In situ nano- to microscopic imaging and growthmechanism of electrochemical dissolution(e.g., corrosion) of a confined metal surfaceC. Merolaa, H.-W. Chenga, K. Schwenzfeiera, K. Kristiansenb, Y.-J. Chena, H. A. Dobbsb, J. N. Israelachvilib,1,and M. Valtinera

aDepartment for Interface Chemistry and Surface Engineering, Max-Planck-Institut für Eisenforschung GmbH, D-40237 Düsseldorf, Germany;and bDepartment of Chemical Engineering, University of California, Santa Barbara, CA 93106

Contributed by J. N. Israelachvili, July 28, 2017 (sent for review May 19, 2017; reviewed by Joelle Frechette and Philippe Marcus)

Reactivity in confinement is central to a wide range of applicationsand systems, yet it is notoriously difficult to probe reactions inconfined spaces in real time. Using a modified electrochemicalsurface forces apparatus (EC-SFA) on confined metallic surfaces, weobserve in situ nano- to microscale dissolution and pit formation(qualitatively similar to previous observation on nonmetallic sur-faces, e.g., silica) in well-defined geometries in environmentsrelevant to corrosion processes. We follow “crevice corrosion” pro-cesses in real time in different pH-neutral NaCl solutions and appliedsurface potentials of nickel (vs. AgjAgCl electrode in solution) for themica–nickel confined interface of total area ∼0.03 mm2. The initialcorrosion proceeds as self-catalyzed pitting, visualized by the suddenappearance of circular pits with uniform diameters of 6–7 μm anddepth ∼2–3 nm. At concentrations above 10 mM NaCl, pitting isinitiated at the outer rim of the confined zone, while below10 mM NaCl, pitting is initiated inside the confined zone. We com-pare statistical analysis of growth kinetics and shape evolution ofindividual nanoscale deep pits with estimates from macroscopic ex-periments to study initial pit growth and propagation. Our data andexperimental techniques reveal a mechanism that suggests initialcorrosion results in formation of an aggressive interfacial electrolytethat rapidly accelerates pitting, similar to crack initiation and propa-gation within the confined area. These results support a generalmechanism for nanoscale material degradation and dissolution(e.g., crevice corrosion) of polycrystalline nonnoble metals, al-loys, and inorganic materials within confined interfaces.

crevice corrosion | surface forces apparatus | pitting dynamics |surface electrochemistry

Reactivity in confinement is central to a wide range of appli-cations and systems, such as geological and technological

processes, for example, pressure solution (1, 2), chemical mechani-cal polishing (CMP) (3), electropolishing (4), and corrosion (5, 6),including biocorrosion, for example, implants and dental fillings (7,8). In some cases, increased dissolution rates are desired, such asCMP; however, in other systems, particularly involving pitting andcrevice corrosion, increased reactivity leads to material degradationand damage. In each of these systems, the confinement (close ap-position) of two asymmetric surfaces in an aqueous environment is acommon factor where increased reactivity is observed. However,unlike systems exposed to bulk solutions, the difficulties introducedby reactivity and transport (e.g., diffusion) in confinement lead tocomplications in developing an understanding of reaction processesunder these conditions. Furthermore, few techniques are able todirectly measure and study reactivity in confinement in situ at relevantlength scales, leaving even more uncertainty surrounding the initiationand early growth of dissolution or corrosion processes in confinement.As a result, the understanding of early-stage reactivity in confinementremains challenging and more detailed nanoscale characterizationmethods are required to develop a deeper understanding.From an electrochemical viewpoint, crevice corrosion––the fo-

cus of this article––and pitting corrosion are related phenomena

(9–11), although there are significant geometrical differences be-tween them. Pitting corrosion is driven by an aggressive ionic en-vironment and oxygen depletion in the stagnant electrolyte ofgrowing pits on a surface exposed to bulk solution (12–14). Crevicecorrosion is a related form of localized/confined corrosion that mayoccur within cracks and crevices where depletion and reducedtransport of species such as oxygen may lead to the development ofaggressive media and local corrosion potentials (15). Within crevicecorrosion, the initiation and growth of pits are observed and,throughout this paper, we will commonly refer to pits within thecrevice (which is a separate phenomenon from pitting corrosion).Crevice corrosion is of particular concern for stainless steels and

nickel alloys in contact with chloride-containing solutions, com-monly found in marine environments (5, 6). Chloride anions can befound in a wide range of environments such as seawater, industrialraw water, and sludge. In particular, the potential gradient thatdevelops in crevice corrosion drives cations out of the crevices andanions, such as chloride, into the crevices (14). The diffusion ofions into or out of crevices depends on their mobility. Chloride hasa very high mobility and it is often present in high concentration increvices. As such, chloride anions tend to accumulate inside thecrevice while the more hydrated and larger cations (e.g., sodium),which typically are less mobile, remain mainly outside the crevice.In addition, the hydrolysis of the dissolving metal ions leads to astark local variation of the pH (11, 12, 16). The dissolution prod-ucts of the material (metallic cations) are also driven out of thecrevice, which leads to a local change of the pH.

Significance

A surface forces apparatus was used with an electrochemicalattachment to directly visualize electrochemical reactions, disso-lution, and pitting on surfaces when confined in nanoscale gaps(or “crevices”) in real time. The results provide unique insightsinto how two closely apposed surfaces degrade under confine-ment, which is different from degradation of surfaces exposed tobulk solution. Degradation proceeds via local nucleation of “pits”at the periphery or center of the “contact area” that rapidly growoutward (in radius) and, more slowly, in depth, until the pitsconverge and the whole area has become one large pit. The re-sults on a metal–inorganic interface are similar to previousobservations on inorganic interfaces, and support a generalmechanism of “pitting corrosion” at confined interfaces.

Author contributions: K.K., H.A.D., J.N.I., and M.V. designed research; C.M., H.-W.C., K.S.,and Y.-J.C. performed research; C.M., H.-W.C., and M.V. analyzed data; and C.M., K.K.,H.A.D., and M.V. wrote the paper.

Reviewers: J.F., Johns Hopkins University; and P.M., Chimie ParisTech-CNRS.

The authors declare no conflict of interest.1To whom correspondence should be addressed. Email: jacob@engineering.ucsb.edu.

This article contains supporting information online at www.pnas.org/lookup/suppl/doi:10.1073/pnas.1708205114/-/DCSupplemental.

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Ex situ macroscopic studies of pitting corrosion, crevice cor-rosion, and uniform corrosion of nickel alloys in differentaqueous media and conditions were reported by various authors(17–19). However, real-time and in situ nanoscale investigationsof localized corrosion (both pitting and crevice corrosion) remain achallenge. In particular, localized corrosion processes take place atvery small length scales and highly localized initiation sites thatmay be buried or hidden within the confined geometry of twoclosely apposed surfaces. A direct, in situ view into the variouscrevice corrosion processes at progressing cracks is experimen-tally difficult. Recently, significant advances have been madeusing X-ray tomography (20); however, resolution limits stillinhibit characterization of the onset of corrosion. Typically,electrochemical pitting corrosion and crevice corrosion tests arecarried out in a system-relevant environment and conditions, andthe sample surfaces are examined ex situ using a variety of dif-ferent materials characterization and imaging techniques. Initialpit formation on an extended surface can also be studied in situby locally resolved probe techniques, such as spectroscopy orX-ray tomography, but it is not straightforward to detect pittingcorrosion since the initiation sites are sparsely distributed andtypically hidden inside a confined area where two surfaces areseparated by nanometer-sized gaps or crevices which are difficultto resolve (temporally and spatially).Previous observation using a modified electrochemical surface

forces apparatus (EC-SFA) on nonmetallic surfaces (e.g., silica) ofelectrochemical reactions, dissolution, and in situ observation of, forexample, pit formation, has been demonstrated as a valid method tostudy reactivity in confinement using multiple beam interferometry(MBI) (21, 22). We further demonstrated angstrom-resolved real-time visualization of the initial corrosion behavior in a confinedenvironment using white-light interferometry to study corrosion ofconfined aluminum surfaces (23). Here, we extend our studies tocrevice corrosion of nickel. In addition to our previous method, acamera was added to visualize and quantify the crevice corrosionprocess with simultaneous optical microscopy and interferometry(MBI) in real time. We studied the effects of concentration andcharging dynamics on the corrosion initiation by changing the sur-face potential/charge ramping rate, and following the initial disso-lution and pit formation of confined nickel surfaces in chloridesolutions. Based on our experimental results and methodology, wepresent a detailed mechanism for the initiation and initial growth ofpits during nickel crevice corrosion, which may provide generalinsights into nanoscale pit formation and growth in other systems.

Materials and MethodsChemicals. All chemicals were purchased from Sigma-Aldrich at the highestpurity available (sodium chloride, 99.9% pure). Sodium chloride (NaCl) so-lutions (1, 10, and 100mM) were prepared with Milli-Q (Millipore) water witha resistivity ≥18 MΩ·cm and a total organic carbon (TOC) value ≤ 2 ppb.Optical grade 1 ruby-mica was obtained from S&J Trading Company assheets of about 20 cm × 20 cm and 2-mm thickness. Surfaces were preparedusing 5–10-cm2 mica sheets with thicknesses ranging from 2 to 5 μm whichwere hand-cleaved from raw sheets and used for the experiments (24).Physical vapor deposition (PVD) back-silvered highly flexible mica sheets of2–6-μm thickness were then glued to cylindrical silica disks of nominal radiusof curvature R = 7–15 mm. Semitransparent Ni thin films with a thicknessTNi = 30 nm were prepared by PVD at deposition rates of ≤0.5 Å/s. Themicrostructure of these films is dominated by grains that are below 20 nm,as estimated by electron microscopy imaging.

SFA. SFA measurements were performed at 21 °C in a cleanroom using the SFA2000 (SurForce LLC). One cylindrical glass disk with an atomically smooth, back-silvered muscovite mica surface and an apposing glass disk with a 30-nm Nilayer were mounted in the SFA in cross-cylinder geometry. This geometry cre-ates a two-layer interferometer with the Ni surface facing the solution side,while the back-silvered mica surface acts as the crevice former. Zero distance(D = 0) is defined as contact between mica and nickel in dry air (see theschematic in Fig. 1B where ΔtE = 0). Any changes of the separation distancebetween the Ag and the Ni “mirror” surfaces due to, e.g., ingress of water, or

shift of the Ni-mirror surface due to corrosion (conversion into oxide or disso-lution), can be tracked in situ by following both the distance shift of the mirrorand the change of the light intensity (due to the thinning of the reflectingnickel layer). The setup is sensitive to ≈25–50-pm absolute shifts of mirror dis-tances, relative to the zero distance, D = 0, defined under dry conditions. Theelectrochemical three-electrode attachment consists of a Ni working electrodeand Pt counter- and reference electrodes. All measurements were performed inNaCl solutions of different concentrations with a pH of 7–7.5. Every Pt pseu-doreference was referenced against a standard 3N AgjAgCl electrode (23),showing a difference of about +400 mV with typical variations of 50–80 mVdepending on the Pt wire used. All data are referenced to the open-circuitpotential, which is set to zero at −450 mV vs. AgjAgCl. Here, Pt pseudorefer-ences are used to limit concentration changes within the small volumes of NaClelectrolytes between the surfaces (e.g., in the crevice). Chloride-based referenceelectrodes potentially leak chlorides and may substantially change the con-centration in small volumes (as in our cell) and, in particular, in small concen-trations (here, e.g., 1 mM). As such, Pt wires are a balance between accuracyand chloride concentration stability.

Apparatus Modifications. The EC-SFA has been adapted for characterization ofcorrosion processes in confinement as shown in Fig. 1. In addition to astandard EC-SFA setup (22, 25), an optical camera connected to a 10× mi-croscope was added to simultaneously record both (i) optical images of thecorroding confined areas and (ii) shifts of the corroding metal mirrors viawavelength shifts in the fringes of equal chromatic order (FECO) (26). Atypical FECO image showing a flattened contact is depicted in Fig. 1B. The 2DFECO pattern or spectrum measured in the spectrometer simultaneouslyprovides a spectroscopic axis to directly measure the distance between thetwo apposing mirrors with a resolution of 0.1 nm and a lateral axis for anoptical image of a section through the contact region with a resolution ofabout 0.7 μm. For more details on the use of FECO, we refer the reader to anarticle by Israelachvili (26). A 50/50 beam splitter equally divides the lightbetween the camera (i) and the spectrometer (ii). A round-shaped confinedarea was established by pressing a back-silvered mica surface, as a creviceformer, onto a nickel-coated glass disk (30-nm Ni film, PVD deposited) aspreviously described (23). In addition, (iii) a three-electrode setup with nickelas the working electrode (WE) and two Pt wires as the reference (RE) andcounter electrodes (CE) are built into this setup.

Experimental Procedure. First, a dry contact (defined as D = 0) was establishedbetween the mica and the nickel surface (direct contact). Afterward, aqueousNaCl solution was injected between the two surfaces at concentrations rangingfrom 1 to 100mM. FECO and real-time videowere recorded simultaneously as theelectrochemical potential was ramped up from open-circuit potential (OCP) =0 mV (i.e., −450 mV vs. AgjAgCl) at varying rates (Fig. 1C). During a prerampperiod of 120 s after the electrolyte was injected, the ingress of electrolyte wastraced and the OCP was recorded to confirm a proper working electrochemicalcell. Any changes of the separation distance between the Ag and the Ni mirrorsurfaces due to, e.g., ingress of water, or a shift of the Ni-mirror surface due tocorrosion (conversion into oxide or dissolution), can be tracked in situ by moni-toring three signals: (i) the distance shift of the mirror (wavelength shift of astanding wave), (ii) the change of the light intensity (i.e., more light passingthrough the nickel mirror due to the thinning of the reflecting nickel layer), and(iii) the width of the FECO fringes, which provides information about theroughness of the reflecting Ni surfaces. As such, all signals and parameters to-gether allow us to distinguish clearly the difference between roughening fromoxidation or water ingress. Water ingress will shift the FECO fringes to longerwavelengths, while an oxidation shifts the FECO fringe wavelength and increasesthe intensity of the light going through the nickel mirror. Moreover, pureroughening leaves the FECO fringe wavelength position unchanged (because theaverage surface height does not change), yet the width of a FECO fringe increasesproportionally with the rms roughness. As of now, the setup is restricted to cor-roding thin films studied in transmission mode; extensions to bulk samples areeasily possible using reflection mode interferometry (27).

Nickel Thickness Determination. For analyzing and determining Ni thicknesses,we can use both white-light interferometry and intensity analysis of the lightpassing through the interferometer. First, to translate FECO wavelengthshifts recorded with the imaging spectrometer using white-light in-terferometry into distance changes, we use an in-house-developed software(SFA Explorer from Max-Planck-Innovation). The SFA Explorer uses themultiple-matrix method, as proposed by Schubert (28), to simulate and fitinterferometric spectra against measured spectra. To fully visualize evolvingcrevice corrosion in 3D, the part of the surface that is projected into thespectrometer slit can be rastered back and forth to obtain FECO for the

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entire crevice area. Lateral scans of the crevices can be performed with alinear motor, moving at a given speed (here, 4.8 μm/s), while the frame ratewas kept at 10 frames per second (fps) providing a nominal 500 nm reso-lution with 1-fps time resolution. With this setting, the spectroscopic di-mension (wavelength and light-intensity changes of 2D slices) can bescanned with frame rates up to 30 Hz. However, the third lateral microscopicdimension can only be scanned at a considerably slower rate (∼1 Hz)––takingabout 100 s for the entire crevice scan. As such, the FECO can only berecorded fast enough to visualize initial corrosion at a static position. Hence,we also implemented a faster Ni-thickness determination method based onusing the intensity changes in the optical image (top-view camera) taken ata given wavelength. For this, we calculated the intensity of typical standingwaves as a function of the Ni thickness using the multiple-matrix method forthe given interferometer setup. This reveals a linear relation of intensity andNi thickness (Fig. S1). Based on this result, a linearized attenuation modelwith an effective extinction coefficient can be used to directly extract the Nithickness from optical images with ∼0.5-nm depth and ∼500-nm lateralresolution. Over the entire crevice area this analysis can be performed in realtime with frame rates up to 0.1 kHz (depending on the top-view camera).

Results and DiscussionA typical corrosion experiment starts with the mica and nickelsurfaces pressed together under dry condition to form a circular-shaped crevice with about 300-μm diameter. The electrolytesolution is added while the nickel electrode remains at OCP. Fig.S2 shows how ∼2–3-nm electrolyte ingresses into the confinedarea (i.e., where the surfaces are in close contact). This can bedirectly inferred from a positive mirror shift, detected by a redshift of the FECO fringes, and was routinely recorded during thewetting of the crevice. One hundred and twenty seconds afterwetting the crevice and establishing a stable OCP, a potentialramp with +1 mV/s was started from the OCP toward anodicpotentials. There are a number of interesting details that weobserve during the corrosion of the 30-nm-thick nickel surface inthis configuration, which we show and discuss below.With the EC-SFA, one can track and quantify in real time the

crevice corrosion mechanism upon application of any (here, an-odic) ramp. The real-time imaging (recorded with the top-viewcamera; see again Fig. 1A) of crevice corrosion initiation shown inFig. 2 indicates significant qualitative differences in the corrosionmechanism depending on the chloride concentration and the ap-plied ramp rate. At high chloride concentration (>10 mM), shownin Fig. 2 A and B, corrosion initiates at the crevice mouth while theinner part of the confined zone remains intact at 1 mV/s ramp rate;this qualitative effect reverses, with corrosion initiating inside the

confined zone, at slower ramp rates of 0.5 mV/s shown in Fig. 2 Cand D. Furthermore, at low chloride concentrations (<10 mM),shown in Fig. 2 E and F, corrosion initiates inside the confinedzone where the surfaces are separated by only a few molecularlayers of solvent (∼2.5-nm liquid film; see again Fig. S2), irre-spective of the ramp rate (from 0.5 to 10 mV/s).The rate-dependent behavior of the crevice corrosion, as seen

in Fig. 2 B and D, shows the connection between active ionmobility due to the applied potential rate. At high potential rates(>1 mV/s), the chloride ions accumulate at the rim (contactboundary) enhancing the crevice corrosion rate while lower po-tentials (<0.5 mV/s) allow ions to diffuse into the crevice beforecrevice corrosion is initiated.As expected, this behavior supports a simultaneous potential

gradient and ion migration kinetic-driven mechanism of crevicecorrosion. In particular, crevice corrosion has been explainedpreviously based on a combination of the ohmic drop (iR drop)into the crevice (29) and the location where an aggressive so-calledcritical crevice solution (CCS) with low pH [formed from the hy-drolysis of Ni due to corrosion (30)] and high chloride concen-tration (ion migration to maintain charge neutrality in the crevice)is created and maintained in relation to diffusion and migration ofsolution species (31). We find that the location of crevice corrosioninitiation moves toward the center of the crevice with an increasingiR drop at lower chloride concentrations. The observed crevicecorrosion behavior in the experiments fits the framework of the iRdrop theory because the potential inside the crevice will be lowerthan outside. This explains why crevice corrosion initiates and alsomaintains higher rates deeper into the crevice at lower solutionconcentrations. Conversely, at lower iR drops––i.e., at higherconcentrations––crevice corrosion is initiated closer to the crevicemouth. With the lower iR drop, faster ion migration into the rimregion (compared with migration into the center region) triggersthe formation of a CCS at a region closer to the crevice mouth(31). As anticipated from macroscopic observations, crevice cor-rosion depends on a delicate interplay of both the transport anddepletion kinetics of active species which profoundly change thequalitative corrosion mechanism.There are two electrochemical processes going on simulta-

neously here that need to be distinguished, which we now refer toas the “active” and “passive” processes. The active process involvesthe applied potential, U, that generated the current flow and isquantified by the iR values and changes as corrosion progresses.

A B

C

Fig. 1. Crevice corrosion setup using the SFA. Simultaneous (A) optical video microscopy and (B) white-light interferometry allow tracking of both opticalimages of corroding confined areas and depths hp of corroding areas via shifts of the corroding metal mirrors via wavelength shifts ðΔλ=ΔλD − λ0Þ of FECO,and via intensity changes due to metal mirror thinning. For establishing a confined round-shaped area (A) back-silvered mica surfaces are used as compliantcrevice formers. Typical contact areas have a diameter dC ∼ 100–300 μm. Depending on the contact force F and the resulting flattening of the compliant micasurface, a controlled pressure can be applied, resulting in a uniform equilibrium electrolyte film thickness ΔtE in the confined zone. In addition, (C) apotentiostat with a three-electrode setup is used to apply external potential ramps to the nickel surface. Here, nickel was used as the WE, and Pt wires wereused as the RE and CE. All data are referenced to OCP, which is −450 mV vs. AgjAgCl.

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However, at zero applied potential there also exists a passive or“resting” potential, U0, as in a capacitor or battery, even in theabsence of any current flow or electrochemical reactions. Thevalue of U0 is determined by the difference in the “double-layer”surface potentials of the two (dissimilar) surfaces––nickel andnonconducting mica in this case––as well as the aqueous distancebetween them and the Debye length (32). This potential, too, willchange as corrosion progresses and presumably also contribute tothe rate of pit growth (32). A simple calculation based on electric

double-layer theory (ref. 33, chap. 14) shows that a 3-nm-deep pitwill cause a change in U0 of 30–60 mV. We are not able toquantitatively establish which of these processes dominate the iRdrop or the change in the different potentials associated with thiscorrosion process.Our data not only provide a unique real-time view of this

expected behavior during crevice corrosion initiation, but an in situEC-SFA experiment enables the ability to track and quantifynanoscopic details of the initial growth mechanism. This includes

50 µm 50 µm 50 µm

c > 10 mM1 mV/s

c > 10 mM0.5 mV/s

NaCl NaClc > 10 mM

1 mV/s

NaCl

Fig. 2. Optical microscopy of confined nickel areas while applying a linear potential ramp with ΔU = 1 mV/s (Fig. 2). Crevice corrosion (CC) of nickel in NaCljust after injection (Top) and after severe corrosion occurred (Bottom): (A and B) At c > 10 mM with a ramp rate of 1 mV/s indicating corrosion along thecrevice rim; (C and D) at c > 10 mM with a ramp rate of 0.5 mV/s, indicating corrosion inside the crevice; and (E and F) at c <10 mM, indicating CC again insidethe confined zone. The contact (e.g., Newton rings) for each experiment appears visually different due to the different angles arising from different radii/curvatures of the metal surfaces. Interestingly, we do not observe any notable changes in the morphologies and rates of CC due to curvature effects, contactarea, and aqueous separation distances (in the range 1–3 nm).

Fig. 3. Reactivation of initially passivated pits above the critical corrosion potential. (A) Optical image (Left) and corresponding Ni thickness contour plot(Right) showing initial repassivating pits inside and outside the confined zone (region marked) after 600-s linear anodic polarization with 1 mV/s starting fromOCP (−450 mV vs. AgjAgCl). (B) Same confined zone after 820-s polarization (i.e., 820 mV vs. OCP) showing fast active pitting inside the confined area.(C) Enlarged Ni-thickness contour plot of pit that reactivates (A) region marked by square, and (D) depth profile as function of time recorded at line indicatedin C. Pit diameter dp(t) and pit depth hp(t) are indicated for the repassivated state. Corrosion will continue until consumption or repassivation of the nickelsurface. (E) Pitting density of repassivating pits referenced to the initiation potential Uinit ∼ 300 mV (vs. OCP) where the first pit was detected in the field ofview. Individual pit counting was terminated once a critical pit started fast corrosion.

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qualitative and quantitative changes of corrosion dynamics intolocal currents with nanoscale-depth (single active site) accuracy andmicroscale lateral resolution.Fig. 3 shows two snapshots of (Left) a microscopic image, and

(Right) a 3D analysis of the Ni thickness at progressing time andincreasing potential (+1 mV/s) (see also Movies S1 and S2). Thedisplayed experiment was recorded at 1-mM solution concen-tration of NaCl.First, we observe that crevice corrosion of mica-confined nickel

surfaces does not initiate as homogeneous, locally uniform corro-sion. One can observe the formation of nanoscale-to-microscalepits that evolve with exceptionally rapid kinetics (in the microsec-ond range) before they become metastable, both inside and outsidethe crevice, before the initiation of nickel crevice corrosion inconfinement.In particular, Fig. 3A shows an optical image indicating pits

with very similar diameters of dp(t) ∼ 6–7 μm. The first pit wasdetected about 300 s after anodic polarization was started (i.e., at300 mV vs. OCP). The average pit depth of these initial pits isabout hp = 2–4 nm and the bottom is flat, as can also be seen inthe depth profile shown in Fig. 3 C and D. These pits appearinside and outside the confined zone within less than 500 μs (onerecorded frame), indicating a localized and extremely rapid, yetself-inhibiting, corrosion process. These initial pits do not im-mediately penetrate the entire Ni layer as expected by typicallyestimated anodic pit-current densities (10). These pits formwithin less than 500 μs (i.e., within one recorded frame); how-ever, once a typical critical diameter dp(t) ∼ 6–7 μm is reached,the formed pits become inactive and the localized corrosionstops (i.e., arrests). We can directly calculate the local chargethat transfers in a single site based on a volumetric analysis ofnickel loss. From the estimated local Ni-thickness loss (Fig. 3 Cand D), we can calculate a charge transfer of 0.97 μC per pit,which estimates a minimum pitting-current density of >5 A/cm2

for these pits that appear within t < 500 ms. These directlymeasured single pit-current densities compare well to estimatesfrom penetration time measurements using thin metal films (34),where a range from 1 to 30 A/cm2 was estimated. These previousstudies found that the anodic pit-current density––derived fromthe penetration time––depends on the thickness of the metalfilm and the applied potential. Our data extend this view, anddemonstrates that initial––nonpenetrating––pits show ratheruniform diameters and depth, without any strong correlationwith the applied potential. Fig. 3E shows that these initial, oftenreferred to as noncritical, pits pop up with an exponentially in-creasing probability with the increasing potential. We argue thatthese noncritical pits may be central to establishing the localsolution conditions that are necessary to catalyze the corrosionrate of a critical pit that will penetrate the thin film. Consideringthe corroded volume of nickel in such a pit, we estimate that thesolution temporarily reaches very high concentrations.In addition, the low aspect ratio of pit height to diameter [hp(t)/

dp(t) ∼ 10−3] of these noncritical pits suggests that the initial lateralcorrosion is at least three orders of magnitude faster than theinitial vertical corrosion into the material. The lateral resolution(∼0.5 μm) constrains the direct observation of the nanoscalemorphology (e.g., structural or chemical defects), preventing anaccurate description of the pit initiation; however, other studieshave observed direct evidence of such defects using in situ near-field microscopy techniques (12, 35, 36). Combination of the ob-served circular-shaped pits and the flat bottom (i.e., counter-boreshape) suggests that a very local oxide breakdown––possibly atdefect sites––triggers the initial pit formation with epitaxial cor-rosion dynamics. Corrosion triggering defects may significantlylower the activation energy for passive layer dissolution. Once alocal defect or impurity in the metal or oxide triggers pit initiation,initially sharp and convex edges of an evolving pit generate a strongelectric-field gradient that accelerates lateral corrosion, rapidly

increasing the lateral diameter of the pit while the depth increasesabout three orders of magnitude more slowly (see Fig. S3 for adetailed proposed mechanism). Due to the confinement of theapposing surface, the solution within the pit becomes saturatedquickly with nickel solutes, which arrests the corrosion reaction.As shown in Fig. 3C, above a critical potential (here, at

+400 mV vs. OCP), an arrested pit within the confined zone (atthe rim or the center depending on the chloride concentration)breaks the initial pit size of dp(t) ∼ 6–7 μm and hp(t) ∼ 2–3 nmand rapidly evolves into a larger depth. Again, from the esti-mated local Ni-thicknesses loss we calculate a charge transfer of2.75 μC for this critical pit, which estimates a pit current densityof ∼25 A/cm2, which is up to 5× faster compared with the growthrate of the repassivating pits. From Fig. 3D the initial depthpenetration rate of this critical pit can be estimated at about25 nm/s or 90 μm/h.Moreover (Fig. 3B), the critical pit triggers and catalyzes a

rapid evolution of new pits as well as reactivates other arrestedpits in its immediate surroundings. This collaborative pit growthis most likely due to a strong variation of solute concentrationsinside the confined zone (see Movies S1 and S2 for movies of therapid evolution of pits) and the formation of a CCS with low pHand high chloride concentration.The collaborative pit-growth behavior supports the hypothesis

that crevice corrosion of nickel in chloride solutions initiates asself-catalyzed pitting corrosion due to chemical changes in thecrevice solution within the confined areas. As expected, theongoing metal dissolution inside the confined region leads to anexcess of positive charge in the crevice solutions, and to a low-ering of the pH due to metal ion hydrolysis. This excess of localcharge is largely compensated by the relatively faster migrationof chlorides into the confined zone, while metal ion diffusion outof the crevice is relatively slower due to greater ion mobility ofchlorides (31). Pits that are formed outside the confined crevicearea are all swiftly arrested, even at higher potentials, which maybe due to the lack of buildup of local aggressive electrolyte en-vironments (31, 37).Based on the corrosion rates found for single pits (together with

their measured dimensions and growth kinetics, as demonstrated inFig. 3), and for the entire confined zone we can finally compareintegral currents (inside and outside of confinement) to currents

Fig. 4. Electrochemically measured current density for experiment shown inFig. 3 (solid line). The current estimated from the pitting density of non-critical pits (triangles) as well as the current density originating from the fastCC within the confined zone are compared with the integral current (see thetext for details). At higher anodic potential of ∼800 mV (∼450 mV vs.AgjAgCl), polarization leads to the expected passivation of the outside re-gion. The solid red line is a guide to the eye for the noncritical pitting curvethat deviates from the Tafel slope, as indicated by the dashed red line.

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estimated from volumetric Ni-thickness loss for localized corrosionsites. Fig. 4 shows excellent agreement, indicating that the de-viation from the characteristic exponential Tafel slope (with i0 ∼8 × 10−7 A/cm2) is caused by the additional current density fromdefect breakdown into the observed noncritical pits. Also, thecorrosion in confinement is fast and consumes the entire confinedarea within ∼100 s. Interestingly, the fast corrosion of the confinedzone does not add significantly to the overall current due to thesmall area of our confined zone compared with the entire samplearea. Locally in confinement, however, the crevice corrosion cur-rent density ranges up to 300 A/cm2, which is a reasonable numbergiven that about 10–20 critical pits with anodic pit-current densitiesof ∼25 A/cm2 (see above) progress at the same time. As such, theEC-SFA provides a unique measure to in situ quantify the localcurrent density in confinement, which is independent of any com-petitive reduction reactions such as local hydrogen evolution.

ConclusionsIn summary, the EC-SFA provides a unique setup to directlyvisualize individual corrosion sites (here, pits) in confined spacesin real time and to analyze their initial growth kinetics. Com-pared with standard optical microscopy with transparent creviceformers, the depth resolution of evolving initial corrosive deg-radations is superior, allowing measurement of pit depths in 3Dwith subnanometer and millisecond time resolution. Also, theEC-SFA uniquely enables measurement of single pit-currentdensities and growth shapes in 3D (both laterally and in depth),as well as time-evolving pit densities within the confined zone.Initial corrosion inside nickel crevices proceeds as preferential

and self-catalyzed pitting corrosion (inhomogeneous localized dis-solution with the formation of pits), presumably at surface defectswithin the confined zone. Our data also confirm that there is a

strong interplay between migration kinetics of chlorides, bulkchloride concentrations, and the related iR drop (voltage drop),which leads to pronounced qualitative changes in the location (atthe rim or center) at which pits form, and their subsequent ki-netics. Importantly, local fast corrosion seems to play a crucial rolein the initial development of a CCS within the crevice that drivesthe observed pit growth and propagation, suggesting that un-derstanding and controlling surface chemistry and morphology areessential to prevent the onset of corrosion in the present system.Pit-current densities are ≥5 A/cm2 for swiftly arresting noncriticalpits, and ∼25 A/cm2 for critical pits.The above conclusions may be general observations for crevice

corrosion systems; however, further studies with other corrodingsystems are needed to potentially extract and verify general rules and/or models as well as system dependencies of local corrosion rates insitu with nano- to microscale resolution. Local currents originatingfrom reactions within the confined zone (with single pit resolution),and currents from outside the zone can be uniquely differentiatedusing the EC-SFA technique. In addition, the EC-SFA provideshighly resolved data suitable for direct comparisons with quantitativetheoretical models of the initiation of crevice corrosion for othermetal and alloys systems, enabling direct characterization of initialcrevice corrosion mechanisms.

ACKNOWLEDGMENTS. J.N.I., K.K., and H.A.D. were supported by the USDepartment of Energy, Office of Science, Basic Energy Sciences, under AwardDE-FG02-87ER-45331 (development of the experimental SFA and white-lightoptical interference imaging techniques, and interpretation of the results).C.M., K.S., H.-W.C., and M.V. thank the Max-Planck-International-Research-School SURMAT for financial support. K.S. and M.V. thank the EuropeanResearch Council (ERC Starting Grant 677663) for support. C.M. and M.V.thank the Deutsche Forschungsgemeinschaft for support (Project VM-03).

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