Role of gradient nanograined surface layer on corrosion behavior of aluminum 7075 alloy

npj Materials Degradation volume 6, Article number: 62 (2022) Cite this article

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Gradient nano-grained structures have been a promising technique to evade the strength-ductility trade-off in metals and alloys. Therefore, in this work, the effect of surface mechanical attrition treatment (SMAT) on the microstructure and corrosion behavior of the high-strength aluminum alloy was investigated. SMAT was performed at room temperature and liquid-nitrogen (LN2) flow conditions to generate two distinctly different initial gradient microstructures. Potentiodynamic polarization, electrochemical impedance spectroscopy, and intergranular corrosion tests were performed. Surface film characterization of untreated and treated samples was performed using X-ray photoelectron spectroscopy and time of flight secondary ion mass spectroscopy techniques. Result reveals significant microstructural changes in SMAT processed samples such as the formation of precipitates and dissolution of inherent phases. In addition, a reduced anodic dissolution rate was observed with the SMAT processed samples. Furthermore, the surface film characterization revealed a thicker oxide film with Cu and SiO2 enrichment in SMAT samples.

Gradient nanostructure (GNS), a class of heterostructure materials has gained considerable attention in the material community due to its ability to achieve a combination of high strength and ductility without altering the overall alloy composition1,2,3,4. These properties are derived from hetero-deformation-induced strengthening and strain-hardening by the synergistic interaction between hard and soft zones5,6. GNS materials have also shown promising improvement in surface-sensitive properties such as fatigue, wear, corrosion-fatigue, and corrosion behavior of materials7,8,9,10. Even with several advantages, the applicability of these materials has not reached its full potential due to limitations in the processing of bulk samples with controlled microstructure for mechanical properties6,10,11.

Gradient nano-grained structure, as its name suggests, consists of a surface structure of nanocrystalline grains whose magnitude in size gradually increases the further from the surface you are. In particular, surface nanocrystallization (SNC) with the nanograins at the surface can be achieved through various severe plastic deformation techniques such as surface mechanical attrition treatment (SMAT)12, ultrasonic shot peening13, sandblasting14, laser shock peening (LSP)15, and fast multiple rotations rolling16. Among these SMAT has been proven to be efficient at producing the smallest possible grain size at the free surface and an appreciable gradient several hundred microns deep into the bulk of the sample. This gradient forms as a result of dynamic Hertzian impacts of milling media with the sample's free surface which induce a cumulatively high strain value. Generally, the milling media utilized during SMAT are composed of different chemistry than the underlying sample to be processed17. The physics and nature of the impacts occur millions of times over thus introducing a way to transfer material. That is the repetitive high rate impacts and local elevation in temperature provide an environment in which atomic diffusion coefficients and chemical reactivity are enhanced facilitating the creation of alloys at the free surface of the sample. If engineered properly the high reactivity of the surface layer and easy diffusion of alloying elements through the grain boundaries of nanograins provides an efficient way to create a corrosion-resistant coating with favorable microstructure12,18. For instance, the temperature at which SMAT is carried out has been correlated with the degree of grain refinement that occurs at the free surface as well as the depth of gradient formed19,20,21,22,23. This is also been found to be true regarding the degree of contaminant and its mixing with the base material. In particular, SMAT at cryogenic temperature has shown a higher grain size reduction in pure copper due to a change in underlying deformation mode19. Aside from cryo-SNC, alloys prepared through other cryogenic thermomechanical processing such as cryo-rolling24 and cryo-extrusion25 have shown improved strength as well as ductility compared to their room-temperature processed counterparts. Largely, the SMAT processing has been carried out in a university setting on small-scaled simplified geometries such as flat plates. However, recently modified versions of SMAT processing equipment utilizing acoustic vibrations versus the traditional shaking of changes are allowing SMAT to be applied to tubes/cylinders and even thin wires9,26,27. Other embodiments of SMAT processing equipment have employed using burnishing techniques mounted to computer numerical controlled machines to allow even greater flexibility, potentially allowing for in-situ processing of parts2,28,29. While industrial systems such as the MELDTM technology30 have been developed to tackle the larger scale and more complex parts, the integration of such technology is still in its infant stages in industrial applications.

The corrosion behavior of nanostructured surfaces obtained through severe plastic deformation (SPD) and surface severe plastic deformation techniques have been studied in many systems, including aluminum alloys31,32,33,34, pure magnesium, and magnesium alloys11,35, titanium36, and stainless steels8,37,38. In the case of aluminum alloy, the use of stainless steel balls as a milling medium has been found to contaminate the nanocrystalline layer with iron. Given the galvanic potential difference between the two, the presence of Fe contamination significantly reduced the overall corrosion resistance34,39. Similar to aluminum alloys, pure magnesium and magnesium alloys also suffer from impurity contaminations and degradation in corrosion resistance with steel medium-based SMAT processing35,40. However, in comparatively harder alloys such as stainless steel and titanium-based alloys, where contamination is minimized, SMAT has been found to enhance corrosion resistance through the creation of a nano-grained surface layer41,42,43 within the bulk. A high density of grain boundary and triple junctions formed within the nano-grained surface was found to significantly enhance the diffusion of chromium to the surface layer in stainless steel. In titanium, processed in a similar manner, this same surface structure allowed oxygen to readily diffuse into the bulk enhancing the thickness of the oxide layer. In light alloys, in which contamination is more favorable during SMAT, contamination of the surface layer can be avoided by coating the milling medium with non-deleterious elements in order to facilitate its transference to the bulk. In the past, many studies have been performed to effectively coat metallic alloys through SMAT-based mechanical alloying44,45. In the case of aluminum 2024 and pure aluminum coated with Ni powder, second phases such as Ni3Al, Al3Ni, and NiAl were formed on the SMAT surface46,47. More such studies have been performed on other systems such as Al coated on steel48, Ti and Al coated on Ti plates49 and Al coated on magnesium alloys45. In each study, the driving force for easy atomic diffusion in defects and grain boundaries and temperature evolution during SMAT were exploited to form intermetallic compounds and coating on the surface. The morphology and thickness of the coating layer can be tuned by controlling process parameters such as impact energy, milling time, and temperature. To the author's best knowledge, no prior study has been performed with a detailed investigation of SMAT-induced inherent second phases dissolution, precipitation of second phases, and their correlation with corrosion response in an aluminum alloy.

Therefore, in the present work, a high-strength aluminum alloy AA7075 is used for SMAT processing with Al-coated stainless steel balls (Fig. 1). The SMAT was performed at the room as well as liquid-nitrogen (LN2) temperature, (sample will be called RT SMAT and LN2 SMAT afterward). A combination of microstructural characterization, electrochemical measurements, and surface spectroscopy characterizations was used to decipher the underlying governing mechanism. The understanding gained through this work suggests a clear route to further enhance corrosion-resistant of aluminum alloys for structural applications in the automotive and aerospace sectors.

a coating stainless steel milling media with pure aluminum powder in liquid nitrogen (LN2) flow environment, b SMAT of aluminum alloy samples in room temperature (RT) and LN2 flow condition, c schematic showing resulted gradient microstructural changes through SMAT.

Prior to SMAT, the optical micrograph of the cross-sectional surface shows a typically rolled microstructure with elongated lamellar grains parallel to rolling directions (Fig. 2a). After SMAT at LN2 and RT, a distinct change in cross-sectional microstructure can be observed (Fig. 2b and c). In LN2 SMAT, the elongated grains were found to be fragmented adjacent to the impact surface along with the plastic flow of grains due to the high strain rate of plastic deformation. In RT SMAT, a similar microstructural change was observed along with a significant plastic flow of grains to 100 µm from the impact surface. In both SMAT samples, the second phase particles along the grain boundaries were found to flow with deformed grains.

a Base untreated alloy, b LN2 SMAT, and c RT SMAT. Scale bars in a, b, and c indicate 100 µm.

Microstructural change on the impacted surface and across the cross-section was further characterized using Scanning electron microscopy (SEM) (Fig. 3). Figure 3a shows a collage of SMAT surfaces of the LN2 SMAT and RT SMAT conditions. The undeformed surface (Fig. 3d) shows two types of second-phase particles. The first one is a bright phase mainly composed of Cu and Fe along with Al and is related to either Al7Cu2Fe or Al23Fe4Cu intermetallic phase similar to earlier work50. Second, a darker color phase is mainly composed of Mg and Si, i.e., Mg2Si inclusion phase50. However, deformed SMAT surfaces showed an evolution of a different second phase with no trace of Al7Cu2Fe/Al23Fe4Cu and Mg2Si on the surface (Fig. 3b and c). These phases are mostly composed of Cu and Si along with Al. Elemental compositions for the second phase were listed in Supplementary Table 1 and Supplementary Fig. 1. The cross-sectional SEM micrographs of SMAT surfaces show an inhomogeneous layer of bright Al-Cu-Si phase with no trace of Al7Cu2Fe and Mg2Si second phase near the impact surface (Fig. 3e and f).

a collage of RT SMAT and LN2 SMAT surfaces; Backscattered scanning electron micrograph (BSE-SEM) of the top surface and enlarged sections of image Fig. 3a; b LN2 SMAT, c RT SMAT, and d base; Cross-sectional images of e LN2 SMAT and f RT SMAT. Scale bars in a indicate 2 mm; in b, c, and d indicate 20 µm and in e and f indicate 10 µm.

To understand the variation in grain size and evolution of second phases close to the SMAT surface, scanning transmission electron microscopy (STEM) characterization was performed and the results are depicted in Fig. 4. The LN2 SMAT sample shows nanocrystalline grains close to the SMAT surface and the grain size gradually increases with an increase in distance from the surface (Fig. 4a and b). In addition, Z (atomic number) contrast in the HAADF image indicates the presence of two bright second phase particles along the grain boundaries, and the grain boundaries were also found to be decorated by a higher Z element compared to the aluminum matrix (Fig. 4a). In the case of RT SMAT, similar nanocrystalline grains can be observed; however, the gradient in grain size was not as prominent as in LN2 SMAT (Fig. 4c and d). This can be due to the higher depth of fragmented nanocrystalline grains in the RT SMAT as observed in Fig. 2. Next, the second phase particles in RT SMAT samples were found to be smaller in size compared to LN2 SMAT samples. The energy dispersive spectroscopy (EDS) line profile and EDS mapping (Fig. 4f–h and Supplementary Fig. 2) illustrate that the bright phases are composed of either Si or Cu. The Cu-based phases are found to be higher in volume fraction than Si-based phases in both SMAT samples. Additionally, Si-based phases were found to be comparatively higher in volume fraction in the RT SMAT sample than in the LN2 SMAT sample. Besides, some of the bright second-phase particles were also found to be rich in Zn, but their volume fraction was found to be smaller than the Cu and Si-rich particles (Fig. 4f and h). Like the SEM images, the STEM image also indicates the dissolution of inherent Al-Cu-Fe and Mg2Si phases. In addition, no trace of MgZn2 (η) precipitates was observed in the STEM micrograph suggesting SMAT-induced dissolution. The average grain size of the top 500 nm layer from the impact surface was found to be 44 and 40 nm for LN2 SMAT and RT SMAT samples, respectively (Fig. 4e).

HAADF and respective MAADF image of a, b LN2 SMAT and c, d RT SMAT sample with the white arrow showing the SMAT surface, red, green, and blue arrows indicate Si-based, Cu based phase and heavier element along grain boundaries, e grain size distribution for top 500 nm depth from SMAT surfaces, f EDS line profile showing composition variation across a dark phase adjacent to a bright phase in LN2 SMAT sample, g magnified HAADF image from Fig. 4c and h respective EDS elemental mapping image of Al, Cu, Si, and Zn. Scale bars in a, b, c, and d indicate 200 nm, and in f and g indicate 50 nm.

Figure 5 shows the XRD patterns of the unimpacted base surface and the two SMAT surfaces (perpendicular to the rolling direction). In Fig. 5a, the unimpacted surface shows a comparatively higher intensity for the (311) planes compared to the (220) planes indicating a presence of initial texture in the rolled plated sample. However, with SMAT, a sharp drop in (311) plane peak intensity can be observed which indicates a nearly random orientation has been achieved on the impacted surfaces. Figure 5b shows a broadening of FWHM of the pure aluminum peaks along with a higher angle shift in peak position. The peak broadening can be assigned to the accumulation of micro-strain and grain refinement caused by SMAT. Next, in the SMAT samples, the peak positions were shifted towards a higher two-theta value indicating a combined effect of solute element saturation in the aluminum matrix and compressive residual stress. Further, between LN2 and RT SMAT samples, the latter has a higher peak broadening and two-theta shift indicating higher micro-strain and residual stress. Lastly, Fig. 5c shows the low-intensity peak for all three surfaces, where the base indicates the presence of Al-Fe-Cu and MgZn2 phases. In contrast, SMAT surfaces illustrate the presence of Al-Cu and Si and dissolution of MgZn2 and Al-Cu-Fe phases. Hence, it implies that the SMAT process has not only created a gradient microstructure but also dissolved the inherent second phases and precipitation of new ones. These microstructural changes can alter the electrochemical response to a large extent and will be characterized through electrochemical tests and discussed in the next section.

a X-ray diffraction (XRD) pattern of base untreated, LN2 SMAT and RT SMAT surfaces; b XRD peaks of three surfaces at different pure aluminum hkl planes, which showed peak full width at half maxima (FWHM) broadening and higher angle shift of peak position with SMAT treatment; c XRD peaks at low intensity of three surfaces to address the dissolution and formation second phase after SMAT.

The corrosion potential for a sample having equilibrium between anodic and cathodic kinetics (minimum current density) can be measured through an Open Circuit Potential (OCP)51. Figure 6a shows the OCP curves for the base, LN2 SMAT, and RT SMAT samples. The OCP of the base increases to a more positive value with time and stabilizes close to 0.8 VSCE after 300 s of immersion. However, SMAT samples showed a relatively stable OCP from the beginning of immersion compared to the base sample. The stabilized OCP value of the RT SMAT sample was found similar to the base alloy. However, for the LN2 SMAT sample, the stabilized OCP value is around 20 mV negative relative to the other two alloys. These changes in OCP behavior can be correlated to changes in the surface microstructure upon SMAT processing. In the case of the base, the increase in OCP (after immersion) before stabilization can be explained by dealloying and dissolution of Mg2Si and η(MgZn2) precipitate52. However, in the case of SMAT samples dissolution of η and Mg2Si phases and formation of Al-Cu-based phases have created a more stable OCP from the beginning of immersion52. Similar dissolution of the η phase due to surface severe plastic deformation has been observed in earlier studies13,53. Comparing the RT SMAT and LN2 SMAT samples, the difference in stabilized OCP can be due to the difference in the volume fraction of precipitated second phase particles and the amount of residual stress on the surface.

a OCP measurements; b Potentiodynamic polarization and c Cyclic potentiodynamic polarization curves after 10 min and 30 min of immersion in 0.6 M NaCl solution with initial near-neutral pH (6.5 ± 0.2) for base untreated, LN2 SMAT, and RT SMAT samples, respectively.

The applied potential-driven surface phenomena can be analyzed through potentiodynamic polarization and cyclic potentiodynamic polarization for all three samples (Fig. 6b and c). In Fig. 6b, the base sample shows a rapid increase in anodic current density after OCP and followed by a second breakdown potential in the curve close to 0.75 VSCE. The sudden increase in current density after OCP is related to the dissolution of the altered surface layer generated by polishing and the second breakdown potential can be co-related to the initiation of localized pitting and intergranular corrosion (IGC)54. In the case of the SMAT samples, the current density in both cathodic and anodic regimes was found to be decreased in comparison to the base sample. The LN2 SMAT sample showed the smallest cathodic current density whereas the anodic current density for the RT SMAT sample was found to be the smallest among the three alloys. Unlike the base sample, neither of the SMAT processed samples displayed a second breakdown potential after OCP indicating the underlying microstructural change due to SMAT has altered both the oxide layer formation as well as the corrosion mechanism. The equilibrium potential (Ecorr) values were found to be in good agreement with the stabilized OCP values shown in Fig. 6a. The corrosion current density (icorr) for the SMAT alloy was smaller than the base alloy, indicating that SMAT has improved the overall corrosion resistance. In cyclic potentiodynamic polarization curves (Fig. 6c), forward cycles for all alloys were similar to the polarization curves (Fig. 6b) and the curve reversed after reaching a pre-defined reverse current density (irev) of 5 mA cm−2. The potential (Esw) at which the curve reached the irev was found to be most positive for the RT SMAT sample followed by the LN2 SMAT and the base samples. However, in the reverse cycle, the protection potential (Eprot) is found to be most negative for the RT SMAT sample followed by the LN2 SMAT and the base samples. This behavior signifies that SMAT treated samples need higher overpotential (Ecorr – Eprot) to repassivate compared to base (untreated alloy) and it seems counter-intuitive with potentiodynamic polarization results where SMAT sample showed better corrosion resistance. However, it should be noted that the Esw for the RT SMAT sample is most positive, hence for the roughly same amount of charge passed (as irev = 5 mA cm−2 and same scan rate) for all samples, the RT SMAT surface can create larger pits compared to the base which may have a large number of smaller pits. Hence, larger pits in SMAT can eventually lead to a higher driving force for repassivation. A similar opposite trend in polarization and cyclic polarization results have been observed in a recent study by Zhou et al.55, where it was observed that for an Al-Zn-Mg alloy lower overpotential was needed for repassivation in higher [Cl−] solutions as compared to lower [Cl−] solutions although the former have higher susceptibility of pitting corrosion. Apart from this, an increase in the surface micro-roughness due to SMAT can also hinder the repassivation behavior.

Electrochemical impedance spectroscopy (EIS) measurements were performed at different immersion times (0.5, 2, 6, 12, 24, 48 h) to understand the time-dependent surface evolution. The measurements were plotted as Nyquist, Bode-bode, and Bode-phase angle plots for the three alloys (Fig. 7). Nyquist plots at 0.5 h for each sample were found to consist of two capacitive and an inductive loop at high, medium, and low-frequency domains. The high-frequency capacitive loop represents a double layer inhomogeneous oxide film (Al2O3) which consists of a porous thick hydrated oxide film at the oxide/solution interface and a thin continuous compact film at the metal/oxide interface56,57. Similarly, the medium to low-frequency capacitive loop represents double-layer capacitance at the film-solution interface and active metal dissolution pits58. This loop is more prominent in SMAT samples. With an increase in immersion time, the overall diameter of the Nyquist plot was found to decrease in each sample with a large depression in the second capacitive loop, which can be attributed to a weakening of a passive oxide layer with exposure to aggressive Cl− ion and the increase in metal dissolution from the oxide-free surface. Furthermore, the low-frequency inductive loop corresponds to the faradic reaction and the adsorption/desorption of charged species and their relaxation process at the interface59,60,61. Lastly, in the base alloy at 24 and 48 h a semi-Warburg type impedance behavior can be seen that can correspond to diffusion-limited resistance provided by Al(OH)3 corrosion product formed on the surface by the below-mentioned reactions (1)–(3)62,63,64:

Nyquist, bode-bode, and bode-phase angle plots for a–c Base untreated; d–f LN2 SMAT; g–i RT SMAT.

The \(\left| Z \right|\) value for all the samples can be observed in Fig. 7b, e, and h where with time, the \(\left| Z \right|\) value is found to decrease and saturate as immersion time increases. Similarly, the Bode phase angle plot (Fig. 7c, f, and i) shows a clearer representation of transition in various components over the frequency range. In between 10 and 1000 Hz, one time constant can be observed in the bode plot representing the oxide layer contribution and at 0.1–0.01 Hz there is a second time constant that represents the contribution from double-layer capacitance.

Equivalent circuits displayed in Fig. 8a were used to fit the measured impedance spectra. In the circuits Rs represents solution resistance, CPEox and CPEdl indicate oxide layer capacitance and double-layer capacitance. Likewise, Rox and Rt are resistances to ion transport in the oxide layer and charge transfer resistance on an active dissolving surface. Lastly, L, RL, and W indicate inductance and resistance corresponding to the inductor and Warburg impedance at low-frequency domains. Parameters obtained through fitting were listed in Supplementary Table 2. Variation of surface film resistance with immersion time for three samples was shown in Fig. 8b. With the increase in immersion time, the resistance values gradually drop followed by stabilization at 12 h and higher immersion time. The RT SMAT sample showed the highest resistance throughout the immersion time compared to the other two samples. Comparing the LN2 SMAT and base samples, the former was found to have smaller resistance till reaching 6 h and again at 24 h and 48 h of immersion, the trend flips. Variation of effective capacitance (Ceff) value for the surface film with immersion time was extracted from CPEf presented in Fig. 8c. The CPEf represents a non-ideal capacitance contribution from the surface film as the phase angle is not exactly −90˚ (Fig. 7c, f, and i). The value of Ceff for the surface film can be expressed by Eq. (4) first proposed by Hsu and Mansfeld65,66:

where Rf represents the surface film resistance, Qf and α are the CPEf parameters. The thickness of the surface film is inversely proportional to Ceff by Eq. (6)67:

where df represents the thickness of the surface film, ε0 is the permittivity of the vacuum and ε is the dielectric constant of the surface film. So, it can be inferred from Fig. 8c that the base sample has a thicker film compared to SMAT samples, and among LN2 SMAT and RT SMAT the latter has a relatively thicker film in all immersion times. Although the base has a thicker film, its lower film resistance makes it more susceptible to the corrosive environment67.

a Equivalent circuits used for fitting EIS curves for different immersion times; b variation of surface film resistance with immersion times obtained through EIS curve fitting; and c variation of effective capacitance with immersion time calculated by Eq. (4).

It is well known that AA7xxx series alloys suffer from IGC when exposed to a corrosive medium due to the anodic dissolution of grain boundary precipitate. To understand the effect of SMAT on IGC susceptibility, ASTM G110-based immersion testing was performed for each sample. After immersion testing, optical micrographs imaging were performed on the alloy cross-section as shown in Fig. 9a–c. For the base alloy (Fig. 9a), it can be observed that grain boundaries parallel to the rolling direction were dissolved after 24 h of immersion. Along with grain boundary, some grain dissolution can be observed in the enlarged micrograph shown in the insert of Fig. 9a. Comparing the IGC profile of LN2 SMAT and RT SMAT samples with the base sample (Fig. 9b and c), it can be inferred that grain boundary dissolution has occurred along the plastically flowed elongated grains due to the high strain rate impact of aluminum-coated stainless-steel balls. Moreover, a severe dissolution of highly deformed grains near the impacted surface was observed in the enlarged micrograph shown in the inserts of Fig. 9b and c. This can be due to the high reactivity of severely deformed nanocrystalline grains formed via SMAT processing. The IGC depths for all three alloys were measured from several similar optical micrographs and the percentage variation of IGC depth in three alloys was shown as a histogram plot in Fig. 9d. It can be observed from the plot that the RT SMAT sample has maximum IGC penetration depths of around 50–75 µm, LN2 SMAT alloy has around 100–125 µm and the base alloy varies from 75 to 150 µm. However, for the base alloy, a high fraction of IGC depth was observed above 200 µm than in SMAT samples. It can be inferred that with the SMAT process, the sacrificial dissolution of high-energy nanograins close to the impact surface and the torturous path provided by shear flowed elongated grains have minimized the average IGC penetration depth in comparison to the base sample.

Optical micrographs of an etched cross-section of AA7075 alloy after intergranular corrosion test for 24 h according to ASTM-G110 standard. a Base untreated alloy, b LN2 SMAT alloy, c RT SMAT alloy, and d A histogram plot comparing intergranular corrosion penetration depth for three alloys. Scale bars in a, b, and c indicate 200 µm, and inserts in respective images indicate 100 µm.

To understand the surface oxide layer modification through SMAT, time of flight-secondary ion mass spectroscopy (ToF-SIMS) characterization was performed on all three alloy surfaces after 15 min of immersion in 0.6 M NaCl solution at open circuit potentials. Figure 10a–c shows the general ion depth profile as a function of Cs+ ion sputtering time. All the concentration profiles were plotted on a logarithmic scale to visualize low concentration change and the variation with time represents the in-depth ion concentration variation from the top alloy surface layer. In each alloy's depth profile, three regions can be identified. The first region has a quasi-constant \(AlO_2^ -\) and \(O_{16}^ -\) the concentration represents the aluminum oxide and the oxygen concentration represents the oxide layer. The \(O_{16}^ -\) was used instead of \(O_{18}^ -\) because of the latter causes saturation of the detector. The first region extends up to the time when the \(AlO_2^ -\) and \(O_{16}^ -\) reaches the maximum. In this region, a small and gradually increasing concentration for \(Al_2^ -\) (represents metallic aluminum) for all three alloys. Cu− which represents the copper concentration revealed different behavior in base and SMAT alloys. In SMAT alloys, Cu− ion maxima were observed in the center of the oxide layer, unlike the base alloy, suggesting copper enrichment through selective oxidation of Al due to the Al-Cu-based second phase generated through SMAT. The Cu enrichment under the Al oxide film has been observed in earlier studies on Al-Cu thin films68,69. The \(SiO_2^ -\), which represents the silicon dioxide concentration shows a sharp drop in the oxide region of base alloy; however, in SMAT alloys, the concentration is steady. The Cl− and OH− represent the contribution of corrosion product and aluminum hydroxide and their concentration reduce drastically after reaching a maximum near the surface. SMAT samples showed a comparatively higher thickness (i.e., 423 and 463 s of sputtering time for LN2 SMAT and RT SMAT, respectively) relative to the base sample (104 s of sputtering time). Following the oxide layer, the interface region begins till the innermost substrate starts. The interface region thickness depends on sample surface roughness and that is the reason the SMAT alloys show higher interface region thickness as compared to the base alloy. In this region Cu− undergo a minimum in concentration for LN2 SMAT and RT SMAT before increasing again till the substrate region where the concentration stabilizes. The \(AlO_2^ -\) concentration was found to decrease in this region for the base alloy; however, for SMAT alloys the \(AlO_2^ -\) concentration is found relatively constant before it decreases from the beginning of the substrate region. In the substrate region a sharp increase in \(Al_2^ -\), stabilization of Cu− and drop-in \(AlO_2^ -\) and \(O_{16}^ -\) was observed.

ToF-SIMS negative ion depth profile after immersion in 0.6 M NaCl solution for 15 min for a Base untreated alloy; b LN2 SMAT alloy; and c RT SMAT alloy. X-ray photoelectron spectrum for all three samples after immersion in immersion in 0.6 M NaCl solution for 15 min; d survey spectrum, deconvoluted high-resolution spectrums e C 1s, f O 1s, g Al 2p; high-resolution h Zn 2p and i Si 2p.

X-ray photoelectron spectroscopy (XPS) scans were performed on three alloys after 15 min of immersion in 0.6 M NaCl solution (Fig. 10d). The main components observed were aluminum, oxygen, and carbon. To capture the presence of these elements Al 2p, O 1s, and C 1s, high-resolution spectrums were obtained (Fig. 10e–g). In addition, the presence of minor elements such as zinc and silicon were obtained through Zn 2p, and Si 2p high-resolution spectra (Fig. 10h and i). The C 1s high-resolution used as a calibration reference spectrum has four components: at 284.6 eV (C–C), 286.1 eV (C–O), 287.2 eV (C-O=C) and 289 eV (presence of CO32−) bonds, respectively (Fig. 10e). The presence of carbon comes from the surroundings as contamination70. Additionally, the C 1s, and O 1s peaks were deconvoluted into three components: 530.53 eV (O2−), 531.83 eV (presence of OH− components), and 532.73 eV (presence of water and C=O bonds) (Fig. 10f)71. In the case of Aluminum, as shown in Fig. 10g, Al 2p was designated as the Al3+ component, which suggests the presence of aluminum oxide and hydroxide (Fig. 10g)72. The chemical composition of the surface film has been tabulated in Table 1. It can be observed that the surface layer is mainly composed of O, C, and Al from which Al is found to be mostly present as aluminum hydroxide and C as contamination from the environment. In addition, the RT SMAT sample shows the absence of a Zn peak indicating the dissolution of η precipitates (Fig. 10h). Surprisingly there was no indication of Cu as a component on the surface layer (Table 1), which may be due to the presence of copper in the inner layer of the film and was observed during sputtering in TOF-SIMS characterization (Fig. 10b and c). Lastly, in SMAT samples, RT SMAT (Fig. 10i) sample shows a comparatively higher fraction of the Si component which can be attributed to a higher fraction of the Si phase compared to LN2 SMAT.

To understand the onset of localized corrosion and its propagation with time, sample surfaces were characterized after immersion in 0.6 M NaCl solution for 6 h and 24 h. Towards this, the same region was tracked for each sample after different immersion times and the results are presented in Fig. 11. Note that the micrographs in Fig. 11 for each sample are magnified images from Supplementary Fig. 3. As prepared base sample (Fig. 11a) shows the presence of bright Al7Cu2Fe and dark Mg2Si intermetallic phase. After 6 h of immersion, the localized dissolution of the aluminum matrix adjacent to the Al7Fe2Cu phase was evident, which is highlighted using red arrows as shown in Fig. 11b. In addition, dissolution of matrix around nanoscale precipitates decorated along grain boundaries was also observed (shown in green arrows). Furthermore, the Mg2Si phase undergoes selective dissolution of Mg73, leaving behind Si which further undergoes passivation to form SiO2 as detected in SIMS ion depth profile (Fig. 10a)74. In contrast, in SMAT specimens, processing-induced surface roughness can be observed with the distribution of sub-micron to few microns range Al-Cu based phases, shown by red arrows in Fig. 11d and g. After 6 h of immersion, the localized dissolution around these cathodic phases was observed in SMAT conditions (Fig. 11e and h). After 24 h of immersion, the deposition of white corrosion products was found around the matrix and over the second phase particles. Deposition of corrosion product with time provides a hindrance effect towards transport of ionic species towards and away from the corroding sites (Fig. 11c, f, and i). Moreover, in the base sample, the intergranular and intragranular matrix dissolution and dealloying of Al7Cu2Fe was observed (Fig. 11c).

Scanning electron micrographs of base a–c, LN2 SMAT d–f, and RT SMAT g–i alloy after immersion in quiescent 0.6 M NaCl solution for different times (shown in corresponding figures). Red, green, and yellow arrows represent the cathodic phases, precipitates at the grain boundary, and corrosion products, respectively. Scale bars in a–i indicate 5 µm.

The experimental results clearly showed that SMAT processing in different environments introduced a large amount of plastic deformation into the base aluminum 7075 alloy surface, which significantly altered the surface microstructure and improved the corrosion resistance of the aluminum alloy. In both of the SMAT samples, an ultra-fine layer of grains was present close to the impact surface and followed by a region of shear flowed elongated grain. In addition to grain size refinement, dissolution of inherent second phases (Al7Cu2Fe/Al23Fe4Cu and Mg2Si), and precipitation of Al-Cu, Si, and Zn phases were observed near the impact surface. The SMAT-treated samples exhibited increased corrosion resistance due to the combined beneficial effect of grain size refinement, absorption, and second-phase precipitation. In the below subsections, microstructure evolution and the observed improvement in corrosion resistance will be discussed separately.

It is well known that the η (MgZn2) phase is the primary hardening precipitate found in Al-Zn-Cu-Mg-based alloys. This strengthening phase is distributed within the matrix and along the grain boundaries75. In the current work, the base alloy was in peak-age (T6) condition, hence the aluminum matrix will have η’ incoherent precipitate and the presence of MgZn2 was confirmed using XRD (Fig. 5c). However, in the SMAT processed samples, the absence of the relevant XRD peaks and the absence of observation from TEM analysis confirmed that SMAT processing has effectively dissolved/solutionized η’ precipitates back into the matrix. Similar, fragmentation and dissolution of precipitates through SPD processing have been observed in Al-1.7 at% Cu alloy76, 7055-T7753, and 703477 alloys. In addition to the dissolution of η’ precipitates, the surface close to the impact surface was also found to be deprived of Al7Cu2Fe and Mg2Si intermetallic phases (Fig. 3). This indicates that the high strain rate impact of aluminum-coated steel balls has also fragmented and dissolved these phases near the impact surface. Hence, the dissolution of multiple phases must have created a supersaturated solid solution of Zn, Mg, Cu, Si, and Fe in the aluminum matrix.

Precipitation of AlCu, Al2Cu(θ), Si, and Zn from the supersaturated solid solution is interesting. Similar phenomena have been observed in the altered surface layer (ASL) of AA7055 alloy after surface abrasion53. In the same study, Wang et al. suggested the following mechanism for the precipitation of AlCu and θ phases. The supersaturated solid solution in ASL after the dissolution of η and η’ precipitates was shown to drive the nucleation of AlCu, θ, and Zn phases through vacancy, dislocations, and grain boundaries assisted diffusion of solute atoms. These precipitates later coarsened at 200 nm in length after natural aging for 42 months. In contrast, in this work, θ precipitates of 2–4 µm in diameter were observed after 30 min of SMAT processing (Fig. 3). In addition, precipitation of pure Si (Fig. 4a and c) was observed in both RT SMAT and LN2 SMAT samples, which to the author's best knowledge has not been observed in the aluminum 7075 alloy reported earlier. This suggests that heavy plastic deformation through the continuous impact of high-velocity aluminum-coated stainless steel balls has not only increased the amount of non-equilibrium defects but also increased the localized surface temperature, which has enhanced the kinetics of precipitation significantly. The preferential formation of pure Si, Al-Cu-based, and Zn phases can be described by the interaction tendency of solute atoms towards the vacancies. The vacancy binding tendency of each solute atom can be arranged in following order Si (0.08 eV) > Zn (0.03 eV) > Cu (0.02 eV) > Fe (∼0.00) > Mg (−0.02)78. Being most favorable Si and Zn were found to form pure elemental phases. Vacancy-mediated diffusion mainly helps to transport solute atoms from the grain interior to the nearby dislocations and grain boundaries.

Collectively, polarization and EIS measurement indicate that with SMAT processing, specifically in RT SMAT, there is a reduction in anodic and cathodic current density and an increase in film resistance. Generally, in the aluminum alloy, the underlying interfacial mechanism for corrosion can be summarized into a series of steps, i.e., adsorption of Cl− ion onto the oxide layer79, the interaction of Cl− with oxide layer80,81, exposure of substrate and formation of metastable pits, and finally metastable pit converting to stable pit depending on pit environment82. So, an attempt to correlate the surface microstructure of each specimen with above mentioned interfacial mechanism was carried out to understand the possible cause of improved corrosion resistance with SMAT.

In the base specimen, the surface microstructure consists of second phases such as Al7Cu2Fe, Mg2Si, and MgZn2 along with solute elements in the matrix. This microstructure dictates the alloy's pitting and intergranular corrosion susceptibility54,83,84. The polarization curve of the base sample showed two breakdown potentials similar to earlier work by Wang et al.83. The two-breakdown potentials are due to the difference in solute content, specifically (Cu and Zn) between the ASL and the underlying matrix. In EIS measurement, after 0.5 h of immersion presence of an inductive loop and the absence of a second capacitive loop indicates the localized breakdown of an oxide film and initiation of anodic dissolution (Fig. 7a). With further progress in corrosion, the overall diameter of a capacitive loop in the Nyquist plot was found to decrease and the pitting corrosion in this specimen can be observed from post 6 h of immersion microstructure (Fig. 11b). Additionally, the dissolved Al3+ ions can interact with OH− ions to precipitate Al(OH)3 product layer, which was confirmed through the SIMS depth profile (Fig. 10a), XPS surface spectrum (Fig. 10f), and post corrosion surface characterization (Fig. 11c). At higher immersion times, the higher volume of these corrosion products (Fig. 11c) can provide resistance to the mass transport of ions and its response can be visualized by Warburg impedance in the Nyquist plot (Fig. 7a).

In the case of RT SMAT surface grain size refinement, precipitation of pure Si and Al-Cu-based phases and dissolution of inherent phases were observed. The average grain size of SMAT samples close to the impact surface is ~30 nm which could enhance the reactivity of oxygen with surface elements by providing nucleation spots and diffusion pathways for elemental and ionic species. It can be correlated to a higher thickness of the oxide layer in the SMAT samples compared to the base sample observed through the ToF-SIMS ion depth profile (Fig. 10a–c). Similarly, through equal channel angular pressing, Ralston et al.12 have achieved an average grain size (refinement) of 125 µm in pure aluminum, which was shown to improve the surface reactivity and ion diffusion through higher grain boundary density. Furthermore, oxidation of pure Si phases in SMAT sample has shown to form SiO2 in the oxide layer and found to be more stable in concentration in RT SMAT sample through ToF-SIMS sputtering (Fig. 10c). SiO2 with lower pH of zero charges (pHpzc~2.0)85 can reduce the overall pHpzc of Al2O3 (~9.5)85 and can slow down the adsorption of Cl− ion on the oxide surface. A thicker and compound oxide layer in RT SMAT sample can contribute to the observed reduction in anodic kinetics and higher CPE1 and CPE2 value in the Nyquist plot for 0.5 h. Next, considering the effect of the dissolution of inherent second phases and the formation of different phases, their size, distribution, and galvanic coupling propensity will affect the cathodic activity of these phases86,87,88. In the base alloy, considering the corrosion potential w.r.t. standard calomel electrode (SCE), Al7Cu2Fe (−654 mVSCE) will act as a cathode, whereas Mg2Si (−1536 mVSCE) and MgZn2 (−1095 mVSCE) will act as an anode for 7075 aluminum matrix (~−800 mVSCE)87. Similarly, in the case of SMAT alloys, Al-Cu phases (−695 mVSCE) and Si (−452 mVSCE) will act as cathode and Zn (−1098 mVSCE) as anode for the aluminum matrix87. It can be noted that after SMAT, Fe and Mg were mostly found retained within the matrix as a solid solution due to a lower vacancy binding tendency as compared to other elements. Among Al7Cu2Fe and Al-Cu phase, later will act as a weaker cathode towards the matrix. However, Si which is a stronger cathode compared to Al7Cu2Fe and Al-Cu will stay inert in a near-neutral solution and can form amorphous SiO2 as suggested by Kairy et al.74. The presence of SiO2 in the SMAT sample was confirmed through the ToF-SIMS ion depth profile (Fig. 10b and c). Furthermore, the size of second phases on SMAT surface were found to be in nanometer range which will reduce their cathodic activity by the ‘small cathode and large anode’ mechanism86,89. For instance, Osorio et al. have found that fine and homogenous Al2Cu (θ) phases with fine dendritic spacing can provide galvanic protection to the aluminum matrix90. Hence the dissolution of Al7Cu2Fe phases with precipitation of small Al-Cu-based phases can be attributed to a reduction in cathodic kinetics in SMAT samples.

In conclusion, SMAT of Al 7075-T6 with aluminum-coated steel balls was performed at room and LN2 temperatures to obtain a gradient nanostructure surface layer. The below-listed points are the key finding from this work:

Microstructural characterization and XRD analysis of the SMAT surfaces revealed a nano-grained surface layer with precipitation of AlCu, Al2Cu (θ), Si, and Zn phases. In addition, dissolution of inherent phases, i.e., Al7Cu2Fe, Mg2Si, and MgZn2 close to SMAT surface was observed.

Potentiodynamic polarization showed a reduction in the anodic and cathodic kinetics in SMAT processed samples relative to the base sample. Among the three conditions, the RT SMAT exhibited the lowest anodic dissolution rate.

EIS measurements for two days of immersion revealed that the RT SMAT has the highest polarization resistance among all samples tested under all immersion times. In addition, the SMAT processed sample showed a comparatively stable surface film in lower immersion times relative to the base alloy.

Surface film characterization through the TOF-SIMS revealed a thicker oxide film with Cu and SiO2 enrichment in SMAT processed samples due to the presence of Al-Cu phases and oxidation of Si.

Overall, the understanding gained through this work suggests a clear route to further improvement of corrosion-resistance of aluminum alloys and the design of corrosion-resistant light alloys for structural applications in the automotive and aerospace sector.

In this work, an aluminum rolled plate from a Midwest steel and aluminum company was used with a T651 temper condition. Samples were sectioned into circular disks from the rolled plate with a thickness of 6 mm and a diameter of 50 mm. The thickness of the disk was parallel to the rolling direction. The composition of the as-received sample is provided in Table 2. Samples were polished through 1200 grit SiC paper followed by cleaning in an ultrasonication bath for 5 min and air-dried before SMAT.

SMAT was carried out using a SPEX 8000M high-energy ball milling machine. The prepared samples were used as an outer lid for the SPEX stainless steel vial along with 50 g of 6.35 mm diameter 440C stainless steel balls. A high-purity argon glove box was used to open and close the vial for each treatment. SMAT was performed at room temperature (RT SMAT) and inside a liquid nitrogen-flowing chamber (LN2 SMAT) for 30 min. A SMAT duration of 30 min was chosen following earlier reported works39. To avoid contamination of iron from the surface of the milling media and the inner wall of the vial, these surfaces were coated with a layer of pure aluminum powder. This coating was achieved through cryogenic milling of 0.5 g of pure aluminum powder, 325 mesh size (Alfa Aesar) for 1 h with 5 min of the interval after every 15 min of milling. Resting in between the runs was used to avoid an increase in temperature that can lead to the agglomeration of aluminum particles. RT SMAT was performed by loading the vial into the SPEX 8000M mill, whereas for LN2 SMAT the vial was inserted into a Teflon sleeve with the inflow and outflow of liquid nitrogen from the dewar.

Optical microscopy was used to visualize the effect of SMAT on cross-sectional surface grain structure. Disks were sectioned perpendicular to the SMAT surface and polished to a 0.05 µm diamond finish. Keller's reagent (95 ml H2O, 2.5 ml HNO3, 1.5 ml HCl, 1 ml HF) was used to etch the polished surface. FEI-XL30 SEM equipped with EDS was used to characterize the microstructural changes with SMAT through the characterization of the SMAT surface and the plane perpendicular to it. Samples for TEM characterization were lifted out from the SMAT surface using a Focused Ion beam Helios 5 UX and eventually thinned to electron transparency. STEM images were captured at various magnifications using the aberration-corrected ARM 200F (JEOL) to assess the microstructure, such as grain size and second phases, close to the SMAT surface of RT and LN2 SMAT specimen. More than 200 grains were sampled to generate an average grain size for each condition using ImageJ software. XRD measurements were performed using Aeris Panalytical X-ray Diffractometer.

A Gamry 600 Potentiostat connected to a standard three-electrode cell was used to perform each electrochemical experiment. SMAT and base sample surfaces with an immersed area of 1 cm2 were used as the working electrode. Platinum wire and SCE were used as counter and reference electrodes respectively. Electrochemical measurements were performed at ambient temperature (i.e., ~23 ± 2 °C) in a naturally aerated 0.6 M NaCl solution with an initial pH of 6.5 ± 0.2. The desired pH was obtained by adding dilute HCl and NaOH solutions. Potentiodynamic polarization (PP) measurements were measured after 10 min of immersion in open circuit potential (OCP). PP measurements were performed from −100 mV (vs OCP) to −700 mVSCE potential with a scan rate of 0.167 mV s−1. Cyclic polarization measurements were performed after 10 min of OCP and all the scans were started at −50 mV (vs OCP) and reversed after it reached 5 mA cm−2 current density. Scans were stopped once it reaches the cathodic branch in the reverse cycle. Similarly, EIS measurements were performed after immersion in OCP conditions for different duration i.e., 0.5, 2, 6, 12, 24, and 48 h. Measurements were performed with a sinusoidal voltage of 10 mV over a frequency range of 100 kHz to 10 MHz. Tafel plots and EIS plots were analyzed by EC-Lab software. IGC tests were performed according to ASTM G11091, with a slight modification for health and safety precautions. Samples were pre-treated in 50 ml nitric acid (HNO3) + 950 ml deionized water for 1 min at 93 °C followed by 1 min in concentrated HNO3 at room temperature. Samples were then cleaned with deionized water and dried with compressed air. All the samples were then immersed in a solution of 57 g NaCl + 10 ml of H2O2 + 990 ml of deionized water for 24 h. After immersion samples were thoroughly cleaned and sectioned and the cross-sections were polished and etched to reveal the microstructure. The propensity of IGC for each sample was calculated for the IGC depth from 30–50 fissures using ImageJ software. Quasi in-situ surface characterization was performed by analyzing the same region of the three samples after performing immersion in unbuffered 0.6 M NaCl solution with initial pH of 6.5 ± 0.2 for 6 h and 24 h.

ToF-SIMS spectrometer (Ion-ToF IV GmbH, Münster, Germany) was used and all the measurements were performed under ultra-high vacuum conditions (10−9 mbar). 25 keV Ga+ ion was used as the primary source and 1 keV Cs+ was used for sputtering. 100 * 100 µm2 was analyzed from a sputtering area of 350 * 350 µm2. The negative polarity ions were chosen for analyzing the surface. X-ray Photoelectron spectroscopy measurements were performed using Kratos Axis 165 Ultra X-ray Photoelectron Spectrometer. Monochromatic AlKα (1486.6 eV) was used and all the runs were carried out in 10−9 torr vacuum conditions. The analyzer pass energy was 100 eV for survey spectra and 20 eV for high-resolution spectra. The step size for wide spectrum was 1 eV and 0.1 eV for high resolution spectra. The photoelectron take-off angle (the angle of the surface with the direction in which the photoelectrons are analyzed) was 90° and a 400 µm spot size was used for each of the scans. Data processing (peak fitting and decomposition) was performed with the Casaxps software using iterative Shirley-type background subtraction and Gaussian/Lorentzian peak shapes.

The data that support the findings of this study are available from the corresponding authors upon reasonable request.

No custom codes or algorithms have been used in this paper.

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V.B., Y.K., and K.S. acknowledge the use of facilities within the LeRoy Eyring Center for Solid State Science at Arizona State University. This work was supported by US Army Research Laboratory under contract W911NF-15-2-0038.

School for Engineering of Matter, Transport, and Energy, Arizona State University, Tempe, AZ, 85287, USA

V. K. Beura, Y. Karanth & K. Solanki

Weapons and Materials Research Directorate, Aberdeen Proving Ground, Aberdeen, MD, 21005, USA

K. Darling

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K.S. and K.D. develop the idea. V.B. carried out processing and experimental studies including SMAT processing, SEM characterization, corrosion measurements (EIS, OCP, PP), and spectroscopy measurements (XPS and ToF-SIMS). Y.K. performed TEM experiments. All authors contributed to drafting and revising the paper. K.S supervise the work.

Correspondence to K. Solanki.

The authors declare no competing interests.

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Beura, V.K., Karanth, Y., Darling, K. et al. Role of gradient nanograined surface layer on corrosion behavior of aluminum 7075 alloy. npj Mater Degrad 6, 62 (2022). https://doi.org/10.1038/s41529-022-00271-z

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Received: 07 March 2022

Accepted: 05 July 2022

Published: 26 July 2022

DOI: https://doi.org/10.1038/s41529-022-00271-z

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