Interfacial voids in aluminum created by aqueous dissolution

Electrochimica Acta 55 (2010) 6093–6100

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Interfacial voids in aluminum created by aqueous dissolution S. Adhikari a , L.S. Chumbley b,c , H. Chen d , Y.C. Jean d , A.C. Geiculescu e , A.C. Hillier a , K.R. Hebert a,∗ a

Department of Chemical and Biological Engineering, Iowa State University, Ames, IA 50011, United States Department of Materials Science and Engineering, Iowa State University, Ames, IA 5001, United States c US DOE, Ames Laboratory, Ames, IA 50011, United States d Department of Chemistry, University of Missouri-Kansas City, Kansas City, MO 64110, United States e CRMD, St. Jude Medical, Liberty, SC 29657, United States b

a r t i c l e

i n f o

Article history: Received 18 January 2010 Received in revised form 14 May 2010 Accepted 24 May 2010 Available online 1 June 2010 Keywords: Aluminum Corrosion Positron annihilation Transmission electron microscopy Voids

a b s t r a c t Nanometer-scale voids in aluminum formed by aqueous room-temperature corrosion were detected and characterized by a combination of electron microscopy techniques, atomic force microscopy, and positron annihilation spectroscopy. Void-containing layers were found within 100 nm of the metal surface, containing voids of 10–20 nm width with a number density of 108 –109 cm−2 . The voids were generated continuously during dissolution. The rapid nucleation and growth of voids suggest elevated concentrations of hydrogen-vacancy defects near the dissolving surface. Using the measured void radius and thickness of the void layer, the hypothesis that voids grow by vacancy condensation led to reasonable calculated values of the diffusion coefficient, and vacancy concentrations in agreement with independent estimates. © 2010 Elsevier Ltd. All rights reserved.

1. Introduction Hydrogen absorbs rapidly into aluminum during roomtemperature alkaline dissolution and cathodic electrochemical polarization. Bulk concentrations of 0.1 at.%, several orders of magnitude higher than the room-temperature solubility, have been measured after extended hydrogen charging [1,2]. Even higher near-surface H concentrations in alloy 7050 were found after humid air exposure at 90 ◦ C [3]. Formation of aluminum hydride during alkaline corrosion, suggestive of large hydrogen chemical potential at the metal surface, was detected with secondary ion mass spectrometry (SIMS), and also inferred from dissolution potential measurements [4–6]. The elevated H chemical potential was confirmed recently with electrochemical potential measurements, in hydrogen permeation cells with AlPd bilayer membranes [7]. Significant microstructural changes result from these high hydrogen absorption rates. Several investigators reported the nucleation of large numbers of 10–100 nm scale metallic voids or hydrogen bubbles during H charging. Transmission electron microscopy (TEM) revealed voids in the bulk of pure Al, after charging for several days [8]. The presence of voids close to the dissolving surface was indicated by positron annihilation spectroscopy (PAS), after aqueous dissolution for times of minutes [9–12]. Supporting

∗ Corresponding author. Tel.: +1 5152946763; fax: +1 5152942689. E-mail address: [email protected] (K.R. Hebert). 0013-4686/$ – see front matter © 2010 Elsevier Ltd. All rights reserved. doi:10.1016/j.electacta.2010.05.073

microscopic evidence for these near-surface voids was obtained, through TEM and atomic force microscopy (AFM) examination of anodic oxide films formed by reaction of the defective metal region [13]. Other TEM studies revealed bubbles and blisters at the metal/oxide interface after extended cathodic charging of Al [14], and blisters at the metal–oxide interface after humid air exposure [15]. Both voids themselves and the process of void formation are relevant to localized corrosion. PAS S-parameter measurements closely resemble those of clean Al surfaces, as opposed to oxidecoated nanopores [17,18]. The oxide-free void surface suggests that near-surface voids can act as highly reactive initiation sites for localized corrosion, if exposed by dissolution of overlying metal. In fact, void populations strongly correlate with nucleation rates of corrosion pits during anodic etching of Al [20], and industrial etching processes rely on dissolution pretreatments to dramatically enhance pit nucleation [21]. Understanding the formation of hydrogen-induced voids can also provide insight into the state of hydrogen in the metal during charging and embrittlement processes, e.g. whether it is bound to vacancies, or occupies interstitial sites [22]. The present article reports characterization of nucleation and growth rates of voids during room-temperature alkaline dissolution of Al. These observations provide the basis of a quantitative understanding of void formation, when coupled with other transient measurements of mechanical stress to be reported separately [19]. The present investigation employed several complementary

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methods with common sample types and dissolution treatments. TEM, scanning transmission electron microscopy (STEM), scanning electron microscopy (SEM) and atomic force microscopy (AFM) identified the defects, and provided estimates of their formation rate and depth distribution at localized areas. New PAS results revealed changes, attributable to dissolution, of the defect depth distribution averaged across the entire Al surface. Non-destructive PAS measurements also reinforce conclusions from microscopy by addressing concerns about artifacts introduced during sample preparation. The results are discussed with respect to recognized mechanisms of voids and hydrogen bubble formation in Al [22]. 2. Experimental The aluminum samples were 110 ␮m thick foils of 99.99% purity (Toyo). The final processing step of these foils is an extended anneal, which results in a grains with typical size 100 ␮m and predominantly (1 0 0) surface orientation. While the details of this treatment are proprietary information, a typical set of annealing conditions for capacitor foils might be 550 ◦ C, 10−3 Pa, and 5.5 h [23]. Prior to the alkaline dissolution, all foils were electropolished in a 20% perchloric acid (70%) and ethanol (98%) bath at 5 ◦ C for 5 min. Electropolishing removed surface layers containing voids present in the as-annealed state of these samples [13]. The alkaline treatment was carried out in aerated 1 M NaOH solution at open circuit, for various times at 21 ◦ C. After the dissolution, samples were removed and rinsed with deionized water to stop the reaction, and then air dried. In all experiments, solutions were made from reagent grade chemicals and nanopure water. Doppler-broadening PAS measurements were carried out using a slow positron beam system at the University of Missouri-Kansas City. The positron beam was produced by a 50 mCi 22 Na source. The beam implanted positrons within the sample to an energydependent mean depth given by zm = 14.8Eb1.6

(1)

where Eb is the beam energy in keV and zm is the mean implantation depth in nm. For each beam energy, the spectrum of Dopplerbroadened annihilation radiation was measured at 2000 cps, using a single Ge detector having an energy resolution of 1.5 keV at the annihilation photopeak energy of 511 keV. The acquired spectrum included about 106 photon counts in the photopeak. Near 511 keV, the spectrum is determined by the momentum distribution of electrons participating in annihilation events. We characterized the peaks by the two Doppler-broadening peak shape parameters S and W. S represents the fractional area of the central portion of the peak close to the maximum, and is due to annihilation with low-momentum valence electrons. W is the fractional area at the extremes of the photopeak, contributed by annihilation with highmomentum core electrons. S and W were calculated by the system software, to within a relative accuracy of 0.001. Defective regions of the sample are associated with high S and low W values, relative to defect-free regions. Readers interested in greater detail about PAS are referred to recent reviews [24–26]. To prepare samples for TEM and STEM observations, NaOHtreated or as-electropolished samples were thinned from the unreacted back side using a single-jet electropolisher, until the sample perforated. The treated side was not exposed to the polishing solution during the thinning process. The electron-transparent regions close to the perforation were then imaged using TEM (Philips CM30) and STEM (FEI-Tecnai G2-F20). The thickness of these regions was roughly 100 nm. SEM and AFM observations were carried out after first removing controlled thicknesses of metal from the surface. After electropolishing and alkaline dissolution, Al foils were anodically oxidized in

a solution of 0.1 M boric acid and 0.05 M sodium borate (pH 8.8), at a constant applied current of 1 mA/cm2 , up to voltages of 7, 31 and 69 V against the Pt counter electrode. This procedure forms planar anodic alumina films, characterized by a constant “anodizing ratio” of 1.2 nm/V, of their thickness to the anodizing voltage. Since all the anodizing charge contributes to film growth at these experimental conditions, the charge determines the reacted depth of metal. The selected anodizing times correspond to reacted metal depths of 5, 23, and 50 nm. After anodizing, the oxide was stripped in a 5% chromic–20% phosphoric acid bath at 70 ◦ C for 2 min. This stripping bath efficiently removes the surface oxide, and replaces it with a protective chromium-containing film inhibiting dissolution of the underlying metal. For SEM, samples were examined using a field emission instrument (Hitachi S4800). The accelerating voltage was 15–25 kV and the beam current was 10 ␮A. Atomic force microscope images were acquired using a Dimension 3100 scanning probe microscope with Nanoscope IV controller (Veeco Metrology). Imaging was performed in air using tapping mode with silicon AFM tips (TESP7) having a spring constant of 70 N/m and a resonance frequency of 280 kHz. 3. Results 3.1. Positron annihilation spectroscopy The present PAS measurements of alkaline-treated electropolished samples followed procedures of earlier experiments using as-annealed foils [9,10,12]. The latter samples contained many large near-surface voids along rolling lines, which were removed by electropolishing. The reduced population of near-surface voids in electropolished samples permitted improved characterization of defect formation during the initial stages of dissolution. Electropolishing was used as a common surface pretreatment in all the analytical techniques used to characterize hydrogen-charged samples, including the electron optical methods in the present article and also SIMS measurements of hydrogen-containing species reported earlier [4,5]. Plots of S parameter vs. beam energy revealed near-surface defective regions in each sample tested. Fig. 1 shows examples of such S-energy profiles for as-electropolished Al foil, and for samples treated in 1 M NaOH for 10 s and 15 min. All annihilation shape parameters were normalized against the bulk value, which represents a defect-free reference state in our annealed samples. The top axis shows the depth corresponding to the beam energy, according to Eq. (1). As was the case for as-annealed Al foils [9,10,12],

Fig. 1. PAS lineshape parameter S vs. beam energy of Al after different times of sodium hydroxide dissolution. Error bars on S are smaller than the symbol size. The top scale represents the mean implantation depths of the positrons.

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Fig. 3. Defect layer parameters (Sd and Bd ) vs. NaOH dissolution time. Fig. 2. Plot of experimental W and S parameters of Al for various NaOH dissolution times. Error bars on S and W are smaller than the symbol size.

the S-energy profiles display S values less than one close to the surface, and regions of S larger than one at greater depths. The near-surface S values less than one are attributable to a surface oxide of a few nanometers thickness [27]. The region with S greater than one extends to implantation depths of about 600 nm, much greater than the oxide thickness. S values larger than one are attributable to annihilation with low-momentum valence electrons in the Al metal, in open-volume defects such as vacancies and voids. Comparable evidence of defect-containing regions were revealed by S-energy profiles of as-electropolished foils, and after alkaline treatments up to 30 min. Open-volume defects can be identified by plots of W vs. S parameters [28]. Such a plot is shown in Fig. 2, containing data for all experiments for dissolution times through 16 min. All points fall along two straight-line segments connecting three vertices. The vertices represent the S and W values of annihilation “states,” i.e. material phases or defects with characteristic electron momentum distributions. At points along the lines connecting the vertices, the positrons annihilate at either of the two endpoint states. In Fig. 2, the leftmost vertex with low S and high W represents annihilation within the oxide layer (low energy in Fig. 1), and the vertex with S and W close to one corresponds to annihilation in the defect-free bulk Al (large-energy limit in Fig. 1). The rightmost vertex with S and W parameters of 1.045 and 0.940 is due to the near-surface open-volume defects identified in Fig. 1. In similar dissolution experiments using as-annealed foils and a higher-resolution detector, S–W plots identified larger defect S parameters in the range 1.06–1.08 [10,11]. These values pertained to all dissolution times up to 30 min, when it is unlikely that the measurements would be influenced by the initial surface condition of the foil. Hence, the present defects may be the same as those detected earlier with S of 1.06–1.08. The S parameter of vacancies in Al has been reported as 1.027, significantly lower than the present measurements [29]. S values much larger than those of vacancies can be explained by the presence of positronium (Ps), an electron–positron bound state which can be formed in nanoscale cavities or at open surfaces [24]. Ps decays by a combination of twophoton annihilation (p-Ps), and three-photon annihilation (o-Ps). p-Ps produces photons very close to 511 keV and thus yields very large S, such as the value of 1.10 obtained on the clean Al(1 1 0) surface in vacuum at 400 K [17]. The somewhat smaller defect S parameter after dissolution can be explained by confinement of p-Ps in nanoscale voids [16]. Unlike the results on Al(1 1 0), the present spectra revealed no evidence of off-peak annihilation indicative of the three-photon o-Ps process. However, o-Ps in voids often exhibits two-photon “pick-off” annihilation with electrons of atoms lining the void surface [24]. Pick-off o-Ps decay creates

photons at energies within the 511 keV photopeak, but much more evenly distributed than those formed by p-Ps. The defect S parameter significantly larger than one is distinct from expected S values for oxide-lined nanopores, which are smaller than one [18]. Bulk aluminum oxide has a low S of 0.9 relative to Al [27], and S in oxide pores is less than that of Al, even when enhanced by p-Ps formation. Since the S parameter is highly surface-sensitive, thin passive oxide layers on Al have S values similar to those of an oxide phase (such as the leftmost vertex in Fig. 2). Thus, we conclude that the high-S defects in Fig. 2 are voids in Al with oxide-free surfaces. Such voids could not have been produced by dissolution, since oxide would have formed on any Al surface momentarily exposed to solution. In order to obtain quantitative information about the depth distribution of defects, we simulated the PAS results with the diffusion–annihilation equation for positrons in a solid. The simulations utilized the software application VEPFIT [30]. The simulation accounts for the beam energy-dependent depth distribution of implanted positrons, diffusion of thermalized positrons in the solid, annihilation in the bulk metal and trapping into defects. All these processes have been independently characterized in Al, except trapping, which depends on the sample-specific defect depth distribution. A two-layer model was used to characterize our samples, consisting of a bulk defect-free metal layer and a surface-adjacent defect layer. The defect concentration in the latter layer was independent of both depth and lateral position. The corresponding fitting parameters were the defect layer thickness (Bd ), characteristic S parameter (Sd ) and diffusion length (Ld , determined by the trapping rate constant and positron diffusion coefficient), along with the surface S value which accounted for the oxide layer. The solid lines in Fig. 1 represent calculated S-energy profiles, and may be seen to follow the measurements precisely. Fig. 3 shows the fit values of defect layer S parameter and thickness, for dissolution times up to 16 min. Sd increases with both the defect S value and their concentration within the defect layer. Since the S–W plot indicates the presence of only one type of defect, the variations of Sd in the figure reflect changes in the concentration. Thus, Fig. 3 indicates that the defect concentration increases sharply in the first 10 s, and reaches a maximum at 3 min, before decreasing gradually. The significant initial increase of Sd shows clearly that dissolution creates open-volume defects. The defect layer thickness varied between 50 and 80 nm. Because the defect size is unknown, we cannot determine whether the defects are large voids which span the entire defect layer, or small vacancy clusters which permeate throughout the layer. Since the Al foil dissolves at a rate of about 120 nm/min [4], the metal in the defect layer present at any given time is dissolved in approximately 30 s. The decrease of Sd between 3 and 10 min implies that the rate of continuous defect creation is smaller than that of removal by dissolution.

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Fig. 4. Open-circuit potential transient during NaOH dissolution, along with defect − layer S parameter Sd From Fig. 3. The inset shows AlD depth profiles measured by − SIMS after dissolution of electropolished foil in 1 M NaOD (D2 O) [4,5]. AlD counts − were normalized with the counts of Al2 in the bulk of the metal.

Alkaline dissolution is accompanied by the formation of AlH3 [4,5]. Hydride at the metal-oxide interface was detected by SIMS as AlD− , after dissolution of electropolished foils in 1 M NaOD (D2 O). Fig. 4 displays AlD− depth profiles during the first minute of dissolution, along with Sd from Fig. 3, and the open-circuit potential transient [5]. The Nernst potential for hydride oxidation is approximately −1.85 to −1.95 V for these conditions, close to the range of potentials in Fig. 4[4,6]. Thus, hydride is electrochemically stable at the Al surface at times after the potential minimum. The SIMS profiles indicate that hydride is present at the interface during the rapid increase of Sd in the first 10 s of dissolution. The possible relationship between surface hydride and the process of void formation, as suggested by Fig. 4, is explored in separate articles [7,19]. 3.2. Transmission and scanning transmission electron microscopy TEM and STEM imaging were carried out to reveal the morphology, geometry and location of the defects. TEM is especially suitable for voids, since the specific imaging properties of voids enable their direct identification. The TEM samples were electropolished and then immersed for 3 min in the alkaline bath. This dissolution time

was selected because it corresponds to the maximum Sd value (Fig. 3), and hence the highest defect concentration. Fig. 5 shows TEM micrographs of the same region of a sample, at two slightly different angles relative to the Bragg diffraction condition. The image reveals the presence of small circular features 10–20 nm in size (marked with white arrows). Tilting the sample about the Bragg diffraction condition produced changes in contrast consistent with what is expected for small voids [31], i.e. the features remained in view as the contrast changed from light to dark. Tilting in this manner is equivalent to obtaining through focus images, a method commonly used to identify voids. The response of the voids to tilt was substantially different from the contrast exhibited by nearby dislocations, which could be made to become invisible, while the voids were always visible. A dislocation showing typical alternating black/white contrast is indicated by the black arrow in Fig. 5(b), which undergoes extinction (absence of contrast) in Fig. 5(a). The PAS results indicate that while defects are present after electropolishing, substantially more are introduced by the alkaline treatment. Images of as-electropolished samples contained a few voids like those in Fig. 5. The number of voids, in each type of sample, was estimated by examining a total imaged area of approximately 25 ␮m2 , in about 40 micrographs. The order of magnitude of the void number density was 108 cm−2 in the NaOH-treated sample, and at least ten times smaller in the as-electropolished foil. This dramatic increase of the void population due to alkaline dissolution parallels the increase of defect layer S parameter from 0 to 3 min in Fig. 3. It is clear that the voids result from alkaline dissolution treatment, as opposed to electropolishing during TEM sample preparation. To provide additional evidence ruling out beam-induced void formation, the TEM study was repeated using a lower acceleration voltage of 150 keV, as compared to 200 keV as in Fig. 5. Again, as shown in Fig. 6, similar void features 10–20 nm in size were observed in the NaOH-treated sample (marked with arrows). It was much more difficult to find any void features in the TEM samples of the untreated sample at the same voltage, suggesting that the overall number of voids was smaller. Further imaging of the TEM samples was carried out with STEM, because the latter technique samples larger areas, and can reveal correspondences between void sites and other features. Fig. 7 shows a STEM image of the same general region of the sample shown in Fig. 5. Although the image is less directly interpretable

Fig. 5. TEM images of Al treated in NaOH for 3 min. Images (a) and (b) are of the same region at slightly different tilt angles relative to the Bragg condition. Accelerating voltage was 200 keV.

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both electropolishing and alkaline etching. The small white circular objects in the image are distinct from the more diffuse contrast of the mosaic pattern. While it might be argued that the features are spherical precipitates, the high purity of the foils used and the corresponding TEM observations from the same regions, that revealed no signs of any second phase, argue against this possibility. Thus, these features are deduced to be voids, with contrast produced due to scattering of the scanned transmitted beam from the sides of the voids. The voids marked by arrows in Fig. 7 are located near or on the lighter gray ridge areas bordering depressions. Several other less prominent objects with similar shape may be found in the image, also at ridge sites. 3.3. Topographic imaging

Fig. 6. TEM image of Al treated in NaOH for 3 min. Accelerating voltage was 150 keV.

than those obtained from TEM, Fig. 7 indicates the presence of several roughly 20 nm diameter white circular objects, six of the most prominent of which are marked by white arrows. The image also reveals a pattern of light gray regions on the edges of the dark circular areas. This pattern is produced by contrast associated with the sample topography, which, as shown clearly in the SEM and AFM images in Section 3.3, consists of shallow ∼100 nm wide depressions bordered by ridges. In the dark field mode, the higher brightness of the light gray regions indicates greater scattering, and therefore would correspond to the ridges seen in the SEM images. Patterns of scallop-shaped depressions bordered by ridges are characteristically found after many Al dissolution processes, including

Fig. 7. STEM image of Al treated in NaOH for 3 min, for the same region as shown in the TEM image of Fig. 5. The white arrows mark features identified as voids.

The TEM and STEM study supports the identification of the PAS-detected defects as voids, as both measurements reveal parallel increases due to dissolution. These images, though, do not determine the void depth, because some metal may have been lost from the original surface during TEM sample preparation. Information about the void depth distribution was obtained by SEM and AFM examination of the foil surfaces, after removal of controlled nanometer-scale thicknesses of metal. As explained in the Experimental section, removal of small thicknesses of Al was accomplished by anodic oxidation to depths of several nanometers, followed by chemically stripping the oxide. Figs. 8 and 9 show field emission SEM images of Al foils which had been treated by the same procedure as in Figs. 5–7, i.e. electropolishing and 3 min alkaline dissolution. The reacted Al depths were 5 nm for Fig. 8,

Fig. 8. FE-SEM images of Al treated in NaOH for 3 min, with 5 nm metal anodized and the anodic oxide chemically stripped. (a) 45◦ stage tilt and (b) 0◦ stage tilt.

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Fig. 9. FE-SEM images of Al treated in NaOH for 3 min, with (a) 23 nm and (b) 50 nm metal anodized and the anodic oxide chemically stripped (0◦ stage tilt).

23 nm for Fig. 9(a) and 50 nm for Fig. 9(b). For the 5 nm depth, images at 45◦ and 0◦ stage tilt (Fig. 8(a) and (b), respectively) are presented, to indicate the three-dimensional surface topography. The same ridge-scallop surface pattern found in Fig. 7 is evident, which becomes less apparent with increasing extents of oxidation. The SEM images Fig. 8(a) and (b) show a large number of circular objects, with an estimated number density of 109 cm−2 . No such features were visible on samples which were not anodically oxidized, but which received the same electropolishing and alkaline dissolution treatments. Most of the objects are positioned near ridges created by dissolution, but some are also found in the interior of the scalloped regions. The majority of the circular features are smaller than 30 nm in diameter, similar to the size of voids in the TEM and STEM images. There are also a small number of larger objects (2–3 per image) with diameter approaching 50 nm. The number of circular 10–30 nm diameter features in SEM images decreased with anodic oxidation. Evidence of only a few features are found after oxidizing a 23 nm thick layer of Al (Fig. 9(a)), and none are apparent after oxidation to a depth of 50 nm (Fig. 9(b)). Thus, the features apparently originated from objects buried to depths no greater than 23–50 nm, as compared to the 40–50 nm defect layer thickness detected by PAS at similar dissolution times (Fig. 3). The similarity of these depths, as well as the close resemblance between the size, shape and location of the objects to those of voids in the TEM and STEM micrographs, support their identification as voids. However, it is not clear from SEM alone that the

Fig. 10. AFM image of Al treated in NaOH for 3 min, with 5 nm metal anodized and the anodic oxide chemically stripped. (a) Top view image. (b) Topographic section along the line in (a).

objects are concave cavities, as would be expected if the features are voids identified in TEM. Further imaging of these samples by AFM was carried out, to discern topographic details relevant to the interpretation of SEM. Fig. 10 shows an AFM top view image of an Al surface prepared by identical procedures as that in Fig. 8, which exhibits the same ridge-scallop topographic pattern found in SEM. The rectangular facets in the AFM image are probably caused by interaction of the surface geometry with the pyramid-shaped probe tip used in tapping-mode imaging. The height sensitivity of AFM highlights two distinct populations of scallop “cells”: relatively shallow cells less than 100 nm in width, and deeper ones with width roughly 200 nm. Interestingly, the smaller cells are similar to those comprising surface patterns formed by electropolishing [4], while the larger cells closely resemble those produced by alkaline dissolution [10]. This correspondence suggests that the alkaline dissolutiongenerated surface pattern may initiate at selected points, from which it spreads to cover the Al surface. Fig. 10 reveals a number of distinct recessed cavities with width and depth of 20–40 nm,

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located near the ridges of the larger scallop cells. The size of the cavities, along with their number density and typical locations near ridges, correspond well with the circular objects in Fig. 8. AFM detected no other features which might possibly explain the objects found in SEM. The agreement between SEM and AFM supports the view that topographic imaging revealed cavities corresponding to the voids detected by PAS, TEM, and STEM. The estimated number density of cavities is about an order of magnitude larger than the void number density from TEM. The number density from SEM is probably more reliable, as a larger area was sampled. Also, some near-surface voids may have been lost by electropolishing, during sample preparation for TEM. Additional voids may have formed during anodic oxidation of the SEM samples, as found in earlier work [12,27], but in this case, more voids should have been apparent in the samples oxidized to depths of 23 and 50 nm (Fig. 9(a) and (b)). The similarity of the void diameter and thickness of the voidcontaining layer suggests that upper surfaces of many voids are within a few nanometers of the metal–oxide interface. Therefore, the oxide-free surfaces of voids could be exposed during the course of uniform corrosion, explaining their ability to initiate rapid localized dissolution [20]. Fig. 9(a) indicates a number of additional roughness features smaller than the circular objects in Fig. 8. These small features may originate at other types of defects within the metal, but since no such defects were detected by TEM, it is more likely that they resulted from surface roughening during sample preparation. For example, roughening may have been induced by anodic oxidation through the void-containing metal layer.

4. Discussion The hypothesis that voids are produced by alkaline dissolution coherently explains the results of the five microscopic and analytical methods presented here. The TEM and STEM images show that the defects detected by PAS are voids with widths smaller than about 20 nm, and confirm the high rate of void formation detected by PAS in the first 3 min of dissolution. Estimates of the void layer depth from SEM are consistent with the defect layer thickness of 40–50 nm determined by PAS. The non-destructive nature of PAS excludes the possibility that voids are artifacts of sample preparation for microscopy. The number density of voids is 108 –109 cm−2 , and they are found preferentially near ridges on the ridge-depression topography created by dissolution. This correspondence indicates that vacancy defects and voids are formed as parts of the dissolution mechanism. For example, nucleation at ridges may be favored by the stress distribution in the metal during dissolution. Extensive studies of hydrogen-induced voids and bubbles in Al have identified three general nucleation and growth mechanisms: gas-containing blisters created by near-surface plastic flow, pressurized gas bubbles formed by dislocation loop punching, and clustering of vacancy-hydrogen defects [22]. Each mechanism requires elevated hydrogen chemical potential in the metal, as is present in our experiments. Bubbles in the first two mechanisms derive from interstitial H atoms formed by the cathodic reaction. No evidence of surface protrusions suggesting blisters was found in this work. In the third mechanism, dissolution of Al atoms creates H-vacancy defects, due to the large H-vacancy binding energy in Al [32,33]; vacancies then aggregate in the metal to form gas-containing voids. We cannot distinguish between the loop punching and vacancy mechanisms on the basis of morphology alone, as voids of similar size and shape as those reported here are found, irrespective of vacancy supersaturations [34–36]. The best way to differentiate these mechanisms would be through transient stress measurements, since a buildup of compressive or tensile

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stress, respectively, would accompany bubble or void nucleation. Tensile shifts of the bulk stress are already indicated by reported lattice parameter contractions [1,2], evidence which favors the vacancy-hydrogen defect agglomeration mechanism. We may ask whether the measured void sizes and distributions are consistent with the vacancy-clustering mechanism. If voids grow by vacancy diffusion, the void radius Rv should be approximately Rv ≈



2xv Dv t

(2)

where xv is the mole fraction of vacancies in the subsurface region and Dv is the vacancy diffusion coefficient, and the product xv Dv is the same as the diffusivity of Al atoms. The vacancy concentration profile may be modeled in terms of diffusion from a moving interface, in which case the concentration would decay exponentially according to xv ≈ xvs exp

 −v z  D Dv

(3)

where xvs is the vacancy concentration at the dissolving surface,

vD is the dissolution velocity, and z is the distance from the interface. Thus the estimated thickness of the void-containing layer is Bd ≈ 2Dv /vD . Using Rv = 10 nm, vD = 2 nm/s [4], t = 100 s, and Bd = 100 nm, we obtain xv = 0.005 and Dv = 10−12 cm2 /s. The order of magnitude of the vacancy concentration agrees with estimates from length change and lattice parameter measurements [1,2], and that of the vacancy diffusion coefficient is the same as that extrapolated from high-temperature data [37]. Thus, the microscopic observations seem to be consistent with void formation from vacancy-hydrogen defects produced by dissolution. The vacancy-clustering mechanism could also explain the elevated hydrogen chemical potential during dissolution, as the trapping energy of 69 kJ/mol would not permit H to diffuse out of the metal [32]. Otherwise, it would be difficult to reconcile the large subsurface H concentration with the high diffusivity of H interstitials [32]. Observations of hydrogen permeation accompanying dissolution could be explained by absorption of interstitial H from voids. As mentioned above, the vacancy mechanism implies that the subsurface region is under tensile stress. This stress can be estimated as the differential pressure between the interior and exterior of the void, according to the Young–Laplace equation, P = 2/Rv . Using the typical void radius of 10 nm and a value of 1 J/m2 for the specific surface energy , we obtain P = 200 MPa.

5. Conclusions We report evidence from electron microscopy and positron annihilation spectroscopy for the formation of nanoscale voids in aluminum during dissolution in alkaline solutions. Imaging using TEM, STEM, SEM, and AFM revealed voids of about 20 nm diameter, at depths up to tens of nanometers, which appeared in the first 3 min of after the metal was immersed. The number density of voids at 3 min was 108 –109 cm−2 . PAS indicated similar void layer thicknesses, and further demonstrated that voids form continuously during dissolution. The hypothesis that voids form by a vacancy condensation mechanism was evaluated by calculating the vacancy diffusivity and concentration from the void diameter and void layer thickness determined in this work. A reasonable estimate of the diffusion coefficient of 10−12 cm2 /s was obtained, and the calculated concentration of 0.5 at.% agrees with estimates based on length change and lattice parameter measurements [1,2]. The vacancies may be hydrogen-vacancy defects produced by the elevated hydrogen concentration in the metal during dissolution.

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