Mechanisms of Localized Corrosion Inhibition of AA2024 by Cerium

Article pubs.acs.org/JPCC

Mechanisms of Localized Corrosion Inhibition of AA2024 by Cerium Molybdate Nanowires Kiryl A. Yasakau,*,† Joaõ Tedim,† Maria F. Montemor,‡ Andrei N. Salak,† Mikhail L. Zheludkevich,† and Mário G. S. Ferreira† †

CICECO, Department of Materials and Ceramic Engineering, University of Aveiro, Campus Santiago, 3810-193, Aveiro, Portugal ICEMS, Department of Chemical Engineering, Instituto Superior Técnico, UTL, Lisbon, Portugal

‡

ABSTRACT: The cerium molybdate nanowires were recently reported as an efficient inhibiting pigment for aluminum alloys. In the present work, the inhibition mechanism of localized corrosion of S-phase intermetallics in AA2024 was studied in detail using a complementary combination of localized and analytical techniques. A significant suppression of dealloying of S-phase was demonstrated in the presence of cerium molybdate nanowires. Microscopic observations clearly show the formation of a conversion layer on the entire alloy surface after immersion in nanowire-containing solutions. A noticeable Volta potential difference (VPD) increase up to around −0.25 V vs Ni reference was measured on alloy after immersion in inhibited solutions. Such VPD changes have been related to the presence of Mo oxides on the alloy surface. Analysis performed by energy-dispersive spectroscopy (EDS) and X-ray photoelectron spectroscopy (XPS) showed that the surface oxide film is mainly composed by Mo(VI) and/or Mo(IV) oxides, and cerium(III) and cerium(IV), and aluminum oxides/hydroxides. A model galvanic couple made of aluminum and copper wires was used to simulate corrosion inhibition processes on S-phase intermetallics and alloy matrix. An enhanced inhibition efficiency of cerium molybdate was observed in electrolytes with higher concentration of sodium chloride. This was associated with the structural transformation of amorphous cerium molybdate nanowires into crystalline (NaCe)0.5MoO4 in concentrated NaCl solution, thereby triggering the release of cerium(III). This active feedback release can be used for development of “smart” self-healing coatings with inhibition triggered by the presence of corrosive salts in environment.

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INTRODUCTION Aluminum alloys such as AA2024 are widely used in aerospace industry for structural applications and as fuselage skins. The specific properties of Al alloys such as weight to strength ratio are valuable from the standpoint of mechanical resistance and cost efficiency.1 However, the highly heterogeneous microstructure obtained after specific heat-treatment procedures is also responsible for high corrosion susceptibility, especially in chloride containing environments. AA2024 may contain up to 5 wt % of copper as the major alloying element. The alloy microstructure is very heterogeneous due to the presence of other alloying elements and additives such as Fe, Si, Mn, etc. resulting in the formation of various intermetallic inclusions.2,3 One of the most abundant phases that occupy around 3% of surface area of alloy is Al− Cu−Mg with Al2MgCu nominal composition, called S-phase.3 Aluminum alloys are subjected to different types of corrosion such as intergranular corrosion, stress corrosion cracking, and pitting corrosion.4,5 Intermetallic particles such as S-phase have been previously identified as a main source of localized pitting attacks.6−10 Electrochemical measurements performed on Sphase showed that its corrosion potential in NaCl solution is more active compared to Al−4Cu solid solution.11 The S-phase © XXXX American Chemical Society

is prone to very fast dealloying, with Mg dissolved while copper stays inside the S-phase. The remaining copper acts as good cathode for oxygen reduction reaction promoting anodic dissolution of the surrounding aluminum matrix. Copper is redeposited along the alloy surface on later stages of corrosion, enlarging the effective cathodic area and increasing the corrosion attack.12,13 The localized corrosion activity of AA2024 can be effectively suppressed using corrosion inhibitors. Cr(VI) compounds have been known as effective corrosion inhibitors for different metallic substrates. However, due to high toxicity, the use of chromates is limited in industry as a result of imposing health regulations. Many works have demonstrated effective corrosion inhibition of aluminum alloys by alternative species including rare earth elements13−20 and specifically cerium-containing compounds,21,22 molybdates, vanadates,23−25 and organic species.7,26,27 Nonetheless, in many cases it is not possible to achieve the inhibiting level of chromates using only one inhibitor. One possibility to improve corrosion inhibition is via Received: December 18, 2012 Revised: February 19, 2013

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dx.doi.org/10.1021/jp3124633 | J. Phys. Chem. C XXXX, XXX, XXX−XXX

The Journal of Physical Chemistry C

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measurements were correlated with the chemical composition of the S-phase particles as determined by EDS and surface film composition as given by X-ray photoelectron spectroscopy.

combination of different corrosion inhibitors to achieve a synergistic inhibitive action.21,28−30 For example, Taylor and Chambers have demonstrated that combinations of rare earth cations and molybdates or silicates provide synergistic effect in corrosion inhibition of AA2024.28 However, the synergistic effect was not always found for a binary mixture at different mixing proportions. In some cases, certain mixing proportions resulted in antagonistic action which led to accelerated corrosion. Other groups reported positive inhibiting effect of combination of rare-earth anions with organophosphate anions such as dibutyl phosphate and diphenyl phosphate.21,29 It is interesting that the inhibiting effect of such compounds was attributed to the formation of rare earth metal−organic complexes with Al rather than simple cerium hydroxides precipitation.29 Recently, amorphous cerium molybdate nanowires were suggested as corrosion inhibitor for AA2024.31 The experimental results revealed higher inhibiting efficiency probably because of the combined action of cerium and molybdate species. However, the localized inhibiting nature of such compound on S-phase intermetallics in AA2024 has not been completely understood. Recent reports have shown very high potential of localized methods such as scanning Kelvin probe force microscopy (SKPFM) based on atomic force microscopy (AFM) coupled with Volta potential difference (VPD) mapping to study the corrosion mechanisms of different alloys and metals.8,9,32−37 Schmutz and Frankel showed a linear relationship of measured VPD of different metals in air and the corrosion potential in electrolyte.8,9 However, it has been debated that results obtained by SKPFM measurements cannot be used in a straightforward way, comparing to the conventional SKP. As has been mentioned by Rohwerder et al., the correlation between the SKPFM signal and the corrosion potential is not of general validity.38 The presence of surface oxide films or layers of the corrosion products may significantly affect the measured VPD on the intermetallic inclusions of aluminum alloy 2024.9,32,39 Cleaning of S-phase particles by ion sputtering lowers their Volta potential.32,33 Such changes of VPD can be explained in terms of surface adsorption, charge accumulation, changes in oxide structure,32 chemical composition, and the presence of hydroxyl content.40 Contamination on the surface of the tip,41 conditions of sample preparation which change the chemistry of the surface, tip, and cantilever geometry,42 roughness of the substrate,37 and configuration of electronic equipment43 can also be issues that influence Volta potential measurements. Nonetheless, SKPFM is a prospective and useful technique to study localized corrosion of alloys or their inhibition since the corrosion processes involve certain transformations of the metallic phases and formation of surface oxide films, which mostly influence the measured VPD. The results provided by this technique are essential for obtaining information on the localized corrosion activity and electrochemical nature of metals and intermetallic inclusions. The present work is devoted to understand the synergistic inhibition mechanism conferred by cerium molybdate. The amorphous cerium molybdate nanowires were dispersed in 0.05 and 0.5 M NaCl solutions and studied using X-ray diffraction (XRD) and scanning electron microscopy (SEM). AFM/ SKPFM measurements were performed to investigate the localized inhibition mechanism on polished AA2024 surface having well-defined S-phase particles before and after immersion in different NaCl solutions containing dispersed cerium molybdate powder. The scanning probe microscopy

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MATERIALS AND METHODS Materials and Samples Preparation. Aluminum alloy 2024-T3 with the following elemental composition representing wt % of alloying elements and additives was used in this work: Cu 3.8−4.9; Fe 0.5; Cr 0.1; Mg 1.2−1.8; Mn 0.3−0.9; Si 0.5; Ti 0.15; Zn 0.25; other 0.15; Al rest. The alloy specimens (dimensions 0.1 cm × 1 cm × 1 cm) were ground down to 4000 grit SiC sandpaper in water and finished with nonaqueous diamond slurries down to 2 μm followed by cleaning in acetone and 2-propanol under ultrasonic agitation. After cleaning, the samples were gently wiped by filter paper and stored in a desiccator. For galvanic corrosion tests, aluminum and copper wires with 1 mm diameter were embedded in epoxy resin. The prepared cell was abraded perpendicular to the wires arrangement from one side using SiC sand papers until grade 4000. A set of suspensions based on cerium molybdate (Table 1) was prepared. Cerium molybdate nanowires (CMN) were Table 1. Composition of Solutions Used in Corrosion Studies composition

reference name

0.05 or 0.5 M NaCl 0.7 g of CMN in 100 mL of 0.05 M NaCl 0.7 g of CMN in 100 mL of 0.5 M NaCl

0.05 M NaCl or 0.5 M NaCl solution A solution B

synthesized according to the procedure described elsewhere.31 Suspensions consisting of 0.7 g (dry weight) of cerium molybdate slurry (20 wt % in ethanol) in 100 mL of 0.05 or 0.5 M NaCl solution were prepared and sonicated for 10 min. These systems were used without further treatment after 1 day of preparation for AFM/SKPFM tests performed on AA2024. Corrosion Testing. Polished AA2024 samples were immersed in the solutions indicated in Table 1 for 2 h and 2 days in the case of 0.05 M NaCl based solutions and for 2 h and 1 day in the case of 0.5 M NaCl based solutions. A shorter immersion time has been chosen for 0.5 M NaCl solution because significant localized corrosion produced large accumulation of precipitates that negatively affect AFM/SKPFM imaging. Polished alloy samples have been investigated by AFM and SKPFM before and after immersion in corrosive solutions for different periods of time. After immersion, samples were gently rinsed with deionized water and dried in a desiccator before measurements. Galvanic corrosion measurements were performed using zero resistance ammetery. The galvanic current was measured between the two wire electrodes during immersion in pure 0.5 and 0.05 M NaCl solutions or solutions modified with CMN during approximately 2 days. Inhibited solutions were filtered prior to galvanic corrosion measurements. Characterization. X-ray Diffraction Analysis. X-ray diffraction measurements were performed with a Philips X’Pert MPD diffractometer (Ni-filtered Cu Kα radiation, tube power 40 kV, 50 mA, X’celerator detector, and the exposition corresponded to 2 s per step of 0.02). AFM/SKPFM. A commercial AFM Digital Instruments NanoScope III system with Extended Electronic Module was used to study the evolution of the topography of alloy surface B



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ex situ after immersion in different solutions. SKPFM was performed to study the evolution of Volta potential difference of the alloy surface after immersion in NaCl electrolytes. For SKPFM measurements the AFM was operating in the interleave mode with two pass scans. The first scan acquired topography of the surface, while during the second scan the tip was lifted up from the surface to a distance of 100 nm, the piezoelectric actuator was switched off, and an ac voltage of 1000 mV was applied between the tip and sample to cause electrostatically induced oscillation of the cantilever. Using the nulling technique, the Volta potential difference between the sample surface and the tip was measured over the alloy surface to obtain the Volta potential map. Values of the Volta potential were normalized with respect to the Volta potential measured on cleaned Ni substrate.8,9 For all SKPFM measurements, silicon probes covered with Cr−Pt layers were used. SEM/EDS. The microstructure and elemental composition of the metal surface was studied by Hitachi S-4100 SEM system with electron beam energy of 15−25 kV or Hitachi SU-70 with beam energy 15−20 kV combined with energy-dispersive spectroscopy analysis. XPS Measurements. The X-ray photoelectron spectroscopy measurements were performed using a Microlab 310 F (from Thermo Electron, former Vg Scientific). The spectra were taken in CAE mode (20 eV), using an Al (nonmonochromate) anode. The accelerating voltage was 15 kV. The quantitative XPS analysis was performed using the Avantage software. The spectra were corrected for the C 1s at 285.0 eV. The relative atomic concentration (Ax) was calculated using the following relation

Figure 1. XRD patterns (a) and SEM micrographs of cerium molybdate nanowires before (b) and after 1 day immersion in 0.05 M NaCl (c) and 0.5 M NaCl (d) solutions.

initial cerium molybdate nanowires indicates that these nanomaterials are amorphous, without any defined crystalline phase. However, the nanowires revealed structural changes after immersion in NaCl solutions, with a phase consisting of Ce2(MoO4)3·4.5H2O,44 formed after immersion in 0.05 M NaCl. Apparently, the NaCl solution induces reorganization of the amorphous phase into a more stable and crystalline form. The transformation is accompanied by changes of the precipitate color to a more yellowish grade. A microstructural image of these precipitates is presented in Figure 1c. The micrograph shows platelet-like particles and agglomerates of smaller round shaped particles with sizes in the range of 200− 300 nm. The XRD diffractogram after immersion in 0.5 M NaCl shows a well-defined crystalline phase corresponding to Na0.5Ce0.5(MoO4),45 different from the one formed in more diluted electrolyte. The suspension prepared in 0.5 M NaCl is not stable and quickly precipitates at the bottom of the beaker, forming a light yellow deposit. SEM image (Figure 1d) shows only agglomerates of particles with diameters around 200−300 nm as in the case of 0.05 M NaCl. In 0.5 M NaCl solution the Na cations were partially exchanged with Ce cations, forming a new crystalline phase. Therefore, cerium molybdate demonstrates cationic exchange properties that depend on the concentration of NaCl. According to inductively coupled plasma optical emission spectroscopy (ICP-OES) measurements, the concentration of cerium species found in 0.5 M NaCl modified with cerium molybdate is 451 ± 50 ppm unlike for 0.05 M NaCl modified with cerium molybdate, which showed a concentration of cerium species around 37 ± 11 ppm. The concentration of molybdate species was found to be 10 ± 4 ppm for 0.5 M NaCl and 0.05 M NaCl solutions. These electrolytes loaded with the amorphous cerium molybdate nanowires were used to study corrosion protection of the aluminum alloy and the results are described below. AFM/SKPFM Results. Localized Corrosion of AA2024 in NaCl Solutions. AA2024 surface contains various types of intermetallic inclusions composed of nobler elements than aluminum such as copper and iron, among others. It is known that the localized corrosion of S-phase, which is a copper-rich phase, is very strong in the presence of NaCl solution.2,5−13 Sphase dealloying plays an important role in pitting corrosion of AA2024. Thus, effective inhibitors must prevent S-phase

Ax = (normalized peak area)100 /(∑ normalized peak areas) i

where subscript x refers to quantified species and the subscript i refers to the other species detected in the XPS spectra. The normalized peak area was obtained by dividing the intensity of XPS peak of the species (after background subtraction) by the sensitivity factor of the corresponding species. The background subtraction was performed using the Shirley algorithm, which gives a S-shaped curve and assumes that the intensity of the background is proportional to the peak area on the higher kinetic energy side of the spectrum. The quantification was performed after peak fittings. The peak fitting function used was a Gaussian−Lorentzian product function, and the algorithm was based on the Simplex optimization as used in the Avantage software.

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RESULTS Characterization of Cerium Molybdate Nanowires. Previous work has shown the presence of cerium and molybdate species in NaCl solutions after addition of CMN powder.31 Therefore, the primary idea was to investigate possible structural changes of the nanowires after immersion in NaCl solutions aiming at unveiling the mechanism of the release of cerium and molybdate species from relatively insoluble compound. The cerium molybdate nanowires were immersed during 1 day in NaCl containing electrolytes and then analyzed by XRD (Figure 1a) and SEM and compared with nanowires before immersion, which are about 50−100 nm in diameter and 0.5−2 μm long (Figure 1b). The XRD pattern corresponding to the C



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Figure 2. AFM (a,c) and SKPFM (b,d) images of alloy surface before immersion (a,b) and after 2 h immersion in 0.05 M NaCl (c,d); topography and VPD profiles made across the intermetallic inclusion before and after 2 h are presented in panel e.

Figure 3. AFM (a,c) and SKPFM (b,d) images of alloy surface before (a,b) and after 2h immersion in 0.5 M NaCl (c,d); topography and VPD profiles (c) made across the intermetallic inclusion before and after 2 h and 1 day. Arrows indicate S-phase intermetallic zone.

of immersion, a deeper pit (about 1.5 μm) is formed at the intermetallic place (Figure 2e). VPD profile also shows a higher positive shift up to about −0.05 V vs Ni reference. These experiments demonstrate that the localized corrosion even in diluted NaCl solution is fast and destructive. The localized corrosion attack becomes even stronger when 0.5 M NaCl electrolyte is used. Figure 3a,b presents AFM and SKPFM maps acquired on alloy substrate at the place of Sphase intermetallic before immersion. Topography shows that significant localized corrosion appeared in the S-phase after 2 h of immersion in 0.5 M NaCl. A pit having depth around 300 nm and corrosion product deposits up to 1−1.5 μm high are clearly visible (Figure 3c,e). The VPD map presented in Figure 3d shows a white region corresponding to high Volta potential (around −200 mV vs Ni) in a pit zone. Such high potential corresponds to the place of the pit and to the alloy matrix within 5 μm surrounding the pit. Compared to results obtained in 0.05 M NaCl, the larger alloy surface has higher VPD level because of more intensive copper redeposition resulting from the enhanced localized corrosion activity in intermetallics. Aluminum matrix outside the pit has higher VPD level around −400 mV vs Ni reference. After 1 day of immersion the corrosion process has significantly progressed, inducing substantial dissolution of S-phase intermetallic up to 3 μm in depth and high increase of the VPD over the whole alloy surface (Figure 3e). The VPD level measured at the alloy surface after 1 day of immersion is around 0 V vs Ni. Such increase is consistent with the VPD measurements on pure copper after immersion in NaCl solution made by Frankel et al.9 They reported a VPD value measured on copper sample

dealloying significantly reducing the corrosion susceptibility of the alloy. In the present work, the localized corrosion behavior of such intermetallics was studied in NaCl solutions with and without amorphous cerium molybdate nanowires, which were previously reported to confer a good corrosion inhibition on AA2024.31 S-phase particles were identified on the polished alloy surface prior to corrosion tests and localized measurements. Topography, VPD maps, and corresponding profiles across the maps of the alloy surface prior to and after 2 h immersion in 0.05 M NaCl solution are presented in Figure 2. The topography scan shows a darker region in the middle of the map corresponding to S-phase intermetallic inclusion (Figure 2a). AFM image reveals that the intermetallic was slightly overpolished about 50 nm compared to the alloy matrix, which can be associated with the lower hardness of the intermetallic inclusion.35 VPD map of untreated surface shows higher potential (about 100 mV) in the place of intermetallic compared to the alloy matrix (Figure 2, b and e). Such increase can be correlated to electrochemically nobler S-phase (Al2CuMg) containing higher amount of nobler element, copper. After immersion in 0.05 M NaCl solution, drastic changes occur on the topography and VPD maps (Figure 2c,d). A localized corrosion attack caused partial dissolution of intermetallic (about 170 nm) in depth. Moreover, VPD above the intermetallic has been significantly shifted to a higher value around −200 mV vs Ni. It can be noted that VPD has also increased to about −500 mV vs Ni in the intermetallic surrounding zone, indicating modification of the alloy matrix after corrosion exposure to 0.05 M NaCl solution. After 2 days D



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Figure 4. AFM (a,c) and SKPFM (b,d) images of alloy surface before immersion (a,b) and after 2 h immersion in solution A electrolyte (c,d); topography and VPD profiles (e) made across the intermetallic inclusion before and after 2 h and 2 days. Arrows indicate S-phase intermetallic zone.

Figure 5. AFM (a,c) and SKPFM (b,d) images of alloy surface before (a,b) and after 2 h immersion in solution B electrolyte (c,d); topography and VPD profiles (e) made across the intermetallic inclusion before and after 2 h and 1 day. Arrows indicate S-phase intermetallic zone.

around −0.1 V, which is close to the values obtained in this work. Therefore, these results lead to conclusion that copper is being redeposited during the dealloying process of S-phase covering almost the entire surface as has been reported before.12,13,33 Localized Corrosion of AA2024 in the Presence of CMN. Localized corrosion of S-phase is quite different when cerium molybdate nanowires are added to the corrosive environment. Figure 4a,b demonstrates the topography and VPD maps of polished AA2024 before immersion indicating S-phase intermetallic. Several important changes of both topography and VPD maps can be observed after immersion in solution A (0.05 M NaCl + CMN) (Figure 4c,d). A topography map shows increase of the Z-range due to the presence of numerous precipitates formed on the surface. The precipitates having a height about 50−100 nm can be detected on both the alloy matrix and the intermetallic inclusion (Figure 4e). The VPD map also shows important changes in the Volta potential distribution across the surface. The VPD difference between the alloy matrix and intermetallic inclusion stays almost the same after 2 h of immersion (around 150 mV) (Figure 4e). Nevertheless, the relative position of VPD profile after 2 h of immersion is shifted to more positive values about −450 and −600 mV vs Ni for intermetallic and alloy matrix respectively as can be seen in Figure 4e. Such VPD increase on metal is most probably related to the formation of a layer of precipitates on the alloy surface. The longer immersion leads to formation of larger precipitates both on intermetallic and on alloy matrix. The precipitate’s height increases up to 150−200 nm after 2 days

compared to fresh alloy surface. The VPD level has been also increased until around −270 mV vs Ni metal. Moreover, after 2 days of immersion, the difference between the VPD values on intermetallic inclusion and the alloy matrix disappears in contrast to the fresh surface or when compared to the sample after only 2 h of immersion (Figure 4e). These results show that there is a correlation between the deposits on the alloy surface and increase of the measured VPD. Moreover, the results also do not show dissolution of S-phase as was found in the case of pure NaCl solutions (Figures 2 and 3). The AFM/SKPFM measurements presented in Figure 5 show evolutions of topography and VPD on the alloy surface and S-phase indicated by black arrows before (Figure 5a,b) and after immersion (Figure 5c,d) in solution B based on 0.5 M NaCl and inhibitor. After 2 h of immersion in electrolyte the topography map shows the formation of smaller precipitates (100−500 nm) on the alloy matrix and larger precipitates about 1 μm exactly at the places of intermetallics (Figure 5c,e). The VPD map shows that potential difference between the matrix and intermetallic place is decreased compared to the as polished conditions (Figure 5b,d). Moreover, the VPD profile has a higher potential across the deposits zone being around −340 mV vs Ni and the level at the surrounding matrix being lower about −400 mV vs Ni (Figure 5e). A longer immersion time causes the formation of large precipitates more than 4 μm atop the intermetallic zone (Figure 5e). The VPD level on large precipitates is slightly shifted to negative values around −340 mV vs Ni which could be caused by the topography effect hampering SKPFM measurements.37 The VPD across the alloy matrix is around −250 mV vs Ni having similar values obtained E



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Figure 6. SEM micrographs corresponding to the places with S-phase studied by AFM analysis (Figures 2 and 4) of alloy surface after 2 days immersion in pure 0.05 M NaCl (a) and solution A electrolyte indicating S-phase (b) and surrounding alloy matrix (c); EDS spectra (d) taken locally at the S-phase intermetallic place (a,b) and on the surrounding alloy matrix (c).

Figure 7. SEM micrographs corresponding to the places with S-phase studied by AFM analysis (Figures 3 and 5) of alloy surface after 1 day immersion in 0.5 M NaCl (a) and solution B electrolyte (b); EDS spectra (c) taken at the S-phase intermetallic place (a, b) and on the surrounding alloy matrix (away from S-phase).

and VPD values of both alloy and intermetallic grow until approximately −250 mV vs Ni. In the case of pure NaCl solutions, VPD levels obtained on the alloys surface after 2 days and 1 day immersion in 0.05 and 0.5 M NaCl solutions respectively are close to 0 V vs Ni reference. The surface of samples after corrosion has been further evaluated by SEM/ EDS techniques in order to clarify the composition and microstructure of the deposit layers. The results are presented below. SEM/EDS Results. Microstructural analysis was performed in order to understand better the results obtained by AFM/ SKPFM and to measure the elemental composition of the

on sample surface after measurements in solution A electrolyte (−270 mV vs Ni) (Figure 4e). Although local topography changes are more pronounced in more concentrated corrosive electrolytes such as solution B compared to solution A, VPD levels measured on different samples after corrosion are very similar in both cases. Such behavior may be associated with similar inhibiting mechanism in cerium molybdate based electrolytes. The experimental findings presented above demonstrate that the VPD value for the alloy surface is related to the coverage of surface by a layer of precipitates formed in CMN-containing solutions. The increase of immersion time apparently produces a thicker film F



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large peaks of Mo and smaller peaks of Ce are detected on Sphase after immersion in solution B (Figure 7c). Spectrum acquired on alloy matrix demonstrates mostly Mo signal and signals associated with elemental composition of the alloy matrix (mainly Cu, Mg, Al). It should be noted that the Mo signal on alloy matrix is much more intense compared to spectrum acquired in solution A electrolyte (Figure 6d). Therefore, it can explain the presence of a truly conversion layer at the alloy surface (Figure 7b). The results of SEM/EDS analysis are in accordance with the AFM observations presented above. Higher amounts of deposits on intermetallics compared to alloy matrix were identified on the sample after immersion in inhibiting solutions using both techniques. The deposited film acts as a protective barrier and inhibits local dissolution of S-phase intermetallics. XPS Results. The chemical composition of the surface after immersion in different corrosive solutions was investigated by XPS analysis. The elements found on the alloy surface after immersion in pure NaCl solution are Al and Cu and small quantities of Cl (not shown) (Table 2). The XPS spectra for Al

precipitates and underlying intermetallic inclusions. Figure 6 and 7 present SEM micrographs of S-phase zone after immersion in pure 0.05 M NaCl solution and in solution A (Figure 6) and after 1 day immersion in 0.5 M NaCl and solution B electrolytes (Figure 7). The micrographs were taken directly on the surface that was studied using AFM/SKPFM techniques (Figures 2−5). A significant degradation of intermetallic precipitates occurs after immersion in pure NaCl solution. The cracks and crevices formed around the S-phase can be seen after localized corrosion attack and dissolution of more active elements after immersion in 0.05 and 0.5 M NaCl solutions (Figures 6a and 7a). Immersion in 0.5 M NaCl causes higher localized attack on S-phase compared to 0.05 M NaCl, though the difference is not that evident when comparing both SEM pictures (Figures 6a and 7a). AFM results correlate well with the previous observation demonstrating an extensive dissolution of S-phase up to 2.5 μm in 0.5 M NaCl compared to 1.5 μm in 0.05 M NaCl which is more clearly visible in the AFM profile images across the S-phase intermetallics (Figures 2e and 3e). EDS spectra have been acquired on intermetallics after immersion in NaCl solution in order to evaluate changes in elemental composition. Figures 6d and 7c demonstrate that magnesium signal disappeared and the copper signal is significantly increased compared to Al being in line with typical corrosion behavior of S-phase and confirming the preferential dealloying of more active elements. Higher copper content has been measured on intermetallic after immersion in 0.5 M NaCl solution which can be attributed to faster intermetallic dealloying. After immersion in solution A, the surface of S-phase and surrounding alloy matrix has a layer of round-shaped deposits having dimensions about 100−400 nm, which is clearly seen in both micrographs (Figure 6b,c). The layer of deposits on Sphase is somewhat more compact compared to that on the alloy surface (Figure 6b,c). From another hand, a thicker layer is formed in the electrolyte B. S-phase places are covered by solid white precipitates and the surrounding alloy matrix is covered by round-shaped deposits after immersion in solution B (Figure 7b). The alloy matrix seems to be completely covered by a thicker conversion layer as can be clearly seen in Figure 7b, which shows a crack between the white deposit and the covered surrounding matrix. Such crack could be formed after shrinking of deposits caused by heating of sample during preparation for SEM analysis or heating under electron beam during SEM observation. In order to verify the nature of the deposits, EDS analysis was performed in both the S-phase intermetallics and alloy matrix (Figures 6d and 7c). EDS spectrum of inhibited sample clearly shows well-defined peaks for Mg in S-phase intermetallics (Figure 6d). The calculated relative proportions of elements from EDS spectrum Al:Mg:Cu (2.2:1:1) are close to nominal composition of S-phase.3 This suggests that intermetallics do not suffer localized attack after 2 days of immersion in solution A, clearly demonstrating an inhibiting effect. The results show that, apart from the elements present in the alloy and S-phase such as Al, Cu, and Mg, signals of molybdenum and oxygen are also visible on the spectra taken at S-phase intermetallic. The cerium signal is weak and not quite evident. These results suggest that alloy surface after 2 days immersion is covered mainly by a layer consisting of molybdenum and oxygen species. EDS analysis performed on the alloy matrix exposed to inhibited electrolyte shows very weak signals coming from Mo, which are almost not visible in the spectrum. In contrast to that,

Table 2. Element Concentrations Calculated from the XPS Spectra element concentrations (atom %) element

0.05 M NaCl

Cl Cu Al Mo

1.0 6.7 33.9

Ce

solution A

solution B

25.78 Mo(IV) 6.6 Mo(VI) 5.4 Ce(III) 2.7 Ce(IV) 3.4

20.39 Mo(IV) 4.4 Ce(III) 1.0 Ce(IV) 2.7

and Cu are presented in Figure 8. The spectra of Al 2p are represented mainly by a small peak at 73 eV46 and larger peak at 74.6 eV47 that can be assigned to the presence of elemental aluminum (lower binding energy (BE)) and to the presence of Al−O bonds in the form of either aluminum oxides or hydroxides on the sample surface (higher BE), respectively (Figure 8a). The existence of a peak coming from pure aluminum is not expected on the highly corroded alloy surface covered by corrosion precipitates. In the present study, the alloy surface did not contain a large amount of corrosion products as follows from the SEM micrographs (Figures 6a and 7a). One of the plausible explanations is that the alloy surface has places which are not covered by corrosion products and places where native oxide film is thin enough for XPS analysis to detect Al. The Cu 2p envelope for the sample exposed to blank NaCl reveals the contribution of different forms of copper. Although the shakeup satellites above 938.0 eV and characteristic of Cu(II) are not shown (Figure 8b), Cu(0), Cu(I), and Cu(II) species cannot be discarded at the surface.48 Cl signal has also been identified in the XPS spectrum (not shown). Therefore, the existence of Cu(I) can be related with Cu(I)−Cl complexes and Cu(I)−O species. The presence of metallic copper can be expected as a part of the S-phase dealloying and copper redeposition process.39 In the case of alloy sample immersed in 0.5 M NaCl solution, the elemental composition of surface is very similar, demonstrating high copper coverage and distribution of aluminum oxides (not presented). G



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Figure 8. XPS spectra representing Al 2p (a) and Cu 2p (b) ionizations obtained on the AA2024 after immersion in 0.05 M NaCl solution.

compounds which precipitate during corrosion activity compared to 0.05 M NaCl (Table 2). On the other hand, signals of Cu and Cl have not been found on the surface, supporting that copper-containing intermetallics remain intact during corrosion of alloy. Figure 9 shows XPS spectra of Mo 3d and Ce 3d for alloy after exposure to cerium molybdate based electrolytes. For the sample exposed to solution A during 2 days, the Mo 3d spectrum is complex and includes two chemical states of Mo having Mo 3d5/2 and Mo 3d3/2 overlapped ionizations (Figure 9a). The Mo 3d5/2 ionization can be deconvoluted into two main contributions: Mo(IV) (lower binding energy) and Mo(VI) (higher binding energy).49 However, the presence of cerium molybdate species cannot be excluded. Generally, these species overlap with the Mo(VI) peaks. Surface oxide film shows atomic concentrations of Mo(IV) and Mo(VI) species around 6.6 and 5.4 atom %, respectively. Unlike the alloy exposed to solution A, the Mo 3d ionization spectrum of alloy surface after immersion in solution B shows only one oxidation state Mo(IV) (Figure 9b) with atomic concentration about 4.4 (Table 2). Figure 9c,d depicts Ce 3d ionization spectra for AA2024 samples exposed to inhibiting electrolytes. It seems that Ce 3d peak at low BE is composed of two peaks that can be ascribed to Ce(III) and Ce(IV) oxidation states present as oxides and/or hydroxides.29 The satellite peak at high BE is a signature of the Ce(IV) oxidation state and is clearly visible on both spectra. The results also show that concentration of Ce(IV) at the surface is higher compared to that of Ce(III).

The Al 2p ionization spectra for the samples after immersion in the inhibiting solutions demonstrate that the BE of Al 2p peaks correspond to oxides/hydroxides (not shown). However, surface oxide film shows lower concentration of Al-rich

DISCUSSION As expected from previous works,2,3,6−12,35 the results of corrosion tests in NaCl-based solutions show that localized corrosion of AA2024 occurs at high rates and S-phase intermetallics are the most active sites, with dealloying of this

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Figure 9. XPS spectra representing Mo 3d (a,b) and Ce 3d (c,d) ionizations obtained on the AA2024 after immersion in solution A (a,c) and solution B (b,d). H



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of the alloy surface was found to change after immersion in nanowire-containing electrolytes. An increase of VPD with respect to alloys exposed to CMN-free electrolytes has been found for all samples after immersion in cerium molybdate modified solutions shown in Figures 4 and 5. The VPD value around −0.25 V vs Ni reference is found for samples with the thickest surface oxide film. In contrast in the case of actively corroding AA2024 the VPD increases only locally close to the active pitting. This increase is associated with the copper redeposition around the pits increasing nobility of the surface.13,50 However, such an explanation is not valid for the inhibited solutions because of absence of copper redeposition as proved by XPS. Thus, the VPD increase observed in the entire alloy surface is most probably related with the growth of a protective layer with different electronic properties. The deposition of Ce hydroxides/oxides cannot make the alloy surface significantly nobler as reported in previous works.13,19 Therefore, the most probable reason of Volta potential increase is formation of Mo oxide based surface layers which is consistent with results from SEM and AFM. It is well-known that the Volta potential difference measured in SKPFM mode in air is related to the difference in the work function (WF) of the probe and analyzed surface. The reported WF values of different Mo oxides having Mo oxidation state +4, +5, and +6 lie in the range of 5−6.7 eV with the lower values attributed to Mo(IV) species.51−54 The WF of molybdenum oxides is significantly higher than that of metallic aluminum (around 4.2 eV).55 This can partially explain the observed increase of Volta potential of the whole alloy surface after immersion in the electrolyte with nanowires, but the increase of VPD is less than expected for a surface covered with a dense layer of pure Mo oxides and can be related to coprecipitation of complex oxide films of Mo, Al, and Ce oxides/hydroxides in the presence of CMN. Indeed, the performed XPS measurements support this assumption, revealing different types of Al, Mo, and Ce species on the surface. Therefore, the existence of Al, Ce(III,IV) oxides and hydroxides and Mo oxides apparently influences the electronic properties of the protecting oxide film. A similar case has been reported by Liang et al.56 Accordingly, a mixed oxide system containing MoO2 and SiO2 oxides has been prepared by coevaporation in order to tune electronic properties of MoO2 oxide. The addition of SiO2 has resulted in gradual decrease of the work function of a mixed system from 6.5 eV to about 5.5 eV where maximal content of SiO2 was 33%. The important point for corrosion inhibition in the present work is that the oxide film increases the nobility of the Al alloy surface. Moreover, the potential difference between different phases present in alloy also diminishes significantly, reducing the driving force for localized corrosion attack. This leads to efficient inhibiting performance. In spite of the obvious inhibiting effect of the formed oxide film, the particular roles of molybdate and cerium ions in the inhibiting process are not straightforward. EDS spectra suggest that the oxide film on alloy surface is mostly composed of Mocontaining species having a small intensity of Ce (Figures 6d and 7c). A surface analysis performed by XPS also demonstrates higher intensity of Mo 3d ionization when compared to Ce 3d (Figure 9). Then, additional electrochemical model experiments have been performed in order to clarify the specific roles of molybdate and cerium ions. Galvanic corrosion measurements have been performed in a cell consisting of separate aluminum and copper wire electrodes embedded in an epoxy resin. Such a cell was chosen as a model

type of intermetallics and copper refining/redeposition playing an essential role during pitting initiation and stabilization. In contrast, the corrosion processes in the presence of cerium molybdate nanowires are significantly suppressed. A deeper analysis based on information obtained and presented in previous sections is performed to clarify the mechanism of inhibition associated with CMN. The dealloying/redeposition processes in the S-phase intermetallic inclusions are suppressed by the formation of an oxide-based layer atop of both the alloy matrix and the intermetallics as demonstrated by SEM (Figures 6 and 7). However, it would be essential to reveal if there are any changes of intermetallics underneath the oxide-based deposits such as those presented in Figure 7b. Figure 10a shows SEM

Figure 10. SEM image of alloy surface after 1 day immersion in solution B and EDS spectra taken at the S-phase intermetallic (#1) shown in picture (a) and on intact S-phase (#2) after alloy polishing (not shown).

micrograph of the alloy surface with intermetallic which does not have the oxide dome deposit after 1 day immersion in solution B. The deposit was removed by rinsing the sample after immersion. The important point is that the intermetallic did not suffer localized corrosion and the respective EDS spectrum clearly shows the presence of Al, Cu, and Mg elements (Figure 10b). The EDS analysis demonstrates that relative concentration of Mg and Al (elemental ratio Al:Cu:Mg is 2.4:1:1) in this particle is barely changed when compared to the initial composition before immersion tests (elemental ratio Al:Cu:Mg is 2:1:1). The obtained proportions are very close to the ratio calculated from EDS spectrum on S-phase intermetallic after immersion in solution A being around 2.3:1:1 (Figure 6d). The redeposition of Cu on the alloy surface is also not observed by XPS even after long immersion tests. Additionally, the growth of deposits during immersion in inhibited solutions is observed by AFM (Figures 4 and 5). It can also be considered as an important factor contributing for the damping of corrosion processes kinetics. The performed SKPFM study contributes to a better understanding of the electrochemical nature of the formed layer. The Volta potential I



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with the XRD findings indicating an exchange of Ce by Na cations occurring with nanowires immersed in 0.5 M NaCl (Figure 1a). The released cerium cations in solution apparently confer an efficient corrosion protection to copper electrode forming oxide/hydroxide precipitate on the surface (Figure 11b). In the case of solution A, the concentration of cerium species available in electrolyte is low, and as a consequence the efficiency of copper coverage and inhibition of cathodic process is less pronounced. The results discussed above contribute to clarification of a possible inhibiting mechanism of cerium molybdate nanowires, which has been partially demonstrated in the previous work.31 The cathodic inhibiting properties of cerium ions on copper surface are related to a well-described process of insoluble hydroxide formation on active copper- and iron-containing intermetallics.13−19,57 Intermetallics support oxygen reduction reaction producing hydroxyl ions and thus increasing a pH in the cathodic zone. The formation of cerium(III) hydroxide precipitates occurs in such places hindering the cathodic activity. The mechanism of oxygen reduction on coppercontaining intermetallics may involve formation of peroxide species which can oxidize Ce(III) cations or oxy/hydroxides to Ce(IV) followed by the precipitation of Ce(IV) oxy/ hydroxides13,57,58 with very low pKsp (47.7).59 In addition, the dissolved oxygen and the pH of the solution may influence Ce oxidation state as has been shown by Davenport et al.60 XPS results presented here prove the fact that cerium in both valent states is present in the deposits. The ratio between Ce(IV) and Ce(III) is higher in more concentrated electrolyte (solution B), suggesting that electrochemical mechanism is probably involved. A clear deposition of Ce oxide/hydroxides has been found on copper due to high cathodic activity of Cu electrode in galvanic coupling with Al (Figure 11). On the other hand, it has been shown that there is no dealloying of S-phase intermetallics in inhibited cerium molybdate solutions (Figure 10). Therefore, cathodic oxygen reduction process during initial stages should be weak and the formation of large amounts of Ce precipitates on the alloy because of cathodic processes on S-phase is not expected. Nevertheless, the deposition of Ce oxides/hydroxides is possible in most active spots considering corrosion susceptibility of the alloy. In contrast to cerium, molybdenum signals have been identified on both aluminum and copper wires surface (Figure 11b). Such behavior may be attributed to several reactions pathways that lead to precipitation of molybdenum oxides.61 A simple reduction of molybdate is possible via reaction 1:

for localized galvanic corrosion in AA2024 in the later stages after dealloying of S-phase and refining of copper cathodes. The cell was immersed in different corrosive electrolytes containing 0.05 or 0.5 M NaCl and a current passing between the two electrodes was measured during 2 days using a zero resistance ammeter (ZRA). Figure 11a presents the coupling current

Figure 11. Evolution of galvanic current (a) measured using ZRA flowing through aluminum and copper wires in pure 0.05 and 0.5 M NaCl and in solution A and solution B electrolytes; EDS analysis of Al and Cu electrodes (b) after immersion in solution A and solution B electrolytes.

evolution with time. The values are higher at the beginning, which is possibly related to instability of the system before it reaches steady-state current. After several hours, the current stabilizes in all the cases. The highest current is observed for the blank concentrated solution, as expected. The inhibition effect of nanowires is already well visible when 0.05 M NaCl inhibited solution is used. However, the difference is significantly higher in more concentrated inhibited corrosive electrolyte. Surprisingly, the galvanic corrosion current density in 0.5 M NaCl electrolyte prepared using CMN is about 7 times lower than that in the less aggressive solution (0.05 M NaCl). Thus, an active triggering of inhibition by concentration of corrosive electrolyte can be assumed. This effect is explained in terms of enhanced release of Ce ions from CMN in the presence of 0.5 M NaCl, as confirmed by ICP analyses (see previous sections). The surface composition on Al and Cu electrodes was analyzed by EDS after immersion and the results are presented in Figure 11b. The aluminum electrode shows well-defined Mo and O signals and no signals of Ce after immersion in both inhibiting solutions (Figure 11b). This observation correlates well with the anodic nature of the electrochemical processes on aluminum which does not create conditions for deposition of cerium hydroxide species. The copper electrode surface has strong peaks of Ce and Mo and O peaks after being immersed in concentrated solution B. On the contrary, only weak signals of Ce and O are present on copper electrode after immersion in solution A. A high inhibiting efficiency of solution B is correlated with the additional formation of cerium precipitates on the copper electrode surface. This corroborates very well

MoO4 2 − + 4H+ + 2e− → MoO2 + 2H 2O

(1)

On the other hand, the reduction process as was suggested in ref 62 may also occur on expenses of oxidation of aluminum: 3MoO4 2 − + 6H+ + 2Al → 3MoO2 + Al 2O3 ·3H 2O

(2)

In these ways, molybdenum oxides can be formed on copper or aluminum surface. However, the corrosion process in the alloy substrate is more complex. XPS measurements indicated the presence of Mo(IV) and Mo(VI) on the surface in case of solution A (Table 2), which can be caused by several intermediate processes that include transformation of molybdate anions to Mo(III) state and then formation of MoO2 or MoO3 oxides according to reactions 3 and 4:61 J



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Mo3 + + 2H 2O → MoO2 + 4H+ + e−

(3)

Mo3 + + 3H 2O → MoO3 + 6H+ + 3e−

(4)

Localized corrosion inhibition of AA2024 has been investigated in NaCl solutions modified by cerium molybdate nanowires using a combination of localized AFM/SKPFM analysis supported by a microstructural SEM/EDS and surface XPS analysis. It has been found that S-phase intermetallics stay stable in both inhibited solutions unlike for pure NaCl solution where a significant dealloying of S-phase intermetallics occurs. SKPFM measurements revealed a noticeable VPD increase (about −0.05 V vs Ni reference) on the alloy surface after immersion in corrosive solutions. The VPD increase after immersion in pure NaCl solutions has been attributed to the copper redeposition on the entire alloy surface. On the other hand, VPD values (around −0.25 V vs Ni reference) were measured on alloy after immersion in both inhibited solutions (0.05 and 0.5 M NaCl). Such VPD changes have been related to the presence of Mo oxides on the alloy surface. Electrochemical experiments performed on Al−Cu galvanic couple indicate lower galvanic currents in 0.5 M NaCl unlike for 0.05 M NaCl solution based on cerium molybdate. Such difference is attributed to higher concentration of Ce species, which are released in 0.5 M NaCl solution during the modification of crystalline structure of cerium molybdate as denoted by ICP-OES and XRD. In NaCl solutions based on cerium molybdate nanowires, the molybdate plays an important role suppressing the active dissolution of Al and Mg from S-phase decelerating dealloying and copper redeposition. The cerium cations have major effect on the cathodic processes occurring on the surface of partially dealloyed intermetallics at the later stages of corrosion.

XPS results performed on alloy surface after immersion in solution B showed only one oxidation state of Mo(IV) (Table 2). The existence of both oxidation states on the alloy surface will be governed by different factors such as pH, oxygen concentration, concentration of reactive species, etc. The potentials of half-reactions 3 and 4 are around −1.23 V and −0.64 V vs SCE, respectively, assuming pH 7 and the concentration of Mo3+ to be 10−6 M. Then it becomes clear that Mo(VI) is more stable at more positive potentials. The value of open circuit potential measured on AA2024 in 0.5 M NaCl electrolyte containing cerium molybdate (−0.7 V vs SCE) is lower than that in 0.05 M NaCl with cerium molybdate (−0.56 V vs SCE). Moshier et al. have stated from their XPS analysis results that the Mo4+ signal obtained at OCP around −1 V vs SCE was high, whereas when the electrode was polarized until breakdown potential (more than −0.6 V (SCE)) Mo6+ intensity significantly increased.63 It may be suggested that in the case of lower alloy potential Mo(IV) state is more thermodynamically stable compared to Mo(VI), which explains the presence of only Mo(IV) on the surface. The difference in the corrosion potential in electrolytes with different concentrations can also be responsible for the change of the ratio between Ce(III) and Ce(IV) species. The performed analysis highlights the importance of having both cerium and molybdate species available in solution, prepared using CMN pigment, for efficient inhibition of localized corrosion processes on AA2024. Also, it can be deduced that cerium and molybdate ions play different roles at different stages of the corrosion process. The initial stage of corrosion on AA2024 surface involves dealloying of Al and Mg active elements from copper-rich intermetallics. The copperenriched remnants become active cathodes in the oxygen reduction process that support anodic dissolution of the surrounding alloy matrix. In addition to the local corrosion process, copper redeposition follows the dealloying process, and the spreading of copper on the surface can increase the effective cathodic area and accelerate further the corrosion of alloy. The molybdate plays an important role at the initial stage suppressing the active dissolution of Al and Mg from S-phase, thereby decelerating dealloying and redeposition of copper. The deposits formed during initial immersion of alloy in inhibited solution support this statement. The role of cerium in the inhibition process is different compared to molybdate. The cerium cations have major effect on the cathodic processes occurring on the surface of partially dealloyed intermetallics only at later stages of corrosion. When the oxygen reduction occurring on copper sites is strong enough, the precipitation of cerium oxides/hydroxides occurs, as demonstrated in the galvanic corrosion experiments using pure aluminum and copper wires.

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AUTHOR INFORMATION

Corresponding Author

*Tel.: 351-234378146. Fax: 351-234378146. E-mail: [email protected]. Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS K.A.Y. thanks for Post-Doctoral grant (ref SFRH/BPD/80754/ 2011), J.T. thanks FCT for Post-Doctoral grant (ref SFRH/ BPD/64335/2009). The authors acknowledge Aleksey Lisenkov for the SEM analysis and Dr. Alexandre Bastos for helpful discussion. MULTIPROTECT (ref NMP3-CT-2005-011783) and FCT (ref PTDC/CTM/65632/2006) projects are acknowledged as well for financial support and INM (Germany) in particular for supply of nanowires.

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REFERENCES

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CONCLUSIONS Amorphous cerium molybdate nanowires were used as inhibiting pigment for 0.05 and 0.5 M NaCl electrolytes. It has been shown that the amorphous CMN transform into crystalline phases Ce2(MoO4)3·4.5H2O and (NaCe)0.5(MoO4) during immersion, respectively, in 0.05 and 0.5 M NaCl electrolytes, releasing cerium and molybdate species in solution. K



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Mechanisms of Localized Corrosion Inhibition of AA2024 by Cerium

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