Institute for Advanced Materials and Technology, University of Science and Technology Beijing, Beijing 100083, China
2)
Southwest Institute of Technology and Engineering, Chongqing 400039, China
3)
National Center of Analysis and Testing for Nonferrous Metals and Electronic Materials, Guobiao (Beijing) Testing & Certification Co., Ltd., Beijing 101407, China
Understanding the influence of purities on the electrochemical performance of pure aluminum (Al) in alkaline media for Al–air batteries is significant. Herein, we comprehensively investigate secondary phase precipitate (SPP)-induced localized corrosion of pure Al inNaOH solution mainly based on quasi-in-situ and cross-section observations under scanning electron microscopy coupled with finite element simulation. The experimental results indicate that Al–Fe SPPs appear as clusters and are coherent with the Al substrate. In alkaline media, Al–Fe SPPs exhibit more positive potentials than the substrate, thus aggravating localized galvanic corrosion as cathodic phases. Moreover, finite element simulation indicates that the irregular geometry coupled with potential difference produces the non-uniform current density distribution inside the SPP cluster, and the current density on the Al substrate gradually decreases with distance.
The development of renewable energy technologies with less CO2 release to replace traditional fossil energy to cope with global climate change has attracted considerable attention [1–4]. The efficient conversion of chemical energy into electricity via electrochemical reactions has been regarded as a promising approach in this area. Metal–air batteries (MABs), which have simple configurations and high energy densities, exhibit great potential among various sustainable energy resources. Being abundant, recyclable, environment-friendly, and economical, aluminum (Al) is an acceptable anode material for MABs [5–9]. Nevertheless, the relatively low coulombic efficiency induced by the severe corrosion of Al anode in alkaline electrolytes is accompanied by the generation of a large amount of hydrogen, which considerably limits the practical applications of Al–air batteries [10–12].
To minimize the dissolution loss of Al anode during standby, high purity grade cast Al plates (i.e., 4N, 99.99% or 5N, 99.999%) without impure elements, such as Fe, are usually employed for commercial Al–air batteries. However, the high expense of 4N and 5N Al seriously hinders the scale-up of the products in the battery field despite the aforementioned merits of Al–air batteries. As a consequence, the development of high-efficiency Al anode with slightly lower purity grades to replace the higher purity grades, which rely on the comprehensive understanding of impurity-induced effects on the corrosion mechanism of the Al, is of vital importance [13]. Thus far, several studies have been conducted to analyze the electrochemical corrosion behavior of Al anode in strong alkaline media; notably, Al with lower purity tends to suffer from a more drastic corrosion process ascribed to the depolarization effect caused by Fe-bearing impurities [14–17]. Moreover, several researchers demonstrated that a higher purity could lead to the enhancement of the operating voltage and discharge efficiency. However, the inherent relationship between the microstructure of Al and its electrochemical behavior has not yet been completely established. In particular, early-stage localized corrosion mechanisms triggered by secondary phase precipitates (SPPs, i.e., Al–Fe and Al–Si–Fe phases) remain unclear because of the lack of sufficient evidence on the in situ morphology evolution observation in the region adjacent to the precipitates, which can shed light on the long-term corrosion fundamentals of precipitate-induced corrosion.
Localized corrosion induced by impurities on Al alloys has been widely investigated [18–20]. Notably, the impurities in such alloys tend to initiate localized corrosion because of the geometric and physiochemical inhomogeneities in the local regions. The most common phenomenon is galvanic corrosion triggered by the potential difference between the impurities and the substrate. For instance, the T1 phase in Al–Li alloys preferentially dissolves and causes intergranular and exfoliation corrosion [21]. Similarly, Mg2Si precipitates in 7xxx series Al alloys act as anodic phases and become the origination sites for pitting corrosion [22]. By contrast, cathodic precipitates, such as Al7Cu2Fe, Al2Cu, Al2CuMg, and Al3Fe, trigger localized corrosion on the Al alloy matrix because of the galvanic coupling between them [23–27]. Except for the potential difference, the impurity-induced localized corrosion can be influenced by the size and morphology of the impurity particles. Currently, most of the previous works mainly focus on localized corrosion driven by precipitate particles, where the interface between the matrix and the impurity particles before being corroded is ascribed to the crevice or galvanic effects [28–29]. Nonetheless, the mechanism of localized corrosion induced by precipitate clusters is rarely reported. Furthermore, the previous studies are centered on precipitate-induced corrosion when the materials are subject to natural environments, such as the atmosphere or seawater. The corresponding corrosion behavior in strong alkaline media and its mechanism remain unexplored.
Herein, we comprehensively analyzed the effect of SPP clusters on the initial corrosion behavior of pure Al in alkaline media utilizing multiple technologies. The microstructural characteristics of SPPs were initially investigated, followed by the exploration of the elemental composition and other physical properties. The corrosion behavior of SPPs was further determined based on the quasi-in-situ scanning electron microscopy (SEM) observations on one typical SPP cluster. Furthermore, additional information, including the cross-section view of the SPP cluster and the overall surface analysis, was provided to determine the localized corrosion mechanism induced by the SPP clusters. The corresponding schematic, which can be used to develop more effective pathways to reduce SPP-induced corrosion on pure Al in alkaline media, was also proposed at the end of the paper.
2.
Experimental
2.1
Materials and surface preparation
The materials used in this work are pure Al ingots with different purity grades of 99.70%, 99.85%, 99.95%, and 99.99%. Their chemical compositions were determined by inductively coupled plasma optical emission spectroscopy and listed in Table 1.
Table
1.
Chemical compositions of the experimental materials wt%
Specimens with sizes of 10 mm × 10 mm × 6 mm were mechanically ground with silicon carbide papers down to 4000 grit and polished with a 1 µm diamond paste. Then, the specimens were ultrasonically rinsed in ethanol. To avoid the possible influence of the mechanical polishing process, several specimens were prepared with the ion polishing preparation procedure. The specimens were cut into slices of 0.4 mm, manually ground to a thickness of 30 µm with silicon carbide papers, and reduced with argon ions using a Gatan 691 precision ion polishing system.
2.2
Electrochemical measurement
After each side was ground, the Al samples were welded with a copper wire to prepare the test samples with an exposed area of 1 cm2 as the electrochemical test surface. The electrochemical tests were performed using a macroscopic three-electrode system, in which a pure aluminum anode (PAA) specimen was used as the working electrode, a 1 M mercuric oxide electrode (Hg/HgO) was used as the reference electrode, and a platinum sheet with the size of 20 mm × 20 mm × 0.3 mm was used as the counter electrode. Then, 1 M sodium hydroxide solution was utilized as the electrolyte. The electrochemical measurements included open circuit potential (OCP) and polarization curve measurements. The Autolab PGSTAT302N electrochemical workstation was connected to the three-electrode system, and the sample was immersed in 1 M sodium hydroxide electrolyte for 60 min. Then, the OCP was measured. After the OCP was stabilized, the polarization curve test was performed from −2.0 to −0.5 V vs. Hg/HgO with a scan rate of 1 mV/s. Nova 2.1 was used to perform the tests and fit the results.
2.3
Material characterization
The original material was cut into pieces, and Al was smelted by the GZ-0.01 vacuum induction melting furnace (40 kW, 4000 Hz) with a high-purity graphite crucible and protected by argon gas. When the temperature reaches approximately 800°C, casting is conducted, and the pure Al ingot used in the experiment is obtained after cooling to room temperature. First, the samples were polished to a mirror finish. Then, an area was selected using the optical microscope of the hardness tester and calibrated. The sample was immersed in 1 M NaOH at different times, and the corrosion process of the secondary phase and grain boundary (GB) in solution was observed in situ using the Regulus8100 cold field emission scanning electron microscope. The sample composition was analyzed based on the energy spectrum. The parameters were as follows: the accelerating voltage was 20 kV, the working distance was 15 mm, the secondary electron mode was adopted, and the beam spot diameter was 4.5 nm. Scanning kelvin probe force microscopy (SKPFM) was used to measure the surface morphology and surface potential change of the secondary phase in the sample. The SKPFM-AM mode in the atomic force microscope (MultiMode 8) was used to measure the sample surface potential with a silicon nitride probe, a scanning speed of 0.5 Hz, and a scanning interval range of 1–50 µm. During the entire scanning process, the distance between the probe and the sample surface was (100 ± 3) µm, and height calibration was conducted before scanning. The electron backscatter diffraction (EBSD) technique was used to determine the crystal orientation of pure Al samples. EBSD data were collected using the Oxford AZtec device connected to the GeminiSEM 500 field emission scanning electron microscope. During data acquisition, the EBSD operating parameters were as follows: acceleration voltage of 20 kV, sample stage tilt angle of 70°, spot size of 3.0, exposure of 3.2, working distance of 10 mm, and the scan step and test time determined according to the sample state. To obtain high-quality EBSD photographs, the samples were electrochemically polished in a solution of perchloric acid/anhydrous ethanol with the volume ratio of 10:90 at a constant voltage of 25 V for 10–15 s after the mechanical polishing process, as described previously, to obtain samples with a mirror surface for data acquisition. EBSD data were analyzed and processed using the Oxford Channel 5 software. The morphology, structure, and composition distribution of the second phase in the samples were characterized and analyzed by a transmission electron microscope (JEM-2010). The electron excitation power is 200 keV. The ESCALAB 250 X-ray photoelectron spectrometer was used to characterize the surface composition and elemental valence states of the samples. The Mg Kα X light source is adopted, and the passing energy is 30 eV. The focused ion beam (FIB) function of the Tescans9000x high-resolution and high-sensitivity surface analysis system was used to analyze the overall morphology of the second phase after immersion at different times. The ion source is Xe plasma, and the acceleration voltage is 5 kV.
2.4
Finite element simulation
The localized corrosion was simulated based on the finite element method using the COMSOL Multiphysics 5.2a software. In detail, the secondary current distribution interface is employed to solve the electrolyte potential over the electrolyte domain according to Ohm’s law. The phase-field interface was used to track the dissolution of the Al phase. At the phase-field interface, the two-phase flow dynamics is governed by the Cahn–Hilliard equation. The multiphase geometric parameters used in the model construction were based on the experimental test results.
3.
Results and discussion
3.1
Characterization of microstructure
Fig. 1(a)–(d) shows the low-magnification field emission scanning electron microscopy (FESEM) images to characterize the original morphologies of four Al samples with different grades. Notably, SPPs with various shapes and sizes are randomly distributed on the surface of two low-purity Al (i.e., 99.70% and 99.85%). By contrast, fewer SPPs with much smaller sizes can be detected on the surface of 99.95% Al, and SPPs are hardly observed on the 99.99% Al sample, which has a smooth surface with a uniform contrast. To further characterize the SPPs in pure Al, high-magnification SEM observations were conducted. Notably, SPPs mainly appear in two different forms, i.e., spheres and lines, which both consist of numerous small SPP particles. The former SPPs are usually precipitated within the Al grains, and the obtained spheres are usually precipitated within the minimum surface energy of SPPs. The latter SPPs are deposited along the GBs to form a discontinuous line. To further investigate the chemical composition of SPPs, energy-dispersive X-ray spectroscopy (EDX) mapping was employed, as shown in Fig. 1(e)–(f). The results indicate that the two types of SPPs are both composed of Al and Fe, similar to those reported in previous work [30].
Fig.
1.
Original morphologies of four aluminum (Al) samples with different grades: (a) 99.70%, (b) 99.85%, (c) 99.95%, and (d) 99.99%; elemental distribution of secondary phase precipitates (SPPs) (e) within the grain and (f) at the grain boundary (GB).
Dark-field transmission electron microscopy (TEM) was also adopted to characterize the detailed microstructures of SPPs. As illustrated in Fig. 2(a), the large spherical cluster is the assembly of many SPP particles with diverse shapes and sizes ranging from microns to nanometers. The high-magnification EDX mapping at region A of Fig. 2(a) also indicates that the SPP particles are composed of Al and Fe. Notably, Fe and Al almost show reverse intensity distribution. The higher atomic number of Fe leads to its stronger intensity than Al in SPPs, and the surrounding region is identified to be the Al substrate, where only the Al signal is reflected in the absence of Fe. The crystalline structure relationship between the SPPs and the Al matrix is further revealed by selected area electron diffraction (SAED), as displayed in Fig. 2(c). The SAED patterns generated from region B of Fig. 2(a) include two sets of coherent crystals, i.e., pure Al and Al9.75Fe3. The diffraction patterns of Al and Al9.75Fe3 are indexed to be along the zone axis of [101] and[ˉ120], respectively, where the lattice planes(ˉ111), (020), and (11ˉ1) of Al are correspondingly coherent to the (422), (006), and (ˉ4ˉ24) planes of Al9.75Fe3. The well-overlapped diffraction patterns indicate that lattice deformation is barely produced at the interface between the Al9.75Fe3 precipitates and the Al substrate within the spherical SPP, which can be further proven by the kernel average misorientation mapping in Fig. 2(e), where no clear local misorientation is detected inside the SPP sphere. Nonetheless, a relatively high local misorientation was measured at the interface between the spherical SPP cluster and the Al matrix, indicating the existence of stronger plastic deformation in this region. This high local misorientation could induce the alteration of the locally electrochemical homogeneity and cathode/anode area ratio, thus decreasing the stability of the lattice distortion region and facilitating the initiation of corrosion when exposed to aggressive media [31–33].
Fig.
2.
(a) Dark-field transmission electron microscopy image of a spherical SPP cluster; (b) enlarged image of the square region in (a); (c) and (d) elemental mappings corresponding (b); (e) selected area electron diffraction (SAED) of the circular region in (a); (f) SEM image of a spherical SPP cluster; (g) Kernel average misorientation map of the cluster in (f) (the inserted image shows misorientation distribution).
Subsequently, the localized topography and electrochemistry surrounding the SPP clusters were measured by atomic force microscopy coupled with SKPFM mode. Fig. 3(a) and (b) shows the topography and Volta potential mapping of a typical spherical SPP cluster within the grain, respectively. Notably, the SPP cluster exhibits high values both in height and Volta potential, which can be further confirmed by the line scan profiles shown in Fig. 3(c). Because of the higher hardness value of Al9.75Fe3 precipitates than that of Al, a topographical hump with an average height of approximately 70 nm is left after the mechanical polishing process. A more positive Volta potential (approximately 400 mV) indicates that Al9.75Fe3 serves as the cathodic phase with higher chemical stability than the Al substrate, which could exacerbate the localized corrosion tendency when subjected to a corrosive atmosphere. In comparison, the SPP cluster deposited at the GB region exhibits a similar altitude and Volta potential difference to the substrate, indicating its identical nature to the spherical SPP cluster.
Fig.
3.
Atomic force microscopy topographies (a, d), scanning kelvin probe force microscopy maps (b, e), and corresponding line scan profiles (c, f) of the SPPs grown (a–c) within the grain and (d–f) at the GB.
The corrosion behavior of Al anodes with four grades of purity was initially evaluated by OCP measurement and potentiodynamic polarization tests. The OCP vs. time profiles of Al samples in 1 M NaOH solution are illustrated in Fig. 4(a). Notably, all of the OCPs experience a rapid drop followed by a gradual increase before reaching a stable plateau till approximately 300 s. Moreover, the increase in the purity of the Al anode could result in a lower OCP value, which is mainly determined by the impurity element level of the Al samples. Fe is well-known to be the primary impurity element in Al, which contributes to the formation of Al–Fe impurity particles, such as Al9.75Fe3 in Figs. 1–3. Because the standard electrode potential of Fe is more positive (Fe/Fe3+ = −0.44 vs. NHE, where NHE is normal hydrogen electrode) than that of Al (Al/Al3+ = −1.66 vs. NHE), the measured mixing potential of Al anodes with lower purity is consequently higher than that of Al anodes with higher purity. The potentiodynamic polarization curves of the Al samples with different purity grades are shown in Fig. 4(b), and the relevant fitting results are listed in Table 2. Different Al samples exhibit similar electrochemical behaviors, but the corrosion potential (Ecorr) and corrosion current density (icorr) between low-purity Al (99.70% and 99.85%) and high-purity Al (99.95% and 99.99%) are quite different. The icorr value of low-purity Al is larger than that of high-purity Al, indicating that the low-purity Al has a larger corrosion rate. Thus, the purity of Al significantly affects the corrosion resistance of Al anode, and the lack of Fe-rich precipitates in high-purity Al improves its corrosion resistance.
Fig.
4.
Open circuit potential profiles (a) and potentiodynamic polarization curves (b) of four Al samples with different purity grades in 1 M NaOH.
The initial corrosion behavior of Al with four grades of purity in 1 M NaOH solution was investigated by immersion tests for 5 min. According to the FESEM images shown in Fig. 5(a), the Al specimen with low purity (99.70%) appears to suffer from severely SPP-induced localized corrosion, where a large number of micropits with different depths are produced after the immersion test. Notably, most of the pits are wide and shallow with micron size, but some pits are narrow and deep, as shown in the inset of Fig. 5(a). According to the size and shape of the pits, the formation of the pits can be reasonably deduced to be associated with the separation of SPP clusters from the substrate, which can be further supported by the trench left along the GBs (marked by a yellow dashed circle). Compared with the 99.70% specimen, the sample with slightly higher purity (99.85%) shows a similar corrosion morphology, except that the pits are more shallow and smaller. Moreover, the SPP cluster that appears along the GBs becomes more convex compared with the Al substrate, which can be attributed to the higher dissolution rate of Al than Al9.75Fe3 precipitates, as demonstrated by the measured Volta potential shown in Fig. 3. As the Al purity increases, the SPP-induced corrosion tends to be milder. The surface of the 99.99% specimen remains smooth, except for the existence of a few tiny pits. Based on these observations, we can conclude that Al with a higher purity could effectively reduce the localized corrosion caused by SPPs.
Fig.
5.
Top view FESEM images of four Al samples after immersion in 1 M NaOH for 5 min: (a) 99.70%, (b) 99.85%, (c) 99.95%, and (d) 99.99%. Enlarged images of corrosion pits are inserted in (a) and (b).
To determine the mechanism of SPP-induced localized corrosion in strong alkaline media, quasi-in-situ SEM observation was conducted to monitor one typical spherical SPP cluster during the immersion test. Fig. 6(a) shows the original morphology of the SPP cluster on a 99.85% Al specimen. Notably, the microstructure of the SPP cluster is identical to those shown in Figs. 1 and 2, where numerous particles and lamellas are interlaced together to form a sphere in Al. After being immersed for 30 s, the edges of SPPs are not as sharp as the original ones because of the electrochemical reactions between the SPPs and the alkaline electrolyte. Furthermore, the clear scratches on the Al substrate indicate that the specimen only experienced a weak corrosion attack. With the prolongation of immersion up to 60 s, the surface of the SPP cluster is covered by a layer of corrosion product, making it difficult to distinguish the microstructure within the SPP cluster. Moreover, the sharp white outline of the sphere indicates a strong underfocus (Fig. 6(c)), which shows that the distance between the edge of the spherical SPP cluster and the Al substrate is larger than the initial status, illustrating that the dissolution rate of Al is faster than that of SPPs during the first 60 s. When the immersion time increases to 5 min, the surface layer on top of the SPP cluster becomes thicker so that its internal structure can hardly be recognized. Moreover, the scratches on the Al substrate disappear because of its rapid consumption in strong alkaline media. With the further decrease in the Al surface level, the contact area between the SPP cluster and the substrate promptly diminishes, leading to a dramatic weakening of the binding force between them until the SPP cluster floats in the solution. As shown in Fig. 6(f), a wide and shallow pit is left on the surface when the SPP cluster drifts away after 10 min. This phenomenon explains the existence of numerous pits on the surface of Al with low purity after being immersed in the alkaline solution.
Fig.
6.
Quasi-in-situ FESEM observation on one typical SPP cluster in 1 M NaOH up to 10 min: (a) original; (b) 30 s; (c) 60 s; (d) 2 min; (e) 5 min; (f) 10 min.
To obtain more information about the internal structure of the SPP cluster during the corrosion process, the FIB–SEM technique was utilized to characterize the cross-section morphology of the spherical SPP cluster after being immersed for 2 min. The top view SEM image in Fig. 7(a) shows that the cluster has a porous structure, and the surface of SPPs is decorated with many product nanoparticles. Notably, Al inserted into the SPP sphere has been removed, exposing the top part of SPPs with a higher level than the Al substrate. Fig. 7(b) illustrates the cross-section view of the SPP cluster embedded in the Al substrate. The embedded part of the SPP cluster is easily distinguished to have a semicircle shape, which is intertwined with a certain amount of SPPs. Different from the corrosion morphology caused by solid precipitate particles [27,29,34], the porous SPPs are compactly affiliated to the Al substrate without any crack or trench, which could be interpreted as the coherent lattice relationship between them as characterized by high-resolution transmission electron microscopy. Moreover, no crevices or pits are produced at the interface between the cluster and the surrounding substrate, indicating that galvanic corrosion is not dominant during the intense dissolution of Al in dense alkaline electrolytes. The elemental distribution is illustrated in the yellow rectangular region shown in Fig. 7(b), and the mapping results are shown in Figs. 7(c)–(e). As expected, the signal of Al can be fully detected in the region because the base and the SPP cluster are both mainly composed of Al. The intensity contrast can be identified because of the Al content difference between the Al9.75Fe3 SPPs and the substrate. The EDX mapping of Fe is denoted by the semicircle outline of SPPs, which is highly consistent with the SEM image shown in Fig. 7(b). Furthermore, the surface region is enriched with oxygen, which can be ascribed to the formation of corrosion products from complex anodic reactions [15].
Fig.
7.
Top view (a) and cross-section (b) FESEM images of one typical spherical SPP cluster after being immersed for 2 min; (c–e) elemental distribution of the SPP cluster after immersion in 1 M NaOH for 2 min.
The SPP-induced localized corrosion was simulated by establishing a finite element model at the cross-section interface, including the NaOH electrolyte, Al substrate, and Al–Fe SPP cluster. The secondary current distribution interface is adopted to solve electrolyte potential over the electrolyte domain according to Ohm’s law. The relationship between the local current density and the electrolyte potential is incorporated into the model using a piecewise cubic interpolation function for the experimental polarization data. Combined with the determined relationship and the interpolation function ensuring that the local current density is only applied to different phases, the local current density for the Al and Al–Fe phases at the electrode surface can be set. The finite element simulation assumed that the anodic dissolution reaction occurs at the Al phase surface, and the cathodic hydrogen evolution reaction (which is not associated with any loss of material) occurs at the Al–Fe phase surface. Furthermore, the dissolution of the Al phase is violent, whereas that of the Al–Fe phase is less violent. When solved using the COMSOL Multiphysics 5.2a software, the phase-field interface is utilized to analyze the dissolution of the Al phase, and the two-phase flow dynamics is governed by the Cahn–Hilliard equation for tracking the diffuse interface separating the immiscible phases. Fig. 8 illustrates the evolution of current density distribution at the interface region as the corrosion proceeds. Notably, only a small current density difference is produced at the initial stage. As shown in Fig. 8(b), the current densities at the tri-phase interfaces (i.e., Al, Al9.75Fe3, and liquid phases) are slightly higher than those at the other regions, where the primary microbatteries are formed. With the occurrence of localized corrosion, the current density distribution becomes nonuniform because of the complex hierarchical microstructure of the SPP cluster, leading to the fierce dissolution of the small Al particles that are surrounded by the cathodic Al–Fe precipitates with large sizes, as indicated by the two separate dark red regions in Fig. 8(c) and (d). With the further decrease of the surface level, the exposed area of the tri-phase interface sharply increases, and more local regions with higher current densities are formed within the SPP frame. Notably, the current densities of the substrate near the SPP are not homogeneously distributed either. As depicted in Fig. 8(d)–(f), local current densities at several points of the Al substrate start to surge, and these small nuclei would integrate as a large spot with severe corrosion. The phenomenon is probably caused by the small humps on the coarse surface of the Al substrate after a serious corrosion attack. This finding is consistent with that shown in the SEM images, where numerous small pits around the SPPs are observed. Furthermore, Fig. 8(f) illustrates that the current density of the Al substrate gradually decreases with distance to the SPP, which can be attributed to the weakened galvanic effect in remote regions.
Fig.
8.
Simulation of the localized corrosion at the region near one typical SPP cluster: (a) original; (b, c) initial period of corrosion; (d, e) development of corrosion; (f) later period of corrosion.
To further understand the composition of the corrosion products on the surface of Al, X-ray photoelectron spectroscopy analysis was employed to characterize the chemical and electronic states of the Al samples after being immersed in the alkaline solution for 5 min. The high-resolution Al 2p spectrum can be deconvoluted into two distinct peaks positioned at 74.8 and 72.5 eV (Fig. 9(a)) corresponding to Al3+ and Al0, respectively [35]. The strong intensity of peak Al0 from the substrate indicates that the thickness of the surface film is insufficient to block the penetration of X-ray photoelectron irradiation. Moreover, the existence of the Al3+ peak indicates the formation of aluminum oxide species (i.e., Al2O3 and Al(OH)3). This finding can be further confirmed by the M–O–H peak of O 1s shown in Fig. 9(b), where the peak assigned to metal oxides overlaps the peak assigned to hydroxides [36]. Fig. 9(c) depicts the high-resolution spectrum of Fe 2p. The noisy profile indicates the low Fe content in the surface layer because it is merely contributed by the signal of oxidized SPPs. Although the intensity is low, it can still be deconvoluted into several peaks corresponding to Fe3+, Fe2+, and satellites, which demonstrates the oxidation of Fe in SPPs [37–38]. Except for the overall oxidation information previously presented, the compositional evolution of the Al–Fe precipitates was investigated based on the EDX measurements on dozens of the SPP clusters after being immersed at different times. According to the results shown in Fig. 9(d), the ratio of Al experiences a dramatic linear drop in the first 5 min, which can be attributed to the rapid dissolution of Al within the SPP cluster frame. In comparison, the O and Fe contents both increase because of the oxidation of SPPs. Further prolongation of immersion leads to gradually stable ratios of Al and Fe, probably resulting from the growth of oxide layers on SPPs to slow down the dissolution and diffusion rates of the metals, which can be further certified by the gradual increase in O content.
Fig.
9.
X-ray photoelectron spectroscopy analysis results of the 99.70% Al after immersion in 1 M NaOH for 5 min: (a) Al, (b) O, and (c) Fe; (d) atomic ratios of Al, Fe, and O in SPPs during immersion.
The schematic of the effect of SPPs on the initial corrosion behavior of pure Al in alkaline media is illustrated in Fig. 10. Initially, the porous Al9.75Fe3 SPP is embedded in the Al substrate. With the increase in the immersion time inNaOH solution, Al is rapidly dissolved in the electrolyte, and the level of the Al surface significantly decreases, leading to the exposure of SPP to the electrolyte. Meanwhile, the surface of SPP is also covered by a layer of corrosion products (i.e., the purple layer in Fig. 10(b)) because of the reaction between the Al9.75Fe3 SPP and the dense OH− ions. Because of the edge effect, the corrosion rate around the SPP is slightly higher than the general Al substrate. With the further prolongation of the immersion time, the horizontal plane of the Al matrix continues to drop, leading to further exposure of the SPP. Finally, the continuous dissolution of the Al matrix facilitates the formation of a weak contact between the SPP and the Al matrix, resulting in the SPP drifting from the substrate and leaving a pit on the surface, as shown in Fig. 10(d). Electrochemical reactions associated with the corrosion process are expressed as follows.
Fig.
10.
Proposed schematic of SPP-induced localized corrosion in alkaline media.
The anodic and cathodic reactions are shown in Eqs. (1) and (2), respectively:
Al→Al3++3e−
(1)
2H++2e−→H2
(2)
Then, Al3+ reacted with OH− in the solution to form metallic hydroxides:
Al3++3OH−→Al(OH)3
(3)
Further decomposition of unstable Al(OH)3 leads to the formation of Al2O3:
2Al(OH)3→Al2O3+3H2O
(4)
4.
Conclusion
In this work, the effect of porous SPPs on the initial corrosion behavior of pure Al in alkaline media was comprehensively investigated based on experiments coupled with simulations. The results indicate that Al–Fe SPPs in pure Al are hardly dissolved in an alkaline solution, thus serving as the cathodic phase during the corrosion process, leading to galvanic corrosion around them. The low-grade Al samples suffer from severe corrosion because of the intense galvanic corrosion caused by the large SPP clusters both within the grains and at the GBs. By contrast, this phenomenon is weaker in highly pure Al samples because of the smaller size and lower content of SPPs. Quasi-in-situ SEM observations show that the rapid dissolution of the Al substrate further contributes to the exposure of the SPP clusters. Furthermore, the simulated current density distributions based on the finite element model indicate that the SPPs can accelerate the corrosion rates inside the SPP cluster and the substrate near the SPP. The findings of this work can serve as references for the design of anode materials for Al–air batteries.
Acknowledgements
This work was financially supported by the National Natural Science Foundation of China (No. 51901018), the Young Elite Scientists Sponsorship Program by the China Association for Science and Technology (YESS, No. 2019QNRC001), the Fundamental Research Funds for the Central Universities, China (No. FRF-AT-20-07, 06500119), the Natural Science Foundation of Beijing Municipality, China (No. 2212037), the National Science and Technology Resources Investigation Program of China (No. 2019FY101400), and the Southwest Institute of Technology and Engineering Cooperation Fund, China (No. HDHDW5902020107).
Conflict of Interest
The authors declare that they have no conflict of interest.
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