at ca. 1.05 V. The examination of the ME-PB electrode after the

Fig. 2. Pitting polarization curves of the unmodified, ME-NN and ME-PB 316LVM surfaces recorded at 22 °C in a 0.16 M PBS solution pH 7.4. Scan rate = 1 mV s− 1. Modification of the surfaces was done by potentiodynamic cyclic polarization of the electrode at 22 °C in 0.1 M sodium nitrate (ME-NN) or phosphate buffer (ME-PB) between −0.8 Vand 0.9 V at a scan rate of 100 mV s− 1 by applying 300 sweeps.

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polarization measurement revealed that the increase in current observed at ca. 1.05 V is actually not due to pitting but due to the crevice corrosion at the electrode/resin interface. The surprising result was that the polarization curve for the ME-NN surface shows the absence of a pitting-loop, which demonstrates that the surface passive film formed in NN is completely resistant to pitting under the experimental conditions applied. The rise in current at ca. 1.0 V is not due to pitting, but due to oxygen evolution. In addition to the high pitting resistance, the other positive effects observed on the ME-NN surface are a decrease in passive current by ca. 1.5 orders of magnitude, and a shift in OCP to more noble values by ca. 160 mV. The experiments performed in Hank's solution also demonstrated the absence of pitting on the ME-NN surface, while the surface on which the passive film was grown naturally (unmodified surface) pitted already at 0.16 V (not shown). The same experiment was also performed with high-sulfur SS 316 L, and the results demonstrated that the surface on which the passive film was grown naturally pitted at 0.4 V, while the surface on which the passive film was formed electrochemically in NN offered a considerably higher pitting corrosion resistance, and pitted only at potentials above 0.96 V (not shown). The results in Fig. 2 and those in Hank's solution demonstrate that the passive film formed electrochemically in sodium nitrate (NN) offers higher corrosion resistance than the film formed in phosphate buffer (PB). Therefore, further in the text the ME-NN surface will be termed as the modified surface, unless otherwise stated. Literature on the modification of SS surfaces by various methods report a wide range of controversial results. For example, when a 316 L SS surface was thermally modified, no improvement in pitting corrosion was noted in Ringer's solution at 37 °C [18]. On the other hand, electropolishing of the same surface resulted in a shift in pitting potential by about 0.25 V. When the surface was coated with an amorphous oxide, no pitting was observed. However, it should be noted that the authors in [18] performed these pitting polarization measurement at a very high scan rate, 167 mV s− 1, which might have been too high to induce pitting in the time scale of the measurement in the pitting-susceptible region. Hence, our opinion is that these results cannot be taken as completely conclusive. Further, the authors in [17] showed that passivation of a 316 surface in a nitric acid solution yielded a shift in the pitting potential by about 0.15 V when tested in a neutral aqueous solution containing 1 M of chlorides at 70 °C, while the anodic modification of a 304 SS surface in a metasilicatecontaining solution [55] resulted in a positive shift in pitting potential by about 0.3 V at room temperature. Treatment of the same surface under the same anodic conditions in a basic hydroxide solution showed no improvement in the material's pitting resistance [55]. A significant increase in the pitting potential (by ca. 0.7 V) was achieved when a 304 L SS surface was nitrided, but the drawback was that the resulting passive current was higher by one order of magnitude with respect to the unnitrided surface [44]. Modification of a 304 SS surface by UV radiation improved the pitting resistance of the surface in 0.6 M chloride solution by about 0.09 V, with a decrease in the pit

generation rate by half an order of magnitude [35]. A slightly higher positive displacement in the pitting potential (ca. 0.17 V) was obtained in 0.1 M chloride solution when a 316 SS surface was anodically polarized at a constant potential and UVilluminated for 5 h [56]. Using an excimer laser surface treatment method in an air stream, a shift in pitting potential by about 0.15 V was achieved in 0.6 M chloride solution at room temperature for a that of 316LS SS (the chemical compassion of 316LS is close to that of 316LVM) surface [39]. In addition, a decrease in passive current by about one order of magnitude was achieved. However, the same treatment in a nitrogen atmosphere resulted in a significant increase in both the pitting corrosion susceptibility and passive current. On the other hand, the results published in [41] showed that the laser surface melting of 304 SS in a nitrogen atmosphere resulted in an increase in pitting potential by about 0.3 V, and by 0.32 V when the surface was treated in an argon atmosphere [40]. This short literature review on the surface modification of SSs for the improvement of their resistance to pitting corrosion shows that the simple electrochemical modification procedure presented in this paper offers a very good and convenient alternative method that could be used to increase the resistance of the 316LVM SS surface to pitting corrosion. It should be noted that the adsorption of phosphate ions (present in PBS and Hank's solution) on SSs is known to contribute to corrosion inhibition [57,58]. Hence, the pitting corrosion resistance of the ME-NN and unmodified surfaces were also tested in pure saline solutions in the absence of phosphate ions (discussed in the forthcoming sections). It should also be noted that the concentration of chlorides in the pure saline solution is the same as that in PBS and Hank's solution (0.16 M), but due to the absence of inhibiting anions, pure saline is a more pitting-corrosion aggressive solution. 3.3. Influence of number of sweeps The electrochemical passivation process was discussed in details earlier in the text. It was also explained that the reversibility of the anodic/cathodic surface processes increases with an increase in the number of passivation sweeps (Fig. 1). This was then hypothesized to be related to the observed increase in the surface pitting resistance. In order to verify this hypothesis, the pitting resistance of the surface modified by applying different number of sweeps was investigated. Fig. 3 shows the variation of the pitting potential of 316LVM with the number of passivation sweeps applied during modification of the surface in the NN solution. There is an obvious correlation between the pitting potential recorded in saline solution and the applied number of sweeps. With an increase in the number of sweeps, the pitting potential also increases from ca. 0.23 V (naturally grown passive film) to ca. 1.07 V for the passive film formed by applying 300 sweeps. This is in agreement with the data presented in the inset to Fig. 1, and indicates that with an increase in the number of passivation sweeps, the compactness and thickness of the passive film also increases, thus contributing to an increase in pitting resistance of the material. No further improvement was observed by increasing the number of polarization sweeps above 300. Thus, all the data on the

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Fig. 3. Dependence of the pitting potential on the number of sweeps applied during the modification of the 316LVM surface. The modification of the surface was done by the potentiodynamic cyclic polarization of the electrode at 22 °C in 0.1 M NaNO3 between − 0.8 V and 0.9 V at a scan rate of 100 mV s− 1 by applying the specified number of sweeps. The corresponding pitting corrosion measurements were done in 0.16 M NaCl at a scan rate of 1 mV s− 1 and at 22 °C.

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(curve 2), and the recorded passive current sharply decreased reaching a quazi-stable value after ca 10 min. Even when the surface was polarized for a longer time (1 h), its stability remained good and no metastable pits were formed. On the other hand, curve 1 (Fig. 4b) shows that the current on the unmodified surface first started to decrease due to the initial passivation of the surface, and then, after approximately 1 min, suddenly started to increase as a result of the localized dissolution (pitting) of the passive film. No current spikes could be observed on this curve, indicating that stable pits were formed. A steady-state (localized) dissolution of the film, characterized by high current density (ca. 10 mA cm− 2) was reached already after ca. 15 min. The results in Fig. 4 are in agreement with those presented in Figs. 2 and 3, and clearly demonstrate that the electrochemically formed passive film (ME-NN surface) offers a significantly higher pitting corrosion resistance than the naturally grown passive film (unmodified surface). 3.5. Influence of chloride concentration

modified surface presented in the remaining sections of the paper refer to this condition, if not otherwise stated. 3.4. Pitting choronoamperometric measurements It is a general agreement that pitting corrosion is triggered from defects and/or susceptible spots on the surface [25–28]. Therefore, as the number of the active spots on the surface increases, the probability of the stable pits formation also increases. One of the means of comparing the susceptibility of materials to pitting corrosion (in addition to measuring the pitting potential) is to measure the pit initiation frequency, using chronoamperometry. Generally, potentiostatic polarization of SS in a metastable pitting potential region (between the OCP and pitting potential) results in the formation of metastable pits on the material's surface, followed by their repassivation, which is in chronoamperometry manifested as a series of sharp current transients (spikes) [59]. Fig. 4a shows the chronoamperometry curve of the unmodified (curve 1) and of the ME-NN (curve 2) 316LVM surface recorded at 0.35 V, which is slightly below the pitting potential (by ca. 50 mV) of the unmodified surface (Fig. 2). There is an obvious difference between the behavior of the unmodified and ME-NN surface. The graph demonstrates that no current spikes are observed on the curve recorded on the ME-NN surface (curve 2). However, large current spikes, representing local dissolution of the passive film followed by its repassivation, are visible on the unmodified surface (curve 1). Although the observed current spikes (Fig. 4a, curve 1) seem to be quite sharp, the actual time length of each spike (from the initiation to death) is ca. 3 s. This time covers the pit initiation stage, followed by a short pit propagation period. However, at 0.35 V (Fig. 4a, curve 1) the propagation (growth) of the pit to form a critical-pit is not thermodynamically and kinetically favourable, and the pit eventually repassivates. For comparison, the results recorded at 0.45 V, which is slightly above the pitting potential (by ca. 50 mV) of the unmodified surface, are presented in Fig. 4b. Again, the formation of metastable pits on the modified surface is completely absent

It was mentioned earlier that the ME-NN modified surface did not pit in PBS and in Hank's solution. Also, increasing the NaCl concentration in those solutions to even 1 M did not result in pitting of the ME-NN surface. Therefore, in order to evaluate the pitting resistance of the 316LVM surface in more severe conditions, experiments were performed by testing the unmodified and ME-NN surfaces in pure saline solutions of high chloride concentrations (up to 1 M). Fig. 5 shows the dependence of pitting potential, Epit, of the unmodified and ME-NN surfaces on the chloride concentration in the testing electrolyte. The variation of the pitting potential with chloride concentration shows a similar trend for both surfaces, but there is a large difference in the corresponding absolute pitting potential values. For the unmodified surface, the pitting potential at a chloride concentration of 0.16 M is 0.23 V. An increase in the chloride concentration to 0.5 M and then to 1 M results in a decrease in Epit to 0.20 V and 0.08 V, respectively.

Fig. 4. Choronoamperometric curves of the (1) unmodified and (2) modified 316LVM surface recorded at 22 °C in 0.16 M PBS at 0.35 VSCE (plot a) and 0.45 VSCE (plot b) Modification of the surface was done by potentiodynamic cyclic polarization of the electrode at 22 °C in 0.1 M NaNO3 between − 0.8 V and 0.9 V at a scan rate of 100 mV s− 1 by applying 300 sweeps.

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Fig. 5. Dependence of pitting potential of the ME-NN and unmodified surfaces on the chloride concentration in the NaCl testing solution. Modifications of the surface was done by potentiodynamic cyclic polarization of the electrode at 22 °C in 0.1 M NaNO3 between −0.8 V and 0.9 V at a scan rate of 100 mV s− 1 by applying 300 sweeps. The corresponding pitting measurements were done in 0.16 M NaCl solution at 22 °C and at a scan rate of 1 mV s− 1.

On the other hand, the modified surface shows a significantly higher pitting corrosion resistance, characterized by a very high pitting potential, Epit = 1.07 V, which remains constant with an increase in chloride concentration even to 0.5 M. By further increasing the chloride concentration to very high values, i.e. 1 M, the pitting potential decreased to 0.5 V. Nevertheless, it should be noted that even at such a high concentration of chlorides in the solution, the modified surface offers a significantly higher pitting resistance than the unmodified surface, 0.5 V versus 0.08 V, respectively. This shows that the electrochemical modification of 316LVM results in the formation of a passive film that offers a significantly higher pitting corrosion resistance than the naturally grown passive film even at chloride concentrations that are more than six times higher than those in the human body. This demonstrates that the applied electrochemical passivation procedure could also be used to increase the pitting corrosion resistance of 316LVM for use in some other applications such as marine, industrial, etc.

respectively. This small decrease in the pitting potential with an increase in the film formation temperature could be due to the difference between the passive film formation temperature and testing temperature (38 °C). Namely, with an increase in the film formation temperature the specific volume of the formed passive film also increases (the films expands). Then, after the sample is cooled down to the testing solution temperature (23 °C), the passive film shrinks. This can result in the formation of nano-sized cracks (pores) in the film, which represent pitting-susceptible surface sites. With an increase in the film formation temperature, the surface density of these cracks also increases and consequently, the pitting potential decreases. However, Fig. 6 demonstrates that this phenomenon is almost marginal in the temperature range investigated, i.e. it does not have a significant influence on the resulting pitting resistance of the surface passive film. Nevertheless, Fig. 6 shows that even the passive film formed at 60 °C offers a significantly higher pitting potential (0.95 V) than the naturally grown passive film (0.22 V). Pallotta et al. [49] and Cristofaro et al. [64] have shown that the formation of a passive layer on SS at higher temperatures promotes the chance of replacement of Cr(III) by Fe(II). They have postulated that this, in turn, results in a decrease in the Cr/Fe ratio in the film and, consequently, in a decrease in the pitting resistance and a considerable increase in the passive current of the material (0.55 VSCE and 0.3 VSCE at 15 °C and 60 °C, respectively). However, as we will demonstrate later in the text, the increased pitting resistance of the electrochemically modified 316LVM surface does not originate from the increased Cr/Fe ratio in the passive film, which thus explains the difference between our results (Fig. 6) and those in [49,64]. The previous section describes the influence of the modification solution temperature on the corresponding pitting resistance of the material, which we have used to optimize the passive film formation conditions. However, it would also be interesting to investigate the influence of the corrosive solution temperature on the material's pitting resistance, especially in the region of physiological importance (36–42 °C). This is of a

3.6. Influence of temperature The importance of temperature on the SS pitting corrosion resistance has been frequently reported in the literature [59–61]. It has been shown that an increase in temperature results in a significantly higher frequency of metastable pit formation [49,62,63]. An increased pitting corrosion susceptibility of SS at higher temperatures can considerably lower the material's suitability in biomedical applications, especially when used as an implant material. Therefore, it is important to investigate the temperature-dependent pitting behavior of 316LMV studied in this work. Fig. 6 demonstrates the influence of the modification solution temperature on the resulting pitting corrosion behavior of the electrochemically formed passive film (ME-NN). The pitting measurements were performed at 22 °C in 0.16 M NaCl. The graph shows that an increase in the modification solution temperature from 22 °C to 60 °C resulted only in a slight decrease in the pitting potential, 1.07 V, 1.05 V and 0.95 V,

Fig. 6. Dependence of pitting potential on the modification solution temperature. The modification of the surface was done by potentiodynamic cyclic polarization of the electrode in 0.1 M NaNO3 at the specified temperature, between − 0.8 V and 0.9 V at a scan rate of 100 mV s− 1 by applying 300 sweeps. The corresponding pitting measurements were done in 0.16 M NaCl solution at 22 °C and at a scan rate of 1 mV s− 1.

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practical importance for the use of 316LVM as an implant material due to possible temperature changes occurring in the implanted material's surrounding as the consequence of inflammation. Fig. 7 shows the dependence of pitting potential on the testing solution temperature. The corresponding passive film was formed at 22 °C. The graph shows that an increase in the testing solution temperature resulted in a decrease in the pitting potential for the electrochemically formed passive film (ME-NN). A similar trend was also recorded for the naturally grown passive film (unmodified), although this is not clearly visible on the graph due to the difference in the pitting potential values recorded on the two surfaces. The enhanced susceptibility of SSs toward pitting at higher temperatures has been reported in the literature [59,61,65]. The trend observed in Fig. 7 is in accordance with an Arrheniuslike type of kinetics, i.e. with an increase in temperature the rate of the corrosion process also increases, resulting in a decreased resistance of the material towards pitting corrosion. Nevertheless, even at 60 °C the electrochemically formed passive film (MENN) still offers higher pitting resistance than the unmodified surface, Fig. 7. It is important to mention that when the same experiments were performed in Hank's solution, no pitting was recorded in the whole temperature region studied. Hence, in the physiologically important temperature range (36–42 °C) the modified surface is completely resistant to pitting corrosion when tested in a physiological electrolyte (Hank's). In order to determine the influence of sterilization of the MENN modified 316LVM surface on its pitting resistance, the corresponding samples were sterilized in an autoclave and 70% ethanol. In both cases the film maintained its high corrosion resistance, and pitted at very high anodic potentials, 1.11 V and 1.37 V, respectively (not shown). The results presented so far demonstrate that the cyclic potentiodynamic polarization of the 316LVM surface results in the formation of a surface passive film that provides a considerably higher resistance to pitting than the naturally grown film. The film maintains its high pitting resistance at high

Fig. 7. Dependence of pitting potential on the testing solution temperature. The corresponding pitting measurements were done in 0.16 M NaCl solution at a scan rate of 1 mV s− 1 and at the specified temperature. The modification of the surface was done by potentiodynamic cyclic polarization of the electrode at 22 °C in 0.1 M NaNO3 between − 0.8 V and 0.9 V at a scan rate of 100 mV s− 1 by applying 300 sweeps. The passive film on the unmodified surface was grown naturally at OCP by keeping the surface in 0.1 M NaNO3 at 22 °C for the same length of time as the ME-NN surface.

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testing solution temperatures and after sterilization. It would be now interesting to determine the major factors that are responsible for the observed increase in pitting resistance of the material. Possible reasons for the very high pitting corrosion resistance of the electrochemically formed passive film could be related to either a change in the semiconducting properties of the passive film, and/or an increase in the film thickness, and/or, as commonly believed [17–20], enrichment of the passive film with chromium. In order to investigate which effects are responsible for the observed high pitting corrosion resistance, XPS and Mott–Schottky measurements were made. 3.7. XPS results The chemical composition of the passive films formed electrochemically in sodium nitrate (ME-NN) and phosphate buffer (ME-PB), and that of the naturally grown passive film (unmodified) was analyzed by XPS. Fig. 8 shows an example of deconvoluted XPS spectra of the outer part of the passive film for the (a) unmodified and (b) MENN sample (chromium and iron responses). The agreement between the modeled and experimental data is very good. The comparison of the corresponding spectra reveals that Cr(VI) species were detected only in the electrochemically formed passive film, while Fe(0) species were detected only in the naturally grown passive film. The predominant state of chromium in both passive films is Cr(III) in Cr(OH)3, with a small contribution of Cr(III) in Cr2O3. Both passive films contain Fe(II) and Fe(III) mostly in Fe3O4, with some contribution of Fe(II) in FeO. The peaks at 715.3, 713.4 and 717.2 eV are well-known shake-up satellite peaks [66–69]. The shake-up satellite is caused by an incident X-ray photon transferring a discrete portion of its energy to the excitation of a second electron rather than imparting its entire quantum of energy to the primary, photo-ejected electron. This photoelectron looses a small amount of energy and appears on the XPS spectrum at a slightly higher binding energy [66]. Iron-oxide satellite structures are frequently used as fingerprints to identify iron-oxide phases. The shake-up satellite in Fig. 8 at 715.3 eV can be assigned to Fe(II) in FeO [66,67], while the other two satellite peaks (713.4 and 717.2 eV) are attributed to Fe(II) in chromite (FeCr2O7) [70]. No shake-up satellite structures are visible on the chromium spectra. The modeled XPS spectra were further analyzed to evaluate the influence of the depth distribution of various Cr and Fe species on the corresponding pitting behavior of the passive film. Fig. 9(a) shows the atomic Cr/Fe ratio in the passive film for the unmodified surface and for the surfaces modified in two different electrolytes (ME-NN and ME-PB). For each case, an increase in the passive film's Cr/Fe ratio with respect to the bulk material is recorded. The highest relative increase in Cr was obtained in the ME-PB passive film, while the Cr/Fe ratio in the ME-NN passive film was even lower than that in the naturally grown film (unmodified). Nevertheless, the ME-NN surface offered the highest pitting resistance, with a complete absence of pitting in PBS and Hank's solution. The ME-PB surface also offered high resistance to pitting, but unlike the ME-NN

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Fig. 8. Examples of deconvoluted XPS spectra recorded on (a) the 316LVM surface on which the passive film was grown naturally (unmodified), and on (b) the 316LVM surface on which the passive film was formed electrochemically at 22 °C by potentiodynamic cyclic polarization of the electrode in 0.1 M NaNO3 between − 0.8 V and 0.9 V at a scan rate of 100 mV s− 1 by applying 300 sweeps. The response was recorded after removing the outer passive oxide layer by sputtering for 1 min. All the components in the graphs are labeled based on binding energies of the peaks.

surface, it eventually pitted in PBS when the polarization curve reached high anodic potentials, ca. 1.1 V (Fig. 2). This clearly demonstrates that the Cr/Fe ratio in the passive film is not the major factor responsible for the high pitting resistance of the electrochemically formed passive films, as usually thought [17–20]. Fig. 9(b) shows the oxygen profile in the electrochemically formed (ME-NN and ME-PB) and naturally grown (unmodified) passive film. The oxygen content in all three passive films increases going from the metal/oxide interface toward the outer passive film surface. This trend is in accordance with the distribution of oxidized Cr and Fe species in the film (i.e. Cr(III + VI) and Fe(II + III)). Namely, the total amount of oxidized Fe and Cr species in the film increases going from the metal/oxide to the outer oxide passive film surface, while the amount of

metallic Fe and Cr decreases and reaches zero at the film thickness of ca. 12 nm and 6 nm for the two electrochemically formed films and the naturally grown film, respectively. This could be seen in Fig. 10 for chromium, and a similar trend was also obtained for iron (not shown). Further, a significant difference between the maximum oxygen content in the electrochemically formed passive films (ME-NN and ME-PB) and the naturally grown passive film (unmodified) is evident, Fig. 9(b). One of the reasons for this is the formation of Cr(VI) species in the form of Cr2O72− in the two electrochemically formed passive films (Fig. 10). This species yields a higher oxygen/Cr ratio than Cr(OH)3 and Cr2O3 species that constitute the naturally grown film (Figs. 8 and 10). Further, the normalized oxygen content (i.e. the total amount of oxygen in the passive film divided by the total film thickness) in the ME-PB and ME-

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Fig. 9. (a) Cr/Fe atomic ratio, and (b) oxygen depth profile in the 316LVM passive film grown naturally (unmodified), and formed electrochemically at 22 °C by potentiodynamic cyclic polarization of the electrode in 0.1 M NaNO3 (ME-NN) and 0.1 M phosphate buffer pH 7.4 (ME-PB) between −0.8 V and 0.9 V at a scan rate of 100 mV s− 1 by applying 300 sweeps.

NN passive films is 2.9 and 1.8 times that in the naturally formed passive film, respectively. This indicates that the two electrochemically formed films should be characterized by a higher density of metal vacancies than the naturally grown film, which could be the main reason for the observed difference in the pitting behavior. Indeed, the Mott–Schottky measurements presented in the following section of the paper will prove this assumption. The Fe oxidation state depth profile analysis reveled that there are not any distinct differences in the distribution of Fe(II), Fe(III) and Fe(0) species in the three films. On the other hand, the results in Fig. 10 demonstrate that the two electrochemically formed passive films (ME-NN and ME-PB) show the accumulation of Cr(VI) species in the outer part of the film. However, Cr(VI) was not detected in the naturally grown passive film (Fig. 10, unmodified). This is quite understandable considering that the two electrochemically grown films are formed by the potentiodynamic polarization of the 316LVM surface to high anodic potentials, where the oxidation of Cr(III) to Cr(VI) occurs (shoulder A2 in Fig. 1). Unlike Cr(VI), Fig. 10 demonstrates that Cr(III) species are present in all three passive films, and that their corresponding depth distribution is very similar. Taking this into account, and also taking into account that there are not any distinct differences in the distribution of Fe species in the three films, and that the normalized oxygen content in the electrochemically formed passive films is higher than that in the naturally grown film (as discussed in the previous paragraph), it seems that the formation of Cr(VI) species in the passive film is the origin of the observed difference in pitting corrosion behavior presented in Figs. 2 and 4. Further, the fact that the thinner passive film formed in sodium nitrate (NN) offered higher pitting resistance than the thicker film formed in phosphate buffer (PB), Fig. 2, indicates that the film thickness is not responsible for the difference in pitting corrosion resistance between the two electrochemically formed films. However, we do not want to eliminate the possibility that

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the film thickness could be partially responsible for the difference in pitting resistance between the naturally and electrochemically formed films. Since possible incorporation of N- and P-species into the passive layer can also contribute to the increased corrosion resistance of the investigated films, their presence in the formed passive films was investigated by XPS. However, N was not detected in the passive film formed in the sodium nitrate solution, while, on the other hand, P was detected in the film formed in the phosphate buffer solution. Hence, the presence of P in the ME-PB film could contribute to the observed increase in the film resistance (compared to the naturally grown film). Yet, the (N- and P-free) passive film formed in the sodium nitrate solution offers the highest pitting corrosion resistance. This indicates that the presence of P in the ME-PB film makes only a minor contribution to the increased corrosion resistance of the film. 3.8. Semiconducting properties It is known that passive films on SSs exhibit semiconducting behavior [65,71–73], which significantly influences the resulting corrosion properties of these materials. Relations between semiconducting properties and susceptibility of passive oxide films to localized corrosion has been well documented [74–76]. Usually, electronic/semiconducting properties of these films are investigated using a capacitance technique [71,72,77]. Then, in the interpretation of capacitance measurements, the Mott– Schottky approach is used, based on the assumption that the capacitance response is controlled by the band bending. Thus,

Fig. 10. Depth profile of different oxidation states of chromium in the passive films formed under different conditions. The data were obtained by modeling XPS spectra recorded at different depths of the passive film.

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the capacitive behavior of the passive film/electrolyte interface is assumed to be similar to that of a semiconductor/electrolyte interface, and the total measured capacitance includes the space charge capacitance, CSC, and the Helmholtz double layer capacitance, CH in series: 1 /CT = 1/CSC +1/ CH. The Helmholtz double layer capacitance is considered to be constant [73]. The variation of space charge capacitance in the passive film with applied potential could be described by the well-known Mott–Schottky equation:   1 2 kT ¼ E  Efb  2 eeo eNdop e CSC

ð1Þ

where ɛ is the dielectric constant of the oxide film, ɛo is the vacuum permittivity (F cm− 1), e is the elementary charge of an electron (C), Ndop is the concentration of dopants (cm− 3), Efb is the flat-band potential (V), k is the Boltzmann constant (JK− 1), and T is the temperature (K). If an oxide film behaves as a semiconductor that can be described by the Mott–Schottky 2 model, the dependence between 1 / CSC and E should give a straight line (positive for an n-type and negative for a p-type semiconductor). An n-type behavior of SS passive films has been correlated with the response of iron-oxides in the passive film, which are characterized by a non-stoichiometric composition resulting in oxygen vacancies that contribute to the n-type semiconductivity [74,78,79]. On the other hand, a p-type behavior has been attributed to the response of chromiumoxides, and is characterized by the non-stoichiometry resulting in metal vacancies that contribute to the p-type semiconductivity [74,78,79]. The Mott–Schottky response of a naturally grown passive film (unmodified surface) and of passive films formed by applying a specific number of polarization (modification) cycles (Fig. 11) shows the existence of several potential regions characterizing the semiconducting behavior of the corresponding passive films. The existence of multiple Mott– Schottky regions has been explained by the existence of donor/ acceptor states of different energy levels (deep and shallow) [71,78]. Fig. 11 shows that each passive film formed on the 316LVM surface gives the same response in Regions I, II, III and V. These regions have been discussed elsewhere [80] and, thus, will not be discussed here. If we compare the Mott– Schottky (Fig. 11) and pitting (Fig. 2) curves of the unmodified sample, we will see that the pitting potential coincides well with the n-type behavior in Region IV. Hence, it appears that on the unmodified surface (naturally formed passive film), oxygen vacancies act as pitting initiation sites. This is quite in agreement with the point defect theory [81,82], which states that the initial pitting reaction that occurs at the film/solution interface involves the adsorption of chloride ions into these oxygen vacancies, followed by a Schottky-pair type reaction. This adsorption of chlorides leads to the generation of cation vacancies at the film/solution interface, followed by their flux through the passive film to the film/metal interface. If the flux of these cation vacancies is so high that it cannot be compensated by the generation of cations at the metal/film

Fig. 11. Mott–Schottky plots of 316LVM passive films. The graphs show the response of the unmodified 316LVM surface and the surfaces modified at 22 °C in 0.1 M NaNO3 between − 0.9 V and 0.9 V at a scan rate of 100 mV s− 1 by applying the specified number of potentiodynamic polarization sweeps. The measurements were done at 22 °C in 0.1 M phosphate buffer solution pH 7.4 under the negative potential bias. Frequency = 5 kHz, ac amplitude (rms): ±10 mV. The data was corrected for the contribution of double layer capacitance.

interface, cation vacancy condensates are formed. This, in turn, results in thinning of the passive film, or local detachment from the metal. Once the condensates have grown to a critical size, the film ruptures and rapid local pitting attack occurs. Taking into account the point defect theory, our assumption is that an increase in pitting corrosion resistance of the material could be achieved by increasing the concentration of metal vacancies in the passive film. On the Mott–Schottky plot (Fig. 11), this change would be manifested as a transition from an n- to a p-type semiconductivity in the pitting susceptible Region IV. Indeed, Fig. 11 shows that by modifying the 316LVM surface by cyclic potentiodynamic polarization, the n-type semiconductivity in Region IV gradually transforms into the p-type semiconductivity, finally resulting in merger of Regions IV and V (Fig. 11, diamond symbols). Consequently, the resulting pitting potential also gradually increases, as demonstrated in Fig. 3. The gradual transition from the n-type to the p-type semiconductivity is related to the gradual enrichment of the passive film with metal vacancies. This process occurs by the cyclic potentiodynamic formation of the passive film under the conditions presented in Fig. 1. Namely, by polarizing the 316LVM surface at high anodic potentials, Cr(VI) species are formed. With cyclization, progressively more of these species remain ‘arrested’ in the growing passive film (inset to Fig. 1 and Fig. 10). On the other hand, only Cr(III), but not Cr(VI), species are formed during the natural growth of the film (unmodified sample), which is clearly demonstrated in Fig. 10. The formation of Cr(VI) species in the form of Cr2O72−increases the oxygen content in the passive film (Fig. 9b), producing an increase in the density 3− of metal vacancies, VM , according to [83]: CrðIIIÞYCrðVIÞ þ VM3 þ 3e

ð2Þ

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This, in turn, leads to the change in the Mott–Schottky behavior in Region IV (Fig. 11), and finally to the increased pitting resistance of the passive film (Fig. 2). 4. Conclusion It was demonstrated that the modification of the biomedical grade 316LVM stainless steel surface by cyclic potentiodynamic polarization in sodium nitrate or phosphate buffer results in the formation of passive oxide films that offer a significantly higher pitting corrosion resistance than the naturally grown passive film. The film formed in the nitrate electrolyte showed to be completely resistant to pitting in physiological simulating electrolytes (PBS and Hank's) even at high temperatures and after sample sterilization, and to offer very high pitting corrosion resistance in the pure saline electrolyte, even at elevated chloride concentrations. It is also worth to mention that the film formed in the nitrate electrolyte maintained the same (high) pitting resistance even after two months of constant exposure to the pure saline solution. The capacitance analysis demonstrated that the major difference between the electrochemically formed and naturally grown passive film is in the type of semiconductivity in the potential region where pitting on the unmodified surface occurs. The XPS measurements showed that this is due to the presence of electrochemically formed Cr(VI)-species in the outer part of the electrochemically formed passive film. Namely, the electrochemical formation of the passive film results in a decrease in the density of oxygen vacancies, which act as pitting initiation sites, and their ‘replacement’ by metal vacancies formed by the oxidation of Cr(III) to Cr(VI). In this configuration the outer Cr(VI)-rich oxide layer behaves as cation selective which, in turn, results in the increased pitting corrosion resistance of the corresponding passive film. Acknowledgements Grateful acknowledgment is made to the Natural Science and Engineering Research Council of Canada and the McGill Centre for Biorecognition and Biosensors for support of this research. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11]

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