**3. Methodology**

*3.1. Microstructure Characterization*

Microstructure characterization was performed in two areas (undeformed and bent areas) using scanning electron microscopy (SEM, Hitachi SU8000, Tokyo, Japan) in backscattered electron mode (BSE) with an accelerating voltage of 5 kV. The microstructure was observed on the cross-section of the Sheet in areas A and B shown in Figure 1b. Specimens for SEM observations were prepared by grinding samples with up to 1200-grit SiC paper and ion polishing with a Hitachi IM-4000 ion milling system. The ion milling is a damagefree process that does not cause surface deformation, eliminates stresses and oxide layers; therefore, the surface quality is good enough to observe the structure under channeling contrast using SEM. Structural investigations were combined with an energy dispersive Xray (EDX) analysis of precipitates and intermetallic particles. To more thoroughly describe the role of intermetallic particles, the density of coarse particles was calculated by dividing the number of counted particles (*N*) by the studied area (*A*) of the specimen. Moreover, their size (*A*—area) and standard deviation (*S*D) were calculated using computer software as per references [49,50].

Detailed microstructural observations of selected areas were performed using a transmission electron microscope (TEM, JEOL 1200, Tokyo, Japan) operating at an accelerating voltage of 120 kV, and the EDX spectra of the nanoparticles were obtained using a high-resolution scanning transmission electron microscope (STEM equipped with an EDX spectrometer, Hitachi S5500). The TEM samples were prepared using twin-jet electropolishing at 3 ◦C/20 V with an electrolyte composed of perchloric acid (HClO4), ethanol and 2-butyxoethanol (type A2 purchased from Struers, Ltd., Ballerup, Denmark) Based on the observations, the dislocation density (*ρ*) was calculated using the formula [51]:

$$
\rho = \frac{N}{L\_r h} \tag{1}
$$

where *N* is the number of dislocation-line intersections, *Lr* is the total length of all lines, and *h* is TEM sample thickness (*h* = 150 nm).

#### *3.2. Corrosion Tests and Characterization of Corroded Surfaces*

Electrochemical tests were performed in an aerated 0.5 M NaCl acidified to pH 3 by the addition of hydrochloric acid (HCl) as per reference [52]. The electrolyte was made

from analytical-grade reagents and distilled water. A FAS1 Gamry potentiostat (Gamry Instruments, Warminster, PA, USA) was equipped with three electrodes: Pt as the counter electrode, Ag/AgCl as the reference electrode (Ag/AgCl wire immersed in a saturated KCl solution), and the measured sample as the working electrode. The corrosion potential (*E*corr) was recorded during 3 or 6 h of immersion under open-circuit conditions. After immersion, the corrosion rate of samples was determined using linear polarization resistance (LPR), and potentiodynamic tests were conducted. LPR data were recorded over a potential range of *E*corr ± 10 mV at a scan rate of 5 mV/min. After LPR, the samples were conditioned for 3 min and potentiodynamic tests were performed at a scan rate of 5 mV/s, starting from −0.5 V/OCP to 2 V/REF at room temperature. To prevent noise, all measurements were conducted in a Faraday cage. The polarization curves were fitted using Gamry Elchem software (Gamry Instruments, Warminster, PA, USA) in Tafel mode. To ensure the reproducibility of the results, all measurements were repeated at least three times on the same samples. The surfaces of the samples after immersion under open-circuit potential were observed using a Hitachi SU8000 SEM.

### **4. Results**

#### *4.1. Microstructure Characterization*

Figures 2 and 3 show BSE images of the AA5754 alloy in an undeformed area and a deformed area, respectively. The BSE image shown in Figure 2a at a lower magnification shows a typical granular microstructure with many large, irregularly-shaped intermetallic particles randomly distributed in the alloy. The SEM-EDS analysis reveals that there are two kinds of particles: the white particles marked as P\_1 are Al, Mg, Mn, and Fe-rich particles (Figure 2a,c), and the dark grey precipitates P\_2 are composed of Mg, Al, and Si (Figure 2b,c). In the deformed region (Figure 3), the subsequent bending altered the microstructure of the material, and the distribution of the intermetallic particles also changed; more coarse white precipitates were observed in the bending area (Figure 3a). Additionally, during deformation, a significant strain was introduced into the material and resulting from this, higher dislocation density occurred [42]. Therefore, a granular structure is not clearly distinguishable in the BSE contrast image; however, many dislocations, especially multiple slip bands, are visible (Figure 3b). Similar to the undeformed area, the coarse intermetallic particles shown in Figure 3c marked as P\_1 are mainly enriched in Mg, Al, Mn, and Fe (Si was detected in the bending area), while those marked as P\_2 are composed of Al, Mg, and Si (Figure 3d).


**Figure 2.** SEM images of the microstructure of the AA5457 sheet in the undeformed area (A): (**a**), an image showing the overall distribution of intermetallic particles, (**b**), image showing the morphology and surrounding area of intermetallic particles, and (**c**), chemical analysis of the intermetallic particles P\_1 and P\_2 marked in panels a and b.

**Figure 3.** SEM images of the microstructure of the AA5457 sheet in the deformed area (B): (**a**) and (**b**), images showing the overall distribution of intermetallic particles, (**c**), image showing the morphology and area surrounding the intermetallic particles P\_1 and P\_2, and (**d**), the chemical analysis of the intermetallic particles P\_1 and P\_2 marked in panel c.

The density of P\_1 particles in the undeformed and deformed areas was comparable, *<sup>N</sup>*<sup>A</sup> = 0.01 and *<sup>N</sup>*<sup>A</sup> = 0.02 (given in 1 × <sup>μ</sup>m−2), for the undeformed and deformed areas, respectively; however, after deformation, they showed smaller sizes (Table 2). The SEM images indicate that there was an abrupt change in the density of P\_2 particles before and after deformation (compare Figures 2a and 3a). The calculations demonstrate that the density of the P\_2 particles changed to *<sup>N</sup>*<sup>A</sup> = 0.001 and *<sup>N</sup>*<sup>A</sup> = 0.01 (given in 1 × <sup>μ</sup>m−2) in the undeformed and deformed areas, respectively (Table 2). As a result of bending, the observed particles were slightly smaller in size (Table 2).

**Table 2.** The mean surface area (*A*), standard deviation (*S*D), and number per 1 μm<sup>2</sup> (*N*A) of coarse intermetallic particles P\_1 and P\_2.


SEM structural analysis was complemented by TEM observations. Figure 4 presents TEM images of the undeformed (Figure 4a–c) and deformed regions (Figure 4d–f). The coarse grains (Figure 4a) contained only a few dislocations in the grain interior (Figure 4b). TEM investigations also revealed the presence of small uniformly-distributed nanoparticles, as illustrated in Figure 4c. The observed particles have sizes from 50–200 nm and are rich in Si and Mn (the chemical composition is presented below Figure 4f).

**Figure 4.** TEM images of the microstructure of the AA5457 sheet in the undeformed area (A): (**a**,**b**), images showing the overall grain boundaries and dislocation structure, and (**c**), image showing the overall distribution of nanosized intermetallic particles. TEM images of the microstructure of the AA5457 sheet in the deformed area (B): (**d**,**e**), images showing the increase in the dislocation density and the formation of cellular dislocation structures and slip bands, and (**f**), image showing an increase in the dislocation density around nanosized intermetallic particles. The representative chemical composition of a nanoparticle is presented below panel (**f**).

The results from the deformed region clearly show an increase in the number of dislocations compared with the undeformed section. The dislocation density increased by nearly an order of magnitude, from 4.5 × 1012 <sup>m</sup>−<sup>2</sup> in the undeformed region (Figure 4b) to 4.4 × 1013 <sup>m</sup>−<sup>2</sup> in the deformed region (Figure 4d,e). This is a typical phenomenon for deformed metallic materials. It should be noted that dislocations are distributed rather uniformly; however, in some regions, the formation of characteristic cellular structures (Figure 4d) and slip bands (Figure 4e) was observed. Moreover, an increase in the dislocation density was also observed around Si and Mn-rich nanoparticles (Figure 4f). Hard intermetallic particles pin dislocations during deformation, resulting in the formation of many dislocations around such particles.

#### *4.2. Corrosion Testing*

#### 4.2.1. Electrochemical Testing

*E*corr measurements recorded in naturally-aerated 0.5 M NaCl adjusted to pH 3 are shown in Figure 5. The evolution of *E*corr in both samples exhibits a similar trend, with a rapid initial increase followed by a longer-term stabilization. The *E*corr of the undeformed sample started at −0.78 V/REF, and increased to around −0.70 V/REF after the first 5 min of immersion. It maintained a near steady-state value of −0.70 V/REF for the remainder of the experiment. The initial increase of *E*corr in the deformed area began from −0.94 V/REF and reached −0.72 V/REF after 10 min immersion. Then, a small but stable decline in *E*corr to −0.78 V/REF was observed at the end of the experiment. The initial rapid increase in *E*corr for both alloys may be related to the extensive corrosion and passivation processes [53], while the slight fluctuations observed in the *E*corr curves were attributed to the pitting process caused by the constituents formed in the alloy [54–56].

**Figure 5.** The *E*corr measurements recorded during 6 h of immersion in 0.5 M NaCl (pH 3) in the undeformed and deformed areas of AA5754. SEM observations were made after 3 and 6 h of immersion, as marked by the dashed lines.

The potentiodynamic polarization curves recorded for the undeformed and deformed samples of AA554 after 3 and 6 h of immersion under open-circuit conditions are given in Figure 6. All curves exhibit similar trends with passive regions, which suggests an anodic control of the corrosion processes. All curves exhibit a wide passive region when the current density plateaued, extended by 0.04 V in the undeformed area (both 3 and 6 h immersion), to around 0.05 V for the deformed area after 3 h of immersion, and 0.07 V for the deformed area after 6 h of immersion [57] Simultaneously, there is a shift in the pitting current, *i*p, which is the lowest for the undeformed area (after both periods of immersion). The pitting current, *i*p, shifted to higher values for the deformed areas after 3 h of immersion and continued to increase over time. Afterwards, an abrupt change in the current density was observed, which indicates the position of the pitting potential (*E*pit) [25,42,58]. The pitting potential for both undeformed samples has the same value of *E*pit = −0.68 V/REF (Table 3). Slightly more positive pitting potentials were recorded for the deformed samples, with *E*pit = −0.64 V/REF after 3 h of immersion. Increasing the immersion time to 6 h decreased the pitting potential of the deformed area to *E*pit = −0.66 V/REF (after 6 h of immersion). The numerical values of pitting corrosion resistance (*R*pit) may be described by the difference between pitting potential and corrosion potential, *E*pit and *E*corr [59,60]

**Figure 6.** The electrochemical curves obtained after 3 and 6 h of immersion in 0.5 M NaCl (pH 3) for undeformed and deformed areas of the AA5754: (**a**), potentiodynamic polarization curves, and (**b**), pitting susceptibility of analyzed materials calculated using Formula (1).


**Table 3.** The corrosion potential (*E*corr), pitting potential (*E*pit), corrosion current density (*i*corr), and corrosion rate calculated after 3 and 6 h of immersion in naturally-aerated 0.5 M NaCl (pH 3) for the undeformed and deformed areas of AA5754.

$$R\_{\rm jet} = \left| E\_{\rm corr} - E\_{\rm jet} \right| \tag{2}$$

Considering the data presented in Figure 6b, the smallest difference between *E*corr and *E*pit was calculated for the undeformed area of the sample for both immersion periods (*R*pit = 0.16 after 3 h of immersion, and *R*pit = 0.14 after 6 h of immersion), suggesting that the undeformed area is less-resistant to localized corrosion.

Regardless of immersion time, the deformed area of the alloy showed the greatest *R*pit = 0.22V. The higher *R* and *i*p values of the deformed areas were attributed to the weakened passivation properties [59,60].

The extrapolated from Tafel plot data show corrosion potentials of *E*corr = −0.84 V/REF and *E*corr = −0.82 V/REF for the undeformed areas immersed for 3 and 6 h, respectively. Corrosion potentials registered for the deformed areas had slightly lower values than the corrosion potentials of the undeformed areas, with values of *E*corr = −0.88 V/REF and *E*corr = −0.86 V/REF after 3 and 6 h of immersion, respectively (Table 3). To further view the overall corrosion behavior of the analyzed materials, LPR data were recorded after 3 and 6 h of immersion. The corrosion rates after 3 h of immersion were 1.8 mpy for the undeformed region, and almost twice as high (2.6 mpy) for the deformed region of the material (Table 3). After 6 h of immersion, the corrosion rate in the undeformed area increased slightly to 2.6 mpy, and the corrosion rate of the deformed area reached a significantly higher value of 9.5 mpy.

#### 4.2.2. SEM Observations after Immersion

A comparison of the *E*corr measurements under open-circuit conditions and surface analysis can be used to characterize the effect of immersion time on the microstructuredependent corrosion processes. The corrosion morphologies of the undeformed and deformed areas after immersion in 0.5 M NaCl at pH 3 are presented in Figures 7 and 8, respectively. The SEM images in Figures 7 and 8 clearly show that the corrosion of AA5754 is a complex process that involves several microstructure-dependent corrosion mechanisms. At the beginning of the experiment, the main corrosion attack on both the undeformed and deformed samples occurred around the Si-rich intermetallic particles (marked as P\_2 in Figures 7a and 8b), while Fe-rich particles remained unreactive (marked as P\_1 Figures 7a and 8b). The local corrosion attack around the Si-rich constituents led to cavitation around the particles, as previously observed [61]. Moreover, SEM investigations clearly indicate that after 3 h of immersion, trenching was more intense around the particles on the deformed area (compare Figures 7a and 8b). In both cases, trenching became more severe over time (Figures 7d and 8d), and when two P\_2 particles were relatively close to each other, trenching spread from one particle to the other, as shown in Figure 7d.

**Figure 7.** Post-corrosion observations of AA5754 in the undeformed area (A) after immersion under open-circuit conditions. Images taken after 3 h of immersion showing: (**a**), the propagation of crystallographically-grown pits, (**b**), trenching around Si-rich (P\_2) particles, and (**c**) high-magnification image of tiny pits. Images taken after 6 h of immersion showing: (**d**), the propagation of crystallographically-grown pits, (**e**), trenching around Si-rich (P\_2) particles, and (**f**) the formation of tiny pits.

**Figure 8.** Post-corrosion observations of AA5754 in the deformed area (B) after immersion under open-circuit conditions. Images taken after 3 h of immersion showing: (**a**), the propagation of crystallographically-grown pits, (**b**), the trenching around Si-rich (P\_2) particles, and (**c**), a high-magnification image of tiny pits. Images taken after 6 h of immersion showing: (**d**), trenching around Si-rich (P\_2) particles, (**e**), corrosion products around Fe-rich (P\_1) particles, and (**f**), the formation of tiny pits.

SEM observations confirmed that after immersion in 0.5 M NaCl (pH 3) trenching occurred around Si-rich particles over the entire analyzed sample surface. Besides intense local corrosion around Si-rich particles, after 3 h of immersion, highly-pronounced local pitting corrosion presented in the form of several superficial crystallographically-grown pits in the original and deformed areas of AA5754 (Figures 7b and 8a) was formed. As the experiment duration increased, more crystallographically-grown pits were observed on the original material compared with the deformed area. Crystallographically-grown pits limited by perpendicular standing walls of the crystallographic lattice have been previously observed on Al and Al alloys, and the pit growth kinetics have been previously explained [62–66]. It was also demonstrated that in acidic chloride-containing solutions, the general mode of pit propagation is corrosion tunneling [62]. Considering the deformation effect on corrosion mechanisms on AA5754 after 3 h of immersion, minor differences were observed regarding the shape of crystallographically-grown pits (compare Figures 7a and 8b), but not their amount. As the experiment duration increased, more crystallographicallygrown pits were observed on the surface of the original material, and some underwent further lateral spreading, while on the deformed area after 6 h of immersion, only a few crystallographically-grown pits were observed, suggesting that no new pits were grown on the deformed area. The locally-formed crystallographic pits on the undeformed material spread laterally on the surface, whilst the pits on the deformed sample seemed to propagate into the depth of the material. Some pits were created around Si-rich particles (type P\_2), as shown in Figure 7e and the inset in Figure 8c. In many areas of the observed samples, these kinds of pits were created after the trenches formed around P\_2 particles.

Apart from the microgalvanic coupling between Si-rich particles and the matrix, or the large crystallographically-grown pits, very tiny metastable ellipsoidal pits formed on the surface of both areas (Figure 7c,f—undeformed areas after 3 and 6 h of immersion, respectively; Figure 8c,f—deformed areas after 3 and 6 h of immersion, respectively). These pits grew larger with the experiment duration (Figures 7f and 8f); however, when comparing Figures 7c and 8c, it is visible that more and deeper pits were created on the deformed area of the sample. Moreover, after 6 h of immersion, loosely-adhered corrosion products were formed on the deformed area of the sample, especially in the locations where two Fe-rich particles (P\_1) were located relatively close to each other (Figure 8e).

### **5. Discussion**

The results of this work show that the observed trenching around Si-rich coarse particles depends on their number and distribution, and the increased amount of these particles in the deformed area resulted in more localized corrosion initiation, which was one of the reasons for its lower corrosion resistance.

As demonstrated in this work, the corrosion that occurred around Si-rich particles in the deformed area of the material was more severe than the corrosion around the same type of particles in the undeformed area of AA5457. This phenomenon is related to the increased number of particles in the material after deformation and due to their mutual interaction with the matrix, particularly particle-dislocation interactions. In our opinion, the cumulative stress occurring during deformation was sufficient to cause particle fragmentation. The increased number of hard coarse particles provided more locations where trenching could occur, and the intensity of such trenching is related to particledislocation interactions. As is commonly known, the hard coarse particles are the places where deformation-induced dislocation movement is blocked, leading to the formation of a high concentration of dislocations around them [67].

The chemical composition of the coarse particles is also important, considering their corrosion behavior in chloride-containing solution. The Si-rich coarse intermetallics observed in AA5754 also contained Mg and Al. Previous reports have clearly indicated that the rapid dissolution of Mg from intermetallics occurs during exposure, which leads to the rapid dealloying of these particles [21,24,61,68,69]. The results of previous studies [68,69] explain the mechanism of Mg and Al dissolution from the coarse particles composed mainly of Mg and Si in Al-Mg-Si alloys. The results of this research show that at the beginning of immersion, the preferential and selective dissolution of Mg occurs, and thus, the Volta potential of the particle changes, which also changes the nature of Mg-containing particles from anodic to cathodic, forming a galvanic couple with the Al matrix. The results of this work show that the increased trenching around Si-rich particles in the deformed area was a result of the cracking of these particles induced by the cumulative stresses generated during deformation. During deformation, the coarse hard particles blocked the dislocation

movement; therefore, a higher dislocation density was formed near such particles. As the locally-formed galvanic cells approach equilibrium, the cathodic particle forces the galvanic dissolution of the anodic matrix, leading to matrix corrosion. The increased dislocation density near the coarse particles promotes the corrosion reactions, leading to more intense trenching around the particles rich in Mg and Si. The reason for this is a higher dislocation density formed in the bent area of the sample, which promoted pit propagation in the depth due to their less-ordered structure [58,70]. This hypothesis is indisputably confirmed by the results presented in this work. It is also worth noting that Fe-rich particles remained unreactive throughout the entire experiment; however, their presence, especially when located relatively close to each other, may enhance corrosion product formation.

Another discrepancy that needs to be explained is that Aballe et al. [71] did not observe any relationship between the formation of crystallographic pitting and the existence of intermetallic particles, but Neetu et al. [42] claimed that pit initiation occurred at particles rich in Fe and Mn. We observed spontaneously-formed pits on the undeformed samples; however, in contrast to both works, our observations indicate that crystallographic pit growth is often initiated in the areas occluded by secondary particles enriched with Si. Moreover, this observation is supported by the fact that corrosion attack around Si-rich particles occurred before crystallographically-grown pits were formed. In a separate set of experiments, we observed trenching around Si-rich particles, but we did not observe any crystallographically-grown pits on the samples immersed for one hour in the same solution (data not published). Since pitting attack is considered to be an autocatalytic process in the active areas, it is reasonable to assume that the higher number of Si-rich coarse particles provided more sites were localized corrosion may be initiated. These places are prone to coalescence and form laterally-spreading local corrosion.

The second type of corrosion that occurred on the AA5754 alloy is microgalvanic corrosion between Si-rich nanoparticles and the alloy matrix. The observations showed that tiny pits were created around Si-rich nanoparticles, and the corrosion damage around them was more severe in the deformed area. Although the deformation-induced stress was not sufficient to change the size or distribution of the observed nanoparticles, our experiments clearly show that the nanoparticles in the deformed area with a high dislocation-stacking fault density around them promoted the formation of many deep tiny pits.

#### **6. Conclusions**

Based on the results of this work, the following conclusions can be drawn:


**Author Contributions:** Conceptualization, corrosion investigations, A.D.; surface characterization, A.D.; microstructure characterization, P.B. and A.S.; methodology, A.D. and A.S.; validation, H.G. and J.M.; investigation, A.D., A.S. and P.B.; writing—original draft preparation, A.D.; writing—review and editing, A.S.; supervision, H.G.; project administration, H.G.; funding acquisition, J.M. All authors have read and agreed to the published version of the manuscript.

**Funding:** This study was funded by the National Centre for Research and Development under grant STRATEGMED1/248664/7/NCBR/14.

**Institutional Review Board Statement:** Not applicable.

**Informed Consent Statement:** Not applicable.

**Data Availability Statement:** Data available in a publicly accessible repository.

**Conflicts of Interest:** The authors declare no conflict of interest.
