# Intermetallic particles distribution in 7010-T7451 and 7175-T6 aluminium alloys substrate impacts cathodic stability of silicate sealed anodic oxide film

Florian Raffin, Maxime Magnan, Antoine Seyeux, Sandrine Zanna, Jacques Echouard, Jolanta Światowska, Polina Volovitch

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Florian Raffin, Maxime Magnan, Antoine Seyeux, Sandrine Zanna, Jacques Echouard, et al. In-termetallic particles distribution in 7010-T7451 and 7175-T6 aluminium alloys substrate impacts cathodic stability of silicate sealed anodic oxide film. Corrosion Science, 2024, 237, pp.112270. 10.1016/j.corsci.2024.112270. hal-04799089

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Intermetallic particles distribution in AA7010-T7451 and AA7175-T6
aluminium alloys substrate impacts cathodic stability of silicate sealed
anodized oxide film

Florian Raffinª, Maxime Magnanª, Antoine Seyeuxª, Sandrine Zannaª, Jacques Echouardᵇ,
Jolanta Światowskaª, Polina Volovitchª,\*

ªChimie ParisTech-CNRS, PSL Research University, Institut de Recherche de Chimie Paris (IRCP), 11 rue Pierre et Marie Curie, 75005 Paris, France

ᵇSafran Landing Systems, 7 Rue Général Valérie André, 78140 Vélizy-Villacoublay, France

*Corresponding author: [polina.volovitch@chimieparistech.psl.eu](mailto:polina.volovitch@chimieparistech.psl.eu)

Abstract

Effect of silicate sealing on tartaric sulphuric acid anodized AA7010-T7451 and AA7175-T6 alloys with different distributions of intermetallic particles (IMPs) in aluminum substrates was studied. Surface characterisations evidenced similar morphology and chemical composition of the sealed layers on both substrates. Electrochemical and time-lapse microscopy results demonstrated the layer's stability under anodic polarization and its localized pitting under cathodic polarization in NaCl electrolyte. No chemical or morphological modifications were detected outside of the pits. The number and size of the pits depended on the distribution of IMPs in the alloy, buried under 6 µm thick oxide. A degradation mechanism explaining these observations was proposed.

Keywords: A. aluminium alloys; B. EIS; B. GD-OES; B. XPS; B. SEM; C. buried interfaces.

Highlights

→ Procedure for intelligent ranking of sealants for anodized Al alloys is proposed.

→ In studied sealant/substrate systems, sealing improves immersion stability.

→ All studied systems show similar degradation mode (localized pitting).

→ Pits distribution depends on the alloy process (machining/rolling) for all sealants.

→ Pits size/number correlates with intermetallics distribution in the substrate alloy.
# 1. Introduction

In the aerospace industry, the 7xxx series aluminium alloys (enriched in Zn, Mg, Cu and Fe) are of particular interest because of their excellent mechanical strength-to-density ratio [1,2]. The presence of intermetallic particles (IMPs), such as Al₂Cu, Al₂CuMg or Al₇Cu₂Fe etc., makes however these alloys susceptible to corrosion. Surface treatments are therefore essential to enhance the corrosion resistance for their applications. One of the most popular treatments for surface protection of aluminium alloy is anodizing [3], which leads to the formation of a two-layered oxide structure on the aluminium alloy substrate: a compact thin (tens of nanometres) inner oxide covered by a porous thick (ranging from several to several tens micrometres) outer oxide layer. This porous layer being sensitive to the penetration of aggressive species, the pores must be sealed. This is achieved by dipping the sample in a sealing bath, leading to pore sealing and a further increase in the corrosion resistance of the oxide layer [3]. Hitherto, both the sealing and the anodizing processes have used a solution based on hexavalent chromium (Cr(VI)), known for its effectiveness due to its self-healing properties [4,5]. However, Cr(VI) is highly toxic and will be banned from industrial applications by 2024 according to the REACH regulation [6,7]. As a Cr(VI)-free alternative for anodizing, tartaric sulphuric acid (TSA) has been recognized as a promising solution [8–11]. Several Cr(VI)-free sealants have been proposed in the literature, such as common hot water sealing (HWS) or cerium salt bath sealing [11–17]. Among these alternatives, silicate sealants are notable for the capacity of silicate anion to be adsorbed inside the aluminium pores [17–20]. The microstructure and the corrosion resistance of these newly sealed layers formed with emerging solutions require further characterization. Many studies have focused on the influence of the sealing on the overall corrosion rate in electrochemical or accelerated climatic tests [11, 12, 14, 16], as well as the mechanical properties of anodized and sealed layers. Some studies examined the effect of machining/rolling on the formation and properties of the anodic film formed by anodizing [21,22]. However, to the best of our knowledge, no study has specifically been devoted to the effect of the alloy processing (machining, rolling etc.) on the efficiency of sealing treatments and/or the stability of the sealed oxide, present in the external part of the porous layer and separated by several micrometres from the metallic substrate. Although the alloy processing does alter the distribution of intermetallic particles (IMPs) near the surface [23,24] and copper-containing IMPs in 7xxx series alloys may contribute to the
decreased corrosion resistance of unprotected alloy [25,26], the effect of these buried IMPs under the oxide on the microstructure, chemical composition, and stability of the top sealed layers has never been considered.

In this study, the primary objective is to determine, for the first time, whether variations in the distribution of intermetallic particles in the alloy substrate, buried under several micrometers of thick anodized oxide, affect the chemical composition, microstructure, and aqueous stability of silicate-sealed anodized oxide.

To achieve this, the microstructures and the immersion stability of TSA-anodized and sealed layers obtained with two different concentrations of alkaline silicate-based sealings using machined 7010-T7451 and rolled 7175-T6 alloys are compared. Different sealing procedures and different concentrations of silicates were tested to investigate whether potential differences in stability evolution between substrates are really dependent on the substrate alloy.

To elaborate a stability test and correctly rank the substrates, different degradation mechanisms must be considered and the procedure, which accelerates the selected degradation mechanism, needs to be chosen. Specifically, for anodised and sealed aluminium degradation may be intensified by cathodic dissolution of the oxide [27] or by anodic degradation resulting from pitting caused by $Cl^-$ adsorption [28]. Both mechanisms can occur during open circuit immersion in a concentrated chloride-rich solution. Cathodic enhancement of degradation processes could occur during the cathodic delamination of organic paints from silicate-sealed anodized Al alloys or in the case of galvanic coupling between anodized and sealed alloy with a more active material, such as, for example, galvanized steel. The anodic mechanism is expected on the open surface where $Cl^-$ adsorption is highly probable. This study primarily focuses on the cathodic initiation of the degradation and investigates whether, in the case of anodized silicate-sealed layer, degradation is enhanced by the weakening of the sealed oxide or by the degradation of the metallic matrix-oxide interface. This phenomenon is usually observed during delamination under "weak" paints that offer relatively low protection due to the relatively high diffusion of water, oxygen and other aggressive species [29–31]. In this mechanism, oxygen (equation 1a) and water (equation 1b) are reduced to form hydroxide ions (and, in some cases, hydrogen gas). In an aerated alkali medium, up to circa -1 V vs SCE, the reaction is driven by oxygen reduction (1a),
but below -1 V vs SCE, the reaction is limited by the transport of oxygen (1b) [32]. In all cases, the intensive cathodic reaction lead to the increased basicity of the solution [33] and subsequent dissolution of the anodized layer, as can be schematically presented by reaction (2). Similarly, a silicate-based layer could be dissolved at very high pH [34]:

$$2H_2O + O_2 + 4e^- \rightarrow 4OH^- \quad (1a)$$

$$2H_2O + 2e^- \rightarrow 2OH^- + H_2 \quad (1b)$$

$$Al(OH) + OH^- + H_2O \rightarrow Al(OH)_4^- \quad (2)$$

To accelerate the cathodic degradation mechanisms of the sealed oxide layer and evaluate the stability of the samples, procedures similar to the AC-DC-AC test used for the evaluation of paints and/or conversion coatings [27,35,36] could be adapted with combination of electrochemical and time lapse microscopy measurements. Mechanistic understanding can be achieved by complementary surface characterisations.

This study applies AC-DC-AC based ageing and characterisation procedure in combination with advanced chemical and morphological surface analysis at different stages of degradation in order to investigate whether the distribution of IMPs resulting from the alloy processing can affect the cathodic stability of silicate-based sealant on the TSA-anodized aluminum alloys of the 7XXX series.

## 2. Experimental

### 2.1. Materials

The AA7010-T7451 and AA7175-T6 alloys were produced by machining (7010) or rolling (7175), resulting in thicknesses of 0.5 and 0.2 mm, respectively, and cut into specimens of 10 x 12 cm². The nominal chemical composition is given in Table 1. Surface pre-treatments, anodizing and sealing were made in pilot conditions (200 L bath). Surface pre-treatments, anodizing and sealing were conducted under pilot conditions (200 L bath). Before anodizing, the plates were degreased in a 10% vol. alkali-based bath (silicate-free, pH = 9) for 10 min at 45 °C. Rinsing with deionized water was performed between each treatment. Etching was then carried out in a mixture of 42% vol. acid-based solution (sulphuric and nitric acid) and 10% vol. ferric sulphate-based solution (bath pH < 1) for 10 min at 50 °C. Anodizing in a tartaric/sulphuric acid (TSA) bath (80 g/L of C₄H₄O₆ and 40 g/L of H₂SO₄) was then performed
at 37 °C for 35 min. These parameters allowed the formation of an oxide layer with a thickness
higher than 6 µm, which left unsealing, exhibited superior corrosion resistance in immersion
and salt spray tests compared to oxide layers obtained under standard conditions [37]. The
applied voltage was set at 14 V, a value reached after 5-min voltage ramp, as mentioned in
the specified time frame above. Once the 5-min voltage ramp is passed, the voltage stabilized
at 14 V and the current density at 1 A/dm².

<Table 1>

A two-step sealing process was performed. The first step involved a pre-treatment consisting of immersion in two different solutions, containing the active agents shown in Table 2 and further referred to as impregnation A or B. Impregnation A was performed in a single bath at 40 °C for 10 minutes. Impregnation B was performed in two successive baths: 10 minutes in a first bath containing active agents at 40 °C, followed by a 5-min immersion at 25 °C in a hydrogen peroxide-based bath. The second step included the application of an alkaline silicate-based sealing agent, applied during 20 minutes in a bath at 98 °C. Two different concentrations of silicate bath, 20 g/L or 80 g/L, were tested. The used combinations of the substrate alloys and post-treatments are listed in **Table 2**. Since the anodizing is the same for all samples, it will not be specified in the name used to describe the samples. For instance, an anodized 7010 alloy, post-treated with impregnation A and sealed in Si-based (20 g/L) will be referred to as 7010-A-Si20. Supplementary results and comparison with HWS samples are available in **Supplementary information 1**. <Table 2>

## 2.2. Scanning Electron Microscope (SEM) conditions

Surface morphology and cross-sections of the samples were examined using a field emission gun scanning electron microscope (FEG-SEM Zeiss Supra 35). EDS data were acquired with an XFlash© 6-60 from Bruker, and quantification data were obtained with the Phi-Rho-Z calculation mode. An accelerating voltage of 15 kV was used for bare alloys and cross-section observations. It was reduced to 1 kV to limit charging effect when the sealed layer was observed. No carbon deposition was applied to the surface to prevent contaminations of the porous microstructure. Cross-section were obtained by mechanical polishing, using abrasive disc of SiC from 1200, 2400, and 4000 grade. To obtain a mirror-polished surface, a colloidal silica suspension was used. Samples were rinsed in milli-Q water and ethanol and then dried
with nitrogen before SEM observations. One cross-section was made by ionic polishing, using
a PECS II model 685 by Gatan. Firstly, the samples were polished for 15 h at 5 keV, followed
by 5 h at 1 keV.

## 2.3. 3D-microscopy

A Keyence VHX-5000 microscope, equipped with a VHZ500T objective, allowing magnification from ×500 to ×500,0 was used for optical characterization. Images were taken with the ×500 magnification. The depth of defects was measured using the 3D-recomposition function provided by the microscope: by capturing multiple images at various working distances, it generates a 3D representation of the image, enabling quantitative description.

## 2.4. Glow discharge – optical emission spectroscopy (GD-OES) depth profiles

Elemental depth profiles of the anodized and sealed layers were obtained with a GD-Profiler 2™ (Horiba). Homogeneous sputtering was obtained with a 4 mm-diameter anode and the following plasma conditions: Ar plasma, a pressure of 900 Pa and a power of 20 W. The emission lines of 48 species were recorded simultaneously on the polychromator but only the relevant elements are presented and discussed below. The following elements and characteristic optical emission lines were selected: oxygen (130 nm), aluminium (396 nm), copper (325 nm), magnesium (285 nm), zirconium (360 nm), chromium (425 nm) and silicium (288 nm). In order to represent the in-depth variation of the relative intensity of the displayed elements, all the signals were normalized by the total light emitted by the plasma (Fi). Between the measurements, the anode underwent the mechanical cleaning to ensure effective characterisation of the extreme surface.

## 2.5. Time of flight – secondary ion mass spectrometry (ToF-SIMS) depth profiles

In-depth chemical analysis was also performed by ToF-SIMS using a ToF-SIMS 5 spectrometer (IONTOF, GmbH, Münster, Germany) operating under ultra-high vacuum conditions (10<sup>-9</sup> mbar). The primary source dedicated to analysis was a Bi<sup>+</sup> beam of 25 keV, delivering a target current of 1.2 pA, over a 100 × 100 µm<sup>2</sup> area. Depth profiling was performed by interlacing the surface analysis with sputtering performed using a 2 keV Cs<sup>+</sup> beam, providing a target current of 100 nA over a 300 × 300 µm<sup>2</sup> area. All measurements were
performed in negative polarity, which enhances the yield to fragments from oxidized species (e.g. AlO⁻). Data processing was performed with the Ion-Spec software. All signals were normalized by dividing their intensities by AlO⁻ signal intensity, as this signal is intense and almost constant. All signals, except AlO⁻, were multiplied by 100 and AlO⁻ signal was divided by 100.

The following ions were selected: AlO⁻ and AlO₂⁻ characteristic of the Al oxide, MO₂⁻, of the M-type oxides (with M = Zr, Si, Cr), O₂⁻ of the oxide film, and alkali⁺ as well as alkaliO⁻, characteristic of the alkali and alkali oxide, respectively, of the sealed layer.

## 2.6. X-Ray photoelectron spectroscopy (XPS)

Analysis was performed with an Escalab™ 250Xi from Thermo Scientific™, under ultra-high vacuum conditions (10⁻⁹ mbar). The source was a mono-chromated Al Kα X-ray (hv = 1486.6 eV). The X-ray spot size was 900 µm. Survey spectra was acquired at a pass energy of 100 eV and a step of 1 eV while high-resolution spectra were recorded at pass energy of 20 eV and a step of 0.1 eV. Because of the charging effect, a flood gun was used on the surface of the sample. Etching of the surface was performed during two minutes with an Ar⁺ ion beam at a voltage of 2 kV. Data treatment was performed on the Thermo Scientific™ Avantage software. Binding energy of all spectra was corrected with reference to the C 1s peak of adventitious carbon at 285.0 eV.

## 2.7. Raman spectroscopy

Raman spectra were obtained with a Renishaw Confocal Raman Microscope InVia™ piloted with Wire4.0 software under green laser excitation (doubled Nd-YAG, 532 nm, using 50 mW power). Observation and focus were made using a Leica microscope objective at x50 magnification and numerical aperture (NA) of 0.75. Both broad scan measurements (up to 5000 cm⁻¹) and low frequency region (100 – 1800 cm⁻¹) were studied. The exposure time was set at 1 s, and 50 acquisitions were recorded for each measurement. The absence of any change in the spectra and/or surface appearance during scanning was carefully verified in order to ensure there was no sample burning. Additionally, the absence of burning was confirmed by the lack peaks of corresponding to amorphous carbon, which could indicate the decomposition of residual oxalate species from the TSA treatment.
## 2.8. Degradation test procedure

The degradation cycle test was inspired by the AC-DC-AC procedure developed by Sanchez et al. [27,35] in order to analyse conversion coatings on Zn-Ni alloys. The optimized degradation procedure is schematized in **Figure 1**. Following 2-hours stabilization of the open circuit potential (OCP) and a first EIS measurement before cycling, the cycle starts with a 30-min cathodic polarization at a selected potential (-1.7 V vs the reference electrode in the presented example), followed by a 30-min relaxation at OCP, and then concludes with another EIS measurement. This potential enabled to obtain repeatable degradation cycle in less than 24 h and facilitated the visual observation of cathodic reactivity according to reaction (1b). Impedance measurements were conducted applying the parameters described in section 2.9. This cycle was repeated up to 15 times, resulting in a total maximal duration of approximately 23 h. All the results presented in the main text and figures of the main paper were obtained with the cycle using the applied potential -1.7 V vs Ag/AgCl reference electrode, however other potentials were also tested, in particular, for anodic degradation an applied potential of +0.5 V vs Ag/AgCl electrode was tested. After cycling and prior to subsequent surface characterizations, the samples were rinsed with milli-Q water and dried with compressed air.

An USB camera was positioned above the electrochemical cell to perform time lapse microscopy (TLM), capturing one image per minute throughout the entire test duration [35]. The electrochemical measurement was correlated with supplementary visual information such as the surface morphology evolution or the local pH evolution, achieved by incorporating a pH indicator (saturated phenolphthalein) in the electrolyte. The pink coloration of the phenolphthalein at alkaline pH served to reveal the intensity of cathodic reaction and location of the cathodic reactivity centers. In the current study, degradation tests were conducted using 15 mL of a 0.01 M NaCl solution. Each experiment was repeated at least 2-3 times, and the results were reproducible, thus in the figures of the results sections only one curve per experiment is presented for clarity.

<Fig. 1>

## 2.9. Electrochemical impedance spectroscopy (EIS)

A Gamry 3000 potentiostat, equipped with the Gamry Framework™ and Gamry Echem Analyst™ software was used for EIS. A classic three-electrode configuration was employed,
consisting of a Pt wire as a counter electrode (CE), a silver/silver-chloride (Ag/AgCl) as a reference electrode (RE) and the sample as the working electrode (WE). A home-made electrochemical cell was designed and 3D-printed using inert polylactic acid (PLA). The active surface of 3.8 cm² was delimited by an O-ring. Samples were immersed in 15 mL of 0.01 M NaCl without stirring. The impedance measurements were performed at the OCP from 100 kHz to 0.01 Hz, with a 15 mV root mean square (rms) amplitude and 7 points per frequency decade.

## 3. Results

### 3.1. Substrate microstructure

To better understand the role of the IMPs on the immersion stability of sealed anodized layer, the distribution of the IMPs in the substrates were compared. Firstly, the backscattered electron (BSE) images of the surface of the machined 7010 alloy and the rolled 7175 are presented in **Figure 2**. Additionally, cross-sections of the substrate alloys were examined for the intact 7010-B-Si80 and 7175-B-Si80 configurations prior to applying degradation procedure illustrated in **Figure 1**. Representative secondary electron (SE) and backscattered electron (BSE) cross-sections are shown in **Figure 3**. Although several types of Cu-containing intermetallics were present for both alloys, only the composition of the biggest particles are indicated in the figures. The identification was made by considering the ratios of atomic fractions of the present elements, while also taking into account the existing knowledge about the composition of IMPs in 7010 and 7175 alloys. In particular, in AA7X75 and AA7475 $Al_7Cu_2Fe$, $(Al,Cu)_6(Fe,Cu)$ are the most abundant IMPs but $Al_{23}CuFe_4$, $Al_6Fe$, $Al_{12}(Fe,Mn)_3Si$, $Al_2CuMg$ and $Al_{40}Cu_{15}Mg_{25}Zn$ were also reported while in AA7010 only the $Al_7Cu_2Fe$ intermetallics are usually present in large quantities [38].

<Fig. 2>

<Fig. 3>

The SEM cross-section analysis reveals that, despite surface pre-treatments (degreasing, etching) used to remove the intermetallic particles (IMPs) before the first step of anodizing, IMPs are still present near the interface at a depth less than several micrometers. Although Al-Cu-Fe containing particles were detected in both substrates, the nature of the
biggest IMPs, identified by EDS, seems to be different between the two alloys. Isolated very big (up to tens of µm long) Al<sub>7</sub>Cu<sub>2</sub>Fe IMPs are present in machined 7010, while numerous round shaped Al<sub>2</sub>Cu and some AlCuFe containing IMPs with more complex shape in rolled 7175 are much smaller. This difference is also evident on the SEM surface of the bare alloys (Figure 2). The average size of IMPs was calculated using ImageJ software. For the bare machined 7010, the average size was 4.1 ± 1.3 µm<sup>2</sup>, whereas for the bare rolled 7175, it was 1.7 ± 0.4 µm<sup>2</sup>. Focusing IMPs larger than 1 µm<sup>2</sup>, the difference in average size was even more apparent: 13.2 ± 5.2 µm<sup>2</sup> for the 7010 and 4.3 ± 0.3 µm<sup>2</sup> for the 7175. The total surface analysed for both samples was around 0.118 mm<sup>2</sup>. On this surface, the total of 82 IMPs with the area higher than 1 µm<sup>2</sup> were found for the 7010 while the total reached 139 IMPs with the area higher than 1 µm<sup>2</sup> for the 7175. This image analysis effectively illustrates the difference in size and distribution of IMPs between the two Al substrates. Therefore, regardless of surface treatments (including the same anodizing, impregnation, and sealing procedures), the nature, distribution and size of intermetallic particles in the substrates stays different.

## 3.2. Surface morphology before degradation

A comparison of the surface microstructure between silicate-sealed samples with various substrate and impregnation is displayed in **Figure 4**. A similar cauliflower morphology is observed for all sealed samples. It is clear from this figure that the morphology of the sealing is not affected by the substrate (machined 7010 or rolled 7175) or the impregnation solution (A or B).

<Fig.4>

## 3.3. Chemistry of the sealed oxide layer before degradation

Although the EDS analysis of the surface was complicated by surface roughness, oxide porosity and strong charge accumulation, it revealed the presence of major Al and O elements and confirmed the presence of minor elements from the treatment bath (S, Cr, Zr and Si). One should note that because of possibly uneven distribution of these elements inside the layer and a relatively high penetration depth of the EDS analysis. Therefore, the description of the layer requires to understand their distribution in depth of the layer, which was analysed by GD-OES. A representative depth profile of a sealed anodized sample (7010-A-Si2O) is shown in **Figure 5**. The elemental profile is supported by a SEM cross-section generated through ionic
polishing to elucidate the fine microstructure of the layer. The microstructure of the sealed anodized layer can be divided in two zones. The first one, close to the surface up to around 200 s of erosion, exhibits a compact microstructure attributed to the sealed layer. This is corroborated by the GD-OES profile, where the main elements of the pre-treatment bath and Si, indicating the silicate-based sealing, are found in this area. Sealing elements from the pre-treatment bath (Zr, Cr) reach their maximum signal in the initial 30 seconds of erosion. The Si signal is exclusively present at the extreme surface, within the first seconds of erosion. An interesting observation lies in the position of the alkali element present in the silicate sealing bath. The alkali element is not present in the top-surface region (unlike Si, Cr or Zr) but is buried deeper in the surface layer. The second zone, from around 200 s of erosion to around 700 s of erosion, is attributed to the anodized layer not affected by sealing due to the presence of a much finer porous microstructure and the absence of sealing markers.

<Fig.5>

Comparing the elemental depth profiles obtained on two different substrates (see examples of the 7010-A-Si20 and 7175-A-Si20 samples in **Figure 6**), one can conclude that the layer thickness and element distribution in depth of the layer is similar for the substrates. Indeed, in both samples, the sealed anodized layer, identified by a plateau formed by the O signal, lasts during the first 600 seconds. The alloy substrate is reached when the Al signal stabilizes. The interface between the layers, defined by the transition zone where the O signal decreases and the Al signal increases, depends on the interface roughness, so similar erosion time indicates comparable roughness of the interface.

<Fig.6>

It is noteworthy that the signal intensities of the major alloying elements (Zn, Mg, Cu) (see **Figure 6c and d**) are very low (almost zero) in the sealed oxide and follow the signal of Al at the interface and in the substrate for both substrates. This confirms the presence of these elements in the substrate and implies that almost no intermetallic particles are present in the sealed oxide layer.

In depth element distribution in the sealed layer was confirmed by a ToF-SIMS depth profiling. One should note that in contrast to the GD-OES profiles (**Figure 6**) showing the entire layer thickness from the surface up to the bulk alloy, the ToF-SIMS profiles (**Figure 7**) represent
only the first 0.5 µm depth from the surface. The displayed signals are noisy due to the low conductivity of the eroded oxide layer, despite the use of charge compensation during the measurement.

<Fig. 7>

Elements from the sealing pre-treatment, represented by the Zr and Cr oxide profiles ($ZrO_2$⁻ and $CrO_2$⁻, respectively), are visible close to the surface with a maximum reached during the first 500 seconds of erosion, followed by a slight and progressive decrease. The Si oxide is present at the extreme surface, as indicated by its maximum intensity from the first seconds of sputtering and its rapid decrease (logarithmic scale). These observations are consistent with the GD-OES depth profiles (Figure 6), showing the Cr and Zr in the oxide layer and the Si at the extreme surface. Moreover, the alkali ion profiles (alkali⁻ and alkali$O^{+}$) also confirm the GD-OES results. The intensities of these two species, increase progressively from 500 seconds of sputtering to reach a maximum intensity plateau. Similarly, to the GD-OES profiles, the alkali element is not present at the extreme surface but slightly deeper.

The chemical form of different elements in the layer was verified by Raman spectroscopy. Figure 8 demonstrates the mains attributions on the example of the 7010-A-Si20 sample. The presence of sulphate ions from the TSA anodizing is confirmed by two peaks at 981 cm⁻¹ and 1060 cm⁻¹, attributed to the symmetric stretching $v_1$ of $SO_4^{2-}$ and the symmetric stretching vibration $v_S$ of $HSO_4^-$ respectively [39]. The broad band between 500 – 600 cm⁻¹ could be due to a mixture of Cr(III) and Zr(IV) oxides [40]. The presence of a distinct peak between 4000 and 5000 cm⁻¹ and a smaller peak at around 750 cm⁻¹ could be attributed as characteristic for spinel-like structures [41,42], which leads us to advance a hypothesis of a spinel-type aluminate, likely formed with the alkali element present in the sealing. The presence of a spinel-like compound combined with the fact that the alkali element is buried deeper under the surface, suggests that during its diffusion within the pores, it facilitates the transformation of oxides into a spinel-type aluminate. The consequences of the spinel formation spinel on the susceptibility to localized corrosion in case of local damage of the oxide film could be further studied.

<Fig. 8>
The results of the analysis of the top surface chemistry of intact samples by XPS are presented in section 3.6, where they are compared with the results obtained for the electrochemically degraded surface.

### 3.4. Accelerated ageing of the sealed layer

#### 3.4.1. Electrochemical parameters evolution

To estimate the stability of the sealing, the accelerated ageing procedure described in section 2.8 was applied to different configurations of aluminium alloys/anodizing/impregnation/sealing.

The evolution of the system during cycling can be first followed by the OCP and cathodic current. **Figure 9a** displays the evolution of the OCP value with cycling, measured for each degradation cycle at the end of the 30-min of the “OCP after” step shown in **Figure 1**. **Figure 9b** presents the cathodic current value (measured for each degradation cycle at the end of the 30 min of the “Cathodic polarization” step shown in **Figure 1**) as a function of the cycle’s number. The OCP decreases in both cases, reaching the values close to -1.15 V for 7010-B-Si80 and -0.95 V for 7175-B-Si80. For 7010 substrate, a slight current decrease down to -0.11 mA after 15 cycles can be observed, contrary to 7175 substrate, which shows a significant decrease to -0.82 mA after 15 cycles.

<Fig.9>

Although the cathodic stability of silicate sealed anodized layers is in the centre of the present work, we would like to comment that similar degradation cycle but with application of anodic polarisation was also tested. While cathodic degradation of the surface and modification of the OCP and cathodic current became evident during less than 5 - 6 cycles, no visual or electrochemical modification of the interface occurred after 6 cycles with anodic polarization (illustrated in **Supplementary information**). This suggests good anodic stability of silicate-based sealing and supports our choice to continue deeper understanding of cathodic degradation.

The evolution of the impedance modulus value at 0.01 Hz as a function of the number of degradation cycles is displayed in **Figure 10**. The figure clearly indicates that for each
configuration, the modulus value decreases with the number of cycles, indicating a loss of
corrosion resistance. For all sealed samples, the modulus values show a similar trend, with a
relatively stable value during the initial cycles (typically around 5 cycles but varying between
2 and 8 cycles depending on the sealing and the substrate combination), followed by a
pronounced decrease. The protective effect of the sealing during the whole ageing process is
evident here, as indicated by the higher impedance modulus observed for the sealed samples
compared to the unsealed. Furthermore, the impact of the sealing bath concentration is
noticeable. In equivalent configuration (same alloy, same anodizing, same impregnation), the
7175 samples treated within the higher concentrated sealing bath display a higher impedance
at 0.01 Hz, suggesting better corrosion resistance. The two configurations expected to have
the best corrosion resistance after 15 cycles of degradation, based on **Figure 10**, are the 7010
and the 7175, both treated with the same impregnation B associated with Si-based sealant in
an 80 g/L bath. Among them, the sample based on the machined 7010 substrate exhibits the
highest corrosion resistance.

<Fig.10>

EIS Bode plots obtained for 7010-B-Si80 before cycling, after 4 and 15 cycles of ageing, and for 7175-B-Si80 before cycling, after 6 and 15 cycles of ageing are displayed in **Figure 11**. The profile obtained before cycling reveals a shape typical for sealed anodized layers [8,10,11,43–49]. Two well-defined time constants can be distinguished, where the mid/high frequency part characterizes the sealed oxide layer, while the low frequency part the barrier layer. After 4 cycles of ageing for the configuration with 7010 substrate (and after 6 cycles for the configuration with 7175), a modification of the plot appears in the mid/high frequency region, indicating a degradation of the sealed oxide layer. This can be interpreted as a perforation of the sealed oxide layer. In the spectra after 15 cycles, a new contribution becomes apparent in the low frequency region. This can be attributed to the complete perforation of the barrier layer, resulting in the exposure of the aluminium alloy substrate. Such a contribution was previously attributed to the electrochemical processes at the metal/electrolyte interface [11]. A significant loss of corrosion resistance is expected at this time because the value of the impedance modulus at 0.01 Hz decreased from 9.9 MΩ.cm² (before cycling) to 0.22 MΩ.cm² (after 15 cycles) for the layers formed on 7010 substrate. For the configuration with 7175 substrate, the impedance decreased from 8.8 MΩ.cm² to 0.15
MΩ.cm². While the impedance modulus at 0.01 Hz provides information on the overall corrosion resistance, examining the modulus at 100 Hz is useful for monitoring the evolution of the sealed oxide layer, in particular. The evolution of the modulus at 100 Hz as a function of the number of cycles for 7010-B-Si80 and 7175-B-Si80 is displayed in **Figure 12**. The modulus at 100 Hz evolution shows a similar trend to that at 0.01 Hz for each sample. For the 7010-B-Si80 sample, the impedance modulus at 100 Hz decreases from 0.14 MΩ.cm² (before cycling) to 0.01 MΩ.cm² (after 15 cycles). For the 7175-B-Si80 sample, the impedance modulus at 100 Hz before cycling is 0.24 MΩ.cm² and drops to 0.002 MΩ.cm² after 15 cycles. It is noteworthy that despite a stronger impedance modulus for the sealed oxide part before cycling, the degradation of this contribution is higher for the sample 7175-B-Si80. Nyquist plot for these two samples, 7010-B-Si80 and 7175-B-Si80, are displayed in **Figure 13**. Before degradation, the presence of a semi-circle shape at high/mid frequency is associated to the sealed oxide layer, while the contribution at low frequency is characteristic of a CPE and associated with the barrier layer of anodizing.

<Fig. 11>

<Fig. 12>

<Fig. 13>

### 3.4.2. Deeper interpretation of impedance data for intact samples

To further clarify the impedance data, equivalent circuits were suggested for fitting the experimental results. Among the numerous models presented in the literature, the most suitable equivalent circuit is depicted in **Figure 14** [8,44,45,49–51]. To improve the fitting accuracy and express the non-ideal behaviour of the capacitor (because of surface roughness, for example), constant phase elements (CPE) were employed instead of pure capacitors. The impedance of a CPE is given by the following relation [52]:

$$Z_{CPE} = \frac{1}{(j2\pi f)^{\alpha_Q}} \quad (3)$$

Q and α represent the CPE parameters where α is the frequency dispersion factor, the coefficient that characterizes the non-ideal behaviour. If α = 1, the CPE is a pure capacitance, while if α = 0 the CPE is a pure resistance. R_e corresponds to the electrolyte resistance, R_p and
$Q_p$ are the sealed layer resistance and its CPE parameter, respectively, $Q_{pw}$ is the CPE parameter of the pore walls. The pore walls resistance $R_{pw}$, in parallel with $Q_{pw}$, can be neglected because it is usually so high due to its non-conductivity that the passage of current across the walls is not possible [44]. $R_b$ and $C_b$ are respectively the barrier layer resistance and capacitance. As the barrier layer is not expected to be affected by the sealing, no CPE is used to adjust its parameters.

<Fig. 14>

After fitting, the quality of the data was good ($\chi^2 \approx 10^{-4}$), but the error on $R_p$ was significant (same order of magnitude or even higher). Despite attempting variations in the elements of the equivalent circuit (such as switching capacitance by CPE), no convincing fitting was obtained for our experimental data. Additional details on the tested equivalent circuit and fitting are available in **supplementary information 2**. To gather more information from the impedance data, graphical and computational methods were employed. Assuming a pure capacitive behavior at high frequency, the contribution can be modelized by a resistance $R_p$ in parallel with a capacitance $C_p$. The low frequency contribution is defined by a CPE $Q_b$ in parallel with $C_b$.

For the high frequency semi-circle, the value of $R_p$ can be read by looking at the circle diameter intercept with the $Z_r$ axis and subtracting the value of the electrolyte resistance $R_e$ [53]. Since the semi-circle diameter value is hidden by the CPE, one can look at the highest $Z_{im}$, $f_{max}$ point of the semi-circle to obtain the radius of the circle $Z_{r,max}$. $C_p$ can then be calculated using the following formula:

$$C_p = \frac{1}{2\pi f_{max}^2 Z_{r,max}} \quad (4)$$

where $f_{max}$ is the frequency at which $Z_{im}$, $f_{max}$ and $Z_{r,max}$ were found. Once the value of $C_p$ is known, it is then straightforward to calculate the thickness of the sealed oxide layer using the following expression:

$$e_p = \frac{\epsilon_0 \epsilon_r S}{C_p} \quad (5)$$
where $ε_0$ represents the dielectric permittivity of vacuum (8.854 x 10⁻¹⁴ F.cm⁻¹), $ε_r$ is the alumina dielectric constant (20), and S is the exposed area in cm² [44]. Because the sealed oxide layer is not pure alumina, as the sulfur ions can be incorporated in the layer during anodizing, a value of 20 for the dielectric constant is considered instead of usually 10 for pure alumina [44].

For the low frequency part, it is necessary first to separate $R_pC_p$ (semi-circle) from $Q_bC_b$ (the CPE). To do so, an extrapolation of the values at very low frequency is performed, to find at the intercept with the real axis, a value that we named $Z_{rCPE}$. Then, $Z_{im} = f(Z_{r-} Z_{rCPE})$ is plotted for the points at very low frequency in order to find the value of the CPE $Q_b$. The plotted line should pass through the origin and form an angle with the real axis that we named θ. The value of θ is found using the expression:

$$ \theta = \tan^{-1}(\frac{Z_{im}}{Z_{r}-Z_{rCPE}}) \quad (6) $$

The value of $Q_b$ is found using formula (3.1), where α is found by dividing θ by 90 and f is selected at the lowest frequency measured. To convert $Q_b$ into a capacitance $C_b$, the following equation is used [54]:

$$ C_b = Q_b^{\frac{1}{\alpha}} Z_{rCPE}^{\frac{1-\alpha}{\alpha}} \quad (7) $$

Using the formula (5), it is then easy to calculate the thickness of the barrier layer $e_b$. The obtained values are summarized in Table 3. The values of $R_e$ are low, which is coherent with the expectations. The thicknesses of the sealed oxide layer are in the order of magnitude of the expected thickness aimed by the anodizing process (4-7 µm). As the thickness of the barrier layer usually grows at the rate of 1.25-1.4 nm.V⁻¹, the values of $e_b$ are coherent with the reality of the process, as the anodizing voltage was 14 V [3]. The overall improvement in corrosion resistance seems to come from the increased barrier effect of the sealed oxide layer.

<Table 3>
### 3.5. Morphological evolution of the surface during ageing

#### 3.5.1. Live imaging

The in situ solid-liquid interface morphology observed after 15 cycles of ageing, typical for the samples with configuration 7010-B-Si80 and 7175-B-Si80, is shown in **Figure 15a** and **Figure 15b**, respectively. Several other configurations of anodizing/impregnation/sealing were tested and showed the identical degradation mode. As seen in **Figure 15** and **Figure 16**, the degradation indicates a localized surface attack, wherein bubbles evolve under cathodic polarization, and several black pits appear during the succeeding polarization cycles. The generation and bursting of bubbles observed on the surface are caused by the cathodic reaction (1b), favoured during the degradation cycle. A short-accelerated video of the surface evolution is available in **supplementary information 3**.

<Fig. 15>

<Fig. 16>

Quantitative analysis of the in situ optical micrographs for the 7010-B-Si80 configuration revealed that the number of pits increases almost linearly with the number of cycles, as shown in **Figure 17a**. Tracking the number of pits during the ageing for the 7175-B-Si80 was not possible due to their excessively high number on the surface.

<Fig. 17>

#### 3.5.2. Ex situ optical images

The morphology of the black pits was analysed after the degradation cycles using ex situ 3D optical and scanning electron microscope. The pit depth is comprised between 12 and 30 µm as estimated by 3D optical microscopy (see **Figure 18**). This indicates that the sealed oxide layer is locally completely removed, considering the expected thickness of this layer is at most 7 µm. The pit depth measurement demonstrated that the first pits formed after 5 cycles of ageing already had a depth of about 10 µm, thus perforating the anodized and sealed layer. Simultaneously, the size of average pit area did not increase significantly with cycling. The relatively stable size of pits, but the increase of number of pits, is coherent with the fact that cathodic current increases.

<Fig. 18>
Ex situ optical observations demonstrated also significant differences in the size and distribution of the pits between the two configurations. As displayed in **Figure 18**, the machined 7010-B-Si80 configuration presented only a few large pits, while the rolled 7175-B-Si80 surface was covered by numerous smaller pits. The size of defects on the 7010 substrate was around 200 µm, while it was around 100 µm for the configurations with the 7175 substrate. This is consistent with the results of live imaging showing only few bubbles on the 7010-B-Si80 contrary to the 7175-B-Si80 with high number of bubbles appearing over the whole surface during degradation. The pits were also observed by SEM images, displayed in **Figure 19**, confirming the information given by optical microscopy.

<Fig. 19>

Considering the areas between pits, which appeared non-modified, SEM observations revealed no morphological changes before and after degradation, as illustrated in **Figure 20**. The characteristic cauliflower morphology is preserved, similar to those observed on the intact samples (**Figure 4**).

<Fig. 20>

In our experiments, pits distribution on the surface after degradation varied significantly for configurations with different substrates: the 7010-B-Si80 presented only a few pits, while the 7175-B-Si80 surface was covered with numerous pits. This correlates well with the respective distribution of IMPs (see **Figure 2** and **Figure 3**) in the alloys before degradation: a few, large IMPs in the machined 7010 and numerous smaller IMPs in the rolled 7175. This clearly shows that the distribution of pits was independent of the surface treatments (impregnation/sealing) but depended on the substrate alloy processing: machining or rolling.

### 3.6. Chemical evolution of sealed samples after ageing

To verify if the sealed layer was chemically attacked in the areas between pits, EDS analysis was performed on the surface of sample exposed to 15 cycles of degradation and on the non-exposed surface (**Table 4a**). One should note that the EDS analysis shows the average ratios between the elements for the whole analysed depth which is in order of micrometre for light materials. The results show no difference in the chemical composition of the layer.
Additionally, XPS measurements were conducted on the pristine surface (not exposed to the electrolyte) and on the surface that was exposed to the electrolyte but appeared intact after 5 cycles of degradation. The objective was to verify a possible evolution of the elemental composition and/or chemical specification of the elements present in the top layer of the sealed anodized oxide. Quantitative results are displayed in **Table 4b**. Si, Cr, Al and O are detected as the main elements, which is coherent with the expected composition of the sealed surface and with the ToF-SIMS results. The detection of Si agrees with the previous observations made by GD-OES and ToF-SIMS that Si is mainly present on the top of the sealed oxide layer.

<Table 4>

XPS revealed also that there were no major differences in the chemical states between the extreme surfaces. For example, the XPS high resolution spectra of Si 2p for 7010-A-Si20 before and after 5 cycles of degradation (Figure 21) shows two peaks attributed to R₂-SiO₂ (102.1 ± 0.2 eV) and SiO₂ (103.2 ± 0.2 eV) [55]. The focus was made on the Si 2p because, as this element was observed almost exclusively at the extreme surface of the silicate-sealed sample (see Figure 6 and Figure 7), a chemical change could be expected. There are almost no differences between the spectrum before and after 5 cycles of degradation, indicating that the top surface chemistry of Si is not changed. Similar results were observed for other surface elements (Cr, Zr), not shown here. One should note that no Cr(VI) was detected by XPS on the extreme surface. Table 4.2. demonstrates also that for the sample 7010-B-Si80, obtained with the most concentrated in Si bath, wt. % fraction of Si is significantly higher and Al significantly lower than on the top surface layer of 7010-A-Si20 sample. This indicates a thicker silicate layer for 7010-B-Si80, consistent with the bath composition and other analyses. High thickness of the top-silicate layer can explain too weak concentration of Cr on the top surface and/or an in-detph detection limit of the XPS analysis technique (expected to be around 5 nm maximum).

<Fig. 21>

SEM-EDS mapping of the pit and its surrounding zone is displayed in the **Figure 22** for the sample 7010-B-Si80. It is evident that the pit depth reaches the alloy, as for example, elements mainly found in the sealing such as the Cr are not visible in the pit. Moreover, alloying
elements such as Cu are clearly visible inside the pit, demonstrating that the alloy is reached and supporting optical microscope observation from **Figure 18**.

To sum up sections 3.5. and 3.6, SEM and optical microscopy results demonstrated that the sealing was attacked only locally in places where pits were developed up to the underlying metallic substrate. Moreover, the results of EDS and XPS suggest that there is no significant global surface chemistry modification in the sealed layer outside the pits. Therefore, the degradation of the sealed layer is expected to be a result of the underlying oxide dissolution.

## 4. Discussion

The presented results underscore the impact of silicate-based anodizing post-treatments on the microstructure and immersion stability of TSA-anodized, 7XXX series Al alloys, as well as the influence of the alloy substrate, particularly its processing, on these properties. For all the systems the sealed microstructure effectively prevented penetration from the aggressive species in the substrate, aligning with the higher impedance modulus observed for the sealed samples during the ageing cycle. Regardless of the alloy or the surface treatments, samples resisted 15 cycles of anodic polarization without any visual modification or loss of the electrochemical characteristics but a localized attack occurred under cathodic degradation, leading to the formation of black pits on the surface. The number of defects escalated almost linearly with the cycle number, which is in accordance with the increasing current measured at the end of each polarization step. The surface morphology of degraded samples presented in **Figure 15** clearly reveals a distinct distribution of pits in the TSA-anodized sealed layers formed on different substrates. Layers formed on machined 7010 sample presents only few large pits, while the layers formed on the rolled 7175 shows smaller but numerous pits. Interestingly, a similar tendency can be observed for the distribution of the IMPs in the substrates reflects - only a few large IMPs are present in the machined 7010 and numerous smaller IMPs in the rolled 7175 substrates, which could be attributed to the differences in alloy processing. As the 7010-B-Si80 samples demonstrated better corrosion resistance than the 7175-B-Si80 samples, it can be proposed that it is not necessarily the size of the pits but more their density that influences the corrosion behaviour. This suggestion is also in agreement with the cathodic current evolution measured during ageing at the end of each polarization cycle. On one side, cathodic current increased much more for the 7175 than the 7010 samples. Moreover, considering the linear increase of the measured cathodic current
with the number of cycles (Figure 9b), one could expect that the number of pits is proportional to the measured cathodic current during the test, as the linearity appears from the cycle 8 (see Figure 17b). The linear growth of the number of pits with cycling suggests that the distribution of pits is independent of the surface treatments (impregnation/sealing) and depends on the substrate alloy or processing: machining or rolling. It is well-established that the alloy processing strongly impacts the intermetallic particles (IMPs) distribution and size in the bulk, modifying the properties of the alloy [1,23,24]. To our knowledge, it has however never been mentioned in the literature that the distribution of the IMPs in the buried substrate can affect the performance of sealing through the 4 to 7 µm thick oxide layer. Our results demonstrate, for the first time, such an effect because the only difference between the two samples was the types of the alloy processing (machining for the 7010 substrate and hot rolling for the 7175 alloy).

A mechanism, explaining the localized degradation of sealed anodized layer observed during the ageing procedure and its dependence on the distribution of IMPs is proposed in Figure 23. After the electrolyte penetration into the sealed layer (step 1), the pH increases by the reduction of water and oxygen to hydroxide ions (reaction 1) during cathodic polarization. The pH increase can contribute to the partial dissolution of the oxide layer, favouring increased diffusion of species to the substrate. If both matrix and IMPs are exposed to the electrolyte, stronger cathodic reactivity is expected on the IMPs because of their more noble potential than the matrix during the OCP period. The dissolution of the anodized oxide layer by reaction (2) will be hence favoured over the IMPs, leading to a rapid pit development (step 2). As a result of the pit formation, a galvanic couple matrix - IMPs is exposed to the electrolyte and anodic dissolution of the matrix surrounding the IMPs leads to its detachment from the substrate (step 3).

<caption><Fig. 23></caption>

The IMPs therefore act as preferential sites for the cathodic reactivity, elucidating the observed localized attack. The emergence of a new contribution in the Bode plots, corresponding to the reactivity at the metal/electrolyte interface, is consistent with the fact that the aluminium alloy is exposed in the pits because their depth exceeds the thickness of the anodized layer. It should be however considered that the proposed schema (in Fig. 23) applies well to the cases of relatively thin anodized layers studied in the present work, for
which, the absence of the intermetallic particles in the anodised / sealed layer can be ensured.
For thick anodised layers, it could be probably impossible to avoid incorporation of IMPs into
the oxide, thus making the mechanisms would more complex.

The discrepancies in the OCP evolution with cycling among the samples can be
attributed to the distinct distribution of IMPs in the substrates. For the 7010 alloy, due to the
IMPs distribution, only a few large pits were formed. The formation of these pits is expected
to be favoured by an isolated big intermetallic particle, and when this IMP is detached, the
exposed area in the pit is expected to be close to pure Al. For this alloy, the measured final
OCP of around -1.2 V vs Ag/AgCl approaches the values reported in the literature for Al alloy
in an alkali environment [56]. In contrast, the 7175 is richer in IMPs, so even if some of them
detach from the matrix, it is still expected that some remain in contact with the electrolyte.
The measured corrosion potential in the pit can therefore reflect galvanic coupling between
the IMPs (mainly Al<sub>2</sub>Cu, with a potential around -0.74 V vs saturated calomel electrode (SCE)
(so -0.695 V vs Ag/AgCl [57]) in alkali medium) and the matrix [26]. Indeed, the measured value
of corrosion potential for the anodized and sealed 7175 is more noble (-0.95 V vs Ag/AgCl)
than for the anodized and sealed 7010 (-1.2 V vs Ag/AgCl).

5. Conclusions

This study focused on the effect of the substrate alloy microstructure on the
composition, morphology and cathodic stability of anodized layers with silicate-based
sealants. The main results can be sum up as follows.

1. The morphology and composition of the sealed layer remained unaffected by the substrate.

2. To access the stability of the sealed layer, a novel accelerated ageing procedure was proposed, which was employed to compare various sealants on two distinct substrates.

3. The accelerated tests enhancing cathodic degradation, demonstrated similar degradation mode of the sealed layer involving pits formation across all sample configurations, including significant variations in IMPs distribution within the substrate. However, no modifications were detected in the sealed layer within intact areas between pits.
4. Electrochemical and time lapse microscopy results demonstrated a beneficial effect of the sealing on the corrosion resistance of all studied configurations and suggested that this enhancement stems from the improved barrier effect of the sealant.

5. The number of the formed pits increased over time, correlating with the measured cathodic current and the evolution of the OCP. Nevertheless, the distribution and size of the pits varied between the two substrates and correlated with the distribution of IMPs within the aluminium alloy buried under the 6 µm thick oxide layer.

6. A mechanism, linking the localized attack on a thick oxide layer to the distribution of the IMPs in the buried substrate is proposed.

Our findings imply that the cathodic degradation of TSA-anodized and sealed Al alloys of AA7010-T7451 and AA7175-T6 series results from the destabilization of the metal-anodized oxide interface rather than the degradation of the sealed layer.

## Data Availability Statement

The data presented in this study are available on request from the corresponding author.

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<figure><img src="image_2.png" /><figcaption>**Figure 1.** Degradation procedure used during this study.</figcaption></figure>

<figure><img src="image_3.png" /><figcaption>**Figure 2.** SEM surface micrographs (BSE images) of the bare machined 7010 alloy and the rolled 7175. Only the composition of the biggest intermetallics is indicated.</figcaption></figure>
<figure><img src="image_4.png" /><figcaption>**Figure 3.** SEM micrographs (SE and BSE images) of cross-sections of intact 7010-B-Si80 and 7175-B-Si80 configurations. Only the composition of the biggest intermetallics is indicated.</figcaption></figure>

<figure><img src="image_5.png" /><figcaption>**Figure 4.** SEM top view of the surface for sealed samples.</figcaption></figure>
<figure><img src="image_6.png" /><figcaption>**Figure 5.** GD-OES profile of 7010-A-Si20, supported by a SEM cross-section obtained by ionic polishing.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>**Figure 6.** GD-OES elemental depth profile of: Global profile of 7010-A-Si20 (a) and 7175-A-Si20 (b) Profile focus on alloying elements of 7010-A-Si20 (c) and 7175-A-Si20 (d).</figcaption></figure>
<figure><img src="image_8.png" /><figcaption>**Figure 7:** ToF-SIMS depth profiles of the 7010-A-Si20.</figcaption></figure>

<figure><img src="image_9.png" /><figcaption>**Figure 8.** Raman spectra recorded on the surface of 7010-A-Si20, centred at 1000 cm⁻¹ and 4700 cm⁻¹.</figcaption></figure>
<figure><img src="image_10.png" /><figcaption>**Figure 9.** Evolution of (a) OCP value (V. vs Ag/AgCl) measured at the end of each degradation cycle and (b) cathodic current measured at the end of each polarization step as a function of the number of the degradation cycles.</figcaption></figure>

<figure><img src="image_11.png" /><figcaption>**Figure 10.** Evolution of impedance modulus at 0.01 Hz as a function of the number of ageing cycles. Dots symbols represent the samples with AA7010 substrate, triangle symbols are given for the samples with AA7175 substrate. For the samples with sealing pre-treatment A, the symbol is empty and for the sealing pre-treatment B, it is full. A darker colour indicates a higher concentration of sealing bath for alkaline silicates. The sample names are given in the legend.</figcaption></figure>
<figure><img src="image_12.png" /><figcaption>**Figure 11.** a) Bode plots for 7010-B-Si80 configuration obtained before cycling, after 4 cycles and after 15 cycles. B) Bode plots for 7175-B-Si80 configuration obtained before cycling, after 6 cycles and after 15 cycles.</figcaption></figure>
<figure><img src="image_13.png" /><figcaption>**Figure 12.** Evolution of impedance modulus at 100 Hz as a function of the number of ageing cycles for 7010-B-Si80 and 7175-B-Si80.</figcaption></figure>

<figure><img src="image_14.png" /></figure>

<figure><img src="image_15.png" /></figure>

<figure><img src="image_16.png" /></figure>

<figure><img src="image_17.png" /></figure>

<figure><img src="image_18.png" /></figure>

<figure><img src="image_19.png" /><figcaption>**Figure 13.** Nyquist plots for 7010-B-Si80 configuration obtained a) before cycling, b) after 6 cycles and c) after 15 cycles; for 7175-B-Si80 configuration obtained d) before cycling, e) after 6 cycles and f) after 15 cycles.</figcaption></figure>
<figure><img src="image_20.png" /><figcaption>**Figure 14.** Equivalent circuit used to fit the impedance data of 7010-B-Si80 and 7175-B-Si80 before degradation.</figcaption></figure>

<figure><img src="image_21.png" /><figcaption>**Figure 15.** Typical in situ optical microscope images of the surface (top-view). Example of two configurations after 15 cycles of ageing: a) 7010-B-Si80 (machined) and b) the 7175-B-Si80 (rolled). The samples differ only by the substrate alloy.</figcaption></figure>
<figure><img src="image_22.png" /><figcaption>**Figure 16:** Typical in situ optical microscope images of the surface (top-view). Example of 7010-B-Si80 (machined) after 3, 10 and 15 cycles of degradation.</figcaption></figure>

<figure><img src="image_23.png" /><figcaption>**Figure 17.** a) Number of pits as a function of the number of cycles for 7010-B-Si80. b) Cathodic current measured at the end of each polarization step as a function of the number of pits for 7010-B-Si80.</figcaption></figure>
<figure><img src="image_24.png" /><figcaption>**Figure 18.** Optical microscope observation of the pit and its associated depth profile, after 15 cycles of degradation.</figcaption></figure>

Top: 7010-B-Si80, a) biggest pit, b) representative pit, c) smallest pit.

Bottom: 7175-Si80, d) biggest pit, e) representative pit, f) smallest pit.

<figure><img src="image_25.png" /><figcaption>**Figure 19.** SEM surface top-view of two configurations after 15 cycles of ageing: a) 7010-B-Si80 (machined) and b) the 7175-B-Si80 (rolled). The samples differ only by the substrate alloy.</figcaption></figure>
<figure><img src="image_26.png" /><figcaption>**Figure 20.** SEM top view of the surface for 7010-B-Si80 (a) before degradation and (b) after 15 cycles of degradation.</figcaption></figure>

<figure><img src="image_27.png" /><figcaption>**Figure 21.** High resolution XPS spectra for the Si 2p of 7010-A-Si20 sample, a) before degradation and b) after 5 cycles.</figcaption></figure>
<figure><img src="image_28.png" /><figcaption>**Figure 22.** SEM-EDS mapping of 7175-B-Si80, showing a pit and its surrounding zone.</figcaption></figure>

<figure><img src="image_29.png" /><figcaption>**Figure 23.** Speculative mechanism proposed for the localized attack explanation.</figcaption></figure>
Tables

<table><caption>Table 1: Nominal chemical composition of the alloys (wt.%).</caption><thead><tr><th>Alloy</th><th>Si</th><th>Fe</th><th>Cu</th><th>Mn</th><th>Mg</th><th>Cr</th><th>Zn</th><th>Ti</th><th>Others</th><th>Al</th></tr></thead><tbody><tr><td>AA7010</td><td>0.12</td><td>0.15</td><td>1.5-2.0</td><td>0.1</td><td>2.1-2.6</td><td>0.05</td><td>5.7-6.7</td><td>0.06</td><td>0.15</td><td>Bal.</td></tr><tr><td>AA7175</td><td>0.15</td><td>0.2</td><td>1.2-2.2</td><td>0.1</td><td>2.1-2.9</td><td>0.18-0.28</td><td>5.1-6.1</td><td>0.1</td><td>0.15</td><td>Bal.</td></tr></tbody></table>

<table><caption>Table 2: List of the samples and abbreviations used for their description.</caption><thead><tr><th>Name</th><th>Alloy</th><th>Anodization</th><th>Impregnation</th><th>Sealing</th></tr></thead><tbody><tr><td>7010-A-Si20</td><td rowspan="2">7010</td><td rowspan="2"></td><td>A (F<sub>6</sub>H<sub>8</sub>N<sub>2</sub>Zr, CrOHSO<sub>4</sub>)</td><td>Si-based (20 g/L)</td></tr><tr><td>7010-B-Si80</td><td>B (K<sub>2</sub>ZrF<sub>6</sub>, KCr(SO<sub>4</sub>)<sub>2</sub>, H<sub>2</sub>O<sub>2</sub>)</td><td>Si-based (80 g/L)</td></tr><tr><td>7175-unsealed</td><td></td><td></td><td>/</td><td>/</td></tr><tr><td>7175-A-Si20</td><td rowspan="3">7175</td><td rowspan="3">TSA</td><td rowspan="2">A (F<sub>6</sub>H<sub>8</sub>N<sub>2</sub>Zr, CrOHSO<sub>4</sub>)</td><td>Si-based (20 g/L)</td></tr><tr><td>7175-A-Si80</td><td>Si-based (80 g/L)</td></tr><tr><td>7175-B-Si20</td><td rowspan="2">7175-B-Si80</td><td rowspan="2">TSA</td><td rowspan="2">B (K<sub>2</sub>ZrF<sub>6</sub>, KCr(SO<sub>4</sub>)<sub>2</sub>, H<sub>2</sub>O<sub>2</sub>)</td><td>Si-based (20 g/L)</td></tr><tr><td>7175-B-Si80</td><td>Si-based (80 g/L)</td></tr></tbody></table>
<table><caption>Table 3: Impedance data parameters obtained with calculation and graphical methods, for 7010-B-Si80 and 7175-B-Si80 before degradation.</caption><thead><tr><th></th><th>7010-B-Si_80</th><th>7175-B-Si_80</th></tr></thead><tbody><tr><td>R<sub>e</sub> (Ω.cm<sup>2</sup>)</td><td>735</td><td>245</td></tr><tr><td>R<sub>p</sub> (Ω.cm<sup>2</sup>)</td><td>5.1 x 10<sup>5</sup></td><td>1.0 x 10<sup>6</sup></td></tr><tr><td>e<sub>p</sub> (μm)</td><td>4.2</td><td>6.0</td></tr><tr><td>Θ (°)</td><td>77.8</td><td>78.2</td></tr><tr><td>α</td><td>0.86</td><td>0.87</td></tr><tr><td>Q<sub>b</sub> (F. s<sup>-1-α</sup>.cm<sup>-2</sup>)</td><td>1.1 x 10<sup>-6</sup></td><td>1.3 x 10<sup>-6</sup></td></tr><tr><td>C<sub>b</sub> (F.cm<sup>-2</sup>)</td><td>1.3 x 10<sup>-6</sup></td><td>1.5 x 10<sup>-6</sup></td></tr><tr><td>e<sub>b</sub> (nm)</td><td>26</td><td>22</td></tr></tbody></table>
<caption>Table 4: Chemical composition of the sealed oxide measured by a) EDS on exposed (15 cycles) and unexposed surface of a 7010-B-Si80.</caption>

b) XPS on the intact surface and after 5 cycles of degradation of 7010-A-Si20 and 7010-B-Si80.

<table><thead><tr><th colspan="3">a) EDS</th></tr><tr><th>At. %</th><th>Unexposed area</th><th>Exposed area</th></tr></thead><tbody><tr><td>Oxygen</td><td>70.0</td><td>68.0</td></tr><tr><td>Aluminium</td><td>28.0</td><td>29.0</td></tr><tr><td>Chromium</td><td>1.85</td><td>2.0</td></tr></tbody></table>

<table><thead><tr><th colspan="5">b) XPS, Atomic content (%)</th></tr><tr><th rowspan="2">Element</th><th rowspan="2">Species</th><th colspan="2">7010-A-Si20</th><th colspan="2">7010-B-Si80</th></tr><tr><th>Intact</th><th>5 cycles</th><th>Intact</th><th>5 cycles</th></tr></thead><tbody><tr><td>Al2p</td><td>Al(III)-O<br/>&amp;<br/>Al(III)-OH</td><td>8.5</td><td>6.4</td><td>1.2</td><td>1.1</td></tr><tr><td>Cr2p3</td><td>Cr(III)</td><td>6.9</td><td>7.6</td><td>0</td><td>0</td></tr><tr><td>O1s</td><td>O(II)</td><td>77.3</td><td>77.8</td><td>65.9</td><td>64.1</td></tr><tr><td rowspan="2">Si2p</td><td>R<sub>2</sub>-SiO<sub>2</sub></td><td>5.6</td><td>5.9</td><td>0</td><td>0</td></tr><tr><td>SiO<sub>2</sub></td><td>1.7</td><td>2.2</td><td>32.9</td><td>34.9</td></tr></tbody></table>