Introduction

Global warming and climate change are some of the ongoing severe concerns that the world is facing. The research to reduce fossil fuel consumption via enhancing the fuel efficiency of the vehicles, increasing service life, and reducing material consumption has attracted significant research attention. Therefore, the development of high-strength, lightweight, corrosion-resistant alloys became crucial to meet current technological and environmental challenges. Among most metallic materials in practice, Al alloys exhibit an outstanding strength-to-weight ratio1. However, commercially available Al alloys exhibit a tradeoff between mechanical and corrosion properties, limiting their application in many industries. Corrosion and mechanical properties of Al alloys can be controlled by appropriate processing routes along with suitable composition. For example, nanocrystalline Al alloys produced by non-conventional techniques such as sputter deposition2,3,4 and ion implantation5,6 were supersaturated solid solution and exhibited high corrosion resistance.

High-energy ball milling (HEBM), also known as mechanical, alloying is reported to cause amorphization, grain refinements, formation of non-equilibrium phases, and uniform dispersion of immiscible elements, which impart unique properties including simultaneous improvement of corrosion and mechanical properties7,8,9,10,11. For instance, Esquivel et al. reported that high-energy ball-milled alloys of Al with corrosion-resistant alloying elements (V) resulted in high strength and corrosion resistance-higher than any commercial alloys9. The microstructure of the ball-milled alloys contained fine intermetallics and matrix with high solid solubility of the alloying elements where solid solubility was several orders higher than that predicted by the phase diagrams. The corrosion resistance of the ball-milled Al alloys was high despite a non-uniform microstructure. The improved corrosion resistance of these Al alloys was hypothesized due to one or more of the following mechanisms, which are proposed based on the literature on the sputter-deposited Al alloys, ball-milled alloys Al, and stainless steels4,12,13,14.

  1. 1.

    Enrichment of the passive film with the alloying elements improved passivity and decreased passive film breakdown. The proposed mechanism is motivated by the phenomenon causing stainless behavior in Fe–Cr alloys by selective dissolution of Fe and enrichment of Cr in the form of Cr2O3.

  2. 2.

    Increased repassivation abilities of the alloys by:

    1. (a)

      Prevention of the formation of critical pit solution by lower dissolution rates in the pits and lower hydrolysis of the ions

    2. (b)

      Enrichment of the corroding alloy/electrolyte interphase with a corrosion-resistant alloying element

    3. (c)

      Redeposition of the solute ions on the cathodic particles (such as Fe containing intermetallics) and decrease the cathodic reaction

  3. 3.

    Release of the oxyanions (i.e., molybdate and vanadate) can inhibit corrosion in initiation stages and accelerate the repassivation.

  4. 4.

    Doping of the passive film with solute ions and therefore increased vacancy annihilation.

Experimental or theoretical verification of the aforementioned mechanisms has not been presented in the literature, mainly because of the unavailability of suitable test specimens and limited resolution of analytical techniques used to study the surface film. The samples produced by the cold compaction of ball-milled powder exhibits significant porosity15. Most of the techniques to produce fully consolidated alloys need exposure to high temperatures, which decomposes the supersaturated solid solution to the thermodynamically stable phases16. The complex microstructure imposes challenges in developing a mechanistic understanding of corrosion. Therefore, a technique able to produce fully consolidated ball-milled alloys without exposure to high temperatures is desirable to study the ball-milled alloys and their properties.

High-energy ball milling performed under certain specific conditions has been reported to result in fully consolidated samples in the form of coarse particles, discs, or spheres depending on the material. Such formation of larger particles is termed in situ consolidation (ISC), which has been reported for only a few materials, e.g., Zn17,18, Cu19,20,21, Fe–Cr–Ni22, Al–Pb–W23, and Al–Cr8. Synthesizing in situ consolidated nanocrystalline Al alloys and studying their properties is beneficial in studying the microstructure and the corrosion mechanisms of ball-milled alloys.

The influences of microstructural features on localized corrosion and the passive film characteristics of high-energy ball-milled Al-V alloys have been investigated herein. The validity of the above-mentioned proposed mechanisms (1–4) has been tested using various electrochemical and characterization techniques. The role of secondary phases in localized corrosion and the influence of the solute on the passive film characteristics, corrosion initiation, and propagation mechanism have been discussed.

Results

Scanning electron microscopic (SEM) and scanning transmission electron microscopic (STEM) analysis

In situ consolidated (ISC) disks of Al-2 at.%V alloy (ISC-Al2V) are shown in Fig. 1a. The diameter of the disks varied from 1 to 5 mm, and thickness was typically 0.5 mm. Backscatter electron (BSE) images and energy-dispersive X-ray spectroscopy (EDXS) analysis showed uniform dispersion of V in Al (Supplementary Fig. 1). Fe-rich particles (Fig. 1b), attributed to the abrasion of milling media during the HEBM process, are confirmed by EDXS analysis (Supplementary Fig. 1). The STEM images of the ISC-Al2V alloy are shown in Fig. 1c, d along with the selected area diffraction (SAD) pattern (Fig. 1e), which reveals a nanocrystalline structure and formation of a supersaturated solid solution of Al-V. The inter planer distance of (111) plane obtained from the SAD pattern was used to determine the lattice parameter (4.042546 Å), which is smaller than that of pure Al (4.04958 Å). Contraction in the lattice can be attributed to the incorporation of V as a solid solution. The solid solubility of the V was estimated to be ~0.89 at.% using the lattice parameters and Vegard’s law24. The estimated solid solubility of V in Al is ~105 times higher than the solubility predicted by the equilibrium phase diagram at room temperature25. The high angle annular dark-field (HAADF) images (Fig. 1c, d) revealed the distribution of the V-rich phase and the elongated grains. High magnification HAADF images were used to estimate the grain size (diagonal) and presented in Fig. 1f, which shows a wide variation in the grain size. A small fraction of grains was coarse as 0.5 microns, whereas most grains were <100 nm. STEM analysis revealed homogeneously distributed phases rich in V and appearing bright in Fig. 1c, d.

Fig. 1: Microstructural characterization of ISC-Al2V alloy.
figure 1

a BSE image of the whole samples, b zoomed-in region showing the Fe rich particles marked by the arrows, c HAADF STEM revealing the distribution of V-rich phase, d HAADF image illustrating the grain size, e SAD pattern, f grain size distribution estimated from the SAD pattern shown in (e), g XRD scans of ISC-Al2V alloy comparison with pure aluminum.

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X-ray diffraction (XRD) analysis

X-ray diffraction scans for the ISC-Al2V alloy and pure Al are shown in Fig. 1g, which showed peaks corresponding to FCC Al and Al3V. Peak broadening and decrease in peak intensity were observed in ISC-Al2V compared to pure Al. A peak corresponding to Al3V intermetallic indicates that HEBM facilitates the formation of both supersaturated solid solution and intermetallic. A peak shift towards a higher 2θ value compared to pure Al was observed in the ISC-Al2V alloy (Supplementary Fig. 2) and was attributed to the reduction of the lattice parameter caused by the formation of the solid solution of V in Al. The solid solubility of V in Al was estimated using XRD analysis (0.92 at.%), consistent with that value estimated from the transmission electron microscopic (TEM) analysis. The average grain size was calculated using peak broadening and was ~45 nm, which is close to the Lorentzian peak of the grain size determined by TEM (Fig. 1f). The formation of a nanocrystalline structure was revealed from both XRD and TEM analysis which agrees with the literature on enhancing grain refinement and formation on non-equilibrium microstructure by HEBM15,26,27.

Cyclic potentiodynamic polarization (CPP) test

Typical cyclic potentiodynamic polarization curves for ISC-Al2V and pure Al after 30 min of open circuit potential (OCP) stabilization in 0.01 M NaCl are presented in Fig. 2a. The average corrosion potential (Ecorr), pitting potential (Epit), and corrosion current density (icorr) were derived from the CPP curve fitting and are presented in supplementary Table 1. The average icorr for pure Al in 0.01 M NaCl was 0.12 µA/cm2, which is ~2 times higher than that in ISC-Al2V alloy. The Epit of ISC-Al2V was 479 mVSCE higher than that in pure Al (99.99% purity). The higher corrosion resistance of ISC-Al2V alloy, as evident from higher Epit and lower icorr could be attributed to the uniform dispersion of V in Al, extended solid solution of V in Al, and nanocrystalline structure8,10,11.

Fig. 2: Electrochemical behavior of ISC-Al2V alloy in comparison with pure Al in 0.01 M NaCl.
figure 2

a Representative CPP curves after 30 min of OCP stabilization, b Nyquist curves at OCP for various immersion times, c the equivalent circuit diagram used for EIS data fitting for both ISC-Al2V and pure Al, d Mott–Schottky curves with trend lines (blue dotted) after 30 min of stabilization.

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The formation of critical pit solution (CPS) due to the hydrolysis of the cations within the pit and migration of Cl− from the electrolyte is necessary for the pit growth, and any conditions ceasing the CPS formation would cause repassivation. The addition of V in ISC-Al2V may decrease the dissolution of the alloy to the level that a critical pit solution is not formed2. The addition of V may also increase the stability of the passive film at lower pH, which could cause repassivation inside the pits13. To comprehend this phenomenon about the influence of V, CPP was performed in 0.1 M HCl (Supplementary Fig. 3), which showed considerably lower corrosion current density and anodic current density in ISC-Al2V, confirming obstacles in the formation of CPS.

Electrochemical impedance spectroscopy (EIS) analysis

Electrochemical impedance spectroscopy measures the passive film resistance and ionic transportation through the passive film. EIS measurements were performed for 14 days at regular time intervals. The Nyquist plots for both ISC-Al2V and pure Al are shown in Fig. 2b. Both ISC-Al2V and pure Al are characterized by two-time constants, one at high frequencies and the second at low frequencies. These two-time constants correspond to the passive film and localized corrosion, respectively28,29. The experimental data fitting was carried out using Zview software and the equivalent circuit shown in Fig. 2c, where, RS is the solution resistance, CPEfilm and Rfilm represent the film constant phase element and the resistance of the film, respectively, CPEdl and Rct denote the double layer constant phase element and charge transfer resistance, respectively. A constant phase element represents a deviation from an ideal capacitor.

Both ISC-Al2V and pure Al revealed a continuous increment of the total impedance (|Z|) over the immersion time (Fig. 2b). The Rfilm and Rct of ISC-Al2V after 12 h of immersion were much higher than that in pure Al, indicating higher corrosion resistance in ISC-Al2V than that in pure Al. Both ISC-Al2V and pure Al showed Rfilm increment over the immersion time (Supplementary Table 2). Rfilm for pure Al and ISC-Al2V was almost similar after 336 h of immersion. Nevertheless, ISC-Al2V displayed a significant increment of Rct over the immersion time until 336 h. The improved Rct in ISC-Al2V could be attributed to the formation of corrosion products inside the pit that can partially block the formed pits, which could act as a barrier for ion transportation between electrolyte and substrate30, whereas the Rct after 12 h of immersion of pure Al dropped noticeably after 48 h and increased gradually over the immersion time until 336 h. The CPEfilm and CPEdl were noticeably lower for ISC-Al2V compared to that in pure Al. In ISC-Al2V, the CPEfilm was increased until 96 h and gradually declined with the continuous immersion, probably due to the passive film growth rather than the influence of hydration. The CPEdl for ISC-Al2V dropped from 12 to 96 h followed by gradual increment, whereas in pure Al, CPEfilm and CPEdl increased over the immersion time. The corrosion resistance improvement of ISC-Al2V in the long-term EIS tests is attributed to the enhancement of both Rfilm and Rct compared to that in pure Al. However, Rct enhancement in the alloy is significantly higher that Rfilm increment, which could be suggested to the enhancement of the pitting corrosion resistance with an increase in the immersion time. An increase in pitting corrosion resistance could occur because of (1) change in the surface of the alloy within the pits, i.e., enrichment of the V near electrolyte/alloy interface within the pit, (2) repassivation due to decreased cathodic activities of cathodic sites (initiation sites for pitting), (3) inability to form CPS and therefore repassivation. These factors are discussed in the following sections.

Defect concentration in the passive film using Mott–Schottky analysis

Mott–Schottky analysis of ISC-Al2V and pure Al after immersion in 0.01 M NaCl for 30 min was performed (Fig. 2d) to evaluate the passive film and understand the presence of point defects. A positive slope was noticed in both ISC-Al2V, and pure Al, which revealed that the formed passive films are n-type semiconductors. It has been widely accepted that the donors such as oxygen vacancies and cation interstitials are the dominant charge carriers in n-type passive films31,32. The average donor density of ISC-Al2V was ~(1.906 ± 0.32) × 1018 cm−3 which is ~1.66 times lower than that in pure Al. A lower density signifies higher corrosion resistance.

X-ray photoelectron spectroscopy (XPS) and secondary ions mass spectrometry (SIMS)

The X-ray photoelectron spectroscopy and time-of-flight secondary ions mass spectrometry analysis were performed on ISC-Al2V after polishing and immersion in 0.1 M NaCl solution for 0 min, 5 min, 30 min, and 2 days. XPS survey scans detected Al, O, and V. Supplementary Fig. 4 shows the survey scans, and the elemental distribution is presented in supplementary Table 3. High-resolution regional scans of O, Al, and V are presented in Fig. 3. The O 1s peak was deconvoluted into two peaks named as O 1SA and O 1SB. The O 1SA peak at lower binding energy (~532.5 eV) was attributed to the oxide contribution, and the O 1SB peak at higher binding energy (~533.7 eV) was associated with the hydroxide and/or oxygen bonded carbon contamination. Figure 3b represents Al 2p peak with two binding energies: A peak with lower binding energy at ~72.9 eV corresponded to metallic Al and the higher binding energy (~75.6 eV) associated with Al oxide/hydroxide peak33. Each Al peak was further deconvoluted into two peaks: The metallic Al peak deconvoluted into 2p3/2 (~72.8 eV) and 2p1/2 (~73.2 eV). An Al2O3 peak at ~75.3 eV and Al(OH)3 peak at ~75.9 eV were obtained from the deconvolution of the oxide/hydroxide peak. The intensity of the Al oxide/hydroxide peak increased with the immersion time, whereas the intensity of the metallic peak decreased over the immersion time; it could be attributed to the film growth on the surface. Due to a thick surface layer, the metallic Al peak was not detected after 2 days of immersion. The film thickness, calculated using the method reported in the literature34,35, was 2.97, 3.02, and 4.95 nm for 0, 5, and 30 min of immersion time, respectively. The O/Al atomic ratio was estimated from O 1s and Al 2p air-formed peaks (Supplementary Table 3). The estimated O/Al atomic ratio was 2.43, more significant than the stoichiometric value of 1.5 (i.e., Al2O3), attributed to hydroxide formation. The O/Al ratio was increased over the immersion time and reached 3.10, which is closer to the O/Al ratio in Al(OH)3. The continuous hydration of the oxide layer and formation of an external Al(OH)3 could be attributed to the increase in the O/Al ratio with immersion time. Moreover, oxides or hydroxides of V also contribute to the overall oxygen content of the passive film. Figure 3c shows the experimental V 2p peak and the four deconvoluted peaks representing four different oxidation states of V (i.e., V0+, V3+, V4+, and V5+). The intensity of the metallic V peak was highest in air-formed and gradually decreased over the immersion time. The V4+ content was the highest among all the oxidized V in all the four immersion conditions, followed by V3+ and then V5+. A gradual decrease in V5+ content was noticed over the immersion time.

Fig. 3: Surface film characterization of ISC-Al2V.
figure 3

a–c XPS regional scan of O 1s, Al 2p, and V 2p3/2 peaks following immersion in 0.1 M NaCl for various times. The spectra are shown in thin solid lines, envelope in dotted lines, and background in dashed lines. d, e ToF-SIMS negative ion depth profiles of ISC-Al2V for air-formed and 30 min of immersion in 0.1 M NaCl solution, respectively.

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The secondary ion mass spectrometry (SIMS) analysis for ISC-Al2V shows depth profiles for 18O‾, OH‾, Al2‾, AlO2‾, VO‾, and V‾. The ions Al2‾ (53.9636 amu) and AlO2‾ (58.9719 amu) corresponds to metallic Al and Al oxide, respectively36,37. The oxygen characteristic for the oxide scale is presented by the 18O‾ (17.9997 amu) since the 16O‾ (15.9955 amu) saturates the detector. The metallic and oxidized V is shown by V‾ (50.9445 amu) and VO‾ (66.9394 amu). Two regions were defined as surface film and the metallic substrate, separated by the dotted lines (Fig. 3d, e). Moreover, the metallic substrate begins once the Al2‾ curve reaches the plateau. The surface film thickness was ~2.88, 3.10, 4.06, and 35.63 nm for 0 min, 5 min, 30 min, and 2 days of immersion. 18O‾, OH‾, AlO2‾, V‾, and VO‾ were observed in the surface film in all four conditions exhibit the presence of both metallic and oxidized V in the passive film. A slight increment of V‾ ions was observed at the substrate/surface film interface in both air-formed and 5 min (Supplementary Fig. 5a); immersed samples exhibit the enrichment of metallic V at the substrate/surface film. This metallic V enrichment at the substrate/surface film interface may be a barrier to the further dissolution/oxidation of the aluminum13. After 2 days of immersion, the V enriched layer was shifted towards the surface film/electrolyte interface (Supplementary Fig. 5b). The VO‾ ion enrichment was observed near the middle of the surface film in all four conditions. This incorporation of V in the passive film and development of V enriched layer is expected to contribute to observed enhancement in corrosion resistance.

Characterization of the passive film and pits using STEM

STEM cross-section images of the passive film and pit for ISC-Al2V alloy after 30 min of immersion in 0.1 M NaCl are shown in Figs. 4 and 5. Figure 4a shows the formation of a uniform passive film with an average thickness of 5.6 nm, consistent with the passive film thickness estimated from XPS and SIMS. The EDXS elemental mapping was performed on the region in Fig. 4b (HAADF), revealing Al, O, and V (Fig. 4c–f). The O was distributed on the top of the alloy surface, confirming a uniform passive film (Fig. 4e). Figure 4f shows the enrichment of V underneath the passive film, which agrees with the enrichment of V observed at the surface film/substrate interface through SIMS analysis. The V content was estimated using an EDXS line scan along with the red arrow (Fig. 4f). The highest V content was observed underneath the passive film (Fig. 4g).

Fig. 4: Structure of the passive film of ISC-Al2V alloy.
figure 4

a STEM image showing the uniform passive film, b zoomed-in HAADF STEM image, c–f EDXS elemental mapping of the same region in (b) and, g distribution of V along the red arrow in (f) revealing the V enrichment at the oxide/metal interface.

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Fig. 5: Pit morphology of ISC-Al2V alloy.
figure 5

a STEM cross-section sample prepared by FIB revealing the pit morphology after 30 min of immersion in 0.1 M NaCl, b zoomed-in HAADF image, c–f EDXS elemental mapping of the region in (b), and g distribution of V along the red arrow in (f) revealing the V enrichment underneath pit.

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A cross-sectional image of a pit developed after immersion in 0.1 M NaCl for 30 min is presented in Fig. 5a. The pit growth was insignificant even after 2 days of immersion (Supplementary Fig. 6). The EDXS analysis was performed on the magnified cross-sectional image of the pit (Fig. 5b–f), confirming the presence of Fe-rich particles inside the pit. Fe containing intermetallics are cathodic and initiate corrosion in Al38. V was observed on the Fe-rich particles (Fig. 5f), indicating the V deposition during corrosion inside the pit and, therefore, decreased cathodic activity of Fe-rich pitting and repassivation in the early stages of pitting corrosion. The corrosion inhibition mechanism of V could be analogs corrosion inhibition of vanadate corrosion inhibitor39,40. In addition, enrichment of V under the corrosion product layer and within the corrosion products in the pit is observed (Fig. 5f). The net intensity of V was estimated along with the red arrow indicated in Fig. 5f. The V content underneath the pit was slightly higher than that in the matrix region (Fig. 5g), revealing V enrichment underneath the pit. The observed phenomenon of V deposition on the Fe-rich particles and incorporation of V in corrosion produced within the pits is confirmed by the TEM study of ISC-Al2V alloy after 2 days of immersion (Supplementary Fig. 6). The size of the pits after 30 min and 2 days of immersion was almost similar, indicating inhibition of the corrosion in the early stages of initiation.

Discussion

Both XPS and SIMS revealed the incorporation of metallic and oxidized V into the passive film developed on ISC-Al2V. This oxidized V species (i.e., Vn+ (n = 3,4,5)) in the passive film of ISC-Al2V could interact with oxygen vacancies. Therefore, the number of oxygen vacancies decreased compared to the pure Al, as observed in the Mott–Schottky analysis (Fig. 2d). The doping of solute V with oxygen vacancies present in the passive film could be attributed to the formation of complex species that have neutral and null mobility. Chloride was not detected by the SIMS and XPS. It should be noted that the incorporation of Cl to the passive film is attributed to the localized corrosion of Al41,42. The absence of chloride shows the effectiveness of V in preventing Cl incorporation to the passive film and therefore improving pitting corrosion resistance.

The enrichment of the V-rich layer underneath the passive film was observed by both TEM and SIMS. Moreover, Supplementary Fig. 7 revealed that V-rich phase became a part of the passive film instead of the dissolution during passive film formation. The observed enrichment of V underneath passive film is attributed to higher dissolution kinetics of Al than V. The incorporation of both metallic and oxidized V in the passive film and the enrichment of V underneath the film enhances the passive film stability. Consequently, the passive film resistance (EIS analysis), pitting potential, and the passive window (in the CPP curve) increase in ISC-Al2V compared to pure Al.

The schematic representation of pit initiation and repassivation in the early stages of pitting corrosion is shown schematically in Fig. 6. The pitting corrosion in ISC-Al2V initiates at the interface of Fe-rich particles and matrix. Fe-rich particles are reported to be cathodic in nature38, i.e., act as a site for the cathodic reaction (i.e., oxygen reduction reaction), and introduce an alkaline environment around them and localized dissolution of the passive film43. Subsequently, the interaction between the alloy substrate and the electrolyte caused the rapid oxidation of Al that led to the pit initiation (Fig. 6a). The hydrolysis of the metal cations causes localized acidification within the pit, and chloride ion transportation occurs toward the pit to maintain the charge neutrality inside the pit44. The enriched V layer underneath the pit (Fig. 6b) has two effects (1) low dissolution rate (Supplementary Fig. 3) and therefore inability to sustain a critical pit solution necessary for pit growth and (2) release of oxidized vanadium species (vanadates), which are well-known corrosion inhibitors39,40. The disposition of V (Figs. 5 and 6b) on Fe-rich particles is attributed to the decrement in the cathodic efficiency of the Fe-rich particles and eventually repassivation of the pit at very early stages of corrosion. The deposition of V on Fe-rich particles is a significant finding and indicates that uniform dispersion of the alloying elements that can form corrosion inhibitors could potentially cause repassivation at the early stage of corrosion.

Fig. 6: Schematic representation of the proposed corrosion and inhibition mechanisms.
figure 6

a Localized corrosion initiation around the Fe–Cr rich particle and b corrosion inhibition due to the deposition of V on Fe–Cr particle and enrichment of V under the passive film.

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In summary, ISC-Al2V alloy exhibited enhancement of corrosion resistance, significantly higher than pure Al. A combination of electrochemical and analytical techniques provided an understanding of the corrosion mechanisms. The improved corrosion resistance of ISC-Al2V alloy is attributed to inhibition of pitting corrosion by the disposition V on the cathodic iron-containing particle, the enhancement of passive film stability due to the incorporation of V species in the passive film, and the enrichment of V at the passive film/substrate interface. Mechanisms for the improved corrosion resistance presented herein are expected to facilitate the determination of corrosion-resistant microstructures, selection of alloying elements, choice of manufacturing technologies for desired microstructure, and therefore facilitate the development of high strength and corrosion-resistant Al alloys.

Methods

Alloy synthesis

ISC-Al2V alloy was produced using the Fritsch P-5/4 planetary ball mill. aluminum powder (purity 99.7% with 0.09% of iron, and 0.03% of silicon and mesh size − 50/+100) and vanadium powder (purity 99.8% and mesh size 325) were weighted inside the glove box with a high purity argon environment (oxygen <25 ppm). The steel balls with 10 mm diameter were loaded into the jar. The ball-to-powder weight ratio was kept at 16:1. The weighted alloy powders (aluminum and vanadium) with the process controlling agent (1.5 wt% Stearic acid) were filled into the jar and sealed inside the glove box. High-energy ball milling was performed at 350 RPM for 100 h. The milling was paused for 30 min after each 1 h of milling to avoid overheating.

Characterization

Scanning electron microscopy and energy-dispersive X-ray spectroscopy

The SEM was performed using a field emission scanning microscope FEI Verios 460 L at 20 kV accelerating voltage. The sample was grounded up to 1200 grit using SiC papers. The fine polishing up to 0.05 μm was carried out using colloidal silica suspension followed by ultrasonication in ethanol for 10 min BSE images and EDXS elemental mapping for Al, V, Fe, and O were performed on the sample to understand the microstructure and the elemental distribution.

Scanning transmission electron microscopy

The electron microscopy and spectroscopic analysis were conducted using a Thermo Scientific Titan aberration-corrected STEM operated at 200 kV. HAADF and bright-field (BF) images were obtained for ISC-Al2V alloy samples after 30 min and 2 days of immersion in 0.1 M NaCl. The samples were polished up to 0.05 μm, followed by ultrasonic cleaning. The focused ion beam (FIB) technique was used to prepare the cross-sectional TEM samples. An e-beam deposition of 300 nm Pt layer on the selected region followed by an ion beam deposition of 200 nm Pt layer on top was used to protect the surface during FIB. The EDXS elemental mapping was carried out to analyze the Al, V, O, Pt, and Fe distribution in each sample.

X-ray diffraction

XRD was performed using the Rigaku SmartLab X-ray diffractometer with a Cu Kα radiation (λ = 0.15406 nm) filter. The XRD pattern was recorded in the \(2\theta\) range from 30 to 85° with a step size of 0.02° and a scanning rate 1° per min. The average grain size was calculated using Scherrer’s formula45,46 after subtracting instrumental broadening. The reported literature referred to the calculation of grain size and solid solubility9,15,24,46.

X-ray photoelectron spectroscopy

The passive film chemical composition of ISC-Al2V was studied in 0 min, 5 min, 30 min, and 2 days of immersion in 0.1 M NaCl solution using SPECS XPS. The loose contamination was removed from the sample surface using DI water after the immersion test. The total vacuum chamber pressure was maintained below 10−9 mbar, and Mg anode was used as the X-ray source of 300 W. The experiment was performed on the scan area of 2 × 2 mm with a take-off angle of 60°. The survey scan was obtained from 0 to 1100 eV binding energy to obtain the elements present on the surface. At the same time, high-resolution spectra were performed to identify the present elements in the passive film and their oxidation states. The binding energies were calibrated with the peak of C1s at 285.0 eV. High-resolution spectra of 1s peak for C and O, 2p peak for Al and V were recorded. The spectrum analysis was conducted using CASA software, considering background subtraction and relative sensitivity factors given by Scofield33.

Secondary ions mass spectroscopy

Elemental depth profiling was carried out for ISC-Al2V in 0 min, 5 min, 30 min, and 2 days immersion in 0.1 M NaCl solution using the time-of-flight secondary ion mass spectroscopy (ToF-SIMS). The chamber pressure was maintained below 5 × 10−9 mbar in high mass resolution mode, and sputtering was performed using a Cs+ ion beam with 1 keV and 7.2 nA target current over a 100 × 100 μm raster area. The target current was restored over an area of 20 × 20 μm using a Bi3+ primary ion source with 25 keV and 0.32 pA. The negative ion profiles were recorded due to the higher sensitivity to fragments from the oxide region. The depth scale was calibrated using surface profilometry with the air-formed sample.

Electrochemical measurements

The specimens for corrosion testing were mounted in epoxy and grounded to 1200 grit SiC paper finish followed by ultrasonication in ethanol for 10 min. The sample edges were sealed by epoxy resin (Araldite) and stored in a desiccator for about 12 h before testing (This method avoided crevice corrosion during the electrochemical test). A conventional three-electrode flat cell with the Pt mech counter electrode and saturated calomel reference electrode (SCE) was utilized for the experiment. Potentiodynamic polarization, EIS, and Mott–Schottky tests were performed using VMP-300 potentiostat (Biologic) with the help of the EC-lab software. All the electrochemical tests were repeated at least three times to ensure reproducibility.

Potentiodynamic polarization

ISC-Al2V and pure Al samples were immersed for 30 min in a 0.01 M NaCl for stabilization. The anodic polarization was performed separately with the sweep rate of 0.167 mV/s starting from 50 mVSCE below the (OCP) toward the more positive potential direction. The scanning direction was reversed once the current reached 200 μA/cm2. The same procedure was followed with 0.1 M HCl instead of 0.01 M NaCl to understand the passive film stability at low pH conditions.

Electrochemical impedance spectroscopy

EIS measurements for the ISC-Al2V and pure Al were carried in a 0.01 M NaCl solution using the AC sinusoidal voltage signal with 10 mV amplitude and frequency range of 100 kHz to 1 mHz. The EIS tests were performed over 14 days. The data was collected after every 12 h.

Mott–Schottky analysis

The Mott–Schottky analysis was performed for both ISC-Al2V and pure Al. The sinusoidal voltage signal with 10 mV amplitude and 1 kHz frequency was used for both samples. The Mott–Schottky analysis was carried after 30 min of immersion in 0.01 M NaCl. The scanning was done with 80 potential steps starting from 50 mVSCE below OCP towards the more noble direction. At least three tests of measurements were carried out to ensure reproducibility.